Epitaxial CoSi2 and NiSi2 thin films

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Materials Chemistry and Physics, 32 (1992) 107-133

Invited Review

Epitaxial CoSi, and Nisi, thin films R. T. Tung AT&T

Bell Laboratories

Murray Hill, New Jersey 07974

(USA)

Abstract Single crystal C&i, and NiSiz structures fabricated on silicon possess the most perfect metal-semiconductor interfaces presently available. They represent the best model systems with which to understand the physical, chemical, and electronic properties of metal-semiconductor interfaces. Epitaxial silicide thin films are also of technological importance because of possible applications in contacts, detectors, high speed devices, and threedimensional integration. The fabrication, structures, and properties of epitaxial NiSiz and CM& thin films will be examined in this paper, with special attention paid to the atomic structures and the Schottky barrier heights observed from epitaxial silicide/Si interfaces.

1. Introduction

2. Basic properties

Metal-semiconductor (MS) structures are necessary for virtuaily all semiconductor electronic and optoelectronic devices, although usually only as a passive part of the circuit, e.g. as ohmic contacts and Schottky barrier (SB) clamps. Recent advances in the science and technology of epitaxial growth of semiconductor structures have also made an impact on our capability to fabricate epitaxial MS structures. Epitaxial MS interfaces offer new opportunities for both device application and fundamental research. On the application side, epitaxial MS structures are an essential part of monolithic vertical integration, and they have potential use in high-speed device application. On the fundamental side, epitaxial interfaces provide an unprecedented opportunity to study, from first principles, basic processes at a MS junction, such as the formation mechanism of the Schottky barrier height (SBH). Among all epitaxial MS structures which have been fabricated thus far, NiSiJSi and CoSi,/Si have clearly the highest structural perfection El]_In this paper, the fabrication and various properties of these two silicides are examined.

Both CoSi, and NiSiz have the cubic CaF, lattice structure [2] a schematic of which is shown in Fig. 1. The diamond structure of Si is also shown for comparison. Even though there are similarities, differences in details of these two crystal lattices lead to a variety of possible atomic structures for their interfaces. The difference in crystal symmetry also dictates the presence of certain defects at some silicide/Si interfaces, as will be discussed. At room temperature, the lattice constant of CoSi, is smaller than that of Si by about 1.3% while the Nisi, lattice constant is smaller by about 0.4%. Since the atomic density of Si in either silicide

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lattice is very close to that in bulk Si crystals, one has the interesting situation that, seemingly, metal atoms may be ‘added’ to a Si crystal to convert it into a disilicide crystal with almost no change in its volume. Since both the Si and the disilicide lattices are cubic, the lattice mismatches just described apply to all crystallographic orientations. One should note, however, that the room-temperature lattice mismatches quoted for these silicides should not be directly compared with the observed densities of misfit dislocations in thick silicide layers. Because the thermal expansion coefficient for either NiSiz (-16~10-~ K-l) [3] or CoSi,( - 12 X lo+ K-l) [4] is larger than that of Si (N 3 x 10d6 K-l), the lattice mismatches at elevated temperatures are generally smaller than those at room temperature. A simple calculation indicates zero lattice mismatch between NiSi* and Si at - 380 “C and, for the CoSi&Si system, at - 1100 “C. These temperatures are to be compared with typical growth temperatures of the two siiicides, -500-800 “C for NiS& and -550-1000 “C for CoSiz. Since thick silicide layers are usually relaxed at the growth temperature, one expects the dislocation density to depend on the growth/annealing temperature. In view of the large thermal mismatch, growth of epitaxial silicide layers which are fully relaxed at room temperature does not seem likely. Although CoS& and NiSiz have an identical crystal structure and similar lattice constants, the issues/problems facing the epitaxial growth of these two silicides are ve dissimilar. For the formation of ultrathin (<20 “p;) silicide layers, a slight difference in lattice mismatch plays only a minor role. Therefore, the observed difference of the silicide reactions is probably attributable to other factors. One likely factor is the fundamental difference between the natures of the Ni-Si and the -Co-Si bond [S], leading to a difference in the energetics of the silicide/silicon interfaces. Another factor may be the dramatically different point defect densities of the two crystals, leading to very different diffusivities at typical growth temperature [6]. There has been a number of studies of the NiSiz phase which suggest that the actual equilibrium stoichiometry deviates sign& cantly from the Si:Ni composition ratio of 2:l of the perfect Cl crystal structure. A composition of Ni1.04Si1.92has been deduced from bulk crystals [7], with - 1.4% Si vacancies as well as -2% antisite defects (Ni atoms on Si sites), By comparison, the disilicide phase of cobalt is always found to be nearly stoichiometric, even though its crystal lattice is identical to that of NiSiz 18, 91.

At a given temperature, the diffusion of metals and Si through the Nisi, lattice is expected to be much easier than that through existing CoSiZ crystals. This may have a significant effect on the growth kinetics and the unifo~i~ of thin silicide layers. The difference in stoichiometries of NiSiz and CoSi, is also reflected in their electrical and optical properties. CoSi, has one of the lowest electrical resistivities among all metal silicides, - 14 p.st cm at room temperature, while NiSiz usually displays a high resistivity, N 30-40 fl cm [lo]. The large residual resistivity for Nisi, is indicative of a high density of defects and/or a signi~~ant degree of disorder in this lattice. Raman scattering experiments also revealed disorder-induced scattering which is consistent with Si vacancies and/or Niantisite defects in the Nisi, lattice, but little was observed for the CoSi, [ll]. The electronic str-uctures of Nisi, and CoSi, have been investigated experimentally [8, 121, and are generally in agreement with the calculated band structures. Theoretical calculations on NiSiz [I& 131 and CoSi, [13,14] generally conclude that the band structures of these two silicides are similar, except for a -2 eV shift of the bands (bands are positioned higher for CoSiZ than for Nisi*). Optical properties flS] and the Fermi surface structure [8] experimentally observed from NiSiz and CoSi, crystals are well explained by the calculated band structures. 3. Basic concepts of silicide growth 3.1 Solid phase reactions of Nisi2 and Co& Early studies of silicide formation have concentrated on nucleation and growth behaviors of thick layers of silicide (> 50 nm), using furnaces and evaporators under mediocre vacuum conditions. The simplest technique used to grow silicides is the solid phase reaction (SPR) technique [16], which entails the deposition of a pure metal layer on a silicon substrate at room temperature and the reaction of the metal with the Si substrate during a subsequent anneal resulting in the growth of various silicide phases [16, 171. In the SPR method, the silicon contained in silicide crystals usually comes from the substrate. The observed silicide reaction sequences for a number of metals have been discussed in terms of simple thermodynamic rules [18, 191. Beside thermodynamics, kinetics is sometimes found to play an important role in silicide reaction and nucleation. Usually after the reaction of the first siiicide phase is completed, due to exhaustion of the deposited

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metal, the growth of the next silicide phase commences at the interface between the first silicide and Si substrate. With deposited nickel and cobalt thin films on Si, annealing leads first to the growth of metalrich silicide phases (N&Si and Co,Si) followed by the monosilicides (Nisi and CoSi), and eventually the disilicides (Nisi2 and CoSi,) [20, 211. Marker experiments and oxidation characteristics of these silicides suggest that metal is the dominant diffusing species in the growth of Ni,Si, Nisi and Co,Si [20-221, while diffusion of Si is believed to dominate in the reaction of CoSi thin films 1231. Metal atoms are thought to be the mobile atoms in both the CoSiz and the Nisi, reactions [23, 241. Nucleation plays a dominant role in the Nisi, reaction and to a lesser extent, in the CoSi, reaction [23-251. For Nisi layers with thicknesses over a few hundred A, no reaction with the silicon substrate is observed upon annealing at a temperature below -750 “C. Above this temperature, NiSiz nuclei are generated at the interface which lead to very rapid growth of this phase. Such an unusual growth behavior is indicative of a small the~odynamic driving force for nucleation [25,26]. 3.2 Epitaxial growth of thick fitm CoSiz and NiSit by SPE and MBE Solid phase epitaxy (SPE) was the only technique used for epitaxial silicide growth before the early 1980’s [27, 281. As a result of early works on silicide reaction, the material aspects of the Nisi, and CoSi, systems are currently well understood, at least for reactions involving thick layers (> 50 nm) [ 29-311. However, early applications of SPE suffered from poor control of the processing ambients, and the quality of epitaxial silicide films was mediocre at best. The elimination of impurities during deposition and annealing was shown to have a beneficial effect on the reaction of silicide thin films 1327. This finding, concurrent with the advances in the field of silicon molecular beam epitaxy (MBE), led to attempts to fabricate epitaxial silicides under cleaner conditions [33-351. The application of ultrahigh vacuum (UHV) conditions led to the growth of silicide films which are truly single crystalline. The SPE process may be viewed as consisting of two parts: nucleation and growth. The nucleation of a silicide phase with a crystal structure which is different from that of either the metal or Si is an aspect not encountered in usual MBE growth of semiconductors (The heteroepitaxial growth of any III-V compound semiconductor on III-V material. or the growth of SiGe on Si, requires no

nucleation). After nucleation, the ‘growth’ of silicide crystals require mass transport (diffusion). Metal has to diffuse to the silicide/Si interface, or silicon has to diffuse to the metal/silicide interface, for growth to be sustained. This is another feature which has no parallel in semiconductor heteroepitaxial growth. Since the nucleation of silicide is often heterogeneous, reacted silicide films thinner than 1 pm thick are sometimes inhomogeneous. This is particularly the case for Nisi, and CoSi,, both of which are known to have nucleation problems [23, 251. Co-deposition of metal and silicon, in the stoichiometric ratio, onto a silicon substrate held at an elevated temperature, constitutes the MBE technique for epitaxial silicides. The MBE process is usually done in a UHV environment. MBE growth of silicides directly on clean Si surfaces also faces many problems, most of which are related to the nucleation process. Codeposition of stoichiometric silicide is not the same as deposition of ‘silicide molecules’, at least not in the initial stages of growth. One has to view this process as random depositions of Si atoms and metal atoms on a heated Si clean surface. Deposited Si atoms, in the absence of deposited metal atoms, are essentially engaged on Si homoepitaxial growth. Deposited metal atoms diffuse on the Si surface until they either combine with Si (substrate or adatoms) to nucleate silicide crystals or get incorporated into existing silicide crystals. At low MBE temperatures, nucleation of silicide is more uniform across the surface; however, the surface is already roughened by partial Si overgrowth; therefore, the silicide/Si interface is rough on an atomic scale. At high MBE temperature, high surface mobilities lead to growth under conditions closer to thermal equilibrium, resulting in three dimensional (island) layer morphology. 3.3 The template technique The reactions of monolayers or submonolayers of metal with clean semiconductor surfaces have been studied extensively since the mid 1970’s. It was clearly demonstrated that the first few monolayers of most metals react with Si even at room temperature 136, 371. Identification and explanation of these silicide phases has been the main focus in the early UHV metal-Si studies. Subsequently, it was discovered that the monolayerrange silicide reaction could be used, much to one’s advantage, to grow uniform, ultrathin, high quality, single crystal, epitaxial silicide films [38]. These ultrathin silicide layers may be used as ‘templates’ to grow thicker epitaxial silicide layers

