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Epitaxial technologies in the growth of semiconductors for device applications
Materials Chemistry 4 (1979) 529 - 548 © CENFOR S.R.L. - Printed in Italy
EPITAXIAL TECHNOLOGIES IN THE GROWTH OF SEMICONDUCTORS FOR DEVICE APPLICATIONS*
S. F R A N C H I * , D. M A R G A D O N N A * * , *
CNR.MASPEC,
**
ASSORENI,
PARMA
a n d C. P E L O S I *
- Italy.
MONTEROTONDO
(Roma)
- Italy.
Abstract -- One of the most promising trends in device development foreseen for the near future is the incorporation of multilayer structures in all devices in which performance improvements justify the higher technological costs. In this perspective, the most diffused epitaxial technologies, which are Liquid Phase Epitaxy (LPE), Vapour Phase Epitaxy (VPE), Pyrolysis of Organometallic Compounds (OMP) and Molecular Beam Epitaxy (MBE), are discussed with respect to a number of definite problems. These topics are related to: 1) Interfaces; 2) Composition profiles; 3) Doping profiles and 4) Stoichiometry. Although the present review refers mainly to the III-V compounds and alloys, some attention has been paid to outlining the general aspects of the above problems. Moreover some comments are made on the possible developments and industrial impact of the epitaxial technologies.
INTRODUCTION A great n u m b e r o f e l e c t r o n i c devices w h i c h are c u r r e n t l y p r o d u c ed b o t h in industrial a n d in r e s e a r c h l a b o r a t o r i e s are f a b r i c a t e d using e p i t a x i a l t e c h n o l o g i e s . O n e r e a s o n f o r this is t h a t , generally speaking,
* Presented at the IV Scientific Meeting of the Italian Association for Crystal Growth (AICC), Parma, Italy, 26-28 February 1979.
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epitaxial materials have better properties than the bulk ones, as far as structural and electrical properties are concerned. This is mainly due to the fact that in epitaxy, independently of the particular method used, the growth is carried out at temperatures which are several hundred degrees lower than the melting point of the compound or alloy. The lower temperature growth accounts, at least in part, for the concentration both of chemical and crystalline defects, which is lower in epilayers than in materials grown from a nearly stoichiometric melt. At the same time, the lattice vacancy concentration is expected to decrease. Multilayer structures have gained an increasing interest both in opto-electronic and in microwave devices; the importance of double heterostructures in lasers and LEDs, and of single heterostructures in solar cells and in light detectors is well known. In the microwave field, the interest is presenly focused on homostructures in which the layers have different doping levels, and devices have successfully been fabricated which incorporate three or four layers. However, due to the intrinsic properties of some III-V alloys or compounds other than the usual GaAs, it is conceivable that in the near future heterostructures will be considered even in this area. Many papers have reported that in devices, such as injection la-" sers, Gunn diodes, IMPATTs, MESFETs and mixer diodes, the thickness of the active region should be in the sub-micron range, as it will be discussed below. The above considerations clearly underline the prominent role of the epitaxial processes in the semiconductor technology, even if, in particular cases, the advantages of the ion implantation technique should be taken into account. The more widely used epitaxial technologies are: a) Liquid Phase Epitaxy (LPE); b) Vapour Phase Epitaxy (VPE) and Pyrolysis of Organometallic Compounds (OMP); c) Molecular Beam Epitaxy (MBE); in the following no discussion of apparatuses and processes relative to the above techniques will be made, since a number of excellent review paper on LPE 1, 2, VPE 3 ' and MBE 4' s are available. Considering the present situation, none of the above technologies
531
can be thought as being the most suitable for a broad range of experimental situations, even if each one of them may have distinct advantages in specific cases. In the following the epitaxial technologies will be discussed with respect to specific device problems. These topics are connected with: a) Interface-related problems caused by lattice mismatch; the achievement of: b) doping and c) composition profiles in the growth direction, and d) the control of the stoichiometry of epitaxial layers. Although the present review refers to the III-V compounds and alloys, some attention has been paid to outlining general aspects of the above problems.
INTERFACE-RELATED PROBLEM
In many of the currently produced minority carrier devices, such as, for instance, semiconductor lasers, LEDs and solar cells, the active region is within the minority carrier diffusion length from homo- or hetero-epitaxial interfaces; consequently defects occurring at heterointerfaces strongly influence the properties of the above devices. In this section, the lattice mismatch is considered as a source of interfacial defects. Frank and Van der Merwe have shown that, in order to relieve the elastic strain due to the lattice mismatch at interfaces, a misfit dislocation network is almost always generated6 ; simple considerations indicate that the misfit dislocation linear density is proportional to the lattice mismatch parameter e = Aa/a. Assuming that a nonradiative recombination center is associated to each a t o m having a dangling bond, it follows that the surface density of these centers, as well as the surface recombination velocity, are linearly related to e 7 . The above model is somewhat oversimplified since the plastic relaxation may not be complete and, consequently, part of the lattice mismatch may result in elastic strain; this has been confirmed by Matthews 8 who has shown that a critical thickness exist, below which the epitaxial layer is under strain and no dislocation network is gen-
532 erated. The above thickness has been evaluated to be about one half the reciprocal of the linear dislocation density. The morphology of the misfit dislocation network has been studied in a number of heteroepitaxial structures (InGaP/GaPg, 1 o, 1 1, InGaAs/GaAs 12 , GaAsP/GaP la ) as well as in homoepitaxial interfaces 14, in which the mismatch is caused by doping level differences. Recently the relationsh!p between interfacial misfit dislocations and device opto-electronic properties has experimentally been investigated. Ettemberg et al. have shown that, both in GaAs/GaP 1 s and in InGaAs/ GaAs 1 2 heterostructures, the minority carrier diffusion length h in the epitaxial layer is upper-limited by the mean spacing between misfit dislocations. This result is best understood while considering that, in the first approximation, the luminescence internal quantum efficiency is proportional to X2. Detailed studies of the dependence of the surface recombination velocity S at interfaces on the lattice mismatch have been made using InGaP/GaAs heterojunctions along with different experimental techniques 16, 17. The main result is that S is proportional to the lattice mismatch, and the coefficient is surprisingly close to the one estimated on the basis of the above model. These data definitely confirm that misfit dislocations act as non-radiative recombination centers, and limit the minority carrier diffusion length; consequently the effect of lattice mismatch in devices such as lasers, LEDs and solar cells may be devastating. Within this context, the important role of GaA1As/GaAs structures in opto-electronics should be recalled; since the lattice mismatch, in general, is within the 10 -4 range, this interface is void of misfit dislocations t a and has a surface recombination velocity evaluated to be 450-+ 100 cm/s19: this corresponds to a recombination center concentration of 4 • 109 cm "2 , which compares favourably with that obtained by Lang et al. using the deep level transient spectroscopy 2°. The experimental evidence, both in GaA1As/ GaAs and in InGaP/GaAs structures, definitely points out that, in general, for S ~< 104 cm/s, the lattice mismatch should be within the 10 -4 - 10 -3 range. In the case of epilayers in compression on the substrate, the tel-
533 atively low surface recombination velocity at the active region interface is obtained by using a number of lattice-matching layers between the substrate and the active region itself: by step-grading the composition of the layers, the lattice constant is progressively matched to the active region one. The main results of this technique are that: the misfit dislocations are localized at the interfaces where the composition is step-graded, and the inclined dislocations propagating from the substrate, or generated within the epitaxial layers, are bent into the interfaces z 1, 22. In this way, a dislocation density of 100 cm "2 could be obtained in the constant composition region 2a On the other hand, in cases in which the epilayer is in tension, tile above technique does not give good results, due to the formation of microcracks in the dislocated epitaxial region. In this situation, a continuously graded layer is grown between the substrate and the active region; consequently, the strain-relieving dislocations are generated with a density which is proportional to the composition gradient and, then, propagate through the device I 3. However, if low enough composition gradients are used, the dislocation density is minimized, so that device-quality materials are obtained. In the following, the epitaxial technologies are compared in so far as the handling of the growth of lattice-matched structures is concerned. With the noteworthy exception of GaA1As/GaAs structures, the epilayer composition must be kept under close control; for example, in order to have a lattice mismatch of 2 • 10 -4 in the InGaP/GaAs heterostructure, the epitaxial layer must be grown with an In molar fraction that is 0.496 ± 0.003. It has been shown that VPE lattice-matched structures based on III-V compounds can be reproducibily grown, either when the epilayer is in tension or in compression. A very brilliant example of the VPE uses is the growth of InGaAs/ InGaP/GaAs and GaAsP/InGaP/GaAs heterostructures, from which lasers operating (under RT and CW conditions) within the 1.0-1.3/am 24 and 0.7 jam 2 s spectral regions respectively, have been successfully successfully fabricated. In the case of LPE, some problems arise since the number of dif-
534 ferent layers is limited by the number of bins in the multibin graphite boat, and the growth of continuously graded layers is intrinsically difficult. Anyway, in a few cases, continuously graded matching layers have been obtained, taking advantage of the phosphorous depletion in solutions based on the GaAsSbP 26 and InGaAsP 27 quaternary systems; however, the same constituent source-depletion may be detrimental when very constant composition layers are required. In LPE the precise control of the supersaturation has an important role, expecially when the epitaxial layer is required to have a composition which is different from that of the seed. Infact, in these cases even a slight undersaturation causes the partial dissolution of the substrate and then the composition of the growing layer is different from the desired one. At any rate, LPE has features which are very interesting: firstly, it has been reported that in the case of perfect lattice matching, the epilayer has a dislocation density which is about one order of magnitude less than that of the substrate; the second characteristic is that, at least in the case of InGaP/GaAs 2s and InGaAsP/GaAsP 29 , over a certain range of liquid composition the LPE layers all have a growing composition which is lattice-matched to the substrate, rather than conforming to the bulk phase diagram. This could be explained by the contribution of dislocation- and elastic-stress-energy to the free energy of the system. Recently, however, some doubts have been raised about the occurrence of the above latching in InGaAsP/InP heterostructures 30 MBE is considered to be a very versatile technology, even in connection with interface-related problems: examples of the growth of continuously graded epilayers in GaAsP 3 i, inGaAs 31 and GaAsSb 32 alloys have been reported; furthermore, very interesting results have been obtained in the growth of InGaAs/GaAsSb heterostrnctures and superlattices, where the lattice-matching conditions have been obtained by carefully controlling the composition of both the ternary solid solutions 3 3, 3 4
535 DOPING PROFILES
