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Growth of ultrathin iron films on Cu(001): an ion-scattering spectroscopy study
Surface Science 293 (1993) 227-238 worth-~oIIand
Growth of ultrathin iron films on Cu( 001): an ion-scattering spectroscopy study Th, Detzei, N. Memmel and Th. Fauster
Received 12 February 1993; accepted for pubIication 26 April 1993
Low-energy ion scattering has been applied to study the pseudomorphic growth of ultrathin fee iron films on a Cu(OO1) single-crystal surface. At all coverages - even as low as 0.1 monolayer (ML) - iron is found to about equal amounts in both the surface and the first subsurface layer, since part of the iron atoms are incorporated into the original copper surface. This is particular evident for the 0.1 ML film, which does not exhibit any iron defects, such as steps or adatoms. Instead copper surface defects become visible upon iron deposition, Near 2 ML the substrate is covered for the most part by a relatively smooth b&et indicating the coalescence of the iron islands. Up to a film thickness of around 6 ML the surface defect density of the iron layers decreases with increasing coverage and raises again at larger coverages. Above N 10 ML the pseudomorphic fee growth breaks down aad domains of bee iron with (110) or~entatioo are formed.
The structure and morphology of ultrathin magnetic films has been of growing interest within the last 15 years for both fundamental and technological reasons. These systems serve as an outstanding opportunity to study basic concepts of theoretical physics like two-dimensional phase transitions and can be used to investigate the scaling concepts for the~ody~ami~ quantities. Moreover, thin magnetic fiIms with the easy axis of magnetization normal to the surface are of prominent technological importance in opto-magnetic storage devices. In the monolayer range of epitaxialiy grown fee iron films, such a perpendicular spin orientation was predicted [I] and experimentally observed [2-71. Since the lattice constants of the face centered cubic phase of iron (y-Fe), extrapolated to room temperature, and fee copper differ less than I%, the system Fe on Cu(OO1) seems to offer unique possibilities to produce ultrathin films of pseudomo~hic fee iron. In recent years, several studies by low-energy electron dif~action ~39~6OZS/93/$06.~
fLEEID), in which the overlayer structure is shown to be nearly identical to a ~ntinuation of the fee lattice of the substrate, corroborated this supposition [8-lo]. According to thermodynamic cansiderations one would expect the initial growth of Fe on Cu(OO1) not to proceed via a perfect layer-bylayer manner, Bauer [ill pointed out that the surface free energies of the substrate and the adsorbate together with the interfacial free energy are in a first appro~mation the thermodynamic quantities dete~ining the growth mode. The the~~~arni~ parameters of the system under desideration fulfill the requirement for island formation [12], i.e., the sum of the surface free energy of Fe and the interfacial free energy between Fe and Cu exceeds the surface free energy of the Cu substrate. Furthermore, these energetic circumstances may lead to the diffusion of Cu substrate atoms onto the surface of the adsorbed iayers in order to minimize the total free energy. Thus, we are dealing with the general problem in which a high surface free-energy metal is attempted to be epitaxed on a low surface free-energy metal. Viewing the chart of the
0 1993 - Elsevier Science Publishers B.V. AI1 rights reserved
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measured elemental surface free energies [131 it becomes evident, that this is the rather usual case concerning epitaxy in magnetic ultrathin film research and application. Therefore, the combination Fe/Cu might be considered as a prototype system. Nevertheless, it is worth mentioning that the difference of the two sides in the free-energy balance is not extremely significant. Since the Bauer criterion concerns only thermal equilibrium and because the growth of thin metal films by vapour deposition is a dynamic process usually being far from true equilibrium, microscopic mechanisms might strongly influence the film growth. In addition, any strain energy due to the accommodation of the lattice mismatch between the two materials is neglected in the Bauer criterion. A large amount of work has been dedicated to study the growth mode of ultrathin Fe films on Cu(OO1) substrates in recent years. While there is general agreement now, that iron grows layer-bylayer for films with a thickness between 5 and N 10 ML, the experimental results and their interpretations concerning the initial stage of the growth up to 5 ML are still quite controversial. Germar et al. [14] claimed from their Auger electron spectroscopy @ES) investigations that the growth mode is essentially layer-by-layer, the deviation from perfect monolayer growth being less than 10%. At the same time Steigerwald et al. [12] found the film growth not to proceed via a Frank-van der Merwe mechanism, but discovered Fe agglomeration and Cu surface segregation using XPS forward scattering and the carbon monoxide titration technique. Glatzel et al. [151 report on the formation of bilayers for the first 6 ML. The authors observed the characteristic breakpoints in the Auger uptake curve recorded during evaporation and used Rutherford backscattering spectroscopy (RBS) to determine the nominal film thickness, clearly revealing that the breakpoints occur at the completion of bilayers. Thomassen et al. [16] investigated growth, structure and morphology of ultrathin iron films utilizing low- and medium-energy electron diffraction (LEED, MEED) as well as AES. This group reports on Fe agglomeration up to about 5 ML followed by layer-by-layer growth for coverages
between 5 and 14 ML. Finally, at around 14 ML due to the formation of misfit dislocations the breakdown of epitaxy was observed. The preceding findings tally more or less with results presented by Arnott et al. who applied specular helium atom scattering techniques [17]. In a recent scanning tunnelling microscopy (STM) study an even more complex growth behaviour is stated [HI. Submonolayer iron deposition (0 = 0.1 ML) yields the f?rmation of small features of typical size N 10 A, interpreted as iron patches in the original copper surface layer where iron atoms have replaced copper atoms, and larger monolayer islands, formed by the “expelled” Cu atoms. The experimental results concerning the magnetic properties of thin iron films on Cu substrates seem to be as contradictory as those obtained for the film morphologies. Different groups, for instance, have reported the presence of ferromagnetism on fee iron grown on Cu(OO1) 12,191,whereas the absence of ferromagnetism at room temperature has been stated by others 120,211. In the latter experiments, however, iron was deposited at a substrate temperature of 200°C thus leading to a pronounced copper diffusion to the top of the Fe film [12], which probably is responsible for the different magnetic behaviour observed for these films. This already shows, that a precise knowledge of the epitaxial growth process is required to understand the magnetism of ultrathin films. In this paper we report on a first study of the growth mechanisms of ultrathin Fe films on Cu(OO1) using low-energy ion scattering (LEIS). Special emphasis is put on the low coverage regime. In recent years, the use of ion-scattering methods has been established as an advantageous approach to study surface, interface and thin film phenomena [22-261. The reason for this is manifold: Like electron spectroscopy methods involving core electrons, ion-scattering spectroscopy is element specific as well as quantitative and therefore can be used for a chemical analysis of the components to be found at the surface. Because of its enormous surface sensitivity caused by neutralization and blocking effects, it is possible to probe exclusively the outermost surface layers of the sample. Moreover, due to shadowing effects
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this method provides structural information with monatomic resolution in the surface plane as well as normal to the surface. In contrast to Fourierspace techniques like electron and atom diffraction LEIS supplies mass-selective real-space information on the local atomic arrangement at the surface including surface defects. As a further advantage it is noteworthy that the evaluation schemes are straight forward without requiring sophisticated and time-consuming model calculations.
2. Experiment The experiments were performed in an UHV system with a base pressure below 3 x lop9 Pa. A (OOl)-oriented Cu crystal (miscut < rtO.3”) was prepared prior to each iron deposition by successive cycles of Ar+ bombardment (600 eV) and annealing to 900 K until the contamination levels were below the detection limits of LEIS. After the last annealing the sample was cooled down slowly (< 2 K/min) in order to reduce the number of possible surface defects. The crystallographic structure of the surface was monitored by LEED, showing a well-defined (OOl)-LEED pattern with sharp spots and low background intensity, as well as by impact collision ion-scattering spectroscopy (ICISS) [22-261, revealing a well ordered copper surface. Subsequently, iron was evaporated onto the Cu(OO1) crystal, held at 298 + 2 K, from a small iron pin of purity 99.999% heated by electron bombardment. The vapour source was surrounded by a water-cooled copper tubing in order to remove most of the radiant heat, thus keeping the levels of contaminating gases small. During evaporation the pressure stayed below 1 X 10-s Pa. An analysis of the residual gas by a quadrupole mass spectrometer revealed that the slight increase in pressure was mainly due to an increase in the partial pressure of hydrogen. The coverage was calibrated independently by AES and MEED and is given in units of a nominal monolayer (ML), i.e. 1.53 x 1015 atoms/cm*. The MEED experiment was carried out using a 4 keV electron beam at grazing incidence (N 100).
