PHYSICAL REVIEW B, VOLUME 63, 245311

Microscopic characterization of InAsÕIn0.28GaSb0.72ÕInAsÕAlSb laser structure interfaces

W. Barvosa-Carter,* M. E. Twigg, M. J. Yang, and L. J. Whitman†

Naval Research Laboratory, Washington, Washington, D.C. 20375~Received 26 January 2001; published 4 June 2001!

We have used cross-sectional scanning tunneling microscopy~XSTM! and transmission electron microscopy~TEM! to study InAs/In0.28Ga0.72Sb/InAs/AlSb strained-layer heterostructures designed for use in infraredlasers. The samples came from the same material previously characterized by photoluminescence~PL! andx-ray diffraction @M. J. Yanget al., J. Appl. Phys.86, 1796 ~1999!#. Several structures grown at differenttemperatures and with either III-As or III-Sb-like interfacial bonds have been characterized. Analysis ofhigh-resolution TEM images finds the same degree of interfacial roughness~;1 ML! for both III-As and III-Sbinterfacial bonded heterostructures, despite significantly greater PL intensity in the latter. We also implementand compare two different methods for analyzing the interfacial roughness using XSTM; both show that thecrucial InAs/InGaSb interface is rougher in the samples grown at high temperature. Even in samples grown atthe optimal temperature~;440 °C!, XSTM reveals intermixing at the AlSb-on-InAs interfaces, as well asunexpected differences in the interfacial bond types at the InAs-on-AlSb vs AlSb-on-InAs interfaces. Whereasall layers grown at or below the optimal growth temperature appear defect-free in TEM, threading dislocationsare observed in samples grown at higher temperature.

DOI: 10.1103/PhysRevB.63.245311 PACS number~s!: 68.37.2d, 68.55.Ln, 68.65.Cd

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I. INTRODUCTION

Considerable research on midwavelength~3–5 mm! andlong-wavelength~8–14mm! infrared~IR! diode lasers is being driven by both military and commercial demand for hioutput power and noncryogenic operation.1–3 Molecularbeam epitaxy~MBE! has allowed the creation of novel lasinmaterials, new lasing transitions, and record-setting permances for devices operating in these wavelength range4–6

However, in order to optimize these devices, control omesoscale and nanoscale defects in the material musachieved. The strain in the heterolayers, for instance,lead to strain-driven roughening of the growth surface athe introduction of misfit and threading dislocations. At tnanometer scale, precise control over layer thickness mumaintained, because evenmonolayer-scaleroughness and interdiffusion between the composite layers in the heterosttures can cause degradation in device efficiencyperformance.7,8 Because very little information is currentlavailable concerning such defects in actual device structutheir characterization should hasten further improvementthe device material.

In this article, we report our investigations of the nanscale and mesoscale structural properties of InAs/InGaInAs/AlSb superlattices designed for ‘‘W-structure’’ mid-Ilasers.9 A qualitative sketch of the spatially varying bangaps and offsets is given in Fig. 1~a!. The electron-hole re-combination process that leads to lasing takes place acthe InAs/InGaSb interfaces, making these the most critinterfaces in the structure. The primary function of the Allayer is to confine the electrons and holes in the InAInGaSb/InAs layers, and to suppress the formation oftended three-dimensional electron states between neighing sets of InAs layers. The AlSb/InAs interfaces and tquality of the AlSb layer are therefore deemed less criticaoptimizing the lasing process, although defects in and c

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ductive pathways through the AlSb layer would certainlydeleterious to device performance.

The photoluminescence~PL! and double-crystal x-ray dif-fraction ~XRD! of these laser structures are strongly depdent on MBE growth temperature.10,11As shown in Fig. 1~b!,

FIG. 1. ~a! Composition and band offsets for the W-type lasstructures evaluated in this work. The boxes represent the bandof the different material layers. Carrier recombination leadinglight emission occurs between quantum-confined electron levethe InAs layers~E! and heavy- and light-hole bands in InGaSb~H!.~b! Integrated photoluminescence intensity at room temperaturesamples grown at various temperatures with either III-Sb or III-interfacial bonds~Ref. 11!. The samples characterized in this stuare circled.

