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Atomic collisions, elastic recombination, and thermal diffusion during thin-film growth from low-energy ion beams
326
Nuclear
Instruments
and Methods
in Physics
Research
B59/60
(1991) 326-331 North-Holland
Atomic collisions, elastic recombination, and thermal diffusion during thin-film growth from low-energy ion beams N. Herbots *, 0. Vancauwenberghe
* *, O.C. Hellman + and Y.C. Joo ++
Massachusetts Institute of Technology, Cambridge, MA 02139, USA
Low-energy ( ~1 keV) ions are used in a variety of thin-film techniques. When low-energy ions are used during growth, the atomic mobility is athermally enhanced. This can lead to a significant lowering of the temperature necessary to induce epitaxial
growth and chemical reactions. Athermal enhancement of atomic mobility in semiconductors can be described below the temperature for plastic deformation (T = 540’ C in Si) by classifying the mechanisms involved into three categories according to their respective timescale: collisions, elastic recombination, and thermal diffusion. A quantitative model can then be derived to predict the conditions of temperature, dose rate, and energy to obtain thin film growth, epitaxial growth, and oxidation in techniques such as ion beam deposition (IBD), and ion beam oxidation (IBO). Using computer simulations, the dynamics of defect generation and redistribution, and the resulting thin-film growth rate can be investigated. Energies below 200 eV are found not only to minimize damage and sputtering, but also create defect distributions that favor surface recombination and hence growth. This elucidates the mechanism of thin-film formation with high atomic density, oxidation with a sharp interface with the substrate and epitaxial growth, and experimental findings on the energy dependence of IBD and IBO.
1. Introduction
ical reactions, and effectively lower processing temperatures [2-71. In this article, the role of elastic recombination and point defect diffusivity below the plastic deformation temperature, is described and integrated into an atomistic model of athermal enhancement of diffusivity, growth, and reaction produced by low-energy ion bombardment. Computer simulations of this model are compared to experimental data on epitaxial growth by ion beam deposition (IBD) [2-61 and oxides formed by ion beam oxidation (IBO) [6-81. The results are also discussed for the case where molecular species are present such as in CIMD [9] or ions that are not unidirectional, such as in plasma deposition. It is shown that the insights gained about IBD can provide a better understanding of the relationship between ion parameters and thin-film properties.
The need to control the crystalline quality and the compositional profile in semiconductor heterostructures has led to the development of techniques such as molecular beam epitaxy (MBE) and chemical vapor deposition (CVD). As device requirements on composition and doping distributions is drawing closer to the atomic scale, the lowering of the thermal budget necessary to create layered structures and small geometry devices has become the focus of new developments such as chemical beam epitaxy (CBE), limited reaction processing (LRP) [l] and plasma-enhanced CVD. This trend has spilled to metallic deposition and high-T, superconductor formation where the need for atomic scale design and control of heterostructures has emerged as well. Low-energy (15 eV-1 keV) ion deposition techniques have extended the temperature range for deposition to the lowest limits. The kinetic energy of ions can eliminate the thermal activation of kinetic steps, modify the kinetics of rate-limited diffusion, growth and chem-
2. An atomistic model of thin-film formation from lowenergy ions
* C.R. Soderberg Assistant Professor of Electronic Materials. * * Supported by the National Science Foundation under contract no. 87-19217-DMR. + Supported by the Carl Soderberg Fund. ++ Supported by the Korean Government Overseas Fellowship.
The physical description of atomic collisions in a wide range of energy has reached a high level of computational sophistication. Atomic collisions that do not involve nuclear resonances or reactions can be described accurately by classical mechanics, using the binary collision approximation (BCA) in Monte Carlo techniques, in codes such as MARLOWE [lo] and TRIM [ll]. The BCA remains valid at energies above 50 eV.