The obtaining of precisely controlled doping profiles is a major' problem in the epitaxial growth of structures suitable for microwave devices. For example, in order to reduce the noise figure in Gunn diodes, a cathode-notch doping profile has been proposed3 s, which can be described in terms of a n ÷ / n ' / n / n ÷ homoepitaxial structure. Both theoretical 3 6, 3 7 and experimental studies 3 8 , 3 9 show that higher DC to RF conversion efficiencies can be achieved by using a low-high-low doping profile in GaAs IMPATT diodes. Although most of the above microwave devices are presently prepared by the halide VPE process, interest has been shown for the LPE method, since LPE grown n-GaAs has an electron trap concentration which is in the low 1012cm'3 range, that is generally lower than the VPE one. However, with LPE, it is difficult to obtain layers with precisely controllable thickness in the sub-micron range, due to the problems of experimentally controlling the nutrient solution supersaturation. A number of methods have been proposed for circumventing this problem, such as using one 40 or two 41 d u m m y wafers which precede the substrate and equilibrate the solution, or using a separate source wafer on the top of each melt to maintain the equilibrium throughout the entire growth cycle 42, 43. A different approach for controlling the active layer thickness is the anodic oxidation of the layer itself 44, however this m e t h o d holds only when the layer to be thinned is not sandwiched in multilayer structures. In any case, none of the above techniques has given completely satisfactory results. For obtaining very controllable doping profiles, the problem of the cross-diffusion both of constituent and dopant elements across the interface is important, expecially when very thin layers must be grown on films or substrates which have different doping or composition. Since the diffusion of atoms through a solid is a thermally activated process, it is immediately evident that the growth, and more generally, the processing temperatures are to be kept as low as possible, and this may constitute a prerequisite in the choice of the epitaxial technology. For example, the typical temperature ranges for growing
536 GaAs are 850-900 C, 700-750 C, and 500-440 C in LPE, VPE and MBE respectively. The cross-diffusion related problems are best discussed in connection with the fabrication of GaAs MESFETs, in which a 0.2-0.3/am thick n-doped layer is epitaxially deposited on Cr-doped semiinsulating (SI) material. Low noise MESFETs generally operate at a bias-voltage so that the depletion region is only a few hundred angstroms far from the active] SI-layer interface; the power gain has been shown to be proportional to the majority carrier mobility in the active region. In VPE devices grown directly on SI-GaAs, the I-V characteristics show a typical hysteresis, moreover the electron mobility was shown to degrade drastically towards the interface. These results have been interpretated in terms of deep levels, due to the Cr-diffusion from the substrate into the epilayer, which act as efficient electron traps 4s, 46. In order to avoid these negative effects, an intrinsic low-conductivity buffer layer has successfully been interposed between the substrate and the active region, whose function is to limit the diffusion of Cr into the upper layer. However, Wood has shown that in MBE GaAs MESFETs no degradation of the majority carrier mobility, nor looping in the I-V curves is present, even in the absence of the buffer layer, thus indicating that the Cr-outdiffusion is not a major problem at MBE GaAs growth temperatures 47 Recently, a halide VPE process has been described, in which the growth temperature has been reduced to 630 C4s; however, even if these homostructures have the MESFET-quality without any buffer layer, the epilayer majority carrier mobility is lower than that of MBE samples. Since many silicon devices, such as microwave diodes and transistors, are fabricated by depositing a lightly doped ('~ 1016 cm-3) epilayer on a heavily doped substrate ('~ 1018 cm-3 ), the same sharpness requisites in doping profiles exist, which have already been discussed in connection with III-V devices. Ota has recently used MBE in preparing n-Si/n*-Si structures in which the n-layer Sb-doping level is controlled within the high 1013_ high 1017cm'3 range 49 . To investigate the outdiffusion of the As-dopant from the substrate (n "~ 5 • 1019
537 cm "3) into the epilayer, SIMS profiles of MBE and VPE samples have been measured. The results show that in MBE grown at 800 C for 3 hr, and 1050 C for 2 hr, the As-outdiffusion width is 0..03 and 0.()2 lain respectively49; these values are definitely lower than those generally obtained (~> 1/am) using SiCI4 VPE at 1200 C. In connection with the trend of reducing the epitaxial temperature, the Solid State Epitaxy (SSE) could be potentially interesting s° The autodoping could hamper the achievement of precisely controlled doping profiles in VPE; this p h e n o m e n o n is related to the existence of a stagnant layer, just around the substrate, which acts as a sink for impurities, whether liberated from the substrate or contained in the main gas stream. Both intentional or unintentional dopants are, therefore, available for incorporation in the growing layer for a time determined by the stagnant layer volume. Different methods have been proposed for eliminating this problem, the simplest being the use of long purge times between growth intervals. A quite different approach, consisting of the reduction of the epitaxial reactor pressure, has recently put forward. This m e t h o d has been used both in Si s I and in GaAs OMP s 2 epitaxy. In the first case Sill4, SiHC13, SiH2C12 and SIC14 have been used as Si- sources in a cold wall horizontal reactor. The operating pressure has been mantained in the 10-100 torr range by means of a Roots pump fitted with a liquid nitrogen trap. Defectfree layers have been grown at temperatures lower than 900 C and doping profile has been shown to be very sharp, even at the interface with a highly Te-doped substrate (/9 "" 10"3Ohm cm) sl In the GaAs OMP epitaxy, good crystalline-quality layers have been grown within the 520-750 C temperature and 25-600 torr pressure ranges s2. The carrier concentration reduction by about two orders of magnitude from the 6 • 10 -4 Ohm • cm Te-doped substrate towards the epitaxial layer has been achieved in about 0.05 /am using growth temperatures as high as 750 C. This result definitely suggests that the autodoping can be effectively reduced, while the reactor is used at sub-atmospheric pressure, and that this is independent of which VPE m e t h o d is actually being used.
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COMPOSITION PROFILES The growth of heteroepitaxial structures is very important for obtaining high efficiency opto-eletronic devices such as injection lasers, LEDs, solar cells and photodetectors. For instance, it has been shown that double heterostructure lasers (DHL) have a threshold current density which is markedly lower than those of homojunction lasers, and that this is essentially related to the improved confinement both of injected carriers and of electromagnetic field in the active region s3. As for solar cells, p-GaA1As/p-GaAs/n-GaAs structures have a significantly higher conversion efficiency than GaAs homojunctions: this is understood by recalling that GaA1As/GaAs interfaces have a surface recombination velocity 2-3 orders of magnitude lower than the vacuum/GaAs one, and that half the carriers generated by AM1 radiation are produced in GaAs within 0.2/~m. Ternary or quaternary III-V alloys are used in the active region of opto-electronic devices operating either within the yellow or red regions of the visible spectrum, or within the 1.0-1.3 #m spectral band. Since for pratical reasons the bulk materials are restricted to binary compounds, such as GaAs and InP, lattice-matching-layers of suitable composition can be used between the substrate and the active region. The possibility of using epitaxi~ techniques for growing multilayer structures incorporating different materials has already been discussed in connection with the problem of lattice-matching layers. However, due to the technological importance of Al-containing alloys, such as GaA1As, some further comments on the growth of these solid solutions are in order. Because of the reactivity of A1C1, very poor results have been obtained by VPE using conventional silica reactors. In order to circumvent this difficulty all-allumina VPE reactors have been used 56 to grow n-A1As/p-GaAs heterostructures, which have a sea-level sunlight conversion efficiency of 13-18% 57. Although these results should be compared with those of LPE devices (18-21%), it is noteworthy that VPE cells have significantly larger areas. It has recently been shown that GaAIAs/GaAs heterostructures grown by OMP have
539 been incorporated in highly efficient, large-area solar cells s s, s 9 and in DHLs operating under CW-RT conditions 60, with a threshold current density as low as that of the best LPE GaA1As/GaAs DHLs 61, 62 These results clearly show that OMP is a technique suitable for producing highly efficient solar cells and lasers on a large scale. The growth of really good GaAIAs epilayers is a major problem for MBE: infact, the MBE GaA1As has a photoluminescence efficiency which usually is 10-20 times lower than that of the best LPE layers. This is thought to be related to the presence of large concentr. tions of non-radiative centers, which are caused by residual CO in the ultra-high-vacuum growth environment. This unintentional doping is a problem only for the growth of Al-containing solid solutions, and this is understood by recalling that the CO sticking coefficient is much higher in GaA1As than in GaAs s . Detailed investigation why the threshold current density in MBE GaA1As/GaAs lasers is about twice that of similar LPE devices has pointed out the role of the high surface recombination velocity (~ 4000 cm/s) at the heterointerfaces; this is caused by the tunnelling of the carriers from the active region towards the above mentioned non-radiative centers, within the GaA1As layers 63 As far as the LPE growth of Al-containing compounds is concerned, it should be noted that, untill now, most of the work on GaA1As/ GaAs lasers and solar cells has been carried out with the above technique and the results obtained to date can be considered as being the high point in this field. Stringent requirements on the thickness of the indiyidual layers are to be met both in solar cells and in DHLs. Infact, both calculations 64 and experimental evidence6 t, ~z show that in GaA1As/GaAs DHLs the threshold current density decrease linearly ~vith the thickness d of the active region, down to 0.1/am. Since 2.5 kA/cm 2 is considered as being a practical upper limit for CW operation at RT, it follows that d should be less than 0.5/am 64. Hovel has shown that thickness of the GaA1As window in p-GaA1As/p-GaAs/n-GaAs solar cells should be less than 0.5/am in order to have conversion efficiencies greater than 18 ~ at AMO s s. Each one of the above techniques permits the growth of layers
540 whose thickness is within the 0.1-1.0/am range, however, the reproducibility of the results depends on the particular technique used. Therefore some comments on the ultimate capabilities of LPE, VPE and MBE in achieving very thin layers would be useful. It has been reported that by means of MBE, superlattice consisting of a sequence of GaA1As and GaAs layers can be reproducibily grown, even when the thickness of each single epitaxial film is as thin as few monolayers 6s As for LPE, the intrinsic difficulties of growing reproducible films in the sub-micron range have already been mentioned; moreover in a number of cases, the epitaxial temperature is within the 800-900 C range; thus, cross-diffusion of constituent atoms across heterointerfaces rises serious problems, expecially when abrupt interfaces are required. The interface width of LPE GaA1As/GaAs heterostructures has recently been measured by means of Auger spectroscopy to be about 150 /~66. Consequently, the thickness of a well defined layer has a lower limit of a few hundred angstroms. However, Holonyak's g r o u p 67 has reported on the LPE growth of 200 A thick InGaAsP layers, and on the occurrence of quantum size effects at 77 K in lasers incorporating the above layers. Holonyak et al. 6 s have obtained very interesting results using OMP, since they achieved RT-CW operation in photo-pumped GaA1As/ GaAs lasers with a 200 A thick active region. Little data on very thin layers grown by conventional VPE exist, however it should be recalled that the early results in the superlattice growth were obtained by means of this tecnique, using ASH3, PH3, HCI and Ga 69; in this work superlattices have been obtained which have a 200 A periodicity.
