physica status solidi, 7 April 9 Non-basal plane surfaces: Anisotropic structures and low-dimensional electron systems Ulrich Starke *, Max-Planck-Institut für Festkörperforschung, Heisenbergstrasse, D-7569 Stuttgart, Germany Received XXXX, revised XXXX, accepted XXXX Published online XXXX PACS Corresponding author: e-mail u.starke@fkf.mpg.de, Fax: +49-7-68999 The polytype dependent stacking sequence in is exposed on its non-basal plane surfaces, and thus complex and anisotropic surface reconstructions can be expected. Detailed investigations of the atomic and electronic structure of a-plane ( ) surfaces and diagonal cuts, namely ( ) and ( ) surfaces of 4H- are reviewed. After hydrogen etching the surfaces show large, flat terraces. Preparation in ultra-high vacuum (UHV) leads to the development of well ordered surface phases. On 4H-( ) three unique and distinguishable ( ) phases can be identified by monitoring the LEED spot intensities. On 4H-( ) surfaces three well ordered phases with different periodicity appear. The Si rich ( ) phase is characterized by an ordered array of Si-adatom chains which host an electronic surface state that is confined within the chains. A c( ) phase exists at a surface composition close to bulk. At high temperatures a ( ) phase develops which is carbon terminated with a graphite-like bond configuration. ( ) A C A () a 3 c ( ) a a 4H- [ ] [ ] 4 3 5 5 Orientation of the ( ), ( ) and ( ) planes within the 4H- structure. STM micrograph and atomic model of the ( ) reconstruction on 4H-( ). 6 Introduction Silicon carbide () is a wide band gap semiconductor particularly useful for high power and high frequency devices. For many years problems with material properties and high quality growth were hampering a widespread application of the material in devices. In the effort to resolve these important issues, the basal-plane surfaces of have been investigated intensively aiming towards improving growth conditions and the deposition of homoepitaxial layers [, ]. surfaces are also important as substrate for heteroepitaxial growth such as in the case of GaN on [3]. Furthermore, quite recently was discovered as a perfect support for large scale epitaxial graphene layers [4 ]. In recent years surfaces have also attracted a fundamental interest with respect to their own electronic structure. A remarkable property of surfaces is the strong Coulomb interaction between dangling bonds which can lead to electron correlation effects on many reconstruction phases []. As recently suggested such electron correlation would be particularly interesting for one-dimensional electronic states [, 3]. The corresponding situation can be achieved naturally when strongly anisotropic atomic arrangements are present at the surface.
Crystal structure and bulk truncated surface termination The orientation of the non-basal plane sur- faces within the 4H- bulk is depicted in Figure. In panel, the bulk structure is schematically displayed by a hexagonal prism which actually represents three 4H- bulk unit cells. The basal plane is spanned by the three vectors a, a and a3. c defines the c-axis direction <>. The ( ) plane is sketched in the back of the drawing (red/dark). It runs perpendicular to a3. The ( ) plane is shown on the left side (green/grey), and proceeds parallel to a3. The ( ) plane (grey) is tilted with respect to the caxis and runs diagonally through the hexagon prism. The atomic model of the 4H- crystal in panel displays the alternating layer stacking (ACA) in the hexagonal 4H polytype. A bold black line indicates the bonds parallel to the projection plane and thus depicts the zig-zag structure of the layer stacking with the orientation changing after slabs of two bilayers in 4H-. The view in the A () c ) C ( () In this respect, non-basal plane surfaces of appear to be particularly promising templates since they directly exhibit the long period of the bulk polytype stacking in the direction of the c-axis. The ( ) - or a-plane and ( ) - or m-plane - surfaces of have such a strongly anisotropic atomic structure with a long period in the c-axis direction, and thus promise a good perspective for an intrinsic development of nanostructures. Specifically in a plane orientation situated diagonally in the bulk unit cell of which is tilted by 6 with respect to the basal plane (see below) the polytype stacking sequence induces an anisotropic bulk truncated surface structure consisting of alternating stripes of different bond configuration [4]. Such a stripe configuration is present on both possible surfaces in this orientation, which are the ( ) and the ( ) surfaces in the case of 4H-. We should note, that a-plane substrates have also been used recently for bulk material growth in view of their potential to reduce the incorporation of defects such as micropipes. Incidently, a surface which is very similar to the tilted surface orientation mentioned above namely the (33 8) surface, has been used to grow high quality bulk material [5]. We note in