Journal of Crystal Growth 254 (2003) 316–328

The effect of the wall contact and post-growth cool-downon defects in CdTe crystals grown by ‘contactless’ physical

vapour transport

W. Palosza,*, K. Graszab, K. Durosec, D.P. Hallidayc, N.M. Boyallc,M. Dudleyd, B. Raghothamachard, L. Caid

aUSRA/NASA-Marshall Space Flight Center, SD47, Space Science Laboratory, Huntsville, AL 35812, USAb IF PAS, Al. Lotnikow 32/46, 02-668 Warsaw, Poland

c Department of Physics, University of Durham, South Road, Durham DH1 3LE, UKdDepartment of Materials Science and Engineering, SUNY at Stony Brook, Stony Brook, NY 11794, USA

Received 23 October 2002; accepted 25 March 2003

Communicated by J.B. Mullin

Abstract

A series of cadmium telluride crystals grown by physical vapour transport without contact with the ampoule walls

and cooled at different rates were characterized using synchrotron X-ray topography, photoluminescence, and chemical

etching. Strain from sticking to silica glass and its effect on the dislocation density is shown. It was found that very fast

cool-down (e.g. air or water quenching) increases dislocation density by at least one order of magnitude. None of the

samples had random dislocation distributions, but coarse clumping of dislocations on the scale of more than 100 mm

was more prevalent in slowly cooled crystals. Photoluminescence revealed that slow cooling (e.g. 10�C/h) favoured the

donor–acceptor luminescence involving complex A centres. This was diminished in fast-cooled material, an effect

presumed to be due to dislocation gettering. Fast cooling also enhanced the formation of shallow acceptors.

Implications for Bridgman growth of CdTe and the vapour growth of CdZnTe are discussed briefly.

r 2003 Elsevier Science B.V. All rights reserved.

PACS: 61.10.Yh; 61.72.�y; 81.05.Dz; 81.10.Bk

Keywords: A1. Cool-down; A1. Defects; A1. Dislocation distribution; A1. Synchrotron white beam X-ray topography; B1. CdTe

1. Introduction

The quality of modern electronic devices de-pends critically on the quality of their semicon-ducting components. The quality of electronicmaterials is dependent, to a large extent, on theircrystallographic perfection. Dislocations, sub-grain boundaries, and other extended defects can

*Corresponding author. Tel.: +1-256-544-1272; fax: +1-

256-544-6762.

E-mail addresses: witold.palosz@msfc.nasa.gov

(W. Palosz), grasza@ifpan.edu.pl (K. Grasza),

ken.durose@durham.ac.uk (K. Durose),

mdudley@notes.cc.sunysb.edu (M. Dudley).

0022-0248/03/$ - see front matter r 2003 Elsevier Science B.V. All rights reserved.

doi:10.1016/S0022-0248(03)01183-7

act as trapping and scattering centres affecting theelectrical and optical properties of the materials.Distribution of defects may also be of importance,particularly with continuous progress in thesystems integration where smaller and smallerfragments of the crystal are responsible for aspecific function and even a single extended defectmay render the section useless. This may particu-larly be the case in optoelectronics where bulkcrystals are used as a substrate for functionalepilayers: the substrate imperfections may propa-gate into the epilayer and extend across its entirethickness. The quality of bulk crystals maydepend, among other factors, on their interactionswith the wall of the ampoule during and afterthe growth, and on the rate of the crystal cool-down at the end of the process. Both these factorslead to generation of strains and dislocations, andmay affect the distribution of defects in the finalcrystal.

Two important electronic materials, cadmiumtelluride and cadmium–zinc telluride, are usedin fabrication of g-ray detectors and IR focalplane arrays. We investigated the effect of post-growth cooling and interaction with silica (stick-ing) on concentration of dislocations and othercrystallographic defects, and their distributionin cadmium telluride. The crystals were character-ized using the synchrotron white beam X-raytopography (SWBXT) technique, low-temperaturephotoluminescence (PL), and chemical etching.The distribution of dislocations was assessed bycomparison to theoretically calculated randomdistribution using three different methodologies.

