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Monolayer resolution in medium energy ion scattering
Nuclear Instruments and Methods in Physics Research B 183 (2001) 62±72
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Monolayer resolution in medium energy ion scattering P. Bailey
a,*
, T.C.Q. Noakes a, C.J. Baddeley b, S.P. Tear c, D.P. Woodru
d
a
b
CLRC Daresbury Laboratory, Daresbury, Warrington, WA4 4AD, UK School of Chemistry, University of St. Andrews, St Andrews, KY16 9ST, UK c Department of Physics, University of York, York, YO10 5DD, UK d Department of Physics, University of Warwick, Coventry, CV4 7AL, UK Received 16 November 2000; received in revised form 16 January 2001
Abstract Although medium energy ion scattering (MEIS) could be considered to be a mature technique, it continues to ®nd new applications and new methodologies, especially in exploiting the monolayer resolution capabilities. In this paper we consider four applications which provide information with monolayer resolution and illustrate each with a practical example. Ó 2001 Elsevier Science B.V. All rights reserved. PACS: 61.18.Bn; 61.85.+p; 68.47)De Keywords: Medium energy ion scattering; High resolution; Surface structures; Surface composition
1. Introduction Medium energy ion scattering (MEIS) [1] is a speci®c re®nement of the more common technique of Rutherford backscattering (RBS) but with greatly enhanced energy resolution. The energy range usually associated with MEIS is between 50 and 500 keV [2] although the limits are not rigidly de®ned. In simple terms, MEIS occupies a niche between low energy ion scattering (LEIS) and RBS. LEIS utilises an energy regime where detectors with high spatial and energy resolution are readily available; however, screening eects which
*
Corresponding author. Tel.: +44-0-1925-603-404; fax: +440-1925-603-173. E-mail address: [email protected] (P. Bailey).
are very strong at these low energies are dicult to model and trajectory-dependent neutralisation can substantially complicate the interpretation [3]. Also, because the scattering cross-section is high, multiple scattering can be problematic. RBS does not suer from the multiple scattering problem and the scattering is almost purely Coulombic but, because of the high energies involved, comparable detectors are technically demanding [4±8] and expensive. MEIS occupies an energy regime high enough for multiple scattering and screening to be negligible (or tolerable) but low enough so that a comparatively simple electrostatic detector with high spatial and energy resolution can be used. Whilst it is generally the energy resolution of the analyser that allows monolayer resolution to be achieved, this is not necessarily always the case. The most popular analyser used in MEIS
0168-583X/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 1 ) 0 0 3 8 4 - 6
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experiments is based on the FOM AMOLF toroidal electrostatic design [9] and utilises a 2D detector [10] for ecient data collection by simultaneous detection of ion energy and angle. For many years, the major application of MEIS was the determination of the structure of reconstructed single crystal surfaces [11]. In more recent years it has been applied to studies of surface melting [12], the composition of alloy surfaces [13], oxide growth [14] and the structure of buried layers [15]. In many cases in ion scattering, monolayer depth resolution is realised by employing an ion energy analyser with an energy resolution which is smaller than the inelastic energy loss experienced by ions traversing the layer. The resolution required in a particular instance is then dependent on the stopping power of the layer at the relevant energy and also upon the path length of the ion. Fig. 1 illustrates the variation in stopping power of Ni as a function of energy for H and He ions calculated using SRIM [16]. SRIM adopts a semiempirical approach which assumes a homogeneous substrate. In monolayer resolution applications of MEIS this assumption should be questioned since there is a change from a macroscopic, continuum situation, where inelastic energy loss is simply proportional to the path length, to a `microscopic' situation, where inelastic energy loss is dependent on the interaction of the ion with small numbers of electrons in speci®c states. Speci®cally in the case of ions travelling through channels in highly crystalline materials (as in double-alignment ex-
periments) the exact path taken by the ion can be important. At a simplistic level, it might be expected that an ion which traverses a region of relatively high electron density loses a greater amount of inelastic energy (the so-called trajectory-dependent stopping). This eect would give rise to an angular variation in stopping power which would eectively intermix some of the information from adjacent layers, thereby possibly compromising monolayer resolution. Evidence for trajectory-dependant stopping has been reported [17] and the eect has been modelled [18], with the latter study also modelling the energy loss and straggling from a single atomic layer. However, a further study [19] concluded that the experimental results contradict the simple picture of increased energy loss along trajectories with high core electron density and that simple electron-density arguments are not sucient to explain the observed eects. Hence, due to the complexity of the problem and the relative sparsity of information, noncontinuum inelastic energy loss eects are not generally considered in MEIS experiments. In Fig. 1, the data might seem to suggest that monolayer resolution is easiest to achieve if one uses high energy He ions. However, when the energy dependent resolution of a typical MEIS analyser [9,10,20] is folded in with the stopping power it becomes clear (Fig. 2) that the best resolution is achieved at the lowest energies. For this type of electrostatic de¯ection analyser, the energy window is a ®xed percentage of the pass energy
Fig. 1. Stopping power of Ni calculated from SRIM for H and He ions in the MEIS energy regime.
Fig. 2. Path-length resolution vs. energy for H and He ions (for Ni) in the MEIS energy regime for the HVEE analyser.
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and the resolution a ®xed percentage of the energy window. Its resolution is thus not a constant energy value but is a constant proportion of the detection energy. The data in Fig. 2 have been generated by taking the analyser resolution (a ®xed percentage of the energy) and dividing it by the stopping power data shown in Fig. 1. The ordinate in Fig. 2 is thus the path length which corresponds to the analyser resolution (the pathlength resolution). To illustrate this, the stopping the power of 50 keV He ions is about 25 eV/A, standard resolution at 50 keV is 150 V 0:3% 50; 000 and the path length correspond Siming to this resolution is thus 150/25 6 A. ilarly, for 100 keV He ions the stopping is 37 eV/ the resolution is 300 V and the corresponding A, To convert path-length resolution is now 8.1 A. the ordinate of Fig. 2 into a depth resolution, one must include the path length associated with the scattering geometry; for example, using a 45° in, 45° out scattering geometry, the standard resolu derived tion for 100 keV He ions is about 3 A, from a path-length resolution of 8.1 A and adjusted for the inward and outgoing trajectories by a factor of 2.8. The `high resolution' condition used in Fig. 2 relates to a vertical beam height of 0.1 mm or less, whilst the standard resolution uses a beam height of 0.5 mm. For this analyser it is the vertical beam height which is important because its energy dispersing plane is vertical. Hence, the best depth resolution is obtained at the lowest energies. For example, in Fig. 1, in going down in energy from 200 to 50 keV the resolution decreases by a factor of 4 whilst the H stopping hardly changes and the He stopping decreases by only a factor of 2. Although stopping power curves do vary from material to material, it is the linear energy dependence of the analyser resolution that is the dominant factor. It is quite clear from Fig. 2 that monolayer resolution is, in principle, readily achievable using this analyser in the energy regime shown. Using 50 keV H or He ions with the analyser in high resolution mode one can anticipate a theoretical depth resolution of even for 180° scattering (a path length about 1.5 A correction factor of 2). In this paper we present examples of four dierent ways in which monolayer resolution has been achieved using MEIS.
