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Controlling the density distribution of SiC nanocrystals for the ion beam synthesis of buried SiC layers in silicon
Nuclear Instruments and Methods in Physics Research B 147 (1999) 249±255
Controlling the density distribution of SiC nanocrystals for the ion beam synthesis of buried SiC layers in silicon J.K.N. Lindner *, B. Stritzker Universit at Augsburg, Institut f ur Physik, Memminger Str. 6, D-86135 Augsburg, Germany
Abstract The depth distribution of SiC nanocrystals formed during high-dose implantation of carbon ions into silicon at conditions suitable for the ion beam synthesis of buried SiC layers in silicon is studied in this paper. For implantation temperatures of 400±600°C and dose rates of 1012 ) 1013 C /cm2 s, SiC precipitates in crystalline silicon are observed to be of approximately equal size, independent of the depth position beneath the surface. Ballistic destruction of small precipitates and diculties in precipitate growth are thought to be responsible for the observed narrow size distribution. The destruction of precipitates may lead to the simultaneous release of a superthreshold concentration of carbon atoms resulting in a carbon-induced amorphization of the silicon host lattice. The local reduction of the number density of SiC nanocrystals involved with this amorphization can be used to tailor discontinuous depth distributions of oriented SiC precipitates providing ideal starting conditions for the synthesis of well-de®ned single-crystalline SiC layers in silicon. Ó 1999 Elsevier Science B.V. All rights reserved. PACS: 61.72.Tt; 68.55.Jk; 68.55.Ln; 61.72Qq; 85.40.Ry Keywords: Ion beam synthesis; Nanocrystals; SiC; Silicon; Amorphization; Ballistic eects
1. Introduction In the ion beam synthesis of buried compound layers in silicon, the as-implanted state de®nes the boundary conditions for the structural evolution during annealing. Buried homogeneous epitaxial metal disilicide layers of CoSi2 or NiSi2 in silicon, for instance, have usually been formed from a nearly Gaussian depth distribution of corresponding disilicide precipitates which were epit-
* Corresponding author. Tel.: +49 821 5983445; fax: +49 821 5983425; e-mail: [email protected]
axially incorporated in a crystalline silicon matrix. These precipitates show a pronounced size distribution, where largest particles are present at [1] or close to [2] the maximum of the concentration depth pro®le. It is generally accepted that the layer formation during annealing proceeds by an Ostwald type of ripening process, i.e. by the growth of larger precipitates in the centre of the distribution on the expense of smaller dissolving particles in the distribution wings. Totally dierent starting conditions have been successfully used to form buried polycrystalline a-Si3 N4 layers in silicon [3]. They grow during annealing of buried amorphous Six Ny layers of variable composition, where nitrogen is
0168-583X/98/$ ± see front matter Ó 1999 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 9 8 ) 0 0 5 9 8 - 9
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J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
bonded to silicon atoms but has not formed any crystalline precipitates during implantation. Crystallization takes place during annealing and most probably starts at the moving amorphous/crystalline (a/c) interfaces. From a series of investigations [4±7] on the ion beam synthesis of buried SiC layers in silicon it emerges that the formation of buried single-crystalline 3C±SiC layers in silicon requires a dierent route, if annealing temperatures close to the melting point of silicon are to be avoided. As will be highlighted in this paper, it involves the control of the density distribution of equally sized SiC nanocrystals, which are formed either in a crystalline or an amorphous matrix during high-dose, high-temperature carbon implantation into silicon. First it will be shown that the formation of this amorphous matrix is induced by the presence of a superthreshold concentration of carbon atoms released from ballistically destroyed SiC precipitates. The temperature dependence of this amorphization and the SiC precipitate destruction involved can be exploited to tailor the depth distribution of SiC nanocrystals. 2. Experimental Float zone Si(1 0 0) wafers were implanted with 180 keV C ions at doses of 4.3 or 8.5 ´ 1017 C / cm2 and a ®xed dose rate of 2 ´ 1013 C /cm2 s. Target temperatures were kept constant at values of 400°C, 500°C and 600°C by resistive heating and temperature monitoring with a thermocouple attached to the target holder. In addition, some implantations were performed at 500°C with a reduced dose rate of 2 ´ 1012 C /cm2 s. Moreover, some samples are considered, which were implanted [4±6] at beam heating conditions with the higher dose rate, leading to target temperatures of 330±440°C. Sample characterization in the as-implanted state and after 5±10 h annealing at 1250°C in ¯owing Ar atmosphere was performed with 1.8 MeV He RBS/channeling at Cornell geometry and with cross-sectional TEM at 300 keV. Both high resolution (HREM) and Moire dark ®eld imaging techniques were used to study the microstructure.