addition, that the ( ) plane appears to be naturally stable in view of the fact that it was initially observed in pores developing in after electrochemical etching [6]. esides other pore types found, this etching process can produce a triangular shaped pore channel with surfaces inclined by about 6 with respect to the basal plane, which - in the case of 4H- - exactly correspond to the ( )-surface orientation. So, non-basal plane surfaces represent a strong potential for a variety of applications. Yet, there is little knowledge accumulated so far about their microscopic structure [, 4, 7 7]. The present paper gives an overview of recent investigations of the atomic and electronic structure of ( ) surfaces of 4H-, as well as diagonal cuts of this polytype, namely the ( ) and ( ) surfaces [8 33]. U. Starke: Non-basal plane surfaces A a3 a a 4H- <> Figure Schematic orientation of the ( ), ( ) and ( ) planes within the 4H- structure shown by lattice vector construction. Atomic model of 4H- in side-view perspective along the [ ] direction with bonds within the ( ) plane indicated by a black line and the orientation of the ( ) plane as grey line. Large (green) atoms represent Si, small (black) atoms C. (adapted from [4]) Figure Atomic model of the 4H-( ) surface: Top view, showing the first layer with perspective (colored) balls, the second layer as flat (grey) circles. The ( ) unit cell is indicated by a (red) rectangle, a glide symmetry plane by a vertical (blue) line. Top view, full model, again with unit cell and glide symmetry plane indicated. Side view along the [] direction. Large (green) atoms represent Si, small (blue) atoms C. (adapted from [8]) [ ] projection direction directly shows the ( ) surface. A grey diagonal line indicates the situation of the ( ) plane. For two layers this plane runs along the line of -bonds lying within the a-plane (bold black lines). In the second two layers of the 4H-unit cell the ( ) plane runs almost perpendicular to these black bonds. The bulk truncated 4H-( ) surface contains chains of the above mentioned zig-zag bond trains as displayed in Figure. The corresponding chain in the topmost layer is indicated by the perspectively drawn (Si = green/grey, C = blue/dark) atoms. With a parallel shift corresponding to the tetrahedral bond configuration the next chain follows in the second layer, as indicated by the grey disks (Si = grey/large, C = light/small). The bulk truncated unit cell is outlined by the (red/dark) rectangle. The
pss header will be provided by the publisher 3 Surface unit cell 4H-bulk unit cell [] C Si [] [ ] [ ] [ ] Figure 3 ulk truncated, C terminated 4H-( ) surface, with surface unit cell and 4H-bulk unit cell indicated. C terminated ( ) surface shown in top view. The ( ) unit cell (.4 Å 3.8 Å size) as well as the separation in cubic () type (crosshatched) and ( ) basal plane type (hatched) patches and crystal directions are indicated. Large (green/light) atoms represent Si, small (blue/dark) atoms C. vertical (blue/light) line indicates a glide (mirror) plane, which corresponds to a glide symmetry element present in the bulk structure. The whole structure is visible in the full perspective top view model in Figure. Panel shows a side view of the surface. In this projection along the c-axis it is immediately apparent that the bulk truncated surface is non-polar, or in other words, the Si and C atoms in the zig-zag bond chain mentioned above form the topmost surface layer. Due to this fact, for this bulk truncated configuration there is only one surface termination possible. Of course, for the real surface relaxations, reconstructions or stoichiometric deviations are to be expected. In ( ) orientation the bulk unit cell consists of six atomic layers in slabs of three adjacent Si and three adjacent C layers. Accordingly, already the bulk truncated crystal can have six different surface terminations ranging from fully Si terminated to fully C terminated, depending on which of the layers is terminating the surface. This scenario holds for both the ( ) and the ( ) surfaces. In a side view along the [ ] direction Figure 3 displays a bulk truncated surface with a full carbon layer chosen as surface termination since this configuration would yield the lowest number of dangling bonds. Panel shows a top view of this surface with the ( ) unit cell indicated by the red/dark rectangle (green/grey background). One can immediately see the anisotropic nature of the surface with stripes of two different atomic configurations within each unit cell (hatched, crosshatched) which stems from the two slabs with opposite layer stacking in the 4H- unit cell. The topmost carbon atoms in one stripe are two-fold coordinated (crosshatched), while the others are three-fold coordinated (hatched). The two configurations are structurally similar to a cubic 3C-() and a basal-plane ( ) surface, respectively. Well defined surfaces of this orientation can be prepared by hydrogen etching which in particular yields large flat terraces [4]. Thus, the two regions are characterized by single and double dangling bonds, respectively. Of course, this holds for the case of a bulk truncated surface. Again, relaxations, reconstructions or stoichiometric deviations will be present on the real surfaces. 