0168-583X/91/$03.50
0 1991 - Elsevier
Science Publishers
B.V. (North-Holland)
N. Herbots et al. / Mechanisms during thin-film growth
The particles involved in collision cascades can be tracked by most computers. Important advances have been made in the description of a universal atomic potential [11,12]. At the low end of the energy spectrum, below 10 eV, the BCA breaks down and molecular dynamics (MD) techniques have to be used, so that multiple collisions can be accounted for [13,14]. Atomic potentials representing adequately directionally bonded solids and computing techniques to handle the large number of small atomic displacements to be taken into account are still being developed, however. The energy range 15-500 eV where most IBD and IBO experiments are carried out [2-61 is thus found to be close to both the BCA range of applicability and that of MD. In refs. [6,15], we have demonstrated that atomic collisions alone were not sufficient to explain damage generation at 300 K during low-energy ion bombardment at 200 eV and below. Thermal effects during ion bombardment above 300 K result in microstructures that cannot be reproduced by a post-deposition annealing, demonstrating an important relationship between irradiation and temperature: deep microstructural damage (70 to 250 nm below the initial surface) is found in IBD samples grown at 650 K and above while a sample subjected to the same IBD conditions at 300 K exhibits damage only near the initial surface. Such
0
Incoming Ion
321
observations were made at 200 eV [15], 65 eV [16] and 40 eV [S]. This has led to the conception of a simple model [6] where collisions, elastic response, and thermal diffusivity are accounted for, by taking advantage of their temporal separation. Computer simulations of thin-film growth based on this model are presented here for the first time and analyzed in light of experimental findings. Fig. 1 depicts the proposed model applied to the growth of a thin film from 65 eV 74Ge+ impinging on a Si substrate where the surface region is neither free of atmospheric or chemical contamination nor ordered. This case can be paralleled to the case shown in ref. [6] for epitaxial growth. As shown in ref. [15], ions at energies above 15 eV have a finite depth of penetration beyond the first surface layer. This leads to layer anchoring [5] as opposed to growth processes initiated at the surface such as MBE. The resulting thin film growth is represented in fig. 1, where the location of the three steps of deposition are shown. Deposition starts with shallow implantation (fig. lb) at the depth of anchoring accompanied by a two-step surface conversion mechanism (figs. lc and ld). A consequence of the surface conversion mechanism is that even if the range of the ions is beyond a few monolayers, disorder beyond the first atomic layer will lead to amorphous or
IL
depth (layer)
I
0
l Thin Film Atom 0
substrate Atom (a)
(b) det%h (layer)
New Atoms on ace
4 Surface motion
8 Thin Film Interstitial @ Elastically recombined 0 Vacancy
Interstitial
Fig. 1. Two-dimensional schematic illustrating the growth of a thin film of random structure according to our model (computer simulations are three-dimensional). See text. III. ION-ENHANCED
DEPOSITION
328
N. Herbots et al. / Mechanisms during thin-film growth
polycrystalline growth [16] without registry with the underlying crystal. On the other hand, if contamination and defects are truly confined to the first monolayer, epitaxy can still happen since ordering is driven by the atomic layers directly below the surface. This can be seen, for instance, in ref. [5] where good epitaxial quality is obtained on a substrate whose first atomic layer is obviously contaminated. The collisions are tracked using the BCA model, as shown in the histograms of figs. lb and lc, and take place in a timescale of less than lo-l3 s. In an amorphous or partially disordered solid, damage can be described with vacancies and interstitials. An interstitial is an atom in a nonbonding site in the random or disordered fourfold coordinated network. A vacancy is a vacant bonding site in the network. Experimental characterization of such defects produced by ion bombardment in amorphous silicon have recently been reported [ 171. Once defects have been created local fluctuations of atomic density occur, and hence small regions with high elastic strain are formed (fig. lb). The density of defects is too small to destroy the organization of the original network. Most nearest and second nearest neighbors are still found on their original site. At low energies, interstitials are generated somewhat deeper in than vacancies which are confined to the first few atomic layers of the surface, as seen in the histograms of figs. lb and lc. The far-from-equilibrium concentrations, with interstitials and vacancies in equal number, are thus spatially separated over a short distance while remaining close to the surface. This increases first the probability of elastic recombination between defects, and second recombination with the surface, either with the originally ejected atom or with the atom having caused the collision. The latter event is called a relocation. Elastic forces tend thus to bring back interstitials into bonding sites, which is of great importance in the epitaxial case [6] but also leads to different properties in amorphous IBD films [17]. Elastic recombination (ER) occurs in a timescale of about 3-6 x lo-l3 s [18]. The time separation of ER with collisions is not critical: ER is computed in at the time of the collisions. It is the time separation of ER with thermal diffusion which is of importance. The timescale for diffusion is several orders of magnitude larger, enabling us to treat defect generation (collisions and ER) as instantaneous when thermal effects are computed. In semiconductors at low temperature, interstitials are readily mobile, even at 300 K. If they are close to the surface, they diffuse in majority towards this sink, leading to net surface growth (fig. Id). Vacancies on the other hand can be considered immobile at low temperatures and can only be annihilated by the diffusing interstitials. This leads to thin films that can be amorphous, but of high atomic density since vacancies are refilled first [15]. If the starting surface
region is ordered, this three-step mechanism leads to recombination mostly on lattice sites, leading to epitaxial films with few remaining defects [9,16].
3. Results and discussion We use two codes to calculate the dynamics of point defects generation and distribution, and the resulting growth rate dependency on energy and dose rate. The first code, TRIMCSR, is based on TRIM (transport of ion matter [ll]), and track recoils down to the energy of 0.5 eV. TRIMCSR takes into account simultaneously cascade and sputtering, including relocations [7]. The program computes the statistical distribution of defects generated by the atomic collisions and relocations from the histograms such as in fig. 1. It should be noted that the spatial separation of defects found at low energies from TRIMCSR calculations is at the basis of the difference between ion beam induced and thermally activated atomic mobility. The number of interstitials present is equal to the number of vacancies in the case of ion bombardment which is very far from the case of thermal equilibrium where vacancies, being more stable, dominate. This leads to massive recombination at the surface later on when thermal effects set in, provided interstitials are close enough to the surface to be captured. The presence of vacancies between the surface and the peak of the interstitials distribution ensures that if these interstitials are diffusing towards the surface they will recombine with vacancies first. The remaining interstitials, which are equal in number to the number of ions introduced, end up on the surface by thermal diffusion. In other words, it is low-temperature diffusion, in the end, which is responsible for net thin-film growth. Thermal diffusion is calculated by DRIVIC [7], where the static distribution obtained by TRIMCSR and the dose rate used in deposition are the input data for defect generation. This ensures that during the simulation defects are introduced at the correct rate, determined by the dose rate, and at the correct location, determined by the static distribution and the surface motion. The continuity and diffusion equations implemented are discussed in ref. [7]. The diffusivity for vacancies and interstitials are found in ref. [19]. An example of the surface motion (and growth rate) obtained is shown in fig. 2b, where the location of the surface at intervals of 1, 10 and 20 s during deposition with a dose rate of lOi ions cm-2s-1 of 3oSi on Si, which is typical for data in refs. [5,6,16]. The TRIMCSR distributions used in this calculation are shown in fig. 2a. The surface motion is obtained by calculating the flux of interstitials reaching the surface. Fig. Id depicts the atomistic mechanism for the thermally driven surface motion. Growth occurs by thermal
7%Herbots et al. / Mechanism during thin-film growth
65 eV “Si* - --
->
Si
Vacancies - lnterstitials -.-.-.- Total Ions
.
