The opportunities, successes and challenges for GaInNAsSb

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Journal of Crystal Growth 278 (2005) 3–17

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The opportunities, successes and challenges for GaInNAsSb

James S. Harris Jr.�

Solid State and Photonics Laboratories, Stanford University, CIS-X 328, Via Ortega, Stanford, CA 94305-4075, USA

Available online 2 March 2005

Abstract

Dilute nitride GaInNAs and GaInNAsSb alloys grown on GaAs have quickly become excellent candidates for a

variety of lower cost 1.2–1.6 mm lasers, optical amplifiers and high power Raman pump lasers that will be required for

advanced high-speed communications systems. Two particularly critical devices are (1) vertical cavity surface emitting

lasers (VCSELs) which must operate at high data rates (X10Gbps), uncooled over a broad thermal operating range

and (2) high-power (X500mW) edge emitting lasers for Raman amplifier pumps. While progress has been both rapid

and very promising, there remain some significant challenges. The major materials challenges include the limited

solubility of N, a 2-D to 3-D growth transition with phase segregation, precise control over composition, strain and

energy band offsets at heterointerfaces and elimination of non-radiative defects that are caused by some combination of

N incorporation, low growth temperature and ion damage from the N plasma source. The addition of Sb significantly

improves the epitaxial growth, maintaining 2-D growth over a broader range of growth temperature, larger In and N

compositions, V/III ratio and significantly improves the optical properties at wavelengths longer than 1.3 mm. Thispaper describes progress in overcoming some of the material challenges, particularly low temperature GaInNAsSb

growth, plasma source ion damage in solid source MBE and progress in realizing record setting long wavelength edge

emitting lasers and the first monolithic VCSELs operating at 1.5 mm based on GaInNAsSb QWs grown on GaAs.

r 2005 Elsevier B.V. All rights reserved.

PACS: 78.55.Cr; 78.66.Fd; 78.60.Fi

Keywords: A3. Heteroepitaxy; A3. Molecular beam epitaxy; B1. Dilute nitrides; B2. Semiconducting III–V materials; B3.

Heterojunction semiconducting devices; B3. Laser diodes

1. Introduction

In spite of the incredibly rapid growth andoverall capacity of current optical communicationsnetworks and the Internet, it should come as little

e front matter r 2005 Elsevier B.V. All rights reserve

ysgro.2004.12.050

723 4659; fax: +1 650 723 659.

ss: [email protected] (J.S. Harris Jr.).

surprise that speed on the Internet is now limitedby the same problems of all mature, high speedtransportation systems—airlines, trains and auto-mobiles—the switching hub from the main highspeed/capacity backbone through traffic jams inthe on/off ramps and slow feeder lines to the finaldestination. This bottleneck is now the focus ofconsiderable effort to create truly high-speed

d.

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J.S. Harris Jr. / Journal of Crystal Growth 278 (2005) 3–174

metro (MAN) and local area networks (LAN) todeliver high-speed to the desktop [1,2]. This hasbeen the driving force behind the development oflow cost, 1.3 mm, directly modulated, un-cooledvertical cavity surface emitting lasers (VCSELs)because in contrast to the backbone, where thereare only tens of thousands of lasers and perfor-

mance is the critical issue, high-speed direct accesswill require hundreds of million lasers and cost isthe issue. Compared to current generation lasers,the cost will have to be reduced by at least 100� .While this might seem a very daunting challenge,first generation LANs based on low cost 850 nmGaAs VCSELs have demonstrated that they canmeet the cost, reliability, speed and thermalrequirements for such lasers, the only problembeing that they operate at the wrong wavelengthand 10Gbps transmission is only possible overabout 50m (m NOT km)! [1–3]. Thus, longerwavelength lasers are absolutely essential toachieve useful transmission distances at this datarate. For long-haul applications, 1.55 mm In-GaAsP/InP-based Bragg grating and distributedfeedback (DFB) lasers are commonly used, butthey are prohibitively expensive for high volumeMAN and LAN systems. Despite sizeable effortsto realize low cost, long wavelength VCSELs,limitations in the current alloy systems have madethis an extremely difficult challenge.In the past few years, a new materials system,

GaInNAs and now, GaInNAsSb have beeninvestigated which have the potential to not onlymake long wavelength, low cost VCSELs a reality,but to provide the foundation technology for abroad range of devices to power a truly revolu-tionary broadband network, including; low cost,long wavelength VCSELs, high power pumplasers, semiconductor optical amplifiers and newfunctionality/lower cost ICs by integrating photo-nic crystal structures to produce a new generationof photonic ICs. GaInNAs(Sb) can be grownnearly lattice matched to GaAs and has a bandgapin the telecommunication wavelength range [1–4].Recently, commercial 1.28 mm VCSELs have beendescribed [5]. Being grown on GaAs gives thissystem the distinct advantage of being compatiblewith monolithic growth of GaAs/AlAs distributedBragg reflecting (DBR) mirrors, a significantly

better developed and lower cost fabricationtechnology and the advantages of the uniqueproperty of AlAs oxidation. Other advantagescompared to InP-based solutions include betterthermal stability due to larger conduction bandoffsets between barrier and quantum well (QW)and simplified compositional control over As/Pmixed group V solutions. While growth ofGaInNAs QWs allows fabrication of near 1.3 mmVCSELs, manufacturers are hedging on goingclear out to 1.3 mm or beyond because they havehad yield problems which are similar to ourobserved growth sensitivity for these QW alloysand their reproducibility remains somewhat pro-blematical. We have found that by incorporatingSb to form a quinary alloy, GaInNAsSb, thegrowth window for good 2-D epitaxy is not onlyexpanded, but has enabled the incorporation ofhigher In and N compositions and realization of1.5 mm low threshold current, high power edgeemitting lasers [6–8] and the first monolithic1.5 mm VCSELs grown on GaAs [8,9]. The majorappeal of this quinary alloy is thus a SINGLEmaterials system for growth and processing whichcan produce not only the entire range of tele-communications devices, but enable significantintegration of photonic and electronic devices withphotonic crystal waveguides and resonators whichcould be the foundation technology for truly lowcost, high speed, highly functional photonicintegrated circuits that will provide high band-width networks to the desktop [1,2]. The keyquestion, ‘‘is this quinary alloy a prototype systemthat provides the ultimate opportunity for band-gap engineering and novel devices, or is it a recipefor a disaster of poorly controlled and non-reproducible epitaxial growth?’’