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[39] with structural qualities much better than layers of similar thickness, but grown directly by SPE or MBE. This two-step growth method for heteroepitaxy, known as the ‘ template technique’ [39] benefits from the fact that ‘nucleation’ and ‘growth’ are carried out under two, sometimes very different, conditions, each of which optimized for its specific purpose. Generally, the first step (nucleation), which itself may consist of several processing maneuvers, is done at a low temperature for maximum layer uniformity and the second step (growth) is done at a, usually higher, temperature optimized for best crystalline quality. Of course, the difference between the first and the second steps may be as subtle as a slight difference in the compositions of the evaporants, the deposition rates, the intensities of external excitation sources, etc. Heteroepitaxial systems other than silicides have also benefitted from such a two-step processing scheme [40-42]. The template technique has become the technique of choice in the growth of epitaxial silicide thin films, because ultrathin silicide layers of a few monolayers thickness may be fabricated with much higher structural quality than thicker (> 10 nm) layers. The reduced misfit stress during nucleation, the uniformity of deposition and precursor reaction/inte~~ing at low temperature, the flatness of the starting Si substrate surface (compared with the roughness of the interface between an intermediate silicide layer and Si substrate), and the simple fact that one has at one’s disposal many different, well-controlled, deposition and annealing schedules are the main reasons why ultrathin silicide ‘template’ layers may be grown with near perfect crystalline quality. One may also point out that the reaction sequence and nucleation environment for monolayers of deposited metal with Si may be quite different from those discussed above for the case of thick-film SPE. Thicker silicide layers grown on these template layers then simply inherit the high crystalline quality of the thin layers. 4. Epitaxial NiSi,/Si(lll)

structures

4.1 Eady studies

In the early days of SPE under non-UHV conditions, NiSi,/Si( 111) was seemingly the easiest system for epitaxial growth [1, 301. Uniform, epitaxial Nisi, layers of good quality were grown by furnace annealing, while polycrystallinity seemed unavoidable for other systems, such as NiSiJ Si( 100) and CoSi,/Si( I1 1). However, NiSiz layers contain mixed A and B orientations. Type A NiS&

has the same orientation as the silicon substrate. Type B NiS& shares the surface normal (111) axis with the Si, but is rotated 180” about this axis with respect to the Si. This double-position problem was not solved by simply resorting to a clean UHV environment [l]. SPE growth of NiSiz in UHV from a thick layer of deposited Ni goes through the usual intermediate phases and also requires a high nucleation temperature, just as in a nonUHV environment. Typically, type B grains occupy a slightly larger portion of the Nisi, films grown by SPE than type A orientation [30,31,43]. Based on the idea that the nucleation of Nisi2 is highly heterogeneous at a NiSi/Si interface, Lau and Cheung [44] explained the observed morphology of thick Nisi, layers as a result of very rapid vertical growth of isolated nuclei followed by, also rapid, lateral grain growth. From this model, one may infer that type A and type B Nisi* nucleate with similar likelihood at the NiSi/Si(lll) interface. Presently, the circumstances which lead to the nucleation of A- and B-type Nisi, from the Nisi/ Si(lll)interface are not fully underst~d and the SPE technique is not deemed a growth technique capable of producing single crystal Nisi2 layers. It is possible, however, to grow single crystal type B NiS& films on Si(ll1) using pulsed laser melting and furnace annealing [45]. 4.2 Room temperature Ni reaction on Si(ll1) The reaction of monolayers of nickel with the Si(ll1) surface at room temperature has been studied using various techniques. It is clear that considerable intermixing occurs between the first few monolayers (ML) of deposited nickel and silicon substrate even at room temperature [36, 461, although the identity, or identities, of the ‘silicide phase(s)’ is still not fully resolved. Photoemission studies [36,47,48] consistently measure a Ni 3d binding energy which suggests that the chemical environment su~ounding the nickel atoms is similar to that found in bulk Nisi crystals [47], although attribution to interstitial Ni was also made [36]. The stoichiometry of the room temperature phase, inferred from the number of displaced Si atoms observed in ion channeling studies, is close to Ni,Si [49]. Surface extended X-ray absorption fine structure (SEXAFS) experiments show a local bonding configuration of the roomtemperature phase to be consistent with epitaxial (B-type) Nisi, [50]. The formation of type B Nisi, at room temperature from the deposition of <5 ML Ni is also supported by subsequent impact collision ion scattering spectrometry (ICISS) [Sl] and UHV-TEM [52]. There seems to be a wide

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difference in conclusions obtained by different experimental techniques. However, if one divides all the experimental results into three Ni thickness ranges: < 3ML, 3-10 ML, and > 10 ML, the discrepancies between reports are much reduced. In the smallest thickness range, <3 ML Ni, evidence for an epitaxial phase, most likely Nisi*, has been obtained by electron diffraction [52, 531 and SEXAFS [50]. RHEED [53] and Auger line shape [54] studies also support the formation of Nisi* at very small Ni coverages. Some degree of epitaxy has been observed from these layers by high energy [55] and low energy ion scattering [51] suggesting the formation of epitaxial Nisi,. The observation, by photoemission, of the Ni delectrons at a binding energy usually found in bulk Nisi is actually explainable in terms of interstitial Ni, or Nisi, [56]. UHV-TEM results are also consistent with the formation of type B Nisi:! at low Ni coverages [52]. Furthermore, strong support for type B Nisi, formation was provided by MBE growth experiments [57], shown in Fig. 2. One notes, however, that epitaxial formation of Nisi, at room temperature is limited to the 7 x 7 surface, as recent experiments with the Si(ll1) 1 x 1 surfaces failed to show evidence for epitaxial growth under similar conditions [58].

I

1.2

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Fig. 2. 2 MeV He (111) channeling spectra of Nisi, layers grown and the by co-deposition of -72 8, NiSiI at room temperature indicated amounts of nickel predeposition. A glancing exit angle, -3”, was used.

When more than 3 ML Ni is deposited, there is clear evidence for an inhomogeneous layer morphology of the reacted silicide [52, 49, 541. An ion scattering study showed evidence for excess Si segregating on top of Ni,Si islands [49], while a diffraction study favored the existence of unreacted nickel lying on top of Nisi, islands [52]. This islanded growth morphology proceeds at least until the Ni coverage exceeds - 10 ML. From the standpoint of epitaxial Nisi, growth, it is clear that the rough morphology as a result of deposition of 3-10 ML Ni at room temperature does not make this an attractive thickness range for the growth of high quality films. Above 10 ML Ni, the silicide islands coalesce, the Ni-Si reaction stops, and unreacted nickel starts to accumulate on top of the reacted layer. 4.3 Growth of thin Nisi2 layers 4.3.a Deposition at elevated temperatures The structure of Nisi, layers obviously depends on various growth parameters such as the composition and the amount of deposition, the deposition temperature, and the annealing temperature. Deposition at elevated temperatures usually results in Nisi, layers which are non-uniform. For instance, when pure nickel is deposited at a temperature in excess of -350 “C, epitaxial Nisi2 is grown on Si(ll1). However, these silicide ‘layers’ are essentially made up of disconnected Nisi2 patches with both type A and type B orientations [59, 601. The morphology and the crystallographic habit of these A- and B- type inverted islands have been studied in detail by HREM [60]. The surfaces of the isolated Nisi, ‘islands’ are flush with the uncovered Si surface. The growth rate of A-type islands is shown to be about ten times that of B-type islands [60], because the type A interface is kinetically more mobile. It seems that type A and type B Nisi, nucleate with similar probabilities when nickel is deposited at elevated temperatures. When both Ni and Si are deposited at an elevated temperature (MBE), there is a clear preference for type B orientation in the grown Nisi, layers [61]. Stoichiometric co-deposition often leads to pure type B orientation, although this is not guaranteed. Parameters such as the stoichiometry of co-deposition rate, the deposition temperature, and the preparation of the clean Si surface are all found to have an effect on the epitaxial orientation and the morphology of the Nisi* layer. MBE grown films often have problems with layer inhomogeneity and interface roughness, as revealed by planview TEM in Fig. 3. At a type B interface,

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10 AVERAGE

Fig. 3. Planview, dark-field, TEM image, formed with (111) reflections originating only from the type B orientation, of Nisi2 layers grown by co-deposition of Nisi, on Si(ll1). The average film thickness is 50 8, and the deposition temperatures are (a) N 300 “C and (b) 500 “C, respectively. The contrast of the image is roughly proportional to the thickness of the film, revealing clearly the non-uniformity of these layers.

symmetry requires the existence of a partial dislocation with a Burgers vector of 1/6(112)a, (or 1/3(111)a,) at a step of single (3.13 A> or double atomic height. Therefore, the high density of dislocations, seen in Fig. 3(a) for a thin type B NiSiz layer, MBE-grown at <400 “C, is indicative of a rough interface mo~holo~. At a temperature of 2500 “C, the dislocation density is noticeably reduced, but the layer uniformity is poor as large pinholes and gaps may clearly be seen (Fig. 3(b)). If co-deposition is carried on much beyond these thicknesses, over 300 8, for example, holes in the layer are expected to be filled and uniform, high quality, type B Nisi, layers may be rown. However, it is unlikely that .thin (< 100 1 ) and uniform NiS& layers can be grown by MBE at elevated temperatures. MBE seems to be a useful technique for homoepitaxial growth of thick NiSiz layers, but may be ill-suited for the nucleation of thin silicide layers. 4.3.b Growth by room temperature nickel deposition and annealing When different amounts of nickel are deposited at room temperature on Si(ll1) and then annealed to N 500 “C for a few minutes, epitaxial Nisi, layers with different orientations may be grown [38]. The variation of the NiSiz orientation as a function of deposited Ni thickness is shown in Fig. 4. NiSip layers grown with less than 9 ML Ni have majority type B orientation, while layers with N 16-20 8, Ni have pure type A orientation. It has been suggested that below 3 ML, the structure of the silicide phase may be slightly different from

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OF DEPOSITED

NICKEL

(i,

Fig. 4. Orientation of thin NiSi2 Iayers grown on Si(ll1) deposition of nickel at room temperature and annealing -450 “C.

by to

that of bulk Nisi* [62], The A/B formation characteristics depends somewhat on details of sample preparation such as annealing procedure, deposition rate and temperature, and wafer preparation. It has been shown a faster deposition rate and an immediate annealing step favor the fo~ation of type A orientation [59]. Slower procedures help to obtain type B layers. A small wafer misorientation favors the formation of type A NiS&. The variation of Nisi, orientation as a function of deposited nickel thickness, depicted in Fig. 4, is likely influenced by the room temperature NiSi(ll1) reaction. For example, the growth of type B Nisi, at low coverages is most likely due to the fact that either type B Nisi,, or a structure very similar to the type B NiS&, has already been nucleated at room temperature. As explained previously, the nucleation of type B Nisi* (and type B CoSi,) at room temperature occurs on the 7 X7 surface but not the 1X1 [SS]. Because of this dependence on the structure of the Si surface, it may be argued that kinetics is the main reason for the nucleation of type B oriented NiSiz at room temperature. Presently, it seems appropriate to think of the nucleation of type B Nisi* at room temperature as due to kinetic AND energetic advantages. The formation of type A NiSiz has been proposed as due to a kinetic advantage in its growth 1381. Type A orientation has also been suggested, under certain experimental conditions, to be linked to an intermediate O-Ni$i phase [52, 631. However, one notes that even though the 8NiaSi phase is not part of the reaction sequence between thick films of nickel and Si, type A orientation is observed to occupy roughly half the reacted NiSiz film. Therefore, the nucleation of type A NiSiz is not always mediated by a precursor @Ni,Si phase. Obviously, there are still unanswered