329
as deep (fig. 3~). A cross-over is observed approximately at 200 eV where the ion distribution peaks beyond the interstitials distribution, while the separation between vacancies and interstitials decreases and their distance with respect to the surface increases. By comparing the range ratios of interstitials and vacancies as a function of energy (fig. 3b), one can see that passed the cross-over, these ratios converge rapidly to one. The same is true for the ion range versus the vacancy range, which decreases the likelihood of ions to recombine with the surface and hence growth. This is also seen in fig. 3c where the depth at which the peak concentration of
40Ar -> Si IOttS VacancY Total Interstitial substitutional Intedtial
Ion
Ion
Replacement RdoeariOtl
’ O”
2
0
100 200 300 400 Energy(eV)
(a)
40Ar -B Si
Rp ImvRp Vat Rp InvRp Vat (Rp lo-Rp VayRp 1
Fig. 2. Computer simulations of thin-film growth and surface motion rates as a function of deposition time for a dose rate of lOI ?Si+ ions cm-*s-l on a Si substrate. The instantaneous depth distributions of vacancies and interstitials are shown after 1, 10 and 20 s of deposition. The growth rate is the net displacement of the surface divided by the deposition time.
(Rp In-Rp VayRp \
-0 200 400600 8001000 Energy(eV) 40Ar -> Si 3Om
migration and surface recombination of substrate interstitials and interstitial ions, left behind after atomic collisions and ER.
These events are highly probable but not necessary occurrences. Interstitials and ions can also recombine before reaching the surface, an event whose probability increases as either range (and thus energy) or damage density (related to both energy and dose rate) are increased. This implies that the growth rate will decrease due to less efficient surface recombination as energy is increased, in addition to sputtering effects. Figure 3a shows the ener y dependence of the spatial distribution B of defects for Ar+ on Si as calculated by TRIMCSR, using the classical range notion derived from a Gaussian fit (figs. 3a and 3b) and the true location of the peak concentration (fig. 3~). At low energies, vacancies are found in the first or second monolayer of the substrate, while interstitials are found two to three times
(c) “0 2004006008001000 Energv(eV) Fig. 3. Energy dependence of the spatial distribution of vacancies and interstitials, ions, interstitials and substitutional ions, replacement events and relocation events for the case of %r on Si. (a) depicts the energy dependence of their respective range. In (b) the ratio of ion range versus vacancy range (solid circles), and the range ratio of interstitials versus vacancies (open circles) are shown as function energy. The range separation ratio (filled and empty squares), normalized to the vacancy range are found to quickly converge to 0 past 200 eV. In (c), the location of the true maximum of the ion, interstitials and vacancies distributions rather than the range inferred from a Gaussian fit as in (a) and (b), is shown.
III. ION-ENHANCED DEPOSITION
330
N. Herbots et al. / Mechanisms
18DGrowth:
5
. 4 L
Si on I.. ev
----
200
Si
r.,....,...,
5 Si
ion
flux:
ev
: -.-._._
during thin-film growth
Recombination below the surface should be favored to obtain oxide growth (100 eV ,< E < 1 keV) as found in ref. [8].
4. Conclusions
--__ --__
-1
1 -1
~....I....,....,....~
-*0
5
Time ‘&ee.)
15
20-2
Fig. 4. Thin-film growth for a 30Sif ion beam with a dose rate of lOI4 ions cm-*SC’ as a function of time, with ion energy as parameter. The slope of the curves provide the growth rates discussedin the text.
vacancies is found to be confined within the first few monolayers below 200 eV, and increases parabolically with energy while the depth of the interstitials increases logarithmically. Practically, these calculations show that below 200 eV, little damage is left behind from the Ar bombardment but the rate of Ar retention increases dramatically. This would probably lead to clustering, which is not included in our calculation. The dependence of defect distributions upon ion energy and the effect of thermal diffusion determines the dependence of thin-film growth rate upon energy, as shown in fig. 4. In fig. 4, the surface motion as a function of time, which translates into a linear growth rate, is shown for the case of deposition 3oSi on Si, for different energies. The growth rates stay similar for energies below 200 eV. Above 200 eV, the growth rate drops about twice as fast as sputtering losses [7,15]. The number of interstitials contributing to surface growth is then below the number of ions impinging the target. A larger fraction of interstitials remains trapped within the substrate. This last observation is important for optimizing the energy for growth rate and properties such as epitaxy. Specifically, the selection of energies maintaining the sputtering yield below 1 [6,16] is not sufficient to obtain a thin-film growth rate larger than 0 and is overridden by the necessity of favoring specific defect distribution and recombination mechanism if epitaxial or high-density films are to be grown. These observations are consistent with reports from independent groups that beyond 200 eV ion bombardment during epitaxial growth leads to residual damage and poor electronic properties [20]. In the case of ion beam oxidation (IBO), the reverse is true: surface recombination (and thus possible desorption) is to be avoided.