2. Molecular beam epitaxial growth

Due to the difficulties of incorporating sufficientN in solid solution, growth of GaInNAs andGaInNAsSb are substantially different than allearlier III–V systems [1,2]. Differences in crystalstructure between GaInN and GaInAs, the smallsize of N compared with As and the highelectronegativity of N cause a large miscibility

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Fig. 1. TEM cross-section image of 3-QW 35% In GaInNAs–

GaAs barrier sample where 3-D growth and phase segregation

have occurred in the third QW.

J.S. Harris Jr. / Journal of Crystal Growth 278 (2005) 3–17 5

gap between GaInAs and GaInN. In order to growhomogeneous alloys of GaInNAs or GaInNAsSbwhich will reach the desired wavelengths, growthmust take place under very low temperaturemetastable conditions. GaInNAs and GaInNAsSbhave been grown by solid source molecular beamepitaxy (MBE) with a reactive N RF plasmasource. Plasma sources put out a variety of species(N2

+, N+, N as well as electrons and each of themolecular species in excited states) and exactlywhich species are incorporated and the potentialfor ion and/or electron damage are still poorlyknown [9,10]. We have examined a number ofthese issues and find that the source operatingconditions and substrate temperature are the mostcritical parameters for metastable growth of highquality materials. We use low temperatures,around 425 1C, in order to inhibit phase segrega-tion and improve surface morphology and a biasvoltage applied to the deflection plates of ourplasma source to minimize ion damage. DimericAs and monomeric Sb are supplied using thermalcracking cells. Arsenic overpressure (BEP) wasgenerally fixed at 15–20 times the group III fluxduring growth, unless otherwise noted. A veryimportant property of nitride growth by MBE isthe fact that N incorporates with nearly unitysticking conditions, i.e. the group III growth rateand not the V–V ratio control the N incorporation[11].One of the major challenges for growth of

GaInNAs is avoiding phase segregation andmaintaining good 2-D epitaxial growth for materi-als operating at or beyond 1.3 mm as illustrated inFig. 1. This illustrates a common observation thatthe first or second QWs grow without phasesegregation, but by the third QW, the top interfaceclearly has quantum dot like regions and phasesegregation. This is also immediately visible withRHEED during growth. TEM characterization ofthese regions indicate that it is the result of Inmigration and a local increase in In composition[12].Growth of GaNAs and GaInNAs both suffer to

some degree with the problem illustrated in Fig. 1.Even though GaNAs clearly has no In, Nsegregation or strain induced 3-D islanding mustbe occurring. Surface surfactant modulated

growth is known to aid in control of surfacekinetics in MBE growth [13] and Shimizu et al. [14]and Yang et al. [15] discovered that Sb could beused as a surfactant in InGaAs and GaInNAsgrowth, respectively. When we introduced Sb intothe growth of GaInNAs, we observed a significantenhancement of PL, particularly at wavelengthsgreater than 1.3 mm as illustrated in Fig. 2 [1,2,6].However, we quickly discovered that Sb acts asboth a surfactant and constituent when used inGaInNAs, forming GaInNAsSb [6,15]. Untilrecently, our devices utilizing GaInNAsSb as theQW material had GaNAsSb as the barriermaterial [7,8,17]. Sb was incorporated into thebarriers because it was thought to be a surfacesurfactant that would remain on the surface duringQW and barrier growth, even if the Sb shutter wasclosed. It also seemed that it might improve thequality of GaNAs and the interfaces between theQWs and barriers. Although GaInNAsSb hasbeen extensively studied as the QW material, therehas been very little work reported on GaNAsSb[18,19].Because of the rather extensive recent reviews of

GaInNAs growth and lasers [1,2], I will concen-trate on recent studies on the effect of growthconditions, including the N RF plasma source and

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0.75 0.80 0.85 0.90 0.95 1.00

1600 1500 1400 1300 1200

35% In GaInNAs

31% In GaInNAs

38% In Sb: 7.2e-8 torr GaInNAsSb



Energy(eV)

Wavelength (nm)

100

10-1

10-2

10-3

Inte

nsity

(a.

u.)

Fig. 2. PL of GaInNAs(Sb) comparing the best 1.3 mmmaterial

grown without Sb and the dramatic improvement in PL at even

longer wavelengths by adding Sb.

0 25 50 75 100 125 150

0.00

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0.09

0.10

0.11

Nitrogen

Antimony

Mol

e F

ract

ion

Depth (nm)

Fig. 3. SIMS results of N and Sb for GaNAs and GaNAsSb as

a function of substrate temperature.


growth of GaInNAsSb which have been key toproducing record long wavelength lasers. Theimprovements in materials growth have focusedon (1) growth temperature and the properties ofGaNAs vs. GaNAsSb as barriers for GaInNAsSbQWs for lasers operating in the 1.3–1.55 mm rangeand (2) the N RF plasma source, species itproduces and possible ion damage created duringgrowth of the epitaxial layers.One of the key issues is the precise behavior of

Sb in the growth. It is clear that it serves as asurfactant to modify the surface kinetics andsegregation of In in the QWs, but depending upongrowth conditions, it may be incorporated into thebarrier and cladding regions as well. Such incor-poration dramatically effects both the compositionand strain in the region of the QW, which in turnstrongly effects the bandgap and offsets betweenbarrier and QW. We have studied the GaNAsSbmaterial used in barriers of our previous 1.3 and1.5 mm laser devices [6,16,20].The GaNAs(Sb) SQWs were grown at 0.45 mm/h

(to duplicate the 1.3 mm QW barriers) or 0.30 mm/h(to duplicate the 1.5 mm QW barriers). An As-to-Ga beam equivalent pressure (BEP) overpres-sure of 20� was used and an Sb flux of0.8–1.0� 10�7 Torr BEP was supplied during theQW growth.

