Epitaxial growth of chromium on Cu{111} substrates

Thin Solid Films, 128 (1985)139-148 PREPARATION AND CHARACTERIZATION

139

EPITAXIAL G R O W T H O F C H R O M I U M ON C u { l l l } SUBSTRATES H. L. GAIGHER,N. G. VAN DER BERGAND J. B. MALHERBE Department of Physics, University of Pretoria, Pretoria0002 (South Africa)

(ReceivedAugust22, 1984;acceptedJanuary 28, 1985)

Deposits of chromium (1-40 nm thick) condensed in ultrahigh vacuum (about 10-9 Torr) onto Cu{ 111 } films at room temperature (about 300 K) were examined by transmission electron microscopy, electron diffraction and Auger electron spectroscopy (AES). Diffraction patterns consistent with Nishiyama-Wassermann (NW) and Kurdjumov-Sachs (KS) orientations were observed. Arcing of the chromium reflections indicated a continuous distribution of alignments within a + 6 ° spread about the exact NW orientation. Dark field and defocusing phase contrast techniques revealed small (about 28 nm) irregularly shaped chromium grains within which smaller parallel-oriented subcrystallites, elongated approximately in the Cr(001) direction, exist. AES confirmed the existence of discontinuous chromium deposits for thicknesses of about 6 nm or less and indicated the presence of a thin layer of Cr20 3.

1. INTRODUCTION The epitaxy of f.c.c, metals on {110} b.c.c, metal single crystals has been extensively observed 1-a. For the complementary system, i.e. the epitaxy of b.c.c. metals on { 111} f.c.c, metal surfaces, most of the work has been limited to deposits of iron on f.c.c, substrates 4-t3. Work on deposits other than iron is limited and includes, for example, the observations by Jesser and M atthews 14 and Cinti et al. ~ for chromium deposits on Ni{ 100} and Ag{ 111 } substrates respectively. Renewed interest in the epitaxy of f.c.c.-b.c.c, metal pairs has been stimulated by the recent theoretical considerations of van der Merwe ~6 on the epitaxy of { 111 } f.c.c, overlayers on {110} b.c.c, substrates. Braun and van der Merwe 17 have extended the theory to {110} b.c.c, films On {111 } f.c.c, substrates. These theoretical studies have been aimed at predicting ideal epitaxial configurations for b.c.c.-f.c.c. epitaxial systems. More experimental work on b.c.c.-f.c.c, systems is needed to evaluate the theory. In this paper results on the room temperature epitaxy of chromium deposits on { 111 } single-crystal copper substrates are reported.

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H . L . GAIGHER, N. G. VAN DER BERG, J. B. MALHERBE

EXPERIMENTAL DETAILS

Successive deposition of substrates and deposits was carried out in an ionsublimation-pumped degassed ultrahigh vacuum system. Substrate thicknesses were about 40 nm for thin (1.5-6 nm) deposits and 60-100 nm for thicker deposits. The {111} single-crystal copper substrates were prepared by electron beam evaporation of copper (purity, 99.99~o) onto {111} vapour-deposited sodium chloride films approximately 10 nm thick which in turn were formed on air-cleaved mica 1a. The temperature of the sodium chloride and mica during deposition of the copper was 350 °C and the residual gas pressure was 5 x 10- a Torr. After the system and substrates had been allowed to cool to room temperature, the chromium (purity, 99.98~o) was sublimed offa tungsten boat at a rate of about 0.03 nm s- ~ and at a residual pressure of about 2 x 10 -9 Torr. The thickness and evaporation rate were determined using a quartz crystal oscillator. A shutter positioned between the source and the substrate allowed the deposition of chromium to different thicknesses. For examination in the electron microscope (Philips EM 301), either Cr/Cu bilayers, obtained by simply dissolving the intermediate sodium chloride in water, or stripped chromium films, obtained by dissolving the copper in a solution of trichloroacetic acid and ammonia 19, were used. Very thin (about 4 nm or less) deposits were protected with a carbon film before the substrate was dissolved. Auger electron spectroscopy (AES), combined with ion etching, was used to determine the composition profile of the chromium films and the interface between the chromium and copper films. The AES was performed in a Varian system with a base pressure of about 10-9 Torr using a cylindrical mirror analyser. Argon gas (purity, 99.999~o) was used at a pressure of 5 x 10- s Torr for sputter depth profiling. The 0.5 keV rastered ion beam was approximately normal to the sample surface, while the incidence angle of the excitation electron beam was, unless otherwise stated, about 55 °. 3. RESULTS 3.1. Orientation Selected area electron diffraction patterns from stripped as well as double-layer films showed diffraction spots due to normal b.c.c, chromium for all deposit thicknesses greater than 1.5 nm. Figure 1 shows a typical transmission electron diffraction pattern from a Cr/Cu bilayer film. Evaluation of the deposit structure was facilitated by using stripped films to eliminate scattering from the relatively thick substrate. Figure 2 shows a typical diffraction pattern from such a stripped film. In previous thin film studies in which a very thin f.c.c, layer transformed to b.c.c. grains, it was emphasized that a unique determination of the b.c.c.-f.c.c, orientation relationship from diffraction patterns such as Figs. 1 and 2 is difficult 1°' 11. Similar difficulties will occur when the substrate is faceted. Diffraction patterns were calculated for a fixed [111] f.c.c, beam direction and all variants of the NishiyamaWassermann (NW), Kurdjumov-Sachs (KS), Pitsch and Bain b.c.c.-f.c.c, orienta-