In summary, the mechanism of athermal enhancement mobility by low-energy ion beam can be described at temperatures below plastic deformation (“low temperature”) in semiconductor materials in terms of atomic collisions, ER and point defect thermal diffusion. The model based on this description gives us as intuitive understanding of the mechanism by which a thin film can be grown from an ion beam. It also allows us to compute quantitatively in what conditions of temperature, dose rate, and energy such processes lead to net thin-film growth, epitaxy and compound formation for conceptually simple techniques such as ion beam deposition (IBD) and ion beam oxidation (IBO).
References VI C. King, J.L. Hoyt, C.M. Cronet, J.F. Gibbons, M.P. Scott and J. Tarur, IEEE Trans. Electron. Dev. Lett. EDL-10 (1989) 52. PI P. Tsukizoe, T. Nakai and N. Ohmae, J. Appl. Phys. 42 (1976) 4770. [31 K. Yagi, S. Tamura and T. Tokuyama, Jpn. J. Appl. Phys. 16 (1977) 245.
[41 P.C. Zalm and L.J. Beckers, Appl. Phys. Lett. 41 (1982) 167. 151 N. Herbots, B.R. Appleton, T.S. Noggle, S.J. Pennycook, R.A. Zuhr and D.M. Zhener, Semiconductor-based Heterostructures: Interfacial Structure and Stability, eds. M.L. Green et al. (1986) pp. 335-349. 161 N. Herbots, O.C. Hellman, P.A. Cullen and 0. Vancauwenberghe, Am. Vat. Sot. Series 4, ed. G. Luckovsky, Deposition and growth: limits for microelectronics, AIP Conf. Proc. 167, ed. G.W. Rubloff (New York, 1988) pp. 259-289. 171 0. Vancauwenberghe, N. Herbots and O.C. Helhnan, Sot. Photo-optical Instr. Eng. Proc. 1285 (SPIE, Bellingham, 1990) pp. 47-59. [8] 0. Vancauwenberghe, N. Herbots, H. Manoharan and M. Ahrens, J. Vat. Sci. Technol. (1991) in press; N. Herbots, 0. Vancauwenberghe and H. Manoharan, 36th Am. Vat. Sot. Meeting, AVS Discovery Session, Boston, MA, October 23-27, 1989; N. Herbots, 0. Vancauwenberghe and H. Manoharan, Americ. Phys. Sot. March Meeting, Anaheim, CA, March 12-16, 1990. [91 N. Herbots and O.C. Hellman, US Patent no. 4800100 (1989). WI M.T. Robinson and I.M. Torrens, Phys. Rev. B9 (1962) 5008. [ill J.P. Biersack and L.J. Haggmark, Nucl. Instr. and Meth. 174 (1980) 257;
N. Herbots et al. / Mechanisms during thin-film growth see also: J.P. Biersack and W. Eckstein, J. Appl. Phys. A34 (1984) 13. 1121A. Tenner, Rainbow Scattering, Ph.D. thesis, FOM-Instituut, Amsterdam (1986). (131 B. Dodson, Phys. Rev. B36 (1987) 1068. [14] B.J. Garrison, M.T. Miller and D.W. Brenner, Atomic Scale calculations in Materials Science, eds. J. Tersoff, D. Vanderbilt and V. Vitek, Mater. Res. Sot. Symp. Proc. 141 (1990) 419. [15] N. Herbots, B.R. Appleton, TX NoggIe, R.A. Zuhr and S.J. Pennycook, Nucl. Instr. and Meth. B13 (1986) 250. [16] N. Herbots, B.R. Appleton, S.J. Pennycook, T.S. NoggIe and R.A. Zuhr, in: Beam-Solid Interactions and Phase Transformations, eds. H. Kurtz, G.L. Olson and J.M. Poate, Mater. Res. Sot. Symp. Proc. 51 (1986) 3669.
331
[17] S. Roorda, J.M. Poate, DC. Jacobson, D.J. Eaglesham, B.S. Dennis, S. Dieker, W.C. Sinke and F. Spaepen, Solid State Commun. 75 (1990) 197. [18] U. Landman, W.D. Luedtke and M.W. Ribarsky, J. Vat. Sci. Technol. A7 (1989) 2829. [19] U. Gosele, in: Semiconductor Silicon 1986, eds. H.R Huff, T. Abe and B. Kolbesen (Pemtington, Electrochem. Sot., 1986) p. 541. [20] M.A. Hasen, J. Knoll, S.A. Bamett, J.E. Sundgren, L.C. Marker& A. Rockett and J.E. Greene, J. Appl. Phys. 65 (1989) 172; also: S.S. Iyer and E. Kasper, private communication.
