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The use of ion beams in molecular beam epitaxy
~hinSolidFiZwz.s,80(1981)197-211
197
METELLURGICAL AND PROTECTIVE COATINGS
THE USE OF ION BEAMS IN MOLECULAR
BEAM EPITAXY
R. F. C. FARRO’N Royal Signals and Radar Establishment,
St Andrews Road, Malvern, Worcs. WR14 3PS (Gt. Britain)
Molecular beam epitaxy (MBE) has been developed over the past decade into one of the most versatile techniques for generating thin film semiconductor device structures. This development has been assisted by the use of low energy (0.1-10 keV) ion beams in a number of areas such as substrate surface preparation, film doping, metal-semiconductor contact formation and in situ surface analysis by secondary ion mass spectrometry and Auger depth profiling. Illustrative examples of the applications of ion beams in these areas are given. The use of ion beams in MBE has led to the discovery and exploration of a number of side effects of ion beams on semiconductor surfaces. These include group III metal surface segregation on III-V compound semiconductors, surface n-type channel formation on some III-V and II-VI compound semiconductors and II-VI alloy composition gradation. Some of these effects have been made use of in device structures. Mechanisms responsible for the effects are discussed and future trends in the use of ion beams in MBE are considered.
1. INTRODUCTION
Molecular beam epitaxy (MBE) has been developed I-3 over the past decade into one of the most versatile techniques for generating thin film semiconductor device structures. This development has been assisted by the use of low energy (0. l10 keV) ion beams in a number of areas such as substrate surface preparation, film doping, metal-semiconductor contact formation and in situ surface analysis by secondary ion mass spectrometry (SIMS) and Auger depth profiling. In Section 2, specific illustrative examples of the applications of ion beams in these areas are given. The wide range of powerful in situ surface analysis techniques, which is built into MBE systems, combined with ex situ characterization of MBE-grown films has led to the discovery and exploration of a number of side effects of ion beams on semiconductor surfaces. These include group III metal segregation on III-V -*Paper presented at the 3rd International Conference on Ion and Plasma Assisted Techniques, Amsterdam, The Netherlands, June 30-July 2, 1981. 0040-6090/8 1/OOOO-0000/$02.50
0 Elsevier Sequoia/Printed in The Netherlands
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R. F. C. FARROW
compound semiconductor surfaces, surface n-type channel formation on some IIIV and II-VI compound semiconductors and II-VI alloy composition gradation. These effects are discussed in Section 3. Some of the effects have been used in device structures. Future trends in the use of ion beams in MBE are considered in Section 4. 2.
APPLICATION
OF ION BEAMS IN MOLECULAR
BEAM EPITAXY
2.1. Semiconductor surface preparation A prerequisite for the preparation of epitaxial films of semiconductors, metals or insulators of high crystallographic quality on semiconductor substrates is the removal of surface impurities prior to film deposition. Carbon and oxygen impurities are invariably present on the surfaces of semiconductors after sample loading and pumpdown to a ultrahigh vacuum ambient and are known to affect the morphology of homoepitaxial GaAs 3, InP ’ and silicon adversely. These impurities probably originate from a variety of sources ranging from wet chemical treatments including polishing and etching to adsorption of atmospheric gases such as CO,, CO and 0, prior to sample loading. Careful preparative treatments followed by heating in ultrahigh vacuum significantly reduce the impurity levels as judged by Auger electron spectroscopy (AES) and can often3 result in a coverage of impurity atoms of less than 1% of a monolayer. Congruent thermal etching of the semiconductor is necessary for this process to be effective, however, and whilst this is achievable for silicon and GaAs it is difficult2 to achieve sufficient congruent etching rates for the more unstable compound semiconductors such as InP and InSb. In these cases, low energy (500 eV) Ar+ ion bombardment has been found’ to be highly effective in reproducibly producing impurity-free surfaces as judged by AES. Alternating ion bombardment with annealing is necessary to permit reordering of the semiconductor surface disordered by the ion beam. A particularly striking recent example of the success of this technique is in the preparation of the first epitaxial metastable phase by MBE: c+Sn on InSb4. This example is particularly interesting since it demonstrates how critical the removal of surface contamination is to the nucleation of the metastable phase. Figure 1 shows a series of Auger spectra recorded before and after Ar+ bombardment and annealing treatments of an InSb(OO1)surface in an MBE system. After loading, the surface has considerable oxygen and carbon contamination shown by the presence of O(515 eV) and C(272 eV) peaks in spectrum a. After a 500 eV Ar+ dose of 2 x lOi cm- ’ and a 1 h anneal at 200 “C there are (see Fig. l(b)) significant reductions in the oxygen and carbon peaks and corresponding increases in intensity of the indium and antimony doublets resulting from removal of the impurity overlayer from the InSb surface. With further repeats of this cycle the Auger spectrum d, recorded at a higher resolution, shows that no oxygen and carbon (a coverage of less than 0.5% of a monolayer) remain. Reflection electron diffraction studies of this surface confirmed4 that it was ordered with a (2 x 4) or C(2 x 8) surface reconstruction. Condensation of a beam of tin atoms on this surface resulted in the growth of an epitaxial a-Sn film at substrate temperatures well above the cl-+B transformation temperature of about 13 “C. In contrast, the condensation of a beam of tin atoms onto a contaminated InSb surface exhibiting an Auger spectrum with a coverage of carbon and oxygen of greater than 1% of a monolayer resulted in the growth of a polycrystalline film of p-
199
ION BEAMS IN MBE
Sn. Clearly, intimate contact between the tin film and the isomorphous closely lattice-matched InSb substrate is necessary for metastable c+Sn growth through pseudomorphic growth. This effect has also been demonstrated for ol-Sn on ionbombarded annealed CdTe surfaces and has opened up a new area of epitaxy with potential device applications in IR detection by photovoltaic or photoconductive response in the a-Sn narrow gap semiconducting film.
&anneals
l&b
surface
I
E(ev)
500
I
I
400
300
]C
I 200
Fig. 1. Auger electron spectra of an InSbfOOl) surface: curve a, before in situ cleaning; curve b, after 500 eV Arf ion bombardment at 1.0 pA err-’ for 1 h followed by a 1 h anneal at T = 200 “C; curve c, after a further cycle of the above treatment; curve d, after one-third of a cycle of the treatment, the spectrum is recorded at a higher resolution than spectrum a-c.
The technique of ion bombardment and annealing has also been used as a precursor to homoepitaxial InSb growth by MBES and also to homoepitaxial InP ’ growth. However, the side effect of group III metal segregation which results from the ion bombardment of III-V compound semiconductors6’7 (see Section 3.1) is particularly severe for InP and results in indium island formation at ion doses of less than lO”j cm-’ unlike InSb and the other III-V compounds which require doses greater than 101’ cm-’ for island formation. Low energy ion bombardment of silicon surfaces followed by annealing at 900 “C or below is commonly used’ to prepare clean ordered surfaces prior to MBE growth of silicon or silicide films in ultrahigh vacuum. This process is preferred for some applications since the alternative cleaning technique of briefly flashing the
200
R. F. C. FARROW
silicon to 1200 “C can cause significant diffusion of dopants within the wafer and diffusion of impurities to and from the wafer. Yamada et al.’ have made a careful study of residual damage to ion-bombarded wafers and concluded that Ne+ is the preferred species for ion bombardment cleaning of silicon since residual surface damage from 4 keV Nef can be annealed out after 15 min at 720 “C. 4 keV Ar+ bombardment results in more tenacious damage and can induce gas bubble formation whilst 4 keV Kr+ bombardment requires an anneal at about 900 “C for full recovery of surface order. Ar+ ion bombardment has been used by Bean” to prepare clean ordered sapphire and spine1 substrates for MBE growth of silicon. In this case the substrates are insulators and room temperature exposure to Ar+ ions leads to surface charging which reduces the sputtering rate and renders ion bombardment ineffective in removing impurities. At substrate temperatures of about 1000°C the substrate conductivity is sufficient to avoid surface charging and clean well-ordered surfaces (as judged by 1.8 MeV He+ channelling) are achieved. Silicon films grown on sapphire substrates prepared in this way reach the island coalescence stage earlier than on substrates cleaned simply by heating to 1400 “C. Surface dissociation of Al,O, is, in fact, significant at 1400°C and is likely to result in changes in near-surface stoichiometry with consequent degradation in surface smoothness and crystallinity. 2.2. Ionized beam doping The versatility of MBE as a preparation technique for device structures arises from the ability to control the level and type ofdopant with great accuracy in vertical architectures. For GaAs there are several suitable n-type dopants which are readily incorporated into the growing film when supplied as neutral atoms from Knudsen effusion sources. However, for p-type doping, beryllium is the only dbpant which approaches ideal requirements when supplied as a neutral atom. Beryllium, however, requires special handling procedures because of its high toxicity. Other dopants such as zinc and cadmium have very small sticking coefficients, as neutral atoms, at normal growth temperatures (500-600 “C) and are not incorporated into the GaAs film. In 1975 Naganuma and Takahashi” proposed a solution to this problem. They showed that when zinc was ionized and when it impinged at an energy of not more than 1.5 keV onto the GaAs surface during MBE growth at 55s 580 “C a significant incorporation of zinc acceptors in the film was achieved. The surface morphology remained smooth at zinc doping levels of up to 5 x 10” crnv3. Reasonable values of both the acceptor mobility and the photoluminescence intensity were observed, and Naganuma and Takahashi suggested simple ion burial as the mechanism for zinc incorporation. Bean and Dingle” later confirmed the effect and supported this mechanism for zinc incorporation but showed that asgrown epitaxial layers retained a substantial amount of residual radiation damage. Significant improvements in photoluminescence intensity and acceptor mobility values were demonstrated for post-growth anneals. The technique offers prospects for direct ionized beam writing of laterally defined dopant structures. However, the spreading of lateral dobing profiles by ion surface migration is an unknown factor at this time. For silicon growth by MBE, ionized beam doping has recently been used to overcome a number of problems inherent in doping with neutral species. For
