Ion-induced defects in semiconductors

Nuclear Instruments and Methods 182/183 (1981) 457-476 North-Honand Publishing Company

457

ION-INDUCED DEFECTS IN SEMICONDUCTORS James W. CORBETT, James P. KARINS Physics Department, SUNY/Albany, Albany, New York 12222, U.S.A.

and Teh Y. TAN IBM Thomas J. Watson Research Center, Yorktown lleights, N Y 10598, U.S.A.

The status of our knowledge of ion-induced defects in semiconductors will be reviewed, including the charge-state dependence of defects, novel defect migration mechanism and enhanced damage production mechanisms. The main emphasis will be on defects in silicon where a panorama of defects is emerging which encompasses the evolution of damage from vacancies and interstitials and their aggregates to stacking faults and dislocations to disordered zones and the development of an amorphous layer.

1. Introduction It has been more than 30 years since the introduction of the transistor. Initially germanium was the material o f choice for semiconductor electronics, but quickly silicon became the dominant material. For much of the time during the era of solid state electronics, diffusion technology was the main processing technique, but for a number of years ion implantation has been extensively carried out in semiconductor processing. It might be thought that in all this time that "all that one needed to know about defects in semiconductors" had been learned, and that Defect Science had followed the normal course o f evolution by creating a Defect Engineering, such that an engineer could specify what was going to be done in a process and he would find no surprises; this halcyon state has not been achieved. Certainly enough is known about defects in semiconductors to create a multibillion dollar solid state electronics industry. It is still true, however, that that technology is very much science-limited. Perhaps defect science in metals and in insulators (which topics are also considered in this proceedings) are also science-limited, but there are special features in semiconductors which make this limitation conspicuous. First the technology places increasing demands upon the materials. These demands arise out o f two opposite thrusts. The first is the continuing desire to go to smaller and smaller device dimensions; this means that a smaller and smaller amount o f defects can, if present in the critical volume of the device, destroy the effective0 0 2 9 - 5 5 4 x / 8 1 / 0 0 0 0 - 0 0 0 0 / $ 02.50 © No rth-H olland

ness of the device. Companion with the effort to achieve smaller and smaller device dimensions is an effort to extend the integration o f circuits into larger and larger systems, to the point that the integration encompasses the full width of a large wafer, such as is implied in the very large scale integrated circuit (VLSI) effort [1]. The VLSI effort requires that more of the wafer be defect free (even at the very small critical volumes of the devices) lest the entire wafer be inoperative. In addition to the increasing demands placed upon the materials by the technology, there is another feature which sets semiconductors apart from metals, at least, if not from all insulators as well, namely, the extreme sensitivity o f the semiconductor to defects. In many cases we are concerned with defect concentrations in semiconductors o f the order o f one part per billion. There are many sources o f defects, e.g., crystalgrowth, -doping, -processing, etc. We will concentrate here on the bombardment-produced defects (as is appropriate for an IBMM Conference) but defect processes form a seamless whole; defects first encountered, say, in the precipitation of impurities induced by the heat-treatment o f a crystal may also occur in a hot-implantation. What information, do we feel, is needed on defects in semiconductors? The information falls into six main categories: (1) the defect's atomic and electronic configurations are needed for the proper identification o f the defect and the basic understanding of its properties. (2) The defect's electrical levels in the forbidden l v. DEF ECTS

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J. W. Corbett et al. / lon-bMuced defects in semiconductors

gap must be known, and perhaps the associated levels in the carrier band, since those levels often relate to the following item. (3) The defect's carrier-interaction mechanisms and rates are needed to characterize the influence of the detect on the electrical properties of the material (and attendant device, if any). (4) The defect's formation mechanisms and introduction rates are needed. (5) The defect's diffusion mechanisms and rates are required to understand its role and interactions. (6) The defect's interactions, the interaction mechanisms and rates must be known to complete the picture of the influence of the defect. All this information is not available for any defect, but it is almost complete for some defects (e.g., substitutional phosphorus in silicon), and is emerging for many "others. The experimental and theoretical tools to complete this task exist; it only takes time and sufficient effort devoted to the task. In section 2 we will review the status of our knowledge in defects in semiconductors. Particular emphasis will be placed on silicon where more is known about defects than in any other system; in silicon there has recently emerged a panorama of explicit defect confi~rations ranging from the simple point defects (e.g., the vacancy and the interstitial) to theh aggregates, to the rod-defects, dislocation dipoles, stacking faults, disordered zones and to the anaorphous material itself. We will discuss special features, .e.g., ionization effects, which occur in semiconductors. These features greatly complicate the study of semiconductors, but simultaneously add considerable challenge and interest. In section 3 we will discuss the less complete panorama of the production parameters for such defects.