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The diffraction pattern was displayed on the screen of a reverse view LEED optics and the intensity variations of several reflexes during Fe deposition were recorded simultaneously by a computer controlled video system. The characteristic intensity variations reported by Thomassen et al. [16] were reproduced and the time necessary to deposit one monolayer of Fe was inferred from the period of the regular oscillations observed in the coverage range between 5 and 10 ML. The evaporation rates ranged from 0.5 to 1 ML/min in the different experiments. In an alternative calibration procedure the intensity ratio of the 705 eV iron and 920 eV copper Auger transitions was determined and compared to the values given by Glatzel et al. [151, who used Rutherford backscattering to measure the nominal iron coverage corresponding to distinct AES intensity ratios. Both calibration procedures agree well with each other. However, care has to be taken, as different experimental setups were used in the AES experiments of Glatzel et al. and our work. (Glatzel: Excitation by a 3 keV electron beam at 10” grazing incidence. Detection by a retarding field analyzer with a mean detection angle of 61”. This work: Excitation by a 5 keV electron beam at 10” grazing incidence. Detection by a retarding field analyzer with a mean detection angle of 48”.) From the MEED and AES experiments we specify the coverages given in this publication with an experimental error of about $- 10% but less than +0.3 ML. As will be discussed below, the LEIS experiments themselves also give a lower limit of the amount of Fe deposited onto the surface. All ion-scattering experiments shown in this paper were carried out using a monoenergetic beam of Ne+ ions with a kinetic energy of 5 keV and an electrostatic analyzer at an scattering angle of 160”. In order to avoid considerable damage of the surface by the impinging ions the primary intensity as well as the recording time were kept low, yielding a maximum fluence of 5 X 101* to 5 X 1013 ions/cm*, depending on the kind of spectrum to be taken. No measurable radiation damage was observed under these conditions. Details about the experimental setup are described elsewhere [23].
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3, Results
To study the growth mode of iron on Cu@Ol) at room temperature ion-scattering experiments were performed for various film thicknesses ranging from 0.1 to 14 ML. However, our main interest concentrated on the low coverage regime. We were especially concerned with the possibility of island formation at coverages up to 1 nominal ML and the phenomenon of copper diffusion during deposition. Fig. 1 shows the intensity of backscattered Ne + ions as a unction of their kinetic energy in the energy range predicted by the binary collision model [23]. The film thickness amounted to 1.3 ML and the angle of incidence 2Fr was chosen equal to 40”, where - as will be explained below - only the outermost surface layer is probed. For perfect monolayer growth one would expect the Cu intensity originating from the first layer to vanish at a nominal coverage exceeding I ML. Obviously this is not the case. Instead, the copper signal (at N 1330 eV> is only slightly less than the iron yield (at _ 1100 eVI, even for a coverage somewhat larger than I ML. At this point it should be mentioned that due to the difference in nuclear charge as we11 as due to conceivable deviations in the neutralization probabilities of
backscattered
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Fig. 1. LEE spectrum of 1.3 ML Fe/CuRMU), recorded for a 5 keV Ne’ beam incident at an angle of 40” relative to the surface plane along the [lo01 azimuth. The scattering angle was 160*. The peaks at 1100 and 1330 eV result from scattering by iron and copper atoms, respectively.
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Fig. 2. Normalized intensity of Ne+ ions backscattered from Fe atoms in the outermost layer as a function of coverage. Dotted, solid and dash-dotted tines: signals expected for monohi- and trilayer growth, respective&.
iron and copper the “scattering cross sections” differ to a small extend. This difference was determined experimentally using pure elemental standards. We found a 10% lower “cross section” for iron compared to copper. This difference is small and has essentially no influence on the qualitative and quantitative interpretations drawn below. Therefore, we infer already at this point that one nominal monolayer of Fe is not spread out homogeneously on the Cu substrate, indicating a non-wetting tendency. Similarly to the well-known Auger uptake curves [Zi’] the Fe/(Fe + Cu) LEIS ratio is depicted as a function of nominal film thickness in fig. 2. The signal increases almost linearly with iron coverage up to about 2 ML where a distinct breakpoint in the slope is observed. Finally, beyond 4 ML the residual copper signal amounts to merely a few percent. This contribution may be attributed to tiny non-covered substrate areas as well as to Cu atoms which have diffused onto the film surface. Near a Fe coverage of 2 ML the Cu intensity disappeared for the most part, but there is stiI1 a non-negligible portion of copper visible in the LEIS measurements. Even so, a three-dimensional island formation (Volmer-Weber growth mode) can be excluded since in this case one would expect the adsorbate to leave large parts of the substrate uncovered. As mentioned earlier, Glatzel et al. claim the growth of Fe to proceed in a bilayer fashion.
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Neglecting Cu diffusion and interm~ng this growth mode would lead to an almost linearly increasing LEIS signal ratio reaching its saturation value at a nominal film thickness of 2 ML, as shown by the solid line in fig. 2. The line is not perfectly straight due to the 10% difference in “scattering cross section” between iron and copper. The dotted and dash-dotted lines show the intensity ratios expected for monolayer or trilayer growth, respectively. The experimental data are found between the two model curves for bilayer and trilayer growth, deviating less from the bilayer case. The assumption of u 0.1 ML of copper diffusing onto the film surface during deposition would explain the difference between the measured values and the ideal bilayer uptake curve. Thus, from this first LEIS experiments a distorted bilayer growth mode may be infered for the first two monolayers. In order to gain more insight into the microscopic growth processes, Fe films of different thicknesses were investigated by impact collision ion-scattering spectroscopy WISS). This method works at scattering angles near 180” and is based on the shadow cone concept [ZS]: The primary ion flux is deflected by the target atoms in such a way that there is no flux immediately behind them ((‘shadow cone”, see fig. 3). The ion trajectories pushed out of this shadow cone are focused to the edge of the shadow cone, yielding an increased ion flux in this region. At larger distances from the cone axis the flux decreases to its average value. The experimentai procedure is as follows: the electrostatic energy analyzer is tuned to the residual energy of the ions backscattered by the chemical species to be probed (i.e., N 1100 eV for Fe, N 1330 eV for Cu, see fig. 1). Then the backscattered intensity is measured as a function of the angle of incidence tv. Near grazing incidence (?P = O? on an ideal crystalline surface each target atom is hidden in the shadow cone of its nearest neighbour in the surface plane. No head-on collisions, as required for 180” backscattering can occur and the backscattered intensity vanishes. If the angle of incidence is increased, the atoms of the topmost layer will eventually move out of the shadow cone of the adjacent atoms (see fig. 3a), leading to a steep increase of
on Cu(oO1~
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Fig. 3. The shadow-cone concept of impact-collision ionscattering spectroscopy (ICISS). Only the trajectories of the incoming beam are shown in this figure. (a) Situation where atoms in the outermost layer just “leave” the shadow cone of their nearest neighbours in the surface plane. (b) Atoms in the first subsurface layer “leave” the shadow cone caused by the surface atoms.
the backscattered intensity at the critical angle !P,, (the indices indicate shadowing from a first layer atom onto a first layer atom). This is nicely seen in fig. 4a, where an ICISS spectrum of 5 keV Ne+ ions backscattered from the Cu atoms of the clean Cu@Ol) surface is displayed. The ion beam was impinging along the [loo] azimuth of the sample. The critical angle here amounts to W,, = 13”. Further increase of the angle of incidence P yields a decreasing backscattered intensity as a consequence of decreasing flux away from the shadow cone edge. Upon further increase of W the neighbours in the second Iayer are also Ieaving the cone (fig. 3b) and a second enhancement . . . of the ion yield 1s observed at ?Pylz= 63” (see fig. 4a). Simultaneously Nef ions are focused from atoms in the second layer onto third layer atoms. Ions backscattered from the third layer, however, do not reach the detector since they are blocked on the outgoing path by the substrate atoms directly above in the first layer. As can be seen, the signal originating from the second layer is approximately reduced by a factor of 4 relative to the first layer signa at lyii = 13”, mainly due to the higher neutralization rate for Ne+ ions backscattered from the second layer. It should be mentioned that this rather simple concept, involv-
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ing only a single large scattering event (plus a small angle deflection due to the shadowing atom) works only well for scattering angles near Ho”, since only in this case multiple scattering events, which would complicate the interpretation of the data, can be neglected. Fig. 5 shows a selection of ICISS spectra taken at various iron film thicknesses. The energy analyzer was set to an energy of 1100 eV, so only Ne+ ions scattered from Fe were detected. From these spectra the following statements can be made: (1) The gross features of all spectra up to N 10 ML are the same as on the clean Cu(OO1) surface (fig. 4a). This confirms that the iron films grow indeed in an fee structure with similar interatomic distances as the copper substrate. Since the atomic numbers of Fe and Cu are similar, the shadow cones caused by Cu and Fe atoms, re-
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Fig. 5. KISS spectra for various Fe films on Cu(OO1). The Nef ions are backscattered from iron. The scattering plane was along the [lo01 azimuth. All spectra are normalized to same height at maximum intensity.