2. `Classic' pMEIS psurface structure determination ± Ni 1 1 1 3 3R30°±Pb The determination of the structures of reconstructed single crystal surfaces is a problem which has been addressed by MEIS experimenters for many years. In this application the structural information has generally been extracted from the raw data by integrating in energy the signal from the topmost atomic layers to produce surface blocking curves. The resultant blocking curves can then be compared quantitatively with the results of simulations conducted on trial structures using the VEGAS code [21,22]. p pIn the example presented here, Ni(1 1 1)( 3 3)R30°±Pb, the relevant surface structural information is contained within the topmost six layers. Using a 2D, angle and energy, position-sensitive detector the signal corresponding to six layers is extracted from the raw data as follows. Blocking curves are generated from the 2D data by integrating over a small energy range at each angular data point. The energy range employed corresponds to the inelastic energy loss associated with the penetration of six layers and is angle dependent to account for the variation in path length with angle. It is important to use such a precise method in extracting the blocking curves to ensure that the curves for each experimental geometry contain data from just the desired number of layers and particularly that this does not change within the angular range. For the set-up used in this experiment (100 keV H ions, standard resolution and a high scattering angle), However, the experimental resolution is about 3 A. in amalgamating the top six layers, we eectively ± degrade the resolution to 6 2:03 12 A some four times worse. Yet, such data still allow determination of the surface structure atomic layer by atomic layer with an accuracy of a few per cent. It is informative to consider what makes this possible. Broadly speaking, the approach relies on two things: (1) shadowing which inhibits sub-surface illumination such that the yield from deeper layers is greatly reduced and (2) prior knowledge ± speci®cally a knowledge of the substrate crystal structure, the scattering geometry and an appreciation that any atom movement must be signi®cantly less than the inter-atomic spacing. Of
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course, one might argue that one knows more or less what the structure is at the start, i.e. an ideal fcc(1 1 1) crystal, and one only considers a small number of minute perturbations from the ideal structure. Nonetheless, the approach is robust enough to have been successfully applied to complex problems involving alloys [23,24] and compound surfaces [25±27]. The trial structures covering what is believed to be a sucient region of parameter space are used to produce simulated blocking curves. An optimum structure is identi®ed from within this range. Once the optimum structure has been determined, the layer-by-layer structure is then known from the simulation input parameters. Thus it is the simulation which eectively deconvolves the data and thereby produces the atomic structure of the surface with layer-bylayer resolution. p p The Ni(1 1 1)( 3 3)R30°±Pb surface phase [28] is a convenient example to illustrate this approach. Similar systems, Au and Mn on Cu(1 0 0) [24], have recently been shown to have a substitutionally alloyed surface layer while for Sb on Cu(1 1 1) and Ag(1 1 1) surfaces [23,29] the surface alloy layer is also found to have a stacking fault between the surface layer p and the p substrate. The structure of the Ni(1 1 1)( 3 3)R30°±Pb system is therefore of interest to establish if this stacking p p fault is characteristic for fcc(1 1 1) ( 3 3)R30°±metal overlayer phases. The procedure for determining the optimum surface structure (and the con®dence limits) by comparing the blocking curves with simulations has been described elsewhere [23,24]. The experimental blocking curves together with simulations of the derived p psurface structure for the Ni(1 1 1) ( 3 3)R30°±Pb system are shown in Fig. 3. Here, the experimental signal has been converted into an atomic yield by the use of a system calibration factor. The calibration was determined by measuring yield pthe p from two model systems: the Si(1 1 1)( 3 3)R30°±Au surface and a clean Ni(1 1 1) p surface pas described elsewhere [28]. The Ni(1 1 1)( 3 3)R30°±Pb surface structure was determined to be a Ni2 Pb alloy layer in fcc registry with the substrate in contrast to the hcp (or stacking fault) registry found in the X2 Sb alloy (X Ag, Cu) layer in the Ag(1 1 1)/Sb [29] and
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Fig. 3. Experimental curves (data points) obtained p blocking p from a Ni(1 1 1)( 3 3)R30°±Pb surface using 100 keV H ions in three incidence geometrics together with the corresponding best-®t simulations (solid curves).
Cu(1 1 1)/Sb [23] systems. Thus, a knowledge of the bulk crystal structure and the use of shadowing enable the surface structure to be extracted layerby-layer from raw data integrated in depth.
3. Strain in buried layers ± growth of Au on Cu(1 1 1) Au and Cu are both fcc metals with a fairly large dierence in lattice parameter (4.078 and respectively). The epitaxial growth of Au 3.615 A on any face of a Cu crystal may therefore be expected to result in a strained Au overlayer. On a Cu(1 1 1) face, an in-plane strain of )11% will occur if the Au exactly matches the Cu spacing. If the Au layer density is to be maintained at the bulk value, this in-plane strain will produce an out-ofplane strain of +27%. This large degree of tetragonal strain should be readily measurable using MEIS by determining the angular shift of major blocking features since, because the aspect ratio of the unit cell changes with the amount of strain, so the blocking angles also change. It is perhaps worth noting that an isotropic expansion of a unit cell would maintain the aspect ratio and not result in a change in blocking angles. The strain can be depth pro®led though the layer by monitoring the position of blocking features at dierent scattered ion energies and hence dierent depths within the
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layer. Preliminary investigations were carried out on the growth of Au on Cu(1 1 1) to determine the thinnest Au layer for which the surface was unaffected by interfacial strain (it is desirable to use the thinnest layer to minimise the eect of straggling) and a Au layer thickness of 10 ML was determined to be suitable. Hence, an experiment to pro®le the strain in the Au requires a resolution substantially Using calculations better than 10 ML (24 A). similar to those shown in Fig. 2 but allowing for a (instead of 36 eV/A Au stopping power of 30 eV/A for Ni) we calculate a resolution at the surface of for 100 keV He ions using high resolution 1.5 A mode. This is the best achievable resolution which will degrade with depth due to straggling. For the scattering geometry used ( 1 1 0 in [0 0 1] out) we to calculate the straggling [30] at a depth of 24 A be 3.3 A. Thus the resolution will change from 1.5 at the surface to 3.6 A at a depth of 24 A, an A increase of a factor of 2.4. To keep the resolution more constant over this depth range we used a standard experimental setting to give a resolution (1.3 ML) at the surface which degrades to of 3 A depth, only a 50% change. 4.5 A (1.9 ML) at 24 A The dependence of Au strain upon depth below the Au surface is shown in Fig. 4. No strain has been measured in the ®rst two layers because of the large scattering signal from the surface peak. From Fig. 4 it is apparent that the strain is highly localised at a depth of about 8 ML. The straggling at
this depth is calculated to correspond to 1.5 ML. The ®t shown in Fig. 4 is the result of convolving a Gaussian of FWHM 1.5 ML with a strain delta function of thickness two layers centred at a depth of 8.5 ML. The implication of this ®t is that the strain exists only in two layers, presumably at the Au/Cu interface. This approach does not take into account the decay in strain towards the surface. The strain relief that is evident between layer eight and layer six probably occurs via dislocations which should, in principle, be detectable with MEIS. However, twinning (i.e. random interfacial stacking faults) in the Au layer means that it is not possible to pro®le any disorder relating to strain relief. The maximum measured strain is thus about +8%, much smaller than the calculated value of +27%. We suggest that the reason for the lower strain is due to inter-diusion at the interface of the two elements. Depth pro®les of each element reveal that Au has diused into the Cu to a depth and Cu is present throughout the of about 30 A Au layer. Thus the overlayer/substrate interface is not abrupt but is actually a graded Au/Cu alloy. The binary phase diagram of the Au±Cu system shows that stable Aux Cu1 x alloys exist over the composition range x 0±1 [31]. It is this graded alloy interface that allows the lattice size to vary slowly over many layers. Nonetheless, substantial tetragonal strain is observed in a localised twolayer region and, by deconvolution based on a knowledge of the magnitude of straggling, has been pro®led layer-by-layer. 4. Layer-by-layer composition ± Cu50 Pd50 (1 1 0)
Fig. 4. The measured strain versus depth below surface for a nominal 10 ML of Au grown on Cu(1 1 1). The solid curve is a ®t produced by convolving the calculated resolution with a strain delta function of thickness two layers.
The physical and chemical properties of bimetallic surfaces are of huge technological interest, particularly in the design of bimetallic catalysts for de-chlorination and in the investigation of surface corrosion [32±34]. It is well known that the composition of a bimetallic surface may dier strongly from that of the bulk material, these dierences being driven by a complex interplay of thermodynamic parameters such as relative surface energies, relative bonding strengths of each metallic component with adsorbed atoms and molecules, the tendency of the metal components to form
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ordered alloy phases and kinetic parameters such as diusion barriers. The quantitative analysis of the layer-by-layer composition of a bimetallic surface is a non-trivial analytical task. The shadowing and blocking capabilities of MEIS make this technique unique for obtaining layer-by-layer depth pro®le information. MEIS has been used by two groups to quantify the near surface composition of bimetallic surfaces. These studies investigated the Fe72 Cr28 (1 1 0) surface [35] and the Pt50 Ni50 (1 1 1) surface [36,37]. Whilst these studies are essentially extremely accurate investigations on well-de®ned surfaces, it has recently been shown that it is also possible to glean compositional information of the surface region of Cu50 Pd50 (1 1 0) in the presence of complex adsorbate layers [38]. For clarity we will con®ne our discussions to the Cu50 Pd50 (1 1 0) system. The approach used is to take a measurement in a particular geometry such that the ion yield is from the top layer only and then in a second geometry where the yield is from the top two layers only and so on for three and four layers. This is illustrated schematically in Fig. 5; the choice of out-going geometry will be discussed later. For each of the four geometries the Cu and Pd signal can be extracted since, due to
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their dierent masses, they are well separated in energy. If a, b, c, d are the fractional amounts of Cu in layer 1, 2, 3, 4 respectively and w, x, y, z are the measured Cu signals in the 1-, 2-, 3-, and 4layer geometries, then, for the case of an ideal bulk terminated crystal at 0 K, we can write aCu 1:0wCu ;
1
aCu bCu 2:0xCu ; aCu bCu cCu 3:0yCu ;
2 3
aCu bCu cCu dCu 4:00zCu :
4
These four equations can easily be solved to give the four layer compositions, a, b, c, d. The geometries are chosen so that the ion beam is channelled in a particular incident geometry where the eect of shadowing minimises the illumination of deeper layers. Of course the shadowing can never be complete. Thermal vibrations and surface relaxations cause unwanted additional illumination of deeper layers. However, these eects can be lessened by utilising the eect of blocking ± the socalled double alignment technique [1] whereby ions are channelled in and also channelled out. The concept is illustrated schematically in Fig. 5 where the four scattering geometries are shown. In Fig. 5(a), an incident 1 0 1 geometry has been chosen to achieve the best single layer illumination but even so, simulations show that the total illumination in this geometry is actually 1.8 layers, with the majority of the extra illumination being in the second layer. However, using the double alignment geometry, 1 0 1 in [0 1 1] out, the eect of blocking substantially reduces the total illumination to 1.45 layers. The extra illumination alters Eqs. (1)±(4) to become 1:0aCu 0:42bCu 0:03cCu 0:0dCu 1:45wCu ; 5 1:0aCu 1:0bCu 0:2cCu 0:06dCu 2:26xCu ; 6 1:0aCu 1:0bCu 1:0cCu 0152dCu 3:15yCu ; 7 1:0aCu 1:0bCu 1:0cCu 1:0dCu 4:00zCu :
Fig. 5. Cross-sectional diagrams of the fcc(1 1 0) surface viewed in the h 1 1 2i azimuth: (a) the one layer illumination geometry 1 0 1 in [0 1 1] out, (b) the two layer geometry 1 1 0 in 1 3 4 out, (c) the three layer geometry 3 1 2 in [3 1 2] out and (d) the four layer geometry 2 1 1 in [1 2 1] out.