EPITAXIALGROWTHOF Cr ON Cu{111}

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Fig. 1. Electron diffraction pattern of a Cr/Cu(l 11) bilayer film with a chromium thickness of about 8 nm. Spots (arcs) marked D are due to double diffraction. Fig. 2. Electron diffraction pattern from a stripped chromium film with an average thickness of about 17 nm.

tion relationships 1°. Deviations of approximately 5° from the exact Bragg condition were allowed. Figure 3 shows the pattern expected for the three NW and six KS orientations resulting from

(110) b.c.c. // (111) f.c.c. The three NW orientations correspond to [001] b.c.c, parallel to the three (110) f.c.c, in the (111) substrate plane, while the KS orientations occur when [T11] and [1T1] b.c.c, are parallel to the (110) substrate directions. Comparison of the observed patterns (Figs. 1 and 2) with Fig. 3 shows that the former contain reflections consistent with the above-mentioned NW and KS orientations. The b.c.c. •

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Fig. 3. Electron diffraction pattern expected from Cr(ll0)/Cu(lll) with [001]cr, [Tll]cr and [111]c~ parallel to ( 110)c u in the substrate plane: O, {220}-type copper reflections; o, chromium reflections from NW orientations, i.e. [001 ]cr//( 110)cu; o, chromium reflections from KS orientations, i.e. IT 11]cr// (110)cu and [ 1 T 1 ] c r / / ( l l O ) c .. Unit meshes for the NW ( - - - - - ) and KS ( - - - ) orientations are indicated.

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H . L . GAIGHER, N. G. VAN DER BERG, J. B. MALHERBE

reflections were not sharp spots but extended arcs (Figs. 1 and 2). The arcs indicated a continuous distribution of alignments within a __+6 ° spread about the exact NW orientation and hence include the KS orientations at + 5 ° 16' from the NW relation. The diffraction patterns observed are also consistent with crystals in the abovementioned N W orientations together with six of the Pitsch orientations. The six Pitsch orientations which may have been present are those for which the two ( I 12) b.c.c, directions in the (110) b.c.c, plane are aligned parallel to the (110) f.c.c. directions which are not parallel to the (111) substrate plane, but which are in the three {100} f.c.c, planes, equally inclined at about 54 ° to the substrate plane. Reflections from these six Pitsch orientations would be indistinguishable from the KS reflections in Fig. 3. Whereas the surface of the Cu{111} substrates, grown on NaCl{lll}/mica, are smooth over relatively large areas (about 100nm), the existence of { 100} facets is certainly plausible. Tetrahedral growth features with side faces parallel to {100} were in fact observed on surfaces ofNaCl{ 111 }/mica 2°. Some of these { 100} facets may be replicated in the copper film and serve as sites for b.c.c. nuclei in Pitsch orientations. However, it is not clear why growth of only six of the 12 possible Pitsch variants should occur. Furthermore, Pitsch orientations that grow on { 100} facets should be limited to relatively local (faceted) areas of the specimen. It is suggested that the b.c.c, reflections inferred as KS in Figs. 1 and 2 are indeed primarily due to grains in KS orientations although the possibility of Pitsch orientations cannot be completely ruled out, especially when nucleation of the b.c.c. phase from an extremely thin, initially f.c.c., layer is possible. Finally, it may be remarked that the patterns in Figs. 1-3 cannot be reconciled with b.c.c, grains in any of the N W and KS orientations on { 111 } f.c.c, planes inclined to the (111) substrate plane nor with any of the Bain orientations. No convincing evidence could be obtained for the existence of f.c.c. (pseudomorphic) chromium. Diffraction effects which might have been interpreted in terms of f.c.c, chromium were also observed for substrates without any deposit and can therefore not be directly connected with the chromium deposits. For example, faint spots which occur just inside some of the 002Cr reflections (Fig. 1) are due to copper oxide 21. Streaks through the 220cu spots and/or a distortion of the spots resulted from defects in the substrate. For deposits less than 1.5 nm in thickness b.c.c. chromium reflections were usually absent. However, this does not necessarily imply that pseudomorphism had occurred since these observations were influenced by oxidation as evidenced by reflections due to Cr203 (see Section 3.2). It may in addition be mentioned that the observations of Jesser and Matthews 14 for chromium on Ni{001} and the low energy electron diffraction results obtained by Thomas and Haas 22 for chromium on Ir{lll} and A u { l l l } did suggest the occurrence of thin (less than 2 nm) f.c.c, chromium layers. Further experiments are clearly needed to gain a better understanding of the early stages of the epitaxial growth of chromium on f.c.c, substrates. 3.2. Auger electron spectroscopy

