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Texture development of evaporated nickel films on molybdenum substrates
ELSEVIER
Thin Solid Fitms 304 (1997) 170-i77
Texture development of evaporated nickel films on molybdenum substrates Z. Shi a,* G.R.G. Craib a, M.A. Player a, C.C. Tang a Department of Engineering, Unicersity of Aberdeen, Aberdeen AB24 3UE, UK b Daresbury Laboratory, Warrington WA4 4AD, UK
Received 14 October 1996; accepted i7 February 1997
Abstract UHY electron beam evaporation of nickel thin films on molybdenum substrates was carried out at different temperatures between Zone I and low Zone III of the structural zone model. The film texture was studied by energy-dispersive X-ray diffraction and the residual stresses were measured by monochromatic X-ray diffraction, both using the Daresbury synchrotron radiation source. It was found that the nickel film deposited at ambient temperature (Zone I) had a mixed fibre texture with a strong (11 t) and a weak (002) component. From Zone I-II to Zone II-III, the orientation of the films was dominated by the substrate texture, and granular epitaxy occurred. The films had an imperfect (001)[100] sheet texture, with the (002) and (I 1 l) pole figures, respectively, mirroring the (002) and (01 t) figures of the substrate, which had an imperfect (001)[110] sheet texture. An atomic arrangement model is put forward to explain the pole figure mirroring. The residual stress was found to be tensile and is mainly due to the thermal strain. It is suggested that the minimisation of the free surface energy dominates the texture development of the nickel film at low temperatures, while the anisotropic interfacial energy plays the determining role at elevated temperatures. © 1997 Elsevier Science S.A. Keywords: Epitaxy; Nickel; Physical vapour deposition (PVD); X-ray diffraction
1. Introduction Microstructures of thin films play an important role in determining their properties and thus their applications. Many previous studies have led to the development of the structural zone model (SZM) for the broad description of polycrystalline film structures. The SZM was first proposed by Movchan and Demchishin [1] and has three zones, identified by their characteristic structures, separated by two boundary temperatures T 1 and T:. For metals, T I = 0 . 3 T m and T2 = 0.45Tm, where Tm is the melting temperature. The model was later extended by Thornton [2] to sputter deposited films by adding a third coordinate to account for the influence of the working gas pressure. An additional transitional zone (Zone T) was introduced, situated between Zone I and Zone II. Although the SZM gives a fair description of many film microstructures, it does not provide information on the crystallographic orientation of the films. Vacuum-de-
Corresponding author. Tel: +44 i224 272526; fax: +44 1224 272497. 0040-6090/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S0040-6090(97)001 10-7
posited or sputter-deposited metal films often have some preferred orientation, or texture, which may lead to desired or undesired anisotropic properties and may affect their potential applications. The formation of texture in thin films is a complex phenomenon. Although some work has been carried out, there is still lack of comprehensive understanding as to how and when texture might form in films. Many factors can influence the development of texture, including both the deposition conditions and substrate effects. The substrate temperature, which determines the film microstructures in the SZM, also has the greatest influence on the orientation of films. It is now generally accepted that a texture with the most densely populated plane parallel to the surface is favoured at low substrate temperatures [3]. During the study of ultrahigh-vacuum (UHV) evaporated erbium (hcp) films on molybdenum substrates, it has been found that the substrate temperature strongly influences the texture of the erbium films [4,5]. The erbium films show a mixed (002) and (i01) fibre texture at ambient temperature (Zone I), which changes to a more mixed population at the transition region of Zone I - I I . At
Z. Shi et al. 7Thin-Solid Films 304 ~I997) 1 7 0 - 1 7 7
Table 1 Details of samples Sample No.
Coating
Substrate
~ (°C)
Ts / Trn
SZM zone
Thickness (nm)
I 2 3 3A 4
Ni Ni Ni Ni Ni
Mo Mo Mo Mo Mo
27 200 390 390 510
0.17 0.27 0.38 0.38 0.45
I I-II II II tI-III
700 700 700 700 700
face roughness of Rq = 80 nm. The thickness was measured after deposition using a Rank Taylor Hobson Talystep profilometer. The base pressure was below 10 .8 hPa before coating was started, and the residual gas was composed mainly of HaO, CO/N2, CO: and H : , as shown by mass spectroscopic analysis. During deposition, the growth rate was kept at about 0.67 n m s -1 and the incident angle was approximately 8 ° to the substrate normal. A set of substrate temperatures (T~) was used, corresponding to the region from Zone I to low Zone III in the SZM, i.e. T~ = 27, 200, 390 and 510 °C. Details of the samples are shown in Table 1. Two samples deposited at 390 °C under the same conditions are included to demonstrate reproducibility. Both the texture and stress measurements were carried out at the Daresbury Laboratory synchrotron radiation source (SRS). A newly developed X-ray diffractometer [6] was employed to study the texture of the nickel thin films and the molybdenum substrates, using energy dispersive techniques on Station 9.7. The geometry of the goniometer has been described in Ref. [4] and has the azimuth stage (rh) above the tilt stage (c~), with the roll stage ( ~ ) below. This provides a constant tilt and hence a constant interaction volume and footprint, and is therefore particularly suitable for thin film or surface layer studies at or neargrazing incidence. In the present study, multi-bunch operation was used, with steps of 7.5 ° in roll (~), 45 ° in azimuth (&) and 120 s counting time per frame. Diffraction patterns were recorded at a glancing angle ce = 2 ° and diffraction angle 2 0 - - 10°. The roll range used was 0-60 ° and the azimuth range was 0-315 ° , resulting in 72 data points for one sample. The pole figures were then prepared
mid-Zone II temperatures, the films have a very strong (002) fibre orientation, which then changes to a mixed (002) and (100) texture at even higher temperatures (Zone n-III). In the present study, fcc nickel was selected for comparison with erbium, in view of its different crystal structure but similar melting temperature. Nickel thin films were deposited on molybdenum substrates using electron beam evaporation under UHV conditions. Different substrate temperatures between Zone I and low Zone III were selected in order to study the effects of deposition temperature. The film texture was studied by energy-dispersive X-ray diffraction and the residual stresses were measured by monochromatic X-ray diffraction, both using synchrotron radiation.
