ZnSe-ZnMnSe and CdTe-CdMnTe superlattices

Surface

522

ZnSe-ZnMnSe

AND CdTe-CdMnTe

R.L. GUNSHOR,

L.A. KOLODZIEJSKI,

Science 174 (1986) 522-533 North-Holland, Amsterdam

SUPERLAITICES N. OTSUKA

and S. DATTA

Purdue University, West Lafayette, Indiana 47907, USA Received

29 July 1985; accepted

for publication

15 October

1985

We report the growth and characterization of superlattices of a new wide-gap, zincblende material system, Zn, _,Mn,Se. ZnSe exhibiting dominant free excitonic emission in photoluminescence (PL) is the well material, while wider-band-gap Zn, _,Mn,Se (0.23 i x < 0.66) forms the barrier material. PL measurements show greatly enhanced quantum efficiency compared to films of the ZnSe well material, while transmission electron microscopy shows extremely abrupt interfaces by the presence of seventh-order satellite spots. Previously reported superlattices in the CdTe-CdMnTe material system were grown with the (111) orientation, and exhibit unique excitonic properties believed related to the (111) interfaces. Using various techniques to select (111) or (100) heteroepitaxy of CdTe on (100) GaAs, we report the first (100) superlattices of this material. and compare the optical properties to the previous (111) structures.

1. ZnSe/Zn,

_ x Mn,Se

superlattices

The motivation for investigating the ZnSe/ZnMnSe material system in the form of both films and superlattices originated from the potential application as efficient blue light emitters and display devices. The challenge involved in growing thin films and superlattices of ZnMnSe stemmed from the tendency for bulk crystals of ZnMnSe to exhibit wurtzite phases for Mn mole fractions in excess of 9% Mn. This situation was crucial, as the small variation in band gap with Mn concentration made it necessary to grow barrier layers with high Mn fraction to achieve the required band offset for carrier confinement. In superlattice structures, high-quality ZnSe was intended for the well material; thus high-quality metastable, zincblende ZnMnSe was needed to isolate the ZnSe wells. Although Zn,_,Mn,Se (0.0 < x < 0.57) has been grown in bulk crystals, a predominance of the zincblende phase is only observed up to x = 0.3, whereas above this mole fraction a predominance of the hexagonal phase exists [l]. For the superlattices and epilayers grown at Purdue, only the zincblende (100) phase is present over the (0 < x < 0.66) composition range investigated; no hexagonal phases were observed [2]. In our attempts to increase the Mn fraction, we have succeeded in growing ZnMnSe films with up to 87% Mn; however, in this case both zincblende and hexagonal phases were observed with X-ray diffraction. 0039-6028/86/$03.50 0 Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division) and Yamada Science Foundation

R. L. Gunshoret al. / &Se-

5.850 1

-

-----‘-.-

ZnMnSe and CdTe- CdMn Te superlattices

523

ZINCBLENDE MKED PHAF?ES HEXAGONAL

6.76 (4.13)

5.800 -

6.71 ‘(4.08)

6.66 (4.03)

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

X Mn Fraction Fig. 1. Experimental lattice constant versus Mn mole fraction (x) Zn , .Mn .Se. (Hexagonal and mixed phases apply to bulk crystals.)

for MBE-grown

zincblende

Various epilayers of ZnMnSe and superlattice structures of ZnSe/ZnMnSe have been grown by molecular beam epitaxy (MBE) using a Perkin-Elmer model 400 MBE system. Details of the growth procedure have been previously reported [2]. Fig. 1 shows the lattice-constant variation with the Mn mole fraction for the Zn, _,Mn,Se epilayers. Also shown in fig. 1 is the lattice-constant variation for bulk crystals, emphasizing the change from a predominantly zincblende crystal structure to a hexagonal crystal structure occurring at a Mn mole fraction of 30% [l]. It is worth noting that these ZnMnSe epilayers have thicknesses of 1-3 pm, and thus represent metastable epitaxial layers. Extrapolating the plot of lattice constant versus Mn mole fraction to x = 1.0 indicates an expected zincblende lattice constant of 5.93 A; this contrasts with a reported value of 5.82 A for zincblende bulk crystals of MnSe [3]. Photoluminescence measurements were conducted to determine the variation of band-gap energy (E,) with Mn mole fraction, x. Fig. 2 shows E, versus x for the zincblende ZnMnSe epilayers [4]; for Mn fractions less than x = 0.1, data are also shown for zincblende bulk crystals [5]. Data obtained on bulk crystals consisting of primarily hexagonal crystal structure (x > 0.30) were found to be approximately 35 meV higher in energy [4]. As the Mn mole fraction increases the excitonic emission intensity decreases, while the emission

R.L. Gunshor et al. / ZnSe-ZnMnSe

524

md CdTe-CdMnTe

superluttices

2.80

t 2.70 L

I

0

0.1

0.2 X Mn

Fig. 2. Energy zlncblende Zn,

of dominant ,Mn,Se.