3 Experimental procedures Mechanically polished wafers of different orientations were cut in small sample pieces. The samples were hydrogen etched in order to remove polishing damage and generate atomically smooth surfaces (see next section). The surface morphology achieved by the hydrogen treatment was monitored using atomic force microscopy (AFM). Subsequently, the samples were introduced into different ultra-high vacuum (UHV) chambers for the different analysis methods. In UHV, the samples could be annealed by electron bombardment or resistive heating in the different chambers. The sample temperatures were measured with optical and infra-red pyrometry. For the control of the surface stoichiometry an electron beam heated Si evaporator was available. The Si flux was controlled by a quartz microbalance and typically kept at about.5 ML/min. Surface analysis experiments were performed in four different chambers all equipped with a sample introduction stage. The chamber for scanning tunneling microscopy (STM) experiments was equipped with a reverse view low-energy electron diffraction (LEED) optics and a 5 spherical sector analyser for Auger electron spectroscopy (AES). The samples could be transferred in situ into the STM [34]. The core level spectroscopy experiments were carried out at the MAX synchrotron radiation laboratory using beam line I3 [35] whose end station [36] is equipped with a hemispherical analyzer and LEED. This setup was also used for normal emission valence band (V) spectra. The dispersion of the surface state bands with respect to parallel momentum was studied in our home laboratory by angle-resolved photoemission spectroscopy (ARPES) using UV light of 4.8 ev from a He discharge lamp. Additional angle-resolved valence band spectroscopy measurements were carried out at the MAX-lab at beamline 33 with a goniometer based angle resolving electron analyser [37] All chambers were
4 U. Starke: Non-basal plane surfaces AES data dn(e)/de Si C O 4 6 Energy (ev) Figure 4 Hydrogen etched 4H-( ): AFM image of 8.3 µm 8.3 µm area, measured in contact mode. The full z-scale (black to white) corresponds to 4 nm. A plane alignment and a three-point smoothing was applied for the presentation. AFM line profile along the blue line indicated in panel, showing a step height of 5 Å. AES spectra of the freshly polished surface (i), after RCA [4] cleaning (ii), after hydrogen etching (iii) and after additional outgassing (iv). The spectra are displayed as differentiated AES-signal dn(e)/de (adapted from [8]). Y[µm] Z[nm].5.5 < >.5 X[µm].5..4.5 X[µm].5.4...4 Y[µm].5.5 < >.5 X[µm] 4.5.5.5.5 X[µm] Figure 5 AFM images obtained after hydrogen etching of 4H-( ) and 4H-( ). Line profiles parallel and perpendicular to the [ ] direction are shown below the respective AFM images (adapted from [4]). equipped with LEED which was used quantitatively to compare the quality and state of surface preparations and surface reconstructions in the different chambers. 6 4 Z[nm] (i) (ii) (iii) (iv) 4 Hydrogen etching It has been established on basal plane samples that hydrogen etching efficiently removes polishing damage from the surface. With the proper process parameters applied, the technique leads to atomically flat surfaces with steps of one unit cell hight, that is e.g. nm for 4H- [38,39]. Hydrogen etching can also be used to prepare atomically smooth non-basal plane surface of. On ( ), hydrogen etching at a substrate temperature of 5 C removes the scratches remaining from the polishing procedure [8]. The process provides a smooth, flat surface with steps in the range of a few Å to nm, i.e. steps of one or a few atomic heights as indicated by the AFM images in Figure 4. Panel displays the AFM line profile taken along the (blue) line in panel. A step of 5 Å height is clearly resolved. Atomically flat terraces with a size in the µm regime develop. A quantitative evaluation of a raw AFM image covering a 6 µm 6 µm area yields a roughness of root mean square (RMS) =.44 nm. The Auger spectra (differentiated AES-signal dn(e)/de) shown in Figure 4 demonstrate how the surface is cleaned during the hydrogen etching process. The significant oxygen signal present for the original polished (curve i) as well as an RCA [4] cleaned (curve ii) sample is largely removed by the hydrogen treatment (curve iii). The Si peak transforms from a oxidic form to the typical type shape. Carbon which is largely absent for the as introduced samples (i and ii) assumes a typical carbidic peak form. Yet, annealing the sample to a temperature of around 5 C is necessary before the residual oxygen signal disappears entirely, whereas the Si and C peaks are largely unaffected by the heating. The polished and the RCA cleaned sample show