2. Experimental Nickel thin films approximately 700 nm thick were produced using a Balzers UMS 500 UHV electron beam evaporation plant, on molybdenum substrates with a sur-
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Z Shi et al./Thin Solid Fihns 304 (1997) 170-177
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using an automated peak finder program. A correction for the loss of synchrotron beam current was applied, as was intensity correction using the procedure described in Ref. [4]. The figures were interpolated and plotted using UNIMAP, and plotted out to tp = 60 °. The residual stresses of the films were measured using the high-resolution monochromatic diffractometer on Station 2.3 at the SRS [7], with nominal wavelength set to 0.15 nm. Although not at grazing incidence, the parallel
beam optics used has many advantages compared with a conventional focusing Bragg-Brentano diffractometer [8]. The measurements were carried out using the sin2~ method with a texture goniometer. Use of a texture goniometer greatly increases the quality of tile results compared with a standard two-circle arrangement [9]. Ni (111) diffraction was selected for its higher intensity. A single-bunch beam was used with a 2 0 step of 0.005 ° and a counting time of 5 s. The peak positions (20) were then obtained using the
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7.. Shi et al. / Thin Solid Films 304 (1997) 1 7 0 - 1 7 7
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Fig. 4. Pole figures for sample 3A Ni film and Mo substrate, Ts = 390 °C. (a) Ni (111), (b) Ni (002), (c) Mo (011), (d) Mo (002).
orientation and the fibre texture is peaked at a direction approximately aligned to the deposition beam direction. A comparison of Fig. l(a) and (b) also reveals that the (111) texture is stronger than (002). The pole figures of the underlying molybdenum (Fig. lc,d) show an imperfect (001)[110] sheet texture. There is no obvious relationship between the orientation of the substrate and the film texture at ambient temperature. For the depositions at different elevated temperatures (samples 2, 3, 3A and 4), the pole figures of the nickel films and the underlying molybdenum substrate are shown
Winfit program [10] and the stresses were calculated from the slope of the 20 versus sin2~ plot.
3. Results
Fig. 1 shows the pole figures for the nickel film deposited at ambient temperature and the underlying molybdenum substrate. From the pole figures, it can be seen that the nickel film in Zone I has a mixed fibre texture. Both the Ni (111) and Ni (002) pole figures show a central fibre
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174
Z Shi et al./Thin Solid Films 304 (1997) 170-177
in Figs. 2-5. It can be seen that the pole figures of the molybdenum substrates are essentially as in Fig. 1, an imperfect (001)[l 10] sheet texture. The directions of the substrates were placed randomly and the similar direction of the substrates in Figs. 2 - 4 is merely coincidental. The textures of the films from low Zone II to low Zone III, however, are different from those of the Zone I film in Fig. 1. The orientation of the Mo substrate profoundly influenced the texture of these Ni films. Instead of a mixed fibre texture as at ambient temperature, the nickel films developed a (001)[100] sheet texture when deposited at elevated temperatures. Thus, the pole figures of the films and the substrate show a mirroring at these deposition temperatures, i.e. Ni (002) mirrors Mo (002), and Ni (111) mirrors Mo (011). It must be pointed out that the mirroring of Ni (111) and Mo (011) is only in the appearance of the pole figures and not in the exact pole positions. The Ni (1 t 1) poles are slightly wider apart from each other than the Mo (011) poles, corresponding to the inter-planar angle of 70.5 ° for Ni (111) and (]11) or Ni (1]1) and (~]1), and 60 ° for Mo (011) and (~01) or Mo (101) and (0]1). From these figures, it can also be seen that the change in the substrate temperature across the whole Zone II region does not have a significant effect on the texture of the films. Comparing Figs. 3 and 4 also shows that the two films deposited at 390 °C under the same conditions have the same preferred orientation, demonstrating the consistency of the results. Fig. 6 shows the variation in the (111) peak position against the value of sina~ for the nicket film deposited at 390 °C. The residual stresses were calculated from the slope of the 20 versus sin20 plot, with the following equation:
E o',
0(20)
2(1 -}-/-')
18--"-0rCOt0"
0(sin20)
(1)
where 0 is in degrees, °'4, is the stress parallel to the surface at an angle 4~ from principal stress axis 1 (see also Fig. 7), 0± the diffraction angle for the planes parallel to the surface, E is Young's modulus and v is Poisson's ratio.
g
48,20
~ 481s E
48.12
43.08 0
0.2
0.4
0.8
0.8
sin2¥
Fig. 6. Variationof (11 l) peak position with sin2~bof Ni film deposited at 390 °C (sample 3).
0" 2 = 4 7 5 MPa
~ 4 0 = 5 6 0 MPa
\
i°2 ~ k~ ,C~60=,480 MPa
Fig. 7. Schematic drawing of the residual stress in the Ni film deposited at 390 °C (sample 3). The orientation of the figure is the same as that of Fig. 3(a) and (b). Using the mechanical properties E = t99.5 GPa and v = 0.312 for nickel [11], a tensile stress of 610 + 40 MPa has been obtained from the slope in Fig. 6 for the particular 4, direction whose relationship to the substrate and film orientation is shown in Fig. 7. Measurements for two other directions at + 6 0 ° and - 6 0 ° to the above ~b direction were made with the two-point technique using ~0= 0 and 60 °. The stress values at these two directions were 560 _+ 40 and 480 _+ 40 MPa, respectively. From these results, the principal stresses o% and o-2 were calculated according to [12]. The values were found to be o-1 = 625 + 4 0 MPa, o-2 = 475 ± 40 MPa and ~ = - 19 ± 10°. The results are shown schematically in Fig. 7, where the orientation is the same as the corresponding pole figures (Fig. 3a,b). It is obvious that the residual stresses are tensile and show only weak anisotropy.