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questions concerning the formation of type A and type B Nisi, by a variation of the nickel thickness. Fortunately, the growth procedures for high quality NiSiz of either orientation are already demonstrated and carefully documented, despite whatever uncertainties we may have regarding the physics of their fabrication. Thin type A layers grown by Ni deposition are very uniform in thickness, and contain no misfit dislocations, indicating the growth to be essentially pseudomorphic. The type B oriented Nisi, layers, grown by this technique with less than -5 8, deposited Ni, are not uniform in thickness and furthermore, they often contain a small fraction of type A oriented grains [38,64]. Such morphology obviously is related to the inhomogeneity, at these Ni coverages, of the room temperature Ni/Si(lll) reaction. With slightly higher Ni coverage, near 9-11 A in one study [59] and - 5 A in another [65], the growth of very uniform, pure type B oriented, NiSiz layers have been reported. However, these results are so sensitively dependent on the details of the deposition and annealing that such a grown technique (deposition of a specific amount of nickel) is not very reliable in producing high quality type B oriented layers. Instead, growth techniques involving deposition of Si or co-depositi(~n of Nisi, are preferred. 4.3. c Growth from deposition of Ni and Si When a suitable amount of Si is deposited at room temperature following Ni deposition, a subsequent anneal leads to the formation of uniform type B Nisi;! layers [64, 591. The attainment of pure type B oriented NiSiz by the deposition of Si probably works for the same reason which allows type B Nisi, to grow from co-deposition and multilayered Ni/Si structures. To ensure the growth of an uniform Nisi, layer, deposited nickel should exceed 15 A in thickness to avoid inhomogeneity in the as-deposited layer [59, 66f. Since 1 8, Ni reacts with -3.6 A Si, and a small amount of Si from the substrate is always involved in the reaction with deposited nickel, one expects the optimum thickness of deposited Si to be -3 times the deposited nickel thickness, in good agreement with the conclusion of a detailed experimental study of this technique [59]. A well-tested, reproducible condition for the growth of a type B NiSiz template is the deposition of - 18 A Ni and -50 A of Si [67]. fn practice, the growth of pure type B oriented Nisi, does not seem to depend critically on the exact amount of deposited Si [64, 661. Deposition of multiple layers of Ni and Si, hundreds of angstroms for each individual layer, results in a

structure consisting of alternate layers of Ni,Si and amorphous Si at room temperature [68, 691. Upon annealing to over 500 “C, transformation into uniform, single crystal, type B oriented NiSiz layers is observed f68, 691. Co-deposition of NiSiz at room temperature leads to the epitaxial growth of type B Nisi, (Fig. 2). The quality of the silicide layer is much improved if a small amount of nickel is first deposited (‘predeposition’~ onto the clean Si(lll), as illustrated in Fig. 5(a). Without predeposited nickel, defective Nisi* layers are formed which are essentially type B oriented but contain a small fraction of type A grains. With the pre-deposition of -2-3 A Ni, co-deposition of NiS& leads to the growth of high quality single crystal B-oriented Nisi, at room temperature. The dislocations present at these high quality interfaces are all related to steps on the original Si(ll1) surface, as shown in Fig. 5(a). Low-temperature-grown NiS& layers may be annealed to elevated temperatures for an improved crystalline quality. This is particularly the case for the more defective Nisi, layers grown at room temperature, of which the dislocation density may be significantly reduced after the annealing. For those layers which have been grown with an optimum nickel pre-deposition, annealing leads to no change in the density of the dislocations, even though the dislocations themselves seem to become somewhat straighter. Figure 5(b) shows an example of very high quality type B Nisi, layers grown by co-deposition at room temperature and annealing to over 500 “C. If, however, the stoichiometry of the room-temperature co-deposition is Si-rich, then annealing leads to Si precipitates and pinholes in the NiS& layer. It is, therefore, generally advan-

Fig. 5. Planview, (220) weak beam, TEM images of - 50 A thick type R NiSiz layers on Si(ll1). (a A layer grown at room temperature by deposition of -2 1 Ni and co-deposition of NiSi,, and (b) a layer grown by the same technique as (a), but with an additional anneal at -500 “C.

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tageous to maintain the co-deposition ometry slightly on the Ni-rich side.

stoichi-

4.3.d Stability of ultrathin Nisi, layers

In general, the crystalline quality of Nisi2 layers improves with annealing. Occasionally, annealing, at temperatures >SOO “C, of a thin Nisi, layer which has been grown without the deposition of Si leads to pinhole formation in the layer. However, the discussion of this phenomenon in terms of absolute stability of thin Nisi2 layers [48, 701 is probably unfounded. It is known that ultrathin type B Nisi2 layers grown by co-deposition at room temperature are stable to at least 800 “C. The apparent ‘instability’ of thin type B Nisi*, grown by deposition of Ni alone, is likely related to the specific stoichiometry and morphology of these layers which result from its particular growth kinetics. It is likely that Nisi2 layers grown at low temperatures by Ni deposition alone are slightly Ni-rich compared to the bulk stoichiometry. Upon annealing at higher temperature, the process of layer and surface stoichiometry adjustment takes place, leading to the formation of pinholes, much like the mechanism responsible for the formation of pinholes in CoSi,, to be discussed later. As shown in Fig. 6, the deposition of a small amount of Si on Nisi2 significantly increases the stability of Nisi2 layers, either type A or type B oriented, grown by Ni deposition alone. Generally, if a Nisi2 layer thicker than the template is desired, there

Fig. 6. Planview, (200) dark-field, TEM images of two -55 8, thick type A NiSil layers on Si(ll1). Both layers were grown by the deposition of 17 A Ni and the annealing at 500 “C for 1 minute. On layer (a) 20 8, Si was then deposited, while no further deposition was done on layer (b), before both layers were given a second annealing at 750 “C for 2 minutes. Large pinholes are found in layer (b), but none in layer (a). Neither layer contains dislocations. However, there is a slight modulation of the film thickness due to steps at the interface and on the surface of the layers (from the unintentional misorientation of the Si(ll1) substrate).

is little to be gained by annealing the template at much above 500 “C. Annealing of thin pseudomorphic type A layers at higher than 600 “C leads to a clear ordering of the step array, revealed by a stair-case type variation of the layer thickness (due to wafer misorientation), as shown in Fig. 6. This is indicative of type A interfaces becoming smoother upon annealing. 4.4 Homoepitaxial growth of Nisi,

Thick Nisi2 layers grown in two steps from thin template layers usually have higher crystalline quality than those grown in a single deposition. Essentially, the ‘thickening’ step is a Nisi2 homoepitaxy process. Early results [39] showed that the deposition of only Ni was sufficient for the growth of thick Nisi2 on existing Nisi2 layers. One may deposit a small amount of Ni at low temperature and use a subsequent anneal to induce, presumably, the diffusion of Ni through the existing Nisi2 layer. Generally, one finds that the orientation of the original template layer is preserved in the resulting thicker layer, if the deposited amount of Ni is kept to a small fraction of that contained in the pre-existing Nisi2 layer. Continuous deposition of Ni onto a heated Nisi, surface, sometimes referred to as reactive deposition epitaxy (RDE), also leads to an increase of the Nisi2 thickness without affecting the orientation of the original Nisi2 layer. Obviously, a necessary condition for this growth technique to be successful is that the Ni deposition rate may not exceed the diffusion and reaction rates. At intermediate temperatures, -300-550 “C, a fast deposition rate may lead to the transformation of the surface Nisi2 into Nisi and so forth. Hence, the RDE temperature should be chosen to be high enough to circumvent this problem and yet not too high so as to cause layer instability. A temperature in the range of 650-750 “C seems to work well for a deposition rate of -1 A Ni s-l. To avoid the problems associated with diffusion through NiSi2, deposition of both Ni and Si may be used for the homoepitaxial growth on thin Nisi2 templates. Deposition of alternating layers of Ni and Si, in roughly the correct stoichiometric ratio, at an intermediate temperature, - 300-400 “C, has been shown to lead to epitaxial growth of Nisi, on thin type B templates [48]. Deposition of layers of Ni and Si at room temperature, followed by annealing at >500 “C, also leads to the growth of high quality Nisi, layers. The most reliable and preferred technique for homoepitaxial growth is co-deposition at low temperatures followed by annealing. It has been shown that co-deposition

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on Nisi* leads to the growth of single crystal Nisi, at room temperature [57]. The crystalline quality of thick Nisi, layers may be improved by a high temperature, >700 “C, anneal [71]. The strain in thin layers or layers grown at lower temperatures depends on the processing and the orientation of the Nisi* layer. Type B NiS& layers always relax more easily than type A layers with the same thickness and similar preparation conditions. With a growth temperature of <650 “C, pseudomorphic growth of type A Nisi;? has been demonstrated to at least a 500 8, thickness. The strain in thick NiS& layers has been studied by TEM, XRD, optical, and ion channeling techniques [72, 73-751. Due to thermal mismatch, the lattice mismatch between Nisi, and Si change sign at - 350400 “C, leading to the interesting phenomenon that thick Nisi, layers grown or annealed at high temperatures, >700 “C, tend to have an in-plane lattice parameter even larger than that of Si [72]. Thick NiSiz layers contain a honeycombshaped network of dislocations with Burgersvectors of l/2( liO)a, and l/6( 112)a,, respectively, for type A and type B orientations. For an annealing temperature of 800 “C, ical separations between dislocations are - 1800? and 900 8, for the two orientations, respectively [72]. These numbers are likely to vary with the cooling rate. 4.5 Surface and integace structures of NiSi,l Si(ll1) The surfaces of NiSiJSi(ll1) usually exhibit unreconstructed 1 X 1 LEED patterns [30, 59, 76, 771. Because the LEED patterns of Nisi, are highly asymmetric (3-fold) at certain incident energies, a veiy convenient in-situ determination of the Nisi* layer orientation is afforded [59]. LEED intensityvoltage measurements [76], AES analysis [59], ion channeling and blocking [77] all showed evidence for a bulk-terminated (terminating on a Si-Ni-Si triple layer), unreconstructed, surface structure. However, upon annealing to higher temperatures, evidence for the segregation of Si on Nisi, surfaces has been found [78, 791. Recent STM and high resolution XPS studies of this surface [SO, 811 showed evidence for randomly distributed Si adatoms. The atomic structures of type A and type B NiSi,/Si(lll) interfaces were studied by high resolution electron microscopy (HREM) and found to both have 7-fold structures [64, 82-841. The terminology for the structure of a silicide interface, e.g. 7-fold, is based on the number of nearest Si neighbor atoms to a metal atom at the interface. In a bulk disilicide lattice, each metal atom has

(a)

(c)

TYPE

A

Nisi2

TYPE

El CaSip/Si

/Si

(Ill)

(b)

(III)

(d)

TYPE

NiSiz/

B

Nisi*

/Si

(III)

Si (100)

Fig. 7. Ball and stick models of epitaxial silicide/Si interfaces, viewed in the [liO] direction, (a) the 7-fold type A NiSi,/Si(lll) interface; (b) the 7-fold type B NiSi,/Si(lll) interface; (c) the S-fold type B CoSiJSi(ll1) interface, and (d) the 6-fold Nisi,/ Si( 100) interface.