2. Experimental procedures

The crystals were grown by physical vapourtransport. The source materials were synthesizedfrom 6N elemental Cd and Te. A stoichiometricmixture of the elements (a total of 200–400 g) wassealed in a fused silica ampoule with 0.1 atm (atRT) of hydrogen. The synthesis was performed ina rocking furnace by a gradual increase of thetemperature up to about 1000�C over a period of3–4 days. The synthesized material was groundand sifted (mesh 80), then annealed at 900�C, first

in 0.5 atm of hydrogen and then under vacuum,and sealed in an ampoule with a frit separating thesource from the other part of the capsule. Thematerial was resublimed through the frit at about950�C. The resublimed material was loaded intothe growth ampoule and sealed with a smallamount of elemental tellurium and hydrogencorresponding to the pressure of 0.1 and0.05 atm, respectively, at the growth temperature.The crystals were grown without contact with theside walls of the silica glass ampoule [1]. Aschematic drawing of the growth ampoule isshown in Fig. 1a. Under the temperature gradient,the source material sublimes and is transportedtowards the cooler part of the ampoule where mostof it deposits on the pedestal. Under appropriatethermal field conditions, part of the materialtransported along the ampoule walls passesthrough the slit between the pedestal and theampoule wall without condensation leaving a gapbetween the outer wall and the growing crystal.To prevent polycrystalline growth and form ahigh-quality nucleus, the low supersaturationnucleation (LSN) technique was used [2,3]. Thenucleation process takes 2–3 days, the subsequentgrowth about 4 days. The source temperature was930�C, the under-cooling was a few degrees. Thecrystals, of diameter 25 mm, grew at the rate of afew mm per day. The post-growth cool-down toroom temperature was conducted at differentrates, and lasted from a few minutes to four days(c.f. Section 3). The ampoule after growth and thegrown crystal are shown in Fig. 1b and c,respectively. The crystals were sectioned length-wise and the central slices, about 3mm thick, wereused for characterization. Each sample wasmechanically lapped, then chemo-mechanicallypolished (2% Br in ethylene glycol) taking out0.3–0.5mm layer to remove the surface damage.The samples were then characterized usingSWBXT [4], low-temperature luminescence, andchemical etching. The dislocation (etch pit) densitywas measured and its distribution was analysed bycomparison with a Poisson distribution, by usingthe correlation function [5] and nearest neighbourdistribution analysis [6]. For the statistical analy-sis, dislocation fields measuring 0.5� 0.5 mmwere photographed and the dislocation (etch pit)

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328 317

positions digitized. FORTRAN routines werewritten for the numerical analysis.

3. Results

A total of five crystals have been grown and thefollowing cool-down conditions were applied: A,very fast cooling by dropping the hot (B900�C)ampoule into water; B, removing the ampoulefrom the furnace and letting it cool down in air; C,switching the power to the furnace off, that leadsto a drop from the growth temperature to 150�C in3 h; and D, controlled slow cool-down at the rateof 10�C/h from the growth temperature down toless than 100�C. In the fifth experiment, E, thegrowth process was conducted such as to allow thecrystal to touch the wall midway through thegrowth stage, and to cool down at the intermediaterate of 100�C/h. The crystals detached easily fromthe pedestal and showed a smooth, shiny surfaceindicating an absence of strong (chemical) inter-actions with the pedestal.

3.1. Characterization with SWBXT technique

The X-ray topographs were recorded from thesamples in the reflection geometry and are shownin Figs. 2–6. In these photographs, TB (twinboundary), SB (sub-grain boundary), C (crack),T (twin), P (precipitate), D (dislocation), and S (aslip band) mark different structural features/imperfections, and g the diffraction vector. Twin-ning operations are 180� type about the {1 1 1}plane normals.

Fig. 2 shows a topograph of the crystal cooled inair (B). Twin bands are revealed by orientationcontrast or through residual strain contrast at thetwin boundaries. The crystal has fine-scale sub-grain boundaries, and its dislocation density is inthe mid-105 cm�2 range. The side of the crystaladjacent to the pedestal shows a distortion of thelattice, the strain coming apparently from interac-tion (sticking) with silica.

Fig. 3 shows the crystal cooled in water (A) andits topograph. After cooling the crystal stayedintact, but during subsequent handling it cleaved

Fig. 1. Crystal growth. (a) Configuration of the growth system; (b, c), the ampoule after growth and the crystal, respectively.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328318

when tapped, as shown in Fig. 3a. The crystal is asingle grain with some bulk (T1, missing from thetopograph image due to orientation contrast) andlamellar (T2) type twins. The matrix regionexhibits a fairly high dislocation density (in theupper 105 cm�2 range) with sub-grain boundariesof 0.2–0.5� misorientation. No strain field in theregion adjacent to the pedestal is observed. Thisfact can be explained by high stresses developed inthe crystal during the very fast cool-down. Thestresses lead to formation of numerous micro-cracks in the crystals that are visible by opticalmicroscopy (Fig. 3b). The formation of the cracksapparently provided the mechanism for release ofthe stress, and prevented formation of a visiblestrain field in the region adjacent to the pedestal.