III. ION-ENHANCED
DEPOSITION
Nuclear
Instruments
and Methods
in Physics
Research
B59/60
(1991) 326-331 North-Holland
Atomic collisions, elastic recombination, and thermal diffusion during thin-film growth from low-energy ion beams N. Herbots *, 0. Vancauwenberghe
* *, O.C. Hellman + and Y.C. Joo ++
Massachusetts Institute of Technology, Cambridge, MA 02139, USA
Low-energy ( ~1 keV) ions are used in a variety of thin-film techniques. When low-energy ions are used during growth, the atomic mobility is athermally enhanced. This can lead to a significant lowering of the temperature necessary to induce epitaxial
growth and chemical reactions. Athermal enhancement of atomic mobility in semiconductors can be described below the temperature for plastic deformation (T = 540’ C in Si) by classifying the mechanisms involved into three categories according to their respective timescale: collisions, elastic recombination, and thermal diffusion. A quantitative model can then be derived to predict the conditions of temperature, dose rate, and energy to obtain thin film growth, epitaxial growth, and oxidation in techniques such as ion beam deposition (IBD), and ion beam oxidation (IBO). Using computer simulations, the dynamics of defect generation and redistribution, and the resulting thin-film growth rate can be investigated. Energies below 200 eV are found not only to minimize damage and sputtering, but also create defect distributions that favor surface recombination and hence growth. This elucidates the mechanism of thin-film formation with high atomic density, oxidation with a sharp interface with the substrate and epitaxial growth, and experimental findings on the energy dependence of IBD and IBO.
1. Introduction
ical reactions, and effectively lower processing temperatures [2-71. In this article, the role of elastic recombination and point defect diffusivity below the plastic deformation temperature, is described and integrated into an atomistic model of athermal enhancement of diffusivity, growth, and reaction produced by low-energy ion bombardment. Computer simulations of this model are compared to experimental data on epitaxial growth by ion beam deposition (IBD) [2-61 and oxides formed by ion beam oxidation (IBO) [6-81. The results are also discussed for the case where molecular species are present such as in CIMD [9] or ions that are not unidirectional, such as in plasma deposition. It is shown that the insights gained about IBD can provide a better understanding of the relationship between ion parameters and thin-film properties.
The need to control the crystalline quality and the compositional profile in semiconductor heterostructures has led to the development of techniques such as molecular beam epitaxy (MBE) and chemical vapor deposition (CVD). As device requirements on composition and doping distributions is drawing closer to the atomic scale, the lowering of the thermal budget necessary to create layered structures and small geometry devices has become the focus of new developments such as chemical beam epitaxy (CBE), limited reaction processing (LRP) [l] and plasma-enhanced CVD. This trend has spilled to metallic deposition and high-T, superconductor formation where the need for atomic scale design and control of heterostructures has emerged as well. Low-energy (15 eV-1 keV) ion deposition techniques have extended the temperature range for deposition to the lowest limits. The kinetic energy of ions can eliminate the thermal activation of kinetic steps, modify the kinetics of rate-limited diffusion, growth and chem-
2. An atomistic model of thin-film formation from lowenergy ions
* C.R. Soderberg Assistant Professor of Electronic Materials. * * Supported by the National Science Foundation under contract no. 87-19217-DMR. + Supported by the Carl Soderberg Fund. ++ Supported by the Korean Government Overseas Fellowship.
The physical description of atomic collisions in a wide range of energy has reached a high level of computational sophistication. Atomic collisions that do not involve nuclear resonances or reactions can be described accurately by classical mechanics, using the binary collision approximation (BCA) in Monte Carlo techniques, in codes such as MARLOWE [lo] and TRIM [ll]. The BCA remains valid at energies above 50 eV.