0.7 ML
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Fig. 4. KISS spectra for (a) the clean Cu(OO1) surface and (b)-(e) the surface covered by various amounts of iron. The Nef ions are backscattered from copper. The scattering plane was along the [lOO] azimuth. All spectra are normalized to same height at maximum intensity.
spectively, are practically identical. The associated critical angles differ less than 0.5”, which is about our experimental uncertainty in the determination of the critical angle. The 14 ML spectrum is qualitatively different. Both critical angles are shifted, with the first layer focusing peak at *ii broadened considerably towards higher angles and the second layer feature at 1yi2 split into a double peak structure. This indicates the breakdown of pseudomorphic growth and the formation of bee Fe with a (110) surface normal [7]. As the unit mesh of the bee (110) surface is of two-fold symmetry only, four domains will exist on an fee (001) surface, but only two of them are inequivalent along the [lOOI-azimuthal direction of the Cu(OO1) substrate, where the measure-
Th. Lletzelet al./ Growth of ultrathinironfilms on CufOOl)
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Fig. 6. Normalized intensity of Net ions scattered from Fe atoms in the first subsurface layer as a function of coverage.
ments were performed. This two inequivalent domains produce the observed broadening and splitting of the ICISS structures [29]. Iron films with an fee or bee (OOl)-surface cannot explain these split or broadened features. (2) Both, the intensity increase at PI1 = 13” due to focusing onto top-layer Fe atoms as well as the intensi~ increase at ‘I’ = 63” due to focusing onto subsurface Fe atoms are observed for all coverages. Note that this even holds for coverages as small as 0.1 ML. This clearly demonstrates that even at such low coverages Fe atoms are present in both the top layer and the first subsurface layer. The ratio of the intensity increases at both critical angles is independent of coverage and amounts to - l/4, which indicates that about equal amounts of iron atoms exist in the surface and the first subsurface layer. (3) This statement is further substantiated in fig. 6, where, as a function of film thickness, the Fe/lFe + Cu> ratio of the intensity increase at P,, is plotted. This ratio is essentially proportional to the percentage of Fe atoms in the first subsurface layer. The iron content in this layer increases almost linearly with coverage up to about 80% at a nominal coverage of 2 ML. For higher coverages it converges slowly towards 100%. (4) The LEIS data presented in figs. 2 and 6 also confirm the coverage calibration established by MEED and AES. At a nominal coverage of 2 ML LEIS shows that 90% of the surface atoms
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and about 80% of the subsurface atoms are iron. Therefore at least 1.7 ML of iron have been deposited on the surface at this point. As iron atoms in the third and deeper layers do not contribute to the backscattered signal in our detection geometry, this is a lower limit of the amount of iron present on the surface. Fe atoms in the third and deeper layers may either be due to Fe agglomeration into islands of more than bilayer height or due to Cu atoms which have diffused onto the surface during the growth process. (5) Further info~ation regarding the microscopic growth process can be infered from the backscattered intensity near grazing incidence. As mentioned in the beginning, on a perfect surface no ions are backscattered for incidence angles below the first critical angle P,,. Upon deposition of iron, however, considerable intensity is observed at small angles of incidence (V < lo”, shaded regions in fig. 5). Adatoms, steps or vacancies have to be made responsible for this observation. Common to all of them is, that the back~attering atom has no neighbour along the nOO] direction. As they cannot be distinguished in the experimental spectra, we will summarize the foregoing terms in the following paragraphs by the collective expression “surface defects”. In fig. 7 the backscattered ion yield at W = 5” normalized to the ion yield at 20” (taken from fig. 5) is depicted as a function of film thickness to provide a means for the concentration of Fe surface defects and the quality of the iron films. This value increases strongly up to about 1 ML, thus revealing the poor quality of the Fe films near this coverage. This is in agreement with the pronounced intensity minimum observed in MEED [16] and He-atom scattering 2171. Near 2 nominal ML the defect intensity is reduced remarkably, pointing to a rather smooth film surface, which therefore can be interpreted as an almost completed bilayer. The 6 ML film shows a defect intensity comparable to the clean substrate. Therefore, we deal with an almost perfectly ordered Fe-film. Elevated coverages exhibit again a rougher surface morpholo~ due to the breakdown of pseudomo~hic growth and the formation of bee iron above - 10 ML.
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of ultrathin iron f&s on CufoOll 0.1 ML Fe deposited at 140 K
‘Fe surface defects’
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Fig. ‘7. Iron surface defect ~ncen~ation versus coverage as determined from the ratio of the LEIS intensity of Nef ions backscattered from iron at incidence angles of 5” and 20”, respectively. For comparison the copper defect intensity observed on the clean copper surface is indicated by the dashed line.
(6) The ICISS spectrum of the 0.1 ML film (fig. Sa> is pa~icularly interesting, as there is essentially no intensity observed at low angles of incidence. Note, that by the normalization procedure applied in fig. 5, the trivial coverage dependence of the signal has already been removed. No Fe surface defects are visible at this coverage. This only seems possible if iron forms large perfectly ordered islands or if the iron (rich) regions are completely surrounded by copper atoms, either due to incorporation of the Fe atoms into the substrate surface or due to decoration of the iron (rich) islands by copper atoms. The first explanation must be ruled out, because an only slightly higher amount of evaporated Fe yields already considerable defect densities (fig. 5b). (7) Further evidence for iron incorporation or decoration of the iron (rich) islands comes from the ICISS spectra in figs. 4a-4e. Here the electrostatic analyzer was tuned to an energy of 1330 eV, corresponding to collision of the Nef ions with copper atoms. The attention is immediately attracted to the elevated Cu defect intensity at smaI1 angles of incidence appearing upon iron deposition (shaded regions in fig. 4). This unambiguously proves that the copper surface layer
angle of incidence [“I Fig. 8. Iron-ICISS spectra for 0.1 ML Fe/Cu(OOl) deposited at 140 K. (a) As-deposited. (b) Annealed to room temperature for 10 min.
does not remain unaffected by adsorption of iron. Both scenarios, the formation of additional Cu islands by the “expelled” Cu atoms or decoration of Fe islands by Cu atoms, can explain the observed Cu defect intensity. In any case the data clearly demonstrate, that even at room temperature significant diffusion of Cu atoms, involving considerable mass transport, occurs. 3.2. Temperatureeffects To study the influence of the substrate temperature during Fe deposition on the growth mode, we performed additional evaporation experiments at cold and at slightly heated samples. Fig. 8a shows an iron ICISS spectrum recorded on a 0.1 ML film deposited at 140 K. Here the situation is changed dramatically as compared to room temperature experiments. An enormous defect density is observed since the small angle intensity exceeds even the peak of regular first layer terrace sites. Subsurface iron in fee sites is also detected. This spectrum drastically demonstrates the reduced mobility of the iron atoms at a temperature of 140 K. The ~~esponding copper spectrum (not shown), however, exhibits an
Th. Detzel et al. /Growth of ultrathin iron films on Cu(OO1)
elevated intensity at small angles of incidence, similar to fig. 4b. Therefore - even at a temperature of 140 K - iron induces a rearrangement of the Cu surface atoms. Annealing to room temperature improves the quality of the iron films as can be seen by the strong reduction of the iron defect peak (fig. 8b). However, it seems remarkable that iron surface defects are still visible, in contrast to adsorption at room temperature (fig. 5a). Simultaneously with the decrease of the surface defect peak an increase of the first layer peak near ?Prr is observed upon annealing, thus showing the “conversion” of Fe atoms with no nearest neighbour along [iOO] into atoms with a neighbour along this direction. The second layer contribution is only slightly affected by the annealing procedure. Keeping in mind the desire for a single fee iron monolayer lying flat on the substrate we deposited just under 1 ML onto a cold sample. We found again a rather rough film surface with many atoms buried in the second layer. Annealing the samples yields a spectrum comparable to fig. 5c, with a decreased low-angle intensity (i.e., a reduced number of iron surface defects) and a mostly unchanged second layer feature. Therefore it is not possible to manufacture an iron monolayer, using this preparation procedure. A similar experiment was carried out for a 3 ML film. Again many iron surface defects were visible for the film deposited at low temperature. Annealing to room temperature improved the quality of the film to a level comparable to that of a room-temperature grown 3 ML film. In a further experiment we deposited about 1 ML of iron at a temperature slightly above room temperature (340 K). We observed ICISS spectra with diminished defect intensity but with an excessive intensity raise at ?P’,, due to an enhanced iron percentage in the first subsurface layer. We estimate the number of iron atoms in the subsurface layer to exceed the top layer concentration by about 30%, which points without any doubt to enhanced diffusion of Cu atoms onto the iron film. The diminished defect density indicates a smoothing effect of the segregated substrate atoms, which was also observed in He atom scattering experiments [ 171.