4. D i s c u s s i o n
The formation of texture is driven by the energy minimisation of the system. This often includes anisotropic surface energy, interracial energy and elastic strain energy in the film system. However, kinetics may decide what is actually occurring in a system, which may not be the most stable structure from a thermodynamics point of view. Therefore, the deposition conditions and the properties of the coating and substrate materials will all affect the final orientation and microstructure of films. Nucleation provides the first opportunity for texture formation since the orientation of vapour deposits is affected, and sometimes determined, by the structure of critical nuclei. Coalescence presents another opportunity for texture evolution as it occurs by an extensive rearrangement of materials of the islands that has been described as "liquid-like" or as "sintering" [13]. It is therefore expected that the nucleation and coalescence would play an important role in determining the film structure and orientation, especially in very thin films. As the film thickness increases, the role of grain growth will
Zo S h i e t aI. / Thin Solid Films 304 (1997) 1 7 0 - 1 7 7
become more and more pronounced as diffusion processes proceed. At low substrate temperatures, it is believed that the texture even in thick film is determined in the early stages of film growth and not by a selective growth mechanism during deposition [14]. In this case, the mobility of the vapour atoms on the substrate surface is low, and nucleation plays an important role in determining the film structure. Walton [3] proposed that the orientation of vapour deposits is determined by the structure of critical nuclei. It was predicted, as has also been found in many experiments, that high vapour supersaturation (low substrate temperature a n d / o r high incidence rate) favours the texture with the most densely populated plane parallel to the surface, but minimises the tendency for epitaxy on singlecrystal substrates, potentially leading to randomly oriented nuclei in the absence of a strong substrate effect. For fcc metals, the most likely texture is therefore (111) planes parallel to the substrate surface, followed by (002) and then (01 t). This is consistent with the fact that the free surface energy is lower with denser planes. It has been reported that nickel films evaporated at ambient temperature on amorphous carbon substrates have a slight preferred orientation with (111) planes parallel to the film surface even at a thickness of 2 - 4 mn [15]. '?his demonstrates the importance of nucleation and coalescence in texture formation. It is thought that the surface energy minimisation, possibly coupled with interfacial energy minimisation, plays the determining role in this system. The present results have shown that the nickel films have a stronger (111) and weaker (002) fibre texture at ambient temperature. Similar results have been reported for magnetron-sputtered nickel films on silicon wafers, which show a dominant (111) texture with a weak (002) component [16]. These results are consistent with Walton's model and indicate that free surface energy minimisation is the dominant driving force at ambient temperature for the texture development of nickel films. There is a possibility
175
that some (002) texture originates from epitaxial growth from the textured molybdenum substrate, which becomes dominant under the condition of higher substrate temperature discussed later. Such a case has been reported in as-deposited gold films on vacuum-cleaved mica, which are composed of grains with (111) texture, with a subpopulation of grains with epitaxial orientations [17]. This shows that inteffacial energy also plays an important role in the above situation. Nevertheless, whatever the reason for the (002) texture in nickel films, the present results show the dominant effect of the free surface energy on the overall texture at low temperatures. As said earlier, strain energy is one of the driving forces for the texture evolution. However, this is unlikely to be fully effective during film deposition. Instead, it will play a more prominent part during heat treatment of films where thermal stress can drive preferential grain growth through bulk migration of grain boundaries. Indeed, it has been reported that this had led to the development of (002) texture after annealing in copper, gold and silver films (fcc) on glass substrate and also in silver films on CaF 2 and NaC1 substrates [18-20]. The as-deposited texture at 80 K, however, was (111) and was also attributed to the surface energy minimisation. The texture of the nickel films deposited at elevated temperature are markedly different from that at ambient temperature. They all show an imperfect (001)[100] sheet texture (Figs. 2-5). The mirroring pole figures of the nickel films and the molybdenum substrates suggest that there should be an alignment of the nickel atoms to the molybdenum atoms. A description of this relationship depends on the structure of the interface, and in particular whether or not an oxide film is present. Ignoring for the present the question of an oxide film, it can be proposed that nickel atoms sit in the position corresponding to the valley of four (001) molybdenum atoms, as shown in Fig. 8(a). In this configuration, the positions of nickel atoms are an extension of the molybdenum lattice, and the nickel
Mo[010]
Mo[010]
Ni[010]
i
, ......
• Moil001
(a)
Ni[100] • Mo[I00]
Co)
Fig. 8. Arrangements of Ni atoms (small circle) over Mo (001) atoms (large circle). (a) Epitaxial configuration with four Ni (001) cells; (b) altemative configuration with four Ni (001) ceils.
176
Z Shiet al. / Thin Solid Films 304 (1997) 170-177
films grow epitaxially on the molybdenum substrate. This would lead to an orientation relationship of (002)Ni ]](002)Mo and [100]Ni H[110]Mo and would result in pole figure alignments as shown in Figs. 2-5. This configuration has a quite large lattice mismatch of 21%, using the data aN~ = 0.35238 nm and aMo = 0.31472 nm [21]. But this is not an unprecedented result; even higher misfits of 27% have been observed in gold, silver or copper films grown on sodium chloride cleavage surfaces [3]. However, taking account of lattice mismatch value alone, i.e. considering elastic strain energy as the only driving force, another atomic arrangement would be more favoured. This is shown in Fig. 8(b), with nickel atoms sitting in the position between two nearest (001) molybdenum atoms. It has a lattice mismatch of - t2%, substantially less than that shown in Fig. 8(a). Nevertheless, there are two obvious reasons to suggest that the configuration in Fig. 8(b) should not prevail. First, in the Zone II temperature range in which adatoms have sufficient mobility to adjust their positions, an individual nickel atom is more likely to sit in the valley of four molybdenum atoms than on the ridge between two molybdenum atoms. In other words, the interfacial energy with the configuration in Fig. 8(a) is lower than that with Fig. 8(b), because the valley is an interatomic potential minimum. Second, nickel atoms need to be compressed to fit into the positions in this configuration, which is less likely during early stages of vapour deposition process as a result of the interatomic potential between nickel atoms, especially when it requires the nickel atoms to stay on the ridge. In this case, therefore, variation of interfacial energy with orientation would be expected to dominate. The above model suggests that the alignment of the pole figures is due to the variation in the interfacial free energy with different orientations, rather than lattice matching alone. In fact, for polycrystalline films on single crystal substrates, it has been pointed out that lattice matching between the film and the substrate is not required for epitaxial grain growth, although it certainly promotes it [17]. All that is required is that there be a single orientation that minimises interfacial energy. From the above, it is proposed that the nickel films grow epitaxially on individual grains of the molybdenum substrate at elevated temperatures, due to the interracial energy minimisation. As the film thickens, these epitaxially grown grains continue to grow with the arriving deposit atoms taking up vacant lattice sites on the grain surfaces. This has been termed "granular epitaxy" [14]. In fact, the above atomic arrangement model is probably a simplification of the actual interface. In the present study, it is believed that any adsorbed contaminants are removed by heating under UHV before deposition. However, a thin, stable oxide layer is expected to be present on the molybdenum surface. The detailed characteristics and expected structure of such a film are not known, but the proposed atomic arrangement model can still be taken as a