3. Results and discussion 3.1. Processes aecting the SiC precipitate density distributions in the as-implanted state Fig. 1 shows RBS random and channeling pro®les of silicon implanted at 400°C and 500°C with 180 keV C ions at a dose of 4.3 ´ 1017 C / cm2 . This dose corresponds to approximately half the stoichiometry dose necessary to obtain a peak composition of C/Si 1. As the random spectra for both temperatures are identical, only one of them is displayed. The depth distribution of implanted carbon is re¯ected by a minimum in the Si part of the random spectrum around channel no. 180. The channeling spectrum for 500°C just touches the random spectrum, indicating the onset of amorphization. It is important to note that at this temperature amorphization sets in at the depth of maximum carbon concentration, indicating that amorphous phase formation is related to the presence of carbon impurities. Due to the low nuclear stopping power of carbon ions in silicon, one would not expect pure silicon to turn amorphous at this temperature and dose in the absence of compositional eects. In comparison,
Fig. 1. RBS/channeling spectra of Si(1 0 0) implanted with 4.3 ´ 1017 C /cm2 at 400°C and 500°C, corresponding random spectrum (400°C) and simulated reference spectrum of pure silicon.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
for implantations with the much heavier Si self ions at 177°C a critical dose for amorphization of 4.5 ´ 1017 Si/cm2 has been found [8]. At the lower target temperature of 400°C (Fig. 1), a broader RBS amorphous zone has developed, approximately centred about the peak of the carbon pro®le. The comparison with the random spectrum indicates that the amorphous layer extends over a depth interval, where a ®xed carbon concentration of about 17 at.% C is exceeded. The occurrence of a threshold carbon concentration for the amorphization of silicon is also observed in samples implanted at beam heating conditions and doses between 4 and 15 ´ 1017 C /cm2 (not shown), where the width of the RBS amorphous layer increases with the depth interval at which the threshold concentration is exceeded. The threshold carbon concentration slightly decreases with decreasing temperature. At suciently low target temperatures, the amorphization process is governed by the nuclear energy deposition, and consequently amorphization occurs at the upper wing of the carbon depth pro®le. This is obvious from studies [9] on Si(1 1 1) implanted with 0.8 MeV C ions at a target temperature of 227°C. The presence of amorphous layers is con®rmed by cross-sectional TEM (see below). Electron microscopy for doses of 4 ´ 1017 C /cm2 and above also shows that nanosized 3C±SiC crystallites have formed. Most precipitates are aligned with the silicon lattice and exhibit, when in c-Si surrounding, translational Moire fringes in Moire dark ®eld and high-resolution lattice images. Thus these particles are incoherent precipitates. Examples are given in Fig. 2 for a temperature of 600°C. This temperature has been chosen because the region in which SiC precipitates can be observed extends with increasing temperatures up to smaller depths (see also Fig. 3(d)). Moreover, one would expect the largest variations in particle size at the highest temperature. However, all SiC precipitates in the topmost 400 nm c-Si region are found to be identical in size, approximately 5 nm. This value is identical for all implantation temperatures used and also for the smaller dose rate and 500°C. In Fig. 2, in a depth zone from 380±460 nm, Moire fringes vanish, indicating the absence of c-Si inclusions around the depth of maximum carbon
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Fig. 2. Cross-sectional TEM dark ®eld image of Si(1 0 0) implanted at 600°C with 180 keV C ions at a dose of 8.5 ´ 1017 C /cm2 , showing the depth distribution of SiC precipitates by the Moire fringes they produce, and high resolution image (inset) of individual precipitates, causing translational Moire contrast of {1 1 1} lattice planes. Some precipitates are marked by rings and arrows.