0.3

0.4

0.5

0.6

Fraction near-band-edge

feature

versus

Mn mole fraction

for MBE-grown

due to the excitation of Mn transitions (2.1 eV) increases. For x = 0.51 the Mn excitation dominates, making it difficult to find a near-band-edge feature [6]. In fact, for the ZnMnSe (x = 0.66) epilayer the band-edge feature was not found. Also, extrapolation of the Eg versus x data indicates an expected exciton emission energy of 3.36 eV at 6.5 K for zincblende MnSe. A series of superlattice structures involving the ZnSe/ZnMnSe material system have been grown [2]. In all cases ZnSe was used as the well material, and well dimensions ranged from 28 to 90 p\. The Zn, ~ ,Mn_,Se barrier material consisted D of various Mn mole fractions (0.23 < x < 0.66) with dimensions of 80-150 A. All superlattice structures consisted of 67 periods grown on a ZnSe buffer layer of 0.57-0.93 pm. In order to characterize the barrier material, ZnMnSe epilayers were grown on a ZnSe buffer layer of 0.25 to 0.44 pm and were typically l-3 pm thick. Transmission electron microscope (TEM) observations of cross-sectional specimens were performed with a JEM 200CX electron microscope. Fig. 3 is a dark-field image of a Zn, .Mn ,Se(x = 0.23)/ZnSe superlattice. These films were found to be highly susceptible to the radiation damage induced by the ion beam during sample preparation. Although a very low voltage was used at the final stage of ion thinning, fine structures due to the radiation damage are still seen in the image. No misfit dislocations are observed at interfaces between superlattice layers. The lattice mismatch is totally accommodated by elastic

R.L. Gunshor et al. / ZnSe-ZnMnSe

Fig. 3. Dark-field

TEM image of ZnSe/Zn,,,Mn

and CdTe-CdMnTe

,,$Se

superlattice

superlatrices

(well 63 A, barrier

525

104 A).

strain giving rise to strained-layer superlattices. Both dark-field images and diffraction patterns show the highly regular structure of the superlattice. No sign of disordering at the interfaces is seen in the image. The number of visible satellite spots around the (200) electron diffraction spot [2] is comparable to those of typical AlGaA-GaAs superlattices. Sixth- and seventh-order satellite spots are clearly seen in negative films; this observation suggests that the structural quality of these superlattices is comparable to those of well-developed III-V compound superlattices. In photoluminescence measurements at 6.5 K the ZnSe epilayer exhibited a free-exciton dominant peak at 2.799 eV having a full width at half maximum of 1.5 meV [2]. The observed two free-exciton features at 2.799 and 2.804 eV are attributed to strain splitting of the valence band [7]. (The origin of the strain is the small but finite lattice mismatch between film and substrate.) The feature at 2.796 eV is attributed to an impurity-bound exciton. In addition to the aforementioned exciton features, a broad emission centered at about 2.4 eV is observed with a peak intensity two orders of magnitude below the dominant free-exciton feature.

R. L. Gunshor et al. ,’ ZnSe-ZnMnSe

526

and CdTe-CdMnTe

.superlutrrw.\

Se-ZnSe Superlattice

2.747eV

9.49 Mno.51

2.17eV

FWHM f&

2

,A x 250

446

448

450

Fig. 4. Photoluminescence

452

454

of ZnSe/Zn,,,,Mn,,

545

565

585

605

625

645

,,Se superlattice.

