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The deposition and the subsequent anneal of submonolayer amounts of cerium on Si{111}7*7 surfaces produces silicide structures. A Si{111}rt3*rt3-30-Ce phase can be formed. It is probably a surface-stabilized polymorph of CeSi_2 in the AlB_2 structure, as opposed to either the ThSi_2 or the GdSi_2 bulk structure. A Si{111}2*1-Ce phase can be formed, but it probably arises from either a ThSi_2-, a GdSi_2-, or a CaSi_2-polymorph of CeSi_2. In both cases, rt3*rt3-30 and 2*1, the diffraction patterns are not associated with adsorbate superstructures. Instead, we believe they derive from silicide structures.
Cerium and other rare-earth metals have been deposited on Si surfaces for use as oxidation enhancers [113-117]. The high reactivity of Ce can cause order-of-magnitude increases in the oxidation rate of Si surfaces, even at relatively low temperatures. Fujimori et.al.[y1] noted the presence of distinct Si{111}rt3*rt3-30-Ce and 2*2-Ce surface structures. These surfaces were created by heating (350 C) 1/3 and 1/4 ml of Ce, respectively, deposited on clean Si{111}7*7 substrates. The rt3 phase was attributed to a H_3-site structure. The 2*2 phase was attributed to a buckled-chain structure.
The intent of the current research was to create a Si{111}rt3*rt3-30-Ce surface, collect the LEED spectra, and analyze the structure. However, a purely rt3*rt3-30 diffraction-pattern of such a structure was never observed in the present work. Even though rt3 LEED beams were seen, these beams were always of similar intensity as the Si{111} integral-order beams, in the presence of 2*1 beams, and additional minor beam-clusters around the integer-order beams were present. Owing to the complexity of these patterns, LEED I(V) spectra could not be collected. The Ce/Si{111} surface structures formed were probably those of a cerium silicide. The next sections detail the formation of such surfaces and describe their possible structure.
The experiment was done with a similar sample and techniques as those used in the previous Au and Mg experiments. The prime difference in the Ce experiment was the evaporation of Ce metal.
Cerium is a very reactive metal and quickly tarnishes in air. It has a relatively low vapor pressure (3E-11 Torr) at its melting point (798 C). For this reason, Ce would have to be evaporated from the liquid state at approximately 1100 C with a vapor pressure of 2E-7 Torr. However, since liquid Ce is denser than the very stable oxide, ceria, it was necessary to use oxide-free source material, otherwise the ceria could float above the liquid cerium and hinder any evaporated Ce from reaching the Si sample surface.
High purity Ce was obtained from Ames Laboratory, Iowa State University, sealed in glass ampules. A semi-cylindrical (6 mm radius, 10 mm length) slug was loaded into a Ta tubular evaporation-boat, which was previously outgassed. The boat was mounted in the vacuum system 10-cm away from the sample surface and inclined at a 30 degree angle.
The temperature of evaporation was monitored using an infra-red detector on the side the tubular Ta-boat. Initially, the Ta boat required 60 Amperes to reach roughly 800 C and melt the Ce. At this point there was not a large burst of gas. Instead, the melted Ce uniformly filled the bottom of the Ta boat and effectively shorted-out the source. This was clearly visible in the form of a `silhouette' of Ce in the tube. The decreased resistance of the Ta boat, in the volume filled with Ce, created a sharp drop in temperature at the Ce/vacuum interface. For example, the unfilled part of the boat would be at 1000 C (bright red) and the Ce-filled part of the boat was at 800 C (dull red). From then on, larger current would be necessary to heat the source and overcome the reduced electrical resistance of the source. Typically, more than 100 Amperes would be required to heat the Ce source material to 1100 C. Such high currents caused large initial outgassing (which subsided), but never caused any measurable sample substrate-heating.
The cleanliness of the deposition was monitored using AES. Initial contaminants included S, Cl, C, and O, but oxygen was the only contaminant to remain after sustained outgassing and prolonged use of the Ce source. Typical values of R_(O/Ce) (the ratio of the O_(510 eV) and the Ce_(82 eV) Auger peaks) were 1/50, indicating a proportion of less than one oxygen atom per cerium atom [37]. Interestingly, several AES peaks were observed which are absent in the standard AES spectra [37], at approximately 902, 884, 291, 207, and 109 eV. These peaks probably correspond to transitions involving cerium's M_IV 3d_3/2, M_V 3d_5/2, N_I 4s, N_III 4p_3/2, and N_IV 4d_3/2 electron energy levels and some interfacial states near the Fermi energy.
Cerium thickness was also monitored using AES. The ratio R_Ce, see Table 2.1, was compared to calculated values in the same way as for Mg and Au, where R_inf = I^inf_Ce(82) / I^inf_Si(92)=0.38 and lambda_82 = 4.3 Ang. were assumed [37, 38]. Using this method the deposition rate was determined to be 3.0 Ang./min. This number could not be converted to LE/min., since the structure and the thickness of a layer were not known. Also, it should be noted that the thicknesses were based on the assumption of non-interdiffusion; this was most likely not a correct assumption.