STO!CHIOMETRY An increasing amount of attention is being paid to the influence of the non-stoichiometry on the epitaxial layer properties. In view of this, typical III-V and IV-VI compounds, such as GaAs, PbTe and
541 PbSnTe are briefly reviewed below. The influence of the growth techniques on the layer stoichiometry is clearly shown when the behaviour of an amphoteric dopant, such as Ge, which constitutes a donor and an acceptor in VPE and LPE samples respectively, is considered. These results can be easily interpreted by considering that Ge is likely to be incorporated at Asvacancy sites in Ga-rich samples and viceversa. The relationship between the layer properties and their stoichiometry, determined by the growth process, has been discussed in a number of papers. It has been shown that GaAs: Zn LPE layers have minority carrier diffusion length about twice as long as that of similarly doped VPE samples prepared by the hydride method under the best surface morphology condition, that is when AsH3/GaC1 = 3 7o. VPE G a A s : Z r samples have recently been grown with the AsH3]GaC1 ratio set at 3, 1, 1/3 and it has been shown that stoichiometric- or Ga-rich-vapourphase prepared layers have a minority carrier diffusion length very close to that of LPE samples 71 . Moreover, Miller et al. have reported that unintentionally doped n-GaAs samples have a significantly lower concentration of the 0.82 eV deep level, when grown under stoichiometric or Ga-rich vapour conditions 72. In VPE epilayers the above deep level is, to some extent, connected with the presence of oxygen in the growth ambient 73; on the other hand, it is well known that O behaves as a shallow donor in LPE samples. The above results have been interpreted in terms of the association of an O impurity with a Ga-vacancy which gives a deep donor; this situation seems to be very general for VPE samples, since these same considerations can be also extended to Cr- and Fe-related deep levels. For GaP, Stringfellow has obtained similar results, which indicate an improvement both in electroluminescence efficiency and in minority carrier lifetime, as the vapour composition moves towards the metal-rich side 74. According to the Harris et al. sketch of the GaAs existence region 76, LPE layers are usually grown at temperatures at which the deviation from the stoichiometry should be close to zero. Although the problem of the shape and the extension of the GaAs existence region requires further investigations, the growth temperature range may ac-
542 count for both the relatively high minority carrier diffusion length and the practical absence of deep levels related to Ga-vacancies 77. In this context, it is interesting to note that, according to Ashley et al. 78 it is possible to grow amphoteric GaAs:Si p-n junctions by means of LPE within the 905-8"80 C temperature range: infact, Si behaves as a donor when the growth temperature T is above 895 C, and as an acceptor when T < 895 C. These authors have shown that the minority carrier diffusion lenght is as high as 1 3 / l m and 7/zm respectively, thus indicating that this junction could be very interesting in solar cells. In OMP p-GaAs layers the minority carrier diffusion length equals that of similarly doped LPE samples 7s . This may account for the really good results achieved in minority carrier OMP devices. As far as MBE is concemed, it should be noted that Ga- or Asrich samples can be reproducibly grown simply by adjusting Ga to As flux ratio, according to each definite growth temperature, and that the growing-surface high-energy-electron-diffraction pattern clearly indicates which of the two cases occurs s, 79. Cho et al. have shown that the near-bandgap photoluminescence of MBE unintentionally doped samples is much higher in Ga-rich material than in As-rich one 7 ~, and it should be recalled that the photoluminescence efficiency of MBE GaAs: Sn samples is higher than that of similarly doped LPE material a°. As a consequence of the PbTe and PbSnTe exitence region width a 1, the deviations from the stoichiometry are large enough to effectively determine the carrier type and concentration in unintentionally doped samples. In this case, epitaxial technologies such as MBE and Hot Wall Epitaxy (HWE) may give noteworthy results, since they allow for a precise control of the deviation from the stoichiometry. For instance, Lopez-Otero has reproducibly grown HWE p- or n-type samples by using two sources containing PbTe and Te respectively a2 . Smith and Pickhard have shown that MBE Pb0. 77Sn0.2 a t e films grown by using PbTe and SnTe sources invariably had p-type carriers, while por n-type samples could be grown, in a controllable way, using a third source containing Pb ,a a
543 CONCLUSION In conclusion, a few comments will be made on some possible future developments in the epitaxial technologies. LPE was the first epitaxial technology developed, probably because of the simplicity of the experimental set-up; for these same reason it appears to be very suitable even for future laboratory-scale experiments. Even though intrinsic difficulties exist in such aspects as precise control both of the layer thickness and of the solid solution composition, LPE has primarily been used in epitaxial structure growth for device applications. This may account, at least in part, for the fact that LPE DHLs operating both within the 0.8-0.9 ~m 84 and in the 1.0-1.3/am 8s spectral regions have the longest lifetime obtained up to now, and that the conversion efficiency of LPE solar ceils is still unsurpassed. VPE is a versatile technology, and almost all the compounds, alloys and heterostructures suitable for practical applications can be successfully grown, with the notable exception of Al-containing materials, which cannot be obtained in conventional silica reactors. However, it :is interesting that highly efficient (20% at AM1) solar cells have been obtained by using very simple VPE-n*/p/p ÷ GaAs homostructures, in which the n ÷ window layer is about 1000 A thick and the antireflection coating has been obtained directly by anodic oxidation 86. This result is seen as giving new impetus to VPE for large area solar cell fabrication, since this technique is adaptable for large-scale multislice operation. The OMP variant has gained increasing attention, since it is quite similar to the silicon epitaxy, and thus seems to be suitable for industrial exploitation. The main feature of MBE is that the growth is controlled essentially by kinetics, rather than by diffusion processes through liquid or gaseous layers, as in LPE and VPE respectively; for this reason it is possible to have epitaxial layers with almost any predetermined doping or composition profile in the growth direction. Moreover the relatively low rate allows for a very accurate control of the layer thick-
544 ness. T h e r e f o r e , it seems to be a very suitable t e c h n i q u e w h e n high spatial r e s o l u t i o n in the g r o w t h d i r e c t i o n is required. F u r t h e r advantages can be t a k e n o f the writing s7 and s h a d o w i n g s8 capabilities, w h i c h allow f o r the precise d e p o s i t i o n o f layer s t r u c t u r e s w i t h specific material c o m p o s i t i o n at specified l o c a t i o n s o n a given substrate. This is c o n s i d e r e d interesting for the f a b r i c a t i o n o f multilevel integr a t e d o p t o - e l e c t r o n i c circuits, and in this c o n n e c t i o n b o t h the in-situ f o r m a t i o n o f n o n - a l l o y e d o h m i c c o n t a c t s o n n÷-GaAs layers 89 and single-crystal A1-Schottky-barriers o n GaAs 9 0 have r e c e n t l y b e e n rep o r t e d . These results m a y indicate t h a t the f a b r i c a t i o n o f sophisticated devices during a single U H V - r u n is conceivable.
REFERENCES
1. 2. 3.
4. 5.
6. 7.
8. 9. 10. 11. 12. 13.
L.R. DAWSON -- Progress in Solid State Chemistry, H. Reiss & J.O. Mc Caldin eds, Pergamon Press, N.Y., 1972, Vol. 7, pp. 117-139. E.A. GIESS, R. GHEZ -- Epitaxial Growth, J.W. Matthews ed., Academic Press, N.Y., 1975, part A, pp. 183-209. D.W. SHAW -- Ibidem, pp. 89-106. L.L. CHANG, R. LUDEKE -- Ibidem, pp. 37-69. A.Y. CHO, J.R. ARTHUR -- Progress in Solid State Chemistry, J.O. Mc Caldin & G. Somorjai eds., Pergamon Press, N.Y., 1975, Vol. 10, pp. 157191. F.C. FRANK, J.H. VAN der MERWE --Proc. Roy. Soc., A198, 205, 1949. H. KRESSEL --J. Electr. Mat., 4, 1081, 1975. J.W. MATTHEWS - Epitaxial Growth, J.W. Matthews ed., Academic Press, N.Y., 1975, part B, pp. 560-607. M.S. ABRAHAMS, CJ. BUIOCCHI, G.H. OLSEN -- J. Appl. Phys., 46, 4259, 1975, A.G. SIGAI, C.J. NUESE, R.E. ENSTROM, T.J. ZAMEROWSKY -- J. Elec. trochem. Soc., 120, 947, 1973. CJ. NUESE, A.G. SIGAI, J.J. GANNON, T J . ZAMEROWSKY --J. Electr. Mat., 4, 51, 1975. M. ETTEMBERG, C.J. NUESE, J.R. APPERT, J.J. GANNON, R.E. ENSTROM --J. Electr. Mat., 4, 37, 1975. M.S. ABRAHAMS~ L.R. WEISEMBERG, C.J. BUIOCCHI, J. BLANC -- J. Mat. Sc£, 4, 223, 1969.