a coordination number of 8. Structural models of the two 7-fold NiSiJSi(ll1) interfaces are schematically shown in Fig. 7. These early HREM results were later confirmed by X-ray standing wave (XSW) [73,85] medium energy ion scattering (MEIS) [86], and X-ray interference [87] investigations. 4.6 Electronic properties of NiSi,lSi(ll 1) An intriguing dependence of the SBH on the epitaxial orientation has been observed at the epitaxial NiSi,/Si(lll) interfaces: type A and type B Nisi, have distinctively different SBH’s [88, 891. The SBH’s of Nisi;? and CoSi, interfaces are summarised in Table 1. This dependence of SBH on the epitaxial orientation was briefly challenged [90]. However, very extensive studies by various groups [91-931 have since fully confirmed the original findings [88] that type B Nisi, has a SBH

116 TABLE Silicide

Nisi, Nisi, Nisi* Nisi* Nisi, CoSi, CoSi, CoSi, CoSi,

1. Schottky barrier

heights of epitaxial silicides

Orientation

Type A Type B

(1W W) (110) Type B Type B (100) (110)

Substrate

Si(ll1) Si(ll1) Si( 100) Si(100) Si(ll0) Si(ll1) Si(ll1) Si(100) Si(ll0)

Interface

structure

7-fold 7-fold 6-fold* 7-fold (11 l}A 7-fold (11 l}A 8-fold*** 7-fold*** 8-fold** 8-fold***

SBH (eV) n-type

p-type

0.65 0.79 0.40 0.65 0.65 0.67 0.4 0.71 0.70

0.47 0.33 0.73 0.45 0.45 0.44 0.7 0.41 0.42

*with partial 1 X2 reconstruction. **l x 2 reconstructed. ***Tentative.

about 0.14 eV higher than type A NiSiz on n-type Si(ll1). The origin of the initial disagreement is now understood [67] to be a surface boron contamination problem associated with the experimental conditions employed in the study of Liehr et al. [go]. Under careful procedures which avoid this boron problem, a low SBH, -0.65 eV, for type A Nisi, on n-type Si(ll1) is very consistently measured by every laboratory. I-V studies of type A Nisi* diodes on n-type Si(ll1) indicate that the ideality factor for high quality type A diodes is almost exactly 1.00, independent of the measurement temperature. As recently shown [94], this observation is consistent with a homogeneous SBH. For the type B NiSiz interface, some variation of the SBH data has appeared in the literature, which suggests that the SBH may depend on growth, diode processing, and the method of SBH measurements [69, 91-931. Most notably, it was reported that the SBH of type B Nisi* showed a dependence on film thickness [95] which was attributed to a variation in dislocation density [95, 961. In reality, however, the dislocation density at a type B NiSiz interface is not a simple function of the film thickness. Furthermore, the samples employed in the Kikuchi et al., study [95] likely have very rough layer morphology (e.g. Fig. 3). Inclined facets of a type B Nisi* layer are incoherent twin boundaries with structures different from the planar, 7-fold, structure. Non-planar type B NiS& interfaces most likely have inhomogeneous SBH’s which vary with layer morphology! Uniform type B Nisi* layers do not seem to have a gross SBH inhomogeneity problem, even though the slightly larger ideality factors usually observed at type B Nisi, interfaces are suggestive of some minor inhomogeneities.

There have been many calculations of the electronic structures and the SBH properties at type A and type B Nisi, interfaces [97]. Recent calculations, involving very large supercells [98, 991 have all yielded SBH results which are in excellent agreement with experimental results [88]. Although the densities of metal-induced gap states (MIGS) are high at both interfaces [99], there is no easily identifiable feature in the energetic distribution of these states to suggest a ‘pinning’ of the FL. Actually, the distribution of MIGS at these two interfaces significantly depends on the interface structure. This dependence indicates that the formalisms previously constructed for MIGS [loo, 1011 are too simplified for real MS interfaces. Furthermore, since MIGS depend on interface structure, one may not apply a charge neutrality criterion [102] based on the semiconductor band structure alone. The formation mechanism of the SBH at epitaxial Nisi* interfaces seems, at least in spirit, close to the charge transfer and chemical bonding mechanism previously discussed [ 103,104]. However, it is the specific charge redistribution due to the interfacial bonding, and not silicide bulk bonding, that seems to be more relevant. 5. Epitaxial

NiSi,/Si(lOO)

structures

5.1 Early studies and inte$ace instability It is relatively straightforward to grow epitaxial NiS& on Si(100) by SPE from deposited nickel [27, 301. However, layers grown by this method are only partially epitaxial and, furthermore, are very non-uniform in thickness [30]. Severe interfacial (111) faceting leads to modulation of the Nisi, film thickness with an amplitude comparable to the average thickness. As a result, SPE-grown Nisi* layers (at -800 “C) often contain a high

117

density of pinholes, even when the average film thickness far exceeds 500 A. The reason for interfacial faceting had, for some time, been thought to be due to a thermodynamic driving force [30], namely, that the interface energy of (type A) Nisi,/ Si{lll}, ojrll), is much lower than that of the NiSiJ Si{lOO}, g{lool. However, recent experiments [71] clearly show that the apparent preference for (111) facets is not a result of thermodynamics but is likely related to the growth kinetics. The fact that uniform Nisi, interfaces may be grown by pulsed laser melting [45] also indicates that the Nisi,/ Si(100) interface is stable. 5.2 Growth of ultrathin NiS& layers on Si(100) The crystallinity and morphology of thin Nisi2 layers grown on Si(100) depend on the manner in which the film is grown. First of all, as is generally true with all epitaxial silicide systems, the deposition temperature for the growth of high quality films should be kept as low as possible. This is because Nisi, layers grown by deposition or co-deposition at elevated temperatures are generally very non-uniform. Selected layers of Nisi, grown by the co-deposition of Nisi, at elevated temperatures are shown in Figs. 8(a) and 8(b). Contrast in these (200) dark-field images directly indicates thickness changes in the Nisi, layer, as the (200) reflections are forbidden in Si. Therefore the disconnected rectangular contrast observed in Figs. 8(a) and 8(b) are indicative that co-deposition at elevated temperatures leads to the growth of isolated islands of Nisi*, with the average size of the islands, and also the average separation between islands, increasing with the deposition temperature. Deposition of Ni at elevated temperatures, > 300 “C, directly onto Si( 100) leads to Nisi, morphologies even rougher than that shown in Figs. 8(a) and 8(b). Because of the poor layer

Fig. 8. Planview, (020) dark-field, TEM images of selected nonuniform Nisi* layers, -70-90 8, in average thickness, grown on Si(100). Layers (a) and (b), which were grown by co-deposition of Nisi1 at - 350 “C and 500 “C, respectively, contain discontinuous Nisi* islands. Layer (c), which was grown by deposition of - 13 8, Ni on a -25 A thick Nisi* layer and annealing at 500 “C, is essentially made up of facet bars.

morphology, deposition at elevated temperature directly onto Si is not deemed a suitable technique for the nucleation of high quality Nisi, layers. Depositions at room temperature onto Si(100) does not lead to the spontaneous formation of NiS&, in contrast to results observed on Si(ll1) substrates [57]. An anneal at >400 “C, after the deposition(s), is generally needed to induce the epitaxial growth of Nisi,. Obviously, depositions of either pure Ni or both Ni and Si and the codeposition of Nisi, can be used to grow epitaxial Nisi, thin films. The morphology of ultrathin ( < 30 A) Nisi, layers grown by deposition at room temperature and annealing at -500 “C is found to depend very much on the material(s) deposited [71]. It was noted some time ago that the morphology of Nisi, layers grown by the deposition of Ni alone depends on the amount of deposited Ni [38]. With less than - 10 8, Ni, isolated Nisi, islands and facet bars comprise the morphology of the annealed structure. The epitaxial formation of Nisi, from deposited thin layers of Ni occurs at -400 “C, as revealed by in-situ LEED [71] and TEM observations [52]. When the Ni thickness exceeds -30 I$, higher anneal temperatures are needed and the layers have very uneven morphologies [38]. The most uniform Nisi, layers which may be grown from deposition of Ni alone are layers which are -40-100 8, thick, grown with 10-30 8, of deposited Ni. Essentially continuous Nisi! layers are observed in this thickness range, although these Nisi, layers invariably contain two types of particular defects, as shown in Fig. 9. The short bright streaks and the dark lines seen in Fig. 9 are what is known as ‘facet bars’ and ‘twin-related defects’, respectively [52, 105, 1601.

Fig. 9. Planview, (020) dark-field, TEM images of a Nisi, layer formed by deposition of 15 A Ni at room temperature followed by annealing at 500 “C. Twin-related defects and facet bars are clearly revealed.

118

Facet bars, which are parallel to either [Oil] or [011] directions, are trough-shaped protrusions of Nisi* into Si bounded by (111) planes. Typically, they are roughly SO-100 A in width and a few hundred 8, in length [105]. The orientation of the facet bars is determined by the orientation of the Si bonds at the NiSiJSi interface. The projection of the Si bond orientation at the NiSi,/Si(lOO) interface alternates between a [Oil] direction and a [Oil] direction for regions of interface differing by an odd number of 1/4(100)a, interface steps. As a result, facet bars which ‘reside’ on the same interfacial plane have the same orientation while facet bars residing on planes differing in height by an odd number of step height changes have perpendicular orientations. A region of interface on the same interface plane forms a single domain; thus, within a single domain, the facet bar orientation is the same, see Fig. 9 [52, 10.51. Facet bars are commonly observed in virtually all films of Nisi, on Si(100) prepared by Ni deposition and annealing. The possible mechanism for the formation of facet bars and their elimination will be discussed later. The dark line contrast in the (200) dark-field images shown in Fig. 9 appear as bright lines under (200) bright-field imaging conditions. Earlier investigations have proposed these defects, which form closed loops, to represent an absence of NiS&, possibly consisting of trenches between domains of Nisi*, exposing the Si [105]. However, the expression ‘coreless defect’ [105] turns out to be a misnomer because more recent analyses by cross-sectional and conventional TEM have shown these defects to be related to twinning in the Nisi* film [106]. These twin-related defects border regions of interface differing by an odd number of step height changes, thereby delineating the different interface domains. Due to the difference in crystal symmetry between the Nisi* lattice and the Si lattice, an interface dislocation with a Burgers vector related to 1/4(111)a, would be created at the position of the interface step if the domains were not separated by twin-related defects [107]. Dislocations of l/4( 11 l)a, character have indeed been observed experimentally at the interface of thick Nisi* films grown by Ni deposition and annealing [75, 1081. Thin Nisi, layers grown by co-deposition do not contain twin-related defects even upon extended annealing at high temperatures [71], discussed below, suggesting that the formation of these defects is not a result of energetics alone. The deposition of 5 8, of Si, followed by annealing at -500 “C, causes twin-related defects to be eliminated from thin Nisi:! layers [106]. Presently,

the formation of twin-related defects may be regarded as a by-product of the particular (lateral) reaction kinetics involved in the growth of Nisi;! from the deposited Ni layer. The conversion of the twin-related defects to the energetically more favorable dislocations seems to be met with a significant kinetic barrier under usual growth conditions [106]. The presence of silicon seems to reduce the kinetic barrier, a fact which is suggestive of some local non-stoichiometry at the twin-related defects [106]. The structure of a twin-related defect is still not exactly known. Co-deposition of Nisi, at room temperature followed by annealing is a more reliable technique which allows the fabrication of uniform NiSiz layers with good crystalline quality [71]. The minimum annealing temperature for the formation of epitaxial NiS& is - 400 “C, similar to that for pure Ni deposition. Twin-related defects and facet bars are absent in uniform Nisi* layers grown from stoichiometric co-depositions, as shown in Fig. 10(a). Line defects seen at the interface of these layers are all related to dislocations with b= 1/4(111)a, associated with steps at the interface. Uniform layers with thicknesses as small as - 20 A have been observed. The morphology of codeposited NiS& layers is affected by the stoichiometry of the co-deposition. With Ni-rich codeposition, facet bars (and twin-related defects) are occasionally observed. With Si-rich co-deposition, small rectangular-shaped pinholes in the NiSiz layers are sometimes observed. 5.3 Growth of thicker NiSi, on templates Starting with a thin Nisi, template, it is relatively easy to grow a thicker layer by further deposition

Fig. 10. Planview, (022) weak beam, TEM images of two -28 8, thick NiSi2 layers on Si(100) grown by co-deposition at room temperature and annealed at (a) 500 “C and (b) 700 “C, respectively.