RHEED patterns were monitored duringgrowth. GaNAs produces a streaky 2� 4 surfacereconstruction which became spotty almost in-stantly when the Sb shutter was opened, suggestingthe surface quality was not as good as GaNAs.This result was somewhat surprising since previousstudies have shown Sb had the opposite effect withInGaAs and GaInNAs [15,19].For a fixed N flux, varying the growth rate alters

the amount of N incorporated [11]. However, for afixed Sb-to-As ratio, variations in the growth ratedo not change the Sb composition. Sb incorpora-tion appears to be independent of the group-IIIgrowth rate and N incorporation, suggesting theflux ratio of Sb-to-As is the controlling factor.A SIMS depth profile for a GaN0.029-

As0.873Sb0.098/GaAs QW sample is shown inFig. 3. An interesting feature to note in the SIMSdepth profile is the interface regions of theGaNAsSb layer. The N and Sb profiles do notstart and end at the same location within thesample. This was determined to be real and not ameasurement or sputtering artifact, as it wasrepeatable within the same sample and observedon all other samples measured with SIMS. It isclear that Sb first builds up on the surface beforeappreciable incorporation and continues to incor-porate �5–7 nm beyond the end of N incorpora-tion. Surfactants tend to float on the growth frontand do not incorporate into the material. As

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mentioned in earlier papers relating to GaIn-NAs(Sb), Sb appears to act as a surfactant as wellas a group-V component [5,19]. If this is the case,there will be Sb remaining on the growth frontafter the shutter is closed to the cell and theresidual Sb continues to incorporate or desorbsuntil the supply is exhausted. This growth com-plication could be quite detrimental to devices,since there is then a thin layer of GaAsSb on top ofthe QW which could significantly change theoriginally intended band structure properties dueto changes of both composition and strain.One of the major questions for quaternary and

quinary alloys with three different column Vconstituents is their sensitivity to temperatureand flux as this could make compositional controlvery challenging. In an attempt to study generalgrowth properties of GaNAsSb and potentiallyimprove PL intensity and material quality, a seriesof samples with varying As-to-Ga overpressuresand growth temperature were grown with the samestructure as the previous samples. It is known thatin mixed group-V materials (AsP or AsSb), therelative fluxes of each group-V element play a largerole in composition and growth kinetics. InGaInNAs, there was no significant effect on Nincorporation by different As fluxes due to the‘‘unity’’ sticking properties of N [11]. However inGaNAsSb, it is suspected that the As and Sb fluxesdo indeed affect each other since they do not havethe same sticking properties as N. SIMS measure-ments showed that as the As overpressure in-creased from 15� to 30� , the Sb concentrationdropped from 12% to 9% while the N concentra-tion remained virtually unchanged [20,21]. Theincrease in As flux thus has a direct effect on theSb incorporation rate but no discernable effect onN incorporation (as seen in GaNAs) [21,22]. Inaddition, the change in Sb concentration had noeffect on the N incorporation in agreement withpreviously obtained results [18,21]. The measure-ments also show enhanced N incorporation in theGaNAsSb. GaNAs grown under the same growthconditions yields 1.8% N, much lower than theobserved 2.4–2.9% in GaNAsSb. The PL inten-sities are within measurement error and samplerepeatability, thus revealing no significant changein optical material quality between the different

samples. Going from 15� to 30� As over-pressure, we have only a 3% Sb decrease, thuschanging the As overpressure does not have anymajor effect on GaNAsSb except for the smallchange in Sb incorporation [22]. The materialquality remains the same structurally and opti-cally. For GaInNAs(Sb) QWs with GaNAsSbbarriers, it would be beneficial to increase the Asoverpressure so that the GaNAsSb would have lesscompressive strain and a larger bandgap forincreased confinement in the wells.The second major growth variable is substrate

temperature. The substrate temperature duringGaNAsSb growth was also varied to examine theeffects on crystal quality and composition [22].GaInNAs(Sb) was grown at 425 1C to preventphase segregation and relaxation. One of thedriving factors in segregation in GaInNAs(Sb) isthe clustering of In-rich areas. The GaNAsSbbarriers were also grown at the same temperaturesince it was also thought the material wouldsegregate. However, In is not present in thismaterial and thus raised the possibility thematerial could be grown at a higher temperature.One problem with nitride–arsenide growth is thelow substrate temperature. Such low temperaturesintroduce defects in GaAs, such as As antisites andGa vacancies, hence it is desirable to grow thematerial as close to 580 1C as possible to minimizethese defects, which may cause non-radiativerecombination and reduce luminescence. A seriesof samples with structures and growth conditionsidentical those in the initial study were grown withvarying substrate temperatures: +50 1C (475 1C),+100 1C (525 1C), and +150 1C (575 1C) [22].Another set of samples with no Sb (GaNAs) wasalso grown for comparison. HRXRD scans showthat the N composition remained the same, exceptfor the hottest sample in which there was a slightdecrease in N. The most surprising result from theHRXRD scans was the fact that the samples didnot appear to have any significant segregation,relaxation, or interface degradation when grown athigher temperatures. All samples had very welldefined QW peaks and Pendellosung fringes. Fig. 4is a plot of SIMS scans taken of the GaNAsSbsamples to measure composition and depthprofiles. As seen in the figure, there is a large

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400 420 440 460 480 500 520 540 560

2

4

6

8

10

12

Mol

e F

ract

ion

(%)

Substrate Temperature (C)

Sb (GaNAsSb)N (GaNAsSb)N (GaNAs)

Fig. 4. Nitrogen and antimony SIMS profiles of Ga-

N0.029As0.873Sb0.098 sample illustrating both the large incor-

poration and surface accumulation nature of Sb [22].