545 M.S. ABRAHAMS, C.J. BUIOCCHI - J. AppL Phys., 37, 1973, 1969. M. ETTEMBERG -- J. AppL Phys., 45, 901, 1974. CJ. NUESE --J. Electr. Mat., 6, 253, 1977. M. ETTEMBERG, G.H. OLSEN -- J. AppL Phys., 48, 4275, 1977. G.A. ROZGONY, P.M. PETROFF, M.B. PANISH -- J. Cryst. Growth, 27, 126, 1974. 19. R.J. NELSON, R.G. SOBERS -- AppL Phys. Lett., 32, 761, 1978. 20. D.V. LANG, R.A. LOGAN -- Appl. Phys. Let&, 31, 684, 1977. 21. G.H. OLSEN, M.S. ABRAHAMS, T J . ZAMEROWSKY - J. Electrochem. Soc., 121, 1650, 1974. 22. G.H. OLSEN, M.S. ABRAHAMS, C.J. BUIOCCHI, T J . ZAMEROWSKY J. Appl. Phys., 46, 1643, 1975. 23. G.H. OLSEN --J. Cryst. Growth, 31, 223, 1975. 2.4. CJ. NUESE, G.H. OLSEN, M. ETTEMBERG, J J . GANNON, T J . ZAMEROWSKY - - AppL Phys. Lett., 29, 807, 1976. 25. H. KRESSEL, CJ. NUESE, G.H. OLSEN - J . Appl. Phys., 49, 3140, 1978. 26. R.E. NAHORY, M.A. POLLAK, J.C. DE WINTER --J. AppL Phys., 48, 320, 1977. 27. H. NAGAI, Y. NOGUKI - Appl. Phys. Lett., 26, 108, 1975. 28. G.B. STRINGFELLOW - - J. AppL Phys., 43, 3455, 1972. 29. J.J. COLEMAN, N. HOLONYAK, MJ. LUDOWISE - Appl. Phys. Lett., 28, 363, 1976. 30. J J . HSIEH, M.C. FINN, J.A. ROSSI - Proc. of Sump. on GaAs and related compounds, St. Louis 1976, Institute of Physics, London, 1977, pp. 37-44. 31. K. TATEISHI, M. NOGANUMA, K. TAKAHASHI -- Jap. J. AppL Phys., 15, 785, 1976. 32. T. WAHO, S. OGAWA, S. MARUYAMA - - J a p . J. Appl. Phys., 16, 1875, 1977. 33. H. SAKAKI, L.L. CHANG, R. LUDEKE, C.A. CHANG, G.A. SAI-HALASZ, L. ESAKI -- AppL Phys. Lett., 31, 211, 1977. 34. G.A. SAI-HALASZ, R. TSU, L. ESAKI - Appl. Phys. Lett., 30, 651, 1977. 35. R. CHARLTON, G . S . H O B S O N - - IEEE Trans. Electron. Dev., ED-20, 812, 1973. 36. G. SALMER, J. PRIBETICH, A. FARRAYRE, B. KRAMER - J. Appl. Phys., 44, 314, 1973. 37. D.R. DECKER -- IEEE Trans. Electron Dev., ED-21,469, 1974. 38. C. KIM, R. STEELE, R. BIERIG - - Electron. Lett., 9, 173, 1973. 39. R.E. GOLDWASSER, F.E. ROSZTOCZY - AppL Phys. Lett., 25, 92, 1974. 40. R.L. DAWSON --J. Cryst. Growth, 27, 86, 1974. 41. T.GJ. van OIRSCHOT, WJ. LESWIN, PJ.A. THIJS, W. NIJMAN -- J. Cryst. Growth, 45, 262, 1978. 14. 15. 16. 17. 18.
546 42. 43. 44. 45.
46. 47. 48. 49. 50.
51. 52. 53. 54.
55. 56. 57. 58. 59. 60. 61. 62. 63. 64. 65.
H.F. LOCKWOOD, M. ETTEMBERG -- J. Cryst. Growth, 15, 81, 1972. H.F. LOCKWOOD, H. KRESSEL -- ]. Cryst. Growth, 27, 97, 1974. D.L. RODE, B. SCHWARTZ, J.W. DILORENZO - Sol. State Electron., 17, 1119, 1974. T. NOZAKI, M. OGAWA, H. TERAO, H. WATANABE -- Proc. of Syrup. on GaAs and related compounds, DeauviUe, 1974, Institute of Physics, London, 1975, pp. 46-54. H.M. COX, J.W. DILORENZO - Proc. of Syrup. on GaAs and related compounds, St. Louis, 1976, Institute of Physics, London, 1977, pp. 11-22. C.E.C. WOOD -- Appl. Phys. Lett., 29, 746, 1976. J. HALLAIS, D. BOCCON-GIBOD, J.P. CHANE', J.M. DURAND, L. HOLLAN --J. Electrochem. Soc., 124, 1290, 1977. Y. OTA --]. Electrochem. Soc., 124, 1795, 1977. G. MAJNI, G. OTTAVIANI -- Thin Film Phenomena - Interface and Interactions, J.E. Baglin, J.M. Poate eds., The Electrochem. Soc., Princeton, 1978, pp. 293-299. J,P. DUCHEMIN, M. BONNET, F. KOELSCH ]. Electrochem. Soc., 125, 637, 1978. J.P. DUCHEMIN, M. BONNET, F. KOELSCH, D. HUYGHE -- ]. Cryst. Growth, 45, 181, 1978. M.B. PANISH --Proc. IEEE, 64, 1512, 1976. HJ. HOVEL, J.M. WOODALL, W.E. HOWARD - P r o c . of Syrup. on GaAs and related compounds, Boulder, 1972, Institute of Physics, London, 1973, pp. 205. H.J. HOVEL -- Solar Cells, Vol. 11, of Semiconductors and Semimetals, R.K. Willardson, A.C. Beer eds., Academic Press, N.Y., 1975. W.D.JOHNSTON Jr., W.M. CALLAHAN -- J. Electrochem. Soc., 123, 1524, 1976. W.D. JOHNSTON Jr., W.M. CALLAHAN - Appl. Phys. Lett., 28, 150, 1976. R.D. DUPUIS, P.D. DAPKUS, R.D. YINGLING, L.A. MOUDY - AppL Phys. Lett., 31, 201, 1977. Nd. NELSON, K.K. JOHNSON, R.L. MOON, H.A. VANDER PLAS, L.W. JAMES - AppL Phys. Lett., 33, 26, 1978. R.D. DUPUIS, P.D. DAPKUS -- AppL Phys. Lett., 32, 473, 1978. H. KRESSEL, M. E'I~rEMBERG -- ]. Appl. Phys., 47, 3533, 1976. M. ETTEMBERG -- AppL Phys. Lett., 27, 652, 1975. W.T. TSANG -- AppL Phys. Lett., 33,245, 1978. H.C. CASEY Jr. - J . AppL Phys., 49, 3684, 1978. A.C. GOSSARD, P.M. PETROFF, W. WIEGMANN, R. DINGLE, A. SAVAGE --AppL Phys. Lett., 29, 323, 1976. -
547
66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90.
C.M. GARNER, Y.D. SHEN, C.Y. SU, G.L. PEARSON, W.E. SPICER - J . Vacuum Sci. Techn., 15, 1480, 1978. E.A. REZEK, N. HOLONYAK, B.A. VOJAK, H. SHICHIJO - J. Appl. Phys., 49, 69, 1978. N. HOLONYAK, R.M. KOLBAS, R.D. DUPUIS, P.D. DAPKUS -- Appl. Phys. Lett., 33, 73, 1978. A.J. BLAKESLEE --]. Electrochem. Soc., 115, 118, 1971. M. ETTEMBERG, C.J. NUESE - J . Appl. Phys., 46, 3500, 1975. M. ETTEMBERG, G.H. OLSEN, C.J. NUESE - Appl. Phys. Lett., 29, 141, 1976. M.D. MILLER, G.H. OLSEN, M. ETTEMBERG -- Appl. Phys. Lett., 31, 538, 1977. A. MIRCEA, A. MITONNEAU -- Appl. Phys., 8, 15, 1975. G.B. STRINGFELLOW, H.T. HALL - ]. Electrochem. Soc., 123, 916, 1976. M. ALLENSON, S.J. BASS - Appl. Phys. Lett., 28, 113, 1976. J.S. HARRIS, Y. NANNICH, G.L. PEARSON, G.F. DAY -- ]. Appl. Phys., 40, 75, 1969. D.V. LANG, A.Y. CHO, A.C. GOSSARD, M. ILEGEMS, W. WIEGMANN J. AppL Phys., 47, 2558, 1976. K.L. ASHLEY, S.W. BEAL - Appl. Phys. Lett., 32, 375, 1978. A.Y. CHO, L. HAYASHI -- Sol. State Electron., 14, 125, 1971. H.C. CASEY, A.Y. CHO, P.A. BARNES -- IEEE J. Quantum Electronics, QE-11, 467, 1975. T.C. HARMAN - J . Nonmetals, 1, 183, 1973. A. LOPEZ-OTERO - Appl. Phys. Lett., 2 6 , 4 7 0 , 1975. D.L. SMITH, V.Y. PICKARD -- J. Electron. Mat., 5, 247, 1976. R.L. HARTMANN, N.E. SCHUMAKER, R.W. DIXON -- Appl. Phys. Lett., 31, 756, 1977. C.C. SHEN, J.J. H-ISIEH, T.A. LIND - Appl. Phys. Lett., 30, 353, 1977. J.C.C. FAN, C.O. BOZLER, R.L. CHAPMAN - Appl. Phys. Lett., 32, 390, 1978. W.T. TSANG, A.Y. ClIO -- Appl. Phys. Lett., 32, 491, 1978. W.T. TSANG, M. ILEGEMS - Appl. Phys. Lett., 31, 301, 1977. A.Y. CHO -- Appl. Phys. Lett., 33, 651, 1978. A.Y. CHO, P.D. DERNIER --J. Appl. Phys., 49, 3328, 1978.