8
A numerical solution of these equations is considerably more dicult than the special case of Eqs. (1)±(4) but can be made trivial by employing determinants and a standard software package. The choice of incident ion in these experiments is
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important; generally speaking, He ions are better in this application than H ions for two reasons. Firstly, the larger shadow cone of the heavier ion helps to reduce the unwanted subsurface illumination and, secondly, the greater mass separation of the heavier ion allows one a greater range of scattering angles and hence more freedom in choosing the best geometries. There are a number of aspects of the simulations which warrant discussion: 1. The substrate composition ± because the substrate is bimetallic with random occupancy and a variable layer-by-layer composition, it is not possible to simulate the structure precisely. Instead we have used a crystal composed of a single element, Sr which has an atomic number (38) midway between Cu (29) and Pd (46). 2. Surface relaxation ± by simulations of a surface layer relaxation of magnitude 0±10% we see that its eect on the layer illumination is negligible. The reason for this is the large shadow cone of the 100 keV He ions; for the one layer geometry at the second atom in the chain, the sha and for Pd is dow cone radius for Cu is 0.29 A 0.37 A, both of which are greater than a 10% Blocking further relayer expansion (0.21 A). duces any residual eect. 3. Enhanced surface thermal vibrations ± these were found to have a strong eect on the layer illuminations. Literature values for the surface Debye temperature of Cu(1 1 0) range from 112 to 313 K depending on the orientation of measurement [39±45]. The only reported value for Cu50 Pd50 (1 1 0) is 200±250 K [46]. These values correspond to enhancements of the root-mean-
square vibrational amplitudes from 1.1 to 3.0, giving one layer yields from 1.1 to 1.8 layers. We used a surface enhancement of 100% giving a root-mean-square surface vibrational ampli This enhancement was allowed tude of 0.14 A. to decay exponentially into the bulk. 4. Possible adsorbate induced reconstructions ± effects of variable amounts of adsorbates on the surface were assessed by comparing the experimentally measured yield as a function of Cl coverage. The greatest eect is in the one layer geometry and even then no clear trend is seen. As a test of the assumptions made in the simulations, a comparison can be made between the normalised experimental yield and the predicted yields from simulations for each geometry. Fourteen separate measurements of the clean Cu50 Pd50 (1 1 0) surface were made and their average values were found to match the results of simulations to within 5%. There is, therefore, a good agreement between experiment and the theoretical model used in the simulations. When this technique is applied to the problem of segregation in the Cu50 Pd50 (1 1 0)/C2 H4 x Clx system, we see that hydrocarbon adsorption produces a measurable Pd enrichment in the surface layer, while Cu enrichment is produced by the presence of Cl(ads) and C(ads). Cu50 Pd50 (1 1 0)/ C2 H4 x Clx surfaces were prepared by adsorption of ethene or 1,2-dichloroethene (DCE) on clean Cu50 Pd50 (1 1 0) and subsequent annealing. Table 1 summarises the results. The composition of the ®rst layer of the clean surface is determined to be 65% Cu, in good agreement with the value of 70% quoted for similar preparation conditions [47±49].
Table 1 Cu content of the ®rst four layers of a Cu50 Pd50 (1 1 0) crystal after dierent surface treatments Experiment
First layer (% Cu)
Second layer (% Cu)
Clean Cu50 Pd50 (1 1 0) 10 ML ethene at 323 K 10 ML ethene ¯ashed to 383 K 10 ML ethene ¯ashed to 473 K 10 ML ethene ¯ashed to 573 K 10 ML 1,2-DCE at 323 K 10 ML DCE ¯ashed to 383K 10 ML DCE ¯ashed to 473 K 10 ML DCE ¯ashed to 573 K
65 38 57 53 55 47 72 30 76 28 71 30 72 29 88 7 Data suggest surface restructuring
Third layer (% Cu)
Fourth layer (% Cu)
71 45 61 69 62 66 62 76
50 80 66 71 81 78 73 65
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The values in Table 1 are plotted in Fig. 6(a) and (b). The 300 K data points in Fig. 6 are for the clean surface and serve to show how the clean surface changes when covered by the adsorbate. Referring to Fig. 6(a), from the known chemistry of ethene on Pd(1 1 0) [50] it appears that dissociative ethene adsorption occurs at 323 K on small Pd ensembles to yield an ethynyl (C2 H) species, pinning Pd at the surface and thereby increasing the ®rst layer Pd content at the expense of the second layer. This is evidenced in Fig. 6(a) by the reduction in ®rst layer Cu content and the increase in second layer content. After annealing at 573 K, the ®rst layer is more Cu rich since the ethnyl is known to decompose to C(ads). Thus, Cy Hx spe-
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cies cause Pd enrichment at the surface whilst C(ads) causes Cu enrichment. In the case of DCE adsorption, from a knowledge of the chemistry of DCE on Cu(1 1 0) [51] and Pd(1 1 0) [52] and an appreciation of the anity between Cu and Cl, decomposition is preferred to desorption such that, at 473 K, C(ads) and Cl(ads) are the species present at the surface. The C(ads) and Cl(ads) strongly enhance the ®rst layer Cu content (88% Cu) at the expense of the second layer (7% Cu). The anomalously high MEIS background seen upon annealing this surface to 573 K suggests a restructuring of the surface, possibly caused by the presence of CuClx islands, an eect which has also been seen for Cl2 adsorption on Ag(1 1 0) [53]. It is apparent, therefore, that the judicious use of double alignment in several scattering geometries can provide a clear-cut layer-by-layer composition pro®le of a bimetallic crystal surface. 5. Buried layer structure determination ± twodimensional rare earth silicides and germanides on (1 1 1) substrates
Fig. 6. Variation in the composition of the topmost four layers of the Cu50 Pd50 (1 1 0) surface as a function of annealing temperature following the adsorption of (a) 10 ML ethene and (b) 10 ML 1,2-DCE at 323 K. The 300 K data points represent the clean surface.