2. Observed defect structures 2. l. B a c k g r o u n d

We will confine our considerations to those semiconductors with fourfold coordinated atoms; we tlms omit discussion of selenium and tellurium with their lattices of chains of atoms and the I V - V I semiconductors which have NaCI lattices. The lattices which we will consider may be visualized with the help of fig. 1. Each atom has four nearest neighbors disposed as at the vertices of a tetrahedron. The cubic struc-

tures (the cubic diamond and the zinc-blende lattices) differ from the hexagonal structures (the hexagonal diamond and the wurtzite lattices) in the position of the second nearest neighbors; thus a useful way to visualize the structures is in terms of an eight-membered unit - two neighbors and their three, other nearest neighbors, for in the hexagonal structure in this unit when viewed along the axis of the two central atoms the other neighbors are superimposed and one has the "eclipsed ethane" structure, while in the cubic structure the other neighbors are in planes sixty degrees from each other and one has tile "staggered ethane" structure. Both the cubic and the hexagonal structures have as basic building blocks sixmembered rings. In the cubic lattices these rings are all "chair" rings as shown in fig. la. The hexagonal structures 'also involve the "boat" rings as shown in fig. lb. Amorphous material is thought to preserve largely the fourfold coordination of the atoms but to incorporate, in addition to both "chair" and "boat" rings, five- and seven-membered rings; figs. 2a and 2b show how five-membered rings may be joined to a "chair" ring and a "boat" ring, respectively. The simple defects in the crystal do not disturb the local lattice synlmetry. Thus in the diamond lattice the substitutional site has Tj symmetry; some substitutional impurities preserve that symmetry as may the lattice vacancy in some charge states. Similarly an interstial atom may not disturb the local lattice symmetry. In many instances forms of bonding or distortion take place for even simple point defects so that some substitutional impurities, vacancy and interstitial defects have a reduced symmetry. Experimental techniques which reveal the defect symmetry are particularly useful in establishing defect identification; such techniques are electron paramagnetic resonance (EPR), smaU angle X-ray scattering, some optical techniques and channeling. The evolution of the state of understanding has witnessed a progression of more and more incisive techniques. We will discuss most extensively the status in silicon where these techniques have been applied most, and then briefly review the status in other systems. 2.2. Silicon

There have been measurement on many impurities that have been presumed to be on substitutional sites in silicon but the EPR [2] other measurements which establish the site symmetry are limited; certainly the

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J. W. Corbett et al. / Ion-induced defects in semiconductors

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following often occupy substitutional sites: B, A1, Ga, In, and T1; C, Ge, Sn, and Pb;P, As, Sb, and Bi. Other elements may occupy substitutional sites under some circumstances. For many substitutional elements Ta symmetry is preserved; in others a lower symmetry occurs. Certainly the symmetry has not been established for all elements in the periodic table when in substitutional sites in silicon. We can be more explicit concerning references to impurity interstitials in silicon. In fig. 3 we show those elements for which evidence (EPR, channelling, optical, diffusion) has been advanced in the literature that those elements are interstitial; clearly the symmetry and properties have not been established for all elements in interstitial sites in silicon, even for some shown stipled in fig. 3. Some of these interstitials occupy the so-called tetrahedral (T) interstitial site shown in fig. 4; some the so-called hexagonal (H) site shown in fig. 5. The T and tt interstitials may be thought of as non-bonding interstitials. Some impurities are thought to occupy bonding interstitial sites such as the (linear) bond-centered site (BC), shown in fig. 6, the puckered bond-centered site (PBC) shown in fig. 7, or one of the so-called split-interstitial (or interstitialcy, or dumb-bell interstitial) configurations, shown in fig. 8, where two atoms occupy a single substitutional site, the differences in the conIV. DEFECTS

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Fig. 3. Periodic Table showing (by the stipling) those elements for which some indication exists in the literature that that element can have an interstitial configuration in silicon.

Fig. 6. Linear bond-centered interstitial site in the diamond lattice.

Fig. 7. A puckered bond-centered interstitial site in the diamond lattice.

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Fig. 5. Hexagonal interstitial site in the diamond lattice,

Fig. 8. Thei substitutional' site (upper left) in the diamond lattice and various split interstitials in which two atoms occupy the interstitial site, the various configurations differing in the orientation of the axis of the two atoms.

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J.W. Corbett et al. / Ion-induced defects in semiconductors

figurations in fig. 8 being the orientation of the axis of the pair of atoms. The bonded interstitials clearly interrupt the bonding of the silicon lattice, reforming bonds with the lattice atoms and the "dangling bonds" of the "interstitial" atom, thereby lowering the total energy. Theoretical calculations [3-5] argue that the split(100) confirmation is the stable configuration for the intrinsic interstitial in silicon; direct experimental configuration of this result is not available; although early results suggested that the G25 EPR spectrum [6], which has the right symmetry, was due to this defect, more recent work [7] suggests that it may not be. As can be seen in fig. 9, a number of EPR centers have been attributed to interstitial related defects (P6 [8,91, A5 [10], B3 [11], and 02 [12]) and their annealing properties roughly established. The P6 center is thought to be due to a di-interstitialcy, as shown in fig. 10, but other structures are not established. It has been noted [13] that interstitials can possibly aggregate along (110) axes as shown in fig. 1 I, where it will be seen that the bonding of atoms A and B of the split interstitials will result in a five-membered ring, such as found in amorphous material. The vacancy is observed (fig. 12) with lower symmetry [14] than Td (namely as D2d and C2v) due to bonding and Jahn-Teller distortions in various charge states. Similarly the divacancy [15], (fig. 13), trivacancy [16] (fig. 14), planar- [8,16,17] (fig. 15) and non-planar-tetravacancy [16] (fig. 16) and pentavacancy [16,18] (fig. 17) all exhibit bonding distortions.

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Fig. 10. The proposed di-interstitial configuration in silicon.

(We will return to these planar vacancy chains.) As shown in fig. 18, there is some information about the annealing behavior of (multi-) vacancy defects. We should note that there is not a consensus as to what is the dominant native defect in silicon in equilibrium at high temperatures. It seems clear that some impurity diffusion, e.g., phosphorus [19], can be fully accounted for by the vacancy diffusion mechanism. Van Vechten and Thurmond [20] and Bourgoin and Lannoo [21] have reiterated the view that the vacancy is the dominant equilibrium defect. Others had suggested that the divacancy [22-24] or the split-vacancy [25] dominated, but the main contender versus the vacancy is the interstitial, a point of view vigorously championed by Seeger and his co-workers [26-30]. The problem is that direct observation of the high-temperature defect has not been feasible and quenching studies (as has been surveyed recently [31]) seem dominated by impurity effects (see however, ref. [32]).