EPITAXIAL TECHNOLOGIES IN THE GROWTH OF SEMICONDUCTORS FOR DEVICE APPLICATIONS*
S. F R A N C H I * , D. M A R G A D O N N A * * , *
CNR.MASPEC,
**
ASSORENI,
PARMA
a n d C. P E L O S I *
- Italy.
MONTEROTONDO
(Roma)
- Italy.
Abstract -- One of the most promising trends in device development foreseen for the near future is the incorporation of multilayer structures in all devices in which performance improvements justify the higher technological costs. In this perspective, the most diffused epitaxial technologies, which are Liquid Phase Epitaxy (LPE), Vapour Phase Epitaxy (VPE), Pyrolysis of Organometallic Compounds (OMP) and Molecular Beam Epitaxy (MBE), are discussed with respect to a number of definite problems. These topics are related to: 1) Interfaces; 2) Composition profiles; 3) Doping profiles and 4) Stoichiometry. Although the present review refers mainly to the III-V compounds and alloys, some attention has been paid to outlining the general aspects of the above problems. Moreover some comments are made on the possible developments and industrial impact of the epitaxial technologies.
INTRODUCTION A great n u m b e r o f e l e c t r o n i c devices w h i c h are c u r r e n t l y p r o d u c ed b o t h in industrial a n d in r e s e a r c h l a b o r a t o r i e s are f a b r i c a t e d using e p i t a x i a l t e c h n o l o g i e s . O n e r e a s o n f o r this is t h a t , generally speaking,
* Presented at the IV Scientific Meeting of the Italian Association for Crystal Growth (AICC), Parma, Italy, 26-28 February 1979.
530
epitaxial materials have better properties than the bulk ones, as far as structural and electrical properties are concerned. This is mainly due to the fact that in epitaxy, independently of the particular method used, the growth is carried out at temperatures which are several hundred degrees lower than the melting point of the compound or alloy. The lower temperature growth accounts, at least in part, for the concentration both of chemical and crystalline defects, which is lower in epilayers than in materials grown from a nearly stoichiometric melt. At the same time, the lattice vacancy concentration is expected to decrease. Multilayer structures have gained an increasing interest both in opto-electronic and in microwave devices; the importance of double heterostructures in lasers and LEDs, and of single heterostructures in solar cells and in light detectors is well known. In the microwave field, the interest is presenly focused on homostructures in which the layers have different doping levels, and devices have successfully been fabricated which incorporate three or four layers. However, due to the intrinsic properties of some III-V alloys or compounds other than the usual GaAs, it is conceivable that in the near future heterostructures will be considered even in this area. Many papers have reported that in devices, such as injection la-" sers, Gunn diodes, IMPATTs, MESFETs and mixer diodes, the thickness of the active region should be in the sub-micron range, as it will be discussed below. The above considerations clearly underline the prominent role of the epitaxial processes in the semiconductor technology, even if, in particular cases, the advantages of the ion implantation technique should be taken into account. The more widely used epitaxial technologies are: a) Liquid Phase Epitaxy (LPE); b) Vapour Phase Epitaxy (VPE) and Pyrolysis of Organometallic Compounds (OMP); c) Molecular Beam Epitaxy (MBE); in the following no discussion of apparatuses and processes relative to the above techniques will be made, since a number of excellent review paper on LPE 1, 2, VPE 3 ' and MBE 4' s are available. Considering the present situation, none of the above technologies
531
can be thought as being the most suitable for a broad range of experimental situations, even if each one of them may have distinct advantages in specific cases. In the following the epitaxial technologies will be discussed with respect to specific device problems. These topics are connected with: a) Interface-related problems caused by lattice mismatch; the achievement of: b) doping and c) composition profiles in the growth direction, and d) the control of the stoichiometry of epitaxial layers. Although the present review refers to the III-V compounds and alloys, some attention has been paid to outlining general aspects of the above problems.
INTERFACE-RELATED PROBLEM
In many of the currently produced minority carrier devices, such as, for instance, semiconductor lasers, LEDs and solar cells, the active region is within the minority carrier diffusion length from homo- or hetero-epitaxial interfaces; consequently defects occurring at heterointerfaces strongly influence the properties of the above devices. In this section, the lattice mismatch is considered as a source of interfacial defects. Frank and Van der Merwe have shown that, in order to relieve the elastic strain due to the lattice mismatch at interfaces, a misfit dislocation network is almost always generated6 ; simple considerations indicate that the misfit dislocation linear density is proportional to the lattice mismatch parameter e = Aa/a. Assuming that a nonradiative recombination center is associated to each a t o m having a dangling bond, it follows that the surface density of these centers, as well as the surface recombination velocity, are linearly related to e 7 . The above model is somewhat oversimplified since the plastic relaxation may not be complete and, consequently, part of the lattice mismatch may result in elastic strain; this has been confirmed by Matthews 8 who has shown that a critical thickness exist, below which the epitaxial layer is under strain and no dislocation network is gen-
532 erated. The above thickness has been evaluated to be about one half the reciprocal of the linear dislocation density. The morphology of the misfit dislocation network has been studied in a number of heteroepitaxial structures (InGaP/GaPg, 1 o, 1 1, InGaAs/GaAs 12 , GaAsP/GaP la ) as well as in homoepitaxial interfaces 14, in which the mismatch is caused by doping level differences. Recently the relationsh!p between interfacial misfit dislocations and device opto-electronic properties has experimentally been investigated. Ettemberg et al. have shown that, both in GaAs/GaP 1 s and in InGaAs/ GaAs 1 2 heterostructures, the minority carrier diffusion length h in the epitaxial layer is upper-limited by the mean spacing between misfit dislocations. This result is best understood while considering that, in the first approximation, the luminescence internal quantum efficiency is proportional to X2. Detailed studies of the dependence of the surface recombination velocity S at interfaces on the lattice mismatch have been made using InGaP/GaAs heterojunctions along with different experimental techniques 16, 17. The main result is that S is proportional to the lattice mismatch, and the coefficient is surprisingly close to the one estimated on the basis of the above model. These data definitely confirm that misfit dislocations act as non-radiative recombination centers, and limit the minority carrier diffusion length; consequently the effect of lattice mismatch in devices such as lasers, LEDs and solar cells may be devastating. Within this context, the important role of GaA1As/GaAs structures in opto-electronics should be recalled; since the lattice mismatch, in general, is within the 10 -4 range, this interface is void of misfit dislocations t a and has a surface recombination velocity evaluated to be 450-+ 100 cm/s19: this corresponds to a recombination center concentration of 4 • 109 cm "2 , which compares favourably with that obtained by Lang et al. using the deep level transient spectroscopy 2°. The experimental evidence, both in GaA1As/ GaAs and in InGaP/GaAs structures, definitely points out that, in general, for S ~< 104 cm/s, the lattice mismatch should be within the 10 -4 - 10 -3 range. In the case of epilayers in compression on the substrate, the tel-
533 atively low surface recombination velocity at the active region interface is obtained by using a number of lattice-matching layers between the substrate and the active region itself: by step-grading the composition of the layers, the lattice constant is progressively matched to the active region one. The main results of this technique are that: the misfit dislocations are localized at the interfaces where the composition is step-graded, and the inclined dislocations propagating from the substrate, or generated within the epitaxial layers, are bent into the interfaces z 1, 22. In this way, a dislocation density of 100 cm "2 could be obtained in the constant composition region 2a On the other hand, in cases in which the epilayer is in tension, tile above technique does not give good results, due to the formation of microcracks in the dislocated epitaxial region. In this situation, a continuously graded layer is grown between the substrate and the active region; consequently, the strain-relieving dislocations are generated with a density which is proportional to the composition gradient and, then, propagate through the device I 3. However, if low enough composition gradients are used, the dislocation density is minimized, so that device-quality materials are obtained. In the following, the epitaxial technologies are compared in so far as the handling of the growth of lattice-matched structures is concerned. With the noteworthy exception of GaA1As/GaAs structures, the epilayer composition must be kept under close control; for example, in order to have a lattice mismatch of 2 • 10 -4 in the InGaP/GaAs heterostructure, the epitaxial layer must be grown with an In molar fraction that is 0.496 ± 0.003. It has been shown that VPE lattice-matched structures based on III-V compounds can be reproducibily grown, either when the epilayer is in tension or in compression. A very brilliant example of the VPE uses is the growth of InGaAs/ InGaP/GaAs and GaAsP/InGaP/GaAs heterostructures, from which lasers operating (under RT and CW conditions) within the 1.0-1.3/am 24 and 0.7 jam 2 s spectral regions respectively, have been successfully successfully fabricated. In the case of LPE, some problems arise since the number of dif-
534 ferent layers is limited by the number of bins in the multibin graphite boat, and the growth of continuously graded layers is intrinsically difficult. Anyway, in a few cases, continuously graded matching layers have been obtained, taking advantage of the phosphorous depletion in solutions based on the GaAsSbP 26 and InGaAsP 27 quaternary systems; however, the same constituent source-depletion may be detrimental when very constant composition layers are required. In LPE the precise control of the supersaturation has an important role, expecially when the epitaxial layer is required to have a composition which is different from that of the seed. Infact, in these cases even a slight undersaturation causes the partial dissolution of the substrate and then the composition of the growing layer is different from the desired one. At any rate, LPE has features which are very interesting: firstly, it has been reported that in the case of perfect lattice matching, the epilayer has a dislocation density which is about one order of magnitude less than that of the substrate; the second characteristic is that, at least in the case of InGaP/GaAs 2s and InGaAsP/GaAsP 29 , over a certain range of liquid composition the LPE layers all have a growing composition which is lattice-matched to the substrate, rather than conforming to the bulk phase diagram. This could be explained by the contribution of dislocation- and elastic-stress-energy to the free energy of the system. Recently, however, some doubts have been raised about the occurrence of the above latching in InGaAsP/InP heterostructures 30 MBE is considered to be a very versatile technology, even in connection with interface-related problems: examples of the growth of continuously graded epilayers in GaAsP 3 i, inGaAs 31 and GaAsSb 32 alloys have been reported; furthermore, very interesting results have been obtained in the growth of InGaAs/GaAsSb heterostrnctures and superlattices, where the lattice-matching conditions have been obtained by carefully controlling the composition of both the ternary solid solutions 3 3, 3 4
535 DOPING PROFILES