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4. Discussion Our experimental findings on the growth of ultrathin iron films on Cu(OO1) single crystal surfaces are in good agreement with a previous study by photoelectron diffraction [121, which reports on bilayer island formation in the low-coverage regime as well as on the appearance of Cu segregation already at room temperature. The results published by Germar et al. [14], who claim the growth to proceed via a undisturbed Frank-van der Merwe mechanism are unambiguously disproved because always about 50% of the iron atoms are buried in the subsurface. As described in the introduction, Glatzel et al. [15] assume the formation of perfect bilayers in order to explain their Auger and RBS data. This interpretation is compatible with our data at least for the first two monolayers, even though at 2 ML the film is not perfectly ordered and its formation is accompanied by the possibility of some Cu segregation and intermixing. Thomassen et al. [161 suggest for the Fe-on-Cu system iron agglomeration up to a thickness of 5 ML. Our real-space ion-scattering data, however, indicate that the iron islands coalesce already at about 2 ML, leading to an almost pure (> 90%) iron surface beyond this coverage. Furthermore, the order in the films between 2 and N 9 ML deviates only slightly from that of the clean Cu(OO1) surface. Both findings, coalescence of the iron islands around 2 ML and the formation of rather smooth films above this coverage (up to the breakdown of epitaxial growth) is also reported in the most recent STM study of Brodde and Neddermeyer [30]. Therefore we are led to conclude that the strong intensity variations of various MEED spots between 2 and 5 ML - as reported in ref. [16] and reproduced in the present work - do not reflect changes in the morphology of the Fe films but are related to structural rearrangements of the surface as indicated by the (4 x l), (5 x 1) and (2 x 1) superstructures observed by LEED for coverages of about 2, 4 and 6 ML [161. These structural changes, however, involve only small displacements ( - 0.1 A> of the Fe atoms from their ideal (1 x 1) positions in such a way, that (to first order) the mean
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of ultrathin iron films on Cu(OOl)
nearest-neighbour distances remain unchanged [31,32]. Therefore these structural changes do not change the critical angles in the ICISS spectra, but only lead to a slight (- 1”) broadening effect. At smallest coverages we state a qualitative agreement of our LEIS data to the STM investigations of Chambliss et al. 1181.In particular, they report on Fe incorporation and the emergence of Cu islands on the surface, in agreement with the conclusions based on our LEIS results. Quantitatively, however, some discrepancies exist. According to the interpretation of Chambliss et al. the copper islands are formed by the Cu atoms expelled from the original surface layer by the incorporated Fe atoms. By comparison of the Cu surface defect ICISS intensity measured for the 0.1 ML iron film with that of a miscut Cu crystal [231, we estimate, that about 0.05 ML of copper is visible as surface defects (i.e., as atoms at the steps of islands and terraces). If the expelled Cu atoms agglomerate in rather large islands this means, that - contrary to the suggestion of Chambliss et al. - the Fe-Cu replacement does not simply occur on a one-to-one basis, but that additional Cu atoms are coming from steps, as suggested in the STM work of Brodde and Neddermeyer [30] from the simultaneously observed step roughening. An alternative explanation would be, that the atoms at the perimeter of both the large monolayer islands and the small patches, seen in the STM work, are copper atoms. ’ One of the remaining puzzles of the Fe-on-Cu system is, that on the one hand - in agreement with photoelectron diffraction and CO titration experiments [12] - the LEIS data show unambiguously that Fe atoms exist to about equal amounts in the surface and the first subsurface layer at all coverages. On the other hand recent STM investigations [18,30] only report on islands of monolayer height. The easiest solution of this puzzle is to assume that only about 50% of the deposited iron atoms adsorb on top of the original substrate surface while the remaining Fe atoms are incorporated into the original copper surface layer and are covered by Fe or Cu atoms. The STM results also rule out the above mentioned Cu decoration of the Fe islands as the only cause of the absence of Fe surface defects at 0.1
ML, since in this case islands with bilayer height should be observed by STM as well. The scenario of partial iron incorporation suggested above yields islands of monolayer height only and simultaneously explains the observation of Fe atoms in both the first and second layer. As the shadow cones for Fe and Cu are undistinguishable within our experimental resolution, we are not able to decide whether the Fe atoms appearing in the first subsurface layer are covered by Fe or by Cu atoms, i.e., whether a kind of Fe-Cu surface alloy is formed or whether the Fe atoms reside in (almost) pure iron islands. While a capping by copper atoms, which have diffused onto the incorporated iron atoms, can explain our results for low Fe coverages, it certainly cannot be the exclusive effect at nominal coverages above 1 ML, as in this case more than 50% of the outermost layer consists of iron. The simplest explanation, however, compatible with all data obtained so far, is that iron grows in partly-incorporated doublelayer islands, accompanied by the diffusion of about 0.1 ML of copper onto the deposited (and incorporated) iron films. The mechanism by which the Fe atoms replace the Cu atoms in the surface layer may be similar to that suggested by Feibelman [33] to explain the low barrier to self diffusion on Al(OO1). In this theoretical work it was found, that for the diffusion of an (aluminum) adatom, located in a four-fold hollow position, it is energetically more favourable to move one of the four underlying substrate atoms out of the surface plane onto a neighbouring four-fold hollow site on top of the surface and to fill the created vacancy by the adatom, instead of hopping of the adatom itself to a neighbouring fourfold hollow site via an intermediate two-fold bridge position. This concerted displacement mechanism involves a replacement of substrate atoms by adatoms, just as observed for the Fe-onCu system. The expelled Cu atoms may then diffuse across the surface and form copper islands as suggested in recent STM studies [18,301. 5. Conclusion In summary, our measurements show unambiguously that initially iron does not grow in a
Th. Detzel et al. /Growth of ultrathin iron films on Cu(OO1)
layer-by-layer fashion at room temperature. In particular it was demonstrated that one monolayer of iron does not lie down on the Cu(OO1) substrate reflecting the tendency for island formation, which is anticipated by thermodynamic considerations. The LEIS measurements show that for all coverages (even as low as 0.1 ML) about equal amounts of Fe atoms are visible in the surface and the first subsurface layer, respectively. At smallest coverages (- 0.1 ML) all iron atoms or islands are surrounded by copper atoms. Simultaneously copper defects appear on the surface. The iron films exhibit maximum roughness around 1 ML and smoothen with increasing film thickness. At coverages exceeding 2 nominal ML iron is found to wet the substrate for the most part. For the 6 ML film the morphological and crystallographic order closely resembles that of the clean Cu(OO1) surface. At higher coverages the roughness increases once more and above _ 10 ML the pseudomorphic growth of fee iron breaks down and the stable bee structure of iron, with the surface normal along (llO), is formed. Taking together the present LEIS results and previous STM work, we suggest a distorted bilayer growth with about one half of the iron incorporated in the substrate for coverages up to 2 ML. This automatically requires transport of Cu atoms, as was observed experimentally. Therefore Cu diffusion cannot be neglected for the growth at and especially above room temperature. Finally we would like to mention that the present work shows that LEIS is a versatile tool to study growth and diffusion of ultrathin films. It offers a complementary approach to modern surface microscopy techniques because of its element resolving abilities, but also because the probing depth is not only restricted to the topmost surface layer.
Acknowledgements
The authors would like to thank Staib Instruments and in particular B. Senftinger for placing generously at our disposal a commercial computer controlled video system. We are grateful to
237
V. Dose for encouragement and support and to E. Taglauer and E. Bertel for helpful discussions during the preparation of this publication. N.M. wants to thank H. Diirr for the introduction to the ion-scattering technique.
References [l] J.G. Gay and R. Richter, J. Appl. Phys. 61 (1987) 3362. [2] D. Pescia, M. Stampanoni, G.L. Bona, A. Vaterlaus, R.F. Willis and F. Meier, Phys. Rev. Lett. 58 (1987) 2126. [3] M. Stampanoni, Appl. Phys. A 49 (1989) 449. [4] C. Liu, E.R. Moog and S.D. Bader, Phys. Rev. Lett. 60 (1988) 2422. [S] D.P. Pappas, K.P. Kamper and H. Hopster, Phys. Rev. Lett. 64 (1990) 3179. [6] R. Allensbach and A. Bischof, Phys. Rev. Lett. 69 (1992) 3385. [7] J. Thomassen, F. May, B. Feldmann, M. Wuttig and H. Ibach, Phys. Rev. Lett. 69 (1992) 3831. [8] M. Onellion, M.A. Thompson, J.L. Erskine, C.B. Duke and A. Paton, Surf. Sci. 179 (1987) 219. [9] A. Clarke, P.J. Rous, M. Arnott, G. Jennings and R.F. Willis, Surf. Sci. 192 (1987) L843. [lo] Y. Darici, J. Marcano, H. Min and P.A. Montano, Surf. Sci. 217 (1989) 521. [li] E. Bauer, Z. Krist. 110 (1958) 372, 395. [12] D.A. Steigenvald, I. Jacob and W.F. Egelhoff, Jr., Surf. Sci. 202 (1988) 472. [13] A. Zangwill, Physics at Surfaces (Cambridge University Press, Cambridge, 1988). [14] R. Germar, W. Diirr. J.W. Krewer, D. Pescia and W. Gudat, Appl. Phys. A 47 (1988) 393. [15] H. Glatzel, Th. Fauster, B.M.U. Scherzer and V. Dose, Surf. Sci. 254 (1991) 58. [16] J. Thomassen, B. Feldmann and M. Wuttig, Surf. Sci. 264 (1992) 406. [17] M. Arnott, E.M. McCash and W. Allison, Surf. Sci. 269/270 (1992) 724. 1181D.D. Chambliss, R.J. Wilson and S. Chiang, J. Vat. Technol. A 10 (1992) 1993. 1191M.F. Onellion, C.L. Fu, M.A. Thompson, J.L. Erskine and A.J. Freeman, Phys. Rev. B 33 (1986) 7322. [20] W.A.A. Macedo and W. Keune, Phys. Rev. Lett. 61 (1988) 475. [21] P.A. Montano, G.W. Fernando, B.R. Cooper, E.R. Moog, H.M. Naik, S.D. Bader, Y.C. Lee, Y.N. Darici, H. Min and J. Marcano, Phys. Rev. Lett. 59 (1987) 1041. [22] E. Taglauer, in: Ion Spectroscopies for Surface Analysis, Eds. A.W. Czanderna and D.M. Hercules (Plenum, New York, 1991) p. 363. [23] Th. Fauster, Vacuum 38 (1988) 129. 1241H. Diirr, R. Schneider and Th. Fauster, Phys. Rev. B 43 (1991) 1802.