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Fig. 11. Two split-(100) interstitials on next nearest neighbor sites along a ~110~ chain showing how bonding between atoms A and 13would result in a five-membered planar ring in the diamond lattice. IV. DEFECTS

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J.W. Corbett et al. / hm-induced defects in semiconductors

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This brings us to two types of defects which, although not confirmed structures, are of special relevance to people doing channelling studies, since these defects may provide a lot of lattice disorder. The defects are the extended interstitials and the vacancy sponges. As mentioned in the previous paragraph, the interstitial has been proposed as the dominant native defect at high temperatures; in the early version [27] of this point of view that interstitial was a special interstitial - an extended interstitial, which would provide the substantial lattice disorder and defect entropy, needed to explain the data. Subsequently Jackson [33] (as reported by Kimerling [34]) proposed a model for such an interstitial with five atoms in the place of four, as shown in fig. 19. No specific calculations have been done on this model. Assumh~g that the structure shown in fig. 19 is stable,

Fig. 14. The trivacancy in the diamond lattice.

J. W. Corbett et al. / Ion-induced defects in semiconductors

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J. W. Corbett et al. lion-induced defects in semiconductors

464

we must also recognize that there may yet be other, related, extended interstitial structures. Hornstra [35] introduced the vacancy-sponge defects in discussing low-density vacancy clusters in Ge. The simplest example of such a defect is shown in fig. 20; it consists of two vacancies at the third nearest neighbor site from each other, and a relaxation of the two intervening atoms, which Hornstra proposed would form a double bond. Hornstra showed that a superlattice could be formed from such defects by putting the vacancies on ordered sites or a random network could be formed as well. JudNng from the formation of chemical molecules, the proposed double bonding may occur in diamond, but it is less likely in silicon, germanium, and tin; we note, however, that Sieverts [36] has found that a sponge structure may be the key to the understanding of an existing EPR spectrum in Si. But the double bonding is not necessary. We regard the sponges as simply examples of a broader class of defects - the partially dissociated multivacancy defects. Kleinhenz et at. [37] have already discussed how a divacancy dissociated to a second neighbor site could provide an Si--H2 structure in the crystal (fig. 21) and a similarly dissociated trivacancy an Sill3 structure (fig. 22), both of which are indicated from infrared data (as is the Si--H bond); they also said that the dissociated trivacancy could be viewed as the smallest-sized {11 I} microcrack which (in larger size) have been observed in transmission electron microscopy studies [38]. But the nature of the dissociation of multivacancy defects is not limited to these configurations and such defects may have large relaxations associated with them. Both the self-interstitial and the vacancy have been found to interact with in, purities, but the identified associated defects are limited to a few elements of the periodic table [39]. For the vacancy there are identified pairs with the following substitutional elements: B, A1, Ga; C, Ge, Sn; P, As, Sb. There are vacancy plus oxygen and multivacancy plus multioxygen

Fig. 21. A dissociated divacancy in the diamond lattice with the dangling bonds saturated with hydrogen atoms to show both Si-H and Si-H2 components.

defects which have been found. There is considerable discussion of vacancy and multivacancy plus (multi-) hydrogen defects but no specific configurations have been ftrmly established. The self-interstitial has been found to interact with B, AI, Ga, C, P and O. There are a number of instances [2] where naturally occurring impurity interstitials will form pairs with substitutional impurities, through a donor-acceptor, Coulombic interaction; impurity interstitials created by irradiation have been found to have similar interactions. There are also large structures observable with the electron microscope, wherein the experimental results may be interpreted as observations of early stages of

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Fig. 20. A "sponge" defect in the diamond lattice (see text).

Fig. 22. A dissociated trivacancy in the diamond lattice with the dangling bonds saturated with hydrogen atoms to show both Si-lt and Si-ll3 components.

J. I¢. Corbett et al. / Ion-induced defects in semiconductors

dislocation generations [40-51]. These are the so-called <110) rod-like defects and {113} stacking faults. Annealing to about 800°C of ion-irradiated silicon produces rod-like defects elongated along a <110) direction. These were found with neon [40, 41], boron [41-44], proton [45,46], and phosphorus [47] implantation. Rod-like defects and {113} stacking faults were also observed in thin silicon samples heated up to 800°C and bombarded by electrons in situ in high voltage electron microscopes [46,4850] [as well as in thin germanium [51] samples (heated to 300°C)]. At still higher temperatures, instead of rod-like defects and {113} stacking faults, dislocations were observed. In some cases it was found that dislocations result directly from rod-like defects and {113} stacking faults [50,51]. Thus, the rod-like defects and {113} stacking faults are metastable defects nucleated from condensation of point defects which may be transformed into dislocations. It is generally agreed that these defects are of interstitial type and have a radial displacement field with little or no longitudinal displacement field. There is yet no agreed view upon the atom species forming these defects. Some authors are of the opinion that the defects include impurity atoms, e.g., boron [44, 46], oxygen and carbon [49], while others postulate only host interstitial atoms [50,51]. While it is obvious that the nature of these defects may be rows of point defects [41], dislocation dipoles [44], or "precipitates" [46,49], it is noted that, except for one case [50] (which we view as highly unlikely, since it involves a row of unbonded interstitials), there are no proposed structural models. Recently Tan [52] has made detailed proposals about how point defects aggregate and create these structures, as we will now discuss. Based on the assumption that during the (extended) defect nucleation state, the dominant factor ha the configurational energy of the defect is the number of dangling bonds per point defect incorporated, rather than the more commonly recognized factor of strain energy, he was able to model the defect nucleation process. Tliis has been done for the intrinsic (vacancy) as well as extrinsic (interstitial) cases for the undissociated 90 ° edge-, the 60 °- and the Frank partial-dislocations. It was shown that, in order to minimize the number of dangling bonds, point defects would first condense into a row configuration elongated in (110). The vacancy <110) rowconfiguration is simply the inf'mite extension of the planar multivacancy configuration shown in figs. 1 3 -