As discussed above, the MBE ZnMnSe epilayers exhibit two photoluminescence features, a near-band-edge emission and a Mn emission originating from a Mn2+ transition. A comparison of the ZnSe/ZnMnSe superlattices to their corresponding ZnMnSe epilayers (which are grown with identical growth parameters as the barrier material used in the superlattices), dramatically illustrates a significant degree of carrier confinement. Fig. 4 shows the PL spectrum of the ZnSe/ZnMnSe (X = 0.51) superlattice. The near-band-edge emission is more than 250 x more intense than the emission originating from the Mn excitation. In contrast, the PL spectrum of the ZnMnSe (x = 0.51) epilayer was dominated by the emission due to the Mn excitation making it difficult to detect a band-edge feature. Similar behavior is seen for all the superlattices regardless of the barrier height present (i.e. mole fraction of Mn in the barrier layer) [2]. Photoluminescence intensity comparisons of the superlattice structures with the brightest ZnSe epilayer (described earlier) indicated that the band-edge-related emission was consistently much brighter for the superlattices. The measured lattice-constant variation with Mn mole fraction (fig. 1) for this material system imposes on the ZnSe well material an expansive hydrostatic strain parallel to the interface, and a compressive uniaxial strain normal to the interface. This strain situation results in a narrowing of the hand gap and a splitting of the valence band maximum. (The light-hole band moves up relative to the heavy-hole band.) Thus. although a blue-shift is expected due IO quantum confinement in the superlattices. the PL spectra of these ZnSe/ZnMnSe superlattices show a continual red-shift in energy as the barrier height is increased. The superlattices having well dimensions of 60- 90 A actually exhibit a net red-shift from the PL free-exciton energy of the ZnSe epilayer of 3-52 meV. The magnitude of the strain-induced shift can be varied by either changing the Mn mole fraction or the relative thickness of the well and barrier. For a superlattice with 28 A wells of ZnSe and 81 A barriers of

R.L. Gunshor et al. / .&Se-ZnMnSe

and CdTe-CdMnTe

superlattices

521

ZnMnSe (x = 0.3) a net blue-shift in energy (from the PL free-exciton energy of the Z&e) of 27 meV has been observed. Although the aforementioned superlattices have not been expressly fabricated for use as a laser structure, optically pumped stimulated emission and laser oscillations have been observed [8]. The lasing threshold for this ZnSe/ZnMnSe (x = 0.33) sample was 2.0 X lo5 W/cm2 at 5.5 K. (The superlattice consisted of 67 periods of 61 A ZnSe wells and 133 A ZnMnSe barriers.) At higher excitation densities, lasing was observed from 451.5 to 455 nm and occurred up to a temperature of 80 K. The fabrication of an idealized laser structure (by utilizing optimal growth parameters, including guiding cladding layers, and investigating various layer thicknesses and compositions) is expected to result in a decrease in the lasing threshold. These initial superlattice structures, however, still exhibit threshold powers which are an order of magnitude lower than those reported for bulk single-crystal ZnSe.

2. Heteroepitaxy

of CdTe on GaAs

Despite the existence of a large lattice-constant mismatch (14.6%) highquality single-crystal CdTe films have been grown by a variety of growth techniques on GaAs substrates. GaAs substrates provide an attractive alternative, rather than more closely lattice-matched substrates, for several reasons. One advantage is that large-area GaAs wafers of high structural perfection are readily available at modest cost. The large body of information characterizing the GaAs surface subsequent to various chemical and thermal etch treatments provides additional impetus for the use of GaAs wafers. As a result, research directed toward the use of alternative substrates for use in HgCdTe technology recognizes GaAs as a very viable candidate. Single-crystal CdTe films can be grown on (100) GaAs substrates with either of two epitaxial relations: (111) CdTe ]](lOO) GaAs or (100) CdTe ]](lOO) GaAs. (Numerous references are listed in ref. [9].) High-resolution electron microscope studies strongly suggest that conditions enhancing strong bonding between the epitaxial film and substrate tend to give rise to the (111) epitaxial relation, whereas factors which weaken the bonding interaction of the epitaxial film and substrate result in the (100) orientation. To consistently achieve a specific orientation, various growth methods may be employed which directly influence the film-substrate interface. For the (111) CdTe epitaxial films, the GaAs substrates were preheated at 580°C for 2 min prior to CdTe film growth to desorb the oxide: After preheating, reflection high-energy electron diffraction (RHEED) patterns showed the coexistence of (2 x 6) and (3 X 6) reconstructed surfaces which is characteristic of the transition state from the As-stabilized surface to the Ga-stabilized surface [lo]. The substrate temperature is then reduced to the