©2001 The American Physical Society11-1

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BARVOSA-CARTER, TWIGG, YANG, AND WHITMAN PHYSICAL REVIEW B63 245311

the optical characteristics are optimal within a rather narrgrowth temperature range~410–450 °C!, and are typicallymuch worse outside of that range. In addition, the inclusof III-As vs III-Sb-like interfacial bonds between the InAand InGaSb layers drastically affects the material qualitymeasured by PL, but appears to leave the macroscopic stural quality unchanged as measured by XRD.10,11 The goalof this work is to understand the source of the dependencmaterial quality on temperature and bond type, and to idtify the defects occurring in the material under variogrowth conditions. We attempt to directly correlate tatomic-scale and mesoscale properties of the material,as dislocation density and interfacial roughness, with groconditions and macroscopic properties, by studying samof the same materialpreviously characterized by PL anXRD. To characterize the properties of the layers and infaces over a range of length scales, we used a combinatiotransmission electron microscopy~TEM! and atomic-resolution cross-sectional scanning tunneling microsc~XSTM!.

Whereas TEM is well developed and widely implementfor characterization of electronic devices, XSTM is generaless familiar for this application. It is a potentially powerftechnique for imaging defects in superlattice structurwhere one uses STM to image a$110% cleavage face of apiece of superlattice material, allowing ‘‘edge-on’’ crossectional characterization of the as-grown superlatticeparticular, XSTM can obtainatomic-scaleinformation aboutthe superlattice layers and interfaces, although it has slimitations that have only recently been explored in detai12

One limitation of particular relevance to this work is the fathat only every other atomic layer in the superlattice canimaged. Here we will also describe the impact of thlimitation on our ability to adequately quantify interfaciroughness.

II. EXPERIMENT

The samples studied here came from superlattices grand previously characterized by a varietytechniques.10,11,13 Briefly, the ‘‘W’’ structures were grownon GaSb substrates in an MBE chamber equipped witvalved As cracker, an Sb cracker, and a conventionaleffusion cell. The buffer layer consisted of 0.3mm of GaSbgrown at 530 °C followed by 1.0mm of AlSb grown at580 °C. The substrates were then cooled without Sb2 flux toa lower temperature for superlattice growth. The foconstituent ‘‘W’’ active layer consisted of 20 periods of 5ML InAs/10-ML In0.28Ga0.72Sb/5.5-ML InAs/14-ML AlSb.Subsequently, a capping layer consisting of 0.2mm of AlSband 10 nm of GaSb was grown at 480 °C. Active layers wgrown at a variety of temperatures with either InSb or Gainterfacial bonds.10,11

High-resolution TEM~HRTEM! of cross-sectional TEM~XTEM! samples was performed by imaging along the@001#direction through the corner of cleaved samples.14 OtherXTEM samples were prepared by mechanical lappinglowed by ion milling at liquid-nitrogen temperature. Micronscale features such as threading dislocations were exam

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by imaging cleaved plan-view samples with the incideelectron beam directed 45° from the growth direction.described in detail elsewhere,12 XSTM samples scribed fromthe same wafers were mounted on an STM sample holAfter introducing the samples into the STM vacuum chaber, a sample was then scribedin situ and cleaved to expos

either a~110! or (1̄10) surface. Single-crystal tungsten tipwere prepared by electrochemical etching and cleanedin situby electron-bombardment heating prior to use. All constacurrent images shown are of filled states on~110! surfacesunless otherwise noted.

III. RESULTS AND DISCUSSION

A. Interfacial roughness

To understand the effect of growth temperature onatomic-scale structure of the material, we have used XSto compare several laser structures with InSb interfabonds: one grown at the optimal temperature~based on PL!and another grown at higher temperature.@The samples studied are circled in Fig. 1~b!#. Here we wish to distinguishbetween two sources of interfacial disorder:~1! roughness, asmight arise from incomplete layer completion during epitaand ~2! interlayer mixing at the InAs/InGaSb/InAs interfaces. Figure 2 compares XSTM images from the~110! facesof the two cleaved superlattices~optimal vs higher temperature!. Both samples were grown on the same day using oerwise identical conditions. Close examination of the imagshows that roughness at the crucial InAs/InGaSb interfappears to be somewhat larger at the higher growth tempture. In addition, the InGaSb-alloy layers appear less unifoin the high-temperature samples. In this section, we willcus on characterizing the roughness of the interfaces.