119

and annealing. Obviously, one may deposit Ni, layers of Ni and Si, or (co-deposited) Nisi, on existing Nisi2 templates at a certain deposition temperature, followed by, if necessary, annealing to allow thicker Nisi2 to grow. It was shown very early that Ni deposition on thin Nisi2 templates at elevated temperatures (RDE) leads to the growth of thick uniform Nisi2 layers [64]. However, the success of the RDE process requires higher temperature, > 650 “C, and thicker Nisi2 templates. Since better layer uniformity may be achieved with a lower deposition temperature, it is more advantageous to adopt a method which does not require such a high growth temperature. It is known [109] that co-deposition on Nisi2 templates on Si(100) leads to the (homo)epitaxial growth of Nisi2 at temperatures even as low as room temperature. Therefore, the growth of thicker Nisi2 layers by the co-deposition of stoichiometric Nisi, at low temperatures, followed by, if necessary, additional anneals is a more reliable technique for the fabrication of thicker Nisi, layers [71]. Optimum Nisi2 films are obtained by co-deposition of stoichiometric Nisi,. 5.4 Il4anipulation

of NiSi,

layer morphology

Annealing thin films of Nisi,, grown at <500 “C, at temperatures above 600 “C may cause, in general, three pronounced changes in the film morphology: the density of twin-related defects/ dislocations in the thin Nisi, layer decreases, the density of facet bars, if they had been present, decreases, and in some films pinholes may develop, especially for thin films annealed to very high temperatures. Epitaxial NiSi,( 100) layers, even as thin as 30 A, are remarkably stable upon annealing to high temperatures provided the layers were continuous and pinhole-free upon initial growth. Annealing of Nisi, films at >700 “C leads to a significant improvement in the quality of the film, as shown in Fig. 10(b). The increased domain size (and resulting decreased dislocation density) indicate the presence of a much flatter and more uniform interface than that obtainable through low temperature annealing. Dislocation densities less than lo8 cm-’ are obtainable following annealing at high temperature. RBS and ion channeling experiments clearly demonstrate the increased crystalline quality obtainable with these films. Nisi, layers annealed at 500 “C show typical but annealing at higher temperXmin’Sof -lo%, atures reduces the observed Xmin substantially (to typically 3%) [71]. For films which do not contain pinholes but which do contain facet bars following a low temperature anneal, e.g. the film in Fig. 9,

annealing in the range 700-750 “C usually causes facet bars to be eliminated concurrent with domain growth and a reduction in dislocation density [71]. Annealing Nisi, films at high temperatures may also lead to a degradation of the layer uniformity. Films, which contain pinholes or Si facets after a low temperature anneal, break-up upon annealing to moderately high temperature (> 750 “C). In contrast to what was observed for CoSi, on si(ll1) [llO], the deposition of excess Si on Nisi2 films or Si-rich co-deposition on Si(100) promotes pinhole formation upon high temperature annealing, rather than suppressing it. The high-temperaturestability of pinhole-free Nisi, films de ends on the film thickness. Films as thin as 30 w may be safely annealed to 750 “C, but invariably breakup upon extended annealing longer than one minute) at 800 “C. Films - 100 6 thick are stable against the formation of pinholes for at least up to 850 “C. When a larger amount of Ni, one which is comparable to that already contained in an existing Nisi, layer, is deposited on the thin Nisi, layer and annealed to -500 “C, a high density of facet bars is grown at the interface [71]. This results in nearly completely faceted structures when the starting Nisi, layer has a high density of domain boundaries, as the example of Fig. 8(c) shows. When the starting Nisi2 film has a lower density of dislocations, such as a film which had already been annealed at 750 “C, the effect of deposition of a large amount of Ni, and annealing is to cause facet bars to form at the domain boundaries of the film, while preserving the general domain size and distribution. An examination of Figs. 8-10 shows that the morphology of thin Nisi, layers may be varied (from being nearly completely faceted to being very uniform) by employing additional depositions and anneals in UHY. This rare feat was accomplished because there is a good understanding on the reason for the formation of facet bars. Through various annealing and deposition experiments, including experiments carried out inside a TEM, it was clearly shown [71] that the formation of (111) facet bars is related to the kinetics of Nisi2 growth and, therefore, is only a metastable configuration thermodynamically. Apparently, ~{lllfX

43

>

~{loo}

(1)

and the thermodynamic driving force is towards the formation of flat, facet-free (100) interfaces for Nisi, on Si(lOO), which, as a result, are invariably observed after high temperature anneals.

120

5.5 Su$ace Si(lO0)

and interSace structures of NiSiJ

The NiSi,/Si(lOO) surface has a complicated reconstruction. In addition to the bright integral order spots of LEED, additional spots are observed between the integral order beams and have been labeled as of the l/3, the l/4, the l/5, and the l/10 orders [30], as shown in Fig. 11(a). A LEED I-V analysis was performed on the integral-order spots alone, and a model was proposed consisting of Si adatoms atop four-fold sites of a Si bulklike termination [ill], while a low energy ion scattering study concluded that the surface was bulk-like Si terminated with a 25 to 30% vacancy concentration [112]. Recently, however, it was discovered that carefully annealed surfaces of high quality NiSi,/Si(lOO) layers display robust LEED patterns consisting of two rotated domains of a &X. &&R30.96” reconstruction (the angle of rotation is arctan(3/5)). Under suitable annealing conditions (annealing at 550-600 “C for >4 hours in UHV), all fractional-order spots of these two \/?;;i reconstructions are sharp and clearly resolved, as shown in Fig. 11(b), indicating the presence of large, well-ordered domains. STM shows that this reconstruction arises from an ordered arrangement of double rows of adatoms [113]. When the surface is less well-ordered, which results from annealing at over 650 “C or from anneals for short periods of time, the LEED patterns change to the more customary Fig. 11(a), which, in reality, is nothing more than a disordered &4 pattern. The atomic structure of the planar NiSi,/Si(lOO) interface has been studied experimentally by HREM [64, 1081. Rigid shifts of the phase contrast HREM images were interpreted in terms of a sixfold coordinated structure for the NiSi,/Si( 100)

Fig. 11. LEED patterns, with an electron energy of 63 eV, of a - 100 8, thick NiSir layer grown on Si(100). (a) After annealing at 750 “C for 1 minute, displaying the pattern usually associated with the NiSi,(lOO) surface, and (b) after for 12 hours, displaying two domains of a

interface, as depicted in Fig. 7, even though these experimental results were also consistent with an g-fold interface model. An early dismissal of the eight-fold coordination by HREM was invalid because an unphysical model of the interface atomic structure was assumed [lo]. Recent theoretical calculations are suggestive that the coordination number at this interface should be higher than six [114, 1151. In addition to the uncertainty in the 6-fold model, there is now new evidence for the presence of additional structures at this interface [71]. Transmission electron diffraction (TED) patterns of uniform, well-annealed NiSiz layers (Fig. 12) reveal extra diffraction spots not related to the silicide or Si lattice. In addition to the (022) diffraction spots arising from both Si and Nisi,, there are spots at (020) arising from NiS& alone, and spots at (011) and near (0 44) which are not kinetically allowed in bulk Si or NiS&. The (011) diffraction can be explained by reciprocal lattice rods at the (111) reflections and is indicative of a flat and abrupt NiSiJSi interface. The diffraction near (0 4 3) which is most intense in films annealed at higher temperatures, can only be explained by superstructure(s) at the NiSiJSi interface. When the Nisi* is not thoroughly annealed or when the Nisi* layer is thick, diffuse spots or streaks are observed at positions related to both (0 $4) and (0 4 I), originating from domains related by a 90” rotation (Fig. 12(a)). A higher temperature and/or a longer time anneal leads to the appearance of discrete spots slightly displaced from the (0 4 4) and (0 Q$) positions. By limiting the diffraction to only one domain of NiSiz, only one set of the superstructure spots is seen, see Fig. 12(b). The diffraction spots near (0 34) are

Fig. 12. Transmission electron diffraction pattern taken from NiSiz films, 50-70 8, thick, on Si(100); (a) annealed at 700 “C for - 1 minute, (b) annealed at 800 “C for -30 minutes. Arrows point to positions equivalent to (O&f).

121

reminiscent of the sharp (0 t 4) spots observed in TED patterns from CoSiJSi(100). In the case of CoSiz on Si(lOO), these spots have been correlated to a model of dimerized chains of excess Si running along (011) directions [116] or, possibly, to a missing row model 11177. It is likely that a similar reconstruction exists at the NiSi,/Si(lOO) interface. The streakiness at the (0 4 4) positions in Fig. 12(a) is consistent with narrow domains with a distribution in width and/or separation. At present, the details of this interface reconstruction are not understood.

It has long been shown that the NiSi~/Si(lOO~ interface has a much lower SBH on n-type Si than the SBHs at NiSi,/Si(lll) interfaces [1X3-120]. However, with the omnipresence of facet bars in the early studies, no conclusion on the SBH could be drawn [119]. Early results presented by Kikuchi et al. [95] actually suggested that the SBH of Nisi*/ Si(100) may not be different from that observed at the type A NiSi~/Si(lll) interfaces after all. These discrepancies were resolved in a recent study [121] on the SBHs of NiSi,/Si(lOO) interfaces with a variety of morphologies. It was demonstrated that the SBH measured at a NiSi,/Si(lOO) interface depended very much on the observed mo~holo~ of the particular layer. Interfaces which are almost completely (111) faceted were shown to have SBHs similar to that found at a type A NiSi,/Si(lll) interface [121], as expected. This result is in good agreement with Kikuchi et at. [95], since the growth procedures used by Kikuchi et al. [95] invariably lead to nearly completely faceted interfaces. Uniform NiSiJSi(100) interfaces show a SBH which is much lower on the n-type Si than either of the two NiSi,/Si(lll) interfaces. Since the Nisi*/ Si(100) interface has an entirely different atomic structure from either of the two NiSiJSi(ll1) interfaces, it is perhaps not surprising that the SBH is also different. One notes that the low SBH of 0.4 eV measured from uniform NiSiz layers is also very different from the value of 0.6-0.7 eV usually observed for all phases of polycrystalline nickel silicides on Si fl22, 1231. On n-type substrates, ideal&y factors for I-V measurements are good (n < 1.03) and very consistent results were obtained from either I-V, C-V, or activation energy studies 11211. There is no observed dependence of measured SBH on the doping level [121]. However, on p-type Si(lOO), planar Nisi, diodes showed evidence for a slight inhomogeneity of the SBH, such as large ideal&y factors, n 2 1.08, and discrepancies between SBHs measured by I-V and

C-V techniques 11211. Therefore, even very uniform NiSi,/Si(lOO) interfaces may in fact be electrically inhomogeneous. The origin of this SBH inhomogeneity is likely related to the 1 x 2 reconst~~tion which has been shown to be present at this interface [71]. The presence of a few facet bars at NiSi,/Si(lOO) interfaces, which are otherwise flat, has little effect on n-type SBH, but has a strong influence on the measured SBH on p-type Si. Typical I-V results from a faceted and a uniform NiSi,/p-Si(100) diode are shown in Fig. 13 to illustrate the dependence of SBH on interface morphology. The I-V deduced p-type SBH decreases rapidly as the density of facet bars increases, while a slower, but noticeable, decrease of the C-V SBH is concurrently observed. As a result, the C-V measured SBH for any specific diode significantly exceeds that deduced for I-V. aged-mo~holo~ p-type diodes are ‘leaky’, having poor ideality factors YE > 1.1 in forward bias and displaying reverse currents which do not saturate, There is also a clear dependence of the electron transport on the substrate doping level. As an example, Nisi, layers with similar densities of facet bars, but grown on p-type Si with different doping levels, show almost identical SBHs as determined by C-V, but a sharp decrease of the SBHs determined by I-V, with doping level. These experimental results are a clear indication of an inhomogeneous SB [94]. Applying the analytic theory for inhomogeneous SBH, a semiquantitative I

i

Nisi,

I

I

I

on p-type Si(lO0)

1

.