desired growth temperature of 325°C and the CdTe shutter is opened. Using this growth technique, the (111) invariably nucleates with the [211] of the (111) CdTe film always parallel to the toll], and never to the [OlT] of the (100) GaAs substrate. On the basis of this observation, we have suggested a model for the (11 l)CdTe-( 100)GaAs interface which involves chemical bonding considerations [9]. The [jll] of the (111) CdTe film has a 0.7% lattice mismatch with the [Oil] of the (100) GaAs substrate. With this directional orientation. the Te atoms bond to the dangling Ga bonds at the surface; the dangling bonds arise from oxide desorption in an As-deficient ambient. Additional support for the proposed model is found in the RHEED observation of the initial nucleation of the (111) CdTe film. The initial growth mechanism is observed to be two dimensional, thereby re-enforcing the idea of the existence of a strong epitaxial film-substrate interaction. Such a strong film--substrate interaction is only promoted in the case of a “clean” (completely desorbed oxide) GaAs surface. In the case of the (100) epitaxial relation. the interacti~~n of the epitaxial film and substrate must be inhibited in some manner to produce nucleation with the (100) orientation. Experimentally we have found various growth techniques which do inhibit the film-substrate interaction, thus providing a relative strengthening of the interaction between epitaxial nuclei in the depositing Iayer. This strong nuclei-nuclei interaction is also suggested by the observation of three-din~ension~l nucleation in the case of the (100) epitaxial orientation. The three-dimensional growth mechanism for nucleating the (100) film is observed by RHEED for a variety of growth sequences employed. Two specific types of growth sequences are used which provide a disturbance to the film-substrate interface. Using a conventional growth sequence, where the substrate is preheated, foltowed by reducing the substrate temperature to the growth temperature, results in the (100) orientation only in the presence of a residual 10 A oxide layer [9]. Alternatively a (100) film is nucleated if, after preheating the substrate at an elevated temperature (580°C). the shutters are opened, followed by reducing the substrate temperature to the growth temperature [ll]. In this latter growth sequence a (100) film nucleates in either the presence or absence of an oxide layer [12]; the initial nucleation of (100) CdTo on a 15 A residual oxide as observed by RHEED is identical to the sequence of RHEED patterns observed in nucleation on a completely desorbed oxide surface [12]. For both of these cases, the three-dimensional nucleation, prcsumably associated with a disturbance of bonding to the substrate. is observed. (The sequence of RHEED patterns for the completely desorbed case is shown in ref. fl2].) An additional technique for obtaining (100) CdTe on (100) GaAs involves the predeposition of a thin Cd, , Zn,,Te layer on a completely oxide-desorbed GaAs substrate [13]. Here the conventional growth sequence is employed where heating to 580°C is followed by a reduction to the growth temperature before source shutters are opened 1121.