465

15; the interstitial (110) row configuration is made up of the (110) chain of split interstitials shown in fig. 11. These row-configurations are called intermediate defect configurations (IDC), because they may subsequently evolve into (110) elongated dislocation dipoles. The vacancy IDCs have extended bonds and the interstitial IDCs have dangling bonds; the evolution into dislocation dipoles is driven by forming more perfect bonds. Since the IDCs extend indefinitely along a (110) direction, we can visualize their bonding perpendicular to that direction by viewing the atoms in the plane perpendicular to that (110). In fact the perfect lattice in that projection consists of <1 I0) chains of atoms from each of the two face-centered cubic sub-lattices of the diamond lattice. The bonding of the vacancy IDC (and, say, of the planar tetravacancy) in this projection is shown in fig. 23, where the 5 - 8 - 5 membered ring of the bonding is emphasized. Let us assume that the plane of the projection is a (110) plane; Tan observed that a shear stress on the (111) plane in the [112] direction can provide a driving force for bond rearrangement yielding the structure shown in fig. 24, in which there are no dangling bonds and there are adjacent pairs of 5 - 7 membered rings. Tan showed how this structure could develop into a 90 ° dislocation dipole. The interstitial IDC consists of (110) chains of sprit interstitials (fig. 11) on each sublattice, as shown in fig. 25, which Tan showed would rebond to form the configuration shown in fig. 26, namely, again adjacent pairs of five- and seven-membered rings, which he showed could be separated (in two ways) to form a 90 ° dislocation dipole. These rebonded IDCs yield nothing more than two dislocation core configurations placed right next to each other in the lattice in an appropriate fashion.

Fig. 23. The bonding [projected onto the (110) plane] surrounding a planar multivacancy [ 110] chain. (See text.) IV. DEFECTS

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J. W. Corbett et aL / lon-h~duced defects in semiconductors

Fig. 24. The rebonding of the structure in fig. 23 following a shear. (See text.)

Fig. 25. The bonding [projected onto the (110) plane] surrounding two chains of split-(110) interstitials (as shown in fig. 11). (See text.)

Fig. 26. Rebonding of the structure shown in fig. 25. (See text.)

The IDC as well as closely spaced dipoles have a radial displacement field but a very small (or 11o)longitudinal displacement field. Some (110) elongated dislocation dipoles as well as IDC have been recently observed in As* damaged silicon using the lattice imaging technique of transmission electron microscopy [53,54]. Thus, we now have a better understanding of the dislocation nucleation process from condensations of point defects as well as of the more exact nature of the rodqike defects, namely, that they indeed can be point defect rows and narrowly spaced dislocation dipoles. Tan has constructed the present models by showing the incorporation of silicon (or gemaanium) self-interstitials or vacancies, to emphasize that the nucleation process is dictated by one of the material's intrinsic properties, namely, the bonds, rather than any properties of tile hnpurity atoms which might confuse the issue if they were considered from the very beginning. However, if each self interstitial is replaced by a carbon (or boron) spilt-interstitial, the model is not changed. Also, oxygen atoms in the puckered bond-center may be segregated in (1 I0) chains. Tbus these models may provide corresponding IDCs for impurity purity precipitation. Also the incorporation of C, B, or O lowers the strain energy of the IDCs of the intrinsic defects and may thus play an important role in their formation, as some experiments seem to indicate [46,49,55]. We are also exploring the applicability of these structures to those defects observed in channelling of ionimplanted structures. As we said in section 2.1, the five- and seven-membered rings are thought to be constituents of amorphous material. Gossick [56], of course, noted that a heavily damaged region, whether amorphous or not, would introduce many defect levels into the forbidden gap and that region of the crystal would have the Fermi level in the middle of the gap in silicon. If the surrounding material had a different Fermi level there would have to be an intervening space-charge region. These Gossick zones have been observed; for example; they can be imaged in electron microscopy as has been shown by Bertolotti and co-workers [57]. But such observations simply require a damaged or disordered region, not an amorphous region. We believe that the conclusive observation of the essential equivalence of the isolated amorphous zones in silicon and amorphous silicon (a-Si) prepared by sputtering or deposition was that of Crowder et al. [58] who found the same EPR resonance properties for the iso-

J.W. Corbett et al. / Ion-induced defects in semiconductors

lated amorphous regions produced by unplantation as for that of sputtered material. The EPR results are a sensitive test; e.g., a heavily dislocated crystal which might show random diffraction or channeling results would not give the same EPR results as a-Si. In summaly we see emerging a panorama of specific defect structures for the full range of defects from simple point defects to configurations which are the building blocks of alnorphous material. 2.3. O t h e r materials