no LEED pattern after introduction into the analysis chamber. The hydrogen etched surface after loading into UHV immediately exhibits a ( ) LEED pattern with some background intensity visible due to ambient contamination. After annealing the background is to a large extent reduced. Also for the ( ) and ( ) surfaces, the residual polishing scratches can be removed by hydrogen etching (5 C). The AFM images in Figure 5 reveal step arrays on both surfaces that are aligned parallel and perpendicular to the [ ]-direction and result from the residual mis-orientation (of less than.5 degrees) of the samples. However, the step heights and terrace widths are drastically different for the two surface orientations as demonstrated by the line profiles plotted in the figure. On ( ), a stripe pattern morphology develops with 5 nm µm terraces, as shown in panel. The line profiles displayed underneath depict that the steps perpendicular to the [ ]-direction have a typical height of.4 nm, whereas the steps parallel to the [ ]-direction are larger with a typical height of about.5 nm. On ( ), on the other hand, large terraces (typically.5 µm µm) are found with step heights in the - 5 nm regime [line profiles below the AFM image in panel ]. oth surface orientations display a sharp ( )-LEED pattern again with some residual background from atmospheric contamination [4]. On ( ) in some cases a very faint c(4 ) superstructure was observed [3]. AES indicates a considerable amount of oxygen on the surface after hydrogen etching [4]. This was corroborated by core level spectroscopy for the ( ) surface where the Si p spectrum for the H etched samples shows strong oxidic Si components and the C s spectrum a strong component characteristic for hydrocarbons [3]. Seyller et al. reported hydrogen saturation of the surface dangling bonds after hydrogen etching for ( ) as well as ( ) surfaces based on synchrotron core level photoemission spectroscopy and fourier-transform infrared absorption spectroscopy []. In
pss header will be provided by the publisher 5 both cases also a pure ( ) periodicity of the surface was found. 5 A-plane and m-plane surfaces On basal-plane surfaces a rich surface phase diagram is accessible by varying the surface composition. Si rich surfaces are prepared by Si deposition - in most cases under simultaneous sample heating. So, e.g. on () a Si rich (3 3) reconstruction appears [4]. y further annealing of the sample successively less Si rich reconstructions can be prepared [4] until carbon rich surfaces are reached [, 8, 43]. The same technique can be used on non-basal plane surfaces. The ( ) LEED pattern observed on ( ) after hydrogen etching as outlined in the previous section corresponds to the bulk truncated unit cell shown in Figure. Different from the basal plane surfaces, however, on ( ) the periodicity remains ( ) after Si deposition and heating. Figure 6 shows LEED patterns for different annealing steps subsequent to Si deposition at 85 C. With increasing temperature the ( ) periodic spots become brighter accompanied by a reduction of the background intensity. Note that the glide symmetry plane of the bulk structure, as indicated in Figure is present on the surface. It results in a systematic absence of (,n-) spots in all LEED patterns. While the surface remains in a ( ) periodicity throughout the annealing process, the spot intensities change at certain temperatures. A quantitative inspection of the LEED spots reveals the presence of three unique phases. Figure 6 shows the spot intensities in form of I(E)-spectra, averaged from the (,) and its symmetry equivalent (,) spot. The spots are marked by circles in Figure 6. A transformation of spectral features from low to high temperatures is clearly visible in the plotted intensity curves. These changes in the LEED intensities correspond to structural changes on the surface. It can be seen, that the transformation occurs in two steps as depicted by the spectra plotted with dashed lines, i.e. three different types of spectra (plotted in different colors) and thus three different surface phases are observed, as differentiated: The first phase exists straight after the Si deposition at 85 C, the second phase develops at about 96 C and the final transformation occurs at about C. The changes in the spot intensity spectra can be quantitatively evaluated by determining the Pendry R-factor [44] between spectra acquired after subsequent annealing steps. The R- factor maxima as shown in Figure 6 and emphasized by the (spline) interpolated line indicate the temperatures of the phase transformation. These temperature values are corroborated by the analysis of several different spots in the LEED pattern. However, from a detailed inspection of the intensity data it cannot be ruled out that the intermediate phase may in fact represent a phase coexistence regime between the low and high temperature phase. A quantitative analysis of the structure of these phases by means of LEED calculations remains to be performed. Earlier theoretical investigations of the ( ) surface predicted a ( ) phase with a Si adatom configuration [9]. However, recent experimental work reported a c(4 ), a c( ) and a ( ) phase for this surface [7]. A strongly dispersing surface state was found for the c(4 ) phase only. Structural models for the reconstruction phases on m-plane surfaces on have not been established. 