This ®nal example of achieving monolayer resolution relates more to sensitivity than instrument resolution. Because of the high degree of sensitivity of ion scattering to a heavy impurity in a light substrate, the impurity signal can easily be resolved from that of the substrate. For clarity we will refer to the heavy impurity as HI and the light substrate as LS. In the case of a monolayer of HI on LS, if the mass dierence is suciently large, the HI signal occurs at a much higher energy than the LS signal and the two can easily be resolved. This may still be the case even if the inherent energy resolution of the analysing instrument is not sucient to resolve the energy lost by ions traversing the monolayer. However, in MEIS studies of systems comprising of a monolayer of HI on LS there is generally no structural information to be gained from the signal from the HI. Because the HI lies on the LS surface and is contained within a single layer, there will be no HI±HI or HI±LS blocking and hence no detail in the HI blocking curves, e.g. [23]. However, LS±HI blocking occurs in an unusual family of systems where a HI
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monolayer readily adopts sites beneath LS atoms. These are the 2D rare earth (RE) silicides and germanides on (1 1 1) substrates which are formed when 1 ML of RE is deposited on a clean Si(1 1 1) or Ge(1 1 1) substrate and annealed. In these cases, the surface converts into a well-ordered semi-metallic 2D layer of stoichiometric RESi2 or REGe2 . The RE forms a hexagonal monolayer accommodated underneath a buckled Si or Ge top layer where the buckled top layer is rotated by 180° in the surface plane (Fig. 7). Such 2D alloys are known to occur for Tm, Er, Ho, Dy and Gd on Si(1 1 1) [54,55] and Ho and Dy on Ge(1 1 1) [26,55]. The eect is believed to be limited to those REs which are trivalent at a coverage of 1 ML. Because the RE monolayer is buried beneath an ordered Si or Ge bilayer, blocking features occur in the RE signal and can be used for structural determination. Fig. 8 shows blocking curves for the case of Ge(1 1 1) 1 1±Dy (Ge provides a better illustration than Si as the blocking minima are stronger due to the higher Z of the Ge). The yield calibration procedure has been described earlier. The curves in Fig. 8 have been used to determine the position of the surface Ge atoms with respect to the buried Dy atoms and the results are presented in Table 2. The atom positions are de®ned in Fig. 7. For 100 keV He ions in the Si(1 1 1) 1 1± Dy system, a calculation of the elastic scattering reveals that (for the geometries shown in Fig. 8) there is a separation of 18 keV between Dy and Si
Fig. 8. Experimental blocking curves (data points) obtained from a Ge(1 1 1) 1 1±Dy surface using 100 keV H ions together with the corresponding best-®t simulations (solid curves): (a) the 1 1 0/[0 0 1] scattering geometry, (b) the normal/ 1 1 1 scattering geometry. The insets show the origin of the major blocking dips.
signals. Thus, an instrument with a resolution of 15 kV FWHM at 100 kV could be successfully employed on the Si(1 1 1) 1 1±Dy system because the energy separation of (and thus the sensitivity to) the Dy is so great. A similar calculation shows that the structure determination of this monolayer system could be carried out using 2 MeV He ions and a conventional RBS instrument with a Si barrier detector if this could be scanned through a sucient in-plane angular range. In the case of these 2D RE (1 1 1) silicides and germanides, monolayer resolution could be
Table 2 The outermost layer spacing and Ge±Dy bond lengths for the Ge(1 1 1)1 1±Dy surface Bond length (A) Vertical distance (A) Fig. 7. A cross-sectional diagram of the Ge (1 1 1) surface viewed in the h1 1 2i azimuth illustrating the structure of the Ge(1 1 1) 1 1±Dy 2D germanide.
Ge(1)±Ge(2) Ge(2)±Dy
1:08 0:02 1:71 0:02
2:55 0:03 2:87 0:03
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considered to be inherent when applying any ion scattering technique to these systems. 6. Conclusions Monolayer resolution in MEIS might generally be associated solely with the resolving power of the ion energy analyser. There are, however, several MEIS techniques which can be used to achieve monolayer resolution which do not necessarily rely on an analyser having a resolution in energy comparable to that lost by an ion traversing a single layer. In the classic MEIS surface structure experiments, the signal from a number of layers is amalgamated giving an apparent resolution of perhaps six layers. However, a detailed knowledge of the bulk crystal structure combined with best ®t simulations allow the layer-by-layer structure to be determined from the simulation input data. Even where monolayer resolution is readily available from the ion energy analyser, the straggling associated with penetration of buried layers will ultimately dominate the eective resolution. In this case an ability to calculate the straggling is necessary to perform a deconvolution. If the straggling is only of the order of a few layers we suggest that an accurate deconvolution can be carried out to yield meaningful layer-by-layer information. The layer-by-layer compositional analysis of an alloy surface is an extreme example of an application requiring monolayer resolution. The unique shadowing and blocking capabilities of MEIS are well suited to this problem using the concept of double alignment. It is perhaps surprising that the compositional information can be determined even though the surface is covered by complex and unstable adsorbates. We ®nd that the adsorbateinduced compositional changes are quite large and any uncertainties in the modelling of the surface produce only minor variations in the composition values calculated. In the unusual case of an ordered buried layer of heavy impurity covered by an ordered layer of light substrate atoms, the application of an ion scattering technique confers monolayer resolution almost inherently. The large mass separation of the REs
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compared to the substrate in the 2D RE-(1 1 1) silicide and germanide systems ensure a high degree of sensitivity to the RE. In these systems, monolayer resolution can be achieved without use of an ion energy analyser with high energy resolution. Four examples of very dierent approaches to achieving monolayer resolution in MEIS have been described and illustrated. These examples are not meant to be an exhaustive compilation of all the possible ways to achieve monolayer resolution in MEIS ± other approaches exist and yet more may be discovered.
Acknowledgements The authors would like to acknowledge Kevin Connell and Brian Blackwell for technical help, the FOM Institute for the provision of the VEGAS code and the Engineering and Physical Science Research Council for funding this research.
References [1] J.F. van der Veen, Surf. Sci. Rep. 5 (1985) 199. [2] R.M. Tromp, in: D. Briggs, M.P. Seah, (Eds.), Practical Surface Analysis, 2nd ed., Vol. 2, Wiley, New York, 1992. [3] D.P. Woodru, Nucl. Instr. and Meth. B 194 (1982) 639. [4] W.A. Lanford, S. Bedell, S. Amadon, A. Haberl, W. Skala, B. Hjorvarsson, Nucl. Instr. and Meth. B 161 (2000) 202. [5] K. Kimura, K. Nakajima, M. Mannami, Nucl. Instr. and Meth. B 138 (1998) 1196. [6] D.O. Boerma, W.M. Arnoldbik, W. Wolfswinkel, Mat. Sci. Forum 248 (1997) 451. [7] W.M. Arnoldbik, W. Wolfswinkel, D.K. Inia, V.C.G. Verleun, S. Lobner, J.A. Reinders, F. Labohm, D.O. Boerma, Nucl. Instr. and Meth. B 118 (1996) 566. [8] H.D. Carstanjen, Mat. Sci. Forum 248 (1997) 459. [9] R.G. Smeek, R.M. Tromp, H.H. Kersten, A.J.H. Boerboom, F.W. Saris, Nucl. Instr. and Meth. B 195 (1982) 581. [10] R.M. Tromp, M. Copel, M.C. Reuter, H. Horn von Hoegen, J. Speidell, R. Koujis, Rev. Sci. Instrum. 62 (1992) 2679. [11] P.R. Watson, J. Phys. Chem. Ref. Data 19 (1990) 85. [12] J.F. Frenken, J.F. van der Veen, Phys. Rev. Lett. 85 (1985) 134. [13] D.J. O'Connor, M. Draeger, A.M. Molenbroek, Y.G. Shen, Surf. Sci. 357±358 (1996) 202. [14] E.P. Gusev, H.C. Lu, T. Gustafsson, E. Garfunkel, Phys. Rev. B 52 (1995) 1759.
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[15] T.C.Q. Noakes, P. Bailey, P.K. Hucknall, K. Donovan, M.A. Howson, Phys. Rev. B 58 (1998) 4934. [16] SRIM code, J.F. Ziegler, IBM Yorktown Heights, NY, http://www.research.ibm.com/ionbeams. [17] J.F. Frenken, J.F. van der Veen, Phys. Rev. Lett. 54 (1985) 134. [18] P.F.A. Alkemade, W.C. Turkenburg, W.F. van der Weg, Nucl. Instr. and Meth. B 28 (1987) 161. [19] D.E. Fowler, J.V. Barth, Phys. Rev. B 52 (1994) 2117. [20] HVEE Ltd. manufacturers, Amersfoort, The Netherlands. [21] R.M. Tromp, J.F. van der Veen, Surf. Sci. 133 (1983) 159. [22] J.F. Frenken, J.F. van der Veen, R.M. Tromp, Nucl. Instr. and Meth. B 17 (1986) 334. [23] P. Bailey, T.C.Q. Noakes, D.P. Woodru, Surf. Sci. 426 (1999) 358. [24] D. Brown, T.C.Q. Noakes, D.P. Woodru, P. Bailey, Y. le Goaziou, J. Phys. Condens. Matter 11 (1999) 1889. [25] M. Lohmeier, W.J. Huisman, G. ter Horst, P.M. Zagwijn, E. Vlieg, C.L. Nicklin, T.S. Turner, Phys. Rev. B 54 (1996) 2004. [26] D.J. Spence, S.P. Tear, P. Bailey, T.C.Q. Noakes, Phys. Rev. B 61 (2000) 5707. [27] D.J. Spence, P. Bailey, T.C.Q. Noakes, S.P. Tear, Phys. Rev. B 62 (2000) 5016. [28] D. Brown, P.D. Quinn, D.P. Woodru, P. Bailey, T.C.Q. Noakes, Phys. Rev. B 61 (2000) 7706. [29] S.A. de Vries, W.J. Huisman, P. Goedtkint, M.J. Zwanenburg, S.L. Bennet, I.K. Robinson, E. Vlieg, Surf. Sci. 414 (1998) 159. [30] W.K. Chu, Phys. Rev. A 13 (1976) 2057. [31] ASM Handbook, Vol. 3, Alloy Phase Diagrams, ASM International, 1999. [32] J.A. Rodriguez, Surf. Sci. Rep. 24 (1996) 225. [33] D.W. Goodman, Surf. Sci. 300 (1994) 834. [34] C.T. Campbell, Ann. Rev.Phys. Chem. 41 (1990) 775. [35] C. Xu, D.J. O'Connor, Nucl. Instr. and Meth. B 53 (1991) 326.