relationship of the crystal lattices of the molybdenum substrate and the nickel film, with the oxide layer between having sufficient order to support the epitaxy. Therefore, it can still be concluded that it must be the minimisation of the interracial energy, rather than simply lattice matching, which is dominant in keeping the nickel lattice in such an orientation. A further complication is the high substrate roughness, which might be expected to affect the interfacial structure. For epitaxy based on lattice match to a single crystal substrate, roughness should be largely irrelevant, but it is not clear that the case is so straightforward here, although nuclei will generally form on single grains of the substrate. This question needs further investigation. The development of the texture with (002) parallel to the surface contrasts with the results at ambient temperature and with the observation that (111) planes have lowest surface energy in fcc metals. Nevertheless, it is consistent with the general rule that texture formation is driven by the minimisation of the system free energy, including interfacial free energy and elastic strain energy in addition to the free surface energy. It must be noted that the (002) planes of nickel have the second lowest free surface energy, behind only the (11 t) planes. This fact, together with the minimisation of the interfacial energy, puts (002) texture in a favourable position. The final structure is undoubtedly owing to the effect of anisotropic interfacial energy. Epitaxial growth of nickel film has also been observed in electrodeposition on polycrystalline bcc c~-Fe substrates, in which the orientation relationship is nearly the same as that of the K u r d j u m o v - S a c h s relationship, i.e. (111)Ni[J(110),~, [l~0]NiiI[l'[1] 6 [221. Again, interfacial energy plays an important role. The present results also contrast with previous findings for Er (hcp) films on Mo substrates, where Zone I textures are similar to those found in the present work, but Zone II samples show a strong (002) fibre orientation [5]. In addition, the erbium samples at Zone I-II and II-III transition temperatures are also different from Zone II samples, showing mixed fibre textures. It seems that surface energy minimisation dominates the texture formation of erbium films on molybdenum substrates. It is likely that anisotropic surface energy of the erbium (hcp) system overrides any possible anisotropic interfacial energy effect. From the atomic arrangement of Fig. 8(a), a tensile stress should exist in the nickel films. However, it is expected that the measured residual stress wilt be mainly from the thermal strain effect. The thermal stress, due to the difference in thermal expansion coefficient of nickel and molybdenum, can be estimated using the following equation, assuming a biaxial stress: E
l-
To)
(2)
where the thermal expansion coefficients aNi and aMo
Z. Shi et aL / T h i n Solid Films 304 (1997) 1 7 0 - 1 7 7
have values of 14.88 and 5.6 × t0 .6 K - t , respectively, for the range 0 - 4 0 0 °C [23]. This gives an estimated tensile stress of 980 MPa from 390 °C cooling to room temperature. It can be seen that this value is not far from the measured value. Therefore, the main contribution is indeed from the thermal strain and the stress caused by the lattice mismatch is not revealed. The latter is due to the microstructural adjustment in the very thin layer of the film joining the interface. Owing to the existence of grain boundaries and the introduction of a high density of defects like dislocations under tensile strain, the tensile strain will quickly diminish as the film grows. In fact, if the mismatch were kept throughout the film, it would result in a tensile stress of about 40 GPa, which is far too big for it to sustain without deforming or breaking the :film.
5. Conclusions UHV electron beam evaporated nickel films on textured molybdenum substrates show a mixed fibre texture at ambient temperature (SZM Zone I), with a strong (111) and a weak (002) component. From Zone I - I I to Zone I I - I I I , the orientation of the films is dominated by the substrate texture and granular epitaxy occurs. An atomic arrangement model has been put forward, which has an o r i e n t a t i o n r e l a t i o n s h i p o f (002)NilI(002)Mo and [100]Ni[l[ll0]Mo and would lead to the pole figure mirroring. The residual stress was found to be tensile, which is mainly due to the them~al strain, and the stress caused by the lattice mismatch is not revealed by the present stress measurements. It is believed that the texture formation in nickel films on molybdenum is dominated by the surface energy minimisation at ambient temperature, but is controlled by the anisotropic interfacial energy at elevated temperatures. The form of epitaxy found in this system at higher temperatures supports the theory that the minimisation of the system free energy rather than the lattice match is the detemlining factor in epitaxial film growth, as proposed by Thompson [17]. This is particuhtrly so given the expected presence of an intervening oxide layer.
177
Acknowledgements
W e thank EPSRC for support under G R / J 28032 and CCLRC for access to the Daresbury SRS.
References [1] B.A. Movchan, A.V. Demchishin, Phys. Met. Metallogr. 28 (1969) 83-90. [2] J.A. Thornton, J. Vac. Sci. Technol. 1i(4) (1974) 666-670. [3] D. Walton, Phil. Mag. VIII 7 (1962) 1671-1679. [4] M.A. Player, G.V. Marr, E. Gu, H. Savaloni, N. 0ncan, I.H. Munro, J. Apph Crystallogr. 25 (1992) 700-777. [5] G.R.G. Craib, M.A. Player, M.J. Rodman, C.C. Tang, Mater. Sci. Forum 228-231 (1996), part 2, 469-474. [6] C.C. Tang, M.C. Miller, S.M, Clark, M.A. Player, G.R.G. Craib, J. Synch. Rad. 3 (1996) 6-13. [7] S.P. Collins, Rd. Cernik, P. Pattison, A.M.T. Bell, A.N. Fitch, Rev. Sci. Instrum. 63(1) (1992) 1013-1014. [8] A.N. Fitch, C.R.A. Catlow, A. Atkinson, J. Mater. Sci. 26 (1991) 2300-2304. [9] F. Arrando, J. Esteve, J. Bassas, X. Alcob~, Mater. Sci. Forum 228-23I (1996), part 1,317-322. [i0] S. Krumm, Mater. Sci. Forum 228-231 (1996), part 1, 183-188. [I1] Goodfellow Cambridge Ltd, Catalogue 1995/96. [12] C.S. Barrett, T.B. Massalski, Structure of Metals, 3rd edn., Pergamon, Oxford. 1980. [13] D.B. Knorr, Mater. Sci. Forum 157-I62 (1994) 1327-1336. [14] C.R.M. Grovenor, H.T.G. HentzelI, D.A. Smith, Acta Metall. 32(5) (1984) 773-78I. [15] T. Zheng, Y. Shen, W. Wang, Z. Zhang, Z. Wang, Chin. Phys. Lett. 6(8) (1989) 377-380. [16] M.S. Miller, F.E. Stageberg, Y.M. Chow, K. Rook, L.A. Heuer, J. AppI. Phys. 75(i0) (1994) 5779-5781. [17] C.V. Thompson, Ann. Rev. Mater. Sci. 20 (1990) 245-268. [18] R.W. Vook, F. Witt, J. Vac. Sci. Technol. 2 (1965) 49-57, 243-249. [19] R.W. Vook, F. Witt, in: R. Niedermayer, H. Mayer (Eds.), Basic Problems in Thin Film Physics, Vandenhoeck and Rupprecht, Gbttingen, 1966, pp. 188-193. [20] F. Witt, R.W. Vook, S. Luszcz, Trans. Metallurg. Soc. AIME 242 (i968) 11ii-ii15. [21] JCPDS Powder Diffraction Cards, No. 4-850, 42-I 120, (1994). [22] K. Kamei, J. Jpn. Inst. Metals 52 (1988) 743-748. [23] K. Raznjevic, Handbook of Thermodynamics Tables and Charts, McGraw-Hill, New York, I976.