6 Surface phases on ( ) As noted before, immediately after loading into UHV the H-etched ( ) surface typically displayed a ( ) LEED pattern with very faint c(4 ) contributions in some cases, cf. Figure 7. Upon annealing at around 9 C the residual oxygen contamination is removed and a pure ( ) phase develops. The shape of the ( ) unit cell is indicated in Figure 7 (left). Well-ordered superstructure phases can be prepared by Si deposition with the substrate heated to about 7 C and subsequent annealing. At 83-9 C annealing temperature a Si enriched ( ) phase develops with a LEED pattern shown in Figure 7. The unit cell is depicted in Figure 7 (right). Further annealing around 95-975 C results in a c( ) order, cf. Figure 7(d) with nearly stoichiometric surface composition (see Figure 7 (bottom) for the unit cell shape). Heating to about 5 C yields a surface with ( ) periodicity again, Figure 7(e), however this time with a carbon rich composition. The composition of the different surface phases was determined from core level spectra and AES. 6. The ( ) phase The surface is of Si rich composition as indicated by the core level spectra shown in Figure 8-(d). The C s spectrum contains only a bulk component at a binding energy of 8.6 ev independent from the photon energy, i.e. surface sensitivity (panel a,b). Fitting the Si p core level reveals two prominent surface shifted components (labeled S and S) at -.9 and -.6 ev relative to the bulk peak at a binding energy of.9 ev. This indicates that the reconstruction is entirely restricted to the Si adlayers. Differences in the relative intensity ratio of the S and S components compared to the bulk component for different photon energies, cf. panels (c,d), indicate that the S component belongs to Si adatoms located in the outermost surface layer, while the S peak originates from Si atoms located at an interface between the bulk and the outermost Si atoms (S). The thickness of the topmost Si layers (S and S components) can be estimated to about 3 Å [3]. An empty state STM image of the surface shows single protrusions per ( ) unit cell as indicated by the rectangle in Figure 8(e). They form rows with a period within the row of 6 Å and obviously correspond to atomic chains of Si adatoms. The chain separation in [ ] direction of.. nm corresponds to the surface unit cell length. The chains are arranged in a regular pattern which corresponds to the bilayer stacking in the 4H- polytype as it is projected onto the ( ) surface. Occasional defects visible in the image are concentrated in the channels be-
6 U. Starke: Non-basal plane surfaces 97eV 89 C 97 C 5 C ev Spot intensities (arb. units, vertical shifted) C 9 C 6 C 5 C 3 C C C 98 C 96 C 95 C 93 C 9 C 9 C 89 C Si-dep. 5 5 5 Energy (ev) Pendry R-factor,35,3,5,,5,,5, 9 95 5 Temperature ( C) Figure 6 LEED patterns of the 4H-( ) surface after Si deposition and annealing at different temperatures. (,)- and (,)-spot marked by white circles. Intensity spectra of the averaged (,) and (,) spots during sequential annealing after Si deposition with annealing temperatures indicated. Spectra of different shape (different phases) are plotted with different colors, black spectra indicate the phase transformation temperature. Pendry R-factor between spectra acquired after subsequent annealing steps (red squares) with spline interpolation (blue line) indicating different structural phases. tween the adatom rows while the rows are mostly well ordered. In Figure 8(f) an STM image of the filled states is displayed which reveals further details of the unit cell. Here again the well ordered rows can be identified, but an additional high charge density is resolved in the middle between the adatom rows. In average its z-position appears to be slightly below that of the adatom rows in most cases, although a variation is observed that indicates some disorder in the regions between the rows. Noteworthy, the lateral position of this charge density is asymmetric. As depicted by the arrow in the inset of Figure 8(f) it is neither on the line nor in the hollows between the adatoms but shifted from the unit cell center by about.5 Å along the [ ] direction. As indicated in the main part of Figure 8(f) by the unit cells marked in solid and dotted lines, two domains are formed where this shift points in opposite direction. It should be noted that the two images were acquired from different preparations, which can explain the more inhomogeneous appearance of the surface in Figure 8(f). A comparison of the z-scales indicated in the Figure shows that the z-height differences are much smaller for empty states than