[36] S. Deckers, F.H.P.M. Hadraken, W.F. van der Weg, A.W.D. van der Gon, B. Pluis, J.F. van der Veen, R. Baudoing, Phys. Rev. B 42 (1990) 3235. [37] S. Deckers, F.H.P.M. Hadraken, W.F. van der Weg, A.W.D. van der Gon, J.F. van der Veen, J.W. Geus, Appl. Surf. Sci. 24 (1990) 121. [38] C.J. Baddeley, L.H. Bloxham, S.C. Laroze, R. Raval, T.C.Q. Noakes, P. Bailey, Surf. Sci. 433±435 (1999) 827. [39] D.P. Jackson, Surf. Sci. 43 (1974) 431. [40] Y.N. Devyatko, S.V. Rogozhkin, V.I. Troyan, E.P. Gusev, T. Gustafsson, J. Exp. Theor. Phys. 89 (1999) 1103. [41] H. Durr, R. Schneider, T. Fauster, Vacuum 41 (1990) 376. [42] R. Schneider, H. Durr, T. Fauster, V. Dose, Phys. Rev. B 42 (1990) 1638. [43] T. Fauster, Vacuum 38 (1988) 129. [44] E. van de Riet, J.M. Fluit, A. Niehaus, Vacuum 41 (1990) 372. [45] S. Speller, S. Parascandola, W. Heiland, Surf. Sci. 383 (1997) 131. [46] J. Loboda-Cackovic, Vacuum 46 (1995) 1449. [47] J. Loboda-Cackovic, M.S. Mousa, J.H. Block, Vacuum 46 (1995) 89. [48] J. Loboda-Cackovic, Vacuum 48 (1997) 1405. [49] J. Loboda-Cackovic, Vacuum 48 (1997) 571. [50] J. Yoshinobu, T. Sekitani, M. Onchi, M. Nishijima, J. Electron Spectros. Relat. Phenom. 54 (1990) 697. [51] Y. Jugnet, N.S. Prakash, J.C. Bertolini, S.C. Laroze, R. Raval, Catal. Lett. 56 (1998) 17. [52] L.H. Bloxham, S. Haq, C. Mitchell, R. Raval, J. Chem. Phys., submitted. [53] B.V. Andryushechkin, K.N. Eltsov, V.M. Shevlyuga, Surf. Sci. 433±435 (1999) 109. [54] P. Paki, U. Kadafer, P. Wetzel, C. Pirri, J.C. Peruchetti, D. Bolmont, G. Gewinner, Phys. Rev. B 45 (1992) 8490. [55] D.J. Spence, P. Bailey, T.C.Q. Noakes, S.P. Tear, in preparation.
www.elsevier.com/locate/nimb
Monolayer resolution in medium energy ion scattering P. Bailey
a,*
, T.C.Q. Noakes a, C.J. Baddeley b, S.P. Tear c, D.P. Woodru
d
a
b
CLRC Daresbury Laboratory, Daresbury, Warrington, WA4 4AD, UK School of Chemistry, University of St. Andrews, St Andrews, KY16 9ST, UK c Department of Physics, University of York, York, YO10 5DD, UK d Department of Physics, University of Warwick, Coventry, CV4 7AL, UK Received 16 November 2000; received in revised form 16 January 2001
Abstract Although medium energy ion scattering (MEIS) could be considered to be a mature technique, it continues to ®nd new applications and new methodologies, especially in exploiting the monolayer resolution capabilities. In this paper we consider four applications which provide information with monolayer resolution and illustrate each with a practical example. Ó 2001 Elsevier Science B.V. All rights reserved. PACS: 61.18.Bn; 61.85.+p; 68.47)De Keywords: Medium energy ion scattering; High resolution; Surface structures; Surface composition
1. Introduction Medium energy ion scattering (MEIS) [1] is a speci®c re®nement of the more common technique of Rutherford backscattering (RBS) but with greatly enhanced energy resolution. The energy range usually associated with MEIS is between 50 and 500 keV [2] although the limits are not rigidly de®ned. In simple terms, MEIS occupies a niche between low energy ion scattering (LEIS) and RBS. LEIS utilises an energy regime where detectors with high spatial and energy resolution are readily available; however, screening eects which
*
Corresponding author. Tel.: +44-0-1925-603-404; fax: +440-1925-603-173. E-mail address: [email protected] (P. Bailey).
are very strong at these low energies are dicult to model and trajectory-dependent neutralisation can substantially complicate the interpretation [3]. Also, because the scattering cross-section is high, multiple scattering can be problematic. RBS does not suer from the multiple scattering problem and the scattering is almost purely Coulombic but, because of the high energies involved, comparable detectors are technically demanding [4±8] and expensive. MEIS occupies an energy regime high enough for multiple scattering and screening to be negligible (or tolerable) but low enough so that a comparatively simple electrostatic detector with high spatial and energy resolution can be used. Whilst it is generally the energy resolution of the analyser that allows monolayer resolution to be achieved, this is not necessarily always the case. The most popular analyser used in MEIS
0168-583X/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 1 ) 0 0 3 8 4 - 6
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experiments is based on the FOM AMOLF toroidal electrostatic design [9] and utilises a 2D detector [10] for ecient data collection by simultaneous detection of ion energy and angle. For many years, the major application of MEIS was the determination of the structure of reconstructed single crystal surfaces [11]. In more recent years it has been applied to studies of surface melting [12], the composition of alloy surfaces [13], oxide growth [14] and the structure of buried layers [15]. In many cases in ion scattering, monolayer depth resolution is realised by employing an ion energy analyser with an energy resolution which is smaller than the inelastic energy loss experienced by ions traversing the layer. The resolution required in a particular instance is then dependent on the stopping power of the layer at the relevant energy and also upon the path length of the ion. Fig. 1 illustrates the variation in stopping power of Ni as a function of energy for H and He ions calculated using SRIM [16]. SRIM adopts a semiempirical approach which assumes a homogeneous substrate. In monolayer resolution applications of MEIS this assumption should be questioned since there is a change from a macroscopic, continuum situation, where inelastic energy loss is simply proportional to the path length, to a `microscopic' situation, where inelastic energy loss is dependent on the interaction of the ion with small numbers of electrons in speci®c states. Speci®cally in the case of ions travelling through channels in highly crystalline materials (as in double-alignment ex-
periments) the exact path taken by the ion can be important. At a simplistic level, it might be expected that an ion which traverses a region of relatively high electron density loses a greater amount of inelastic energy (the so-called trajectory-dependent stopping). This eect would give rise to an angular variation in stopping power which would eectively intermix some of the information from adjacent layers, thereby possibly compromising monolayer resolution. Evidence for trajectory-dependant stopping has been reported [17] and the eect has been modelled [18], with the latter study also modelling the energy loss and straggling from a single atomic layer. However, a further study [19] concluded that the experimental results contradict the simple picture of increased energy loss along trajectories with high core electron density and that simple electron-density arguments are not sucient to explain the observed eects. Hence, due to the complexity of the problem and the relative sparsity of information, noncontinuum inelastic energy loss eects are not generally considered in MEIS experiments. In Fig. 1, the data might seem to suggest that monolayer resolution is easiest to achieve if one uses high energy He ions. However, when the energy dependent resolution of a typical MEIS analyser [9,10,20] is folded in with the stopping power it becomes clear (Fig. 2) that the best resolution is achieved at the lowest energies. For this type of electrostatic de¯ection analyser, the energy window is a ®xed percentage of the pass energy
Fig. 1. Stopping power of Ni calculated from SRIM for H and He ions in the MEIS energy regime.
Fig. 2. Path-length resolution vs. energy for H and He ions (for Ni) in the MEIS energy regime for the HVEE analyser.