Thin Solid Fitms 304 (1997) 170-i77
Texture development of evaporated nickel films on molybdenum substrates Z. Shi a,* G.R.G. Craib a, M.A. Player a, C.C. Tang a Department of Engineering, Unicersity of Aberdeen, Aberdeen AB24 3UE, UK b Daresbury Laboratory, Warrington WA4 4AD, UK
Received 14 October 1996; accepted i7 February 1997
Abstract UHY electron beam evaporation of nickel thin films on molybdenum substrates was carried out at different temperatures between Zone I and low Zone III of the structural zone model. The film texture was studied by energy-dispersive X-ray diffraction and the residual stresses were measured by monochromatic X-ray diffraction, both using the Daresbury synchrotron radiation source. It was found that the nickel film deposited at ambient temperature (Zone I) had a mixed fibre texture with a strong (11 t) and a weak (002) component. From Zone I-II to Zone II-III, the orientation of the films was dominated by the substrate texture, and granular epitaxy occurred. The films had an imperfect (001)[100] sheet texture, with the (002) and (I 1 l) pole figures, respectively, mirroring the (002) and (01 t) figures of the substrate, which had an imperfect (001)[110] sheet texture. An atomic arrangement model is put forward to explain the pole figure mirroring. The residual stress was found to be tensile and is mainly due to the thermal strain. It is suggested that the minimisation of the free surface energy dominates the texture development of the nickel film at low temperatures, while the anisotropic interfacial energy plays the determining role at elevated temperatures. © 1997 Elsevier Science S.A. Keywords: Epitaxy; Nickel; Physical vapour deposition (PVD); X-ray diffraction
1. Introduction Microstructures of thin films play an important role in determining their properties and thus their applications. Many previous studies have led to the development of the structural zone model (SZM) for the broad description of polycrystalline film structures. The SZM was first proposed by Movchan and Demchishin [1] and has three zones, identified by their characteristic structures, separated by two boundary temperatures T 1 and T:. For metals, T I = 0 . 3 T m and T2 = 0.45Tm, where Tm is the melting temperature. The model was later extended by Thornton [2] to sputter deposited films by adding a third coordinate to account for the influence of the working gas pressure. An additional transitional zone (Zone T) was introduced, situated between Zone I and Zone II. Although the SZM gives a fair description of many film microstructures, it does not provide information on the crystallographic orientation of the films. Vacuum-de-
Corresponding author. Tel: +44 i224 272526; fax: +44 1224 272497. 0040-6090/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S0040-6090(97)001 10-7
posited or sputter-deposited metal films often have some preferred orientation, or texture, which may lead to desired or undesired anisotropic properties and may affect their potential applications. The formation of texture in thin films is a complex phenomenon. Although some work has been carried out, there is still lack of comprehensive understanding as to how and when texture might form in films. Many factors can influence the development of texture, including both the deposition conditions and substrate effects. The substrate temperature, which determines the film microstructures in the SZM, also has the greatest influence on the orientation of films. It is now generally accepted that a texture with the most densely populated plane parallel to the surface is favoured at low substrate temperatures [3]. During the study of ultrahigh-vacuum (UHV) evaporated erbium (hcp) films on molybdenum substrates, it has been found that the substrate temperature strongly influences the texture of the erbium films [4,5]. The erbium films show a mixed (002) and (i01) fibre texture at ambient temperature (Zone I), which changes to a more mixed population at the transition region of Zone I - I I . At
Z. Shi et al. 7Thin-Solid Films 304 ~I997) 1 7 0 - 1 7 7
Table 1 Details of samples Sample No.
Coating
Substrate
~ (°C)
Ts / Trn
SZM zone
Thickness (nm)
I 2 3 3A 4
Ni Ni Ni Ni Ni
Mo Mo Mo Mo Mo
27 200 390 390 510
0.17 0.27 0.38 0.38 0.45
I I-II II II tI-III
700 700 700 700 700
face roughness of Rq = 80 nm. The thickness was measured after deposition using a Rank Taylor Hobson Talystep profilometer. The base pressure was below 10 .8 hPa before coating was started, and the residual gas was composed mainly of HaO, CO/N2, CO: and H : , as shown by mass spectroscopic analysis. During deposition, the growth rate was kept at about 0.67 n m s -1 and the incident angle was approximately 8 ° to the substrate normal. A set of substrate temperatures (T~) was used, corresponding to the region from Zone I to low Zone III in the SZM, i.e. T~ = 27, 200, 390 and 510 °C. Details of the samples are shown in Table 1. Two samples deposited at 390 °C under the same conditions are included to demonstrate reproducibility. Both the texture and stress measurements were carried out at the Daresbury Laboratory synchrotron radiation source (SRS). A newly developed X-ray diffractometer [6] was employed to study the texture of the nickel thin films and the molybdenum substrates, using energy dispersive techniques on Station 9.7. The geometry of the goniometer has been described in Ref. [4] and has the azimuth stage (rh) above the tilt stage (c~), with the roll stage ( ~ ) below. This provides a constant tilt and hence a constant interaction volume and footprint, and is therefore particularly suitable for thin film or surface layer studies at or neargrazing incidence. In the present study, multi-bunch operation was used, with steps of 7.5 ° in roll (~), 45 ° in azimuth (&) and 120 s counting time per frame. Diffraction patterns were recorded at a glancing angle ce = 2 ° and diffraction angle 2 0 - - 10°. The roll range used was 0-60 ° and the azimuth range was 0-315 ° , resulting in 72 data points for one sample. The pole figures were then prepared
mid-Zone II temperatures, the films have a very strong (002) fibre orientation, which then changes to a mixed (002) and (100) texture at even higher temperatures (Zone n-III). In the present study, fcc nickel was selected for comparison with erbium, in view of its different crystal structure but similar melting temperature. Nickel thin films were deposited on molybdenum substrates using electron beam evaporation under UHV conditions. Different substrate temperatures between Zone I and low Zone III were selected in order to study the effects of deposition temperature. The film texture was studied by energy-dispersive X-ray diffraction and the residual stresses were measured by monochromatic X-ray diffraction, both using synchrotron radiation.