ION BEAMS IN MBE
201
example, aluminium, antimony and gallium have a tendency to surface segregation during growth which tends to distort the shape of the intended doping profiles13. Sugiura l4 has shown that, at least for antimony, this segregation can be overcome by using low energy (130-1000 eV) ionized antimony as the dopant beam during growth at 850 “C. All arriving Sb+ ions are incorporated as donors in the film at this temperature, a doping range of 10’6-1020 cme3 has been demonstrated and Rutherford backscattering data confirmed the absence of profile distortion arising from surface segregation. Clearly the segregation effect is eliminated by shallow (about two silicon atomic layers) implantation of Sb+. Ota” has similarly shown that As+ ion doping (ion energy range, 400-800 eV) during silicon MBE can be used to produce arbitrary doping profiles over the range 10’4-10’g cme3. The quality of As+-doped epitaxial silicon films has proved adequate for majority carrier devices such as millimetre wave p-i-n and varactor diodes but to date no studies have been reported on minority carrier properties or on the distribution of electron traps in ionized-beam-doped films. 2.3. Metal-semiconductor conta.ctformution Metal film epitaxy from atomic beam sources has been developed by Farrow and coworkers16-r8 and others3 for the preparation of metal-semiconductor contacts in situ in an MBE system. For example, the growth of an epitaxial aluminium film at room temperature onto MBE-grown GaAs forms part of a process recently developed by Schneider and Cholg for fabricating arrays of millimetre wave diodes. In this process the epitaxial aluminium film is deposited over the entire GaAs wafer, and post-growth lithography and ion etching are used to define discrete diodes. In the future it will be desirable to eliminate the timeconsuming ex situ lithography and ion-etching process used to define structures such as these. The liquid metal field emission ion sources originated by Clampitt et aL2’ and developed by Pruett and Jefferies 21 offer considerable potential for direct micromachining of metal patterns such as those required for diode arrays. In addition, under some operating conditions the sources are known to emit charged microclusters of metal, suggesting the possibility of patterned metal deposition if the process can be sufficiently well controlled. Seliger et al.” have demonstrated the micromachining of 1000 A lines in gold films 400 A thick on silicon using a focused 57 keV Ga+ ion beam generated from a gallium field emission ion source. Farrow et al.* have studied the microstructure of aluminium and gold films deposited from an unrestricted nozzle type of field emission source in an MBE system. More recently the microstructure and contact characteristics of tin and silver films deposited onto semiconductors from restricted nozzle” sources have been studied in this laboratory 23. In both types of source, operation at low emission currents led to generation of both metal ions and charged microdroplets less than 1000 A in diameter. Auger spectra recorded of semiconductor surfaces exposed to beams generated from gold and tin field emission sources operated at low emission currents (less than 20 PA) confirmed the surface-cleaning effect of the ion component of the beam. Figure 2 shows Auger spectra recorded of an InP(OO1) surface before and during exposure to a gold source operated at 20 pA and 10.5 keV. Spectrum a shows considerable initial oxygen and carbon contamination of the InP surface. Exposure of the surface to the source for 30 min yielded spectrum b which shows considerable
202
R. F. C. FARROW
reductions in the surface impurity peaks and increases in the native species In and P. A further exposure of the sample to the beam for 30 min (see spectrum c) led to further reductions in the impurity peaks and to further enhancement of the peaks due to indium and phosphorus. In addition, the progressive increase in the gold Auger peak is evident. These spectra and parallel scanning electron microscopy studies’* confirm that the semiconductor surface experiences simultaneous ion cleaning, microcluster deposition and probably some shallow Au+ implantation. Increasing the emission current led to increases in both the rate and the size of gold microdroplet condensation. Similar results have recently been observed for a restricted nozzle tin ion source. Metal-semiconductor contacts formed from continuous films of field-emission-deposited Au/InP and Ag/InP show reduced barrier heights for n-type InP but, for silver on p-type InP, near-ideal classical Schottky barrier behaviour has been observed. Since barrier height control and optimization is a key requirement for several semiconductor devices, this result is particularly significant and further exploration of the effects of deposition conditions on barrier height are in progress. Preliminary studies of the exposure of InSb surfaces to the beam from a restricted nozzle tin ion source confirm the existence of p-Sn nuclei which result from the condensation and freezing of microdroplets of liquid tin. Barrier height measurements are in progress.
C
bw
II
c
I__ so5
272
j_
a
500 -
400
Auger
300
electron
200
energy
100
eV
Fig. 2. Auger electron spectra of an InP(OO1) surface before and after exposure to a beam from a gold field emission source operated at an emission current of 20 uA and an extraction potential of 10.5 keV: curve a, spectrum recorded before exposure of surface to a gold field emission source; curve b, spectrum recorded after exposure of the surface to a gold field emission source for 30 min; curve c, spectrum recorded after exposure of the surface to a gold field emission source for 60 min. The progressive decreases in the surface impurity peaks and the increases in the indium, phosphorus and gold peaks should be noted.
203
ION BEAMS IN MBE
At this time it may be concluded that if the single-ion and microdroplet components of field emission sources can be separated and controlled by the addition of an ion lens system then both modes of operation would have applications in microcircuit technology: micromachining or shallow implantation for the ion emission mode and selected area contact deposition for the microdroplet emission mode. This area of development is particularly active, at present, not .only in Great Britain but also in the U.S.A. and Japan. Studies of the mechanism of ion and ion cluster emission from the sources are continuing in Great Britain24925 and France26. 2.4. Ion beams for in situ analysis in molecular beam epitaxy The two areas of in situ diagnosis in MBE where ion beams have been used are in static SIMS and in Auger depth profiling. In SIMS a primary ion beam probe is used to eject secondary ions from the sample surface which are then mass analysed. In the static mode of SIMS the primary ion beam probe has a sufhciently low current density that the time for mass analysis of the secondary ions is much less than the sputtering time for one monolayer. In this mode SIMS is a surface-sensitive technique capable of providing information on surface composition and chemistry. Ploog and Fischer2’ used the static SIMS technique to probe in situ the surface of an MBE-grown tin-doped film of GaAs. The main reason for this work was to establish the growth temperature range over which tin segregated to the surface of the growing film. Positive secondary ion mass spectra were recorded at an erosion rate of about 50 A h- ’ using a primary Ar+ ion current of 10m7A and a beam diameter of 5 mm. Although the absence of an ion extraction and focusing lens between the sample and mass spectrometer greatly reduced the signal-to-noise ratio in the mass spectra, the data showed segregation of tin (for films grown at temperatures higher than 490 “C) in the outer 10-20 A of the film at a level three orders of magnitude greater than in the bulk. Parallel AES studies of the as-grown surface confirmed this surface segregation effect and by sequentially Ar+ ion sputtering the top 30 A of the film away it was established that the tin segregation was confined to within 20 A from the surface. The static SIMS technique has also been used by Dowsett et al.28 to evaluate residual surface impurity levels on InP surfaces prior to MBE. In this work the high sensitivity of a well-designed static secondary ion mass spectrometer to surface impurities was demonstrated. Surface impurities at levels of the order of 100 ppm were clearly evident on Ar+-ion-bombarded InP(OO1) surfaces which appeared clean as judged by AES (impurity detection limit, about 1% of a monolayer). Auger depth profiling has been used extensively3 in analysing the composition of multilayer structures grown by MBE (see Section 3.2 and Fig. 5). 3.
SIDE EFFECTS OF INERT GAS ION BEAMS ON SEMICONDUCTOR
SURFACES
3.1. III-V compounds In studying the microstructure of metal films deposited by MBE onto Ar+bombarded and annealed InP(OO1)surfaces Farrow and coworkers6,7,18 noticed the presence of small ( less than 1000 A in size) islands of indium metal at the deposited metal-InP interface. Careful experiments revealed that these were due to the effect of
204
R. F. C. FARROW
Fig. 3. Scanning electron micrograph of an InP(OO1)surface bombarded with Ar+ ion3 at 300 K and not subjected to any annealing treatment after bombardment (ion energy, 500 eV; ion dose, 1.3 x 10” cme2). The small (about 150 A in diameter) particles covering the surface were confirmed (see text) as metallic indium. The 1000 A particle in this figure is a dust particle used to provide a focus.