T = 300 K

t

I*

lo-i--i FORWARD

BIAS (Volts)

Fig. 13. L-V plots of NiSi,/p-Si(100) SB diodes, showing the dependence of the measured SBH on the density of facet bars observed for the layer.

122

explanation of these electrical data have been demonstrated [121, 1241. Because of the uncertainty concerning the atomic structure at planar NiSi,/Si(lOO) interfaces, the electronic properties of this interface have not been conclusively studied with theoretical means. Very recent calculations by Fujitani and Asano [114] showed that the SBH at an 8-fold planar NiSi,/Si(lOO) interface resembled that of the type A NiSi,/Si(lll) interface. However, very different redistributions of charge were found for these two interfaces [114].

6. Epitaxial

NiSi,/Si(llO)

structures

The growth of single crystal NiSiz on Si(ll0) is very easy, however interface faceting is an unavoidable problem [125]. Due to the poor quality of epitaxy and the lack of obvious applications, the NiSi,/Si(llO) epitaxy has not been carefully studied. Essentially uniform NiS& layers may be grown by deposition of Ni and Si, or by codeposition of Nisi*, at room temperature and annealing at -500 “C. However, upon annealing to high temperature (>550 “C), the interface breaks up into inclined (111) facets. Indeed, interfacial faceting of this system is so complete that, in most well-annealed samples, the entire NiSiJSi interface is made up of inclined (111) facets, as shown in Fig. 14. At the NiSiJSi(ll0) interface, a 1/4(111)a, type dislocation is required at the boundaries between any two (111) facets [125]. The fact that not even a small portion of flat 11101 NiSiJSi interface has been observed suggests that this interface may well be unstable. Since the NiSi,/Si(llO) interface is actually made up of (111) facets, it is not surprising that a SBH similar to that of type A NiSiJSi(ll1) interface, - 0.65 eV on n-type Si, is usually found [118, 1261.

Fig. 14. Planview, (002) dark-field, TEM image of ‘a -7OAthick NiSiz layer on Si(ll0). Long streaks running aIlong the [liO] direction indicate near complete (111) faceting.

7. Epitaxial

CoSi,/Si(lll)

structures

7.1 Early work Despite a significant lattice mismatch, CoS& was the first metal to be grown with a pure and single orientation on a semiconductor. Thick layers, > 700 A, of CoSi, with type B orientation were grown by SPE at -1000 “C on Si(lll), under UHV conditions [34,35]. Co-deposition of stoichiometric CoSi, at 650 “C also led to the growth of pure type B CoSi, layers on Si(ll1) [35]. Even though layers have been demonstrated which are single crystal and essentially uniform over large areas, they always contain structural imperfections which reduces their usefulness. First of all, the density of misfit dislocations in CoSi, layers grown by SPE and MBE is always quite high, with an average separation of N 30 nm between dislocations. Secondly, both the MBE- and the SPE-grown films contain a network of coarser defects, the ‘tartan’, which stimulated some discussions concerning the nature and the origin of these defects [127, 1281. In addition to line defects, a density of randomly distributed large pinholes, typically - 100-300 nm in size, were always present, even for layers more than 700 8, thick. The surface cleanliness and the vacuum condition are usually found to significantly influence the crystalline quality of SPE-grown epitaxial CoSi, [35]. However, careful annealing procedures could lead to the growth of single crystal CoSi, even under modest vacuum conditions [129]. Furthermore, by employing pulsed laser annealing, single crystal type B CoSi, layers could be grown without any elaborate vacuum processing [45]. 7.2 Room temperature ColSi(Il1) reaction The reaction of a few ML Co with a clean Si(ll1) surface has been investigated by many groups. There is good agreement among most of the studies that a uniform, e&axial, tvpe B oriented, CoSi, layer is grown when l>*ML Co is deposited onto clean 7x7 Si(ll1) [130-1321. A bright, although diffuse, 3-fold 1 x 1 LEED pattern, similar to the CoSi,-C pattern of well annealed CoSi, layers, is observed from the room temperature reacted silicide layer. The existence of good long range order for the room-temperature-grown Co/% surface is in sharp contrast to the Ni/Si surface at similar metal coverages. There is, however, evidence for disorder in ultrathin CoSi, layers grown at room temperature [133]. When more than 5 ML Co is deposited at room temperature, the surface becomes more cobalt rich than the CoSiZ phase. At more than 10 ML Co, spectroscopies show evidence for unreacted cobalt near

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the surface. However, recent studies suggest the silicide phase present at this coverage might be a distorted Co,Si phase [132-1341. The homogeneity of the Co/Si surface structures at room temperature is in contrast to the rough morphology observed for Ni/Si. Co-deposition of CoSi, on Si(ll1) leads to the growth of epitaxial CoSi, at room temperature, up to a thickness of - 40 8, [135]. If co-deposition continues beyond this thickness, amorphous silicide begins to grow on top of the epitaxial CoSiZ layer [135]. It was discovered, however, that the deposition of a small amount of Co at room temperature prior to CoSi, co-deposition leads to a very substantial improvement in the quality of the epitaxial type B CoSi, layers [136]. As illustrated in Fig. 15, the crystalline quality of epitaxial CoSi, is the very best when - 2-3 A Co is pre-deposited. Ion channeling measures a typical Xminof 2% [ 1361, indicating the very high crystalline quality of CoSi, layers grown with Co pre-deposition. Actually, the quality of these CoSi, layers is so high that the only extended defects seen at the interfaces are those required by symmetry, due to steps on the original Si(ll1) 7X 7 surface. This property has been utilized very successfully to study the topography of the Si(ll1) surface [137, 1381. With a proper Co pre-deposition, the growth of epitaxial CoSi,, from co-deposition at room temperature is not limited to any finite thickness. Epitaxial layers as thick as a few thousand angstroms have been demonstrated.

Fig. 15. Planview dark-field TEM images of four CoSi, layers grown by co-deposition of -72 8, Co%, at room temperature. Prior to co-deposition, 1 A, 3 8, and 5 8, Co was deposited onto clean Si(ll1) for (b), (c) and (d), respectively. No pre-deposition was used in the growth of layer (a).

7.3 Growth of thin Co& layers on Si(ll1) In order to grow uniform CoSi, layers, a low deposition temperature is preferred. Deposition of Co onto Si at temperatures above -400 “C leads to the spontaneous growth of isolated CoSi, epitaxial islands. When co-deposition is employed, much more uniform CoSi, layers may be grown with a clear preference for type B orientation. Thick CoSi, layers which are essentially uniform may be grown by co-deposition at temperatures less than -550 “C. However, the morphology of ultrathin (<50 A) CoSi, layers grown at elevated deposition temperatures is generally inferior to that grown by deposition at room temperature. Furthermore, without a pre-existing template, it seems unlikely that deposition or co-deposition (of any composition) at temperatures >650 “C could lead to thin CoSi, with a uniform layer thickness. As an extreme example of deposition at very high temperatures, one may point to the deliberate growth of CoSi, columns with co-deposition using a Si-rich stoichiometry [139]. Using SPE and low deposition temperatures, the uniformity of CoSi, layers is significantly improved. Because of the simplicity of the process, SPE from deposited Co layers [140-1421 has been the most popular technique of epitaxial CoSi, growth. With moderate annealing tern eratures (< 500 “C), ultrathin CoSi, layers (< 20 x ) grown from deposited Co are essentially continuous but contain a high density of large pinholes [142 . When the CoSi, film thickness exceeds -40 1 the density of dislocations increases rapidly and the size and density of pinholes generally decrease [142]. With more than - 15 A of deposited Co, an epitaxial CoSi phase is found to precede the epitaxial CoSi, phase from SPE [141-1431. Early reports [141, 1441 of a two-dimensional membrane phase as an intermediate reaction product turns out to be nothing more than one of the two surface structures [145, 1461 CoSi,-C, of CoSi,(lll). Insitu TEM studies demonstrated that for deposited Co exceeding 15 A in coverage, all CoSi, phases are grown via lateral reaction fronts which sweep across the surface [132]. Dislocations are involved, from the very beginning, in the SPE growth of CoSi, in this thickness range [ 1321. This high density of dislocations explains the absence of large strain in even the very thin ( < 100 A) SPE-grown CoSi, films [147-1491 CoSi, layers grown from deposited Co often contain a fraction of type A orientation, which has been suggested to be related to the intermediate phases [142]. The conversion of type A grains into type B CoSi, at times may be related to the process of pinhole formation in the layer

124

[132]. Even though single crystal CoSi, may be grown by SPE from deposited Co, pinholes are almost always present in these films. It has not been possible to avoid pinholes in epitaxial CoSiz layers of any thickness, grown by Co deposition alone, at annealing temperatures of 650 “C. An early claim [150] of the growth of thin CoSi,, layers at 650 “C which are dislocation-free and pinholefree was obviously in error, as more careful studies which employed TEM clearly demonstrated [151, 1521. Much better crystalline quality of CoSi, may be achieved by employing Si deposition in its growth. With the deposition of both Co and Si, the orientation of the CoSiz layer becomes exclusively type B and the density of pinholes in the CoSi, layers is significantly reduced. The sequential deposition of Co and Si at room temperature followed by annealing to > 450 “C has been shown to lead to high quality Co%,, layers [151, 1521. The empirically determined optimal Co:Si ratio of sequential deposition for the growth of CoSi, layers with the best crystalline quality varies slightly with studies, ranging from -1:l [151, 1341 to -1:1.6 [152]. Co-deposition of CoSi, at room temperature followed by annealing has also been shown to lead to uniform CoSiz layers [153, 1541. Since co-deposition of CoSi2 already leads to the growth of epitaxial CoSi, at room temperature [136, 135, 1541, subsequent anneals serve to improve the crystalline order of these (partially) epitaxial CoSi, films [153, 1541. It should be noted, however, that layers grown by co-deposition at room temperature are very defective and contain a very high density of dislocations [154, 1351, as shown in Fig. 16(d). It seems that a precise control of the deposition parameters is required for the growth of high quality CoSi,. With more deposited Si, there is less danger of pinhole formation, but the density of dislocations is increased. The formation of pinholes in CoSi, layers was initially attributed to a smaller surface energy for the Si(ll1) than that of the CoSi,(lll) [155]. The formation of pinholes has also been described as an ‘unwetting’ process driven by the high interface energy. Surface and interface energies obviously are the reason for pinhole formation in CoSiz films. However, straightforward models [155] are not consistent with either the observed growth of Si islands on CoSi* [156] or the observed thermal stability of CoSiz layers with the CoSiTS surface structure. It has also been proposed that pinholes may be created in C&i, layers to relieve misfit stress [142]. Careful measurement of the lattice parameters with and without a high density of pinholes, however, showed