The topograph of the crystal cooled down in thefurnace by turning the power off (C) is shown inFig. 4. Some (long) sub-grain boundaries andprecipitates are present. Dislocation density islow, in the mid-104 cm�2 range, and individualdislocations can be discerned (Fig. 4). Heavydistortion is visible in the region bordering the

pedestal, and slip bands are observed emanatingfrom this region.

Similar lattice quality is observed in the crystalcooled down at 10�C/h (D-Fig. 5): some twins,sub-grain boundaries, a few precipitates, and slipbands with a strain field in the pedestal region arepresent. Dislocation density is low, and individualdislocations can be discerned.

The topograph of the crystal that was allowed totouch the ampoule wall (E) is shown in Fig. 6. Theresulting strain field produced severe distortion,seen superimposed on the image from other partsof the crystal. The sample is characterized by highdislocation density (in the high 105 cm�2 range)and sub-grain structures (0.1–2� misorientation) inmost regions.

3.2. Distribution of dislocations

The crystal wafers prepared as described inSection 2 were etched using FeCl3 solution [7] thatis known to give a good correlation to dislocationson a range of orientations of CdTe and CdZnTe.

Fig. 2. SWBXTs of the crystal cooled in air.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328 319

For each sample four areas of 0.5� 0.5 mm2 in sizewere selected for further assessment. The analysiswas made by comparison of the experimentalresults and those obtained by theoretical calcula-tions. The assessment of the dislocation distribu-tion was based on three methods: comparison witha Poisson distribution, by generation of a radialcorrelation function, and by calculating thenearest neighbour distribution parameter.

In the Poisson distribution method, the sampledarea is divided evenly into 100 squares and thefrequency of occurrence as a function of thenumber of dislocations in the square is obtained[5]. The experimental results are compared to thePoisson distribution function

PðnÞ ¼exp ð�lÞln

n!; ð1Þ

Fig. 3. Crystal cooled in water. (a) The cleaved crystal, diagram of the grain distribution, and white beam topographs; (b) optical

micrograph showing micro-cracks visible as linear features on either side of a grain boundary.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328320

where l is the mean value (average number ofdislocations). As the test of randomness, the w2

criterion was applied:

w2 ¼X

i

ðOi � EiÞ2

Ei

; ð2Þ

where O is the observed, and E is the expectedfrequency of occurrence.

To calculate the correlation function, thedistances between each point dislocation to allother points in the sampled area are calculated andthe frequency of the occurrence of points at a givendistance and distance interval (r and dr, respec-tively, in Fig. 7) is plotted as a histogram. Thisdistribution is normalized using a similar data setobtained from a computer-generated randomdistribution. Hence, if a random distribution istested, the value of the correlation function will be1.0 for all radii. Deviations above 1.0 indicateexcess correlation, while values less than 1.0indicate deficit levels of correlation.

The nearest neighbour distribution parameterRn is calculated using

Rn ¼ 2 %d

ffiffiffiffin

A

r; ð3Þ

where %d is the average nearest neighbour distance,n the sample size, and A the field area. If Rno1;the nearest neighbours are clustered; if Rn > 1there is ordering, while a value of Rn ¼ 1 indicatesrandomness.

The methods of the assessment of randomnessof distribution of the etch dislocations are demon-strated in Figs. 8a–c. Fig. 8a is the representationof the dislocations in the sampled area (obtainedfrom digitized images of the etched surface);comparison of the Poisson distribution and thecorrelation function are shown in Fig. 8b and c,respectively. The tails of the experimental data inFig. 8b, showing higher than random (Poisson)distribution frequency of occurrence for high(>10) and low (o5) pit numbers are due to thepresence of the dislocation clustering/sub-grainboundaries and of the depleted area between them,respectively. In the correlation function (Fig. 8c),the clustering of dislocations is reflected by thepeak at low radius. All samples analysed showsuch a peak and it is an indicator of the degree ofclustering with low dislocation separation such asin sub-grain boundaries. The tail-off is the con-sequence of a concentration of points in thecentre of the sampled area (Fig. 8a). Such coarse

Fig. 4. The grain distribution and white beam topographs of the crystal cooled spontaneously in the furnace; the cooling curve is

shown at the top left.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328 321

Fig. 6. Grain distribution and X-ray topograph of the crystal grown for half of the time in contact with the ampoule walls.