Ar+ ion bombardment and not to the effects of annealing or to the electron beam. This effect was subsequently studied in greater detail using a variety of inert gas ions and other III-V semiconductors. Figure 3 is a scanning electron micrograph of an ion-bombarded InP surface not subjected to any annealing treatment. The surface is covered with particles about 150 A in diameter which transmission electron diffraction studies’ later confirmed were metallic indium. The 1000 A particle in this figure is a dust particle used to provide a focus. Lower ion doses produced smaller islands but at a similar density. Figure 4 is a scanning electron micrograph of an InP surface subjected to an Ar+ ion dose of 1.5 x 1016 cmW2and to a post-bombardment anneal at 250 “C for 1 h. The distribution and size of the islands are quite different in this case. It seems likely that at 250 “C the liquid indium islands are mobile on the clean surface of the semiconductor and coalesce to form the larger islands shown in Fig. 3. This idea is supported by measurements of the density and size of the islands on the assumption of hemispherical cap-shaped nuclei. Direct comparisons between the effect on InP and the other III-V compounds InSb, GaAs and GaP were made by simultaneously exposing the samples to the same ion beam. This confirmed that the effect was much more pronounced in InP than in the,other semiconductors. For InSb, for example, ion doses well in excess of 10” cm-’ were required before the effect was detectable. Similar doses of Kr+ and Xe+ ions for InP produced islands of similar size and density to those for Ar+ although Auger spectra recorded
ION BEAMS IN MBE
205
Fig. 4. Scanning electron micrograph of an InP surface subjected to Arf ion bombardment (ion energy, 500 eV; ion dose, 1.5 x 1016cm-‘) and to a post-bombardment anneal at 250 “C for 1 h. The presence of larger indium islands than in Fig. 3 should be noted. The increase in island size is attributed to surface migration and coalescence of small islands of liquid indium (see text).
immediately after sputtering showed a large enhancement of surface phosphorus. Indeed for Arf ion bombardment and annealing the electron diffraction patterns indicated’ a well-ordered phosphorus-stabilized (2 x 4) surface reconstruction. The contribution to the diffraction pattern from the small surface coverage of hemispherical indiiun islands is in any case negligible. This suggests that differential sputtering of the group V element is not the dominant underlying mechanism for the effect since on momentum transfer grounds we would expect considerable general group V depletion especially for Arf bombardment of InP. Most probably, the mechanism responsible is simply the breaking of III-V bonds with subsequent surface migration and coalescence of the group III metal under the ion beam. Not only is the group III metal more thermodynamically stable than the III-V compound but in addition it is more probable that metal segregation occurs via a simple first-order surface reaction than that the III-V bond is reformed via the more complex III-[V dimer] association reactiot?. The excess of group V atoms resulting from the metal segregation reaction tends to drive the rest of the semiconductor surface towards a group-V-rich surface phase. The post-bombardment anneal temperature therefore determines the phase of the rest of the surface. For example, if anneal temperatures higher than the congruent evaporation temperature limit2 are employed, the surface not occupied by group III metal will be driven towards a group-III-rich surface phase, e.g. C(8 x 2) for InSb annealing at 400°C 5 (about
206
R. F. C. FARROW
100°C above the congruent evaporation limit). Independent confirmation of the formation of free group III metal islands on III-V compounds under a variety of ion bombardment, heating and film growth conditions has been presented by Olivier2g who used the technique of low energy electron loss spectroscopy to detect plasmon losses in the islands. The presence of these islands, which may form localized ohmic contacts beneath an overdeposited metal film, is believed’ to be a contributory factor to the greatly reduced barrier heights observed for a range of metals deposited onto ionbombarded and annealed InP and GaAs. The islands, however, seem to have little effect on the structure and electrical properties of overdeposited InP 3o and GaAs 31 films. They may, however, play a role in generating the cone structures seen by Farrow and others on InP surfaces subjected to extensive Ar+ ion milling during Auger depth profiling. Such structures invariably show indium metallic precipitates near the cone tips. A second and equally important side effect of Ar+ ion bombardment of some III-V compounds is the formation of an n-type surface-conducting channel, probably resulting from electrically active point defects arising from ion-induced lattice damage. This effect has been reported for p-type InAs 32. Also Farrow3’ and Kreutz34 have observed a similar effect for p-type InSb and semi-insulating InP. For p-type InAs (net hole concentration, 2 x lOi cmm3; 77 K) an Ar + (140 eV) dose of only 6.25 x 1013 cm-’ (= 10 PC cmm2) was adequate32 to generate a surface channel of effective depth 0.28 pm with a volume density of electrons in the channel of about 2 x 1016 cm- 3. The mobility of electrons in the surface channel was comparable (2.3 x lo4 cm2 V-r s- ‘) with the bulk mobility value, indicating that few extrinsic scattering defects were produced by the low energy ion bombardment. Thermal annealing at temperatures above 70°C tended to remove the n-type channel and to restore p-type conductivity, presumably as the damage-induced point defects annealed out. The use of a higher energy (10 keV) Ar+ beam improved the thermal stability of the n-type channel but led to lower electron mobilities in the surface channel. Clearly, in this case, extrinsic scattering defects were created by the ion beam. The effect of intermediate ion energies and other inert gas ion species were not explored and it seems likely that an adequately stable surface channel with good electron mobility could be formed at ion energies lower than 10 keV. This would, in principle, permit the technique to be used for fabrication of an IR photodiode array with some on-chip signal processing. We observed a similar surface conversion effect for p-type InSb and it is significant that Fujisawa 35 has recently reported surface conversion of p-type InSb by proton or silicon ion bombardment or by Nd-yttrium aluminium garnet laser irradiation. Clearly, for InAs, InSb and InP, lattice damage induces a near-surface defect state or series of states with associated levels close to the conduction band minimum. The nature of these defects is at present unclear although it is tempting to associate them with the near-surface defects, responsible for Fermi level pinning, which result from metal deposition on the (110) cleavage faces of these semiconductors. This Fermi level pinning occurs at near midgap for GaAs 36 but near the conduction band minimum for InP 36 and InSb 34. This trend is consistent with the known effects of ion-induced lattice damage on surface conductivity. Proton bombardment of GaAs produces semi-insulating material (deep-level near-
207
ION BEAMS IN MBE
midgap damage-induced electron trap states) whereas proton bombardment of InSb produces n-type surface conductivity (shallow damage-induced donor-like states near the conduction band minimum). The identity of the lattice defects responsible for Fermi level pinning is still the subject of controversy. However, group III vacancies are a distinct possibility in view of the strong tendency of group III atoms to form localized metal clusters when displaced from their lattice sites. 3.2. II-VI compounds and alloys Recent studies3’ on MBE growth of CdTe films on CdTe, InSb and Hg, _,Cd,Te substrates have revealed several striking differences in the effects of inert gas ion beams on the surfaces of these semiconductors. For CdTe, scanning and transmission electron microscopy studies of Ar+-bombarded and milled surfaces show smooth featureless surfaces with a complete absence of segregation of either phase. In addition, AES studies of such surfaces show no evidence for preferential sputtering of either cadmium or tellurium. This absence of ion-beam-induced artefacts results in a planar front during Auger and ion microprobe depth profiling using Ar + ions. Figure 5 shows an Auger depth profile of a CdTe epitaxial film 7500 A thick grown by MBE on InSb(OO1) at 120°C. The abrupt interface (with a transition width of about 100 A) between CdTe and InSb is an indication both of the absence of surface topographical modifications during CdTe ion milling and of the absence of interdiffusion at the low growth temperature. ParallelSIMS profiles have confirmed3’ this abrupt interface. Whilst no compositional or topographical effects of CdTe ion bombardment were observed we found that ion bombardment of high resistivity CdTe substrates induces a surface n-type channel. Measurements of the channel depth and conductance are in progress and it is too early to speculate on the mechanism responsible for this effect. Auger signal
t
;te~,.lb*-.........._. ....-....-*........ I-
p”-Q
- . . . . . . . - . . . . - . . . . . . . . . . . . . . . . 1: :,;
:’
/H TRANSITION
ZONE
Fig. 5. Auger depth profile ofan MBE-grown CdTe film on an InSb(OO1)substrat@r+ ion beam at 1 keV and 40 pA cm-‘). The profile shows the unnormalized peak heights for cadmium(O), tellurium ( n), indium (0) and antimony (A) as functions of depth for the last 3500 A of a CdTe film 7500 A thick (transition zone, 100 A).
The effects of ion bombardment and ion beam milling of the narrow gap semiconducting alloy Hg, _,Cd,Te contrast strongly with their effects on CdTe in that gross compositional, topographical and structural changes occur for
208
R. F. C. FARROW
Hg,_,Cd,Te. Table I shows compositional data for Hg&d,,,Te wafers before and after exposure to an Ar+ ion beam (energy, 5 keV; 12pA cm-‘). The data were derived from an in situ study of surface composition using the technique of X-ray photoelectron spectroscopy (XPS) and refer to the average composition of an area about 5 mm x 5 mm of the top 15-20 A of the sample surface. After Ar+ ion milling, the surface composition is significantly depleted in mercury and enriched in cadmium (i.e. x for the surface is greater than x for the bulk). The effect appears to have a dominant underlying thermal mechanism, possibly transient surface heating, since the depletion is of the volatile end member (HgTe) of the alloy rather than of the two lighter elements (cadmium and tellurium) which might be expected on solely momentum transfer grounds. Congruent loss of HgTe from the ahoy is known3* to occur as a result of heating the alloy at temperatures in the range 300-4OO”C. Support for the thermal mechanism comes from the observation of greater HgTe depletion at higher power denSities of the ion beam with the ion beam energy kept constant. The very low thermal conductivity of Hg, _,Cd,Te also tends to favour surface heating despite good sample heat sinking and cooling to - 180 “C (see Table I). Transmission electron diffraction studies3’ of Ar+-milled foils of Hg, _,Cd,Te have revealed considerable structural changes in the surface layers resulting from ion milling. These include inert gas bubble or microvoid formation and phase separation of the alloy into microdomains of HgTe and CdTe. Electrical measurements of ion bombarded Hg,_$d,Te wafers indicate p- to n-type conversion at ion doses much less (less than 1016 cmT3) than those required to give significant changes in the x value of the alloy from its bulk value. This effect has also been observed by Solzbach and Richter4’ and opens up the prospect of diode formation by controlled ion bombardment as for InAs and InSb. TABLE I X-RAY PHOTOELECTRON SAMPLE~SUBJECTEDTO
SPECTROSCOPY DATA ON THE SURFACE IONMILLING(~ keV;l2pAcm-‘)
COMPOSITION
OF AN
Hg,,,Cd,,,Te
Ar+
Composition (at.%)
Before ion milling 40 min; 25 “C 15 min; - 180°C
Hg
Cd
Te
C
0
23.1 34.2 35.6
1.8 18.8 18.4
14.5 47 46
41.7 0 0
12.9 0 0
“The bulk composition of the Hg,,,Cd,,,Te
sample is 4O%Hg-lO%Cd-50%Te.