for filled states (This was verified for images from several sample preparations). Interestingly, the corrugation within the unit cell is quite anisotropic. Along the adatom chains the corrugation is smaller than perpendicular. This anisotropic corrugation certainly reflects the different distance of the adatoms in [ ] and [ ] direction. Yet, also in the filled state image, where the protrusions are more or less equally spaced, this anisotropy is visible, which indicates a delocalization of the electronic density in the direction of the chains. Noteworthy, in the channels between the adatom rows (also [ ] direction) the corrugation is even shallower. Normal emission V spectra at different photon energies indicate the presence of a strong surface state at about.8 ev below the Fermi level (E F ), since its binding energy does not change with photon energy [3]. The dispersion of the state parallel to the surface (energy vs. k ) was determined by angle resolved measurements with photon of 4.8 ev energy from a He discharge lamp. A small dispersion of. ev was found when varying the emission angle along the [ ] direction as shown in Figure 9. Figure 9 shows in contrast that along the [ ] direction the binding energy of the state remains constant. Thus it appears that the state is one-dimensionally delocalized but confined within the atomic chains. However, the dispersion band remains below E F, so that the chains resemble semiconducting nanowires in an ordered array. A full dispersion diagram was acquired using synchrotron radiation of 6 ev photon energy along high symmetry directions of the surface Z and is plotted in Figure 9. The V maximum is found to be.5 ev below E F. Also here,. ev dispersion of the surface state is seen along Γ- X. No dispersion exists along Γ- X. ased on the experimental results a tentative model was derived for the ( ) reconstruction on ( ) which is displayed in Figure. In accordance with the Si enrichment drawn from the core level intensities the surface must be covered by two additional Si layers. The atoms of the first Si layer (green atoms, nr. 5) saturate the carbon dangling bonds and could be viewed as natural continuation of the bulk stacking. However, their lateral positions must be severely displaced and the arrangement of the next Si layer must be even more complex due to the strain resulting from the % longer bond between two Si atoms than between Si and C in. On the basal
pss header will be provided by the publisher 7 Cs 33 ev Sip 4 ev S S 45 ev (d) 33 ev [ ] [ ] S S 86 84 8 8 inding energy (ev) 3 99 97 inding energy (ev) (e) z =.53 Å z =.46 Å (f) (d) E=48 ev (e) E=48 ev [ ] [ ] E=48 ev E=48 ev Figure 7 a) Top view of the bulk truncated 4H-( ) surface with the ( ), ( ) and c( ) unit cells indicated (left, right and bottom). Large balls represent Si, small balls C atoms. (be) LEED patterns of 4H-( ) acquired at 48 ev electron energy after hydrogen etching, for the ( ) phase, for the c( ) phase (d) and for the high temperature ( ) phase (e). (adapted from [3]) Figure 8 and C s core level spectra recorded at photon energies of 33 and 45 ev. and (d) Si p spectra recorded at photon energies of 4 and 33 ev, respectively with fit components indicated by the solid curves. (e) STM micrograph of the ( ) reconstruction on 4H-( ) taken for empty states with - V tip bias and.3 na tunneling current. The surface unit cell and the low Miller index crystal directions are indicated (5 nm nm). (f) Filled state STM image (9 nm 9 nm) with + V tip bias. On the right side domains with a different position of the second protrusion are indicated by solid and dashed unit cells, the inset depicts its lateral shift. The full z-scale as indicated by the scale bar in between is.53 Å for (e) and.46 Å for (f), respectively. (adapted from [3]) plane type stripes the second layer Si atoms (red atoms, nr. and 3) occupy H3 positions [45], while on the cubic type stripes they form double bonded Si dimers in a Si bridging configuration (red atoms, nr. 4). Finally, the basal plane type stripes are terminated by Si adatoms, again in H3 positions (purple atoms, nr.) in ( ) periodicity. The model contains only two dangling bonds per unit cell, one on the adatom and one on the free Si atom in the second layer (red atoms, nr. 3). It should be noted that on the basal plane type stripes, the second layer Si and the Si adatoms may occupy two different three-fold coordinated positions, namely the so-called T4 and H3 site [45]. The H3 position was chosen in both cases to minimize the strain in the Si layer. Moreover, if the T4 position were chosen for the topmost Si adatom, the Si dimers on the cubic stripes would be positioned more asymmetric with respect to the topmost adatom chains, which is clearly in contradiction to the filled state STM image shown in Figure 8(f). Consistent with the positions of the additional charge in the STM data, the dimers are shifted with respect to the center be- tween the adatoms in [ ] direction by /4 unit vector. This