0168-583X/91/$03.50
0 1991 - Elsevier
Science Publishers
B.V. (North-Holland)
N. Herbots et al. / Mechanisms during thin-film growth
The particles involved in collision cascades can be tracked by most computers. Important advances have been made in the description of a universal atomic potential [11,12]. At the low end of the energy spectrum, below 10 eV, the BCA breaks down and molecular dynamics (MD) techniques have to be used, so that multiple collisions can be accounted for [13,14]. Atomic potentials representing adequately directionally bonded solids and computing techniques to handle the large number of small atomic displacements to be taken into account are still being developed, however. The energy range 15-500 eV where most IBD and IBO experiments are carried out [2-61 is thus found to be close to both the BCA range of applicability and that of MD. In refs. [6,15], we have demonstrated that atomic collisions alone were not sufficient to explain damage generation at 300 K during low-energy ion bombardment at 200 eV and below. Thermal effects during ion bombardment above 300 K result in microstructures that cannot be reproduced by a post-deposition annealing, demonstrating an important relationship between irradiation and temperature: deep microstructural damage (70 to 250 nm below the initial surface) is found in IBD samples grown at 650 K and above while a sample subjected to the same IBD conditions at 300 K exhibits damage only near the initial surface. Such
0
Incoming Ion
321
observations were made at 200 eV [15], 65 eV [16] and 40 eV [S]. This has led to the conception of a simple model [6] where collisions, elastic response, and thermal diffusivity are accounted for, by taking advantage of their temporal separation. Computer simulations of thin-film growth based on this model are presented here for the first time and analyzed in light of experimental findings. Fig. 1 depicts the proposed model applied to the growth of a thin film from 65 eV 74Ge+ impinging on a Si substrate where the surface region is neither free of atmospheric or chemical contamination nor ordered. This case can be paralleled to the case shown in ref. [6] for epitaxial growth. As shown in ref. [15], ions at energies above 15 eV have a finite depth of penetration beyond the first surface layer. This leads to layer anchoring [5] as opposed to growth processes initiated at the surface such as MBE. The resulting thin film growth is represented in fig. 1, where the location of the three steps of deposition are shown. Deposition starts with shallow implantation (fig. lb) at the depth of anchoring accompanied by a two-step surface conversion mechanism (figs. lc and ld). A consequence of the surface conversion mechanism is that even if the range of the ions is beyond a few monolayers, disorder beyond the first atomic layer will lead to amorphous or
IL
depth (layer)
I
0
l Thin Film Atom 0
substrate Atom (a)
(b) det%h (layer)
New Atoms on ace
4 Surface motion
8 Thin Film Interstitial @ Elastically recombined 0 Vacancy
Interstitial
Fig. 1. Two-dimensional schematic illustrating the growth of a thin film of random structure according to our model (computer simulations are three-dimensional). See text. III. ION-ENHANCED
DEPOSITION
328
N. Herbots et al. / Mechanisms during thin-film growth
polycrystalline growth [16] without registry with the underlying crystal. On the other hand, if contamination and defects are truly confined to the first monolayer, epitaxy can still happen since ordering is driven by the atomic layers directly below the surface. This can be seen, for instance, in ref. [5] where good epitaxial quality is obtained on a substrate whose first atomic layer is obviously contaminated. The collisions are tracked using the BCA model, as shown in the histograms of figs. lb and lc, and take place in a timescale of less than lo-l3 s. In an amorphous or partially disordered solid, damage can be described with vacancies and interstitials. An interstitial is an atom in a nonbonding site in the random or disordered fourfold coordinated network. A vacancy is a vacant bonding site in the network. Experimental characterization of such defects produced by ion bombardment in amorphous silicon have recently been reported [ 171. Once defects have been created local fluctuations of atomic density occur, and hence small regions with high elastic strain are formed (fig. lb). The density of defects is too small to destroy the organization of the original network. Most nearest and second nearest neighbors are still found on their original site. At low energies, interstitials are generated somewhat deeper in than vacancies which are confined to the first few atomic layers of the surface, as seen in the histograms of figs. lb and lc. The far-from-equilibrium concentrations, with interstitials and vacancies in equal number, are thus spatially separated over a short distance while remaining close to the surface. This increases first the probability of elastic recombination between defects, and second recombination with the surface, either with the originally ejected atom or with the atom having caused the collision. The latter event is called a relocation. Elastic forces tend thus to bring back interstitials into bonding sites, which is of great importance in the epitaxial case [6] but also leads to different properties in amorphous IBD films [17]. Elastic recombination (ER) occurs in a timescale of about 3-6 x lo-l3 s [18]. The time separation of ER with collisions is not critical: ER is computed in at the time of the collisions. It is the time separation of ER with thermal diffusion which is of importance. The timescale for diffusion is several orders of magnitude larger, enabling us to treat defect generation (collisions and ER) as instantaneous when thermal effects are computed. In semiconductors at low temperature, interstitials are readily mobile, even at 300 K. If they are close to the surface, they diffuse in majority towards this sink, leading to net surface growth (fig. Id). Vacancies on the other hand can be considered immobile at low temperatures and can only be annihilated by the diffusing interstitials. This leads to thin films that can be amorphous, but of high atomic density since vacancies are refilled first [15]. If the starting surface
region is ordered, this three-step mechanism leads to recombination mostly on lattice sites, leading to epitaxial films with few remaining defects [9,16].
3. Results and discussion We use two codes to calculate the dynamics of point defects generation and distribution, and the resulting growth rate dependency on energy and dose rate. The first code, TRIMCSR, is based on TRIM (transport of ion matter [ll]), and track recoils down to the energy of 0.5 eV. TRIMCSR takes into account simultaneously cascade and sputtering, including relocations [7]. The program computes the statistical distribution of defects generated by the atomic collisions and relocations from the histograms such as in fig. 1. It should be noted that the spatial separation of defects found at low energies from TRIMCSR calculations is at the basis of the difference between ion beam induced and thermally activated atomic mobility. The number of interstitials present is equal to the number of vacancies in the case of ion bombardment which is very far from the case of thermal equilibrium where vacancies, being more stable, dominate. This leads to massive recombination at the surface later on when thermal effects set in, provided interstitials are close enough to the surface to be captured. The presence of vacancies between the surface and the peak of the interstitials distribution ensures that if these interstitials are diffusing towards the surface they will recombine with vacancies first. The remaining interstitials, which are equal in number to the number of ions introduced, end up on the surface by thermal diffusion. In other words, it is low-temperature diffusion, in the end, which is responsible for net thin-film growth. Thermal diffusion is calculated by DRIVIC [7], where the static distribution obtained by TRIMCSR and the dose rate used in deposition are the input data for defect generation. This ensures that during the simulation defects are introduced at the correct rate, determined by the dose rate, and at the correct location, determined by the static distribution and the surface motion. The continuity and diffusion equations implemented are discussed in ref. [7]. The diffusivity for vacancies and interstitials are found in ref. [19]. An example of the surface motion (and growth rate) obtained is shown in fig. 2b, where the location of the surface at intervals of 1, 10 and 20 s during deposition with a dose rate of lOi ions cm-2s-1 of 3oSi on Si, which is typical for data in refs. [5,6,16]. The TRIMCSR distributions used in this calculation are shown in fig. 2a. The surface motion is obtained by calculating the flux of interstitials reaching the surface. Fig. Id depicts the atomistic mechanism for the thermally driven surface motion. Growth occurs by thermal
7%Herbots et al. / Mechanism during thin-film growth
65 eV “Si* - --
->
Si
Vacancies - lnterstitials -.-.-.- Total Ions
.