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concentration. In the depth zone above, the density of SiC precipitates increases quite monotonously with depth. A much more abrupt density increase is observed for 500 °C, and for 400 °C the depth distribution is even discontinuous; i.e., there is a SiC precipitate denuded zone, which coincides with the upper part of the amorphous layer formed at this temperature. This is displayed in the schematic drawings of Fig. 3, which are based on detailed XTEM investigations. The entire amorphous layer observed after implantation at 400°C (Fig. 3(a)) is subdivided into four dierent zones: (a) the upper, SiC precipitate denuded zone; (b) a layer containing relatively large precipitates, which are either aligned with the substrate (9 nm) or at random orientation (17 nm); (c) a zone of low precipitate density; (d) a zone with a high density of small (4 nm) oriented 3C±SiC precipitates. Based on dose dependent studies [6] using beam heating samples, the occurrence of a majority of
oriented precipitates within an amorphous layer can be explained by the fact that precipitate formation starts at doses beneath the amorphization dose. The amorphization is accompanied by a signi®cant reduction of the number density of SiC precipitates. The small number of precipitates with diameters of less than 5 nm in the upper 400 nm of samples and the complete extinction of SiC nanocrystallites in the upper part of the amorphous zone indicate that ballistic destruction of precipitates may occur. Ofcourse, 5 nm is much larger than the critical size for the nucleation of SiC precipitates at the present temperatures, as signalized by the presence of the few smaller precipitates. In fact, SiC precipitates as small as 2 nm have been observed in low-dose carbon implanted silicon after 850°C annealing [10], and it has been suggested that the nucleation of SiC precipitates is preceeded by the agglomeration of C±Si dimers (T < 850°C).
Fig. 3. Schematical description of the microstructure of Si(1 0 0) implanted at 400°C and 600°C, respectively, with 180 keV C ions at a dose of 8.5 ´ 1017 C /cm2 , in the as-implanted state (a.i.) and after 5 h annealing at 1250°C, as revealed by cross-sectional TEM.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
In order to roughly estimate the strength of ballistic mixing at precipitate/matrix interfaces, TRIM96 [11] simulations of collision cascades generated in Si/SiC/Si systems were performed [7]. From the carbon recoil atom concentration in the vicinity of the interfaces it follows that in the absence of (obviously present) precipitate regrowth mechanisms a dose of the order of few 1016 C /cm2 is required to destroy a 3 nm precipitate. However, the incoherent precipitates in c-Si are unable to grow or regrow signi®cantly, mainly because of their high interfacial energy [12]. Therefore, the strong supersaturation and the ballistic destruction of small particles should lead to a quasisteady-state between the formation and the destruction of precipitates, ``quasi'' because the density of precipitates increases with dose. At a certain precipitate density, the amount of carbon which is simultaneously released from dissolving precipitates, exceeds a certain threshold and the amorphous phase is stabilized. Once an amorphous layer has been formed, however, surviving SiC precipitates in an amorphous surrounding may grow at an increased rate. This is indicated by the occurrence of comparatively large precipitates within the amorphous layer (b) and can be explained by a smaller SiC/aSi interfacial energy compared to particles of the same size in c-Si. The presence of misoriented SiC crystallites points to SiC nucleation within the amorphous matrix. Since the thermal crystallization of ion beam synthesized amorphous SiC layers requires temperatures in excess of 800°C [13], this nucleation should be considered as beam induced. The high density of slightly smaller SiC precipitates in the lowest part (d) of the amorphous layer can be ascribed to the fact that this depth corresponds to the range of 180 keV C ions. Therefore the highest generation rate of SiC precipitates should exist at this depth (at least towards the end of implantation, if one takes sputter eects into account). 3.2. Resulting SiC layers after annealing It has been reported earlier [7] that annealing for 5 h at 1250°C leads to homogeneous, stoichiometric 3C±SiC layers for all implantation tem-
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peratures considered. For a temperature of 600°C, the SiC layers are single-crystalline with a high density of defects. However, the interface with the Si top layer is extremely poor, being governed by a high density of SiC precipitates. Most abrupt interfaces were obtained at 400°C, however, buried SiC layers were then subdivided into a polycrystalline upper and an epitaxial lower part. This is due to the formation of SiC crystallites at random orientation during implantation and possibly also during annealing. A good compromise was the use of an intermediate temperature of 500°C [7]. Single-crystalline SiC layers with even better interfaces can be obtained using a temperature of 450°C. The best layers with respect to both interfacial width and crystal quality, were obtained using experimental conditions which aimed to exploit our understanding of SiC precipitate formation/destruction mechanisms described above for the control of the precipitate density depth distribution. For this purpose, a stoichiometric implantation was performed with 93% of the dose at a high implantation temperature of 500°C and 7% of the dose after reducing the temperature down to 250°C. The purpose of the ®rst implantation step was to generate a large number of oriented SiC precipitates in a limited depth interval, while avoiding an early amorphization with the possibility of misaligned SiC nucleation and growth. The second implantation step was performed at a reduced temperature in order to induce amorphization in the region of the upper pro®le ¯ank. In fact, dierent from single-temperature implantations at 500°C, cross-section TEM reveals the presence of a 70 and a 20 nm thick amorphous layer denuded of SiC precipitates above and underneath the central precipitate layer, characterized by a high density of mostly oriented SiC nanocrystals. After annealing at 1250°C, RBS measurements (Fig. 4) show a rectangular carbon depth distribution corresponding to a well-de®ned buried stoichiometric SiC layer of 155 nm thickness, covered with 290 nm of nearly pure silicon and a 75 nm thermal oxide ®lm resulting from oxygen impurities in the annealing ambient. Crosssection TEM (Fig. 5) shows the presence of a distinct epitaxial 3C±SiC layer, with few isolated SiC precipitates in the top layer. These are mostly
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misaligned, have grown up to 50 nm in size and are therefore hard to dissolve. The buried layer itself gives rise to a single-crystalline diraction pattern of epitaxial SiC in silicon. The contrasts within this layer visible in Fig. 5 are mostly due to a high density of {1 1 1} planar defects, as determined from high-resolution images. 4. Conclusions
Fig. 4. RBS spectrum of Si(1 0 0) implanted at 500°C and 250°C with 180 keV C ions at a dose of 7.9 and 0.6 ´ 1017 C /cm2 , respectively, after 10 h annealing at 1250°C. The dashed line is the simulated reference spectrum of pure silicon, and the solid line is a simulated spectrum of the SiO2 /Si/SiC/Si layer sequence described in the text.