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and the resolution a ®xed percentage of the energy window. Its resolution is thus not a constant energy value but is a constant proportion of the detection energy. The data in Fig. 2 have been generated by taking the analyser resolution (a ®xed percentage of the energy) and dividing it by the stopping power data shown in Fig. 1. The ordinate in Fig. 2 is thus the path length which corresponds to the analyser resolution (the pathlength resolution). To illustrate this, the stopping the power of 50 keV He ions is about 25 eV/A, standard resolution at 50 keV is 150 V 0:3% 50; 000 and the path length correspond Siming to this resolution is thus 150/25 6 A. ilarly, for 100 keV He ions the stopping is 37 eV/ the resolution is 300 V and the corresponding A, To convert path-length resolution is now 8.1 A. the ordinate of Fig. 2 into a depth resolution, one must include the path length associated with the scattering geometry; for example, using a 45° in, 45° out scattering geometry, the standard resolu derived tion for 100 keV He ions is about 3 A, from a path-length resolution of 8.1 A and adjusted for the inward and outgoing trajectories by a factor of 2.8. The `high resolution' condition used in Fig. 2 relates to a vertical beam height of 0.1 mm or less, whilst the standard resolution uses a beam height of 0.5 mm. For this analyser it is the vertical beam height which is important because its energy dispersing plane is vertical. Hence, the best depth resolution is obtained at the lowest energies. For example, in Fig. 1, in going down in energy from 200 to 50 keV the resolution decreases by a factor of 4 whilst the H stopping hardly changes and the He stopping decreases by only a factor of 2. Although stopping power curves do vary from material to material, it is the linear energy dependence of the analyser resolution that is the dominant factor. It is quite clear from Fig. 2 that monolayer resolution is, in principle, readily achievable using this analyser in the energy regime shown. Using 50 keV H or He ions with the analyser in high resolution mode one can anticipate a theoretical depth resolution of even for 180° scattering (a path length about 1.5 A correction factor of 2). In this paper we present examples of four dierent ways in which monolayer resolution has been achieved using MEIS.
2. `Classic' pMEIS psurface structure determination ± Ni 1 1 1 3 3R30°±Pb The determination of the structures of reconstructed single crystal surfaces is a problem which has been addressed by MEIS experimenters for many years. In this application the structural information has generally been extracted from the raw data by integrating in energy the signal from the topmost atomic layers to produce surface blocking curves. The resultant blocking curves can then be compared quantitatively with the results of simulations conducted on trial structures using the VEGAS code [21,22]. p pIn the example presented here, Ni(1 1 1)( 3 3)R30°±Pb, the relevant surface structural information is contained within the topmost six layers. Using a 2D, angle and energy, position-sensitive detector the signal corresponding to six layers is extracted from the raw data as follows. Blocking curves are generated from the 2D data by integrating over a small energy range at each angular data point. The energy range employed corresponds to the inelastic energy loss associated with the penetration of six layers and is angle dependent to account for the variation in path length with angle. It is important to use such a precise method in extracting the blocking curves to ensure that the curves for each experimental geometry contain data from just the desired number of layers and particularly that this does not change within the angular range. For the set-up used in this experiment (100 keV H ions, standard resolution and a high scattering angle), However, the experimental resolution is about 3 A. in amalgamating the top six layers, we eectively ± degrade the resolution to 6 2:03 12 A some four times worse. Yet, such data still allow determination of the surface structure atomic layer by atomic layer with an accuracy of a few per cent. It is informative to consider what makes this possible. Broadly speaking, the approach relies on two things: (1) shadowing which inhibits sub-surface illumination such that the yield from deeper layers is greatly reduced and (2) prior knowledge ± speci®cally a knowledge of the substrate crystal structure, the scattering geometry and an appreciation that any atom movement must be signi®cantly less than the inter-atomic spacing. Of
P. Bailey et al. / Nucl. Instr. and Meth. in Phys. Res. B 183 (2001) 62±72
course, one might argue that one knows more or less what the structure is at the start, i.e. an ideal fcc(1 1 1) crystal, and one only considers a small number of minute perturbations from the ideal structure. Nonetheless, the approach is robust enough to have been successfully applied to complex problems involving alloys [23,24] and compound surfaces [25±27]. The trial structures covering what is believed to be a sucient region of parameter space are used to produce simulated blocking curves. An optimum structure is identi®ed from within this range. Once the optimum structure has been determined, the layer-by-layer structure is then known from the simulation input parameters. Thus it is the simulation which eectively deconvolves the data and thereby produces the atomic structure of the surface with layer-bylayer resolution. p p The Ni(1 1 1)( 3 3)R30°±Pb surface phase [28] is a convenient example to illustrate this approach. Similar systems, Au and Mn on Cu(1 0 0) [24], have recently been shown to have a substitutionally alloyed surface layer while for Sb on Cu(1 1 1) and Ag(1 1 1) surfaces [23,29] the surface alloy layer is also found to have a stacking fault between the surface layer p and the p substrate. The structure of the Ni(1 1 1)( 3 3)R30°±Pb system is therefore of interest to establish if this stacking p p fault is characteristic for fcc(1 1 1) ( 3 3)R30°±metal overlayer phases. The procedure for determining the optimum surface structure (and the con®dence limits) by comparing the blocking curves with simulations has been described elsewhere [23,24]. The experimental blocking curves together with simulations of the derived p psurface structure for the Ni(1 1 1) ( 3 3)R30°±Pb system are shown in Fig. 3. Here, the experimental signal has been converted into an atomic yield by the use of a system calibration factor. The calibration was determined by measuring yield pthe p from two model systems: the Si(1 1 1)( 3 3)R30°±Au surface and a clean Ni(1 1 1) p surface pas described elsewhere [28]. The Ni(1 1 1)( 3 3)R30°±Pb surface structure was determined to be a Ni2 Pb alloy layer in fcc registry with the substrate in contrast to the hcp (or stacking fault) registry found in the X2 Sb alloy (X Ag, Cu) layer in the Ag(1 1 1)/Sb [29] and
65
Fig. 3. Experimental curves (data points) obtained p blocking p from a Ni(1 1 1)( 3 3)R30°±Pb surface using 100 keV H ions in three incidence geometrics together with the corresponding best-®t simulations (solid curves).
Cu(1 1 1)/Sb [23] systems. Thus, a knowledge of the bulk crystal structure and the use of shadowing enable the surface structure to be extracted layerby-layer from raw data integrated in depth.
3. Strain in buried layers ± growth of Au on Cu(1 1 1) Au and Cu are both fcc metals with a fairly large dierence in lattice parameter (4.078 and respectively). The epitaxial growth of Au 3.615 A on any face of a Cu crystal may therefore be expected to result in a strained Au overlayer. On a Cu(1 1 1) face, an in-plane strain of )11% will occur if the Au exactly matches the Cu spacing. If the Au layer density is to be maintained at the bulk value, this in-plane strain will produce an out-ofplane strain of +27%. This large degree of tetragonal strain should be readily measurable using MEIS by determining the angular shift of major blocking features since, because the aspect ratio of the unit cell changes with the amount of strain, so the blocking angles also change. It is perhaps worth noting that an isotropic expansion of a unit cell would maintain the aspect ratio and not result in a change in blocking angles. The strain can be depth pro®led though the layer by monitoring the position of blocking features at dierent scattered ion energies and hence dierent depths within the
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layer. Preliminary investigations were carried out on the growth of Au on Cu(1 1 1) to determine the thinnest Au layer for which the surface was unaffected by interfacial strain (it is desirable to use the thinnest layer to minimise the eect of straggling) and a Au layer thickness of 10 ML was determined to be suitable. Hence, an experiment to pro®le the strain in the Au requires a resolution substantially Using calculations better than 10 ML (24 A). similar to those shown in Fig. 2 but allowing for a (instead of 36 eV/A Au stopping power of 30 eV/A for Ni) we calculate a resolution at the surface of for 100 keV He ions using high resolution 1.5 A mode. This is the best achievable resolution which will degrade with depth due to straggling. For the scattering geometry used ( 1 1 0 in [0 0 1] out) we to calculate the straggling [30] at a depth of 24 A be 3.3 A. Thus the resolution will change from 1.5 at the surface to 3.6 A at a depth of 24 A, an A increase of a factor of 2.4. To keep the resolution more constant over this depth range we used a standard experimental setting to give a resolution (1.3 ML) at the surface which degrades to of 3 A depth, only a 50% change. 4.5 A (1.9 ML) at 24 A The dependence of Au strain upon depth below the Au surface is shown in Fig. 4. No strain has been measured in the ®rst two layers because of the large scattering signal from the surface peak. From Fig. 4 it is apparent that the strain is highly localised at a depth of about 8 ML. The straggling at
this depth is calculated to correspond to 1.5 ML. The ®t shown in Fig. 4 is the result of convolving a Gaussian of FWHM 1.5 ML with a strain delta function of thickness two layers centred at a depth of 8.5 ML. The implication of this ®t is that the strain exists only in two layers, presumably at the Au/Cu interface. This approach does not take into account the decay in strain towards the surface. The strain relief that is evident between layer eight and layer six probably occurs via dislocations which should, in principle, be detectable with MEIS. However, twinning (i.e. random interfacial stacking faults) in the Au layer means that it is not possible to pro®le any disorder relating to strain relief. The maximum measured strain is thus about +8%, much smaller than the calculated value of +27%. We suggest that the reason for the lower strain is due to inter-diusion at the interface of the two elements. Depth pro®les of each element reveal that Au has diused into the Cu to a depth and Cu is present throughout the of about 30 A Au layer. Thus the overlayer/substrate interface is not abrupt but is actually a graded Au/Cu alloy. The binary phase diagram of the Au±Cu system shows that stable Aux Cu1 x alloys exist over the composition range x 0±1 [31]. It is this graded alloy interface that allows the lattice size to vary slowly over many layers. Nonetheless, substantial tetragonal strain is observed in a localised twolayer region and, by deconvolution based on a knowledge of the magnitude of straggling, has been pro®led layer-by-layer. 4. Layer-by-layer composition ± Cu50 Pd50 (1 1 0)
Fig. 4. The measured strain versus depth below surface for a nominal 10 ML of Au grown on Cu(1 1 1). The solid curve is a ®t produced by convolving the calculated resolution with a strain delta function of thickness two layers.