2. Experimental Nickel thin films approximately 700 nm thick were produced using a Balzers UMS 500 UHV electron beam evaporation plant, on molybdenum substrates with a sur-
Co)
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0
85000-4Z)00
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Fig. 1. Pole figures for sample 1 Ni film and Mo substrate, T~ = 27 °C. (a) Ni (111), (b) Ni (002), (c) Mo (011), (d) Mo (002). ~' range is 0 - 6 0 ° (at the centre of each side).
Z Shi et al./Thin Solid Fihns 304 (1997) 170-177
172
(a) • • • • • • • • [] [] []
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Fig. 2. Pole figures for sample 2 Ni film and Mo substrate, T~ = 200 °C. (a) Ni ( l i D , (b) Ni (002), (c) Mo (01 t), (d) Mo (002).
using an automated peak finder program. A correction for the loss of synchrotron beam current was applied, as was intensity correction using the procedure described in Ref. [4]. The figures were interpolated and plotted using UNIMAP, and plotted out to tp = 60 °. The residual stresses of the films were measured using the high-resolution monochromatic diffractometer on Station 2.3 at the SRS [7], with nominal wavelength set to 0.15 nm. Although not at grazing incidence, the parallel
beam optics used has many advantages compared with a conventional focusing Bragg-Brentano diffractometer [8]. The measurements were carried out using the sin2~ method with a texture goniometer. Use of a texture goniometer greatly increases the quality of tile results compared with a standard two-circle arrangement [9]. Ni (111) diffraction was selected for its higher intensity. A single-bunch beam was used with a 2 0 step of 0.005 ° and a counting time of 5 s. The peak positions (20) were then obtained using the
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Fig. 3. Pole figures for sample 3 Ni film and Mo substrate, Ts = 390 °C. (a) Ni (111), (b) Ni (002), (c) Mo (011), (d) Mo (002).
7.. Shi et al. / Thin Solid Films 304 (1997) 1 7 0 - 1 7 7
(a) 81000.
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[ ] I~03-24000 [ ] 8000-I~,:X) [] 0- @C00 [ ] ~LGW 0
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Fig. 4. Pole figures for sample 3A Ni film and Mo substrate, Ts = 390 °C. (a) Ni (111), (b) Ni (002), (c) Mo (011), (d) Mo (002).
orientation and the fibre texture is peaked at a direction approximately aligned to the deposition beam direction. A comparison of Fig. l(a) and (b) also reveals that the (111) texture is stronger than (002). The pole figures of the underlying molybdenum (Fig. lc,d) show an imperfect (001)[110] sheet texture. There is no obvious relationship between the orientation of the substrate and the film texture at ambient temperature. For the depositions at different elevated temperatures (samples 2, 3, 3A and 4), the pole figures of the nickel films and the underlying molybdenum substrate are shown
Winfit program [10] and the stresses were calculated from the slope of the 20 versus sin2~ plot.
3. Results
Fig. 1 shows the pole figures for the nickel film deposited at ambient temperature and the underlying molybdenum substrate. From the pole figures, it can be seen that the nickel film in Zone I has a mixed fibre texture. Both the Ni (111) and Ni (002) pole figures show a central fibre
(b)
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• •
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• [] []
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Fig. 5. Pole figures for sample 4 Ni film and Mo substrate, Ts = 510 °C. (a) Ni (i 11), (b) Ni (1302), (c) Mo (01 1), (d) Mo (002).
174
Z Shi et al./Thin Solid Films 304 (1997) 170-177
in Figs. 2-5. It can be seen that the pole figures of the molybdenum substrates are essentially as in Fig. 1, an imperfect (001)[l 10] sheet texture. The directions of the substrates were placed randomly and the similar direction of the substrates in Figs. 2 - 4 is merely coincidental. The textures of the films from low Zone II to low Zone III, however, are different from those of the Zone I film in Fig. 1. The orientation of the Mo substrate profoundly influenced the texture of these Ni films. Instead of a mixed fibre texture as at ambient temperature, the nickel films developed a (001)[100] sheet texture when deposited at elevated temperatures. Thus, the pole figures of the films and the substrate show a mirroring at these deposition temperatures, i.e. Ni (002) mirrors Mo (002), and Ni (111) mirrors Mo (011). It must be pointed out that the mirroring of Ni (111) and Mo (011) is only in the appearance of the pole figures and not in the exact pole positions. The Ni (1 t 1) poles are slightly wider apart from each other than the Mo (011) poles, corresponding to the inter-planar angle of 70.5 ° for Ni (111) and (]11) or Ni (1]1) and (~]1), and 60 ° for Mo (011) and (~01) or Mo (101) and (0]1). From these figures, it can also be seen that the change in the substrate temperature across the whole Zone II region does not have a significant effect on the texture of the films. Comparing Figs. 3 and 4 also shows that the two films deposited at 390 °C under the same conditions have the same preferred orientation, demonstrating the consistency of the results. Fig. 6 shows the variation in the (111) peak position against the value of sina~ for the nicket film deposited at 390 °C. The residual stresses were calculated from the slope of the 20 versus sin20 plot, with the following equation:
E o',
0(20)
2(1 -}-/-')
18--"-0rCOt0"
0(sin20)
(1)
where 0 is in degrees, °'4, is the stress parallel to the surface at an angle 4~ from principal stress axis 1 (see also Fig. 7), 0± the diffraction angle for the planes parallel to the surface, E is Young's modulus and v is Poisson's ratio.
g
48,20
~ 481s E
48.12
43.08 0
0.2
0.4
0.8
0.8
sin2¥
Fig. 6. Variationof (11 l) peak position with sin2~bof Ni film deposited at 390 °C (sample 3).