The obtaining of precisely controlled doping profiles is a major' problem in the epitaxial growth of structures suitable for microwave devices. For example, in order to reduce the noise figure in Gunn diodes, a cathode-notch doping profile has been proposed3 s, which can be described in terms of a n ÷ / n ' / n / n ÷ homoepitaxial structure. Both theoretical 3 6, 3 7 and experimental studies 3 8 , 3 9 show that higher DC to RF conversion efficiencies can be achieved by using a low-high-low doping profile in GaAs IMPATT diodes. Although most of the above microwave devices are presently prepared by the halide VPE process, interest has been shown for the LPE method, since LPE grown n-GaAs has an electron trap concentration which is in the low 1012cm'3 range, that is generally lower than the VPE one. However, with LPE, it is difficult to obtain layers with precisely controllable thickness in the sub-micron range, due to the problems of experimentally controlling the nutrient solution supersaturation. A number of methods have been proposed for circumventing this problem, such as using one 40 or two 41 d u m m y wafers which precede the substrate and equilibrate the solution, or using a separate source wafer on the top of each melt to maintain the equilibrium throughout the entire growth cycle 42, 43. A different approach for controlling the active layer thickness is the anodic oxidation of the layer itself 44, however this m e t h o d holds only when the layer to be thinned is not sandwiched in multilayer structures. In any case, none of the above techniques has given completely satisfactory results. For obtaining very controllable doping profiles, the problem of the cross-diffusion both of constituent and dopant elements across the interface is important, expecially when very thin layers must be grown on films or substrates which have different doping or composition. Since the diffusion of atoms through a solid is a thermally activated process, it is immediately evident that the growth, and more generally, the processing temperatures are to be kept as low as possible, and this may constitute a prerequisite in the choice of the epitaxial technology. For example, the typical temperature ranges for growing
536 GaAs are 850-900 C, 700-750 C, and 500-440 C in LPE, VPE and MBE respectively. The cross-diffusion related problems are best discussed in connection with the fabrication of GaAs MESFETs, in which a 0.2-0.3/am thick n-doped layer is epitaxially deposited on Cr-doped semiinsulating (SI) material. Low noise MESFETs generally operate at a bias-voltage so that the depletion region is only a few hundred angstroms far from the active] SI-layer interface; the power gain has been shown to be proportional to the majority carrier mobility in the active region. In VPE devices grown directly on SI-GaAs, the I-V characteristics show a typical hysteresis, moreover the electron mobility was shown to degrade drastically towards the interface. These results have been interpretated in terms of deep levels, due to the Cr-diffusion from the substrate into the epilayer, which act as efficient electron traps 4s, 46. In order to avoid these negative effects, an intrinsic low-conductivity buffer layer has successfully been interposed between the substrate and the active region, whose function is to limit the diffusion of Cr into the upper layer. However, Wood has shown that in MBE GaAs MESFETs no degradation of the majority carrier mobility, nor looping in the I-V curves is present, even in the absence of the buffer layer, thus indicating that the Cr-outdiffusion is not a major problem at MBE GaAs growth temperatures 47 Recently, a halide VPE process has been described, in which the growth temperature has been reduced to 630 C4s; however, even if these homostructures have the MESFET-quality without any buffer layer, the epilayer majority carrier mobility is lower than that of MBE samples. Since many silicon devices, such as microwave diodes and transistors, are fabricated by depositing a lightly doped ('~ 1016 cm-3) epilayer on a heavily doped substrate ('~ 1018 cm-3 ), the same sharpness requisites in doping profiles exist, which have already been discussed in connection with III-V devices. Ota has recently used MBE in preparing n-Si/n*-Si structures in which the n-layer Sb-doping level is controlled within the high 1013_ high 1017cm'3 range 49 . To investigate the outdiffusion of the As-dopant from the substrate (n "~ 5 • 1019
537 cm "3) into the epilayer, SIMS profiles of MBE and VPE samples have been measured. The results show that in MBE grown at 800 C for 3 hr, and 1050 C for 2 hr, the As-outdiffusion width is 0..03 and 0.()2 lain respectively49; these values are definitely lower than those generally obtained (~> 1/am) using SiCI4 VPE at 1200 C. In connection with the trend of reducing the epitaxial temperature, the Solid State Epitaxy (SSE) could be potentially interesting s° The autodoping could hamper the achievement of precisely controlled doping profiles in VPE; this p h e n o m e n o n is related to the existence of a stagnant layer, just around the substrate, which acts as a sink for impurities, whether liberated from the substrate or contained in the main gas stream. Both intentional or unintentional dopants are, therefore, available for incorporation in the growing layer for a time determined by the stagnant layer volume. Different methods have been proposed for eliminating this problem, the simplest being the use of long purge times between growth intervals. A quite different approach, consisting of the reduction of the epitaxial reactor pressure, has recently put forward. This m e t h o d has been used both in Si s I and in GaAs OMP s 2 epitaxy. In the first case Sill4, SiHC13, SiH2C12 and SIC14 have been used as Si- sources in a cold wall horizontal reactor. The operating pressure has been mantained in the 10-100 torr range by means of a Roots pump fitted with a liquid nitrogen trap. Defectfree layers have been grown at temperatures lower than 900 C and doping profile has been shown to be very sharp, even at the interface with a highly Te-doped substrate (/9 "" 10"3Ohm cm) sl In the GaAs OMP epitaxy, good crystalline-quality layers have been grown within the 520-750 C temperature and 25-600 torr pressure ranges s2. The carrier concentration reduction by about two orders of magnitude from the 6 • 10 -4 Ohm • cm Te-doped substrate towards the epitaxial layer has been achieved in about 0.05 /am using growth temperatures as high as 750 C. This result definitely suggests that the autodoping can be effectively reduced, while the reactor is used at sub-atmospheric pressure, and that this is independent of which VPE m e t h o d is actually being used.
538
COMPOSITION PROFILES The growth of heteroepitaxial structures is very important for obtaining high efficiency opto-eletronic devices such as injection lasers, LEDs, solar cells and photodetectors. For instance, it has been shown that double heterostructure lasers (DHL) have a threshold current density which is markedly lower than those of homojunction lasers, and that this is essentially related to the improved confinement both of injected carriers and of electromagnetic field in the active region s3. As for solar cells, p-GaA1As/p-GaAs/n-GaAs structures have a significantly higher conversion efficiency than GaAs homojunctions: this is understood by recalling that GaA1As/GaAs interfaces have a surface recombination velocity 2-3 orders of magnitude lower than the vacuum/GaAs one, and that half the carriers generated by AM1 radiation are produced in GaAs within 0.2/~m. Ternary or quaternary III-V alloys are used in the active region of opto-electronic devices operating either within the yellow or red regions of the visible spectrum, or within the 1.0-1.3 #m spectral band. Since for pratical reasons the bulk materials are restricted to binary compounds, such as GaAs and InP, lattice-matching-layers of suitable composition can be used between the substrate and the active region. The possibility of using epitaxi~ techniques for growing multilayer structures incorporating different materials has already been discussed in connection with the problem of lattice-matching layers. However, due to the technological importance of Al-containing alloys, such as GaA1As, some further comments on the growth of these solid solutions are in order. Because of the reactivity of A1C1, very poor results have been obtained by VPE using conventional silica reactors. In order to circumvent this difficulty all-allumina VPE reactors have been used 56 to grow n-A1As/p-GaAs heterostructures, which have a sea-level sunlight conversion efficiency of 13-18% 57. Although these results should be compared with those of LPE devices (18-21%), it is noteworthy that VPE cells have significantly larger areas. It has recently been shown that GaAIAs/GaAs heterostructures grown by OMP have
539 been incorporated in highly efficient, large-area solar cells s s, s 9 and in DHLs operating under CW-RT conditions 60, with a threshold current density as low as that of the best LPE GaA1As/GaAs DHLs 61, 62 These results clearly show that OMP is a technique suitable for producing highly efficient solar cells and lasers on a large scale. The growth of really good GaAIAs epilayers is a major problem for MBE: infact, the MBE GaA1As has a photoluminescence efficiency which usually is 10-20 times lower than that of the best LPE layers. This is thought to be related to the presence of large concentr. tions of non-radiative centers, which are caused by residual CO in the ultra-high-vacuum growth environment. This unintentional doping is a problem only for the growth of Al-containing solid solutions, and this is understood by recalling that the CO sticking coefficient is much higher in GaA1As than in GaAs s . Detailed investigation why the threshold current density in MBE GaA1As/GaAs lasers is about twice that of similar LPE devices has pointed out the role of the high surface recombination velocity (~ 4000 cm/s) at the heterointerfaces; this is caused by the tunnelling of the carriers from the active region towards the above mentioned non-radiative centers, within the GaA1As layers 63 As far as the LPE growth of Al-containing compounds is concerned, it should be noted that, untill now, most of the work on GaA1As/ GaAs lasers and solar cells has been carried out with the above technique and the results obtained to date can be considered as being the high point in this field. Stringent requirements on the thickness of the indiyidual layers are to be met both in solar cells and in DHLs. Infact, both calculations 64 and experimental evidence6 t, ~z show that in GaA1As/GaAs DHLs the threshold current density decrease linearly ~vith the thickness d of the active region, down to 0.1/am. Since 2.5 kA/cm 2 is considered as being a practical upper limit for CW operation at RT, it follows that d should be less than 0.5/am 64. Hovel has shown that thickness of the GaA1As window in p-GaA1As/p-GaAs/n-GaAs solar cells should be less than 0.5/am in order to have conversion efficiencies greater than 18 ~ at AMO s s. Each one of the above techniques permits the growth of layers
540 whose thickness is within the 0.1-1.0/am range, however, the reproducibility of the results depends on the particular technique used. Therefore some comments on the ultimate capabilities of LPE, VPE and MBE in achieving very thin layers would be useful. It has been reported that by means of MBE, superlattice consisting of a sequence of GaA1As and GaAs layers can be reproducibily grown, even when the thickness of each single epitaxial film is as thin as few monolayers 6s As for LPE, the intrinsic difficulties of growing reproducible films in the sub-micron range have already been mentioned; moreover in a number of cases, the epitaxial temperature is within the 800-900 C range; thus, cross-diffusion of constituent atoms across heterointerfaces rises serious problems, expecially when abrupt interfaces are required. The interface width of LPE GaA1As/GaAs heterostructures has recently been measured by means of Auger spectroscopy to be about 150 /~66. Consequently, the thickness of a well defined layer has a lower limit of a few hundred angstroms. However, Holonyak's g r o u p 67 has reported on the LPE growth of 200 A thick InGaAsP layers, and on the occurrence of quantum size effects at 77 K in lasers incorporating the above layers. Holonyak et al. 6 s have obtained very interesting results using OMP, since they achieved RT-CW operation in photo-pumped GaA1As/ GaAs lasers with a 200 A thick active region. Little data on very thin layers grown by conventional VPE exist, however it should be recalled that the early results in the superlattice growth were obtained by means of this tecnique, using ASH3, PH3, HCI and Ga 69; in this work superlattices have been obtained which have a 200 A periodicity.