0 20 40 60

col

25.4

25.2

25.0

24.8

24.60

10

20

30

Fig. 5. Strain map after annealing of a single-QW GaInNAs

sample, without deflection plate bias and constant growth

temperature [23].


decrease in Sb concentration with increasingsubstrate temperature while the N concentrationremained roughly constant. Similar to the SIMSdata from the As overpressure study, the Ncomposition remained constant, even though theSb concentration changed. With this data, onemight be encouraged by the fact that at hightemperatures, the material was coherent and hadsmaller amounts of Sb so that the bandgap waslarger. However, the PL results were very un-expected given the previous results and quicklydestroyed any notion that high temperature wasadvantageous. In addition to the loss of PLintensity, there was also a large shoulder on thelong wavelength side for all three samples whichwas not found in the original substrate tempera-ture sample [22]. This longer wavelength shouldercould be a result of micro-segregation which couldnot be observed in HRXRD. Since there could beareas of clustering of increased N concentrationwithin the QW, it could lead to areas ofluminescence in which the bandgap is smaller,leading to luminescence of longer wavelength. Ifthese regions occurred only inside the QWs, theywould have no effect on the interfaces and thuswould not affect or reduce the diffraction thick-ness oscillations. As evidenced by the poor PLresults, GaNAsSb cannot be grown at high

temperatures without a large decrease in opticalmaterial quality.Further understanding of the challenges for

GaInNAs growth and the role of Sb have beenobtained from examination of samples with the800 keV atomic resolution TEM at the LawrenceBerkeley Lab [23] and AFM QW interfacecharacterization at Phillips University in Marburg,Germany [24] and the TEM samples are polishedto expose the (1 0 0) surface rather than thetraditional cleaving to expose a (1 1 0) surface.The (0 0 2) diffraction condition is then used indark field to provide an image that is very sensitiveto compositional changes in the group III sub-lattice [23–25]. The TEM results confirm theproblems at high In mole fractions, the effect ofSb and that much better control of the growthconditions (particularly substrate temperature)must be maintained as growth under ‘‘nominallyconstant’’ conditions results in changing In com-position in the QWs. Fig. 5 shows a strain mapderived from grid mapping of a high-resolutionTEM image of a single QW (SQW) GaInNAssample. The composition is �30% In and �1.6%N with peak PL at 1.3 mm, which is still in the 2-Dgrowth regime. The strain profile shows very non-uniform In distribution with In strongly segregat-ing near the top surface of the QW. The differencein In composition between lower and uppersurfaces of the well is �6% (i.e. 24% In at thebottom and 30% In at the top). Such a largegradient in In composition would significantlychange the bandgap, strain, QW energy and

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resulting PL linewidth and energy. This is for anunannealed sample, which may be homogenizedsomewhat by the anneal, but this clearly shows theproblem of uniformity and In segregation in theQW which would certainly be best avoided if at allpossible.Fig. 6 shows an atomic resolution TEM strain

map for the top well of a 3-QW GaInNAsSbsample with GaNAsSb barriers. The compositionis �39% In, �1.6% N, �2% Sb in the QWs and�2% N and �6% Sb in the barriers with peak PLat 1.5 mm. Because of the small bandgap of theQW compared to the substrate and radiativesubstrate heating in the MBE, we suspected thatthe QWs were absorbing more IR radiation fromthe heater and the actual growth surface tempera-ture was increasing. For this sample, the substratetemperature was lowered by 6 1C after eachsuccessive QW. Fig. 6 shows that the strain acrossthe QW is very uniform in comparison to theGaInNAs (i.e. no Sb) sample in Fig. 5. Thus the Indistribution across the quantum well is quiteuniform.AFM measurements are often used to charac-

terize the surface roughness of epitaxial films.While these may provide some insight into thefinal grown surface, it is known that the growingsurface has a different surface reconstruction and

Surface

0 20 40 60col

row

19.0

18.8

18.6

18.40

12

25

38

50

Fig. 6. Strain map after annealing of the top well of a three-

QW GaInNAsSb sample, with deflection plate bias applied and

a 6 1C drop in growth temperature for each successive QW

growth [23].

what is most important for QW devices is thequality of the QW interface when the barrier hasbeen grown on it and the growth of the barrier cansignificantly effect this QW interface. A set ofAFM measurements were carried out on similarQW samples to those above to compare the upperand lower interfaces of GaInNAs and GaInNAsSbto observe the effect of Sb during growth. Thesesamples had AlAs barriers grown around theGaInNAs or GaInNAsSb QWs. AlAs can be veryselectively etched from GaAs and GaInNAs(Sb).By etching the AlAs from the interface, theexposed surface is that of the GaInNAs(Sb)interface. Fig. 7 shows a line scan and planarAFM image of the top interface of the thirdGaInNAs QW. The mean surface roughness is1.4 nm and there are no clear single atomic layersteps in the line scan. A similar AFM line scan andplanar image for the top interface of the thirdGaInNAsSb QW is shown in Fig. 8. The meansurface roughness is only 0.75 nm and the line scanexhibits some single atomic layer steps which areindicative of 2-D, layer-by-layer growth. Thisupper surface is the best interface that has beenobserved by this technique, even on samples withmuch lower In or N compositions and illustratesthe important role Sb plays in maintaining the2-D growth regime and preventing micro-phasesegregation [3].Unfortunately, the sample preparation for both

TEM and AFM interface measurements are quitedifficult and time consuming which make thesetechniques unsuitable for routine characterizationof every sample. However, they certainly showthat compositional changes occur across a SQWand also with progression from QW to QW inMQW samples and are thus consistent with thesignificantly broader PL linewidth that one ob-serves in GaInNAs and GaInNAsSb unless thetemperature is quite precisely controlled andpossibly adjusted during growth. These observa-tions may underlie the successful optimization ofGaInNAs QWs based upon PL linewidth [3,7,8,26]rather than PL intensity, which has been theprimary criteria of the past. We are continuing thisinvestigation on samples to find the optimum Sbflux which maintains 2-D growth in the QW, butminimizes other effects on the material.