0" 2 = 4 7 5 MPa
~ 4 0 = 5 6 0 MPa
\
i°2 ~ k~ ,C~60=,480 MPa
Fig. 7. Schematic drawing of the residual stress in the Ni film deposited at 390 °C (sample 3). The orientation of the figure is the same as that of Fig. 3(a) and (b). Using the mechanical properties E = t99.5 GPa and v = 0.312 for nickel [11], a tensile stress of 610 + 40 MPa has been obtained from the slope in Fig. 6 for the particular 4, direction whose relationship to the substrate and film orientation is shown in Fig. 7. Measurements for two other directions at + 6 0 ° and - 6 0 ° to the above ~b direction were made with the two-point technique using ~0= 0 and 60 °. The stress values at these two directions were 560 _+ 40 and 480 _+ 40 MPa, respectively. From these results, the principal stresses o% and o-2 were calculated according to [12]. The values were found to be o-1 = 625 + 4 0 MPa, o-2 = 475 ± 40 MPa and ~ = - 19 ± 10°. The results are shown schematically in Fig. 7, where the orientation is the same as the corresponding pole figures (Fig. 3a,b). It is obvious that the residual stresses are tensile and show only weak anisotropy.
4. D i s c u s s i o n
The formation of texture is driven by the energy minimisation of the system. This often includes anisotropic surface energy, interracial energy and elastic strain energy in the film system. However, kinetics may decide what is actually occurring in a system, which may not be the most stable structure from a thermodynamics point of view. Therefore, the deposition conditions and the properties of the coating and substrate materials will all affect the final orientation and microstructure of films. Nucleation provides the first opportunity for texture formation since the orientation of vapour deposits is affected, and sometimes determined, by the structure of critical nuclei. Coalescence presents another opportunity for texture evolution as it occurs by an extensive rearrangement of materials of the islands that has been described as "liquid-like" or as "sintering" [13]. It is therefore expected that the nucleation and coalescence would play an important role in determining the film structure and orientation, especially in very thin films. As the film thickness increases, the role of grain growth will
Zo S h i e t aI. / Thin Solid Films 304 (1997) 1 7 0 - 1 7 7
become more and more pronounced as diffusion processes proceed. At low substrate temperatures, it is believed that the texture even in thick film is determined in the early stages of film growth and not by a selective growth mechanism during deposition [14]. In this case, the mobility of the vapour atoms on the substrate surface is low, and nucleation plays an important role in determining the film structure. Walton [3] proposed that the orientation of vapour deposits is determined by the structure of critical nuclei. It was predicted, as has also been found in many experiments, that high vapour supersaturation (low substrate temperature a n d / o r high incidence rate) favours the texture with the most densely populated plane parallel to the surface, but minimises the tendency for epitaxy on singlecrystal substrates, potentially leading to randomly oriented nuclei in the absence of a strong substrate effect. For fcc metals, the most likely texture is therefore (111) planes parallel to the substrate surface, followed by (002) and then (01 t). This is consistent with the fact that the free surface energy is lower with denser planes. It has been reported that nickel films evaporated at ambient temperature on amorphous carbon substrates have a slight preferred orientation with (111) planes parallel to the film surface even at a thickness of 2 - 4 mn [15]. '?his demonstrates the importance of nucleation and coalescence in texture formation. It is thought that the surface energy minimisation, possibly coupled with interfacial energy minimisation, plays the determining role in this system. The present results have shown that the nickel films have a stronger (111) and weaker (002) fibre texture at ambient temperature. Similar results have been reported for magnetron-sputtered nickel films on silicon wafers, which show a dominant (111) texture with a weak (002) component [16]. These results are consistent with Walton's model and indicate that free surface energy minimisation is the dominant driving force at ambient temperature for the texture development of nickel films. There is a possibility
175
that some (002) texture originates from epitaxial growth from the textured molybdenum substrate, which becomes dominant under the condition of higher substrate temperature discussed later. Such a case has been reported in as-deposited gold films on vacuum-cleaved mica, which are composed of grains with (111) texture, with a subpopulation of grains with epitaxial orientations [17]. This shows that inteffacial energy also plays an important role in the above situation. Nevertheless, whatever the reason for the (002) texture in nickel films, the present results show the dominant effect of the free surface energy on the overall texture at low temperatures. As said earlier, strain energy is one of the driving forces for the texture evolution. However, this is unlikely to be fully effective during film deposition. Instead, it will play a more prominent part during heat treatment of films where thermal stress can drive preferential grain growth through bulk migration of grain boundaries. Indeed, it has been reported that this had led to the development of (002) texture after annealing in copper, gold and silver films (fcc) on glass substrate and also in silver films on CaF 2 and NaC1 substrates [18-20]. The as-deposited texture at 80 K, however, was (111) and was also attributed to the surface energy minimisation. The texture of the nickel films deposited at elevated temperature are markedly different from that at ambient temperature. They all show an imperfect (001)[100] sheet texture (Figs. 2-5). The mirroring pole figures of the nickel films and the molybdenum substrates suggest that there should be an alignment of the nickel atoms to the molybdenum atoms. A description of this relationship depends on the structure of the interface, and in particular whether or not an oxide film is present. Ignoring for the present the question of an oxide film, it can be proposed that nickel atoms sit in the position corresponding to the valley of four (001) molybdenum atoms, as shown in Fig. 8(a). In this configuration, the positions of nickel atoms are an extension of the molybdenum lattice, and the nickel
Mo[010]
Mo[010]
Ni[010]
i
, ......
• Moil001
(a)
Ni[100] • Mo[I00]
Co)
Fig. 8. Arrangements of Ni atoms (small circle) over Mo (001) atoms (large circle). (a) Epitaxial configuration with four Ni (001) cells; (b) altemative configuration with four Ni (001) ceils.