The high-dose implantation of carbon ions into silicon at 400±600°C leads to the formation of oriented SiC nanocrystals of identical size in crystalline silicon, fairly independent of temperature, beam current density and depth position. Precipitate growth is hampered by the incoherent precipitate interfaces in c-Si, and by the rapid ballistic destruction of small crystallites. A carboninduced amorphization involving the destruction of SiC nanocrystals is observed to occur at a temperature dependent threshold concentration of
Fig. 5. Bright ®eld (a) and (1 1 1) SiC dark ®eld (b) cross-sectional TEM micrographs of the sample of Fig. 4. Inset: selected area diraction pattern of the buried layer.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
carbon. SiC precipitates may grow faster in an amorphous surrounding, with the possibility of being misaligned. Using reduced implantation temperatures, the amorphization can be exploited to cut o the upper tail of the SiC precipitate distribution, which exists at higher temperatures, in order to obtain a box-like depth distribution of oriented SiC nanocrystals ideally suited to sythesize well-de®ned buried single-crystalline 3C±SiC layers in silicon. References [1] E.A.H. Dekempeneer, J.J.M. Ottenheim, D.E.W. Vandenhoudt, C.W.T. Bulle-Lieuwma, E.G.C. Lathouwers, Nucl. Instr. and Meth. B 55 (1991) 769. [2] J.K.N. Lindner, T. Klassen, E.H. te Kaat, Nucl. Instr. and Meth. B 59/60 (1991) 655. [3] K.J. Reeson, Nucl. Instr. and Meth. B 19/20 (1987) 269. [4] J.K.N. Lindner, A. Frohnwieser, B. Rauschenbach, B. Stritzker, Mater. Res. Soc. Symp. Proc. 354 (1995) 171±176.
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[5] J.K.N. Lindner, K. Volz, B. Stritzker, in: S. Nakashima, et al. (Eds.), Silicon Carbide and Related Materials 1995, IOP Conf. Ser. Vol. 142, 145, Institute of Physics Publishing, Bristol, 1996. [6] J.K.N. Lindner, K. Volz, B. Stritzker, Mater. Res. Soc. Symp. Proc. 438 (1997). [7] J.K.N. Lindner, W. Reiber, B. Stritzker, in: G. Pensl, et al. (Eds.), Silicon Carbide, III-Nitrides and Related Materials, Materials Science Forum Vol. 264±268 (1998) 215. [8] J. Belz, K.F. Heidemann, H.F. Kappert, E. te Kaat, Phys. Stat. Sol. (a) 76 (1983) K81. [9] A. Frohnwieser, Thesis, University of Augsburg 1995, unpublished. [10] P. Werner, R. Koegler, W. Skorupa, D. Eichler, in: E. Ishidida et al. (Eds.), Proceedings of the 11th International Conference on Ion Implantation Technology, Austin TX, USA, June 16±21, 1996, p. 675. [11] J.F. Ziegler, J.P. Biersack, U. Littmark, in: J.F. Ziegler (Ed.), The Stopping and Range of Ions in Matter, Vol. 1, Pergamon Press, New York, 1985. [12] W.J. Taylor, T.Y. Tan, U. G osele, Appl. Phys. Lett. 62 (1993) 3336. [13] J.A. Borders, S.T. Picraux, W. Beezhold, Appl. Phys. Lett. 18 (1971) 509.