4. DISCUSSION The use of ion beams in &iBE has made a significant contribution to the development and application of the technique. This is illustrated by the particular examples given in Section 2. Several points are, however, worth highlighting. Metastable phase formation by epitaxy, initiated by the a-Sn work described in Section 2.1, is a field in which reproducible surface cleaning is critical to the growth of the metastable phase. This area will probably develop rapidly in the near future. For silicon MBE the current trend is towards submicron film epitaxy over silicon wafers 2 or 3 in in diameter with an emphasis on the lateral uniformity of film
ION BEAMS IN MBE
209
thickness and doping level. The current availability of inert gas ion beams of uniform intensity in this size regime will permit reproducible wafer cleaning prior to epitaxy. In addition the use of ionized beam doping in silicon MBE shows considerable promise in achieving arbitrary doping profiles. It may also provide a route to selected area doping. However, a considerable amount of ground work needs to be done first in assessing the effects of ionized beam doping on minority and majority carrier properties. Deep level transient spectroscopy and photoluminescence studies of the material should enable the distribution of carrier traps to be explored as a function of doping and annealing conditions, leading eventually to optimization of these conditions. The field emission sources discussed in Section 2.3 show considerable promise both as high brightness sources for micromachining and as metal deposition sources for in situ metal-semiconductor contact deposition. The ultrahigh vacuum compatibility of these sources, which produce a negligible gas load, combined with their compactness makes them highly attractive for applications in MBE. Here also, much work remains to be done in relating the contact electrical properties of fieldemission-deposited metals to the deposition conditions for a variety of semiconductors. Controllable separation of the single-ion and microdroplet components of the beam would certainly add to the versatility of the sources. The side effects of inert gas ion beams on semiconductors are varied and complex, as illustrated by the examples given in Section 3. Some of them appear to have useful applications in device structures. For example, the p- to n-type surface conversion effect has been used to make junction field effect transistor structures in InAs and may provide a means of making similar structures in InSb and CdTe. Low dose Ar+ ion bombardment of n-type Hg, _,Cd,Te followed by in situ ‘metal deposition under clean controlled conditions in an MBE system may provide a convenient way of making metal-n+-n contacts to the alloy. Such contacts are known to act as minority carrier (hole) accumulation contacts with beneficial effects on photoconductive and photovoltaic device performance41~42. The formation and stability of ion-bombardment-induced diodes in the alloy is also worthy of investigation in view of the considerable current interest41 in far IR photodiode arrays in Hg, _,Cd,Te. 5.
CONCLUSIONS
The use of ion beams in MBE has contributed significantly to the development and versatility of the technique. Much work, however, remains to be done in establishing and understanding the mechanisms of the electrical and structural effects which accompany exposure of semiconductor surfaces to low energy ion beams. An improved understanding of these effects should enable full advantage to be taken of the techniques of ionized beam doping, ion-beamqnduced surface-type conversion and field emission deposition of contacts in device fabrication by MBE. ACKNOWLEDGMENTS
The author is grateful to A. Christie, Loughborough Consultants, and to Hugh Bishop, U.K. Atomic Energy Authority, Harwell, for the XPS studies of the alloy
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I-Ig, _,Cd,Te. The work of D. Sykes, Loughborough Consultants, on Auger depth profiling of MBE-grown CdTe/InSb heterostructures is also acknowledged with thanks. REFERENCES 1 A. Y. Cho and J. R. Arthur, Prog. SolidState Chem., lO(1975) 157. 2 R. F. C. Farrow, in E. Kaldis and H. J. Scheel (eds.), Crystal Growth and Materials, Vol. 1, NorthHolland, Amsterdam, 1977, Chap 1.7. 3 K. Ploog, in H. C. Freyhardt (ed.), Crystals-Growth Properties and Applications, Vol. 3, Springer, Berlin, 1980, p. 73. 4 R. F. C. Farrow, D. S. Robertson, G. M. Williams, A. G. Cullis, G. R. Jones and I. M. Young, Proc. 8th Int. Vacuum Congr., Cannes, September 22-26,1980, in Vide, Suppl., I(201) (1980) 109; J. Cryst. Growth, 54 (1981) in the press. 5 K. Ge. S. Ando and K. Sugiyama, Jpn. J. Appl. Phys., 19 (1980) L417. 6 R. F. C. Farrow and A. G. Cullis, ht. Conf. on Solid Films and Surfaces, Tokyo, 1978, unpublished. 7 A. G. Cullis and R. F. C. Farrow, Thin Solid Films, 58 (1979) 197. 8 J. C. Bean and J. M. Poate, Appl. Phys. Left., 37 (1980) 643. 9 I. Yamada, D. Marton and F. W. Saris, Appl. Phys. Lett., 37(1980) 563. 10 J. C. Bean, Appl. Phys. Left., 36(1980) 741. 11 M. Naganuma and K. Takahashi, Appl. Phys. Left., 27 (1975) 342. 12 J. C. Bean and R. Dingle, Appl. Phys. Lett., 35 (1979) 925. 13 J. C. Bean, Appl. Phys. Left., 33 (1978) 654. 14 H. Sugiura, 11th Int. Conf. on Solid State Devices, Tokyo, 1979, Paper C-3-6, in Digest of Technical Papers, unpublished. 15 Y. Ota, 11th Int. Conf. on Solid State Devices, Tokyo, 1979, Paper C-3-5, in Digest of Technical Papers, unpublished., 16 R. F. C. Farrow, A. G. Cullis, A. J. Grant and J. E. Pattison, J. Cryst. Growth, 45 (1978) 292. 17 R. F. C. Farrow and G. M. Williams, Thin SolidFilms, 55 (1978) 303. 18 R. F. C. Farrow, A. G. Cullis, A. J. Grant, G. R. Jones and R. Clampitt, Thin SolidFilms, 58 (1979) 189. 19 M. V. Schneider and A. Y. Cho, Proc. Cornell University Workshop on Molecular Beam Epitaxy, October 21-22,1980, in J. Vat. Sci. Technol., (July-August, 1981) in the press. 20 R. Clampitt, K. L. Aitken and D. K. Jeffries, J. Vat. Sci. Technol., 12 (1975) 1208. 21 P. D. Pruett and D. J. Jefferies, Inst. Phys. Co@ Ser. 54 (1980) Chap. 7, p. 316. 22 R. L. Seliger, J. W. Ward, V. Wang and R. L. Kubena, Appl. Phys. Left., 34 (1978) 310. 23 R. F. C. Farrow, P. Tapstev and D. G. Coates, unpublished, 1981. 24 A. J. Dixon, J. Phys. D, 12(1979) L77. 25 A. R. Waugh, J. Phys. D, 13 (1980) L203. 26 P. Sudraud, C. Colliex and J. De Walle, J. Phys. Left., 40 (1979) L207. 27 K. Ploog and A. Fischer, J. Vat. Sci. Technol., 15 (1978) 255. 28 M. G. Dowsett, R. M. King and E. H. C. Parker, Appl. Phys. Left., 31(1977) 529. 29 J. Olivier, Proc. 8th Int. Vacuum Congr., September 22-26, 1980, Cannes, in Via’e, Suppl., I (21) (1980) 583.
P. W. Sullivan, R. F. C. Farrow, G. R. Jones and C. R. Stanley, Proc. 8th Int. Symp. on GaAs and September 22-24,1980, in Inst. Phys. Conf. Ser. 56 (1981)Chap.l, p.45. 31 A. Y. Cho, J. V. Di Lorenzo, B. S. Hewitt, W. C. Niehaus, W. 0. Schlosser and C. Radice, J. Appl.
30
Related Compounds,Vienna, Phys., 48 (1977) 346.
32 33
34 35 36 37
M. F. Millea, A. H. Silver and L. D. Flesner, Thin Solid Films, 56 (1979) 253. R. F. C. Farrow, unpublished, 1980. E. W. Kreutz, personal communication, 1980. I. Fujisawa, Jpn. J. Appl. Phys., 19 (1980) 2137. W. E. Spicer, P. W. Chye, P. R. Skeath, C. Y. Su and I. Lindau, J. Vat. Sci. Technol., I6 (1979) 1422. R. F. C. Farrow, G. M. Williams, G. R. Jones and I. M. Young, 1st European Workshop on Molecular Bean Epitaxy, Stuttgart, March 30-April I, 1981; submitted to Appl. Phys. Left.