To obtain a quantitative measurement of this roughnewe have followed and extended the methods of Feenstraco-workers15–17 and Harperet al.18,19 Interface profiles aretypically extracted from an STM image by first taking a drivative of the image and then extracting the interface proby tracing a contour of constant slope. This method cansult in a reasonable facsimile of the interface profile providthe individual surface atoms are not well resolved and this sufficient image contrast between the layers. However,approach becomes difficult to implement when the surfatoms are well resolved. In this case, one can extract inface profiles from individual images by hand, by manuainspecting each atom in the image and determining whlayer it belongs to based on its height in the STM imagAlthough the height is largely determined by the identitythe surface atom~e.g., Sb vs As!, the presence of an alloycan confuse this method because the identity of the subface atoms~In or Ga, randomly determined by the alloy composition! can also affect the height.20,21 The effect on themeasured height is greatest for the first subsurface layer,deeper layers contributing successively less. The net effeto make the heights of the atoms in the alloy appear higirregular. This variability is particularly pronounced at ainterface, because the interfacial bond type~i.e., long Sb-like

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MICROSCOPIC CHARACTERIZATION OF . . . PHYSICAL REVIEW B 63 245311

bonds vs shorter As-like ones! can enhance the alloy effecand hinder unambiguous determination of the interfaprofile.12,22

In order to avoid the subjective assessment requiredvisually trace each interface and to additionally enablerapid analysis of multiple images, we have developedalternate approach to tracing the interface profiles. As illtrated in Fig. 3, we start with an image at a resolutionabout 5 pixels per unit cell along the interface. We th

FIG. 2. Atomic-resolution XSTM images (23323 nm) of twolaser structure samples grown with nominal III-Sb interfacial bonat ~a! the optimal temperature~as measured by PL! and~b! at hightemperature. The sample biases and tunneling currents wer~a!22.0 V, 50 pA and~b! 22.0 V, 0.1 nA.

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smoothen the image using a 333-pixel window, and thenreduce it to black and white using a threshold value midwbetween the average pixel intensities on either side ofinterface. Finally, the profile of the interface is extractfrom the two-color image using simple edge detection. Bcause the contrast in each image is normalized and eachage is treated to the same sequence of manipulationsinterface definition is unique across the entire set of dataall surface atoms are treated on a consistent basis. Wimplementing this method, it is useful to oversample theterface profile, i.e., to use more data points than atomicsitions along the interface. Oversampling helps take intocount those surface atoms that, due to alloy or interfastoichiometry effects, have a topographic height intermedbetween the bulk layers on either side. When analyzpower spectra from the oversampled interface profile, hoever, the power spectral density for wavevectors aboveNyquist limit defined by the surface unit cell mesh mustexcluded from subsequent analysis.23

The interfacial power spectrum can be directly obtainfrom a fast-Fourier transform~FFT! of the interface profile.The associated power spectral density can be very accurathe imaged interfacial roughness extends over multiple lers. However, because only every other layer of the crystaimaged, if the roughness includes only two or three growlayers, this method can lead to errors: a significant portionthe roughness may be obscured between observed atlayers.12 To examine the potential impact of this problem, whave analyzed our interface profiles that fall into this cegory using two techniques: the usual method of Feenand co-workers,15–17 which we will denote as the ‘‘direct’’method, and a modification of this approach, which we wcall the ‘‘inspection’’ method. When implementing the diremethod, we treat all interface measurements in units of$110%-row spacing, 6.1 Å~corresponding to 2 ML during

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FIG. 3. Sequence of images illustrating the processing sused to extract interface profiles from the XSTM images.~a! Origi-nal gray-scale image,22 V, 0.1 nA ~at least 5 pixels per surfacunit cell along the interface!. ~b! Image after twice replacing eacpixel with the minimum value of its 333 kernel, and then replacingeach pixel with the 333 maximum. ~c! Two-color image withthreshold midway between the average values of the two layers~d!The interface defined by simple edge detection.

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BARVOSA-CARTER, TWIGG, YANG, AND WHITMAN PHYSICAL REVIEW B63 245311

epilayer growth!. Power spectra fromall interfaces, includ-ing ones with near-zero rms roughness, are averaged wdetermining the final power spectra. In contrast, withininspection method, we deliberatelyexcludeinterface profilesfrom the analysis that exhibit near-zero rms roughness,those where it appears that the actual interface is hidden fobservation. Furthermore, for cases where the interfacvisible across a single row on the$110% surface, we treat theobserved roughness in units of 3.05 Å~the growth mono-layer height!. This makes physical sense because stepsislands on the~001! growth surface are virtually always thihigh.12