A typical AES depth profile for a 12.5 nm chromium deposit is shown in Fig. 4, The mean sputter etching rate was about 0.08 nm m i n - 1. The oxygen concentration

EPITAXIALGROWTHOF Cr oN Cu{111 }

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is a maximum close to the free surface and decreases rapidly with depth to a constant value of about 2 at.~. From a comparison of the AES spectra with those obtained from clean chromium and Cr20 3 it was concluded that the chromium surface consisted of a thin layer of Cr203. The thickness of this surface oxide is estimated at about 1.5 nm. The oxide formed during the transfer time (about 15 min) between the evaporation and AES chambers. No significant impurity peaks were observed in the Cr/Cu interface. AES spectra were also obtained from samples for which the deposit thickness decreased stepwise in the range 6-0 nm. For these observations the sample surface was oriented normal to the excitation electron beam. In the 6 nm chromium layer spectrum, the high energy (914 eV) peak of copper started to appear. This peak increased with decreasing chromium thickness. The low energy copper peak at 59 eV first appeared in the spectrum of the chromium layer 4 nm thick and also increased with decreasing chromium thickness. The finite escape depth of the Auger electrons has to be taken into account when interpreting the above result 2a-2s. No copper peak will be present in an AES spectrum if the copper is covered with a homogeneous chromium layer with a thickness greater than three times the mean escape depth of the Auger electrons of that specific copper peak. Using the predictions of Seah and Dench 24 and of Penn 2s the escape depths corrected for the cylindrical mass analyser geometry are about 0.3 nm for the 59 eV electrons and about 1 nm for the 914 eV electrons. The presence of copper peaks in the AES spectra from bilayers with relatively thick chromium deposits therefore indicates a discontinuous, i.e. island-like, chromium structure for deposit thicknesses up to about 6 nm. The reason why this early appearance of the copper peaks is not visible in the AES depth profile (Fig. 4) may be due to two factors. Firstly, for the depth profiling experiments the incidence angle of the analysing beam was 55°. Since the chromium crystallites are relatively near to each other (about 1 nm or less; see Section 3.3) shadowing effects will prevent the early emergence of the substrate copper peaks. Secondly, the sample surface and near-surface layer may be significantly altered by the atomic mixing induced by sputtering 26.

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H. L. GAIGHER, N. G. VAN DER BERG, J. B. MALHERBE