Controlling the density distribution of SiC nanocrystals for the ion beam synthesis of buried SiC layers in silicon J.K.N. Lindner *, B. Stritzker Universit at Augsburg, Institut f ur Physik, Memminger Str. 6, D-86135 Augsburg, Germany
Abstract The depth distribution of SiC nanocrystals formed during high-dose implantation of carbon ions into silicon at conditions suitable for the ion beam synthesis of buried SiC layers in silicon is studied in this paper. For implantation temperatures of 400±600°C and dose rates of 1012 ) 1013 C /cm2 s, SiC precipitates in crystalline silicon are observed to be of approximately equal size, independent of the depth position beneath the surface. Ballistic destruction of small precipitates and diculties in precipitate growth are thought to be responsible for the observed narrow size distribution. The destruction of precipitates may lead to the simultaneous release of a superthreshold concentration of carbon atoms resulting in a carbon-induced amorphization of the silicon host lattice. The local reduction of the number density of SiC nanocrystals involved with this amorphization can be used to tailor discontinuous depth distributions of oriented SiC precipitates providing ideal starting conditions for the synthesis of well-de®ned single-crystalline SiC layers in silicon. Ó 1999 Elsevier Science B.V. All rights reserved. PACS: 61.72.Tt; 68.55.Jk; 68.55.Ln; 61.72Qq; 85.40.Ry Keywords: Ion beam synthesis; Nanocrystals; SiC; Silicon; Amorphization; Ballistic eects
1. Introduction In the ion beam synthesis of buried compound layers in silicon, the as-implanted state de®nes the boundary conditions for the structural evolution during annealing. Buried homogeneous epitaxial metal disilicide layers of CoSi2 or NiSi2 in silicon, for instance, have usually been formed from a nearly Gaussian depth distribution of corresponding disilicide precipitates which were epit-
* Corresponding author. Tel.: +49 821 5983445; fax: +49 821 5983425; e-mail: [email protected]
axially incorporated in a crystalline silicon matrix. These precipitates show a pronounced size distribution, where largest particles are present at [1] or close to [2] the maximum of the concentration depth pro®le. It is generally accepted that the layer formation during annealing proceeds by an Ostwald type of ripening process, i.e. by the growth of larger precipitates in the centre of the distribution on the expense of smaller dissolving particles in the distribution wings. Totally dierent starting conditions have been successfully used to form buried polycrystalline a-Si3 N4 layers in silicon [3]. They grow during annealing of buried amorphous Six Ny layers of variable composition, where nitrogen is
0168-583X/98/$ ± see front matter Ó 1999 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 9 8 ) 0 0 5 9 8 - 9
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J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
bonded to silicon atoms but has not formed any crystalline precipitates during implantation. Crystallization takes place during annealing and most probably starts at the moving amorphous/crystalline (a/c) interfaces. From a series of investigations [4±7] on the ion beam synthesis of buried SiC layers in silicon it emerges that the formation of buried single-crystalline 3C±SiC layers in silicon requires a dierent route, if annealing temperatures close to the melting point of silicon are to be avoided. As will be highlighted in this paper, it involves the control of the density distribution of equally sized SiC nanocrystals, which are formed either in a crystalline or an amorphous matrix during high-dose, high-temperature carbon implantation into silicon. First it will be shown that the formation of this amorphous matrix is induced by the presence of a superthreshold concentration of carbon atoms released from ballistically destroyed SiC precipitates. The temperature dependence of this amorphization and the SiC precipitate destruction involved can be exploited to tailor the depth distribution of SiC nanocrystals. 2. Experimental Float zone Si(1 0 0) wafers were implanted with 180 keV C ions at doses of 4.3 or 8.5 ´ 1017 C / cm2 and a ®xed dose rate of 2 ´ 1013 C /cm2 s. Target temperatures were kept constant at values of 400°C, 500°C and 600°C by resistive heating and temperature monitoring with a thermocouple attached to the target holder. In addition, some implantations were performed at 500°C with a reduced dose rate of 2 ´ 1012 C /cm2 s. Moreover, some samples are considered, which were implanted [4±6] at beam heating conditions with the higher dose rate, leading to target temperatures of 330±440°C. Sample characterization in the as-implanted state and after 5±10 h annealing at 1250°C in ¯owing Ar atmosphere was performed with 1.8 MeV He RBS/channeling at Cornell geometry and with cross-sectional TEM at 300 keV. Both high resolution (HREM) and Moire dark ®eld imaging techniques were used to study the microstructure.