The roughness of the exposed cleavage surface typicincreases towards the capping layer, so that finding portof the superlattice long enough for analysis becomes mdifficult farther away from the buffer layer. Hence, we halimited our XSTM study to the interfacial roughness of tfirst three to four superlattice periods out of the 20 growFrom these segments of the superlattice, we have extranearly 100 interface profiles of length;215 Å from numer-ous atomic-resolution images~51 surface unit cells sampleat 5 data points per unit cell!.24 We then obtained the powespectral density from each using an FFT power-spectral dsity estimator with Welch windowing.25 The power spectrafor a particular interface—e.g., the InAs-on-InGaSb interfafor the optimal temperature sample—were then averagedfit to the usual Lorentzian power-spectral density plusadditional background term,

uAqu2L52D2L

@11~qL!2#1B3, ~1!

whereAq is the Fourier amplitude at wavevectorq, L is thelength of the interface profiles,D is the roughness amplitudeL is the correlation length, andB represents the ‘‘whitenoise’’ component arising from uncorrelated poidefects.18,19

In Fig. 4, we show the power spectra extracted usingdirect method for the InAs-on-InGaSb and InGaSb-on-Ininterfaces for both the optimized~a! and high-temperature~b! samples. The data have been fit to Eq.~1! with ~solidline! and without~dashed! the white noise background termB. The averaged power spectra extracted for both sampleusing both the direct and inspection methods are showFig. 4~c!, along with the Lorentzian fits~without the whitenoise term!. The results of the fitting are summarizedTable I.

Comparing the results for the direct vs inspection methwe expect the direct method to overestimate the magnitof the actual interfacial roughness, and indeed we find thabe the case: it finds a larger roughness by a factor of;1.5.We also find that the length scale of the roughness issame for both methods—this should generally be true ifinterfacial roughness does indeed only extend over 2 MFor larger excursions, the inspection method will findlength scale that is too short and the direct method becomore accurate. In general, when the roughness extendsmore than 2 ML, the direct method should be more accufor quantifying the interfacial roughness.

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Two striking features arise from the analysis. First, wfind that the interfacial roughness in the high-growttemperature structure is indeed significantly larger thanthe structure grown at the optimal temperature by 0.4–0.rms ~depending on the method used!. Second, although wefind asymmetries between the InAs-on-InGaSb and InGaon-InAs interfaces, they all share a remarkably short colation length, 2–6 Å. The magnitude of the correlatiolength, L, for both samples is similar regardless of thpower-spectrum extraction method used or the interface t~InAs-on-InGaSb or InGaSb-on-InAs!. This similarity indi-cates that the structure of the growth surface does not chasignificantly over the;50 °C growth-temperature span use

FIG. 4. Average power-spectral densities~PSD’s! for the inter-facial roughness obtained for various III-Sb interfaces, samples,analysis methods.~a! Results for the optimal temperature sampfor profiles extracted by the ‘‘direct method’’ for the InGaSb-oInAs and InAs-on-InGaSb interfaces. Fits to Eq.~1! with ~solidlines! and without~dashed lines! the optional white noise parameteB are included. Note that the InGaSb-on-InAs fit is the same withwithout inclusion of white noise.~b! The same analysis for thehigh-temperature sample.~c! Comparison of the overall interfaciaroughness~given by an average of PSD’s for both interfaces! forthe optimal- and high-temperature samples, fit without the whnoise parameter.

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MICROSCOPIC CHARACTERIZATION OF . . . PHYSICAL REVIEW B 63 245311

We also find that the accuracy of the fit~as measured byx2!for the InGaSb-on-InAs interfaces does not improve signcantly when the background term is added, resulting ininsignificant background contribution. In contrast, the fitsimprove for the InAs-on-InGaSb interfaces if this termincluded, resulting in;30% smaller roughnessD androughly doubled correlation lengthL.

The simplest interpretation of the roughness results isthe interface is best characterized as a combination ofdom, short-range correlated clusters. Such clusters carise from either diffusive or anion-exchange related pcesses between the heterolayers, or reconstruction-restoichiometry changes during interface formation. Closespection of the interface profiles reveals this behavior in rspace: when observable, the interface is primarily compoof closely spaced but isolated or paired Sb or As atoHence, the main effect of high growth temperature onInAs/InGaSb/InAs interfaces is to simply increase the rroughness of the interface.