329
as deep (fig. 3~). A cross-over is observed approximately at 200 eV where the ion distribution peaks beyond the interstitials distribution, while the separation between vacancies and interstitials decreases and their distance with respect to the surface increases. By comparing the range ratios of interstitials and vacancies as a function of energy (fig. 3b), one can see that passed the cross-over, these ratios converge rapidly to one. The same is true for the ion range versus the vacancy range, which decreases the likelihood of ions to recombine with the surface and hence growth. This is also seen in fig. 3c where the depth at which the peak concentration of
40Ar -> Si IOttS VacancY Total Interstitial substitutional Intedtial
Ion
Ion
Replacement RdoeariOtl
’ O”
2
0
100 200 300 400 Energy(eV)
(a)
40Ar -B Si
Rp ImvRp Vat Rp InvRp Vat (Rp lo-Rp VayRp 1
Fig. 2. Computer simulations of thin-film growth and surface motion rates as a function of deposition time for a dose rate of lOI ?Si+ ions cm-*s-l on a Si substrate. The instantaneous depth distributions of vacancies and interstitials are shown after 1, 10 and 20 s of deposition. The growth rate is the net displacement of the surface divided by the deposition time.
(Rp In-Rp VayRp \
-0 200 400600 8001000 Energy(eV) 40Ar -> Si 3Om
migration and surface recombination of substrate interstitials and interstitial ions, left behind after atomic collisions and ER.
These events are highly probable but not necessary occurrences. Interstitials and ions can also recombine before reaching the surface, an event whose probability increases as either range (and thus energy) or damage density (related to both energy and dose rate) are increased. This implies that the growth rate will decrease due to less efficient surface recombination as energy is increased, in addition to sputtering effects. Figure 3a shows the ener y dependence of the spatial distribution B of defects for Ar+ on Si as calculated by TRIMCSR, using the classical range notion derived from a Gaussian fit (figs. 3a and 3b) and the true location of the peak concentration (fig. 3~). At low energies, vacancies are found in the first or second monolayer of the substrate, while interstitials are found two to three times
(c) “0 2004006008001000 Energv(eV) Fig. 3. Energy dependence of the spatial distribution of vacancies and interstitials, ions, interstitials and substitutional ions, replacement events and relocation events for the case of %r on Si. (a) depicts the energy dependence of their respective range. In (b) the ratio of ion range versus vacancy range (solid circles), and the range ratio of interstitials versus vacancies (open circles) are shown as function energy. The range separation ratio (filled and empty squares), normalized to the vacancy range are found to quickly converge to 0 past 200 eV. In (c), the location of the true maximum of the ion, interstitials and vacancies distributions rather than the range inferred from a Gaussian fit as in (a) and (b), is shown.
III. ION-ENHANCED DEPOSITION
330
N. Herbots et al. / Mechanisms
18DGrowth:
5
. 4 L
Si on I.. ev
----
200
Si
r.,....,...,
5 Si
ion
flux:
ev
: -.-._._
during thin-film growth
Recombination below the surface should be favored to obtain oxide growth (100 eV ,< E < 1 keV) as found in ref. [8].
4. Conclusions
--__ --__
-1
1 -1
~....I....,....,....~
-*0
5
Time ‘&ee.)
15
20-2
Fig. 4. Thin-film growth for a 30Sif ion beam with a dose rate of lOI4 ions cm-*SC’ as a function of time, with ion energy as parameter. The slope of the curves provide the growth rates discussedin the text.
vacancies is found to be confined within the first few monolayers below 200 eV, and increases parabolically with energy while the depth of the interstitials increases logarithmically. Practically, these calculations show that below 200 eV, little damage is left behind from the Ar bombardment but the rate of Ar retention increases dramatically. This would probably lead to clustering, which is not included in our calculation. The dependence of defect distributions upon ion energy and the effect of thermal diffusion determines the dependence of thin-film growth rate upon energy, as shown in fig. 4. In fig. 4, the surface motion as a function of time, which translates into a linear growth rate, is shown for the case of deposition 3oSi on Si, for different energies. The growth rates stay similar for energies below 200 eV. Above 200 eV, the growth rate drops about twice as fast as sputtering losses [7,15]. The number of interstitials contributing to surface growth is then below the number of ions impinging the target. A larger fraction of interstitials remains trapped within the substrate. This last observation is important for optimizing the energy for growth rate and properties such as epitaxy. Specifically, the selection of energies maintaining the sputtering yield below 1 [6,16] is not sufficient to obtain a thin-film growth rate larger than 0 and is overridden by the necessity of favoring specific defect distribution and recombination mechanism if epitaxial or high-density films are to be grown. These observations are consistent with reports from independent groups that beyond 200 eV ion bombardment during epitaxial growth leads to residual damage and poor electronic properties [20]. In the case of ion beam oxidation (IBO), the reverse is true: surface recombination (and thus possible desorption) is to be avoided.