Fig. 5. Crystal cooled at the rate of 10�C/h. (a) Low-magnification X-ray topograph showing sub-boundaries (SB), slip bands (S), twin

bands (T) and grain boundaries (G); (b) high-magnification images in which dislocations (D) are resolved.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328322

clumping is indicated by peaks and troughs in thecorrelation function at high radii. For comparativepurposes in this work clumping can be said to besignificant when the correlation function fallsoutside the range 0.5–1.5.

The summary of the etch pit studies is given inTable 1. A clear relation between the cool-downprocedures and the (average) dislocation density(DD) exists: the DD values between the fast (Aand B) and the slow (C and D) cool-downexperiments differ by one order of magnitude,the slow cooling giving lower DD. Unexpectedlylow DD values are observed in the crystal cooledby turning the power off (C). Under suchconditions, the cooling is non-linear, and theinitial cooling is quite fast: in the first 15 min thetemperature drops by about 200�C. However,without a controlled power source the thermalgradients, both radial and axial, may be expectedto become smaller, reducing the lattice strains andgeneration of dislocations. DD values for experi-ment E are very high, apparently due to interac-tion of the crystal with the ampoule walls: crystalC was cooled at a much higher rate, particularly inthe high temperature range (Fig. 4), but its

dislocation density is one order of magnitudelower.

Comparison with the Poisson distributionshowed that no samples had dislocation distribu-tions that could be described by random distribu-tions. However, there was a clear relation betweenthe type of dislocation distribution observed andthe cooling rate: clustering was more apparent inthe slowly cooled samples than in the quenchedsamples.

For the fast-cooled samples (A and B), both thenearest neighbour distribution parameter andcorrelation functions indicated relatively weakassociation of dislocations at low radii comparedto the slow-cooled samples (C and D). Fig. 9shows a comparison of the low radius peak in thecorrelation function for all samples-C and D showsignificantly more of low radius correlation. Also,the slow-cooled samples show a greater tendencyfor coarse clustering and empty zones than dothe fast-cooled ones. This is revealed by deviationof the correlation function from 1.0 at largeradii, as is the case in Fig. 8c. Generally, theslow-cooled samples had correlation functions thatwere confined to a band of 1.070.2, while thosefor fast-cooled samples exceeded the limits1.070.5.

The sample cooled slowly, but that whichhad stuck to the ampoule walls (E) shows abehaviour that is intermediate between the twogroups. It has a high dislocation density and lowtendency for low radius dislocation association,but there is some coarse clumping and some emptyzones.

We can conclude that the differences betweenthe observed dislocation densities and distribu-tions are the result of the cooling schemes. Fastcooling (e.g. an air or water quench) introduceshigh dislocation densities and even cracks as aresult of thermal strain (Fig. 3b). Slow cooling(e.g. free or else controlled cooling in a furnace at10�C/h) reduces the dislocation density by anorder of magnitude. Neither fast nor slow coolinggives random dislocation distributions, but slowcooling affords greater opportunity for the dis-locations to become associated both at low radii(e.g. as in sub-grain boundaries) and high radii(e.g. coarse clumping).

Fig. 7. Illustration of the correlation function method of etch

pit distribution characterization in which correlations at radii

r þ dr are evaluated.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328 323

3.3. Photoluminescence

PL was excited using the 457.9 nm line of anargon ion laser with the samples cooled to 10 K.Typical excitation densities were 0.07 W/cm2.Based on the dominant features of the PL spectra,the crystals can be divided into three distinctcategories: samples C and D (slow cooled),samples A and B (fast cooled), and sample E(stressed by contact with the ampoule wall), shownin Figs. 10a–c, respectively.

Fig. 10a shows the PL emission from a slow-cooled sample (crystal D). PL emission is observed

in three distinct regions. The first is a near bandedge emission with peaks at 1.593 and 1.584 eV.These are readily attributed to D0X and D0htransitions, respectively [8,9]. We also observe veryweak LO phonon replicas of these transitions at1.571 and 1.563 eV. The second prominent featureis an eA0 transition at 1.549 eV, also with an LOphonon replica at 1.528 eV [10]. The third promi-nent feature is a DAP transition at 1.45 eV with apronounced series of phonon replicas to lowerenergy. This is due to recombination of a shallowdonor with a complex acceptor, most likely an‘‘A’’ type centre [11]. This centre has a strong

Fig. 8. Evaluation of the etch pit distribution. (a) Image of etch pit distribution obtained from experiments; (b) Poisson distribution

method: solid line-theoretical (random) distribution curve, solid circles-experimental results; (c) correlation function curve in which

deviation from 1.0 indicates excess (>1) or deficit (o1) correlation compared to a random field.