Substantially less is known for other materials than for silicon. The status has been reviewed recently: Ge [59---63] and diamond [64-69], and I I I - V [70-78] and I I - V [77-81] compounds. In Ge there are a number of elements which have substitutional sites and some with interstitial sites. Whan [82] showed that the vacancy was mobile at about 65 K and fermed a vacancy + oxygen pair; Baldwin [83] found an EPR spectrum which he associated with this center. (Neither the Whan nor the Baldwin work has been fully exploited.) It is thought [60,62] that the vacancy is the dominant, high-temperature, native defect. It is also thought that Frenkel pairs can undergo complex, ionization-dependent annihilation/or dissociation at low temperatures (<20 K) with the interstitial mobile and interacting with any of a number of impurities (Group llI, V, Li, Cu, and O), which may possibly experience a replacement. llaller [63] has carried out a number of sophisticated experiments exploiting ultra-high purity Ge to study defects and has established configurations of the L i - O pair and some hydrogen-related defects. Electron microscopy studies have found rods, dislocations and stacking faults (as in Si). The first systematic study of amorphization of a semiconductor was carried out by Parsons [84] who bombarded Ge in situ in an electron microscope with oxygen ions, the sample temperature being at 300 K or at 30 K (!). He observed individual damage regions with an amorphous structure and the onset of the complete amorphous layer as the fluence was increased. Clearly a great deal more work remains to be done concerning germanium but we are particularly encouraged by observations by Erchak and Stel'makh [85] of a number of new EPR spectra in neutron-irradiated germanium: we hope that this work and subsequent work will prove as fruitful as the comparable pioneering work in Si by Nisenoff and Fan [86]. Work in diamond is even more difficult than in the

467

other systems that we consider since diamond is not the stable phase of carbon at normal temperature and pressure. It has proven possible to make diamonds [87,88] and to dope them [89] with a few elements, e.g., B and N. Vavilov and co-workers [66] and others [90-94], have carried out a series of implantation studies using several elements, e.g., B, N, Li, Sb, and Si. Most of the work, however, has been on natura! stones. The vacancy, divacancy, tri-, and planar four. and five-vacancy chains (as discussed earlier for silicon) have been found as well nitrogen-related defects [66,67]. Dislocations are observed in the electron microscope, but the wide panorama of defects available in silicon does not yet exist. The I I l - V compounds of course are important technologically. Being compounds, stoichiometric defects are quite common and not always under control. We can obtain an excess of A in an A - B compound by having A interstitials or B vacancies; the usual stoichiometric defects are vacancies. There is information concerning some elements as substitutional atoms and fewer elements as interstitials. The interstitial configurations for Si are numerous, complex, and many of them occur experimentally; the interstitial configurations for compound systems are even more numerous and complex [13], and it seems prudent to expect that many of them can occur experimentally. Many important defects in l l I - V ' s (including dislocations) have been characterized electricaUy but few have a well identified configuration. In our view the best identified ones are the gallium vacancy in GaP [95] and in GaAs [96] for which EPR spectra have been recently established. These foundations-stones should help build a greater inventory of defects. The I I - V I compounds have the potential of wide technological use, but that potential has been frustrated by the frequent inability to achieve a suitable doping. Being compounds they too have stoichiometric defects which are often difficult to control. Further, the wide band-gap gives rise in some cases to the phenomenon of vacancy self-compensation, i.e., the annihilation of an electron--hole pair yields more energy than the vacancy formation energy so that it is impossible, in equilibrium, to move the Fermi lvel, say, to the conduction band because vacancies will be created to keep the Fermi level in the middle of the gap [97]. Again there are a number of elements which are thought to be substitutional and a smaller number to be interstitial. There are, however, a number of defects which have been identified by EPR in IV. DEFECTS

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J. W. Corbett et al. / Ion-blduced defects in semiconductors

the various compounds [98], so that we should be able to establish a great deal about point defects and their interactions. Finally, as noted by Donnelly [77], the more ionic a material, the more difficult it is to make it amorphous.

2.4. Special features of semiconductors There are several novel features which are prominent in semiconductor studies, although they need not be unique to these materials. These features complicate defect studies in semiconductors and should be borne in mind in the analysis of particle-solid interactions with semiconductors. A number of authors found indications that longdistance migration occurred at very low temperatures in germanium [99-106] and silicon [107-110] and that it may be ionization assisted [99,102--104,106]. Certainly the most direct indication was the observation by Watkins [107] of the creation of aluminum interstitials on an irradiation at about 20 K, presumably by the long-distance migration of the self-interstitial and its replacement of a substitutional aluminum atom. This replacement process has now been observed for B, A1, Ga and C; replacement is not unique to Si (for example, it apparently occurs in the Cu(Be) system [111]) and should be looked for; it may also result in mobile impurity interstitials. Normal migration processes are thermally activated and are unlikely to account for the long-distance migration of the silicon interstitial; the silicon interstitial is too massive to be expected to migrate by tunnelling. In this proceedings recoil-enhanced diffusion * is mentioned a great deal whether caused by the primary recoil or cascades. Similarly we have heard of what was called radiation-enhanced diffusion and which we feel should be called defect~enhanced diffusion (see ref. 13), usually mediated by vacancies. There can also be the normal ionization-enhanced diffusion; in this mechanism the potential energy traverse through the lattice varies with charge state (as shown in fig. 27) so that a change of charge to the favorable charge state will result in a lower migration energy (Era) and an enhanced diffusion. But all these enhanced diffusion mechanisms would require the thermally assisted nfigration of the silicon interstitial (including the recoil enhanced diffusion unless it

* We use the classification of the enhanced diffusion processes used by Corbett and Bourgoin [13].