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of ultrathin iron films on CdOOl)
[25] H. Niehus and R. Spitzl, Surf. Interface Anal. 17 (1991)
[26] [27] [28] [29] [30]
287. Yi-sha Ku and S.H. Overbury, Surf. Sci. 273 (19921 341. C. Argile and G.E. Rhead, Surf. Sci. Rep. 10 (1989) 277. M. Aono, C. Oshima, S. Zaima, S. Otani and Y. Ishizawa, Jpn. J. Appl. Phys. 20 (1981) L.829. N. Memmel and Th. Detzel, to be published. A. Brodde and H. Neddermeyer, Surf. Sci. 287/288 (1993) 988.
[31] H. Landskron, G. Schmidt, K. Heinz, K. Miiller, C. Stuhlmann, U. Beckers, M. Wuttig and H. Ibach, Surf. Sci. 256 (1991) 115. [32] H. Magnan, D. Chandesris, B. Vilette, 0. Heckmann and J. Lecante, Phys. Rev. Lett. 67 (1991) 859. [33] P.J. Feibelman, Phys. Rev. Lett. 65 (1990) 729.
Growth of ultrathin iron films on Cu( 001): an ion-scattering spectroscopy study Th, Detzei, N. Memmel and Th. Fauster
Received 12 February 1993; accepted for pubIication 26 April 1993
Low-energy ion scattering has been applied to study the pseudomorphic growth of ultrathin fee iron films on a Cu(OO1) single-crystal surface. At all coverages - even as low as 0.1 monolayer (ML) - iron is found to about equal amounts in both the surface and the first subsurface layer, since part of the iron atoms are incorporated into the original copper surface. This is particular evident for the 0.1 ML film, which does not exhibit any iron defects, such as steps or adatoms. Instead copper surface defects become visible upon iron deposition, Near 2 ML the substrate is covered for the most part by a relatively smooth b&et indicating the coalescence of the iron islands. Up to a film thickness of around 6 ML the surface defect density of the iron layers decreases with increasing coverage and raises again at larger coverages. Above N 10 ML the pseudomorphic fee growth breaks down aad domains of bee iron with (110) or~entatioo are formed.
The structure and morphology of ultrathin magnetic films has been of growing interest within the last 15 years for both fundamental and technological reasons. These systems serve as an outstanding opportunity to study basic concepts of theoretical physics like two-dimensional phase transitions and can be used to investigate the scaling concepts for the~ody~ami~ quantities. Moreover, thin magnetic fiIms with the easy axis of magnetization normal to the surface are of prominent technological importance in opto-magnetic storage devices. In the monolayer range of epitaxialiy grown fee iron films, such a perpendicular spin orientation was predicted [I] and experimentally observed [2-71. Since the lattice constants of the face centered cubic phase of iron (y-Fe), extrapolated to room temperature, and fee copper differ less than I%, the system Fe on Cu(OO1) seems to offer unique possibilities to produce ultrathin films of pseudomo~hic fee iron. In recent years, several studies by low-energy electron dif~action ~39~6OZS/93/$06.~
fLEEID), in which the overlayer structure is shown to be nearly identical to a ~ntinuation of the fee lattice of the substrate, corroborated this supposition [8-lo]. According to thermodynamic cansiderations one would expect the initial growth of Fe on Cu(OO1) not to proceed via a perfect layer-bylayer manner, Bauer [ill pointed out that the surface free energies of the substrate and the adsorbate together with the interfacial free energy are in a first appro~mation the thermodynamic quantities dete~ining the growth mode. The the~~~arni~ parameters of the system under desideration fulfill the requirement for island formation [12], i.e., the sum of the surface free energy of Fe and the interfacial free energy between Fe and Cu exceeds the surface free energy of the Cu substrate. Furthermore, these energetic circumstances may lead to the diffusion of Cu substrate atoms onto the surface of the adsorbed iayers in order to minimize the total free energy. Thus, we are dealing with the general problem in which a high surface free-energy metal is attempted to be epitaxed on a low surface free-energy metal. Viewing the chart of the
0 1993 - Elsevier Science Publishers B.V. AI1 rights reserved
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of ultrathin iron films on Cu(OOl)
measured elemental surface free energies [131 it becomes evident, that this is the rather usual case concerning epitaxy in magnetic ultrathin film research and application. Therefore, the combination Fe/Cu might be considered as a prototype system. Nevertheless, it is worth mentioning that the difference of the two sides in the free-energy balance is not extremely significant. Since the Bauer criterion concerns only thermal equilibrium and because the growth of thin metal films by vapour deposition is a dynamic process usually being far from true equilibrium, microscopic mechanisms might strongly influence the film growth. In addition, any strain energy due to the accommodation of the lattice mismatch between the two materials is neglected in the Bauer criterion. A large amount of work has been dedicated to study the growth mode of ultrathin Fe films on Cu(OO1) substrates in recent years. While there is general agreement now, that iron grows layer-bylayer for films with a thickness between 5 and N 10 ML, the experimental results and their interpretations concerning the initial stage of the growth up to 5 ML are still quite controversial. Germar et al. [14] claimed from their Auger electron spectroscopy @ES) investigations that the growth mode is essentially layer-by-layer, the deviation from perfect monolayer growth being less than 10%. At the same time Steigerwald et al. [12] found the film growth not to proceed via a Frank-van der Merwe mechanism, but discovered Fe agglomeration and Cu surface segregation using XPS forward scattering and the carbon monoxide titration technique. Glatzel et al. [151 report on the formation of bilayers for the first 6 ML. The authors observed the characteristic breakpoints in the Auger uptake curve recorded during evaporation and used Rutherford backscattering spectroscopy (RBS) to determine the nominal film thickness, clearly revealing that the breakpoints occur at the completion of bilayers. Thomassen et al. [16] investigated growth, structure and morphology of ultrathin iron films utilizing low- and medium-energy electron diffraction (LEED, MEED) as well as AES. This group reports on Fe agglomeration up to about 5 ML followed by layer-by-layer growth for coverages
between 5 and 14 ML. Finally, at around 14 ML due to the formation of misfit dislocations the breakdown of epitaxy was observed. The preceding findings tally more or less with results presented by Arnott et al. who applied specular helium atom scattering techniques [17]. In a recent scanning tunnelling microscopy (STM) study an even more complex growth behaviour is stated [HI. Submonolayer iron deposition (0 = 0.1 ML) yields the f?rmation of small features of typical size N 10 A, interpreted as iron patches in the original copper surface layer where iron atoms have replaced copper atoms, and larger monolayer islands, formed by the “expelled” Cu atoms. The experimental results concerning the magnetic properties of thin iron films on Cu substrates seem to be as contradictory as those obtained for the film morphologies. Different groups, for instance, have reported the presence of ferromagnetism on fee iron grown on Cu(OO1) 12,191,whereas the absence of ferromagnetism at room temperature has been stated by others 120,211. In the latter experiments, however, iron was deposited at a substrate temperature of 200°C thus leading to a pronounced copper diffusion to the top of the Fe film [12], which probably is responsible for the different magnetic behaviour observed for these films. This already shows, that a precise knowledge of the epitaxial growth process is required to understand the magnetism of ultrathin films. In this paper we report on a first study of the growth mechanisms of ultrathin Fe films on Cu(OO1) using low-energy ion scattering (LEIS). Special emphasis is put on the low coverage regime. In recent years, the use of ion-scattering methods has been established as an advantageous approach to study surface, interface and thin film phenomena [22-261. The reason for this is manifold: Like electron spectroscopy methods involving core electrons, ion-scattering spectroscopy is element specific as well as quantitative and therefore can be used for a chemical analysis of the components to be found at the surface. Because of its enormous surface sensitivity caused by neutralization and blocking effects, it is possible to probe exclusively the outermost surface layers of the sample. Moreover, due to shadowing effects
Th. Detzel et al. /Growth
of ultrathin iron films on Cu(OO1)
this method provides structural information with monatomic resolution in the surface plane as well as normal to the surface. In contrast to Fourierspace techniques like electron and atom diffraction LEIS supplies mass-selective real-space information on the local atomic arrangement at the surface including surface defects. As a further advantage it is noteworthy that the evaluation schemes are straight forward without requiring sophisticated and time-consuming model calculations.