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328324

coupling to the lattice characteristic of suchrecombination centres with a Huang–Rhys cou-pling constant of 1.6 (hence giving the strongphonon replica series).

Fig. 10b shows the PL emission from a fast-cooled sample (crystal A). PL emission is observedin two distinct regions. The first is near band edge

emission with D0X and D0h peaks at 1.593 and1.584 eV similar to Fig. 10a. The second is a verypronounced eA0 emission band at 1.547 eV with aphonon replica at 1.526 eV. This is a differentacceptor centre to the one in Fig. 10a. Further-more, there is a weak shoulder on the eA0 peak inFig. 10b: possible evidence of a second acceptorspecies. The DAP features observed in Fig. 10a arenot present in this sample. There is a weaker PLband at 1.470 eV with a series of weakly coupledLO phonon replicas to lower energy. This is shownin more detail in the inset to Fig. 10b. This band isthe ‘‘Y’’ luminescence band and is characterized bya zero phonon line at 1.470 eV and weak phononcoupling, in our case with an S value of 0.2 [12].The ‘‘Y’’ luminescence band is believed to be dueto recombination of excitons at dislocations.

Fig. 10c shows the PL emission from the samplestressed by contact with the ampoule wall (crystalE). This shows PL emission in two distinct regions.The first is near band edge emission at 1.598 and1.586 eV. These correspond to the energy of FErecombination [13] and the D0h recombinationseen in the other samples. The D0h transition inFig. 10c is broader when compared to othersamples, possibly due to the strain that is itselfrevealed by X-ray topography.

Fig. 9. The average (of four locations) values of the correlation

ratio of the two lowest spacing distances for etch pit

distributions.

Table 1

Summary of dislocation density and distribution findings and main photoluminescence features

Crystal Cooling

method

Average

DD/

104 cm�2

Is the

distribution

random?

Short radius association Coarse

association

(clumping)

Photoluminescence summary

Nearest

neighbour

clustering

Correlation

at low radii

A In water 40.973.5 No Low Low Low Weak D0h and D0X; strong

eA0; Y luminescence present

B In air 18.470.8 No Low Low Low Weak D0h and D0X; strong

eA0; Y luminescence present

E 100�C/ha 23.371.4 No Low Low High Weak D0h and D0X; strong

eA0

C In furnaceb 2.870.4 No High High High Strong D0h and D0X; weak

eA0 DAP

D 10�C/h 5.071.0 No High High High Strong D0h and D0X; weak

eA0 DAP

aWith wall contact.bCooling was effected by turning the power off. Cooling from 930�C to 150�C was in 3 h (c.f. the cooling curve in Fig. 4).

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328 325

The concentration of active point defectsparticipating in the PL emission is stronglyinfluenced by cooling rate and strain. It is knownthat high dislocation densities give a reducedconcentration of species participating in DAPluminescence as a result of gettering of impurities[14]. This is consistent with our observations of areduced DAP emission—and associated increaseof eA0 emission—for the samples with increasedcooling rates: the high dislocation density favourseA0 emission at 1.55 eV over DAP emission at1.45 eV. From this, we infer that rapid coolingfavours the formation of isolated acceptor impu-rities (eA0), whereas slow cooling favours theformation of ‘‘A’’ type complex acceptor centrestypically involving a cadmium vacancy–impuritycomplex [11]. It is also worth noting that in themajority of vapour-grown CdTe crystals, the nearband edge PL emission is dominated by one of twoacceptor-bound exciton transitions. These occur at1.5896 and 1.5903 eV and are due to Cu or the Cl‘‘A’’ centre, respectively [8,15]. Absence of eitherof these PL features implies minimal concentra-tions of these two impurities in our samples. Yluminescence is only observed in the fast-cooledsamples with the highest dislocation density,consistent with its assignment to exciton recombi-nation associated with dislocations. Finally, it isworth noting two features of these spectra that areunusual for similar materials: the observation ofstrong D0h transitions and FE transitions, thereasons for the presence of these bands is unclear.