3.3. Morphology Bright field observation of the films for the exactly in focus position provided very little information about the morphology of the chromium deposits. Using the defocusing phase contrast technique resulted in great improvement in the image contrast 4, 27-29. Slight defocusing of the objective lens clearly revealed a network of channel-like regions separating different chromium grains (Fig. 5). These channels appeared light in the underfocused condition and reversed contrast when the condition was changed from underfocus to overfocus. The results obtained for the bilayer films were in all respects confirmed when stripped chromium films were used (Fig. 5(b)). For stripped films, complicating diffraction and contrast effects due to relatively large substrate thicknesses are absent. The micrographs in Fig. 5 are typical for deposit thicknesses of 6 nm or less. Close examination of these micrographs revealed irregular patches of chromium, most of which consisted of smaller subcrystals with an elongated shape. More information was provided by dark field illumination. Figure 6 shows typical dark field electron micrographs obtained using one of the arced chromium reflections for imaging. Irregularly shaped grains of chromium, ranging in size from about 5 to 80 nm, and averaging about 28 nm, were observed. These grains consisted of two or more smaller elongated crystallites, roughly in parallel orientation, which were aligned with their long axes approximately parallel to Cr(001) (Fig. 6). The widths of these elongated subcrystals were on the average about 5.7 nm, Three equivalent orientations of the chromium on the copper (see Section 3.1) allowed three different sets of these chromium grains to be observed depending on which 110Cr arc was used to form the image. For average deposit thicknesses of about 6 nm or more the channels (Fig. 5) between the irregularly shaped chromium patches were almost closed (Fig. 7), indicating that the grains had merged. The substructure within the chromium grains is still clearly visible (Fig. 7). For the thickness range investigated (1-40 nm) the width of the elongated subcrystals did not appear to change with thickness.

Fig. 5. Defocusedbright fieldelectron micrographsof a chromiumdeposit 6 nm thick on Cu{111}: (a) bilayerfilm;(b) stripped chromiumfilm.

EPITAXIALGROWTH OF Cr ON Cu{111 }

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Fig. 6. Dark field electron micrographs formed by using {110}-typechromium reflectionsand showing, within the chromium grains, subcrystalselongated parallel to (001 )c,: (a) bilayer film with a chromium thickness of 8 nm; (b) a stripped chromium film 6 nm thick.

Fig. 7. Defocusedbright field electronmicrograph of a stripped chromium film about 10nmthick. Moir6 fringes consistent with the bulk lattice spacings of both copper and chromium were often observed. Misfit dislocations were never observed. N o r was there any evidence of strain contrast as would be expected if overgrowth islands were strained to match the substrate. However, the results for the thickness range (less than 1.5 nm) in which misfit dislocations and strain contrast effects might be expected to be operative should be regarded with caution; the AES results (Fig. 4) showed that the top layers (about 1.5 nm) of a chromium film immediately oxidized to C r 2 0 3 on exposure to air. The "110cr spot" which has often been used for dark field imaging is really a triplet (Fig. 3) because of the coexistence of N W and KS orientations. Arcing of the spots furthermore indicates a continuous distribution of alignments. D a r k field

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H. L. GAIGHER, N. G. VAN DER BERG, J. B. MALHERBE

images were formed by displacing the beam to the extremities of the objective aperture thus allowing either the complete triplet or either of the two KS spots to form the image. As the beam was displaced many patches ohanged contrast. In this way it was confirmed that many of the large patches actually consisted of an agglomerate of smaller grains which differed in orientation from each other by up to about 10°. The elongated subcrystals within a particular grain appeared to have similar orientations. These conclusions were also verified by using the 2~-dimensional stereoscopic technique3°: the different reciprocal lattice vectors associated with the slightly different orientations result in dark field image shifts during defocusing of the objective lens. A through-focus pair of dark field images allows a distinction to be made between regions with slightly different orientations. 4. DISCUSSION From geometrical considerations Bruce and Jaeger 2 predicted predominantly KS orientations for the Cr{ 110}-Cu{ 111 } interface. In his recent theoretical study of {111 } f.c.c, overlayers on {110} b.c.c, substrates van der Merwe 16 predicted the KS orientation to be the preferred alignment for Cu{111}/Cr{110}. van der Merwe's theory has been extended to {110} b.c.c, overlayers on {111 } f.c.c, substrates 17. For the particular case of Cr{ 110}/Cu{111 } the KS and NW orientations are about equally probable. The present results, in which both NW and KS orientations occur, are in agreement with the latter calculations. However, care should be taken not to generalize since the present results only apply to a substrate temperature of about 300K. Bruce and Jaeger z found the orientations exhibited by a particular f.c.c.{111 }/b.c.c.{ 110} system to be dependent on the temperature T~of the substrate during film growth. For Ag{lll}/W{ll0}, for example, silver was found to be always KS epitaxial at T~ >~ 500 K but a continuous distribution of alignments between NW and KS occurred at T~ ~ 300 K. They concluded that the equilibrium alignment is KS but an appreciable activation barrier has to be overcome in a change from NW to KS alignment. Similar considerations might apply to the present case. Observations at higher substrate temperatures are in progress. A satisfactory explanation for the formation of the elongated suberystals within the irregular chromium islands (Fig. 6) is difficult. Crystals with elongated shape are not an uncommon occurrence in epitaxial growth. Particular examples are the needle-like crystals observed in the epitaxial growth of iron on nickel 31, gold4, 8 or silver 13,32 and the presence of elongated crystals when chromium was condensed on NaCI aa. In these cases the elongated shape was explained in terms of preferential growth along directions of small misfit. It should also be remarked that these deposits were obtained at substrate temperatures where mass transfer due to surface diffusion could occur. The explanation given above is not applicable to the present observations for chromium on copper. The chromium crystals were elongated parallel to the direction [001]cr ff [T01]cu, say, which is not at all a direction of small misfit: in the [001]cr//l'I01]cu direction the misfit is about 139/owhile the misfit for the [il0]c r//[i2i]c u direction is - 8 ~ . Crystals might have grown into the observed elongated shape because of the presence of substrate steps 34 and/or the tendency to adopt a crystallographic form