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Section Analysis

DC Min

Spectrum

0 1.00 2.00–2.00

0

2.00

nm

µm

Fig. 7. AFM image, cross-sectional scan and roughness spectrum of the GaInNAs interior QW interface.


Finally, choice of the barrier material for QWshas both important physical properties withrespect to strain, smoothness, etc., and electronicproperties, such Type-I alignment and sufficientband-offset to localize the electrons and holes inthe QW and prevent leakage out of the QW. If theQW is strained, it would be important to straincompensate with barriers of the opposite signstrain. This will prevent relaxation and theformation of defects by effectively increasing thecritical thickness of the layers. Finally, the materialmust be of good quality or defects present will leadto a higher rate of non-radiative recombination,preventing carriers from radiatively recombiningin the QW. One of the limitations of MBE is theinability to accurately change fluxes duringgrowth. GaNAs was the previous barrier material

utilized in GaInNAs QWs [12,17,22,26,27]. WhenSb was added to GaInNAs, Sb was also added toGaNAs, assuming the expected surfactant natureof Sb and the improvement in material toGaInNAs would also apply to GaNAs. However,the GaNAsSb material has several propertieswhich do not make it the desired barrier materialof choice. Examination of RHEED patternsshowed that the addition of Sb to GaNAs actuallydegraded the surface quality and turns the patternfrom streaky to spotty. This suggests the surfaceroughens and there could be clustering on thesurface which could lead to defects. GaNAsSbdoes not provide any strain compensation to thehighly compressive GaInNAs(Sb) QWs. GaNAs,however, is tensile and provides strain compensa-tion. GaNAsSb barriers also have a bandgap value

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Section Analysis

DC Min

Spectrum

0 1.00 2.00–0.75

0.00

0.75nm

µm

Fig. 8. AFM image, cross-sectional scan and roughness spectrum of the GaInNAsSb interior QW interface.


which is very close to that of the QW. Having asmall band-offset in the valence and conductionbands results in poor hole and electron confine-ment and may even reduce the overlap between thetwo wavefunctions. This would lead to poorluminescence and also poor thermal characteristicsin device performance. It is also suspected thatGaNAsSb has a bandgap region which is shiftedtowards the conduction band edge, since Sb tendsto increase valence band-offsets while not affectingthe conduction band-offset. This could lead to atype-II alignment with the QW. GaNAs has alarger bandgap (�1.1–1.2 eV) and thus providesbetter electron and hole confinement. Finally, theoptical properties of GaNAsSb have been found tobe of relatively poor quality. Although PL is notthe absolute measure of optical and electricalquality of a material, it does give a very good idea.From our investigations, we have found that

GaNAs is the better barrier material comparedto GaNAsSb for GaInNAs(Sb) QWs.The device structures reported in Sections 5 and

6 for our long-wavelength GaInNAs(Sb) optoelec-tronic devices were changed to reflect the newknowledge that GaNAs was a superior barriermaterial. This change led to an improvement inGaInNAsSb PL quality. For GaInNAsSb QWsamples at 1.55 mm, the FWHM of the PL peakwas 57meV with GaNAsSb barriers and 44meVwith GaNAs barriers. The FWHM of the PL peakis a good measure of the improved quality of thematerial and structure.

3. Plasma source ion damage

A second critical growth issue is the potentialfor ion or electron damage from the RF plasma

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Anneal Temperature (C)

None 720 740 760 780 800

No Deflection

–40V Deflection

+18V Deflection

PL

Inte

nsity

45

40

35

30

25

20

15

10

5

0

Fig. 9. PL for samples grown with different voltages applied to

one deflection plate. The other plate was grounded. Note the

higher luminescence over all annealing conditions, and the

higher anneal temperatures reachable before the PL is

quenched.


source used in MBE systems to produce a reactiveN species. Nitrogen is difficult to incorporate intoGaAs, with the equilibrium solubility o0.5%.GaNAs and related dilute nitride materials mustbe grown under non-equilibrium conditions byMBE or OMVPE. In the case of MBE, a RF Nplasma source is used to supply atomic N for thegrowth and alloys with up to 10% N have beengrown [12]. There is little quantitative knowledgeof exactly what species are being generated in suchRF plasma sources, with N, N2

+, N2 radicals andelectrons being the most prevalent. As a result,there has always been a question of whether N2

+

ions or electrons do any damage to the growingepitaxial materials. Another key question being ifradicals incorporate as N–N pairs or as a Ninterstitial or [N]Ga antisite.There are various methods that can be used to

remove ions from the output of a plasma source.The RF power and gas flow rate have a strongeffect on the fraction of ions in the output of theplasma source, with low RF power and high flowgiving the fewest ions but also the least atomic N.Similarly, we verified that using smaller or fewerholes in the aperture at the end of the plasmasource decreased plasma damage and improvedthe stability of the plasma [12,26,27].Another common technique is to add voltage

ion deflection plates at the output of the plasmasource. These are simply parallel, metal plates oneither side of the output beam of the plasma. If abias is applied across the two plates, the electricfield will drive positive ions in one direction andelectrons in the opposite direction. Negative ionsare assumed to be negligible. We initially applied7800V to the deflection plates, for 1600V totalbias and found no consistent improvement in PLof GaInNAs [22]. We now believe that such a highvoltage bias may have ionized additional Nthrough a DC-enhanced plasma in front of thesource or it may have caused sputtering of metal oradsorbed contaminants from the deflection plateswhich landed on the wafer.One of the features of most MBE systems is that

a nude ion gauge is mounted on the backside ofthe substrate holder to measure beam fluxes beforegrowth. The ion gauge can be used to measure theion and electron fluxes emanating from the plasma

source and the effect of deflection plate bias. Byrotating the ion gauge up and down slightly andwith different biases on the filament, one canproduce a reasonable spatial map and energydistribution for both electrons and ions from thesource. By measuring the location of peak electroncurrent, we were able to demonstrate that smallvoltages (o30V) physically deflect the ions andelectrons such that they do not impact thesubstrate.Following the mapping of the electron and ion

distributions with deflection plate bias, we set offto test the hypothesis that minimizing ions incidenton the sample would result in better material. Wegrew three samples with biases of +18, 0, and�40V applied to one deflection plate for therespective samples. The other plate was groundedfor all three samples. Each sample was a single7 nm GaInNAsSb QW with GaNAs barriers.Fig. 9 shows the PL intensity as a function ofanneal temperature for these three samples. ThePL not only shows very clear improvement withdeflection bias, but that improvement appears tokeep rising to the highest anneal temperature ortime. This is in complete contrast to all ourmaterials grown without deflection plate biaswhich peak at some anneal temperature/time andthen decrease when either of these is exceeded. At