Fig. 16. Planview, bright-field, TEM images of ultrathin CoSi, layers with identical average thickness (N 14 A), grown on Si(ll1) from the deposition of (a) 4 8, Co, (b) 4 8, Co+4 8, Si, (c) 2 8, Co+ 7 8, co-deposited CoSi,+4 8, Si, and (d) 14 A &-deposited Co%,+4 A Si. Subsequent to the room-temperature depositions, an anneal was carried out at -500 “C for 5 minutes.

that pinholes have no effect on the measured film strain [134]. The discovery of the presence of two quasi-stable structures of the CoSi,(lll) surface [145, 1461, the Si-rich CoSi,-S and the bulk terminated CoSi,-C, eventually led to an identification of the main driving force for pinhole formation in epitaxial CoSiz films [llO], namely, the energetic difference between these two surface structures. Both surfaces display 1 X 1 periodicity, although they clearly have different atomic structures [146] (see Fig. 18 below). The surface of CoSi, layers grown by a Co-rich deposit and an anneal at low temperature usually has the CoSi,-C structure, which is less stable than the CoSi,S structure. Upon annealing to higher temperatures, other local thermodynamic equilibria are kinetically accessible. The energetic difference between the two surface structures drives the formation of pinholes and the subsequent Si surface diffusion, both of which are necessary for the achievement of the more stable CoSi,-S structure. Therefore, the formation of pinholes is kinetically the easiest process under the driving force of a reduction of surface and interface energies. Since the deposition of - 2 ML Si on Cc&-C leads to the CoSi,-S structure on a CoSi, layer, this procedure essentially eliminates the formation of pinholes in stoichiometric CoS&. Central to the argument given so far is the assumption that the rate of Si from the crystalline substrate being released and diffusing through the epitaxial CoSiz layer is low. The actual rate obviously depends on the quality of the existing CoSi, layer. A recent

125

speculation of Si diffusion through CoSi, seems to have dwelled too much on circumstantial evidence to be reliable [157]. An STM investigation of the surface morphology of CoSi, layers [158] showed that the Si exposed at the position of pinholes in CoS& is significantly indented, indicating a massive amount of surface diffusion had taken place. Such a profile of pinholes in CoSi, is in good agreement with the idea that pinholes serve as Si sources [llO], and is hard to reconcile with models based purely on energetics. Thus, high quality CoSi, layers which are both free of pinholes and low in dislocations can be grown by first forming CoSi, layers at low temperatures (< 500 “C) under conditions which minimize dislocations. However, before this CoSi, layer is annealed at high temperature for improved crystallinity, a suitable amount of Si should be deposited to adjust the surface structure to the stable CoSi,-S, in order to avoid pinholes. Dislocations in very thin CoSi, films grown at room temperature by pre-deposition and co-deposition are only those required by symmetry. Upon annealing to > 550 “C, thicker films may relax by the introduction of additional dislocations. With very precise stoichiometry control, it is possible to avoid the introduction of additional dislocations in very thin films ( < 30 A) during the high temperature anneal, as shown in Fig. 17. However, these highly strained thin films often have a novel form of lateral distortion (the ‘R’). When this

transformation occurs, it involves only about 3 triple layers of CoSi, at the interface which go through a seemingly shear-induced crystallographic is change at below - 100 “C. This transformation reversible as the fluorite lattice is recovered at over 100 “C through annealing experiments inside the TEM [159]. The most prominent demonstration of this phase transformation is by planview TEM observation of patches of different contrast under weak beam conditions [160, 1611 as shown in Fig. 17. The three equivalent azimuthal directions (the three [2111’s) for the shear to occur lead to different phase shifts with respect to the operating diffracted beam, which leads to the observation of the domainlike structure in TEM [159]. Other evidences for this interfacial transformation have come from cross-sectional HREM [159], which shows the presence of an interfacial layer, Rutherford backscattering (RBS) and channeling, which shows an interface dechanneling peak, and from X-ray scattering, which show extra diffraction spots [160, 1591. By a slight change in the procedures of silicide growth, it is possible to prepare CoSi, layers which either undergo, or do not undergo such a phase transition upon cooling to room temperature. The transformation seems to correlate with a Si-rich stoichiometry at the interface and it may be removed by ion beam bombardment [159, 1621. Presently, the structure of this transformed phase at the CoSi,/Si interface is not fully understood.

Fig. 17. Planview, (220) weak beam, images of -23-29 8, thick type B Co%, layers which had been annealed at 600 “C. These layers were originally grown at room temperature by co-deposition of -18 8, CoSi, and pre-depositions of (a) 1.5 A Co, (b) 2 8, Co, c) 2.5 8, Co, and (d) 3 8, Co. Prior to annealing, a thin -4 k Si layer was deposited on the surface of the CoSi,.

7.4 Growth of thicker layers When thick (> 100 A) and uniform CoSi, layers are desired, the most reliable technique is to use co-deposition at low temperature on an existing thin CoSi, template (which could mean, simply, an appropriate pre-deposition), followed by annealing. Growth of thick CoSi, layers in multiple steps always leads to better crystalline quality and layer uniformity than those grown in one step. If a thin CoSi, template is already present, deposition of layers of Co and Si in the 1:2 ratio, followed by annealing leads to the growth of uniform CoSi, layers at >400 “C. Clearly, if co-deposition is employed, epitaxial CoSi, is grown even at temperatures as low as room temperature [136]. For improved electrical conductivity, it is sometimes desirable to anneal thick CoSi, layers to > 650 “C, in which case the surface should be adjusted to CoSi,-S before the anneal to prevent pinhole formation. The RDE technique does not work for CoSi,/Si( 111) because, as already discussed, deposition of cobalt on CoSi, templates leads to pinholes and rough surface morphologies. Thick

126

type B CoSiz layers have a network of honeycomb shaped misfit dislocations with 1/6(112)a, character [35]. Films which have been annealed to > 600 “C are usually nearly fully relaxed [ 134, 1631. 7.5 Su$ace Si(llI)

and inte$ace

structures of CoSiJ

After some initial discussions [145, 164, 1651 there seems now to be a good agreement amongst various groups [146, 166, 1671 regarding the structure of the CoSi,/Si(lll) surface(s). As already discussed, at least two quasi-stable structures, differing in the ‘surface stoichiometry’, have been found for the CoSi,(lll). The bulk-terminated C&i,-C, shown in Fig. 18(a), is often observed on top of CoSi, layers grown at low temperatures, < 500 “C, from Co-rich deposits (such as deposited Co) [146]. The Si-rich CoSi,-S structure, shown in Fig. 18(b), is observed on all CoSi, layers which have been either annealed to high temperatures, >600 “C, or grown from a Si-rich deposition. Deposition and annealing experiments show that these two structures differ by two monolayers of Si. Deposition of 2 ML Si on a CoSi,-C surface leads to CoSi,-S, while deposition of 1 ML Co on a CoSi,-S surface leads to the CoSi,-C surface [146]. Annealing a CoSi, layer with a CoSi,-C surface at over 600 “C also transforms the surface into a CoSi,-S, although this transformation is accompanied by the creation of pinholes in the CoSi, layer [llO]. Finally, one notes that the stacking sequence of the outermost Si bilayer on the CoSi,-S, shown in Fig. 18(b), is the same as that of the CoSi,, lattice. This is the reason for the orientation of the overgrown Si being identical to that of the Co&, starting from the CoSi,-S surface [ 1681. The atomic structure of the type B CoSiJSi(ll1) interface, despite years of investigations, is not fully resolved. Early studies by HREM [169] XSW

7.6 Electronic properties of CoSiJSi(ll1)



(a)

[147,148] and MEIS [170] compared experimental results with just two structural models of the interface, the 7-fold and the 5-fold models, and concluded that the HREM images were consistent with the 5-fold model [169]. However, this conclusion [169], which had since received various confirmations [147,148,170], is now deemed questionable because all the early studies had neglected to consider the possibility of the interface having an 8-fold coordination. Since the rigid shifts associated with the 5-fold model are indistinguishable from that of the 8-fold model, early experimental results were, in fact, also in agreement with the 8-fold model. Two theoretical papers [5, 1711 argued that the high interfacial free energy for the 5-fold model should make it unlikely to occur in nature. The 8-fold model was shown to have lower energy than even the 7-fold model [5, 1711, and, therefore, is like the structure experimentally observed. A SEXAFS study claimed to provide direct evidence for 8-fold coordination at this interface [172], but these results are most probably fortuitous because of the very poor preparation conditions chosen for the growth of silicide layers [173]. Detailed investigations by HREM showed that the 8-fold model is more likely the structure experimentally observed [174, 1751. However, it is also known that the interfacial structure of type B CoSiJSi(ll1) may vary according to preparation - evidence for 7-fold coordinated structure has been obtained from layers which have only been annealed to low temperatures, ~500 “C [162]. Annealed type B CoSi, layers often contain other mysterious defects [162], which are likely related to variations of interface atomic structure. In addition, the noted ‘R’ transformation at some type B CoSiz interfaces are certain to affect the atomic structure. Type A CoSiJSi(ll1) is thought to be 7-fold coordinated [174, 1751, although 8-fold coordinated sections have also been observed occasionally [ 1751.

Si

0’)

Fig. 18. Ball and stick models of the two surface structures of CoSiz, viewed in the [liO] direction. (a) CoSi,-C and (b) CoSi&.

The structure of type B CoSi,/Si(lll) is complicated because of a possible phase transformation, a possible variation of the atomic structure, and a high density of defects. The SBH of type B CoSiz layers grown at - 600 “C by SPE is usually in the range 0.65-0.70 eV on n-type Si [176, 1771. It seems reasonable to attribute this SBH to an interface which has the 8-fold structure, assuming the effect due to dislocations can be ignored. On the other hand, CoS& layers which have been grown at lower temperatures and which have a low density of dislocations show considerable vari-

127

ation in their SBH [162]. SBHs a slow as -0.40 eV have been observed from these layers on ntype Si(ll1). Since a variation of atomic structure and the existence of a phase transformation at this MS interface have already been suggested by experiments, it is perhaps not surprising that the electronic properties of the type B CoSiz interface show such variations. However, these results have not been consistent enough to allow a conclusion to be drawn. An early tight-binding calculation showed that the SBH of type B CoSi,/Si(lll) depends on whether a 5-fold or an S-fold structure is assumed for the interface atomic structure [178]. Calculations involving large supercells [179, 1801, showed that, on n-type Si, the 7-fold coordinated interface should have a SBH which is higher than the 8-fold interface. 8. Epitaxial

CoSi,/Si(lOO)

structures

8.1 Epitaxial growth of thin film CoSi,lSi(lOO) Unlike Nisi* epitaxy on Si(lOO), faceting is not an issue for epitaxial CoSi, on Si(100). Instead, the major issue is multicrystallinity. There are a few competing orientations for CoSi, epitaxy on Si(100). CoSi,( lOO)/Si( 100) and CoSi,( lOO)// Si(100) have frequently been observed [181-1831. There are two variants of the [llO] epitaxy, related by a 90°_rotation, namely, CoSi,(llO)//Si(lOO) with Co&[ 1~O]//Si[Ol~] and CoSi,( 1lO)//Si( 100) with Cc&[ llO]/!Si[Ol l] [ 1831. In addition, variants of a CoSi,(221)NSi(lOO) with CoSi,[llO]//Si[Oll] orientation, which is likely a result of twinning (type B orientation) along inclined (111) planes, have been observed [184]. Misoriented grains are easily identified by TEM as they appear brighter than the background intensity in (020) dark-field [183]. Various room-temperature deposition schedules have been studied in order to reduce the fraction of misoriented, (110) and (221), grains in epitaxial CoSi,(lOO) films, [183-1851. It is discovered that wafer roughness and surface cleanliness also has an effect on the observed area1 fraction occupied by misoriented grains [183, 1841. Presently, the reason for the nucleation of misoriented grains is not clear. It is possible to completely eliminate misoriented grains using deposited cobalt and silicon, as shown in Fig. 19, which reveals a high quality, single crystal CoSi, layer grown on Si(100). However, it is noted that the control of misoriented grains is still difficult due to poor reproducibility from run to run [183,184]. Best results are obtained from precisely oriented Si(100) wafers which have been carefully prepared by Si buffer layer growth

Fig. 19. Planview, (022) weak-beam, TEM image of a -48 8, thick single crystal CoSi, layer, grown in two steps on Si(lOO), with a final annealing at 600 “C.