R.L. Gunshor et al. / .&Se-ZnMnSe

and CdTe-CdMnTe superlattices

529

3. Cd, _ y Mn,Te / Cd t _ xMn xTe superlattices The first superlattices grown in the CdMnTe material system were of the (111) orientation [14,15]; these superlattices exhibited orders of magnitude higher photoluminescence efficiency than obtained from films and bulk samples [14]. The emitted light in photolu~nescen~e originates in the radiative recombination of interface-localized free excitons [16]. These somewhat unusual exciton properties are thought to be related to the (111) heterointerface. Once consistent selection of (100) or (111) epitaxy of CdTe on (100) GaAs was understood, it was possible to grow a series of CdTe/CdMnTe superlattices having interfaces parallel to the (100) planes [12]. The motivation for growing (lOO)-o~ented superlattices was to compare their optical properties to those of previously grown (Ill)-oriented superlattice structures. The superlattices of CdMnTe were grown in the same MBE system as the ZnMnSe films. A detailed description of the growth procedure has appeared elsewhere [14]. A number of superlattice configurations have been fabricated having (111) and (100) orientations. For the (111) orientation, both CdTe and Cd,_,Mn,Te (X = 0.08) have been used for the well material, with various Mn mole fractions in the Cd,_,Mn,Te (0.30 < x < 0.50) barrier material [14,17]. The well dimensions have ranged from 50 to 650 A, with barrier dimensions ranging from 50 to 650 A. These superlattice structures were typically grown on l-2 pm CdMnTe buffer layers on (100) GaAs substrates. For superlattices having the (100) orientation, only CdTe has been used for the well material; the wells are isolated by Cd 0,-,6Mn,,,Te barrier layers. The CdTe well dimension was varied from 57 to 444 A, whereas the barrier dimension ranged from 84 to 512 A. These (100) superlattice structures were grown on 2 pm of Cd,,,Mn,,z,Te deposited on a CdTe buffer on (100) GaAs. The (100) CdTe buffer layer was nucleated on the (100) GaAs substrate using the growth techniques discussed above. In the CdMnTe material system, the sense of variation of lattice constant versus Mn mole fraction is opposite to that for the ZnMnSe system. In the case of CdMnTe the lattice constant decreases as the Mn fraction increases [18]. As a result, the strain subjects the well material to expansive uniaxial strain normal to the interface, and compressive hydrostatic strain parallel to the interface. (The heavy-hole band now moves up in energy relative to the light-hole band.) TEM observations have been made for both (lOO)- and (Ill)-oriented superlattice structures. Fig. 5 is a dark-field image of a (100) superlattice structure grown on (100) GaAs. In this case, two SL structures having different periodicities were grown on the same buffer. The CdTe wells had dimensions of 87 and 444 A with Cd o,,aMnO,,,Te barriers of dimensions 200 and 512 A, As with the ZnMnSe material system, the CdMnTe material suffers from radiation damage incurred from ion milling during the sample preparation for

530

R.L. Gunshor et al. / ZnSe-ZnMnSe

and CdTe-CdMnTc

.wperluttice.s

-Fig. 5. Dark-field

TEM image of CdTe/CdO~,,MnO~,4Te

-

-

_

(100) superlattice.

cross-sectional TEM. Previous dark-field images showed that a large number of defects are caused by the ion bombardment [17]. High-resolution images of the superlattice interfaces were observed, but no discontinuity of lattice fringes were found at these interfaces. The implication is that the lattice mismatch between superlattice layers is accommodated by elastic strain rather than by misfit dislocation networks, resulting in strained-layer superlattices. All samples have shown highly regular superlattice structures in both images and diffraction patterns. The sharpness of the interfaces between (100) superlattice layers is similar to (or better than) that of (Ill)-oriented superlattices. Typical near-surface dislocation densities in the (lOO)-oriented superlattice structures (with total film thicknesses of = 3 pm) are lO’/cm’. The (111 )-oriented superlattice structures of similar total thickness had dislocation densities an order of magnitude higher (108/cm2) [14]. It is unclear whether the reduction in dislocation density is related to the epitaxial orientation. or due to the additional interface between the CdMnTe and CdTe buffer layers present in the (100) case. Extensive measurements of photoluminescence over a range of temperatures down to 1.8 K have been performed on the CdMnTe superlattices and films. A general observation for the (111) orientation is that the quantum efficiency of the superlattices greatly exceeds that of the films. Optical measurements in a magnetic field performed on superlattice structures having CdTe wells exhibited Zeeman splitting orders of magnitude greater than expected from CdTe [16]. This experimental observation provided an important clue, eventually resulting in the conclusion that the emitted light results from interface-localized excitons [16]. The large Zeeman splitting originates in the alignment of the magnetic moments of the Mn ions present in the barrier layers. This opportunity to diagnose the wavefunction localization with the application of a magnetic field is unique to diluted magnetic semiconductor material.

R.L. Gunshor et al. / ZnSe-ZnMnSe

and CdTe-CdMnTe

superlattices

531

b

I

1.65 Energy

1.70

CeV 1

1.65 Energy

I

1.70

I eV 1

Fig. 6. PL spectra for a (100) superlattice (CdTe well 57 A, Cd,,,Mn,,,,Te (111) superlattice (CdTe well 71 A, Cd0,76Mn,,24Te barrier 128 A).