2. Experiment The experiments were performed in an UHV system with a base pressure below 3 x lop9 Pa. A (OOl)-oriented Cu crystal (miscut < rtO.3”) was prepared prior to each iron deposition by successive cycles of Ar+ bombardment (600 eV) and annealing to 900 K until the contamination levels were below the detection limits of LEIS. After the last annealing the sample was cooled down slowly (< 2 K/min) in order to reduce the number of possible surface defects. The crystallographic structure of the surface was monitored by LEED, showing a well-defined (OOl)-LEED pattern with sharp spots and low background intensity, as well as by impact collision ion-scattering spectroscopy (ICISS) [22-261, revealing a well ordered copper surface. Subsequently, iron was evaporated onto the Cu(OO1) crystal, held at 298 + 2 K, from a small iron pin of purity 99.999% heated by electron bombardment. The vapour source was surrounded by a water-cooled copper tubing in order to remove most of the radiant heat, thus keeping the levels of contaminating gases small. During evaporation the pressure stayed below 1 X 10-s Pa. An analysis of the residual gas by a quadrupole mass spectrometer revealed that the slight increase in pressure was mainly due to an increase in the partial pressure of hydrogen. The coverage was calibrated independently by AES and MEED and is given in units of a nominal monolayer (ML), i.e. 1.53 x 1015 atoms/cm*. The MEED experiment was carried out using a 4 keV electron beam at grazing incidence (N 100).
229
The diffraction pattern was displayed on the screen of a reverse view LEED optics and the intensity variations of several reflexes during Fe deposition were recorded simultaneously by a computer controlled video system. The characteristic intensity variations reported by Thomassen et al. [16] were reproduced and the time necessary to deposit one monolayer of Fe was inferred from the period of the regular oscillations observed in the coverage range between 5 and 10 ML. The evaporation rates ranged from 0.5 to 1 ML/min in the different experiments. In an alternative calibration procedure the intensity ratio of the 705 eV iron and 920 eV copper Auger transitions was determined and compared to the values given by Glatzel et al. [151, who used Rutherford backscattering to measure the nominal iron coverage corresponding to distinct AES intensity ratios. Both calibration procedures agree well with each other. However, care has to be taken, as different experimental setups were used in the AES experiments of Glatzel et al. and our work. (Glatzel: Excitation by a 3 keV electron beam at 10” grazing incidence. Detection by a retarding field analyzer with a mean detection angle of 61”. This work: Excitation by a 5 keV electron beam at 10” grazing incidence. Detection by a retarding field analyzer with a mean detection angle of 48”.) From the MEED and AES experiments we specify the coverages given in this publication with an experimental error of about $- 10% but less than +0.3 ML. As will be discussed below, the LEIS experiments themselves also give a lower limit of the amount of Fe deposited onto the surface. All ion-scattering experiments shown in this paper were carried out using a monoenergetic beam of Ne+ ions with a kinetic energy of 5 keV and an electrostatic analyzer at an scattering angle of 160”. In order to avoid considerable damage of the surface by the impinging ions the primary intensity as well as the recording time were kept low, yielding a maximum fluence of 5 X 101* to 5 X 1013 ions/cm*, depending on the kind of spectrum to be taken. No measurable radiation damage was observed under these conditions. Details about the experimental setup are described elsewhere [23].
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3, Results
To study the growth mode of iron on Cu@Ol) at room temperature ion-scattering experiments were performed for various film thicknesses ranging from 0.1 to 14 ML. However, our main interest concentrated on the low coverage regime. We were especially concerned with the possibility of island formation at coverages up to 1 nominal ML and the phenomenon of copper diffusion during deposition. Fig. 1 shows the intensity of backscattered Ne + ions as a unction of their kinetic energy in the energy range predicted by the binary collision model [23]. The film thickness amounted to 1.3 ML and the angle of incidence 2Fr was chosen equal to 40”, where - as will be explained below - only the outermost surface layer is probed. For perfect monolayer growth one would expect the Cu intensity originating from the first layer to vanish at a nominal coverage exceeding I ML. Obviously this is not the case. Instead, the copper signal (at N 1330 eV> is only slightly less than the iron yield (at _ 1100 eVI, even for a coverage somewhat larger than I ML. At this point it should be mentioned that due to the difference in nuclear charge as we11 as due to conceivable deviations in the neutralization probabilities of
backscattered
ion energy [eV]
Fig. 1. LEE spectrum of 1.3 ML Fe/CuRMU), recorded for a 5 keV Ne’ beam incident at an angle of 40” relative to the surface plane along the [lo01 azimuth. The scattering angle was 160*. The peaks at 1100 and 1330 eV result from scattering by iron and copper atoms, respectively.
0
I
1
I
I
I
2
3
4
5
film thickness
6
[ML]
Fig. 2. Normalized intensity of Ne+ ions backscattered from Fe atoms in the outermost layer as a function of coverage. Dotted, solid and dash-dotted tines: signals expected for monohi- and trilayer growth, respective&.
iron and copper the “scattering cross sections” differ to a small extend. This difference was determined experimentally using pure elemental standards. We found a 10% lower “cross section” for iron compared to copper. This difference is small and has essentially no influence on the qualitative and quantitative interpretations drawn below. Therefore, we infer already at this point that one nominal monolayer of Fe is not spread out homogeneously on the Cu substrate, indicating a non-wetting tendency. Similarly to the well-known Auger uptake curves [Zi’] the Fe/(Fe + Cu) LEIS ratio is depicted as a function of nominal film thickness in fig. 2. The signal increases almost linearly with iron coverage up to about 2 ML where a distinct breakpoint in the slope is observed. Finally, beyond 4 ML the residual copper signal amounts to merely a few percent. This contribution may be attributed to tiny non-covered substrate areas as well as to Cu atoms which have diffused onto the film surface. Near a Fe coverage of 2 ML the Cu intensity disappeared for the most part, but there is stiI1 a non-negligible portion of copper visible in the LEIS measurements. Even so, a three-dimensional island formation (Volmer-Weber growth mode) can be excluded since in this case one would expect the adsorbate to leave large parts of the substrate uncovered. As mentioned earlier, Glatzel et al. claim the growth of Fe to proceed in a bilayer fashion.
Th. Detzel et al. /Growth
ofultrathin iron films
Neglecting Cu diffusion and interm~ng this growth mode would lead to an almost linearly increasing LEIS signal ratio reaching its saturation value at a nominal film thickness of 2 ML, as shown by the solid line in fig. 2. The line is not perfectly straight due to the 10% difference in “scattering cross section” between iron and copper. The dotted and dash-dotted lines show the intensity ratios expected for monolayer or trilayer growth, respectively. The experimental data are found between the two model curves for bilayer and trilayer growth, deviating less from the bilayer case. The assumption of u 0.1 ML of copper diffusing onto the film surface during deposition would explain the difference between the measured values and the ideal bilayer uptake curve. Thus, from this first LEIS experiments a distorted bilayer growth mode may be infered for the first two monolayers. In order to gain more insight into the microscopic growth processes, Fe films of different thicknesses were investigated by impact collision ion-scattering spectroscopy WISS). This method works at scattering angles near 180” and is based on the shadow cone concept [ZS]: The primary ion flux is deflected by the target atoms in such a way that there is no flux immediately behind them ((‘shadow cone”, see fig. 3). The ion trajectories pushed out of this shadow cone are focused to the edge of the shadow cone, yielding an increased ion flux in this region. At larger distances from the cone axis the flux decreases to its average value. The experimentai procedure is as follows: the electrostatic energy analyzer is tuned to the residual energy of the ions backscattered by the chemical species to be probed (i.e., N 1100 eV for Fe, N 1330 eV for Cu, see fig. 1). Then the backscattered intensity is measured as a function of the angle of incidence tv. Near grazing incidence (?P = O? on an ideal crystalline surface each target atom is hidden in the shadow cone of its nearest neighbour in the surface plane. No head-on collisions, as required for 180” backscattering can occur and the backscattered intensity vanishes. If the angle of incidence is increased, the atoms of the topmost layer will eventually move out of the shadow cone of the adjacent atoms (see fig. 3a), leading to a steep increase of
on Cu(oO1~
231
Fig. 3. The shadow-cone concept of impact-collision ionscattering spectroscopy (ICISS). Only the trajectories of the incoming beam are shown in this figure. (a) Situation where atoms in the outermost layer just “leave” the shadow cone of their nearest neighbours in the surface plane. (b) Atoms in the first subsurface layer “leave” the shadow cone caused by the surface atoms.