The strikingly short correlation lengths we observe, retive to the findings of Lewet al.26 and Feenstra andco-workers,15–17suggest a simple mechanism for the genetion of interfacial roughness in these structures that isrectly related to the use of MEE to form the interfaces. EaInGaSb-on-InAs interface in these samples was formedterminating the InAs layer with a monolayer of In followeby an Sb soak. The In layer is intended to prevent Sbexchange and to force InSb interfacial bonds. A residualfect of this procedure is that the; 1

2 ML of As terminatingthe InAs~001!-~234) reconstruction must be filled in bSb.12,27,28 In cross section, this would appear as shocorrelated clusters of As and Sb as observed. To formInAs-on-InGaSb interface, a monolayer of In was alsoposited on the Sb-terminated InGaSb surface to force Ibonding. It has been observed that 0.6–0.7 ML of thefrom the terminal InGaSb surface can segregate into thesequently deposited InAs.29 Hence, a significant amount oSb will be incorporated into the first InAs layers grown otop of the InGaSb layer. A short correlation length for the

TABLE I. Fitting parameters for the Lorentzian power-spectdensities plotted in Fig. 4.

RoughnessD ~Å!

Correlationlengthl ~Å!

White noiseB ~Å!

InGaSb-on-InAs, optimalT 2.0160.06 3.1060.36InGaSb-on-InAs, optimalT 2.0060.26 3.1160.83 0.0162.20InAs-on-InGaSb, optimalT 2.6260.08 2.2860.26InAs-on-InGaSb, optimalT 1.8360.20 4.7861.30 7.7861.93InGaSb-on-InAs, highT 2.9160.07 4.7860.37InGaSb-on-InAs, highT 2.7360.11 5.9960.78 2.7161.21InAs-on-InGaSb, highT 3.0260.08 2.2360.24InAs-on-InGaSb, highT 2.0260.20 5.2061.35 11.2262.18Direct, optimalT 2.2960.05 2.6860.23Inspection, optimalT 1.5160.04 2.6460.25Direct, highT 2.9060.06 3.3360.23Inspection, highT 1.8160.04 3.2360.23

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Sb atoms is certainly to be expected, because the diffulength of group-V atoms on anion-terminated surfacesgenerally accepted to be small.

Contrary to prior studies of interfacial roughness in InAInGaSb heterostructures, we find only slight differencestween the InAs-on-InGaSb and InGaSb-on-InAs interfafor either the sample grown at 410 or 450 °C. Lewet al.26

found that the InAs-on-InGaSb interfaces were significanrougher than the InGaSb-on-InAs ones in In0.25Ga0.75Sb/InAssuperlattices grown at 380 °C. In contrast, Feenstraco-workers15–17 observed the reverse to be true for InAGaSb multilayers grown at 380 °C: their InAs-on-GaSbterfaces were found to be smoother than GaSb-on-InAsterfaces. They explained this behavior based onthermodynamics of Sb segregation on InAs. However, thalso observed that the asymmetry gradually disappearesurface diffusion during growth was enhanced either bying growth interrupts at low growth temperatures orgrowing at higher temperatures. Our observations aretainly consistent with this picture, as both growth interrupand comparatively higher temperatures were employed.

Calculations of carrier transport through varioheterostructures7,30–32have shown that scattering causedinterfacial roughness typically occurs for intermediate corlation lengths, on a length scale near the exciton radius,;10nm in this system. Therefore, the impact of the short corlation lengths of our roughness on the device propertiesthis material is probably small, although calculations willneeded to verify the absolute magnitude of the effect. Tnext largest length scale in these materials is the terracewhich at.50 nm is probably larger than the critical lengscale for roughness-induced scattering. Therefore, the mappropriate model for understanding how the roughnessthese interfaces affects the device properties is to treat einterface as a ‘‘diffuse’’ structure,;1-ML wide, havingproperties intermediate between those of InAs, InGaSb,InSb.

B. Impurities and segregation effects

In addition to increased rms interfacial roughness wincreasing temperature, we also observe degradation inuniformity of the InGaSb alloy layer. There appears toincreased clustering, as revealed by XTEM~Fig. 5!. Thiscomposition modulation is revealed by XTEM becauselocal variations in the elastic relaxation at the free surfacaused by the varying lattice constant.33 Similar variationscan also be seen in XSTM images as shown in Fig. 6, whthe uniformity of the optimal- and high-temperature sampis compared. By integrating the observed topographic hewithin the InGaSb layer along the growth direction, we finthat clustering over a 5–10 nm length scale occurs mmore strongly in the high-temperature sample. Howevclustering in InGaSb is also evident in optimal samples,seen in Fig. 7, where both the cation~empty states! and anionlattices~filled states! are shown together.~The clustering isclearest in the empty-state image.! This clustering will de-grade the material properties in a manner similar to thainterfacial roughness. The observed clustering is most lik

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BARVOSA-CARTER, TWIGG, YANG, AND WHITMAN PHYSICAL REVIEW B63 245311

thermodynamic in origin, being driven either by phase seration of InSb and GaSb or the well-known strain-drivinstability predicted theoretically34–36 and observedexperimentally.37 In either case, clustering should becomincreasingly manifest at higher temperatures when latsegregation is less inhibited by the diffusion kinetics.