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100

150

200

10

15

20

ut P

ower

(m

v)

Effi

cien

cy (

%)


high anneal temperatures, the sample with +18Vdeflection had more than 5� greater PL intensityand at low pump powers, the relatively narrow PLlinewidth was reduced even further with deflectionbias and PL linewidth of 32meV was observedafter 800 1C anneal [26].

0 200 400 600 8000

50

0

5CW

Out

p

Current (mA)

Wal

lplu

g

14600.0

0.2

0.4

0.6

0.8

1.0

1.2

1470 1480 1490 1500

Wavelength (nm)

Pow

er (

A.U

.)

1.05xThreshold

4x Threshold

6x Threshold

(a)

(b)

Fig. 10. (a) Light output power and wallplug efficiency vs.

injection current (L2I2Z) characteristic for 20 mm wide

� 1220mm long ridge waveguide edge emitting laser operating

CW at room temperature. The CW threshold current is 580A/

cm2. (b) Emission spectra at �1.05� , 4� and 6� Ithoperating CW at room temperature.

4. GaInNAs high power edge emitting lasers

The challenges for achieving luminescence andlasing beyond 1.3 mm were illustrated in Fig. 1before the addition of Sb. With the addition of Sb,we were able to achieve strong PL out to 1.6 mm,

hence we initiated growth and fabrication of edgeemitting lasers to operate at 1.5 mm. With theimproved growth conditions (low temperature,deflection plate bias and GaNAs barriers), singleQW PL samples exhibited a room temperature(RT) PL peak at 1.42 mm and a linewidth of29.7meV. This is the narrowest PL linewidthreported for such QW structures at wavelengthsgreater than 1.3 mm and is an indication of themuch improved growth conditions for these PLsamples and resulting lasers [7].Lasers were grown on (1 0 0) n-type GaAs

substrates. Solid silicon and CBr4 sources wereused for n- and p-type doping, respectively. Theactive layer was a single, 7 nm GaInNAsSb QWsurrounded by 20 nm thick GaNAs barriers. TheQW active region was embedded in the center of aone wavelength thick, undoped GaAs waveguidesurrounded by 1.8 mm thick Al0.3Ga0.7As claddinglayers. The structure was capped with a highly p+-doped GaAs contact layer. Device fabricationconsisted of lift-off metallization (Ti/Pt/Au) fol-lowed by a self-aligned dry etch through the topAl0.3Ga0.7As cladding layer. The wafers werethinned to 120 mm to allow the cleavage of high-quality mirrors and backside metal (AuGeNi/Au)was evaporated. This was followed by a 1mincontact sinter at 410 1C. Device bars ranging inlength from 800 to 1200 mm were manuallycleaved.Fig. 10(a) shows the CW RT light output and

wallplug efficiency with input current (L2I ; Z2I)for a device measuring 20� 1220 mm. The devicelased at 1.460 mm (Fig. 10(b)) with a CW threshold

of 580 A/cm2. The CW slope efficiency was 0.45W/A, corresponding to 53% external quantumefficiency. These are the lowest threshold currentdensity and efficiency data reported for a dilute-nitride laser operating at l41:4mm [7,27]. Thesedevices showed maximum CW output powers of200mW (unbonded with no heat sink) beforethermal rollover and pulsed threshold currentdensities of 450A/cm2 and output powers (5 mspulse, 1% duty cycle) as high as 1.2W (pulsercurrent limited). The device structure was chosento minimize thermal heating around the activelayer to allow CW operation. A difference of 130

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0

0.5

1

1.5

2

2.5

3

3.5

4

1.15 1.25 1.35 1.45 1.55

< 1.3 um GaIn(N)As

> 1.3 um GaInNAs(Sb)

This Study GaInNAsSb

J th

(kA

/cm

2 )Wavelength (µm)

Newest result

Fig. 12. Threshold current of GaInNAs lasers

(1:154lo1:25mm) from several groups and our earlier

GaInNAsSb lasers (l41:30mm) and most recent result at

1.46mm with optimized growth conditions.


A/cm2 between pulsed and CW threshold currentdensity clearly indicates some heating was occur-ring for CW operation.The turn-on voltage was 0.9V, which compares

favorably to the bandgap energy of 0.83 eV.Additionally, the diode ideality factor was com-puted to be 1.3. These data indicate an excellentbalance between device heating issues and freecarrier absorption.One of the key differences in comparison to our

earlier results is the reduction in non-radiativerecombination. This is clearly illustrated in thecharacterization of recombination mechanismsusing the Z-parameter, which is defined as thederivative of the log of the total laser currentdivided by the derivative of the square root ofmeasured spontaneous emission [28,29]

Z ¼dðlnðIÞ

dðlnðSPE12Þ

.

The Z-parameter vs. current is illustrated inFig. 11 and shows that the newer devices have avalue of 2 or greater at much lower currents thanour earlier devices. This is most likely due to thereduction of ion damage from the plasma sourcedue to the bias voltage applied to the deflectionplates in the newer devices.

0.0 0.5 1.0 1.5 2.01.0

1.5

2.0

2.5

3.0

3.5

4.0

Z-P

aram

eter

Current Density (kA/cm2)

New Devices (55°C)

Previous Devices (55°C)

Fig. 11. Z-parameter extracted from measured spontaneous

emission as a function of bias current for prior devices (no

plasma source bias and GaNAsSb barriers) and newest devices

(plasma source bias and GaNAs barriers) [29].