the backscattered intensity at the critical angle !P,, (the indices indicate shadowing from a first layer atom onto a first layer atom). This is nicely seen in fig. 4a, where an ICISS spectrum of 5 keV Ne+ ions backscattered from the Cu atoms of the clean Cu@Ol) surface is displayed. The ion beam was impinging along the [loo] azimuth of the sample. The critical angle here amounts to W,, = 13”. Further increase of the angle of incidence P yields a decreasing backscattered intensity as a consequence of decreasing flux away from the shadow cone edge. Upon further increase of W the neighbours in the second Iayer are also Ieaving the cone (fig. 3b) and a second enhancement . . . of the ion yield 1s observed at ?Pylz= 63” (see fig. 4a). Simultaneously Nef ions are focused from atoms in the second layer onto third layer atoms. Ions backscattered from the third layer, however, do not reach the detector since they are blocked on the outgoing path by the substrate atoms directly above in the first layer. As can be seen, the signal originating from the second layer is approximately reduced by a factor of 4 relative to the first layer signa at lyii = 13”, mainly due to the higher neutralization rate for Ne+ ions backscattered from the second layer. It should be mentioned that this rather simple concept, involv-
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ing only a single large scattering event (plus a small angle deflection due to the shadowing atom) works only well for scattering angles near Ho”, since only in this case multiple scattering events, which would complicate the interpretation of the data, can be neglected. Fig. 5 shows a selection of ICISS spectra taken at various iron film thicknesses. The energy analyzer was set to an energy of 1100 eV, so only Ne+ ions scattered from Fe were detected. From these spectra the following statements can be made: (1) The gross features of all spectra up to N 10 ML are the same as on the clean Cu(OO1) surface (fig. 4a). This confirms that the iron films grow indeed in an fee structure with similar interatomic distances as the copper substrate. Since the atomic numbers of Fe and Cu are similar, the shadow cones caused by Cu and Fe atoms, re-
I
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S keV Ne+*
I
I
Fe/Cu(OOl)
IRON signal
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0.2 ML 7 I
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I
I
5 keV Ne+ -W Fe/Cu(OOl) COPPER
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signal 1
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20 40 60 angle of incidence
80
1o1
Fig. 5. KISS spectra for various Fe films on Cu(OO1). The Nef ions are backscattered from iron. The scattering plane was along the [lo01 azimuth. All spectra are normalized to same height at maximum intensity.
0.7 ML
0.2 ML
1.1 ML
1, , , , 20 40 60 angle of incidence
4
:Iean
80 [“I
Fig. 4. KISS spectra for (a) the clean Cu(OO1) surface and (b)-(e) the surface covered by various amounts of iron. The Nef ions are backscattered from copper. The scattering plane was along the [lOO] azimuth. All spectra are normalized to same height at maximum intensity.
spectively, are practically identical. The associated critical angles differ less than 0.5”, which is about our experimental uncertainty in the determination of the critical angle. The 14 ML spectrum is qualitatively different. Both critical angles are shifted, with the first layer focusing peak at *ii broadened considerably towards higher angles and the second layer feature at 1yi2 split into a double peak structure. This indicates the breakdown of pseudomorphic growth and the formation of bee Fe with a (110) surface normal [7]. As the unit mesh of the bee (110) surface is of two-fold symmetry only, four domains will exist on an fee (001) surface, but only two of them are inequivalent along the [lOOI-azimuthal direction of the Cu(OO1) substrate, where the measure-
Th. Lletzelet al./ Growth of ultrathinironfilms on CufOOl)
/
I.01
MI.84
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0.83
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L? $ 0.4-
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. 0.2 2’ 0.06
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f
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s 1 5
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Fig. 6. Normalized intensity of Net ions scattered from Fe atoms in the first subsurface layer as a function of coverage.
ments were performed. This two inequivalent domains produce the observed broadening and splitting of the ICISS structures [29]. Iron films with an fee or bee (OOl)-surface cannot explain these split or broadened features. (2) Both, the intensity increase at PI1 = 13” due to focusing onto top-layer Fe atoms as well as the intensi~ increase at ‘I’ = 63” due to focusing onto subsurface Fe atoms are observed for all coverages. Note that this even holds for coverages as small as 0.1 ML. This clearly demonstrates that even at such low coverages Fe atoms are present in both the top layer and the first subsurface layer. The ratio of the intensity increases at both critical angles is independent of coverage and amounts to - l/4, which indicates that about equal amounts of iron atoms exist in the surface and the first subsurface layer. (3) This statement is further substantiated in fig. 6, where, as a function of film thickness, the Fe/lFe + Cu> ratio of the intensity increase at P,, is plotted. This ratio is essentially proportional to the percentage of Fe atoms in the first subsurface layer. The iron content in this layer increases almost linearly with coverage up to about 80% at a nominal coverage of 2 ML. For higher coverages it converges slowly towards 100%. (4) The LEIS data presented in figs. 2 and 6 also confirm the coverage calibration established by MEED and AES. At a nominal coverage of 2 ML LEIS shows that 90% of the surface atoms
233
and about 80% of the subsurface atoms are iron. Therefore at least 1.7 ML of iron have been deposited on the surface at this point. As iron atoms in the third and deeper layers do not contribute to the backscattered signal in our detection geometry, this is a lower limit of the amount of iron present on the surface. Fe atoms in the third and deeper layers may either be due to Fe agglomeration into islands of more than bilayer height or due to Cu atoms which have diffused onto the surface during the growth process. (5) Further info~ation regarding the microscopic growth process can be infered from the backscattered intensity near grazing incidence. As mentioned in the beginning, on a perfect surface no ions are backscattered for incidence angles below the first critical angle P,,. Upon deposition of iron, however, considerable intensity is observed at small angles of incidence (V < lo”, shaded regions in fig. 5). Adatoms, steps or vacancies have to be made responsible for this observation. Common to all of them is, that the back~attering atom has no neighbour along the nOO] direction. As they cannot be distinguished in the experimental spectra, we will summarize the foregoing terms in the following paragraphs by the collective expression “surface defects”. In fig. 7 the backscattered ion yield at W = 5” normalized to the ion yield at 20” (taken from fig. 5) is depicted as a function of film thickness to provide a means for the concentration of Fe surface defects and the quality of the iron films. This value increases strongly up to about 1 ML, thus revealing the poor quality of the Fe films near this coverage. This is in agreement with the pronounced intensity minimum observed in MEED [16] and He-atom scattering 2171. Near 2 nominal ML the defect intensity is reduced remarkably, pointing to a rather smooth film surface, which therefore can be interpreted as an almost completed bilayer. The 6 ML film shows a defect intensity comparable to the clean substrate. Therefore, we deal with an almost perfectly ordered Fe-film. Elevated coverages exhibit again a rougher surface morpholo~ due to the breakdown of pseudomo~hic growth and the formation of bee iron above - 10 ML.
Th. Detzel et al. /Growth
234
of ultrathin iron f&s on CufoOll 0.1 ML Fe deposited at 140 K
‘Fe surface defects’
I
3
1
’
1
’
1
’
’
2
4
6
8
10
12
14
film thickness
[ML]
Fig. ‘7. Iron surface defect ~ncen~ation versus coverage as determined from the ratio of the LEIS intensity of Nef ions backscattered from iron at incidence angles of 5” and 20”, respectively. For comparison the copper defect intensity observed on the clean copper surface is indicated by the dashed line.
(6) The ICISS spectrum of the 0.1 ML film (fig. Sa> is pa~icularly interesting, as there is essentially no intensity observed at low angles of incidence. Note, that by the normalization procedure applied in fig. 5, the trivial coverage dependence of the signal has already been removed. No Fe surface defects are visible at this coverage. This only seems possible if iron forms large perfectly ordered islands or if the iron (rich) regions are completely surrounded by copper atoms, either due to incorporation of the Fe atoms into the substrate surface or due to decoration of the iron (rich) islands by copper atoms. The first explanation must be ruled out, because an only slightly higher amount of evaporated Fe yields already considerable defect densities (fig. 5b). (7) Further evidence for iron incorporation or decoration of the iron (rich) islands comes from the ICISS spectra in figs. 4a-4e. Here the electrostatic analyzer was tuned to an energy of 1330 eV, corresponding to collision of the Nef ions with copper atoms. The attention is immediately attracted to the elevated Cu defect intensity at smaI1 angles of incidence appearing upon iron deposition (shaded regions in fig. 4). This unambiguously proves that the copper surface layer
angle of incidence [“I Fig. 8. Iron-ICISS spectra for 0.1 ML Fe/Cu(OOl) deposited at 140 K. (a) As-deposited. (b) Annealed to room temperature for 10 min.