barrier

94 A) and a

As an indication of the high quantum efficiency, the (111) SL were compared to an (Al, Ga)As double heterostructure grown by LPE for use in an injection laser [17]. It was found that the II-VI superlattice was somewhat brighter in photoluminescence than the III-V sample at 77 K. Further evidence of the optical quality of this material system is supported by the recent demonstration of stimulated emission [19]. All superlattices reported heretofore in this material system have had interfaces normal to the [ill]. Photoluminescence measurements (Ar laser focused to a 100 pm spot) indicate that the optical properties of the (100) superlattices are different from those having a (111) orientation [12]. The (100) SL exhibited a similar quantum efficiency as the (111) SL. However, the spectral shape for the (100) SL is notably different (fig. 6), consisting of a narrow peak and a distinct low-energy shoulder. The amplitude ratio between the two features is dependent on excitation intensity (1 to 10 mW), with the narrow higher-energy peak dominating at high excitation levels. A pronounced difference in the excitation spectra for the (100) and (111) SL was observed. Such comparisons suggest that the differences in the heterointerface for the two orientations play a significant role in determining the optical properties. Similarities in superlattices of the different orientations are that the PL spectra for both the (111) and (100) orientations did not exhibit the large defect band at 1.4 eV which is always present in thin films and bulk samples. Also the measured PL band-edge emission was shifted to higher energy for both orientations due to a combination of quantum-confinement and strain effects.

532

R.L. Gunshor d ai. / ZnSe-ZnMnSe

and CdTe-CdMnTe

supe&~trrces

Acknowledgement The authors wish to emphasize that this paper represents the contributions of several research groups. T. Sakamoto of the Electrotechnical Laboratory, while a visitor at Purdue, contributed greatly to the early growth of DMS films by MBE. At Brown University, Professor A.V. Nurmikko and his co-workers X.-C. Zhang, S.-K. Chang, Y. Heftez, and J. Nakahara have contributed extensive optical measurements and much of the insight into the excitonic behavior of the II-VI superlattices. Professor M. Yamanishi of Hiroshilna University contributed greatly to our understanding of optical properties of the superlattices. We would like to recognize collaborators Professor J.F. Schetzina and R.N. Bicknell at North Carolina State for their help in identifying the importance of a residual oxide in determining the epitaxial orientation of CdTe on GaAs. At Purdue University, Professor W.M. Becker, R.B. Bylsma, and T.C. Bonsett have been responsible for photoluminescence measurements and obtaining stimulated emission. We would like to thank U. Debska for providing the high-purity, vacuum distilled source materials. We also thank M. Udo, C. Choi, R. Venkatasubramanian, and R. Frohne for their contributions to this project. L.A.K. would like to thank IBM for fellowship support. This research was supported by Office of Naval Research Contract NO~l4-82-KO563, Air Force Office of Scientific Research Grant 83-0231, and NSF-MRL Grant DMR-83-16988.

References [l J D. Yoder-Short and J.F. Furdyna, private communication. [Z] L.A. Kolodziejski, R.L. Gunshor, T.C. Bonsett. R. Venkatasubramanian. S. Datta, R.B. Bylsma, W.M. Becker and N. Otsuka, Appl. Phys. Letters 47 (1985) 169. [3] A. Pajaczkowska, Progr. Crystal Growth Characterization 1 (1978) 289. [4] R.B. Bylsma and W.M. Becker. unpublished. [5] A. Twardowski. T. Diet1 and M. Demianiuk, Solid State Commun. 48 (1983) 845. [6] Y. Heftez, J. Nakahara. A.V. Nurmikko, L.A. Kolodziejski, R.L. Gunshor and S. Datta, to be pubIished. ]7] T. Yao, paper presented at the II-VI Intern. Conf.. Aussois. March 1985. [8] R.B. Bylsma, W.M. Becker. T.C. Bonsett. L.A. Kolodziejski, R.L. Gunshor. M. Yamanishi and S. Datta. to be published. [9] N. Otsuka, L.A. Kolodziejski, R.L. Gunshor. S. Datta. R.N. Bicknell and J.F. Schetzma. Appl. Phys. Letters 46 (1985) 860. [IO] A.Y. Cho, .I. Appl. Phys. 47 (1976) 2841. [I l] W.J. Schaffer, private communicati~~n. 1121 L.A. Kolodziejski, R.L. Gunshor, N. Otsuka. X.-C. Zhang. S.-K. Chang and A.V. Nurmikko. to be published. [13] J.P. Faurie. talk presented at the Focal Plane Array Materials Devices and Processing C‘onf.. McLean, Virginia. April 1985. [14] L.A. Kolodziejski, T.C. Bonsett, R.L. Gunshor, S. Datta. R.B. Bylsma. W.M. Becker and N. Otsuka. Appl. Phys. Letters 45 (1984) 440.