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R. F. C. Farrow, G. R. Jones, G. M. Williams, P. W. Sullivan, W. J. 0. Boyle and J. T. M. Wotherspoon, J. Phys. D, 12 (1979) Ll17. 39 A. G. Cullis and R. F. C. Farrow, unpublished, 1980. 40 U. Solzbach and H. J. Richter, Surf. Sci., 97 (1980) 191. 41 K. J. Riley, ht. Electron Devices Meet., December 8-10,1980, Washington, DC, Paper 19.1. 42 Y. J. Shacham-Diamand and I. Kidron, ZnfraredPhys., 21(1981) 105. 38
197
METELLURGICAL AND PROTECTIVE COATINGS
THE USE OF ION BEAMS IN MOLECULAR
BEAM EPITAXY
R. F. C. FARRO’N Royal Signals and Radar Establishment,
St Andrews Road, Malvern, Worcs. WR14 3PS (Gt. Britain)
Molecular beam epitaxy (MBE) has been developed over the past decade into one of the most versatile techniques for generating thin film semiconductor device structures. This development has been assisted by the use of low energy (0.1-10 keV) ion beams in a number of areas such as substrate surface preparation, film doping, metal-semiconductor contact formation and in situ surface analysis by secondary ion mass spectrometry and Auger depth profiling. Illustrative examples of the applications of ion beams in these areas are given. The use of ion beams in MBE has led to the discovery and exploration of a number of side effects of ion beams on semiconductor surfaces. These include group III metal surface segregation on III-V compound semiconductors, surface n-type channel formation on some III-V and II-VI compound semiconductors and II-VI alloy composition gradation. Some of these effects have been made use of in device structures. Mechanisms responsible for the effects are discussed and future trends in the use of ion beams in MBE are considered.
1. INTRODUCTION
Molecular beam epitaxy (MBE) has been developed I-3 over the past decade into one of the most versatile techniques for generating thin film semiconductor device structures. This development has been assisted by the use of low energy (0. l10 keV) ion beams in a number of areas such as substrate surface preparation, film doping, metal-semiconductor contact formation and in situ surface analysis by secondary ion mass spectrometry (SIMS) and Auger depth profiling. In Section 2, specific illustrative examples of the applications of ion beams in these areas are given. The wide range of powerful in situ surface analysis techniques, which is built into MBE systems, combined with ex situ characterization of MBE-grown films has led to the discovery and exploration of a number of side effects of ion beams on semiconductor surfaces. These include group III metal segregation on III-V -*Paper presented at the 3rd International Conference on Ion and Plasma Assisted Techniques, Amsterdam, The Netherlands, June 30-July 2, 1981. 0040-6090/8 1/OOOO-0000/$02.50
0 Elsevier Sequoia/Printed in The Netherlands
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compound semiconductor surfaces, surface n-type channel formation on some IIIV and II-VI compound semiconductors and II-VI alloy composition gradation. These effects are discussed in Section 3. Some of the effects have been used in device structures. Future trends in the use of ion beams in MBE are considered in Section 4. 2.
APPLICATION
OF ION BEAMS IN MOLECULAR
BEAM EPITAXY
2.1. Semiconductor surface preparation A prerequisite for the preparation of epitaxial films of semiconductors, metals or insulators of high crystallographic quality on semiconductor substrates is the removal of surface impurities prior to film deposition. Carbon and oxygen impurities are invariably present on the surfaces of semiconductors after sample loading and pumpdown to a ultrahigh vacuum ambient and are known to affect the morphology of homoepitaxial GaAs 3, InP ’ and silicon adversely. These impurities probably originate from a variety of sources ranging from wet chemical treatments including polishing and etching to adsorption of atmospheric gases such as CO,, CO and 0, prior to sample loading. Careful preparative treatments followed by heating in ultrahigh vacuum significantly reduce the impurity levels as judged by Auger electron spectroscopy (AES) and can often3 result in a coverage of impurity atoms of less than 1% of a monolayer. Congruent thermal etching of the semiconductor is necessary for this process to be effective, however, and whilst this is achievable for silicon and GaAs it is difficult2 to achieve sufficient congruent etching rates for the more unstable compound semiconductors such as InP and InSb. In these cases, low energy (500 eV) Ar+ ion bombardment has been found’ to be highly effective in reproducibly producing impurity-free surfaces as judged by AES. Alternating ion bombardment with annealing is necessary to permit reordering of the semiconductor surface disordered by the ion beam. A particularly striking recent example of the success of this technique is in the preparation of the first epitaxial metastable phase by MBE: c+Sn on InSb4. This example is particularly interesting since it demonstrates how critical the removal of surface contamination is to the nucleation of the metastable phase. Figure 1 shows a series of Auger spectra recorded before and after Ar+ bombardment and annealing treatments of an InSb(OO1)surface in an MBE system. After loading, the surface has considerable oxygen and carbon contamination shown by the presence of O(515 eV) and C(272 eV) peaks in spectrum a. After a 500 eV Ar+ dose of 2 x lOi cm- ’ and a 1 h anneal at 200 “C there are (see Fig. l(b)) significant reductions in the oxygen and carbon peaks and corresponding increases in intensity of the indium and antimony doublets resulting from removal of the impurity overlayer from the InSb surface. With further repeats of this cycle the Auger spectrum d, recorded at a higher resolution, shows that no oxygen and carbon (a coverage of less than 0.5% of a monolayer) remain. Reflection electron diffraction studies of this surface confirmed4 that it was ordered with a (2 x 4) or C(2 x 8) surface reconstruction. Condensation of a beam of tin atoms on this surface resulted in the growth of an epitaxial a-Sn film at substrate temperatures well above the cl-+B transformation temperature of about 13 “C. In contrast, the condensation of a beam of tin atoms onto a contaminated InSb surface exhibiting an Auger spectrum with a coverage of carbon and oxygen of greater than 1% of a monolayer resulted in the growth of a polycrystalline film of p-
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Sn. Clearly, intimate contact between the tin film and the isomorphous closely lattice-matched InSb substrate is necessary for metastable c+Sn growth through pseudomorphic growth. This effect has also been demonstrated for ol-Sn on ionbombarded annealed CdTe surfaces and has opened up a new area of epitaxy with potential device applications in IR detection by photovoltaic or photoconductive response in the a-Sn narrow gap semiconducting film.
&anneals
l&b
surface
I
E(ev)
500
I
I
400
300
]C
I 200
Fig. 1. Auger electron spectra of an InSbfOOl) surface: curve a, before in situ cleaning; curve b, after 500 eV Arf ion bombardment at 1.0 pA err-’ for 1 h followed by a 1 h anneal at T = 200 “C; curve c, after a further cycle of the above treatment; curve d, after one-third of a cycle of the treatment, the spectrum is recorded at a higher resolution than spectrum a-c.
The technique of ion bombardment and annealing has also been used as a precursor to homoepitaxial InSb growth by MBES and also to homoepitaxial InP ’ growth. However, the side effect of group III metal segregation which results from the ion bombardment of III-V compound semiconductors6’7 (see Section 3.1) is particularly severe for InP and results in indium island formation at ion doses of less than lO”j cm-’ unlike InSb and the other III-V compounds which require doses greater than 101’ cm-’ for island formation. Low energy ion bombardment of silicon surfaces followed by annealing at 900 “C or below is commonly used’ to prepare clean ordered surfaces prior to MBE growth of silicon or silicide films in ultrahigh vacuum. This process is preferred for some applications since the alternative cleaning technique of briefly flashing the
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R. F. C. FARROW
silicon to 1200 “C can cause significant diffusion of dopants within the wafer and diffusion of impurities to and from the wafer. Yamada et al.’ have made a careful study of residual damage to ion-bombarded wafers and concluded that Ne+ is the preferred species for ion bombardment cleaning of silicon since residual surface damage from 4 keV Nef can be annealed out after 15 min at 720 “C. 4 keV Ar+ bombardment results in more tenacious damage and can induce gas bubble formation whilst 4 keV Kr+ bombardment requires an anneal at about 900 “C for full recovery of surface order. Ar+ ion bombardment has been used by Bean” to prepare clean ordered sapphire and spine1 substrates for MBE growth of silicon. In this case the substrates are insulators and room temperature exposure to Ar+ ions leads to surface charging which reduces the sputtering rate and renders ion bombardment ineffective in removing impurities. At substrate temperatures of about 1000°C the substrate conductivity is sufficient to avoid surface charging and clean well-ordered surfaces (as judged by 1.8 MeV He+ channelling) are achieved. Silicon films grown on sapphire substrates prepared in this way reach the island coalescence stage earlier than on substrates cleaned simply by heating to 1400 “C. Surface dissociation of Al,O, is, in fact, significant at 1400°C and is likely to result in changes in near-surface stoichiometry with consequent degradation in surface smoothness and crystallinity. 2.2. Ionized beam doping The versatility of MBE as a preparation technique for device structures arises from the ability to control the level and type ofdopant with great accuracy in vertical architectures. For GaAs there are several suitable n-type dopants which are readily incorporated into the growing film when supplied as neutral atoms from Knudsen effusion sources. However, for p-type doping, beryllium is the only dbpant which approaches ideal requirements when supplied as a neutral atom. Beryllium, however, requires special handling procedures because of its high toxicity. Other dopants such as zinc and cadmium have very small sticking coefficients, as neutral atoms, at normal growth temperatures (500-600 “C) and are not incorporated into the GaAs film. In 1975 Naganuma and Takahashi” proposed a solution to this problem. They showed that when zinc was ionized and when it impinged at an energy of not more than 1.5 keV onto the GaAs surface during MBE growth at 55s 580 “C a significant incorporation of zinc acceptors in the film was achieved. The surface morphology remained smooth at zinc doping levels of up to 5 x 10” crnv3. Reasonable values of both the acceptor mobility and the photoluminescence intensity were observed, and Naganuma and Takahashi suggested simple ion burial as the mechanism for zinc incorporation. Bean and Dingle” later confirmed the effect and supported this mechanism for zinc incorporation but showed that asgrown epitaxial layers retained a substantial amount of residual radiation damage. Significant improvements in photoluminescence intensity and acceptor mobility values were demonstrated for post-growth anneals. The technique offers prospects for direct ionized beam writing of laterally defined dopant structures. However, the spreading of lateral dobing profiles by ion surface migration is an unknown factor at this time. For silicon growth by MBE, ionized beam doping has recently been used to overcome a number of problems inherent in doping with neutral species. For