We find that several noteworthy defects occur even insuperlattices grown at the optimal temperature. As obserin Figs. 2~a!, 7, and 8, the AlSb layers also exhibit some sof contamination that appears as a series of topographichigher atoms in the XSTM, particularly near the AlSb-oInAs interface. The most likely source of this contaminatiis In segregation into the AlSb layer associated withMEE technique: each InAs layer was terminated withmonolayer of In, followed by an Sb soak to force this inteface to be InSb-like. Other possible sources of contaminacould be As from the interface, or the transient As baground pressure, or excess Ga from the InGaSb floatingthe surface of InAs.

In filled-state STM images of the$110% surfaces, one im-ages only the anion sublattice, and image contrast prima

FIG. 5. Dark-field XTEM images of optimal temperatusamples imaged using~a! a ^004& refection and~b! a ^220& reflec-tion. Because the strain effects due to clustering are visible ousing the latter imaging conditions, it is clear that compositiondulations are parallel to the growth plane. The diffraction vectorg isindicated.

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arises from two sources: anion identity and projected bolengths perpendicular to the surface. It has been shownexperimentally12,22 and theoretically38 that the local III-Vbond-length dominates the contrast for point defects. Forample, As residing on a GaSb anion sublattice~a local‘‘GaAs’’ defect! appears topographically lower than Sb,19 soit is extremely unlikely that the source of the contrast is dto As cross contamination in the AlSb layer. If the observfeatures were due to misplaced cations in the AlSb, the lobond length would be slightly shorter for Ga, but longer fIn ~as observed!. Similarly, we expect that local InSb structures should also appear topographically higher than bInAs and AlSb in empty-state images, as observed@Fig.7~a!#. Therefore, we conclude that In contamination is tonly type consistent with the XSTM contrast observed.

The In segregation creates discontinuous InSb layers uabout 4-ML thick sandwiched between the InAs and Allayers. The primary function of the AlSb layer is to separaout the superlattice periods to prevent the formation ofdesirable extended three-dimensional electronic statestween neighboring InAs layers. Therefore, this type of defwould not generally be expected to have a large impactthe electronic properties of the material. However, if tInSb layers locally perturb the symmetry of the superlatticausing slightly different electron energy levels in the adcent InAs layers, it could increase the width of the photominescence peak and reduce the total luminescence.

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FIG. 6. XSTM images and average-height profiles for InGaSalloy layers in an~a! optimal- and~b! high-temperature sample.~a!22.0 V, 50 pA and~b! 22.5 V, 0.5 nA.

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C. Interfacial bonding

Another potential materials problem observed even inoptimal samples is the apparent growth-order dependencthe structure of the two InAs layers in each superlatticeriod. This asymmetry can be seen most clearly in Fig. 8. Tfirst InAs layer in each ‘‘W’’ period tends to appear darkerfilled-state gray-scale images than the second. Upon clinspection, the source of this asymmetrical appearancseen to be the structure of the InAs-AlSb interfaces. Whereach InAs layer below AlSb has a relatively uniform appeance~gray arrows in the figure!, the InAs layer above the

FIG. 7. XSTM images of optimal-temperature sample comping simultaneously recorded~a! empty states~cation sublattice,12.0 V, 50 pA! and ~b! filled states~anion sublattice,22.5 V, 50pA!.

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AlSb on the top of each ‘‘W’’ has an InAs-on-AlSb interfacthat appears darker than the rest of the epilayer~black ar-rows!. One mechanism for such contrast asymmetry betwthe two interfaces is different interfacial bond lengths normto the surface associated with different interfacial botypes. For example, InAs-on-AlSb interfaces would apptopographically lower in the images if the interfacial bonare AlAs rather than InSb as intended. An AlAs-like bondsignificantly shorter than either an InAs or AlSb bond, awhen oriented perpendicular to the surface, the terminalatoms will be topographically lower than adjacent As surfaatoms bonded to In. More detailed discussions of this effcan be found elsewhere.12,22,38