The physical and chemical properties of bimetallic surfaces are of huge technological interest, particularly in the design of bimetallic catalysts for de-chlorination and in the investigation of surface corrosion [32±34]. It is well known that the composition of a bimetallic surface may dier strongly from that of the bulk material, these dierences being driven by a complex interplay of thermodynamic parameters such as relative surface energies, relative bonding strengths of each metallic component with adsorbed atoms and molecules, the tendency of the metal components to form
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ordered alloy phases and kinetic parameters such as diusion barriers. The quantitative analysis of the layer-by-layer composition of a bimetallic surface is a non-trivial analytical task. The shadowing and blocking capabilities of MEIS make this technique unique for obtaining layer-by-layer depth pro®le information. MEIS has been used by two groups to quantify the near surface composition of bimetallic surfaces. These studies investigated the Fe72 Cr28 (1 1 0) surface [35] and the Pt50 Ni50 (1 1 1) surface [36,37]. Whilst these studies are essentially extremely accurate investigations on well-de®ned surfaces, it has recently been shown that it is also possible to glean compositional information of the surface region of Cu50 Pd50 (1 1 0) in the presence of complex adsorbate layers [38]. For clarity we will con®ne our discussions to the Cu50 Pd50 (1 1 0) system. The approach used is to take a measurement in a particular geometry such that the ion yield is from the top layer only and then in a second geometry where the yield is from the top two layers only and so on for three and four layers. This is illustrated schematically in Fig. 5; the choice of out-going geometry will be discussed later. For each of the four geometries the Cu and Pd signal can be extracted since, due to
67
their dierent masses, they are well separated in energy. If a, b, c, d are the fractional amounts of Cu in layer 1, 2, 3, 4 respectively and w, x, y, z are the measured Cu signals in the 1-, 2-, 3-, and 4layer geometries, then, for the case of an ideal bulk terminated crystal at 0 K, we can write aCu 1:0wCu ;
1
aCu bCu 2:0xCu ; aCu bCu cCu 3:0yCu ;
2 3
aCu bCu cCu dCu 4:00zCu :
4
These four equations can easily be solved to give the four layer compositions, a, b, c, d. The geometries are chosen so that the ion beam is channelled in a particular incident geometry where the eect of shadowing minimises the illumination of deeper layers. Of course the shadowing can never be complete. Thermal vibrations and surface relaxations cause unwanted additional illumination of deeper layers. However, these eects can be lessened by utilising the eect of blocking ± the socalled double alignment technique [1] whereby ions are channelled in and also channelled out. The concept is illustrated schematically in Fig. 5 where the four scattering geometries are shown. In Fig. 5(a), an incident 1 0 1 geometry has been chosen to achieve the best single layer illumination but even so, simulations show that the total illumination in this geometry is actually 1.8 layers, with the majority of the extra illumination being in the second layer. However, using the double alignment geometry, 1 0 1 in [0 1 1] out, the eect of blocking substantially reduces the total illumination to 1.45 layers. The extra illumination alters Eqs. (1)±(4) to become 1:0aCu 0:42bCu 0:03cCu 0:0dCu 1:45wCu ; 5 1:0aCu 1:0bCu 0:2cCu 0:06dCu 2:26xCu ; 6 1:0aCu 1:0bCu 1:0cCu 0152dCu 3:15yCu ; 7 1:0aCu 1:0bCu 1:0cCu 1:0dCu 4:00zCu :
Fig. 5. Cross-sectional diagrams of the fcc(1 1 0) surface viewed in the h 1 1 2i azimuth: (a) the one layer illumination geometry 1 0 1 in [0 1 1] out, (b) the two layer geometry 1 1 0 in 1 3 4 out, (c) the three layer geometry 3 1 2 in [3 1 2] out and (d) the four layer geometry 2 1 1 in [1 2 1] out.
8
A numerical solution of these equations is considerably more dicult than the special case of Eqs. (1)±(4) but can be made trivial by employing determinants and a standard software package. The choice of incident ion in these experiments is
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important; generally speaking, He ions are better in this application than H ions for two reasons. Firstly, the larger shadow cone of the heavier ion helps to reduce the unwanted subsurface illumination and, secondly, the greater mass separation of the heavier ion allows one a greater range of scattering angles and hence more freedom in choosing the best geometries. There are a number of aspects of the simulations which warrant discussion: 1. The substrate composition ± because the substrate is bimetallic with random occupancy and a variable layer-by-layer composition, it is not possible to simulate the structure precisely. Instead we have used a crystal composed of a single element, Sr which has an atomic number (38) midway between Cu (29) and Pd (46). 2. Surface relaxation ± by simulations of a surface layer relaxation of magnitude 0±10% we see that its eect on the layer illumination is negligible. The reason for this is the large shadow cone of the 100 keV He ions; for the one layer geometry at the second atom in the chain, the sha and for Pd is dow cone radius for Cu is 0.29 A 0.37 A, both of which are greater than a 10% Blocking further relayer expansion (0.21 A). duces any residual eect. 3. Enhanced surface thermal vibrations ± these were found to have a strong eect on the layer illuminations. Literature values for the surface Debye temperature of Cu(1 1 0) range from 112 to 313 K depending on the orientation of measurement [39±45]. The only reported value for Cu50 Pd50 (1 1 0) is 200±250 K [46]. These values correspond to enhancements of the root-mean-
square vibrational amplitudes from 1.1 to 3.0, giving one layer yields from 1.1 to 1.8 layers. We used a surface enhancement of 100% giving a root-mean-square surface vibrational ampli This enhancement was allowed tude of 0.14 A. to decay exponentially into the bulk. 4. Possible adsorbate induced reconstructions ± effects of variable amounts of adsorbates on the surface were assessed by comparing the experimentally measured yield as a function of Cl coverage. The greatest eect is in the one layer geometry and even then no clear trend is seen. As a test of the assumptions made in the simulations, a comparison can be made between the normalised experimental yield and the predicted yields from simulations for each geometry. Fourteen separate measurements of the clean Cu50 Pd50 (1 1 0) surface were made and their average values were found to match the results of simulations to within 5%. There is, therefore, a good agreement between experiment and the theoretical model used in the simulations. When this technique is applied to the problem of segregation in the Cu50 Pd50 (1 1 0)/C2 H4 x Clx system, we see that hydrocarbon adsorption produces a measurable Pd enrichment in the surface layer, while Cu enrichment is produced by the presence of Cl(ads) and C(ads). Cu50 Pd50 (1 1 0)/ C2 H4 x Clx surfaces were prepared by adsorption of ethene or 1,2-dichloroethene (DCE) on clean Cu50 Pd50 (1 1 0) and subsequent annealing. Table 1 summarises the results. The composition of the ®rst layer of the clean surface is determined to be 65% Cu, in good agreement with the value of 70% quoted for similar preparation conditions [47±49].