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example, aluminium, antimony and gallium have a tendency to surface segregation during growth which tends to distort the shape of the intended doping profiles13. Sugiura l4 has shown that, at least for antimony, this segregation can be overcome by using low energy (130-1000 eV) ionized antimony as the dopant beam during growth at 850 “C. All arriving Sb+ ions are incorporated as donors in the film at this temperature, a doping range of 10’6-1020 cme3 has been demonstrated and Rutherford backscattering data confirmed the absence of profile distortion arising from surface segregation. Clearly the segregation effect is eliminated by shallow (about two silicon atomic layers) implantation of Sb+. Ota” has similarly shown that As+ ion doping (ion energy range, 400-800 eV) during silicon MBE can be used to produce arbitrary doping profiles over the range 10’4-10’g cme3. The quality of As+-doped epitaxial silicon films has proved adequate for majority carrier devices such as millimetre wave p-i-n and varactor diodes but to date no studies have been reported on minority carrier properties or on the distribution of electron traps in ionized-beam-doped films. 2.3. Metal-semiconductor conta.ctformution Metal film epitaxy from atomic beam sources has been developed by Farrow and coworkers16-r8 and others3 for the preparation of metal-semiconductor contacts in situ in an MBE system. For example, the growth of an epitaxial aluminium film at room temperature onto MBE-grown GaAs forms part of a process recently developed by Schneider and Cholg for fabricating arrays of millimetre wave diodes. In this process the epitaxial aluminium film is deposited over the entire GaAs wafer, and post-growth lithography and ion etching are used to define discrete diodes. In the future it will be desirable to eliminate the timeconsuming ex situ lithography and ion-etching process used to define structures such as these. The liquid metal field emission ion sources originated by Clampitt et aL2’ and developed by Pruett and Jefferies 21 offer considerable potential for direct micromachining of metal patterns such as those required for diode arrays. In addition, under some operating conditions the sources are known to emit charged microclusters of metal, suggesting the possibility of patterned metal deposition if the process can be sufficiently well controlled. Seliger et al.” have demonstrated the micromachining of 1000 A lines in gold films 400 A thick on silicon using a focused 57 keV Ga+ ion beam generated from a gallium field emission ion source. Farrow et al.* have studied the microstructure of aluminium and gold films deposited from an unrestricted nozzle type of field emission source in an MBE system. More recently the microstructure and contact characteristics of tin and silver films deposited onto semiconductors from restricted nozzle” sources have been studied in this laboratory 23. In both types of source, operation at low emission currents led to generation of both metal ions and charged microdroplets less than 1000 A in diameter. Auger spectra recorded of semiconductor surfaces exposed to beams generated from gold and tin field emission sources operated at low emission currents (less than 20 PA) confirmed the surface-cleaning effect of the ion component of the beam. Figure 2 shows Auger spectra recorded of an InP(OO1) surface before and during exposure to a gold source operated at 20 pA and 10.5 keV. Spectrum a shows considerable initial oxygen and carbon contamination of the InP surface. Exposure of the surface to the source for 30 min yielded spectrum b which shows considerable
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R. F. C. FARROW
reductions in the surface impurity peaks and increases in the native species In and P. A further exposure of the sample to the beam for 30 min (see spectrum c) led to further reductions in the impurity peaks and to further enhancement of the peaks due to indium and phosphorus. In addition, the progressive increase in the gold Auger peak is evident. These spectra and parallel scanning electron microscopy studies’* confirm that the semiconductor surface experiences simultaneous ion cleaning, microcluster deposition and probably some shallow Au+ implantation. Increasing the emission current led to increases in both the rate and the size of gold microdroplet condensation. Similar results have recently been observed for a restricted nozzle tin ion source. Metal-semiconductor contacts formed from continuous films of field-emission-deposited Au/InP and Ag/InP show reduced barrier heights for n-type InP but, for silver on p-type InP, near-ideal classical Schottky barrier behaviour has been observed. Since barrier height control and optimization is a key requirement for several semiconductor devices, this result is particularly significant and further exploration of the effects of deposition conditions on barrier height are in progress. Preliminary studies of the exposure of InSb surfaces to the beam from a restricted nozzle tin ion source confirm the existence of p-Sn nuclei which result from the condensation and freezing of microdroplets of liquid tin. Barrier height measurements are in progress.
C
bw
II
c
I__ so5
272
j_
a
500 -
400
Auger
300
electron
200
energy
100
eV
Fig. 2. Auger electron spectra of an InP(OO1) surface before and after exposure to a beam from a gold field emission source operated at an emission current of 20 uA and an extraction potential of 10.5 keV: curve a, spectrum recorded before exposure of surface to a gold field emission source; curve b, spectrum recorded after exposure of the surface to a gold field emission source for 30 min; curve c, spectrum recorded after exposure of the surface to a gold field emission source for 60 min. The progressive decreases in the surface impurity peaks and the increases in the indium, phosphorus and gold peaks should be noted.
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ION BEAMS IN MBE
At this time it may be concluded that if the single-ion and microdroplet components of field emission sources can be separated and controlled by the addition of an ion lens system then both modes of operation would have applications in microcircuit technology: micromachining or shallow implantation for the ion emission mode and selected area contact deposition for the microdroplet emission mode. This area of development is particularly active, at present, not .only in Great Britain but also in the U.S.A. and Japan. Studies of the mechanism of ion and ion cluster emission from the sources are continuing in Great Britain24925 and France26. 2.4. Ion beams for in situ analysis in molecular beam epitaxy The two areas of in situ diagnosis in MBE where ion beams have been used are in static SIMS and in Auger depth profiling. In SIMS a primary ion beam probe is used to eject secondary ions from the sample surface which are then mass analysed. In the static mode of SIMS the primary ion beam probe has a sufhciently low current density that the time for mass analysis of the secondary ions is much less than the sputtering time for one monolayer. In this mode SIMS is a surface-sensitive technique capable of providing information on surface composition and chemistry. Ploog and Fischer2’ used the static SIMS technique to probe in situ the surface of an MBE-grown tin-doped film of GaAs. The main reason for this work was to establish the growth temperature range over which tin segregated to the surface of the growing film. Positive secondary ion mass spectra were recorded at an erosion rate of about 50 A h- ’ using a primary Ar+ ion current of 10m7A and a beam diameter of 5 mm. Although the absence of an ion extraction and focusing lens between the sample and mass spectrometer greatly reduced the signal-to-noise ratio in the mass spectra, the data showed segregation of tin (for films grown at temperatures higher than 490 “C) in the outer 10-20 A of the film at a level three orders of magnitude greater than in the bulk. Parallel AES studies of the as-grown surface confirmed this surface segregation effect and by sequentially Ar+ ion sputtering the top 30 A of the film away it was established that the tin segregation was confined to within 20 A from the surface. The static SIMS technique has also been used by Dowsett et al.28 to evaluate residual surface impurity levels on InP surfaces prior to MBE. In this work the high sensitivity of a well-designed static secondary ion mass spectrometer to surface impurities was demonstrated. Surface impurities at levels of the order of 100 ppm were clearly evident on Ar+-ion-bombarded InP(OO1) surfaces which appeared clean as judged by AES (impurity detection limit, about 1% of a monolayer). Auger depth profiling has been used extensively3 in analysing the composition of multilayer structures grown by MBE (see Section 3.2 and Fig. 5). 3.
SIDE EFFECTS OF INERT GAS ION BEAMS ON SEMICONDUCTOR
SURFACES
3.1. III-V compounds In studying the microstructure of metal films deposited by MBE onto Ar+bombarded and annealed InP(OO1)surfaces Farrow and coworkers6,7,18 noticed the presence of small ( less than 1000 A in size) islands of indium metal at the deposited metal-InP interface. Careful experiments revealed that these were due to the effect of
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R. F. C. FARROW
Fig. 3. Scanning electron micrograph of an InP(OO1)surface bombarded with Ar+ ion3 at 300 K and not subjected to any annealing treatment after bombardment (ion energy, 500 eV; ion dose, 1.3 x 10” cme2). The small (about 150 A in diameter) particles covering the surface were confirmed (see text) as metallic indium. The 1000 A particle in this figure is a dust particle used to provide a focus.