A comparison of devices from several groups forlasers with l41:25mm is shown in Fig. 12. Thesignificant improvement in materials technology isclearly evident in the factor of 8� lower thresholdcurrent density compared to other reported devicesat comparable wavelengths. While the thresholdcurrent is very impressive, these devices are moretemperature sensitive. The characteristic tempera-ture (T0) under pulsed conditions was 62 1K asmeasured from 25 to 60 1C. This is likely due toreduced electron confinement in GaInNAsSb/GaNAs structures as compared to GaInNAs/GaAs [6] structures and will be tested in the nearfuture.

5. GaInNAs long wavelength VCSELs

The VCSEL structure was described in Ref. [8]and consists of a bottom mirror, top mirror andone-lambda cavity with three QWs. The bottommirror consisted of 29 alternating pairs of silicon-doped Al0.92Ga0.08As and GaAs, making a DBR.The cavity was a one-lambda thick layer of GaAs,designed for 1.485 mm, with three QWs at thecenter. The QWs were based on edge emitting

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lasers which operated at wavelengths from 1.49 to1.51 mm. The QWs were 7 nm Ga0.62In0.38-N0.016As0.958Sb0.026 with 20 nm GaNAs barriers.The top DBR was p-doped with carbon, and 24pairs thick. The Sb flux was 1.15� l0�7 Torr, andthe BEP was 20 times the Group-III flux in theQWs and 15 times elsewhere. Plasma conditionswere optimized to minimize plasma damage duringgrowth. A conventional liftoff process was used todefine metal rings for the top contacts, however,the liftoff failed and the metal and wafer surfacewere severely damaged, leading to rough side wallsand surfaces with significant scattering loss.Because of processing problems, the top DBRwas oxidized uniformly about 10 mm inwards. Thisadded optical scattering losses and series resistanceto the top DBR. Despite all the above difficulties,the VCSELs lased in pulsed mode when cooled toa chuck temperature of �10 1C. Fig. 13 shows thefiber-coupled power, as a function of wavelength,at 500 and 800mA of peak current, showing theonset of stimulated emission. The VCSELs werepulsed at 0.1% duty cycle, with 2 mm pulses at a500Hz repetition rate. Multiple transverse modesare visible above threshold, due to the large 66 mmdiameter current aperture of the VCSELs. Thethreshold current Ith, was 543mA (pulsed),corresponding to current density, Jth of 16 or5.3 kA/cm2 per QW. The VCSEL QWs were

2nm

5dB

800mA500mA

Center 1459nm

Fig. 13. Spectra of the first monolithic l41:45mm VCSEL on

GaAs operating just below and about 1.6� threshold.

identical to the single QW from our earlier CW1.49 mm edge emitting lasers. Due to the loweroperating temperature and a short cavity, theVCSELs lased at 1.46 mm, a significantly shorterwavelength. Several other devices produced up to33 mW peak power. This modest power is surpris-ingly high given the problems outlined above andthe mismatch between the cavity resonance andthe actual emission wavelength (gain peak). It isexpected that even a simple improvement inmounting technique would reduce the need forcooling the devices. This is the first reportedmonolithic VCSEL grown on GaAs to lase atwavelengths greater than 1.45 mm [8].

6. Summary

The key challenge for expansion of widebandwidth communications to the desktop isgreatly reduced cost of lasers, semiconductoroptical amplifiers, pump lasers and monolithicintegration of all components into photonicintegrated circuits. I believe that GaInNAsSb onGaAs will be the foundation technology that willenable low cost, wide bandwidth MAN/LAN/SAN networks, optical switching and routers.Dilute nitride GaInNAsSb grown epitaxially onGaAs can produce active QW regions that coverthe full 1.2–1.6 mm wavelength region with a singlealloy materials system. This not only greatlysimplifies processing, but enables incorporationof the tremendous processing advantage of oxi-dized AlAs to form high index contrast photoniccrystal structures as well as the existing AlAs/GaAs DBR mirror and VCSEL manufacturingtechnology. This will become an enabling technol-ogy for not only long wavelength lasers, butadvanced photonic integrated circuits.The major challenge of this materials system has

been to understand the differences of dilutenitrides compared to other III–V alloys and toproduce low threshold lasers at any desiredwavelength between 1.3 and 1.6 mm. The mostrecent results incorporating Sb to form a quinaryalloy, GaInNAsSb, appear to overcome many ofthe prior problems with phase segregation. Under-standing the role of damage from plasma sources

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has dramatically reduced the threshold current oflong wavelength lasers. I believe that GaInNAsSbwill be the active gain material of choice because ithas significantly higher gain for VCSELs, is closerto the existing QW technologies than InAs QDsand it has fundamental energy band advantagesover its other competitors. While MBE has beenutilized for production of very low cost, edgeemitting CD-lasers, it has not been utilized in theproduction of VCSELs, although it has been thetool of choice for most of the research anddevelopment of VCSELs. The newest versions ofproduction MBE systems with their greaterversatility in number of liquid metal sources couldeasily change the production role of MBE. Theadvances in equipment combined with the sig-nificantly easier growth of GaInNAsSb by MBEwill likely make MBE the choice for production ofVCSELs, high power edge emitting lasers andphotonic integrated circuits. Progress has beenrapid for such a complex materials system and thefuture for this materials system and the potentialfor its inclusion as a major technology for opticalnetworks is indeed bright.