STO!CHIOMETRY An increasing amount of attention is being paid to the influence of the non-stoichiometry on the epitaxial layer properties. In view of this, typical III-V and IV-VI compounds, such as GaAs, PbTe and
541 PbSnTe are briefly reviewed below. The influence of the growth techniques on the layer stoichiometry is clearly shown when the behaviour of an amphoteric dopant, such as Ge, which constitutes a donor and an acceptor in VPE and LPE samples respectively, is considered. These results can be easily interpreted by considering that Ge is likely to be incorporated at Asvacancy sites in Ga-rich samples and viceversa. The relationship between the layer properties and their stoichiometry, determined by the growth process, has been discussed in a number of papers. It has been shown that GaAs: Zn LPE layers have minority carrier diffusion length about twice as long as that of similarly doped VPE samples prepared by the hydride method under the best surface morphology condition, that is when AsH3/GaC1 = 3 7o. VPE G a A s : Z r samples have recently been grown with the AsH3]GaC1 ratio set at 3, 1, 1/3 and it has been shown that stoichiometric- or Ga-rich-vapourphase prepared layers have a minority carrier diffusion length very close to that of LPE samples 71 . Moreover, Miller et al. have reported that unintentionally doped n-GaAs samples have a significantly lower concentration of the 0.82 eV deep level, when grown under stoichiometric or Ga-rich vapour conditions 72. In VPE epilayers the above deep level is, to some extent, connected with the presence of oxygen in the growth ambient 73; on the other hand, it is well known that O behaves as a shallow donor in LPE samples. The above results have been interpreted in terms of the association of an O impurity with a Ga-vacancy which gives a deep donor; this situation seems to be very general for VPE samples, since these same considerations can be also extended to Cr- and Fe-related deep levels. For GaP, Stringfellow has obtained similar results, which indicate an improvement both in electroluminescence efficiency and in minority carrier lifetime, as the vapour composition moves towards the metal-rich side 74. According to the Harris et al. sketch of the GaAs existence region 76, LPE layers are usually grown at temperatures at which the deviation from the stoichiometry should be close to zero. Although the problem of the shape and the extension of the GaAs existence region requires further investigations, the growth temperature range may ac-
542 count for both the relatively high minority carrier diffusion length and the practical absence of deep levels related to Ga-vacancies 77. In this context, it is interesting to note that, according to Ashley et al. 78 it is possible to grow amphoteric GaAs:Si p-n junctions by means of LPE within the 905-8"80 C temperature range: infact, Si behaves as a donor when the growth temperature T is above 895 C, and as an acceptor when T < 895 C. These authors have shown that the minority carrier diffusion lenght is as high as 1 3 / l m and 7/zm respectively, thus indicating that this junction could be very interesting in solar cells. In OMP p-GaAs layers the minority carrier diffusion length equals that of similarly doped LPE samples 7s . This may account for the really good results achieved in minority carrier OMP devices. As far as MBE is concemed, it should be noted that Ga- or Asrich samples can be reproducibly grown simply by adjusting Ga to As flux ratio, according to each definite growth temperature, and that the growing-surface high-energy-electron-diffraction pattern clearly indicates which of the two cases occurs s, 79. Cho et al. have shown that the near-bandgap photoluminescence of MBE unintentionally doped samples is much higher in Ga-rich material than in As-rich one 7 ~, and it should be recalled that the photoluminescence efficiency of MBE GaAs: Sn samples is higher than that of similarly doped LPE material a°. As a consequence of the PbTe and PbSnTe exitence region width a 1, the deviations from the stoichiometry are large enough to effectively determine the carrier type and concentration in unintentionally doped samples. In this case, epitaxial technologies such as MBE and Hot Wall Epitaxy (HWE) may give noteworthy results, since they allow for a precise control of the deviation from the stoichiometry. For instance, Lopez-Otero has reproducibly grown HWE p- or n-type samples by using two sources containing PbTe and Te respectively a2 . Smith and Pickhard have shown that MBE Pb0. 77Sn0.2 a t e films grown by using PbTe and SnTe sources invariably had p-type carriers, while por n-type samples could be grown, in a controllable way, using a third source containing Pb ,a a
543 CONCLUSION In conclusion, a few comments will be made on some possible future developments in the epitaxial technologies. LPE was the first epitaxial technology developed, probably because of the simplicity of the experimental set-up; for these same reason it appears to be very suitable even for future laboratory-scale experiments. Even though intrinsic difficulties exist in such aspects as precise control both of the layer thickness and of the solid solution composition, LPE has primarily been used in epitaxial structure growth for device applications. This may account, at least in part, for the fact that LPE DHLs operating both within the 0.8-0.9 ~m 84 and in the 1.0-1.3/am 8s spectral regions have the longest lifetime obtained up to now, and that the conversion efficiency of LPE solar ceils is still unsurpassed. VPE is a versatile technology, and almost all the compounds, alloys and heterostructures suitable for practical applications can be successfully grown, with the notable exception of Al-containing materials, which cannot be obtained in conventional silica reactors. However, it :is interesting that highly efficient (20% at AM1) solar cells have been obtained by using very simple VPE-n*/p/p ÷ GaAs homostructures, in which the n ÷ window layer is about 1000 A thick and the antireflection coating has been obtained directly by anodic oxidation 86. This result is seen as giving new impetus to VPE for large area solar cell fabrication, since this technique is adaptable for large-scale multislice operation. The OMP variant has gained increasing attention, since it is quite similar to the silicon epitaxy, and thus seems to be suitable for industrial exploitation. The main feature of MBE is that the growth is controlled essentially by kinetics, rather than by diffusion processes through liquid or gaseous layers, as in LPE and VPE respectively; for this reason it is possible to have epitaxial layers with almost any predetermined doping or composition profile in the growth direction. Moreover the relatively low rate allows for a very accurate control of the layer thick-
544 ness. T h e r e f o r e , it seems to be a very suitable t e c h n i q u e w h e n high spatial r e s o l u t i o n in the g r o w t h d i r e c t i o n is required. F u r t h e r advantages can be t a k e n o f the writing s7 and s h a d o w i n g s8 capabilities, w h i c h allow f o r the precise d e p o s i t i o n o f layer s t r u c t u r e s w i t h specific material c o m p o s i t i o n at specified l o c a t i o n s o n a given substrate. This is c o n s i d e r e d interesting for the f a b r i c a t i o n o f multilevel integr a t e d o p t o - e l e c t r o n i c circuits, and in this c o n n e c t i o n b o t h the in-situ f o r m a t i o n o f n o n - a l l o y e d o h m i c c o n t a c t s o n n÷-GaAs layers 89 and single-crystal A1-Schottky-barriers o n GaAs 9 0 have r e c e n t l y b e e n rep o r t e d . These results m a y indicate t h a t the f a b r i c a t i o n o f sophisticated devices during a single U H V - r u n is conceivable.
REFERENCES
1. 2. 3.
4. 5.
6. 7.
8. 9. 10. 11. 12. 13.
L.R. DAWSON -- Progress in Solid State Chemistry, H. Reiss & J.O. Mc Caldin eds, Pergamon Press, N.Y., 1972, Vol. 7, pp. 117-139. E.A. GIESS, R. GHEZ -- Epitaxial Growth, J.W. Matthews ed., Academic Press, N.Y., 1975, part A, pp. 183-209. D.W. SHAW -- Ibidem, pp. 89-106. L.L. CHANG, R. LUDEKE -- Ibidem, pp. 37-69. A.Y. CHO, J.R. ARTHUR -- Progress in Solid State Chemistry, J.O. Mc Caldin & G. Somorjai eds., Pergamon Press, N.Y., 1975, Vol. 10, pp. 157191. F.C. FRANK, J.H. VAN der MERWE --Proc. Roy. Soc., A198, 205, 1949. H. KRESSEL --J. Electr. Mat., 4, 1081, 1975. J.W. MATTHEWS - Epitaxial Growth, J.W. Matthews ed., Academic Press, N.Y., 1975, part B, pp. 560-607. M.S. ABRAHAMS, CJ. BUIOCCHI, G.H. OLSEN -- J. Appl. Phys., 46, 4259, 1975, A.G. SIGAI, C.J. NUESE, R.E. ENSTROM, T.J. ZAMEROWSKY -- J. Elec. trochem. Soc., 120, 947, 1973. CJ. NUESE, A.G. SIGAI, J.J. GANNON, T J . ZAMEROWSKY --J. Electr. Mat., 4, 51, 1975. M. ETTEMBERG, C.J. NUESE, J.R. APPERT, J.J. GANNON, R.E. ENSTROM --J. Electr. Mat., 4, 37, 1975. M.S. ABRAHAMS~ L.R. WEISEMBERG, C.J. BUIOCCHI, J. BLANC -- J. Mat. Sc£, 4, 223, 1969.