does not remain unaffected by adsorption of iron. Both scenarios, the formation of additional Cu islands by the “expelled” Cu atoms or decoration of Fe islands by Cu atoms, can explain the observed Cu defect intensity. In any case the data clearly demonstrate, that even at room temperature significant diffusion of Cu atoms, involving considerable mass transport, occurs. 3.2. Temperatureeffects To study the influence of the substrate temperature during Fe deposition on the growth mode, we performed additional evaporation experiments at cold and at slightly heated samples. Fig. 8a shows an iron ICISS spectrum recorded on a 0.1 ML film deposited at 140 K. Here the situation is changed dramatically as compared to room temperature experiments. An enormous defect density is observed since the small angle intensity exceeds even the peak of regular first layer terrace sites. Subsurface iron in fee sites is also detected. This spectrum drastically demonstrates the reduced mobility of the iron atoms at a temperature of 140 K. The ~~esponding copper spectrum (not shown), however, exhibits an
Th. Detzel et al. /Growth of ultrathin iron films on Cu(OO1)
elevated intensity at small angles of incidence, similar to fig. 4b. Therefore - even at a temperature of 140 K - iron induces a rearrangement of the Cu surface atoms. Annealing to room temperature improves the quality of the iron films as can be seen by the strong reduction of the iron defect peak (fig. 8b). However, it seems remarkable that iron surface defects are still visible, in contrast to adsorption at room temperature (fig. 5a). Simultaneously with the decrease of the surface defect peak an increase of the first layer peak near ?Prr is observed upon annealing, thus showing the “conversion” of Fe atoms with no nearest neighbour along [iOO] into atoms with a neighbour along this direction. The second layer contribution is only slightly affected by the annealing procedure. Keeping in mind the desire for a single fee iron monolayer lying flat on the substrate we deposited just under 1 ML onto a cold sample. We found again a rather rough film surface with many atoms buried in the second layer. Annealing the samples yields a spectrum comparable to fig. 5c, with a decreased low-angle intensity (i.e., a reduced number of iron surface defects) and a mostly unchanged second layer feature. Therefore it is not possible to manufacture an iron monolayer, using this preparation procedure. A similar experiment was carried out for a 3 ML film. Again many iron surface defects were visible for the film deposited at low temperature. Annealing to room temperature improved the quality of the film to a level comparable to that of a room-temperature grown 3 ML film. In a further experiment we deposited about 1 ML of iron at a temperature slightly above room temperature (340 K). We observed ICISS spectra with diminished defect intensity but with an excessive intensity raise at ?P’,, due to an enhanced iron percentage in the first subsurface layer. We estimate the number of iron atoms in the subsurface layer to exceed the top layer concentration by about 30%, which points without any doubt to enhanced diffusion of Cu atoms onto the iron film. The diminished defect density indicates a smoothing effect of the segregated substrate atoms, which was also observed in He atom scattering experiments [ 171.
235
4. Discussion Our experimental findings on the growth of ultrathin iron films on Cu(OO1) single crystal surfaces are in good agreement with a previous study by photoelectron diffraction [121, which reports on bilayer island formation in the low-coverage regime as well as on the appearance of Cu segregation already at room temperature. The results published by Germar et al. [14], who claim the growth to proceed via a undisturbed Frank-van der Merwe mechanism are unambiguously disproved because always about 50% of the iron atoms are buried in the subsurface. As described in the introduction, Glatzel et al. [15] assume the formation of perfect bilayers in order to explain their Auger and RBS data. This interpretation is compatible with our data at least for the first two monolayers, even though at 2 ML the film is not perfectly ordered and its formation is accompanied by the possibility of some Cu segregation and intermixing. Thomassen et al. [161 suggest for the Fe-on-Cu system iron agglomeration up to a thickness of 5 ML. Our real-space ion-scattering data, however, indicate that the iron islands coalesce already at about 2 ML, leading to an almost pure (> 90%) iron surface beyond this coverage. Furthermore, the order in the films between 2 and N 9 ML deviates only slightly from that of the clean Cu(OO1) surface. Both findings, coalescence of the iron islands around 2 ML and the formation of rather smooth films above this coverage (up to the breakdown of epitaxial growth) is also reported in the most recent STM study of Brodde and Neddermeyer [30]. Therefore we are led to conclude that the strong intensity variations of various MEED spots between 2 and 5 ML - as reported in ref. [16] and reproduced in the present work - do not reflect changes in the morphology of the Fe films but are related to structural rearrangements of the surface as indicated by the (4 x l), (5 x 1) and (2 x 1) superstructures observed by LEED for coverages of about 2, 4 and 6 ML [161. These structural changes, however, involve only small displacements ( - 0.1 A> of the Fe atoms from their ideal (1 x 1) positions in such a way, that (to first order) the mean
236
7k. Detzel et al. /Growth
of ultrathin iron films on Cu(OOl)
nearest-neighbour distances remain unchanged [31,32]. Therefore these structural changes do not change the critical angles in the ICISS spectra, but only lead to a slight (- 1”) broadening effect. At smallest coverages we state a qualitative agreement of our LEIS data to the STM investigations of Chambliss et al. 1181.In particular, they report on Fe incorporation and the emergence of Cu islands on the surface, in agreement with the conclusions based on our LEIS results. Quantitatively, however, some discrepancies exist. According to the interpretation of Chambliss et al. the copper islands are formed by the Cu atoms expelled from the original surface layer by the incorporated Fe atoms. By comparison of the Cu surface defect ICISS intensity measured for the 0.1 ML iron film with that of a miscut Cu crystal [231, we estimate, that about 0.05 ML of copper is visible as surface defects (i.e., as atoms at the steps of islands and terraces). If the expelled Cu atoms agglomerate in rather large islands this means, that - contrary to the suggestion of Chambliss et al. - the Fe-Cu replacement does not simply occur on a one-to-one basis, but that additional Cu atoms are coming from steps, as suggested in the STM work of Brodde and Neddermeyer [30] from the simultaneously observed step roughening. An alternative explanation would be, that the atoms at the perimeter of both the large monolayer islands and the small patches, seen in the STM work, are copper atoms. ’ One of the remaining puzzles of the Fe-on-Cu system is, that on the one hand - in agreement with photoelectron diffraction and CO titration experiments [12] - the LEIS data show unambiguously that Fe atoms exist to about equal amounts in the surface and the first subsurface layer at all coverages. On the other hand recent STM investigations [18,30] only report on islands of monolayer height. The easiest solution of this puzzle is to assume that only about 50% of the deposited iron atoms adsorb on top of the original substrate surface while the remaining Fe atoms are incorporated into the original copper surface layer and are covered by Fe or Cu atoms. The STM results also rule out the above mentioned Cu decoration of the Fe islands as the only cause of the absence of Fe surface defects at 0.1
ML, since in this case islands with bilayer height should be observed by STM as well. The scenario of partial iron incorporation suggested above yields islands of monolayer height only and simultaneously explains the observation of Fe atoms in both the first and second layer. As the shadow cones for Fe and Cu are undistinguishable within our experimental resolution, we are not able to decide whether the Fe atoms appearing in the first subsurface layer are covered by Fe or by Cu atoms, i.e., whether a kind of Fe-Cu surface alloy is formed or whether the Fe atoms reside in (almost) pure iron islands. While a capping by copper atoms, which have diffused onto the incorporated iron atoms, can explain our results for low Fe coverages, it certainly cannot be the exclusive effect at nominal coverages above 1 ML, as in this case more than 50% of the outermost layer consists of iron. The simplest explanation, however, compatible with all data obtained so far, is that iron grows in partly-incorporated doublelayer islands, accompanied by the diffusion of about 0.1 ML of copper onto the deposited (and incorporated) iron films. The mechanism by which the Fe atoms replace the Cu atoms in the surface layer may be similar to that suggested by Feibelman [33] to explain the low barrier to self diffusion on Al(OO1). In this theoretical work it was found, that for the diffusion of an (aluminum) adatom, located in a four-fold hollow position, it is energetically more favourable to move one of the four underlying substrate atoms out of the surface plane onto a neighbouring four-fold hollow site on top of the surface and to fill the created vacancy by the adatom, instead of hopping of the adatom itself to a neighbouring fourfold hollow site via an intermediate two-fold bridge position. This concerted displacement mechanism involves a replacement of substrate atoms by adatoms, just as observed for the Fe-onCu system. The expelled Cu atoms may then diffuse across the surface and form copper islands as suggested in recent STM studies [18,301. 5. Conclusion In summary, our measurements show unambiguously that initially iron does not grow in a
Th. Detzel et al. /Growth of ultrathin iron films on Cu(OO1)
layer-by-layer fashion at room temperature. In particular it was demonstrated that one monolayer of iron does not lie down on the Cu(OO1) substrate reflecting the tendency for island formation, which is anticipated by thermodynamic considerations. The LEIS measurements show that for all coverages (even as low as 0.1 ML) about equal amounts of Fe atoms are visible in the surface and the first subsurface layer, respectively. At smallest coverages (- 0.1 ML) all iron atoms or islands are surrounded by copper atoms. Simultaneously copper defects appear on the surface. The iron films exhibit maximum roughness around 1 ML and smoothen with increasing film thickness. At coverages exceeding 2 nominal ML iron is found to wet the substrate for the most part. For the 6 ML film the morphological and crystallographic order closely resembles that of the clean Cu(OO1) surface. At higher coverages the roughness increases once more and above _ 10 ML the pseudomorphic growth of fee iron breaks down and the stable bee structure of iron, with the surface normal along (llO), is formed. Taking together the present LEIS results and previous STM work, we suggest a distorted bilayer growth with about one half of the iron incorporated in the substrate for coverages up to 2 ML. This automatically requires transport of Cu atoms, as was observed experimentally. Therefore Cu diffusion cannot be neglected for the growth at and especially above room temperature. Finally we would like to mention that the present work shows that LEIS is a versatile tool to study growth and diffusion of ultrathin films. It offers a complementary approach to modern surface microscopy techniques because of its element resolving abilities, but also because the probing depth is not only restricted to the topmost surface layer.
Acknowledgements
The authors would like to thank Staib Instruments and in particular B. Senftinger for placing generously at our disposal a commercial computer controlled video system. We are grateful to
237
V. Dose for encouragement and support and to E. Taglauer and E. Bertel for helpful discussions during the preparation of this publication. N.M. wants to thank H. Diirr for the introduction to the ion-scattering technique.
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