X~ Fig. 27. Potentinl energy (E) traverses through the lattice (x) for two charge states showing Normal Ionization Enhanced Difussion.

yields an extraordinary recoil range and/or extraordinary (ballistic) cross-section for the capture of the interstitial by an impurity). There are possible athermal migration mechanisms. The first is the Bourgoin mechanism [112-118] for which the potential energy traverse for two charge stages are shown in fig. 28; the equilibrium position for one charge state is the saddle point for the other and vice versa, and a charge state change from one charge state to the other and back to the f'trst can result ha motion through the lattice without thermal activation (i.e., motion driven by the two charge state changes). The second athermal enhanced diffusion mechanism we term the energy-release mechanism (see ref. 13). It arises from the fact that the capture of a carrier by a level in the forbidden gap requires the release of the energy from the carrier band to the

Ei' ~

f

-

~

x Fig. 28. Potential energy (E) traverses through the lattice (x) for the Bourgoin m e c h a n i s m in which the saddle-point o f one charge state is the equilibrium site for t h e other charge state and vice versa.

J. I¢. £brbett et al. ~Ion-induced defects in semiconductors

[G']"[e-]

[G]

469

[G']+[,-]

"C--F/toJ Q Fig. 29. Total cncrgy (E) versus configuration coordinate (Q) for two charge states for the case where the curves intersect (a) or do not intersect (b) in the range of interest. energy level. If the energy release is radiative, the energy leaves the defect site and is lost to that defect process. If the energy is released as phonons, it can enhance defect processes. There can be two types of phonon energy release - coherent and incoherent. The coherent phonon emission may be described in terms of the configuration coordinate (Q) diagram shown in fig. 29a where is shown the total energy (E) versus Q for, say, an ionized state [G÷] plus the electron [e-], and for the un-ionized state [G]; the E(Q) curves for the two charge states intersect within the range of interest and a transition can occur to state [G] at the energy 2xE above the [G] ground state. Classically we would describe the [G] state as having oscillated within its well to the energy z2tE (plus the ground state energy) and to the corresponding value of Q; and the classical oscillator would oscillate from turning point to turning point until the energy is dissipated. Quantum mechanically we say that the [G] state, immediately after the transition at AxE, is in a prepared state of coherent phonon wave-packets so that the state mirrors the classical localization at that value of Q; the coherent wavepacket evolves in time mirroring the classical oscillation until that coherence decays by the emission of (incoherent) phonons which brings the system to the ground state of G. The coherent motion of the defect can impart a substantial impulse. In fact this is the origin of the defect creation mechanism envisioned in alkali halides [118-123] in which AE exceeds the energy required for defect creation along the direction of the impulse. (If equipartition of energy holds this energy will be 1/3 the corresponding equilibrium value.) For migration we simply require that AE exceed the corresponding migration energy for motion along the impulse.

If the two E versus Q curves do not intersect, (see fig. 29b) in the region of interest, the transition from [G ÷] to [G] can still occur (although probably with a reduced transition probability) but now the [G] is simply left in one of many wave packets after the transition which yield [G] with a value of Q within the ground state of [G÷]. There is then an incoherence in the prepared initial states of [G] and the same incoherence in the decay to the ground state of [G] which was encountered previously. The incoherent phonons are simply heat and they can enhance defect processes. The Bourgoin mechanism and the energy release mechanisms are athermal. The initial evidence in favor of the possibility of the Bourgoin mechanism was purely theoretical [4,5,117], but Troxell and Watkins [124] have argued that the boron interstitial in silicon migrates by this mechanism. The energyrelease mecahnism (presumably the coherent form) has been found to occur widely in I I I - V compounds [125-127] and for the A1 interstitial in silicon [128]. It must be presumed that these processes play a role in ion-beam interactions with solids. There have been a number of papers indicating that ionization processes can dramatically alter the displacement damage process in semiconductors [129-134], but it has been argued [135] that, since molecular ion damage studies, e.g., those of Mitchell et al. [136], only roughly 50% more damage occurs than with an atomic beam, the role of ionization cannot be too dramatic. There are also a number of papers dealing with enhanced damage (presumed to be due io "spikes") for low energy ion bombardment [137-141], but it has been noted that a surface oxide may play a role in this effect [141], that it is a surface effect [ 142], and that the RBS and channelling IV. DEFECTS

470

J.W. Corbett et al. / 1on-induced defects fl~ semiconductors

techniques used in the studies may be seeing the distortion around defects at surfaces [143]. In any event we should be alert for such effects in ion-beam interactions with semiconductors.

3. Damage production by ions 3.1. Background From electron-irradiation experiments [ 13,98] we know that displacement collisions result in the creation of Frenkel pairs, that this process is crystalorientation dependent, Fenni-level-ionization- and temperature-dependent, and that high recoil energies can result in the direct production of multivacancy defects. We know that interstitials and vacancies can have ionization-enhanced migration mechanisms and that the interstitial migration, and possibly the vacancy migration may be athermal, i.e., occur at 0 K in the presence of ionization. We know that substantial numbers of mobile interstitials and mobile vacancies can escape annihilation and interact with impurities; the interstitial interaction with an impurity may result in the replacement of a substitutional impurity by the host-interstitial creating a possibly mobile impurity interstitial. We know that from high-voltage electron microscope experiments that vacancies can coalesce into voids [144], and, as mentioned before, interstitials and vacancies can aggregate into (110) rods dislocation-dipoles and -loops and into stacking faults. There is no evidence (that we have found) in the literature that even a very high electron fluence irradiation will cause a semiconductor to go from crystalline to amorphous material (see, for example, ref. [45]).