3. Results and discussion 3.1. Processes aecting the SiC precipitate density distributions in the as-implanted state Fig. 1 shows RBS random and channeling pro®les of silicon implanted at 400°C and 500°C with 180 keV C ions at a dose of 4.3 ´ 1017 C / cm2 . This dose corresponds to approximately half the stoichiometry dose necessary to obtain a peak composition of C/Si 1. As the random spectra for both temperatures are identical, only one of them is displayed. The depth distribution of implanted carbon is re¯ected by a minimum in the Si part of the random spectrum around channel no. 180. The channeling spectrum for 500°C just touches the random spectrum, indicating the onset of amorphization. It is important to note that at this temperature amorphization sets in at the depth of maximum carbon concentration, indicating that amorphous phase formation is related to the presence of carbon impurities. Due to the low nuclear stopping power of carbon ions in silicon, one would not expect pure silicon to turn amorphous at this temperature and dose in the absence of compositional eects. In comparison,
Fig. 1. RBS/channeling spectra of Si(1 0 0) implanted with 4.3 ´ 1017 C /cm2 at 400°C and 500°C, corresponding random spectrum (400°C) and simulated reference spectrum of pure silicon.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
for implantations with the much heavier Si self ions at 177°C a critical dose for amorphization of 4.5 ´ 1017 Si/cm2 has been found [8]. At the lower target temperature of 400°C (Fig. 1), a broader RBS amorphous zone has developed, approximately centred about the peak of the carbon pro®le. The comparison with the random spectrum indicates that the amorphous layer extends over a depth interval, where a ®xed carbon concentration of about 17 at.% C is exceeded. The occurrence of a threshold carbon concentration for the amorphization of silicon is also observed in samples implanted at beam heating conditions and doses between 4 and 15 ´ 1017 C /cm2 (not shown), where the width of the RBS amorphous layer increases with the depth interval at which the threshold concentration is exceeded. The threshold carbon concentration slightly decreases with decreasing temperature. At suciently low target temperatures, the amorphization process is governed by the nuclear energy deposition, and consequently amorphization occurs at the upper wing of the carbon depth pro®le. This is obvious from studies [9] on Si(1 1 1) implanted with 0.8 MeV C ions at a target temperature of 227°C. The presence of amorphous layers is con®rmed by cross-sectional TEM (see below). Electron microscopy for doses of 4 ´ 1017 C /cm2 and above also shows that nanosized 3C±SiC crystallites have formed. Most precipitates are aligned with the silicon lattice and exhibit, when in c-Si surrounding, translational Moire fringes in Moire dark ®eld and high-resolution lattice images. Thus these particles are incoherent precipitates. Examples are given in Fig. 2 for a temperature of 600°C. This temperature has been chosen because the region in which SiC precipitates can be observed extends with increasing temperatures up to smaller depths (see also Fig. 3(d)). Moreover, one would expect the largest variations in particle size at the highest temperature. However, all SiC precipitates in the topmost 400 nm c-Si region are found to be identical in size, approximately 5 nm. This value is identical for all implantation temperatures used and also for the smaller dose rate and 500°C. In Fig. 2, in a depth zone from 380±460 nm, Moire fringes vanish, indicating the absence of c-Si inclusions around the depth of maximum carbon
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Fig. 2. Cross-sectional TEM dark ®eld image of Si(1 0 0) implanted at 600°C with 180 keV C ions at a dose of 8.5 ´ 1017 C /cm2 , showing the depth distribution of SiC precipitates by the Moire fringes they produce, and high resolution image (inset) of individual precipitates, causing translational Moire contrast of {1 1 1} lattice planes. Some precipitates are marked by rings and arrows.
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J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
concentration. In the depth zone above, the density of SiC precipitates increases quite monotonously with depth. A much more abrupt density increase is observed for 500 °C, and for 400 °C the depth distribution is even discontinuous; i.e., there is a SiC precipitate denuded zone, which coincides with the upper part of the amorphous layer formed at this temperature. This is displayed in the schematic drawings of Fig. 3, which are based on detailed XTEM investigations. The entire amorphous layer observed after implantation at 400°C (Fig. 3(a)) is subdivided into four dierent zones: (a) the upper, SiC precipitate denuded zone; (b) a layer containing relatively large precipitates, which are either aligned with the substrate (9 nm) or at random orientation (17 nm); (c) a zone of low precipitate density; (d) a zone with a high density of small (4 nm) oriented 3C±SiC precipitates. Based on dose dependent studies [6] using beam heating samples, the occurrence of a majority of
oriented precipitates within an amorphous layer can be explained by the fact that precipitate formation starts at doses beneath the amorphization dose. The amorphization is accompanied by a signi®cant reduction of the number density of SiC precipitates. The small number of precipitates with diameters of less than 5 nm in the upper 400 nm of samples and the complete extinction of SiC nanocrystallites in the upper part of the amorphous zone indicate that ballistic destruction of precipitates may occur. Ofcourse, 5 nm is much larger than the critical size for the nucleation of SiC precipitates at the present temperatures, as signalized by the presence of the few smaller precipitates. In fact, SiC precipitates as small as 2 nm have been observed in low-dose carbon implanted silicon after 850°C annealing [10], and it has been suggested that the nucleation of SiC precipitates is preceeded by the agglomeration of C±Si dimers (T < 850°C).