Ce was deposited onto a room-temperature Si{111}7*7 surface several times. Increasing thicknesses of Ce masked the fractional-order 7*7 LEED beams, created higher background, and eventually masked the integral-order beams. When visible, the integral-order beams were always sharp and there was never any indication of a room-temperature ordered-phase of Ce forming on the substrate surface.
The sample was also heated progressively from 100 C to 1100 C with various Ce thicknesses, and was then allowed to cool to room temperature. LEED observations at no time revealed the presence of a separate and distinct Si{111}rt3*rt3-30-Ce structure. Instead, when 1/2 Ang. of Ce was deposited the pattern remained 7*7 with an increase in the background. Anneal at 200, 250, ........, 500, 550 C for 2 min. caused no apparent change of the LEED pattern. Anneal at 600 C caused the appearance of single-domain 2*1 streaks in the LEED pattern. Further progressive anneal only caused the streaks to sharpen, but the streaks never completely transformed into individual beams. AES measurements indicated a halving of the surface Ce content as a consequence of the annealing treatment.
Deposition of 1.5 Ang. of Ce onto a clean substrate caused the almost complete disappearance of (1/7)-order beams into the increasing LEED background, but the integral-order beams remained sharp. No new ordered-phases were seen. Therefore, the surface was covered by an amorphous Ce film, which increased the amount of incoherent scattering. Anneal caused no change until 400 C, when the integral-order beams broadened and inter-integral-order streaking became visible. With increasing anneal temperature, up to 500 C, the streaking became more pronounced and formed into a 2*1 pattern. However, the pattern was not uniform across the sample. rt3 beams became visible after anneal to 600 C. These beams were of equal intensity as the integral-order beams, and the ratio of intensity between the 1*1 and rt3*rt3-30 beams did not change appreciably across the sample. Additional minor beam-clusters located around the integral-order beams were visible. It should be noted that the R_Ce ratio had not changed, and the distribution of 2*1 and rt3*rt3-30 domains on the sample was not uniform. Anneal to 850 C (in 50 C steps) sharpened the pattern, anneal to 1000 C caused a uniform distribution of 2*1 and rt3 beams. These beams had intensity comparable to the original Si{111}1*1 beams. The satellite clusters were more distinct, and there was a reduction in the Ce thickness by a factor of four.
Higher thicknesses (up to 12 Ang.) of Ce on Si{111}7*7 caused a complete obliteration of the LEED pattern into a very high background. Anneal of such surfaces yielded a similar progression of results: the appearance of broad 1*1 beams, inter-integral-order streaks, 2*1 beams, rt3 beams, and minor beam-clusters. Owing to the complexity of these phases, LEED I(V) spectra were not collected.
Once these structures were formed, it was not straightforward to return to the clean Si{111}7*7 structure. Anneal alone to 1100-1200 C did not remove the Ce. Although slight ion-bombardment removed the Ce atoms, anneal brought back the above mentioned Ce/Si{111} structures. The sample often needed to be bombarded for more than 5 hours in order to completely remove the subsurface Ce atoms. After this process, anneal to 1100 C produced a clean Si{111}7*7 surface.
In contrast to Fujimori et.al.'s[y1] report of a low-coverage 2*2-Ce phase, the present observations show that the surface unit-mesh was actually 2*1. Although a single-domain 2*1 phase was never seen, its precursor was seen: as mentioned above, inter-integral-order beam streaking was often visible, and such streaking often only occurred in one domain, i.e., in one direction. These streaks would then sharpen with additional heat and coalesce into (1/2)-order beams. One cannot distinguish a domain-averaged 2*1 from a 2*2 LEED pattern by the mere presence of fractional beams, but one can do so by examining the relative intensity of (1/2)-order beams. Referring to Figure 6.1, patterns a, b, and c are schematic LEED patterns for three single domains of a Si{111}2*1 surface. Pattern d is a combination of the previous three patterns. Pattern e is a schematic diffraction pattern for a Si{111}2*2 surface. The (1/2,0 beam of pattern a, the (-1/2,1/2) beam of pattern b, and the (0,-1/2) beam of pattern c are degenerate under rotation of the surface unit-mesh (domains) with respect to the bulk-like surface unit-mesh. These three beams, and others, would combine to form pattern d, if all three the domains were present on the surface, but they would only have equal intensity if the domains were present in equal concentration. The same three beams in pattern e are degenerate under the symmetry operations of the 2*2 surface unit-mesh and should always have equal intensities at normal incidence. In the present experiment, these beams did not always have equal intensities at normal incidence. Therefore, the surface unit-mesh was in fact 2*1.