3.2. Damage caused by ions Very early in the study of radiation damage, Brinkman [145,146] recognized that a high-energy heavy particle could create a highly damaged region in the crystal and that the dynamics of the collisions creating that region dictated that the central core of the region would be rich in vacancy-related defects while the surrounding layer would contain the interstitials which had recoiled from the damaging collisions; the composite damage structure became known as a Brinkman zone (BZ). A heavy ion penetrating a surface loses energy by both ionization and "nuclear" losses, the latter meaning displacement collisions with

the lattice; the attendant damage region may extend to the surface or be buried, but it will tend to be longer in extent than in diameter. A traverse either through the length or across the BZ will encounter a region rich in interstitials, then one rich in vacancies, then one rich in interstitials. For light ions, such as protons, the damage is much more dispersed and may consist of isolated Frenkel paks, each separated by several lattice spacings in a low density collision cascade [148]. Presumably as we progress for protons to heavier and heavier ions the damaged regions become more contiguous (and interacting) and more complex. The early workers in ion-implantation, as was ably reviewed by Nelson [147], established that a sufficient fluence (~c) would cause a semiconductor to go completely amorphous in the implanted layer. They found that 4~c depended on the ion mass, the ion energy and the sample temperature, the latter indicating that vacancy migration controlled the amorphization process. They also established that the dmnage was much less in channelling directions. As discussed in section 2.1, EPR and Raman effect studies subsequently established that, before complete amorphization, there are damaged regions in the crystal which are amorphous in all respects. Out of this, and subsequent data, several models of amorphization evolved. In terms of defect production there are those models [ 149,150] which focus on the production of simple, mobile defects and their aggregation with the homogeneous nucleation of an amorphous phase. At the other extreme is the Morehead-Crowder [ 151 ] model which assumes that a small amorphous region is produced in the stopping of a single ion and that the accumulation of these regions creates the amorphous layer. Intermediate is Gibbons [ 152,153] model which permits the damage region produced by the initial ion not to be amorphous but to be made amorphous by the overlap of one, two ..... or N damage regions; the case of N = 0 includes the case of the initial region being amorphous. Other models [154,155] focus on the amount of energy deposited in a damage event and argue that a critical energy density is necessary for amorphization. Dennis and Hale [156] considered a composite involving overlap and the energy density; they also (in the Appendix to the cited paper [156])considered a model in which a damage event converts a portion of the crystal into an amorphous region and a portion of the crystal into a damaged state which can be converted to the amorphous state by subsequent overlapping damage events.

471

J.W. Corbett et al. / 1on-induced defects in semiconductors

Table 1 Critical fluence (¢~c) versus ion mass and implantation energy. All fluences X1014. Measurements are at room temperature except for the value denoted by an asterisk which refers to 200°C. E (keV) 20 30 40 60 66 80 95 100 120 150 160

B

Ne

30 b,c)

Si

P

6.0 e)

6.0 e)

10 c)

30 c) 42 c)

12.5 c)

47 c)

17.5 c)

As

4.0 b) -3.0 f)

Kr

Sb

0.55 c), 1.7 b) 1.7 e)

Xe 1.7 c) 2.0 e) 1.0 h)

~1.0 m)

<100 k) 22.5 c) ~4.0 f) 40 g)

12 c)

85 c)

25 c) 6.0 e) 4.0 f)

1.0 a)

-20 i)

1.0 d)

0.5 i) 151. 153. 156. 159.

e) ref. 160. f) ref. 161. g) ref. 162.

h) ref. 163. i) ref. 164. J) ref. 165.

Dennis and Hale [ 1 5 6 - 1 5 8 ] , using EPR measurements of the a-Si resonance, carried out extensive tests of the various models. In tables 1, 2 and 3 we summarize their results for the critical fluence (~c) Table 2 Critical fluence (¢c) versus ion mass and temperature. All fluences X1014. B, P - 200 keV; N, Ar - 20 keV; Sb 300 keV

8O 100 200 250 3OO 350 355 381 390 410 425 428 475 5OO 508 684

Ga

*100 1) 55 k) 88 c)

800 e)

T(K)

Ar

0.6 J)

180

ref. ref. ref. ref.

O

60 a)

2OO 28O 3OO 8O0 1000 a) b) c) d)

N

B[1591 N[1561 p[1591 Ar[1561 Sb[1591 2.5 10 20 80

6.0 30

2.0 2.0 6.0 20

0.3 1.6 4.0

1.0

6O 85 105 115 5.0 12.5 35 >30 70 270

k) ref. 166. I) ref. 167. m)ref. 168.

versus ion mass, ion energy and sample temperature, along with the comparable data that we have been able to find from other workers. Dennis and ttale [156] concluded [169] that the homogeneous nucleation was not in agreement with the data and that the overlap model was required. There are EPR measurements which monitor more of the defect inventory than just that of a-Si. In fig. 30 we show a composite of data by Brower and Beezhold [162], Gerasimenko and co-workers [164, 170,171] and G6tz and co-workers [172,173]. Brower and Beezhold observed the loss of the phosphorus resonance on oxygen implantation, the growth (and loss) of the divacancy resonance and some unidentified resonances, and the ultimate growth of the a-Si resonance. Gerasimenko et al. [164] found the growth and decline of the planar four-vacancy resonance and the growth of the a-Si resonance on Ar ÷ implantation; Beezhold and Brower [174] found a comparable growth and loss of the four-vacancy resonance for B÷ implantation, where they observed only a trace of the a-Si resonance when the four-vacancy resonance is c o m p l e t e l y gone. Wesch et al. [173] correlated RBS, optical and EPR (a-Si) data for B (150 keV), N (150 keV), P IV. DEFECTS