Acknowledgements

This work has been the result of efforts by anumber of truly great graduate students, includingSeth Bank, Mark Wistey, Homan Yuen, VincenzoLordi, Tihomir Gugov, Lynford Goddard, HopilBae, Vincent Gambin, Wonill Ha, Sylvia Spruytte,Chris Coldren, Mike Larson and Dr. Kerstin Volz.I would also like to thank Dr. Akihiro Moto,Sumitomo Electric for supporting this work, Prof.Alfred Adams at University of Surrey for helpfuldiscussions on absorption, carrier leakage, recom-bination mechanisms and Z-parameter and StephenSmith, Charles Evans and Associates for SIMSprofiles. This work was supported both financiallyand technically by Drs. Bob Leheny and Elias Toweat DARPA and Dr. Yoon-Soo Park at ONR over anumber of years through the OptoelectronicsMaterials Center, contracts MDA972-00-1-0024,DAAG55-98-1-0437, N00014-01-1-0010, the MAR-CO Interconnect Focus Center and the StanfordNetwork Research Center.

References

[1] J.S. Harris Jr., Semicond. Sci. Technol. 17 (2002) 880 and

references therein.

[2] J.S. Harris Jr., in: I. Buyanova, W. Chen (Eds.), Physics

and Applications of Dilute Nitrides, Taylor & Francis,

London, 2004 and chapters therein.

[3] J.S. Harris Jr., IEEE J. Sel. Top. Quantum Electron. 6

(2000) 1145.

[4] M. Kondow, K. Uomi, A. Niwa, T. Kitantai, S. Watahiki,

Y. Yazawa, Jpn. J. Appl. Phys. 35 (1996) 1273.

[5] D.W. Kisker, L.M.F. Chirovsky, R.L. Naone, J.M. Van

Hove, J.M. Rossler, M. Adamcyk, N. Wasinger, J.G.

Beltran, D. Galt, SPIE Photonics West Conference

Proceedings, San Jose, CA, January 2004, to be

published.

[6] W. Ha, V. Gambin, S. Bank, M. Wistey, H. Yuen, S. Kim,

J.S. Harris Jr., IEEE J. Quantum Electron. 38 (2002) 1260.

[7] S.R. Bank, M.A. Wistey, H.B. Yuen, L.L. Goddard, W.

Ha, J.S. Harris Jr., Electron. Lett. 39 (2003) 1445.

[8] J.S. Harris Jr., H. Yuen, M. Wistey, S. Bank, V. Lordi, T.

Gugov, H. Bae, L. Goddard, in: M. Henini (Ed.), Dilute

Nitride Semiconductors, Elsevier, Amsterdam, 2005 Chap-

ter 1 and references therein.

[9] A.R. Kovsh, J.S. Wang, L. Wei, R.S. Shiao, J.Y. Chi, B.V.

Volovik, A.F. Tsatsul’nikov, V.M. Ustinov, J. Vac. Sci.

Technol. B 20 (2002) 1158.

[10] M.A. Wistey, S.R. Bank, H.B. Yuen, J.S. Harris, Appl.

Phys. Lett. (2004) to be published.

[11] S.G. Spruytte, M.C. Larson, W. Wampler, C.W. Coldren,

H.E. Petersen, J.S. Harris, J. Crystal Growth 227–228

(2001) 506.

[12] A. Hasse, K. Volz, A.K. Schaper, J. Koch, F. Hohnsdorf,

W. Stolz, Cryst. Res. Technol. 35 (2000) 787.

[13] J. Massies, N. Grandjean, Phys. Rev. B 48 (1993)

8502.

[14] H. Shimizu, K. Kumada, S. Uchiyama, A. Kasukawa,

Electron. Lett. 36 (2000) 1379.

[15] X. Yang, J.B. Heroux, L.F. Mei, W.I. Wang, Appl. Phys.

Lett. 78 (2001) 4068.

[16] S.R. Bank, W. Ha, V. Gambin, M.A. Wistey, H.B. Yuen,

L.L. Goddard, S.M. Kim, J.S. Harris Jr., J. Crystal

Growth 251 (2003) 367.

[17] L.H. Li, V. Sallet, G. Patriarche, L. Largeau, S.

Bouchoule, K. Merghem, L. Travers, J.C. Harmand,

Electron. Lett. 39 (2003) 519.

[18] G. Ungaro, G. LeRoux, R. Teisseir, J.C. Harmand,

Electron. Lett. 35 (1999) 1246.

[19] J.C. Harmand, G. Ungaro, L. Largeau, G. LeRoux, Appl.

Phys. Lett. 77 (2000) 2482.

[20] K. Volz, V. Gambin, W. Ha, M.A. Wistey, H.B. Yuen,

S.R. Bank, J.S. Harris Jr., J. Crystal Growth 251 (2003)

360.

[21] J. Tournie, N. Grandjean, A. Trampert, J. Massies,

K. Ploog, J. Crystal Growth 251 (2003) 383.

[22] H.B. Yuen, S.R. Bank, M.A. Wistey, A. Moto, J.S. Harris

Jr., J. Appl. Phys. 96 (2004) 6375.

ARTICLE IN PRESS


[23] T. Gugov, V. Gambin, M. Wistey, H. Yuen, S. Bank,

J.S. Harris Jr., J. Vac. Sci. Technol. B 22 (2004) 1588.

[24] P.J. Klar, K. Volz, J. Phys. Cond. Matter 16 (2004) S3053.

[25] J.-M. Chauveaua, A. Tramperta, M.-A. Pinaultb, E.

Tournie, K. Duc, K.H. Ploog, J. Cryst. Growth 251

(2003) 383.

[26] S.R. Bank, M.A. Wistey, L.L. Goddard, H.B. Yuen, J.S.

Harris, Electron. Lett. 40 (2004) 1186.

[27] M.A. Wistey, S.R. Bank, H.B. Yuen, J.S. Harris, J. Cryst.

Growth (2005) (this issue, MBE 2004 Proc.).

[28] R. Fehse, S. Tomic- , A. Adams, S. Sweeney, E. O’Reilly, A.

Andreev, H. Riechert, IEEE J. Sel. Top. Quant. Elect. 8

(2002) 801.

[29] L.L. Goddard, S.R. Bank, M.A. Wistey, H.B. Yuen, Z.L.

Rao, J.S. Harris Jr., J. Appl. Phys. 97 (2005).

Documents

The opportunities, successes and challenges for GaInNAsSb