Fig. 3. Schematical description of the microstructure of Si(1 0 0) implanted at 400°C and 600°C, respectively, with 180 keV C ions at a dose of 8.5 ´ 1017 C /cm2 , in the as-implanted state (a.i.) and after 5 h annealing at 1250°C, as revealed by cross-sectional TEM.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
In order to roughly estimate the strength of ballistic mixing at precipitate/matrix interfaces, TRIM96 [11] simulations of collision cascades generated in Si/SiC/Si systems were performed [7]. From the carbon recoil atom concentration in the vicinity of the interfaces it follows that in the absence of (obviously present) precipitate regrowth mechanisms a dose of the order of few 1016 C /cm2 is required to destroy a 3 nm precipitate. However, the incoherent precipitates in c-Si are unable to grow or regrow signi®cantly, mainly because of their high interfacial energy [12]. Therefore, the strong supersaturation and the ballistic destruction of small particles should lead to a quasisteady-state between the formation and the destruction of precipitates, ``quasi'' because the density of precipitates increases with dose. At a certain precipitate density, the amount of carbon which is simultaneously released from dissolving precipitates, exceeds a certain threshold and the amorphous phase is stabilized. Once an amorphous layer has been formed, however, surviving SiC precipitates in an amorphous surrounding may grow at an increased rate. This is indicated by the occurrence of comparatively large precipitates within the amorphous layer (b) and can be explained by a smaller SiC/aSi interfacial energy compared to particles of the same size in c-Si. The presence of misoriented SiC crystallites points to SiC nucleation within the amorphous matrix. Since the thermal crystallization of ion beam synthesized amorphous SiC layers requires temperatures in excess of 800°C [13], this nucleation should be considered as beam induced. The high density of slightly smaller SiC precipitates in the lowest part (d) of the amorphous layer can be ascribed to the fact that this depth corresponds to the range of 180 keV C ions. Therefore the highest generation rate of SiC precipitates should exist at this depth (at least towards the end of implantation, if one takes sputter eects into account). 3.2. Resulting SiC layers after annealing It has been reported earlier [7] that annealing for 5 h at 1250°C leads to homogeneous, stoichiometric 3C±SiC layers for all implantation tem-
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peratures considered. For a temperature of 600°C, the SiC layers are single-crystalline with a high density of defects. However, the interface with the Si top layer is extremely poor, being governed by a high density of SiC precipitates. Most abrupt interfaces were obtained at 400°C, however, buried SiC layers were then subdivided into a polycrystalline upper and an epitaxial lower part. This is due to the formation of SiC crystallites at random orientation during implantation and possibly also during annealing. A good compromise was the use of an intermediate temperature of 500°C [7]. Single-crystalline SiC layers with even better interfaces can be obtained using a temperature of 450°C. The best layers with respect to both interfacial width and crystal quality, were obtained using experimental conditions which aimed to exploit our understanding of SiC precipitate formation/destruction mechanisms described above for the control of the precipitate density depth distribution. For this purpose, a stoichiometric implantation was performed with 93% of the dose at a high implantation temperature of 500°C and 7% of the dose after reducing the temperature down to 250°C. The purpose of the ®rst implantation step was to generate a large number of oriented SiC precipitates in a limited depth interval, while avoiding an early amorphization with the possibility of misaligned SiC nucleation and growth. The second implantation step was performed at a reduced temperature in order to induce amorphization in the region of the upper pro®le ¯ank. In fact, dierent from single-temperature implantations at 500°C, cross-section TEM reveals the presence of a 70 and a 20 nm thick amorphous layer denuded of SiC precipitates above and underneath the central precipitate layer, characterized by a high density of mostly oriented SiC nanocrystals. After annealing at 1250°C, RBS measurements (Fig. 4) show a rectangular carbon depth distribution corresponding to a well-de®ned buried stoichiometric SiC layer of 155 nm thickness, covered with 290 nm of nearly pure silicon and a 75 nm thermal oxide ®lm resulting from oxygen impurities in the annealing ambient. Crosssection TEM (Fig. 5) shows the presence of a distinct epitaxial 3C±SiC layer, with few isolated SiC precipitates in the top layer. These are mostly
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misaligned, have grown up to 50 nm in size and are therefore hard to dissolve. The buried layer itself gives rise to a single-crystalline diraction pattern of epitaxial SiC in silicon. The contrasts within this layer visible in Fig. 5 are mostly due to a high density of {1 1 1} planar defects, as determined from high-resolution images. 4. Conclusions
Fig. 4. RBS spectrum of Si(1 0 0) implanted at 500°C and 250°C with 180 keV C ions at a dose of 7.9 and 0.6 ´ 1017 C /cm2 , respectively, after 10 h annealing at 1250°C. The dashed line is the simulated reference spectrum of pure silicon, and the solid line is a simulated spectrum of the SiO2 /Si/SiC/Si layer sequence described in the text.