472

J. IV. Corbett et al.

/ Ion-induced

defects in semiconductors

Table 3 Measurements at critical fluence (q~c) versus ion mass and implantation energy (fluences Xl0t4). Measurements are at 80 K E(keV) 20 40 50 60 80 100 120 150 180 200

Li[156]

B1167]

N[1561

16 20

Ne[156]

pitS61

AI1561

3.2 3.2

1.4

Kr[1571

Sb11651

0.4

1.6 28 42 45

2.0 5.5 4.7 6.0

8.0 60

4.0

1.0

1.8 0.5

order damage besides a-Si. The production rate of the a-Si resonance is higher than linear, a point Dermis and Hale [156] also noted as support for the overlap damage model; we note, however, that none of this data precludes the possibility of there being an unobserved linear production rate of a-Si at lower fluences, i.e., that regions of a-Si are directly produced in a single damage event, but by the higher fluence overlap occurs and dominates the a-Si production. (It is appropriate here to note that the EPR signal requires the defect to be in the right charge state for observation, i.e., if the Fermi level is not correct, EPR will not provide a linear monitor of defect contration; we have no reason to think that the measurements cited are troubled by this problem.) We thus come to what is in essence the model considered by Dennis and Hale [156], namely a damage region the outer portion of which is interstitial-rich and the inner portion is composed of a region which

(200 l~eV), As (110 keV) and Sb (70 keV). Sobolev et al. [172] measured the production of the EPR spectra for the four-vacancy, the di-interstitial (for two spectra without identified configurations) and a-Si for B and P implantations at 77 and 300 K. These data make some very specific requirements on the state of the damage in the samples being investigated. The observation of the divacancy and four-vacancy resonances requires that those defects be present in essentially crystalline material; a great deal of disorder nearby should distort the defect or, in any event, greatly alter the resonance. The divacancy and four-vacancy concentrations are observed to grow and decline in a way consistent with either the aggregation of defects or the evolution of defects through overlapping damage regions creating successively higher order defects, but the loss of the fourvacancy in the B+ implantations [174] clearly indicates that there is another possible state of higher

I0 I~' _

/ TOTAL

s'~

-

7----J-"-""

- sO

~ i0 j~ I I0 IO

I iOle

I

I 10~4

I

I t016

[OXYGEN] /cmz Fig. 30. Dependence of various EPR spectra on fluence of the bombarding ions. (See text.) ]'he dashed lines show spectra for which no defect identification has been made.

J. W. Corbett e t al. / Ion-induced defects in semiconductors

is crystalline but vacancy-rich and which may surround a central region which is a-Si. We note that this type of structure in which at least a portion of the vacancy-rich region is essentially crystalline is implied in the Morehead-Crowder model in which the size of the damage region is altered by vacancy diffusion, and by the data which apparently observe the temperature dependence of the critical fluence to be given [147,156] by a migration energy for the vacancy in the crystalline lattice. Whether the central region is directly-produced a-Si is not determined by the present data, although Wesch et al. [173] find that they can fit their damage profile data at 300 K with the overlap model with zero overlap (i.e., direct production) with a damage area of 400 A 2 for 70 keV Sb ions. In summary ions enter the sample and produce damage; protons produce dispersed damage while heavy ions produce more spatially dense damage. There is a prompt rearrangement of the damage regions as whatever bonding which is possible takes place. Depending on the temperature and state of ionization, defect mobility may take place which restructures the damage; the defect motion may simply result in the annihilation of a Frenkel pair, but it will be more effective if it releases a damaged region (such as a complex of dissociated multivacancies) so that the substantial defect rearrangement takes place. As successive ions produce damage near a given damage region the outer, interstitial-rich regions will encounter each other and encounter vacancy-rich regions from other damage regions; continual reconstruction of the damaged regions will take place, presumably with the defects evolving into more and more complex structures. If sufficient defect diffusion takes place, the central damage region is modified, presumably through restructuring, so that it no longer provides the nucleus for the formation of the amorphous layer. If the central region is sufficiently preserved, the amorphous layer evolves, with the interstitial sheath surrounding the length of the damage region progressive annihilated leaving interstitital regions only above and below the a-Si layer. The structures discussed by Tan (and above) provide natural stepping stones for this evolution. Gorelkinskii et al. [ 175] have confirmed this structure in silicon in that they used the EPR of an interstitial related defect (Si-B3) and profiled the depth distribution following hydrogen ion implantation and anne'allng; they found two maxima in the B3 concentration, one on each side of the end of the proton

473

range. To our knowledge comparable studies have not been done for heavier ion hnplantation. This picture of the spatial distribution of defects has also been supported by other measurements [171,176]. Clearly the amorphization models are quite successful in describing the phenomenology. Determination of further details of the process will require further experiments. Clearly EPR experiments for 4 K (or even 20 K) implantations (in a clean vacuum on samples with a variety of dopins) where vacancy migration will be inhibited, may tell whether the divacancies and four-vacancies are being directly produced and possibly whether the a-Si regions are directly produced. Other techniques may contribute as well. Clearly much progress has been made. Clearly there is still lots to learn.

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J.W. Corbett et al. / Ion-induced defects in semiconductors

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