To help identify the interfacial bonds at the InAs-on-AlSinterfaces, we have grown two InAs/AlSb/InAs test hetestructures: the first with deliberately mixed As-Sb interfacbonds and the second with uniform InSb bonds, prepaanalogously to those in the laser structures. As previoureported, the appearance of interfacial bonds depends oncleavage face examined, i.e.,~110! vs (1̄10).12,22XSTM im-ages of the mixed interfacial bonds indeed show topogracally lower AlAs bonds at the InAs-on-AlSb interface on th~110! cleavage face@Fig. 9~a!# and at the AlSb-on-InAs in-terface on the (1̄10) face@Fig. 9~b!#. In principle, a similarexamination of the different cleavage faces would elucidthe AlSb-on-InAs interfacial bonds in the laser structusamples. Because of the limited material available, we wunfortunately not able to image both cleavage faces on thsamples, so we cannot definitively state whether AlAs intfacial bonds are present or not at the AlSb-on-InAs intfaces. However, given the presence of the In contaminain the AlSb layer~indicative of an In-rich, InSb interface!, itis unlikely that AlAs bonding is significant at that interfac

The observation of AlAs-like interfacial bonds in the lasstructures is very surprising given the growth proceduused. The 2-s Sb interrupt applied at each InAs-on-AlSbterface has been previously shown by x-ray superlatticefraction to create an InSb-like interface.39,40 To test the in-fluence of interface formation techniques on the interfacbond type, we grew the second test superlattice with no

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FIG. 8. XSTM image of multiple periods of an optimatemperature sample highlighting the difference in the InAs laywhen they are InAs-on-AlSb~dark arrows! vs AlSb-on-InAs~grayarrows!. The image is of a~110! surface at22.0 V, 50 pA.

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nally identical conditions and shutter sequences used to gthe optimized laser structures. Curiously, in these samwe do not see any evidence of anomalous AlAs interfabonds. In contrast to the laser structures, in these struct

FIG. 9. XSTM image of InAs/AlSb superlattices grown witdifferent interfacial bonds.~a! mixed AlAs/InSb interfacial bonds a

seen on a~110! surface~22 V, 0.1 nA!. ~b! The (1̄10) surface ofthe same superlattice~23 V, 0.12 nA!. ~c! The ~110! surface of asuperlattice with InSb interfacial bonds grown using the sagrowth procedure nominally used for the laser structure sam~22 V, 0.1 nA!.

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the interfacial bonds are InSb-like, as expected, appearinthe XSTM image@Fig. 9~c!# as a slightly raised row of Sbatoms along the AlSb/InAs interface. Similarly, in XSTMimages of structures grown with shutter sequences intento create AlAs interfacial bonds~not shown!, a topographi-cally lower row of As atoms is observed at each interfaceexpected. What is surprising is that the AlAs-like interfacin that case look nearly the same as the InAs-on-AlSb infaces in the laser structure~cf. Fig. 8!, which were nominallygrown with InSb interfacial bonds.

To date, we have been unable to determine why weserve this apparently anomalous interfacial bonding inlaser structures. However, the presence of such AlAs-bonds at the InAs-on-AlSb interfaces would almost certaibe deleterious to the electronic quality of the material, dgrading the performance of the device. Previous workshown that As-like bonding is generally harmful for devimaterial.41 For instance, it drastically decreases the electand hole mobility in AlSb/InAs/AlSb channels and increasthe carrier concentration.42 The decrease in carrier mobilithas been attributed to increased interfacial roughnesAlAs-like interfaces,43 although in our material the InAs-onAlSb interfaces are not significantly rougher than the oth~e.g., the InGaSb-on-InAs interfaces!. The increased carriedensity has been primarily attributed to a higher concention of arsenic antisite defects near the interface.44 In addi-tion, the valence-band offsets have been shown to varyfunction of interfacial bond type, altering the electronstructure and further perturbing device performance.45,46

One alternate interpretation of the interfacial-bondianomaly is that, in this case, we are incorrectly characteing the ‘‘dark’’ rows as AlAs interfacial bonds. Perhaps thtopography is caused by charge depletion at the InAs-AlSb interface, somehow associated with the asymmetr

FIG. 10. High-resolution XTEM images (737 nm) of inter-faces with~a! III-As and ~b! III-Sb interfacial bonds. Images processed to reveal the composition-sensitive interface width areshown~Ref. 14!.

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MICROSCOPIC CHARACTERIZATION OF . . . PHYSICAL REVIEW B 63 245311

InGaSb-on-thin InAs-on-AlSb structure. If this were thcase, we might not see the same effect in simple~andthicker! AlSb/InAs superlattices. Clearly, resolution of thissue will require further investigation.