The high-dose implantation of carbon ions into silicon at 400±600°C leads to the formation of oriented SiC nanocrystals of identical size in crystalline silicon, fairly independent of temperature, beam current density and depth position. Precipitate growth is hampered by the incoherent precipitate interfaces in c-Si, and by the rapid ballistic destruction of small crystallites. A carboninduced amorphization involving the destruction of SiC nanocrystals is observed to occur at a temperature dependent threshold concentration of
Fig. 5. Bright ®eld (a) and (1 1 1) SiC dark ®eld (b) cross-sectional TEM micrographs of the sample of Fig. 4. Inset: selected area diraction pattern of the buried layer.
J.K.N. Lindner, B. Stritzker / Nucl. Instr. and Meth. in Phys. Res. B 147 (1999) 249±255
carbon. SiC precipitates may grow faster in an amorphous surrounding, with the possibility of being misaligned. Using reduced implantation temperatures, the amorphization can be exploited to cut o the upper tail of the SiC precipitate distribution, which exists at higher temperatures, in order to obtain a box-like depth distribution of oriented SiC nanocrystals ideally suited to sythesize well-de®ned buried single-crystalline 3C±SiC layers in silicon. References [1] E.A.H. Dekempeneer, J.J.M. Ottenheim, D.E.W. Vandenhoudt, C.W.T. Bulle-Lieuwma, E.G.C. Lathouwers, Nucl. Instr. and Meth. B 55 (1991) 769. [2] J.K.N. Lindner, T. Klassen, E.H. te Kaat, Nucl. Instr. and Meth. B 59/60 (1991) 655. [3] K.J. Reeson, Nucl. Instr. and Meth. B 19/20 (1987) 269. [4] J.K.N. Lindner, A. Frohnwieser, B. Rauschenbach, B. Stritzker, Mater. Res. Soc. Symp. Proc. 354 (1995) 171±176.
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[5] J.K.N. Lindner, K. Volz, B. Stritzker, in: S. Nakashima, et al. (Eds.), Silicon Carbide and Related Materials 1995, IOP Conf. Ser. Vol. 142, 145, Institute of Physics Publishing, Bristol, 1996. [6] J.K.N. Lindner, K. Volz, B. Stritzker, Mater. Res. Soc. Symp. Proc. 438 (1997). [7] J.K.N. Lindner, W. Reiber, B. Stritzker, in: G. Pensl, et al. (Eds.), Silicon Carbide, III-Nitrides and Related Materials, Materials Science Forum Vol. 264±268 (1998) 215. [8] J. Belz, K.F. Heidemann, H.F. Kappert, E. te Kaat, Phys. Stat. Sol. (a) 76 (1983) K81. [9] A. Frohnwieser, Thesis, University of Augsburg 1995, unpublished. [10] P. Werner, R. Koegler, W. Skorupa, D. Eichler, in: E. Ishidida et al. (Eds.), Proceedings of the 11th International Conference on Ion Implantation Technology, Austin TX, USA, June 16±21, 1996, p. 675. [11] J.F. Ziegler, J.P. Biersack, U. Littmark, in: J.F. Ziegler (Ed.), The Stopping and Range of Ions in Matter, Vol. 1, Pergamon Press, New York, 1985. [12] W.J. Taylor, T.Y. Tan, U. G osele, Appl. Phys. Lett. 62 (1993) 3336. [13] J.A. Borders, S.T. Picraux, W. Beezhold, Appl. Phys. Lett. 18 (1971) 509.