EPITAXIALGROWTHOF Cr ON Cu{ 111}

147

with facets of {110} and {100} type 35'36. However, difficulties arise when explanations based on the above-mentioned ideas are attempted. (i) There is no evidence that the substrate surfaces contained steps or other defects with a scale, orientation and periodicity that could explain the observed substructure. (ii) At the relatively high evaporation rate and low substrate temperature applicable to this study, surface diffusion is not expected to be significant. Surface diffusion is essential for both preferential growth at substrate steps and the development of particular growth habits. A highly speculative explanation for the observed microstructure is based on the above-mentioned misfits ( - 8 % parallel to [T10]cf and +13% parallel to [001 ]cr) for the Cr(110)/Cu(111) system. A coherent or partially coherent chromium island will be in tension, i.e. stretched, parallel to [T 10]c ~(parallel to [T2T]cu) and in compression parallel to [001]Cr. Relief of the tensile stress can occur by the formation of mechanical cracks roughly parallel to [001]c~, i.e. the direction normal to the tensile axis. The cracks will accordingly divide the chromium islands into separate (001)-elongated grains. These cracks, associated with the process in which coherency with the substrate is lost, will evidently appear at a very early stage of growth. Because of the limited adatom mobility and possibly shadowing effects37, filling up of the cracks occurs only very slowly with further deposition. For thicker films, when partial filling up of the channels has occurred, stresses, either intrinsic or induced when the films are separated from the substrate, may once again favour the formation of mechanical cracks along the original cracks. Owing to the relatively high evaporation rate, low substrate temperature and substrate surface defects, the distance over which atoms can rearrange when the stretched chromium relaxes to normal b.c.c. ([i 10]cr//[T2I]cu) will probably he less than one atomic spacing. The distance between the cracks will then equal the vernier period of misfit. It is interesting to note that the observed average distance between cracks (about 5.7 nm) in fact agrees with the vernier period of misfit (5.2 nm) in the direction [i 10]c~//[i21]cu. 5. CONCLUSIONS

(1) Chromium films (1-40 nm thick) vacuum deposited (at about 2 x 10 -9 Torr) onto Cu(lll) substrates at room temperature exhibit both NW and KS orientations with a + 6 ° spread of alignments about the exact NW orientation. (2) The chromium films consist of small (approximately 28 nm) flat irregularly shaped grains. The grains merge at a very early (thickness about 6 nm or less) stage of growth. (3) Within these grains there is a substructure of small crystals (width about 5.7 nm) elongated parallel to (001)cf and separated from each other by narrow channels. (4) These channels are interpreted as mechanical cracks, possibly caused by tensile stresses which result from the misfit between chromium and copper. (5) Only the normal b.c.c, structure was observed for all deposit thicknesses exceeding 1.5 nm.