We have also used atomic-resolution HRTEM to invesgate the structural quality of superlattice layers deliberagrown with Sb-like vs As-like interfacial bonds~Fig. 10!.These samples had dramatically different optical propertwith the As-bonded samples exhibiting very low Peven when grown at the optimal temperature~Fig. 1!.After processing the HRTEM images to determine tinterface width,14 we find that the widths of the interfaceare the same in the two samples, 1.1–1.2 ML. In contto previous assertions,43 in our samples we see no evidenthat the III-As-like interfaces are rougher than the InSb-lones. It is possible that there are differences in the strucof these interfaces at longer length scales, but the lim~10–20!-nm lateral field of view for cleaved HRTEMsamples of optimal thickness makes this difficult for usdetermine. However, XTEM observations indicate that bsamples have dislocation densities,107/cm22. Therefore,we conclude that the degradation in PL associated wIII-As interfacial bonds is primarily an optoelectronic effeof the bond type.

FIG. 11. Nearly plan-view TEM images of cleaved~a! optimaland ~b! high-temperature samples. The samples were imathrough the cleaved edge of the sample. Misfit dislocations~whichend at the cleave face between the substrate and epilayer! appear toterminate below the epilayer, whereas threading dislocationsetrate through the epilayer. The curved morphology of the cleaedge in~b! may be due to strain effects.

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D. Meso-scale defects

We find that one additional source of material degradatoriginates at the mesoscale. As illustrated in Fig. 11, epilers grown in the optimum temperature range appear deffree in plan-view TEM images of cleaved samples, with tmisfit dislocations (,105 cm22)—caused by the AlSb buffelayer/GaSb substrate lattice mismatch~0.66%!—confined tothe buffer layer. In contrast, in samples grown at higher teperature, the dislocations (.106 cm22) thread through theepilayers, suggesting that the growth temperature exceethe elastic-plastic transition for some of the componentsthe superlattice. Although the presence of these dislocatis almost certainly deleterious to device performance, throle in the PL degradation is unclear. The electron and hpairs that recombine to produce the measured photonsbelieved to migrate<1 mm from the position where theywere generated before recombining, a distance somewshorter than the typical threading dislocation separatiTherefore, it would be very surprising if the threading dislcations are primarily responsible for the dramatic reductin PL intensity. However, we have identified no other copelling sources of material degradation that would accofor the PL data.

IV. SUMMARY

We have used XSTM and TEM to investigate the strutural quality of several InAs/In0.28Ga0.73Sb/InAs/AlSbstrained-layer heterostructures as a function of growth teperature~optimal vs high! and interfacial bond type~III-Sbvs III-As!. Implementing various methods for characteriziinterfacial roughness, we find that the roughness at thecial InAs/InGaSb interfaces is generally larger for the highgrowth temperature. We deduce from the surprisingly shcorrelation length of the interface profiles that the primasource of the roughening is most likely a direct result of tMEE technique used to form the interfaces. Even sampgrown at the optimum temperature show significant In intmixing at the AlSb-on-InAs interfaces. This intermixingaccompanied by anomalous AlAs bonding at the InAs-oAlSb interfaces; surprisingly, neither the intermixing nor tAlAs bonding is observed in InAs/AlSb test structures growunder nominally identical conditions. We find that the pmary source of material degradation in layers grown athigh temperature appears to be threading dislocationspropagate from the substrate-epilayer interface throughheterostructure.

This work highlights how XSTM and TEM can be applied in a complementary fashion to characterize III-V herostructures over a range of length scales. In particuwhen XSTM is performed on a$110% cleavage surface, thesame samplesof superlattice material can be analyzed usiboth techniques. The atomic and nanoscale structures~in-cluding point defects! are captured using XSTM, and thsparser mesoscale structures~such as dislocations and precipitates! can be easily identified using TEM. Although im

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BARVOSA-CARTER, TWIGG, YANG, AND WHITMAN PHYSICAL REVIEW B63 245311

aging of individual atomic columns is straightforward in HRTEM, the actual imaging of individual atoms~at least in thecross-sectional samples discussed here! is only achievable inXSTM. Accordingly, for samples that can be easily cleavin vacuum, XSTM provides a more complete analysespecially if detailed quantitative roughness spectra areinterest.

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ACKNOWLEDGMENTS

This work was supported by the Office of Naval Rsearch, the Air Force Research Laboratory, and an NNRL Postdoctoral Research Associateship~WB-C!. We alsothank Dr. B. Z. Nosho and Dr. J. R. Meyer for helpful dicussions.

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*Present address: HRL Laboratories, 3011 Malibu Canyon RMalibu, CA 90265.

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