545 M.S. ABRAHAMS, C.J. BUIOCCHI - J. AppL Phys., 37, 1973, 1969. M. ETTEMBERG -- J. AppL Phys., 45, 901, 1974. CJ. NUESE --J. Electr. Mat., 6, 253, 1977. M. ETTEMBERG, G.H. OLSEN -- J. AppL Phys., 48, 4275, 1977. G.A. ROZGONY, P.M. PETROFF, M.B. PANISH -- J. Cryst. Growth, 27, 126, 1974. 19. R.J. NELSON, R.G. SOBERS -- AppL Phys. Lett., 32, 761, 1978. 20. D.V. LANG, R.A. LOGAN -- Appl. Phys. Let&, 31, 684, 1977. 21. G.H. OLSEN, M.S. ABRAHAMS, T J . ZAMEROWSKY - J. Electrochem. Soc., 121, 1650, 1974. 22. G.H. OLSEN, M.S. ABRAHAMS, C.J. BUIOCCHI, T J . ZAMEROWSKY J. Appl. Phys., 46, 1643, 1975. 23. G.H. OLSEN --J. Cryst. Growth, 31, 223, 1975. 2.4. CJ. NUESE, G.H. OLSEN, M. ETTEMBERG, J J . GANNON, T J . ZAMEROWSKY - - AppL Phys. Lett., 29, 807, 1976. 25. H. KRESSEL, CJ. NUESE, G.H. OLSEN - J . Appl. Phys., 49, 3140, 1978. 26. R.E. NAHORY, M.A. POLLAK, J.C. DE WINTER --J. AppL Phys., 48, 320, 1977. 27. H. NAGAI, Y. NOGUKI - Appl. Phys. Lett., 26, 108, 1975. 28. G.B. STRINGFELLOW - - J. AppL Phys., 43, 3455, 1972. 29. J.J. COLEMAN, N. HOLONYAK, MJ. LUDOWISE - Appl. Phys. Lett., 28, 363, 1976. 30. J J . HSIEH, M.C. FINN, J.A. ROSSI - Proc. of Sump. on GaAs and related compounds, St. Louis 1976, Institute of Physics, London, 1977, pp. 37-44. 31. K. TATEISHI, M. NOGANUMA, K. TAKAHASHI -- Jap. J. AppL Phys., 15, 785, 1976. 32. T. WAHO, S. OGAWA, S. MARUYAMA - - J a p . J. Appl. Phys., 16, 1875, 1977. 33. H. SAKAKI, L.L. CHANG, R. LUDEKE, C.A. CHANG, G.A. SAI-HALASZ, L. ESAKI -- AppL Phys. Lett., 31, 211, 1977. 34. G.A. SAI-HALASZ, R. TSU, L. ESAKI - Appl. Phys. Lett., 30, 651, 1977. 35. R. CHARLTON, G . S . H O B S O N - - IEEE Trans. Electron. Dev., ED-20, 812, 1973. 36. G. SALMER, J. PRIBETICH, A. FARRAYRE, B. KRAMER - J. Appl. Phys., 44, 314, 1973. 37. D.R. DECKER -- IEEE Trans. Electron Dev., ED-21,469, 1974. 38. C. KIM, R. STEELE, R. BIERIG - - Electron. Lett., 9, 173, 1973. 39. R.E. GOLDWASSER, F.E. ROSZTOCZY - AppL Phys. Lett., 25, 92, 1974. 40. R.L. DAWSON --J. Cryst. Growth, 27, 86, 1974. 41. T.GJ. van OIRSCHOT, WJ. LESWIN, PJ.A. THIJS, W. NIJMAN -- J. Cryst. Growth, 45, 262, 1978. 14. 15. 16. 17. 18.
546 42. 43. 44. 45.
46. 47. 48. 49. 50.
51. 52. 53. 54.
55. 56. 57. 58. 59. 60. 61. 62. 63. 64. 65.
H.F. LOCKWOOD, M. ETTEMBERG -- J. Cryst. Growth, 15, 81, 1972. H.F. LOCKWOOD, H. KRESSEL -- ]. Cryst. Growth, 27, 97, 1974. D.L. RODE, B. SCHWARTZ, J.W. DILORENZO - Sol. State Electron., 17, 1119, 1974. T. NOZAKI, M. OGAWA, H. TERAO, H. WATANABE -- Proc. of Syrup. on GaAs and related compounds, DeauviUe, 1974, Institute of Physics, London, 1975, pp. 46-54. H.M. COX, J.W. DILORENZO - Proc. of Syrup. on GaAs and related compounds, St. Louis, 1976, Institute of Physics, London, 1977, pp. 11-22. C.E.C. WOOD -- Appl. Phys. Lett., 29, 746, 1976. J. HALLAIS, D. BOCCON-GIBOD, J.P. CHANE', J.M. DURAND, L. HOLLAN --J. Electrochem. Soc., 124, 1290, 1977. Y. OTA --]. Electrochem. Soc., 124, 1795, 1977. G. MAJNI, G. OTTAVIANI -- Thin Film Phenomena - Interface and Interactions, J.E. Baglin, J.M. Poate eds., The Electrochem. Soc., Princeton, 1978, pp. 293-299. J,P. DUCHEMIN, M. BONNET, F. KOELSCH ]. Electrochem. Soc., 125, 637, 1978. J.P. DUCHEMIN, M. BONNET, F. KOELSCH, D. HUYGHE -- ]. Cryst. Growth, 45, 181, 1978. M.B. PANISH --Proc. IEEE, 64, 1512, 1976. HJ. HOVEL, J.M. WOODALL, W.E. HOWARD - P r o c . of Syrup. on GaAs and related compounds, Boulder, 1972, Institute of Physics, London, 1973, pp. 205. H.J. HOVEL -- Solar Cells, Vol. 11, of Semiconductors and Semimetals, R.K. Willardson, A.C. Beer eds., Academic Press, N.Y., 1975. W.D.JOHNSTON Jr., W.M. CALLAHAN -- J. Electrochem. Soc., 123, 1524, 1976. W.D. JOHNSTON Jr., W.M. CALLAHAN - Appl. Phys. Lett., 28, 150, 1976. R.D. DUPUIS, P.D. DAPKUS, R.D. YINGLING, L.A. MOUDY - AppL Phys. Lett., 31, 201, 1977. Nd. NELSON, K.K. JOHNSON, R.L. MOON, H.A. VANDER PLAS, L.W. JAMES - AppL Phys. Lett., 33, 26, 1978. R.D. DUPUIS, P.D. DAPKUS -- AppL Phys. Lett., 32, 473, 1978. H. KRESSEL, M. E'I~rEMBERG -- ]. Appl. Phys., 47, 3533, 1976. M. ETTEMBERG -- AppL Phys. Lett., 27, 652, 1975. W.T. TSANG -- AppL Phys. Lett., 33,245, 1978. H.C. CASEY Jr. - J . AppL Phys., 49, 3684, 1978. A.C. GOSSARD, P.M. PETROFF, W. WIEGMANN, R. DINGLE, A. SAVAGE --AppL Phys. Lett., 29, 323, 1976. -
547
66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90.
C.M. GARNER, Y.D. SHEN, C.Y. SU, G.L. PEARSON, W.E. SPICER - J . Vacuum Sci. Techn., 15, 1480, 1978. E.A. REZEK, N. HOLONYAK, B.A. VOJAK, H. SHICHIJO - J. Appl. Phys., 49, 69, 1978. N. HOLONYAK, R.M. KOLBAS, R.D. DUPUIS, P.D. DAPKUS -- Appl. Phys. Lett., 33, 73, 1978. A.J. BLAKESLEE --]. Electrochem. Soc., 115, 118, 1971. M. ETTEMBERG, C.J. NUESE - J . Appl. Phys., 46, 3500, 1975. M. ETTEMBERG, G.H. OLSEN, C.J. NUESE - Appl. Phys. Lett., 29, 141, 1976. M.D. MILLER, G.H. OLSEN, M. ETTEMBERG -- Appl. Phys. Lett., 31, 538, 1977. A. MIRCEA, A. MITONNEAU -- Appl. Phys., 8, 15, 1975. G.B. STRINGFELLOW, H.T. HALL - ]. Electrochem. Soc., 123, 916, 1976. M. ALLENSON, S.J. BASS - Appl. Phys. Lett., 28, 113, 1976. J.S. HARRIS, Y. NANNICH, G.L. PEARSON, G.F. DAY -- ]. Appl. Phys., 40, 75, 1969. D.V. LANG, A.Y. CHO, A.C. GOSSARD, M. ILEGEMS, W. WIEGMANN J. AppL Phys., 47, 2558, 1976. K.L. ASHLEY, S.W. BEAL - Appl. Phys. Lett., 32, 375, 1978. A.Y. CHO, L. HAYASHI -- Sol. State Electron., 14, 125, 1971. H.C. CASEY, A.Y. CHO, P.A. BARNES -- IEEE J. Quantum Electronics, QE-11, 467, 1975. T.C. HARMAN - J . Nonmetals, 1, 183, 1973. A. LOPEZ-OTERO - Appl. Phys. Lett., 2 6 , 4 7 0 , 1975. D.L. SMITH, V.Y. PICKARD -- J. Electron. Mat., 5, 247, 1976. R.L. HARTMANN, N.E. SCHUMAKER, R.W. DIXON -- Appl. Phys. Lett., 31, 756, 1977. C.C. SHEN, J.J. H-ISIEH, T.A. LIND - Appl. Phys. Lett., 30, 353, 1977. J.C.C. FAN, C.O. BOZLER, R.L. CHAPMAN - Appl. Phys. Lett., 32, 390, 1978. W.T. TSANG, A.Y. ClIO -- Appl. Phys. Lett., 32, 491, 1978. W.T. TSANG, M. ILEGEMS - Appl. Phys. Lett., 31, 301, 1977. A.Y. CHO -- Appl. Phys. Lett., 33, 651, 1978. A.Y. CHO, P.D. DERNIER --J. Appl. Phys., 49, 3328, 1978.