Ar+ ion bombardment and not to the effects of annealing or to the electron beam. This effect was subsequently studied in greater detail using a variety of inert gas ions and other III-V semiconductors. Figure 3 is a scanning electron micrograph of an ion-bombarded InP surface not subjected to any annealing treatment. The surface is covered with particles about 150 A in diameter which transmission electron diffraction studies’ later confirmed were metallic indium. The 1000 A particle in this figure is a dust particle used to provide a focus. Lower ion doses produced smaller islands but at a similar density. Figure 4 is a scanning electron micrograph of an InP surface subjected to an Ar+ ion dose of 1.5 x 1016 cmW2and to a post-bombardment anneal at 250 “C for 1 h. The distribution and size of the islands are quite different in this case. It seems likely that at 250 “C the liquid indium islands are mobile on the clean surface of the semiconductor and coalesce to form the larger islands shown in Fig. 3. This idea is supported by measurements of the density and size of the islands on the assumption of hemispherical cap-shaped nuclei. Direct comparisons between the effect on InP and the other III-V compounds InSb, GaAs and GaP were made by simultaneously exposing the samples to the same ion beam. This confirmed that the effect was much more pronounced in InP than in the,other semiconductors. For InSb, for example, ion doses well in excess of 10” cm-’ were required before the effect was detectable. Similar doses of Kr+ and Xe+ ions for InP produced islands of similar size and density to those for Ar+ although Auger spectra recorded
ION BEAMS IN MBE
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Fig. 4. Scanning electron micrograph of an InP surface subjected to Arf ion bombardment (ion energy, 500 eV; ion dose, 1.5 x 1016cm-‘) and to a post-bombardment anneal at 250 “C for 1 h. The presence of larger indium islands than in Fig. 3 should be noted. The increase in island size is attributed to surface migration and coalescence of small islands of liquid indium (see text).
immediately after sputtering showed a large enhancement of surface phosphorus. Indeed for Arf ion bombardment and annealing the electron diffraction patterns indicated’ a well-ordered phosphorus-stabilized (2 x 4) surface reconstruction. The contribution to the diffraction pattern from the small surface coverage of hemispherical indiiun islands is in any case negligible. This suggests that differential sputtering of the group V element is not the dominant underlying mechanism for the effect since on momentum transfer grounds we would expect considerable general group V depletion especially for Arf bombardment of InP. Most probably, the mechanism responsible is simply the breaking of III-V bonds with subsequent surface migration and coalescence of the group III metal under the ion beam. Not only is the group III metal more thermodynamically stable than the III-V compound but in addition it is more probable that metal segregation occurs via a simple first-order surface reaction than that the III-V bond is reformed via the more complex III-[V dimer] association reactiot?. The excess of group V atoms resulting from the metal segregation reaction tends to drive the rest of the semiconductor surface towards a group-V-rich surface phase. The post-bombardment anneal temperature therefore determines the phase of the rest of the surface. For example, if anneal temperatures higher than the congruent evaporation temperature limit2 are employed, the surface not occupied by group III metal will be driven towards a group-III-rich surface phase, e.g. C(8 x 2) for InSb annealing at 400°C 5 (about
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R. F. C. FARROW
100°C above the congruent evaporation limit). Independent confirmation of the formation of free group III metal islands on III-V compounds under a variety of ion bombardment, heating and film growth conditions has been presented by Olivier2g who used the technique of low energy electron loss spectroscopy to detect plasmon losses in the islands. The presence of these islands, which may form localized ohmic contacts beneath an overdeposited metal film, is believed’ to be a contributory factor to the greatly reduced barrier heights observed for a range of metals deposited onto ionbombarded and annealed InP and GaAs. The islands, however, seem to have little effect on the structure and electrical properties of overdeposited InP 3o and GaAs 31 films. They may, however, play a role in generating the cone structures seen by Farrow and others on InP surfaces subjected to extensive Ar+ ion milling during Auger depth profiling. Such structures invariably show indium metallic precipitates near the cone tips. A second and equally important side effect of Ar+ ion bombardment of some III-V compounds is the formation of an n-type surface-conducting channel, probably resulting from electrically active point defects arising from ion-induced lattice damage. This effect has been reported for p-type InAs 32. Also Farrow3’ and Kreutz34 have observed a similar effect for p-type InSb and semi-insulating InP. For p-type InAs (net hole concentration, 2 x lOi cmm3; 77 K) an Ar + (140 eV) dose of only 6.25 x 1013 cm-’ (= 10 PC cmm2) was adequate32 to generate a surface channel of effective depth 0.28 pm with a volume density of electrons in the channel of about 2 x 1016 cm- 3. The mobility of electrons in the surface channel was comparable (2.3 x lo4 cm2 V-r s- ‘) with the bulk mobility value, indicating that few extrinsic scattering defects were produced by the low energy ion bombardment. Thermal annealing at temperatures above 70°C tended to remove the n-type channel and to restore p-type conductivity, presumably as the damage-induced point defects annealed out. The use of a higher energy (10 keV) Ar+ beam improved the thermal stability of the n-type channel but led to lower electron mobilities in the surface channel. Clearly, in this case, extrinsic scattering defects were created by the ion beam. The effect of intermediate ion energies and other inert gas ion species were not explored and it seems likely that an adequately stable surface channel with good electron mobility could be formed at ion energies lower than 10 keV. This would, in principle, permit the technique to be used for fabrication of an IR photodiode array with some on-chip signal processing. We observed a similar surface conversion effect for p-type InSb and it is significant that Fujisawa 35 has recently reported surface conversion of p-type InSb by proton or silicon ion bombardment or by Nd-yttrium aluminium garnet laser irradiation. Clearly, for InAs, InSb and InP, lattice damage induces a near-surface defect state or series of states with associated levels close to the conduction band minimum. The nature of these defects is at present unclear although it is tempting to associate them with the near-surface defects, responsible for Fermi level pinning, which result from metal deposition on the (110) cleavage faces of these semiconductors. This Fermi level pinning occurs at near midgap for GaAs 36 but near the conduction band minimum for InP 36 and InSb 34. This trend is consistent with the known effects of ion-induced lattice damage on surface conductivity. Proton bombardment of GaAs produces semi-insulating material (deep-level near-
207
ION BEAMS IN MBE
midgap damage-induced electron trap states) whereas proton bombardment of InSb produces n-type surface conductivity (shallow damage-induced donor-like states near the conduction band minimum). The identity of the lattice defects responsible for Fermi level pinning is still the subject of controversy. However, group III vacancies are a distinct possibility in view of the strong tendency of group III atoms to form localized metal clusters when displaced from their lattice sites. 3.2. II-VI compounds and alloys Recent studies3’ on MBE growth of CdTe films on CdTe, InSb and Hg, _,Cd,Te substrates have revealed several striking differences in the effects of inert gas ion beams on the surfaces of these semiconductors. For CdTe, scanning and transmission electron microscopy studies of Ar+-bombarded and milled surfaces show smooth featureless surfaces with a complete absence of segregation of either phase. In addition, AES studies of such surfaces show no evidence for preferential sputtering of either cadmium or tellurium. This absence of ion-beam-induced artefacts results in a planar front during Auger and ion microprobe depth profiling using Ar + ions. Figure 5 shows an Auger depth profile of a CdTe epitaxial film 7500 A thick grown by MBE on InSb(OO1) at 120°C. The abrupt interface (with a transition width of about 100 A) between CdTe and InSb is an indication both of the absence of surface topographical modifications during CdTe ion milling and of the absence of interdiffusion at the low growth temperature. ParallelSIMS profiles have confirmed3’ this abrupt interface. Whilst no compositional or topographical effects of CdTe ion bombardment were observed we found that ion bombardment of high resistivity CdTe substrates induces a surface n-type channel. Measurements of the channel depth and conductance are in progress and it is too early to speculate on the mechanism responsible for this effect. Auger signal
t
;te~,.lb*-.........._. ....-....-*........ I-
p”-Q
- . . . . . . . - . . . . - . . . . . . . . . . . . . . . . 1: :,;
:’
/H TRANSITION
ZONE
Fig. 5. Auger depth profile ofan MBE-grown CdTe film on an InSb(OO1)substrat@r+ ion beam at 1 keV and 40 pA cm-‘). The profile shows the unnormalized peak heights for cadmium(O), tellurium ( n), indium (0) and antimony (A) as functions of depth for the last 3500 A of a CdTe film 7500 A thick (transition zone, 100 A).
The effects of ion bombardment and ion beam milling of the narrow gap semiconducting alloy Hg, _,Cd,Te contrast strongly with their effects on CdTe in that gross compositional, topographical and structural changes occur for
208
R. F. C. FARROW
Hg,_,Cd,Te. Table I shows compositional data for Hg&d,,,Te wafers before and after exposure to an Ar+ ion beam (energy, 5 keV; 12pA cm-‘). The data were derived from an in situ study of surface composition using the technique of X-ray photoelectron spectroscopy (XPS) and refer to the average composition of an area about 5 mm x 5 mm of the top 15-20 A of the sample surface. After Ar+ ion milling, the surface composition is significantly depleted in mercury and enriched in cadmium (i.e. x for the surface is greater than x for the bulk). The effect appears to have a dominant underlying thermal mechanism, possibly transient surface heating, since the depletion is of the volatile end member (HgTe) of the alloy rather than of the two lighter elements (cadmium and tellurium) which might be expected on solely momentum transfer grounds. Congruent loss of HgTe from the ahoy is known3* to occur as a result of heating the alloy at temperatures in the range 300-4OO”C. Support for the thermal mechanism comes from the observation of greater HgTe depletion at higher power denSities of the ion beam with the ion beam energy kept constant. The very low thermal conductivity of Hg, _,Cd,Te also tends to favour surface heating despite good sample heat sinking and cooling to - 180 “C (see Table I). Transmission electron diffraction studies3’ of Ar+-milled foils of Hg, _,Cd,Te have revealed considerable structural changes in the surface layers resulting from ion milling. These include inert gas bubble or microvoid formation and phase separation of the alloy into microdomains of HgTe and CdTe. Electrical measurements of ion bombarded Hg,_$d,Te wafers indicate p- to n-type conversion at ion doses much less (less than 1016 cmT3) than those required to give significant changes in the x value of the alloy from its bulk value. This effect has also been observed by Solzbach and Richter4’ and opens up the prospect of diode formation by controlled ion bombardment as for InAs and InSb. TABLE I X-RAY PHOTOELECTRON SAMPLE~SUBJECTEDTO
SPECTROSCOPY DATA ON THE SURFACE IONMILLING(~ keV;l2pAcm-‘)
COMPOSITION
OF AN
Hg,,,Cd,,,Te
Ar+
Composition (at.%)
Before ion milling 40 min; 25 “C 15 min; - 180°C
Hg
Cd
Te
C
0
23.1 34.2 35.6
1.8 18.8 18.4
14.5 47 46
41.7 0 0
12.9 0 0
“The bulk composition of the Hg,,,Cd,,,Te
sample is 4O%Hg-lO%Cd-50%Te.