configuration in addition allows for the presence of two domains with the dimer located either above or below that center. The π-orbitals of the dimer may also account for the shallower corrugation in the channels between the rows, although of course an effect of the deeper tunneling position cannot be ruled out. 6. The c( ) phase The stoichiometry of the c( ) phase as indicated by AES is relatively close to a bulk composition. This is corroborated by core level data. Normal emission C s and Si p core level spectra collected at three different photon energies are displayed in Figure. Components obtained from peak fitting are shown for the bottom spectra. The C s core level is composed of the bulk peak located at a binding energy of 8.6 ev and a shifted component (C) at about.7 ev higher binding energy which can be assigned to atoms in the outermost carbon layer, cf. Figure. A very small carbon component can be detected (C with a shift
8 U. Starke: Non-basal plane surfaces [ ] (-º) X (º) (x)..6.e.(ev) M X X [ ] (-º) X X (x) (+º)..6.E.(eV).5..5..5 E (ev) M E F VM X X M Figure 9 (a,b) Angle resolved valence band spectra for the 4H- ( )-( ) surface recorded at emission angles from - to + along the [ ] direction with step and from - to + along the [ ] direction with steps, respectively using a photon energy of 4.8 ev from a He lamp. E(k ) dispersion curves. The blue and red symbols are experimental data for the surface state and the bulk bands, respectively, taken from ARPES measurements with a photon energy of 6 ev using synchrotron radiation. E F and the V maximum (VM) are indicated. The surface rillouin zone is shown at the top of panel. (adapted from [3] and [3]) Intensity (arb. units) Cs C C h (ev) 6 45 33 Sip h (ev) 85 83 8 3 99 inding energy (ev) 6 33 SS SS 4 [ ] A [ ] Figure Cs and Sip core level spectra recorded from the 4H-( )-c( ) surface at different photon energies. Individual components and results obtained after fitting are shown (as lines) for the lowest photon energies. STM micrograph of the c( ) reconstruction on 4H-( ) taken with -.5 V tip bias and.3 na tunneling current. Image size 5 nm 5 nm. The surface unit cell and the low Miller index crystal directions are indicated. A,,C indicate lines along the crystal directions [ ], [ 4 3] and [ ] where the corrugation is discussed in the text. (adapted from [3]) top view: C (x) unit cell top view 3 4 5 4 3 4 3 4 5 5 side view 4 3 5 5 side view: Outermost C-atom Outermost Si-atom Si adatom Dangling bond [ ] [ ] Figure Top and side view of the structural model suggested for the c( ) reconstructed 4H-( ) surface with the primitive unit cell indicated. (adapted from [3]) 5 3 4 Si adatom st Si layer nd Si layer 6 bulk Si atom bulk C atom [ ] 6 [ ] Figure Top and side view of the structural model suggested for the ( ) reconstructed 4H-( ) surface with the unit cell indicated. In the top view only the topmost carbon layer of the bulk is displayed. In the side view the topmost and one carbon layer are shown. The ( ) unit cell (solid rectangle) and the basal plane and cubic type stripes (light and dark shaded/yellow and pink) are indicated. (adapted from [3]) of about.8 ev) in the most surface sensitive spectra at 33 ev photon energy. It was assigned to graphite like carbon resulting from annealing at a slightly too high temperature, corresponding to a beginning transformation to the graphitic ( ) phase. The Si p spectra contain three components, cf. Figure. The major contribution at.9 ev corresponds to bulk. Two surface shifted components (labeled SS and SS) have binding energies of.3 and 99.6 ev, respectively. The SS component is found to be more surface sensitive and therefore corresponds to the outermost Si position, e.g. adatoms. The SS peak should correspond to - at least parts of - the Si atoms of the topmost layer. From the intensity ratios of the different peaks, the SS type Si atoms should have a concentration lower than a third monolayer. The SS component represents less than a full layer of Si atoms. A
pss header will be provided by the publisher 9 comparison of the entire Si p and C s peaks indicates a more or less stoichiometric surface composition. The LEED pattern of the c( ) phase on ( ) is shown in Figure 7(d) together with the centered c( ) and the primitive unit cell. This ( primitive ) unit cell is correctly denoted by the matrix. The shape of the primitive unit cell is drawn at the bottom of Figure 7. We see that for this phase symmetry equivalent points on the surface should be arranged in staggered rows. And indeed this is what one observes in STM as depicted in Figure. One protrusion per unit cell (blue/light diamond in the STM image) is observed for all bias voltages used, which we thus interpret as adatom positions. They form staggered rows along the [ ] direction with a mutual atomic distance of 6.6 Å within the row and a spacing of.4 Å. Under the tunneling conditions applied (-.5 V tip bias,.3 na) the total corrugation within the unit cell is about.7 Å between the