Table 1 Cu content of the ®rst four layers of a Cu50 Pd50 (1 1 0) crystal after dierent surface treatments Experiment
First layer (% Cu)
Second layer (% Cu)
Clean Cu50 Pd50 (1 1 0) 10 ML ethene at 323 K 10 ML ethene ¯ashed to 383 K 10 ML ethene ¯ashed to 473 K 10 ML ethene ¯ashed to 573 K 10 ML 1,2-DCE at 323 K 10 ML DCE ¯ashed to 383K 10 ML DCE ¯ashed to 473 K 10 ML DCE ¯ashed to 573 K
65 38 57 53 55 47 72 30 76 28 71 30 72 29 88 7 Data suggest surface restructuring
Third layer (% Cu)
Fourth layer (% Cu)
71 45 61 69 62 66 62 76
50 80 66 71 81 78 73 65
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The values in Table 1 are plotted in Fig. 6(a) and (b). The 300 K data points in Fig. 6 are for the clean surface and serve to show how the clean surface changes when covered by the adsorbate. Referring to Fig. 6(a), from the known chemistry of ethene on Pd(1 1 0) [50] it appears that dissociative ethene adsorption occurs at 323 K on small Pd ensembles to yield an ethynyl (C2 H) species, pinning Pd at the surface and thereby increasing the ®rst layer Pd content at the expense of the second layer. This is evidenced in Fig. 6(a) by the reduction in ®rst layer Cu content and the increase in second layer content. After annealing at 573 K, the ®rst layer is more Cu rich since the ethnyl is known to decompose to C(ads). Thus, Cy Hx spe-
69
cies cause Pd enrichment at the surface whilst C(ads) causes Cu enrichment. In the case of DCE adsorption, from a knowledge of the chemistry of DCE on Cu(1 1 0) [51] and Pd(1 1 0) [52] and an appreciation of the anity between Cu and Cl, decomposition is preferred to desorption such that, at 473 K, C(ads) and Cl(ads) are the species present at the surface. The C(ads) and Cl(ads) strongly enhance the ®rst layer Cu content (88% Cu) at the expense of the second layer (7% Cu). The anomalously high MEIS background seen upon annealing this surface to 573 K suggests a restructuring of the surface, possibly caused by the presence of CuClx islands, an eect which has also been seen for Cl2 adsorption on Ag(1 1 0) [53]. It is apparent, therefore, that the judicious use of double alignment in several scattering geometries can provide a clear-cut layer-by-layer composition pro®le of a bimetallic crystal surface. 5. Buried layer structure determination ± twodimensional rare earth silicides and germanides on (1 1 1) substrates
Fig. 6. Variation in the composition of the topmost four layers of the Cu50 Pd50 (1 1 0) surface as a function of annealing temperature following the adsorption of (a) 10 ML ethene and (b) 10 ML 1,2-DCE at 323 K. The 300 K data points represent the clean surface.
This ®nal example of achieving monolayer resolution relates more to sensitivity than instrument resolution. Because of the high degree of sensitivity of ion scattering to a heavy impurity in a light substrate, the impurity signal can easily be resolved from that of the substrate. For clarity we will refer to the heavy impurity as HI and the light substrate as LS. In the case of a monolayer of HI on LS, if the mass dierence is suciently large, the HI signal occurs at a much higher energy than the LS signal and the two can easily be resolved. This may still be the case even if the inherent energy resolution of the analysing instrument is not sucient to resolve the energy lost by ions traversing the monolayer. However, in MEIS studies of systems comprising of a monolayer of HI on LS there is generally no structural information to be gained from the signal from the HI. Because the HI lies on the LS surface and is contained within a single layer, there will be no HI±HI or HI±LS blocking and hence no detail in the HI blocking curves, e.g. [23]. However, LS±HI blocking occurs in an unusual family of systems where a HI
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monolayer readily adopts sites beneath LS atoms. These are the 2D rare earth (RE) silicides and germanides on (1 1 1) substrates which are formed when 1 ML of RE is deposited on a clean Si(1 1 1) or Ge(1 1 1) substrate and annealed. In these cases, the surface converts into a well-ordered semi-metallic 2D layer of stoichiometric RESi2 or REGe2 . The RE forms a hexagonal monolayer accommodated underneath a buckled Si or Ge top layer where the buckled top layer is rotated by 180° in the surface plane (Fig. 7). Such 2D alloys are known to occur for Tm, Er, Ho, Dy and Gd on Si(1 1 1) [54,55] and Ho and Dy on Ge(1 1 1) [26,55]. The eect is believed to be limited to those REs which are trivalent at a coverage of 1 ML. Because the RE monolayer is buried beneath an ordered Si or Ge bilayer, blocking features occur in the RE signal and can be used for structural determination. Fig. 8 shows blocking curves for the case of Ge(1 1 1) 1 1±Dy (Ge provides a better illustration than Si as the blocking minima are stronger due to the higher Z of the Ge). The yield calibration procedure has been described earlier. The curves in Fig. 8 have been used to determine the position of the surface Ge atoms with respect to the buried Dy atoms and the results are presented in Table 2. The atom positions are de®ned in Fig. 7. For 100 keV He ions in the Si(1 1 1) 1 1± Dy system, a calculation of the elastic scattering reveals that (for the geometries shown in Fig. 8) there is a separation of 18 keV between Dy and Si
Fig. 8. Experimental blocking curves (data points) obtained from a Ge(1 1 1) 1 1±Dy surface using 100 keV H ions together with the corresponding best-®t simulations (solid curves): (a) the 1 1 0/[0 0 1] scattering geometry, (b) the normal/ 1 1 1 scattering geometry. The insets show the origin of the major blocking dips.
signals. Thus, an instrument with a resolution of 15 kV FWHM at 100 kV could be successfully employed on the Si(1 1 1) 1 1±Dy system because the energy separation of (and thus the sensitivity to) the Dy is so great. A similar calculation shows that the structure determination of this monolayer system could be carried out using 2 MeV He ions and a conventional RBS instrument with a Si barrier detector if this could be scanned through a sucient in-plane angular range. In the case of these 2D RE (1 1 1) silicides and germanides, monolayer resolution could be
Table 2 The outermost layer spacing and Ge±Dy bond lengths for the Ge(1 1 1)1 1±Dy surface Bond length (A) Vertical distance (A) Fig. 7. A cross-sectional diagram of the Ge (1 1 1) surface viewed in the h1 1 2i azimuth illustrating the structure of the Ge(1 1 1) 1 1±Dy 2D germanide.
Ge(1)±Ge(2) Ge(2)±Dy
1:08 0:02 1:71 0:02
2:55 0:03 2:87 0:03
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considered to be inherent when applying any ion scattering technique to these systems. 6. Conclusions Monolayer resolution in MEIS might generally be associated solely with the resolving power of the ion energy analyser. There are, however, several MEIS techniques which can be used to achieve monolayer resolution which do not necessarily rely on an analyser having a resolution in energy comparable to that lost by an ion traversing a single layer. In the classic MEIS surface structure experiments, the signal from a number of layers is amalgamated giving an apparent resolution of perhaps six layers. However, a detailed knowledge of the bulk crystal structure combined with best ®t simulations allow the layer-by-layer structure to be determined from the simulation input data. Even where monolayer resolution is readily available from the ion energy analyser, the straggling associated with penetration of buried layers will ultimately dominate the eective resolution. In this case an ability to calculate the straggling is necessary to perform a deconvolution. If the straggling is only of the order of a few layers we suggest that an accurate deconvolution can be carried out to yield meaningful layer-by-layer information. The layer-by-layer compositional analysis of an alloy surface is an extreme example of an application requiring monolayer resolution. The unique shadowing and blocking capabilities of MEIS are well suited to this problem using the concept of double alignment. It is perhaps surprising that the compositional information can be determined even though the surface is covered by complex and unstable adsorbates. We ®nd that the adsorbateinduced compositional changes are quite large and any uncertainties in the modelling of the surface produce only minor variations in the composition values calculated. In the unusual case of an ordered buried layer of heavy impurity covered by an ordered layer of light substrate atoms, the application of an ion scattering technique confers monolayer resolution almost inherently. The large mass separation of the REs
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compared to the substrate in the 2D RE-(1 1 1) silicide and germanide systems ensure a high degree of sensitivity to the RE. In these systems, monolayer resolution can be achieved without use of an ion energy analyser with high energy resolution. Four examples of very dierent approaches to achieving monolayer resolution in MEIS have been described and illustrated. These examples are not meant to be an exhaustive compilation of all the possible ways to achieve monolayer resolution in MEIS ± other approaches exist and yet more may be discovered.
Acknowledgements The authors would like to acknowledge Kevin Connell and Brian Blackwell for technical help, the FOM Institute for the provision of the VEGAS code and the Engineering and Physical Science Research Council for funding this research.
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