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H . L . GAIGHER, N. G. VAN DER BERG, J. B. MALHERBE

(6) A E S c o n f i r m e d t h e e x i s t e n c e o f a d i s c o n t i n u o u s c h r o m i u m s t r u c t u r e f o r t h i c k n e s s e s o f 6 n m o r less a n d s u g g e s t e d t h e p r e s e n c e o f a t h i n ( a p p r o x i m a t e l y 1.5 n m ) a i r - f o r m e d s u r f a c e l a y e r o f C r 2 0 3. T h e T E M o b s e r v a t i o n s o f v e r y t h i n d e p o s i t s (1.5 n m o r less) m o s t p r o b a b l y m a n i f e s t t h e i n f l u e n c e o f t h i s o x i d a t i o n . REFERENCES 1 E. Grfinbaum, in J. W. Matthews (ed.), Epitaxial Growth, Part B, Academic Press, New York, 1975, p. 611. 2 L.A. Bruce and H. Jaeger, Philos. Mag. A, 38 (1978) 223. 3 E. Bauer, Appl. Surf. Sci., 11-12 (1982) 479. 4 M. Cahoreau and M. Gillet, Surf. Sci., 26 (1971 ) 415. 5 M. Cahoreau and M. Gillet, C.R. Acad. Sci., S~r. B, 268 (1969) 1024. 6 P. Gueguen, C. Camoin and M. Gillet, Thin Solid Films, 26 (1975) 107. 7 J . A . A . Engelbrecht, J. S- Vermaak and H. C. Snyman, Thin Solid Films, 48 (1978) 127. 8 J . A . A . Engelbrecht and H. C. Snyman, Thin Solid Films, 69 (1980) 301. 9 U. Gradmann, W. Kiimmerle and P. Tillmanns, Thin Solid Films, 34 (1976) 249. 10 G.H. Olsen and W. A. Jesser, Acta Metall., 19 (1971) 1009. 11 W.A. Jesser and J. W. Matthews, Philos. Mag., 15 (1967) 1097. 12 E.F. Wassermann and H. P. Jablonski, Surf. Sci., 22 (1970) 69. 13 G.H. Olsen and H. C. Snyman, Acta Metall., 21 (1973) 769. 14 W.A. Jesser and J. W. Matthews, Philos. Mag., 17 (1968) 475. 15 R. Cinti, J. Devenyi, P. Escudier, R. Montmory and A. Yelon, C.R. Acad. Sci., 260 (1965) 6849. 16 J.H. van der Merwe, Philos. Mag. A, 45 (1982) 127, 145, 159. 17 M. Braun and J. H. van der Merwe, AppL Surf. Sci., 22-23 (1985) 545. 18 F.A. Koch, C.T. H o r n g a n d R . W. Vook, J. Vac. Sci. TechnoL, 9 (1972) 511. 19 D.L. Carr, Rev. Sci. Instrum., 40 (1969) 1965. 20 F. A. Koch and R. W. Vook, Thin Solid Films, 14 (1972) 231. 21 E. G. Clarke, Jr., and A. W. Czanderna, Thin Solid Films, 12 (1972) 443. 22 R.E. Thomas and C. A. Haas, J. AppL Phys., 43 (1972) 4900. 23 S. Hofmann, in G. Svehla (ed.), Comprehensive Analytical Chemistry, Elsevier, Amsterdam, 1979, p. 89. 24 M.P. Seah and W. A. Dench, Surf. Interface Anal., 1 (1979) 2. 25 D.R. Penn, J. Electron Spectrosc., 9 (1976) 29. 26 P. Williams and J. E. Baker, Nucl. lnstrum. Methods, 182-183 (1981) 15. 27 R.H. Wade and J. Silcox, Phys. Status Solidi, 19 (1967) 57, 63. 28 S. Nakahara, J. Electrochem. Soc. C, 129 (1982) 20t. 29 R.B. Pettit and J. Silcox, J. Appl. Phys., 45 (I 974) 2858. 30 W.L. Bell, J. AppL Phys., 47(t976) 1676. 31 J.W. Matthews and W. A. Jesser, Philos. Mag., 20 (1969) 999. 32 H.C. SnymanandG. H. Olsen, J. Appl. Phys.,44(1973)889. 33 Y. Fukano and C. M. Wayman, J. Cryst. Growth, 15 (1972) 32. 34 G . D . T . Spiller, P. Akhter and J. A. Venables, Surf. Sci., 131 (1983) 517. 35 H. Poppa, E.H. LeeandR. D.Moorhead, J. Vac. Sci. Technol.,15 (1978) llO0. 36 E. Bauer, in M. H. Francombe and H. Sato (eds.), Single-crystalFilms, Pergamon, New York, 1964, p. 43. 37 A . G . DirksandH. J. Leamy, ThinSolidFUms, 47(1977)219.