adatoms and the three-fold hollows in between. Interestingly, within the dense rows along the [ ] direction (line A) the corrugation is about.4 Å, within the more open rows along the edges of the primitive unit cell, i.e. the [ 4 3] direction (line ) it is between.5 Å and.6 Å. The total corrugation as noted above is measured along the [ ] direction (line C) to be.7 Å with a saddle point at.3 Å in the gap between the atoms in the [ ] rows [3]. Even though these values represent the electronic corrugation they indicate that a single adatom per unit cell must be positioned prominently above the surface with the atoms in the remaining area positioned in a reconstructed layer below. Atomic resolution of this deeper layer, however, could not be achieved. Above the V of the c( ) surface a surface state is also found at approx.. ev below E F. However, in this case it appears to have no dispersion parallel to the surface [3] which indicates that the electrons are more localized to the adatoms. A complete dispersion measurement for the c( ) phase using different photon energies from synchrotron light revealed in addition a number of surface resonances with quite large dispersions [3] which, however, are difficult to interpret in view of the complex surface atomic structure of this phase. ased on the experimental data a model was also developed for the c( ) phase on ( ) as discussed in detail in [3]. The results of the core level line shape analysis allow only for relatively simple reconstruction features on a surface with nearly bulk like composition. The model which is displayed in Figure considers a carbon terminated bilayer on the basal plane type stripe of the unit cell. From the two high-symmetry positions feasible on this bilayer type structure, a Si adatom is chosen in the so-called H3 site, where the adatom is truly three-fold coordinated. This choice is based on comparison to the established model for ( ) [46], and is also the most reasonable model from a bond length point of view [3]. On the cubic type half of the unit cell the model extrapolation is less straight forward than in the case of the basal plane stripes. The experimental data suggest a model with staggered, triple bonded carbon dimers in Si bridging sites oriented in [ ] direction as the most likely choice. However, recent density functional calculations propose a different configuration [47]. While the Si adatom in H3 position on the basal-plane stripe is confirmed, on the cubic stripe the calculations finds carbon dimers oriented in [ ] direction to be energetically favorable. Here, however, the dimers are not staggered. 6.3 The carbon rich ( ) phase Finally, as noted above the ( ) phase resulting from 5 C annealing shows only a strong graphite-like carbon component at the higher E side of the bulk peak in the C s spectrum. The V spectra, however, are fairly similar to those from the c( ) surface. They are still quite different from those obtained for thicker graphite films [48] or bulk graphite [49], which indicates that the graphite-like C signal in the C s core level corresponds to a thin atomic layer only, which is also corroborated by the remaining energy gap. The ( ) phase requires further detailed investigations. 7 The ( ) surface On the ( ) surface, hydrogen etching also leads to a well ordered surface (see above). Initial experiments indicated, that only a ( ) structure exists [4]. However, more detailed recent work reveals that also here several phases exist, namely a ( ) and a c( ) phase, that are currently under investigation [5]. 8 Summary In summary, the electronic and atomic structure of non-basal plane surface on reveal a number of surprises. However, there is still a lot of work to do in order to fully understand all features of this class of surfaces. Similar to basal-plane surfaces, hydrogen etching can be used to generate atomically smooth surfaces. In UHV the surfaces display a rich structural phase diagram much so as on () and ( ). Most intriguing are the ordered phases that were observed on 4H-( ) where a Si rich ( ) phase develops which contains atomic chains at the surface with a onedimensional semiconducting surface state. At about 95-975 C, a c( ) phase evolves with a bulk-like stoichiometry and a ( ) phase shows graphitic features after annealing to even higher temperature of 5 C. The a-plane appears to be ( ) periodic for a wide range of surface compositions. However, depending on stoichiometry distinct structural phases apparently develop. The m-plane again is rich of different phases which still await an analysis of their atomic structure. Acknowledgements I would like to thank R.F. Davis, S. ishop and W.J. Choyke for providing different samples. S.E. Saddow thankfully helped considerably with the hydrogen etching procedure. Support by the European Community - Research Infrastructure Action under the FP6 Structuring the European Research Area Programme through the Integrated Infrastructure
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