4. Summary and conclusions

Both interaction with silica and cool-down rateaffect the quality of the crystals. The strain fieldfrom interaction with the silica pedestal extendsover a distance of a few millimeters. In the case ofcontact with the ampoule wall (and the pedestal),the strain field can extend over a wholeboule leading to higher dislocation densities(B105 cm�2) than in the case of growth free fromthe walls (B104 cm�2). The apparent strongsticking to the wall may be caused by the rough-ness of the surface: the pedestal was flame polishedand had a very smooth surface, while such

1.30 1.35 1.40 1.45 1.50 1.55 1.60

0

1000

2000

3000

4000

PL

Inte

nsity

(co

unts

)

0

1000

2000

3000

4000

5000

6000

1.40 1.44 1.48

Energy (eV)

1.30 1.35 1.40 1.45 1.50 1.55 1.60

0

1000

2000

3000

(a)

(c)

(b)

D0X

D0h

eA0

DAP

DAP-LO

eA0-LO

Y

Y-LO1

Y-LO2

Y

D0X

D0h

eA0

eA0-LO

FE

D0h

eA0

Fig. 10. Representative PL spectra of the crystals. (a) Slow

cool-down (C and D); (b), fast cool-down (A and B); (c) crystal

grown with wall contact (E).

W. Palosz et al. / Journal of Crystal Growth 254 (2003) 316–328326

treatment was not feasible with the ampoule wallswhich might have had a rougher surface and, thus,attached stronger to the crystal. High concentra-tion of dislocations (105 cm�2 range) is also causedby very fast (air or water quench) post-growthcool-down. Dislocation densities lower by oneorder of magnitude or even more are present afterslower coolings (e.g. 10�C/h). High-quality crys-tals with low dislocation densities were obtainedeven after a relatively fast, natural cool-down withthe furnace (B900�C-150�C in 3 h). None of thecrystals had dislocation distributions that wererandom. While all of the crystals showed evidenceof short-range clustering of dislocations (scaleo20 mm), this was more pronounced for slowlycooled crystals. Similarly, coarse clumping ofdislocations (scale >100 mm) was more prevalentin the slowly cooled crystals.

No apparent effect of the cooling rate on thepresence of twins was found. This indicates thatthe twins form during the growth stage of theprocess, what is consistent with our earlierconclusions [16]. A presence of occasional pre-cipitates in slower-cooled crystals is consistentwith expectations: in our experiments the crystalsgrew under conditions close to the stoichiometriccomposition of the solid; therefore only a verysmall amount of tellurium could precipitate out atlower temperatures. Due to slow diffusion and lowconcentration of the excess tellurium, a longer timeis necessary to form a microscopic size precipitatein the crystal.

PL studies of the crystals revealed that thecooling rate and contact with the walls could berelated to the type and activity of optically activecentres in the CdTe. Slow cooling favoured theformation of complex acceptors (A centres) givingdeep DAP luminescence. Rapid cooling leads tohigher dislocation densities and the resultingreduction of the DAP luminescence was presumedto be due to gettering of the species responsible forthe dislocations. For samples with the highestdislocation densities, the dislocation-related Yluminescence was also observed. Fast cooling alsoencouraged the formation of simple shallowacceptor centres, which gave brighter luminescencethan that of shallow acceptors in the slowly cooledmaterial.

The present studies give the basis for somecomments on the likely effects of cooling onBridgman-grown CdTe. The more commonly usedBridgman method operates at a higher tempera-ture-about 1070–1090�C. Reduced hardness of thematerial at such temperatures may be expected tofacilitate the formation of dislocations. Undersuch conditions, the effect of cool-down rate onthe crystalline quality of the crystals may beexpected to be stronger than in our samples, andlower cooling rates than acceptable for our crystalsmay be needed. Slow cool-down may be expectedto enhance polygonization process as well. Furtherstudies, including growth by ‘contactless’ physicalvapour transport (to eliminate the overlappingeffect of sticking) at higher temperatures couldclarify those issues.

The effect of sticking and cool-down oncadmium–zinc telluride crystals may be expectedto be different than that for cadmium telluride. Onthe one hand, addition of zinc increases thehardness and, thus, a resistance to defect forma-tion. On the other hand, possible compositionalnon-homogeneities of the material lead to strainsand increased defect concentration [16,17]. Whenhomogeneous cadmium–zinc telluride crystals aregrown, even a rather fast spontaneous cool-downin the furnace and from a higher growth tempera-ture (about 1000�C) produced high-quality crys-tals with dislocation density in the low 104 cm�2

range [10]. Probably, the effect of reduced hard-ness at higher temperature was compensated bylattice hardening by the zinc component in thematerial.

Acknowledgements

Support of this work by NASA, Office ofBiological and Physical Sciences is greatly appre-ciated.

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