What surface structure of Ce and Si would form a 2*1 LEED pattern? First, experimental observations (thermal stability, permeance of Ce, intensities of beams) indicate that the surface is not a simple adsorbate-superstructure, as in the case of Al, Ga, In, etc.... on silicon. A silicide is most likely. Second, the stable silicides of cerium are CeSi_2, CeSi, Ce_5Si_4, Ce_3Si_2, and Ce_5Si_3 [118]. The CeSi_2 is the most prevalent and occurs in bulk form in two modifications, a tetragonal ThSi_2 and an orthorhombic GdSi_2 structure. The latter is a distortion of the former's unit-cell, from (4.3*4.3*13.8 Ang.) to (4.2*4.1*13.9 Ang.). Figures 6.2.a and b are schematics of the unit-cell of the ThSi_2 structure and a resultant LEED pattern. The pattern was calculated using a kinematic structure factor, the same symbols represent degenerate intensities, the pattern is domain-averaged, the reciprocal unit-mesh of the ThSi_2 structure is indicated with dotted lines, and the reciprocal unit-mesh of the Si{111}1*1 surface is indicated with solid lines. This calculated pattern accurately represents the observed 2*1 diffraction patterns discussed in the previous section. The calculated diffraction pattern indicates several things: a) the ThSi_2 pattern is not commensurate with the Si{111}1*1 pattern, b) the (1/4)-order beams are weak, c) the (1/2)-order beams are most prominent, and d) there are small beam-clusters where the Si{111}1*1 beams would be located. The most important fact is, however, the presence of a ``2*1-like'' diffraction pattern created by this ThSi_2 or a GdSi_2 unit-mesh. Hence, Ce on Si{111} may form a silicide.
2*1 LEED patterns on Si{111} have also been reported for Gd, Yb, Eu, Ca, and Ba silicides [114, 119]. The 2*1 patterns for Gd and Yb silicide can be associated with the previous ThSi_2 discussion. The other three silicides probably have 2*2 patterns created from a bulk-terminated CaSi_2 structure. Interestingly, CaSi_2 can also be formed in the GdSi_2 structure, under high-pressure and temperature 120].
There are five main disilicide-structures that rare-earth metals may form: tetragonal ThSi_2, orthorhombic GdSi_2, hexagonal AlB_2 (where the Al and B atoms are replaced by rare-earth atoms and Si atoms, respectively), hexagonal AlB_2-x (where x varies between 0 and 0.3), and trigonal/rhombohedral CaSi_2 (where the Ca atoms are replaced by rare-earth atoms). ThSi_2, see Figure 6.2.a, has a quasi-hexagonal mesh of Si atoms, with the rare-earth atoms in the interstices. The building block of the unit cell is indicated with solid lines. This subunit contains four Si atoms (filled squares) at the corners of a rectangle (a*b) and one rare-earth atom (empty circle) in the center, below the plane of Si atoms. Alternate stacking (in and out of the plane) and translation (by 1/2 a x^hat) of these subunits create the ThSi_2 unit-cell, with glide-plane symmetry. The AlB_2 unit-cell, see Figure 6.2.c, is created from the same subunit when the ratio of the rectangular sides (a/b) is rt3 and the translation vector is 1/2 a x^hat + 1/2 b y^hat. The AlB_2-x structure has one-in-six B (Si) atoms removed, depicted in the figure by the empty triangular symbol. This structure has a rt3*rt3-30 surface unit-mesh when compared to Si{111}1*1. The CaSi_2 structure is similar to AlB_2, but the stacking is of alternate Si, Si, and Ca 1*1 hexagonal meshes. This structure has a 2*2 surface unit-mesh when compared to Si{111}1*1.
Everyone of the rare-earth elements is reported to form in of the above structures with Si, and most form at least two polymorphs. The silicides of Sc, Y, Nd, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu have all been reported to crystallize in the AlB_2 or the AlB_2-x structure. The silicides of (Ca, Y, La, Ce, Pr, Nd, Sm, Gd, Tb, Dy, and Ho) and (Y, La, Ce, Pr, Nd, Sm, Eu, Gd, and Dy) have been reported [121] to form the GdSi_2 and the ThSi_2 structures, respectively. Additionally, the silicides of Y, Gd, and Er have rt3*rt3-30 LEED patterns (mixed with 2*1) arising from the AlB_2-x surface unit-mesh [19, 114, 115, 116, 117]. Since there exist so many interconnected polymorphic structures, we believe that the Si{111}rt3*rt3-30-Ce LEED pattern arises from a surface-stabilized AlB_2-x-polymorph of CeSi_2. We cannot rule out, however, the presence of a rt3*rt3-30 adsorbate-superstructure on a bulk-terminated CeSi_2 surface.
In conclusion, the (2*1) and the rt3*rt3-30 LEED patterns of Ce on Si{111} probably arise from silicide structures. We find no evidence to support the presence of Si{111}-Ce adsorbate superstructures, as opposed to the cases of most group-III, -IV, and -V metal-atoms on Si{111}.
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