4. DISCUSSION The use of ion beams in &iBE has made a significant contribution to the development and application of the technique. This is illustrated by the particular examples given in Section 2. Several points are, however, worth highlighting. Metastable phase formation by epitaxy, initiated by the a-Sn work described in Section 2.1, is a field in which reproducible surface cleaning is critical to the growth of the metastable phase. This area will probably develop rapidly in the near future. For silicon MBE the current trend is towards submicron film epitaxy over silicon wafers 2 or 3 in in diameter with an emphasis on the lateral uniformity of film
ION BEAMS IN MBE
209
thickness and doping level. The current availability of inert gas ion beams of uniform intensity in this size regime will permit reproducible wafer cleaning prior to epitaxy. In addition the use of ionized beam doping in silicon MBE shows considerable promise in achieving arbitrary doping profiles. It may also provide a route to selected area doping. However, a considerable amount of ground work needs to be done first in assessing the effects of ionized beam doping on minority and majority carrier properties. Deep level transient spectroscopy and photoluminescence studies of the material should enable the distribution of carrier traps to be explored as a function of doping and annealing conditions, leading eventually to optimization of these conditions. The field emission sources discussed in Section 2.3 show considerable promise both as high brightness sources for micromachining and as metal deposition sources for in situ metal-semiconductor contact deposition. The ultrahigh vacuum compatibility of these sources, which produce a negligible gas load, combined with their compactness makes them highly attractive for applications in MBE. Here also, much work remains to be done in relating the contact electrical properties of fieldemission-deposited metals to the deposition conditions for a variety of semiconductors. Controllable separation of the single-ion and microdroplet components of the beam would certainly add to the versatility of the sources. The side effects of inert gas ion beams on semiconductors are varied and complex, as illustrated by the examples given in Section 3. Some of them appear to have useful applications in device structures. For example, the p- to n-type surface conversion effect has been used to make junction field effect transistor structures in InAs and may provide a means of making similar structures in InSb and CdTe. Low dose Ar+ ion bombardment of n-type Hg, _,Cd,Te followed by in situ ‘metal deposition under clean controlled conditions in an MBE system may provide a convenient way of making metal-n+-n contacts to the alloy. Such contacts are known to act as minority carrier (hole) accumulation contacts with beneficial effects on photoconductive and photovoltaic device performance41~42. The formation and stability of ion-bombardment-induced diodes in the alloy is also worthy of investigation in view of the considerable current interest41 in far IR photodiode arrays in Hg, _,Cd,Te. 5.
CONCLUSIONS
The use of ion beams in MBE has contributed significantly to the development and versatility of the technique. Much work, however, remains to be done in establishing and understanding the mechanisms of the electrical and structural effects which accompany exposure of semiconductor surfaces to low energy ion beams. An improved understanding of these effects should enable full advantage to be taken of the techniques of ionized beam doping, ion-beamqnduced surface-type conversion and field emission deposition of contacts in device fabrication by MBE. ACKNOWLEDGMENTS
The author is grateful to A. Christie, Loughborough Consultants, and to Hugh Bishop, U.K. Atomic Energy Authority, Harwell, for the XPS studies of the alloy
210
R. F. C. FARROW
I-Ig, _,Cd,Te. The work of D. Sykes, Loughborough Consultants, on Auger depth profiling of MBE-grown CdTe/InSb heterostructures is also acknowledged with thanks. REFERENCES 1 A. Y. Cho and J. R. Arthur, Prog. SolidState Chem., lO(1975) 157. 2 R. F. C. Farrow, in E. Kaldis and H. J. Scheel (eds.), Crystal Growth and Materials, Vol. 1, NorthHolland, Amsterdam, 1977, Chap 1.7. 3 K. Ploog, in H. C. Freyhardt (ed.), Crystals-Growth Properties and Applications, Vol. 3, Springer, Berlin, 1980, p. 73. 4 R. F. C. Farrow, D. S. Robertson, G. M. Williams, A. G. Cullis, G. R. Jones and I. M. Young, Proc. 8th Int. Vacuum Congr., Cannes, September 22-26,1980, in Vide, Suppl., I(201) (1980) 109; J. Cryst. Growth, 54 (1981) in the press. 5 K. Ge. S. Ando and K. Sugiyama, Jpn. J. Appl. Phys., 19 (1980) L417. 6 R. F. C. Farrow and A. G. Cullis, ht. Conf. on Solid Films and Surfaces, Tokyo, 1978, unpublished. 7 A. G. Cullis and R. F. C. Farrow, Thin Solid Films, 58 (1979) 197. 8 J. C. Bean and J. M. Poate, Appl. Phys. Left., 37 (1980) 643. 9 I. Yamada, D. Marton and F. W. Saris, Appl. Phys. Lett., 37(1980) 563. 10 J. C. Bean, Appl. Phys. Left., 36(1980) 741. 11 M. Naganuma and K. Takahashi, Appl. Phys. Left., 27 (1975) 342. 12 J. C. Bean and R. Dingle, Appl. Phys. Lett., 35 (1979) 925. 13 J. C. Bean, Appl. Phys. Left., 33 (1978) 654. 14 H. Sugiura, 11th Int. Conf. on Solid State Devices, Tokyo, 1979, Paper C-3-6, in Digest of Technical Papers, unpublished. 15 Y. Ota, 11th Int. Conf. on Solid State Devices, Tokyo, 1979, Paper C-3-5, in Digest of Technical Papers, unpublished., 16 R. F. C. Farrow, A. G. Cullis, A. J. Grant and J. E. Pattison, J. Cryst. Growth, 45 (1978) 292. 17 R. F. C. Farrow and G. M. Williams, Thin SolidFilms, 55 (1978) 303. 18 R. F. C. Farrow, A. G. Cullis, A. J. Grant, G. R. Jones and R. Clampitt, Thin SolidFilms, 58 (1979) 189. 19 M. V. Schneider and A. Y. Cho, Proc. Cornell University Workshop on Molecular Beam Epitaxy, October 21-22,1980, in J. Vat. Sci. Technol., (July-August, 1981) in the press. 20 R. Clampitt, K. L. Aitken and D. K. Jeffries, J. Vat. Sci. Technol., 12 (1975) 1208. 21 P. D. Pruett and D. J. Jefferies, Inst. Phys. Co@ Ser. 54 (1980) Chap. 7, p. 316. 22 R. L. Seliger, J. W. Ward, V. Wang and R. L. Kubena, Appl. Phys. Left., 34 (1978) 310. 23 R. F. C. Farrow, P. Tapstev and D. G. Coates, unpublished, 1981. 24 A. J. Dixon, J. Phys. D, 12(1979) L77. 25 A. R. Waugh, J. Phys. D, 13 (1980) L203. 26 P. Sudraud, C. Colliex and J. De Walle, J. Phys. Left., 40 (1979) L207. 27 K. Ploog and A. Fischer, J. Vat. Sci. Technol., 15 (1978) 255. 28 M. G. Dowsett, R. M. King and E. H. C. Parker, Appl. Phys. Left., 31(1977) 529. 29 J. Olivier, Proc. 8th Int. Vacuum Congr., September 22-26, 1980, Cannes, in Via’e, Suppl., I (21) (1980) 583.
P. W. Sullivan, R. F. C. Farrow, G. R. Jones and C. R. Stanley, Proc. 8th Int. Symp. on GaAs and September 22-24,1980, in Inst. Phys. Conf. Ser. 56 (1981)Chap.l, p.45. 31 A. Y. Cho, J. V. Di Lorenzo, B. S. Hewitt, W. C. Niehaus, W. 0. Schlosser and C. Radice, J. Appl.
30
Related Compounds,Vienna, Phys., 48 (1977) 346.
32 33
34 35 36 37
M. F. Millea, A. H. Silver and L. D. Flesner, Thin Solid Films, 56 (1979) 253. R. F. C. Farrow, unpublished, 1980. E. W. Kreutz, personal communication, 1980. I. Fujisawa, Jpn. J. Appl. Phys., 19 (1980) 2137. W. E. Spicer, P. W. Chye, P. R. Skeath, C. Y. Su and I. Lindau, J. Vat. Sci. Technol., I6 (1979) 1422. R. F. C. Farrow, G. M. Williams, G. R. Jones and I. M. Young, 1st European Workshop on Molecular Bean Epitaxy, Stuttgart, March 30-April I, 1981; submitted to Appl. Phys. Left.
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R. F. C. Farrow, G. R. Jones, G. M. Williams, P. W. Sullivan, W. J. 0. Boyle and J. T. M. Wotherspoon, J. Phys. D, 12 (1979) Ll17. 39 A. G. Cullis and R. F. C. Farrow, unpublished, 1980. 40 U. Solzbach and H. J. Richter, Surf. Sci., 97 (1980) 191. 41 K. J. Riley, ht. Electron Devices Meet., December 8-10,1980, Washington, DC, Paper 19.1. 42 Y. J. Shacham-Diamand and I. Kidron, ZnfraredPhys., 21(1981) 105. 38