[183, 1841. Most of the dislocations seen in (lOO)oriented CoSi, areas are 1/4(111)a, in character, related to single interface steps. A careful study of the deposition parameters revealed that purely (lOO)-oriented thicker ( > 300 A) CoSi, layers may be grown at -500 “C with the co-deposition of CoSi,, [184, 1851. With a lower deposition temperature or a different codeposition stoichiometry, mixed orientations are found in the CoSiz films. This discovery seems to violate the general ‘optimum’ nucleation conditions found for all epitaxial silicides discussed so far, namely, low deposition temperature and subsequent annealing for maximum layer uniformity. However, there is no contradiction because the reason for the elimination of mixed orientation in CoSi,(lOO) films is likely a post-nucleation process due to the kinetic advantage of the (100) oriented grains [184]. For this reason, even though thick single crystal CoSi, layers may be grown by co-deposition at 500 “C, it is unlikely that uniform and thin ( < 100 A) CoSi, layers can be grown by the same technique. The dislocation density in MBE-grown CoSi,(lOO) layers is very high, similar to MBE-grown silicide films on other Si substrates. For certain applications, MBE at 500 “C is a convenient and satisfactory technique to fabricate single crystal CoSi,. The homoepitaxial growth of CoSi, on existing CoSi,/Si( 100) templates occurs at room temperature from co-deposited CoSi, [log]. 8.2 Sugace Si(lO0)

and inteqface structures of CoSiJ

The CoSiJSi(100) surface also exhibits more than one quasi-stable structure [186]. Surfaces which are cobalt rich usually display a C(2~2) LEED pattern while those surfaces which are silicon rich display a (&x 3fi)-R45” pattern. These two structures are analogous to the CoSi,-C and

Fig. 20. Transmission electron diffraction pattern (120 keV) from a thin single crystal CoSi, layer on Si(100).

CoSiz-S of the (111) surface, because they can be ‘transformed’ into one another by the deposition of Co and Si, respectively [186]. The atomic structures of these two surfaces are largely unknown. A 2X 1 reconstruction is often observed at the annealed interfaces of CoSiJSi(100). The occurrence of this reconstruction is dependent on the preparation of the silicide layer, being most prominently seen in samples which have been annealed to higher temperatures. A transmission electron diffraction pattern from a thin CoSi, layer is shown in Fig. 20. The (0 f J)-related spots are from this superstructure at the interface. The atomic structure responsible for this reconstruction has been proposed to be a dimerization of excess Si at the S-fold CoSi,/Si(lOO) interface [116]. Recently, it was argued that a ‘missing row’ model is more likely the structure for the observed 2~ 1 reconstruction, because it better fits the experimentally observed HREM images [117]. With such a model, of which there are also two possible domains, the coordination number of interface Co is 7, instead of 8. Further studies are needed to resolve this issue. The electronic properties of single crystal CoSi, interfaces on Si(100) has only been briefly studied [185, 1941, where a SBH of -0.7 eV was indicated on n-type Si.

9. Epitaxial

CoSiJSi(ll0)

structures

9.1 Epitaxial growth of Co.%, on Si(ll0) CoSi, grows with the regular epitaxial orientation on Si(ll0) [187]. Again, in sharp contrast to NiSiz epitaxy on this surface, faceting is not an issue for Co&. Uniform layers of CoSi, may be grown with cobalt deposition and CJoSiz co-deposition at

Fig. 21. Planview, (220) weak beam, TEM image of a -18 8, thick Co.%, layer on Si(llO), grown by deposition of 2 A Co and co-deposition of -11 %, CoSi, at room temperature and annealing at 480 “C. Most line defects are boundaries of antiphase domains.

room temperature, followed by annealing. A TEM image of a thin CoSi, layer grown on Si(ll0) is shown in Fig. 21. There is no phase difference across a step at a CoSi,/Si(llO) interface. Therefore, there are, in principle, no ‘symmetry-required’ defects at this epitaxial system. However, most defects seen at the CoSi,/Si(llO) interfaces, e.g. Fig. 21, are phase domain boundaries [188]. The different phases arise from particular interface structures, discussed below, rather than from interfacial steps [188]. As the CoSi, film thickness increases, additional dislocations are generated at the interface to relieve misfit stress. The kinetics of strain relief along the two orthogonal directions, [liO] and [OOl], are markedly different [187]. For a fixed growth and annealing schedule, the generation of dislocations behaves as if different critical thicknesses are operative for these two directions; strain along [OOl] is relieved at a smaller film thickness than that along [liO] [187]. 9.2 Su$ace Si(ll0)

and inte$ace

structure of CoSiJ

Similar to other surfaces of CoSiz, more than one quasi-stable structure are observed at the CoSiJSi(ll0) surface. Two prominent LEED patterns were found which correspond to Co-rich and Si-rich surfaces of the CoSi,(llO) [187]. The structures of these Co& surfaces are largely unknown. The atomic structure of the CoSi,/Si(llO) interface is very intriguing. The CoSiz crystal is apparently shifted laterally with respect to the Si lattice.

129

Because regions of the interface may have two possible lateral shifts, differing by 1/4[1iO]a, [188], domain structures and different contrasts are observed under various TEM imaging conditions, as shown in Fig. 21. RBS and channeling spectra from thin CoSi, layers are shown in Fig. 22. A clear dechanneling peak, which only consists of a signal from Si atoms at the silicide interface, is observed for the CoSi,/Si(llO) sample shown in Fig. 22(b). This interface peak is consistent with a lateral relaxation between the two crystals. Such a peak is usually absent at other silicide interfaces, such as that, shown in Fig. 22(a), from a CoSi,/ Si(100) sample. The lateral shift at the CoSiJ Si(ll0) interface, indicated by planview TEM and ion channeling experiments [188], is obviously part of the structure at this interface. Presently, it seems reasonable to view such lateral shifts as being driven by the energetics of the interface atomic structure. Based purely on energetic argument, a ‘bridge site’ model of the interface structure, which allows 8fold coordination, has been proposed [ 1881. The junction characteristics of epitaxial CoSi,/Si(llO) has been studied, where a SBH of -0.7 eV on n-type Si and a SBH of - 0.4 eV on p-type Si have been measured. There have not been any theoretical investigations of the energetics and the SBH of this epitaxial interface.

10. Si/silicide/Si structures: recent development Uniform Si/silicide/Si structures, especial1 those with a silicide thickness of less than 100 K , have potential applications in high speed devices [l, 33, 61, 1891. High quality Si/NiSi,/Si(lll), Si/NiSiJ Si(lOO), and Si/CoSi,/Si(lll) structures have been grown by MBE, with the help of a ‘Si template’ technique [151, 152, 190-1931. However, effects due to band structure and scattering make it impractical to pursue ballistic transistors from these structures. Recently, an entirely different approach to silicide growth, mesotaxy, was introduced by White et al. [194] and has shown great potential, especially in the fabrication of buried silicide heterostructures. High-energy high-dose implantations of metal into silicon single crystals and high temperature annealing lead to the formation of buried silicide layers with excellent structural and electrical properties. Mesotaxy is an attractive technique for device application and for fundamental understanding of silicide energetics. Buried Si/CoSi,/Si(lOO) structures, which are of considerable technological interest, are available from this technique; however, the thickness of the buried CoSi, layer is limited to over 200 8, [195]. Another possible drawback of the mesotaxy technique is the residual damage due to implantation. Recently, through their investigation of CoSi, columnar structures, Fathauer et al. [154] were able to demonstrate the growth of essentially continuous, buried Si/ CoSi,/Si(lll) structures by an ‘endotaxial’ process [196]. Employing an MBE chamber, Mantl and Bay [197] demonstrated that thick, essentially continuous, Si/CoSi,/Si(lOO) structures may be fabricated by a co-deposition and annealing process, called ‘allotaxy’. The new concepts proposed in these recent works bring forth possibilities of novel structures and considerably broaden our views toward silicide fabrication. However, because of the energetics of the various CoSiJSi interfaces, it is still doubtful whether interesting structures such as the uniform and thin Si/CoSi,/Si(lOO) structure may be grown by one or any combination of the new techniques.

11. Conclusions 1.0

1.2

1.4 ENERGY

1.6 ( MeV)

Fig. 22. Channeling (open circles) and random (closed circles) ion scattering spectra, obtained with a glancing exit angle, of thin C&i, layers. (a) A 100 A thick layer on Si(100) and (b) a 110 8, thick layer on Si(ll0).

In this article, the basic concepts on the fabrication of epitaxial silicide structures and some results on the structural and electronic properties of epitaxial silicide layers and interfaces are described. The field of epitaxial silicides emerged in the late 1970s with the expectations of new device

130

applications and possibly some new physics. It seems that thus far epitaxial silicides have fallen short of their expectations in terms of applications but have generated an amount of exciting physics far exceeding the original expectations. As a matter of fact, some of the concepts which have long been considered fundamental in the field of SB formation have been shown to be in error, as a result of studies conducted at epitaxial silicide interfaces. Because of their unparalleled structural perfection, epitaxial silicide interfaces are the best vehicle for the study of fundamental mechanisms at MS interfaces. These studies will likely continue for some time until a satisfactory explanation of the SB emerges, and when it does, it is expected that epitaxial silicide and other epitaxial MS interfaces will have played a key role! Studies with epitaxial silicides have delivered two important messages to the solid state physics/ materials science community. The first message is that specially-designed and carefully prepared samples can make any investigation of basic physical phenomena simpler and more meaningful. The time spent in understanding and improving the quality of the samples under study is probably rewarded, in terms of improved quality of the experiment, much more than time invested anywhere else. The second message is that the time when some ‘magical’ fabrication conditions may be arrived at through empiricism is gone. The best way to improve the material quality of a sample (e.g. an epitaxial silicide interface) is through an understanding of the physics of its fabrication, i.e. thermodynamics and kinetics.

12. Acknowledgements I would like to thank my colleagues J. P. Sullivan, F. Schrey, D. J. Eaglesham, J. M. Gibson, A. F. J. Levi, A. E. White, J. M. Poate, S. M. Yalisove, F. Hellman, R. S. Becker, and J. L. Batstone for collaborations and discussions.

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