In summary, the mechanism of athermal enhancement mobility by low-energy ion beam can be described at temperatures below plastic deformation (“low temperature”) in semiconductor materials in terms of atomic collisions, ER and point defect thermal diffusion. The model based on this description gives us as intuitive understanding of the mechanism by which a thin film can be grown from an ion beam. It also allows us to compute quantitatively in what conditions of temperature, dose rate, and energy such processes lead to net thin-film growth, epitaxy and compound formation for conceptually simple techniques such as ion beam deposition (IBD) and ion beam oxidation (IBO).
References VI C. King, J.L. Hoyt, C.M. Cronet, J.F. Gibbons, M.P. Scott and J. Tarur, IEEE Trans. Electron. Dev. Lett. EDL-10 (1989) 52. PI P. Tsukizoe, T. Nakai and N. Ohmae, J. Appl. Phys. 42 (1976) 4770. [31 K. Yagi, S. Tamura and T. Tokuyama, Jpn. J. Appl. Phys. 16 (1977) 245.
[41 P.C. Zalm and L.J. Beckers, Appl. Phys. Lett. 41 (1982) 167. 151 N. Herbots, B.R. Appleton, T.S. Noggle, S.J. Pennycook, R.A. Zuhr and D.M. Zhener, Semiconductor-based Heterostructures: Interfacial Structure and Stability, eds. M.L. Green et al. (1986) pp. 335-349. 161 N. Herbots, O.C. Hellman, P.A. Cullen and 0. Vancauwenberghe, Am. Vat. Sot. Series 4, ed. G. Luckovsky, Deposition and growth: limits for microelectronics, AIP Conf. Proc. 167, ed. G.W. Rubloff (New York, 1988) pp. 259-289. 171 0. Vancauwenberghe, N. Herbots and O.C. Helhnan, Sot. Photo-optical Instr. Eng. Proc. 1285 (SPIE, Bellingham, 1990) pp. 47-59. [8] 0. Vancauwenberghe, N. Herbots, H. Manoharan and M. Ahrens, J. Vat. Sci. Technol. (1991) in press; N. Herbots, 0. Vancauwenberghe and H. Manoharan, 36th Am. Vat. Sot. Meeting, AVS Discovery Session, Boston, MA, October 23-27, 1989; N. Herbots, 0. Vancauwenberghe and H. Manoharan, Americ. Phys. Sot. March Meeting, Anaheim, CA, March 12-16, 1990. [91 N. Herbots and O.C. Hellman, US Patent no. 4800100 (1989). WI M.T. Robinson and I.M. Torrens, Phys. Rev. B9 (1962) 5008. [ill J.P. Biersack and L.J. Haggmark, Nucl. Instr. and Meth. 174 (1980) 257;
N. Herbots et al. / Mechanisms during thin-film growth see also: J.P. Biersack and W. Eckstein, J. Appl. Phys. A34 (1984) 13. 1121A. Tenner, Rainbow Scattering, Ph.D. thesis, FOM-Instituut, Amsterdam (1986). (131 B. Dodson, Phys. Rev. B36 (1987) 1068. [14] B.J. Garrison, M.T. Miller and D.W. Brenner, Atomic Scale calculations in Materials Science, eds. J. Tersoff, D. Vanderbilt and V. Vitek, Mater. Res. Sot. Symp. Proc. 141 (1990) 419. [15] N. Herbots, B.R. Appleton, TX NoggIe, R.A. Zuhr and S.J. Pennycook, Nucl. Instr. and Meth. B13 (1986) 250. [16] N. Herbots, B.R. Appleton, S.J. Pennycook, T.S. NoggIe and R.A. Zuhr, in: Beam-Solid Interactions and Phase Transformations, eds. H. Kurtz, G.L. Olson and J.M. Poate, Mater. Res. Sot. Symp. Proc. 51 (1986) 3669.
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[17] S. Roorda, J.M. Poate, DC. Jacobson, D.J. Eaglesham, B.S. Dennis, S. Dieker, W.C. Sinke and F. Spaepen, Solid State Commun. 75 (1990) 197. [18] U. Landman, W.D. Luedtke and M.W. Ribarsky, J. Vat. Sci. Technol. A7 (1989) 2829. [19] U. Gosele, in: Semiconductor Silicon 1986, eds. H.R Huff, T. Abe and B. Kolbesen (Pemtington, Electrochem. Sot., 1986) p. 541. [20] M.A. Hasen, J. Knoll, S.A. Bamett, J.E. Sundgren, L.C. Marker& A. Rockett and J.E. Greene, J. Appl. Phys. 65 (1989) 172; also: S.S. Iyer and E. Kasper, private communication.
III. ION-ENHANCED
DEPOSITION