176
Z Shiet al. / Thin Solid Films 304 (1997) 170-177
films grow epitaxially on the molybdenum substrate. This would lead to an orientation relationship of (002)Ni ]](002)Mo and [100]Ni H[110]Mo and would result in pole figure alignments as shown in Figs. 2-5. This configuration has a quite large lattice mismatch of 21%, using the data aN~ = 0.35238 nm and aMo = 0.31472 nm [21]. But this is not an unprecedented result; even higher misfits of 27% have been observed in gold, silver or copper films grown on sodium chloride cleavage surfaces [3]. However, taking account of lattice mismatch value alone, i.e. considering elastic strain energy as the only driving force, another atomic arrangement would be more favoured. This is shown in Fig. 8(b), with nickel atoms sitting in the position between two nearest (001) molybdenum atoms. It has a lattice mismatch of - t2%, substantially less than that shown in Fig. 8(a). Nevertheless, there are two obvious reasons to suggest that the configuration in Fig. 8(b) should not prevail. First, in the Zone II temperature range in which adatoms have sufficient mobility to adjust their positions, an individual nickel atom is more likely to sit in the valley of four molybdenum atoms than on the ridge between two molybdenum atoms. In other words, the interfacial energy with the configuration in Fig. 8(a) is lower than that with Fig. 8(b), because the valley is an interatomic potential minimum. Second, nickel atoms need to be compressed to fit into the positions in this configuration, which is less likely during early stages of vapour deposition process as a result of the interatomic potential between nickel atoms, especially when it requires the nickel atoms to stay on the ridge. In this case, therefore, variation of interfacial energy with orientation would be expected to dominate. The above model suggests that the alignment of the pole figures is due to the variation in the interfacial free energy with different orientations, rather than lattice matching alone. In fact, for polycrystalline films on single crystal substrates, it has been pointed out that lattice matching between the film and the substrate is not required for epitaxial grain growth, although it certainly promotes it [17]. All that is required is that there be a single orientation that minimises interfacial energy. From the above, it is proposed that the nickel films grow epitaxially on individual grains of the molybdenum substrate at elevated temperatures, due to the interracial energy minimisation. As the film thickens, these epitaxially grown grains continue to grow with the arriving deposit atoms taking up vacant lattice sites on the grain surfaces. This has been termed "granular epitaxy" [14]. In fact, the above atomic arrangement model is probably a simplification of the actual interface. In the present study, it is believed that any adsorbed contaminants are removed by heating under UHV before deposition. However, a thin, stable oxide layer is expected to be present on the molybdenum surface. The detailed characteristics and expected structure of such a film are not known, but the proposed atomic arrangement model can still be taken as a
relationship of the crystal lattices of the molybdenum substrate and the nickel film, with the oxide layer between having sufficient order to support the epitaxy. Therefore, it can still be concluded that it must be the minimisation of the interracial energy, rather than simply lattice matching, which is dominant in keeping the nickel lattice in such an orientation. A further complication is the high substrate roughness, which might be expected to affect the interfacial structure. For epitaxy based on lattice match to a single crystal substrate, roughness should be largely irrelevant, but it is not clear that the case is so straightforward here, although nuclei will generally form on single grains of the substrate. This question needs further investigation. The development of the texture with (002) parallel to the surface contrasts with the results at ambient temperature and with the observation that (111) planes have lowest surface energy in fcc metals. Nevertheless, it is consistent with the general rule that texture formation is driven by the minimisation of the system free energy, including interfacial free energy and elastic strain energy in addition to the free surface energy. It must be noted that the (002) planes of nickel have the second lowest free surface energy, behind only the (11 t) planes. This fact, together with the minimisation of the interfacial energy, puts (002) texture in a favourable position. The final structure is undoubtedly owing to the effect of anisotropic interfacial energy. Epitaxial growth of nickel film has also been observed in electrodeposition on polycrystalline bcc c~-Fe substrates, in which the orientation relationship is nearly the same as that of the K u r d j u m o v - S a c h s relationship, i.e. (111)Ni[J(110),~, [l~0]NiiI[l'[1] 6 [221. Again, interfacial energy plays an important role. The present results also contrast with previous findings for Er (hcp) films on Mo substrates, where Zone I textures are similar to those found in the present work, but Zone II samples show a strong (002) fibre orientation [5]. In addition, the erbium samples at Zone I-II and II-III transition temperatures are also different from Zone II samples, showing mixed fibre textures. It seems that surface energy minimisation dominates the texture formation of erbium films on molybdenum substrates. It is likely that anisotropic surface energy of the erbium (hcp) system overrides any possible anisotropic interfacial energy effect. From the atomic arrangement of Fig. 8(a), a tensile stress should exist in the nickel films. However, it is expected that the measured residual stress wilt be mainly from the thermal strain effect. The thermal stress, due to the difference in thermal expansion coefficient of nickel and molybdenum, can be estimated using the following equation, assuming a biaxial stress: E
l-
To)
(2)
where the thermal expansion coefficients aNi and aMo
Z. Shi et aL / T h i n Solid Films 304 (1997) 1 7 0 - 1 7 7
have values of 14.88 and 5.6 × t0 .6 K - t , respectively, for the range 0 - 4 0 0 °C [23]. This gives an estimated tensile stress of 980 MPa from 390 °C cooling to room temperature. It can be seen that this value is not far from the measured value. Therefore, the main contribution is indeed from the thermal strain and the stress caused by the lattice mismatch is not revealed. The latter is due to the microstructural adjustment in the very thin layer of the film joining the interface. Owing to the existence of grain boundaries and the introduction of a high density of defects like dislocations under tensile strain, the tensile strain will quickly diminish as the film grows. In fact, if the mismatch were kept throughout the film, it would result in a tensile stress of about 40 GPa, which is far too big for it to sustain without deforming or breaking the :film.
5. Conclusions UHV electron beam evaporated nickel films on textured molybdenum substrates show a mixed fibre texture at ambient temperature (SZM Zone I), with a strong (111) and a weak (002) component. From Zone I - I I to Zone I I - I I I , the orientation of the films is dominated by the substrate texture and granular epitaxy occurs. An atomic arrangement model has been put forward, which has an o r i e n t a t i o n r e l a t i o n s h i p o f (002)NilI(002)Mo and [100]Ni[l[ll0]Mo and would lead to the pole figure mirroring. The residual stress was found to be tensile, which is mainly due to the them~al strain, and the stress caused by the lattice mismatch is not revealed by the present stress measurements. It is believed that the texture formation in nickel films on molybdenum is dominated by the surface energy minimisation at ambient temperature, but is controlled by the anisotropic interfacial energy at elevated temperatures. The form of epitaxy found in this system at higher temperatures supports the theory that the minimisation of the system free energy rather than the lattice match is the detemlining factor in epitaxial film growth, as proposed by Thompson [17]. This is particuhtrly so given the expected presence of an intervening oxide layer.
177
Acknowledgements
W e thank EPSRC for support under G R / J 28032 and CCLRC for access to the Daresbury SRS.
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