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SiC precipitates formed in Si by simultaneous dual beam implantation of C+ and Si+ ions
Nuclear Instruments and Methods in Physics Research B 206 (2003) 989–993 www.elsevier.com/locate/nimb
SiC precipitates formed in Si by simultaneous dual beam implantation of Cþ and Siþ ions €gler R. Ko a
a,*
, F. Eichhorn a, A. M€ ucklich a, H. Reuther a, V. Heera a, W. Skorupa a, J.K.N. Lindner b
Forschungszentrum Rossendorf, Institut f€ur Ionenstrahlphysik und Materialforschung, P.O. Box 510119, 01314 Dresden, Germany b Universit€at Augsburg, Institut f€ur Physik, 86135 Augsburg, Germany
Abstract Nanometer-sized SiC precipitates were synthesized at 450 C in Si by simultaneous dual beam implantation of Cþ and Siþ ions and subsequent annealing. The results are compared with those of sequential dual beam implantation and of single beam implantation. Two types of SiC precipitates were found. Precipitates of type I with a diameter of d ¼ 4–5 nm consist of 3C-SiC epitaxially oriented with the Si matrix. They were formed already in the as-implanted state and do not grow further during subsequent annealing. The SiC precipitates of type II with d 10 nm are not oriented with the Si matrix and grow exclusively during the subsequent annealing. The high growth velocity, the misorientation in regard to the Si matrix and the lower concentration of type II precipitates can be explained by the assumption that these precipitates were formed in an amorphous substrate which modifies their interface energy. 2003 Elsevier Science B.V. All rights reserved. PACS: 61.72.T; 81.10.Jt Keywords: Silicon; SiC; Ion beam synthesis; Dual beam implantation
1. Introduction The ion beam synthesis (IBS) of SiC in Si has been investigated for many years because of the outstanding physical material properties of SiC making it applicable for devices working under extreme conditions. The basic processes of the formation and evolution of SiC precipitates during IBS and subsequent thermal treatment, their nu-
*
Corresponding author. Tel.: +49-351-260-3613; fax: +49351-260-3411. E-mail address: [email protected] (R. K€ ogler).
cleation, growth and ion induced destruction were widely described in the literature [1–6]. Recently the IBS of nanometer-sized SiC precipitates in Si was performed for the first time by simultaneous implantation of two beams, one Cþ ion beam and a second Siþ ion beam [7]. The effect of the simultaneous dual beam IBS was compared with the single ion beam IBS and with the sequential dual beam IBS of SiC. It was shown that the additional 1.5 MeV Siþ ion beam stimulates the SiC precipitate formation due to excess vacancy generation in the region where SiC is synthesized. The present study is directed to the investigation of the thermal evolution of such SiC precipitates
0168-583X/03/$ - see front matter 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0168-583X(03)00908-X
990
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
which were formed by single Cþ ion implantation or by dual beam IBS using two beams of Cþ and Siþ ions.
(XTEM) and high resolution electron microscopy (HREM) were performed using a Philips CM300.
2. Experimental
3. Results and discussion
Samples of (1 0 0) CZ-Si were implanted at a temperature of TI ¼ 450 C with 360 keV Cþ ions and with Siþ ions of 1.5 MeV. The mean projected ion ranges are RP ðCþ Þ ¼ 0:75 lm and RP ðSiþ Þ ¼ 1:55 lm. Both implantations were performed under an angle of a ¼ 22:5 simultaneously as well as subsequently using two ion implanters. In the simultaneous implantation mode the ion beams are synchronized in phase and amplitude that the beam spots permanently overlap during beam scanning. The time averaged ion current densities were about jðCþ Þ ¼ 7 1012 cm2 s1 and jðSiþ Þ ¼ 1:5 1012 cm2 s1 , and they were stable in the range of ±15%. SiC is synthesized mainly in the depth range at 0:55 < x < 0:85 lm (C-range) where the highest C concentration of cC P 2 at% is located. The Siþ ions come to rest in much deeper position. The energy deposition into collisions by Siþ ions inside the C-range amounts to 3.8 · 1020 eV/cm2 a value somewhat smaller than the energy deposition of 1.3 · 1021 eV/cm2 calculated for the Cþ ions themselves [7]. The implantation parameters are summarized in Table 1. A thermal treatment at 1150 C for 2 h in Ar was performed for one part of each sample. The C concentration profiles were analyzed by Auger electron spectroscopy (AES). Depth profiling was performed by sputtering with 3 keV Arþ ions of 1 lA/mm2 . The SiC content in the samples was determined by high resolution X-ray diffraction (XRD) using synchrotron radiation (0.154 nm) of ROBL beamline at ESRF Grenoble [8]. Cross-sectional transmission electron microscopy
Under the implantation conditions used a buried layer of nanometer-sized SiC precipitates aligned with the Si matrix is formed already during implantation [1,3,7]. The C depth distribution calculated by the computer code Transport of Ions in Matter (TRIM) [9] is shown in Fig. 1(a) and correlates well with the measured C depth profiles in Fig. 1(b). Fig. 2 shows XTEM images of the samples after annealing. Dislocation loops are observed at a depth of around x ¼ 0:8 lm at the C-range. Numerous dark spots, the SiC precipitates, are visible. The size distribution of these precipitates is bimodal. Smaller precipitates (type I) with a diameter of d ¼ 4–5 nm are hard to see due to the low magnification in Fig. 2 and are mainly located in the C-range close to the dislocations. The depth distribution of the bigger precipitates (type II) with d 10 nm is broader and their concentration is essentially lower. The simultaneously dual implanted sample C + Si is an exception as the type II precipitates are exclusively observed at x P 0:6 lm. Fig. 3 shows crosssectional HREM images of the precipitates. Type I precipitates are 3C–SiC precipitates epitaxially oriented with the Si matrix. Such precipitates were also reported in the literature [1,2,4,6]. In Fig. 3(a) the SiÆ1 1 0æ lattice image is presented of an area where precipitates of type I and II coexist. The type II precipitate is misoriented to the Si matrix. Two images of the same type II precipitate are shown in Fig. 3(b). The bright contrast in the SiÆ1 1 0æ lattice image on the left side is originated by this precipitate. Another orientation shown on
Table 1 Experimental details Sample
Cþ fluence / (cm2 )
Cþ energy E (keV)
Siþ fluence / (cm2 )
Siþ energy E (keV)
Mode
SiC C + Si SiC + Si
8.41 · 1016 8.41 · 1016 8.41 · 1016
360 360 360
– 2.00 · 1016 1.64 · 1016
– 1500 1500
Single Simultaneous Sequential
40
precipitates II
30
C profile
0.0003 0.0002
20
0.0001
10 0
991
C atoms / ion Å
Number of precipitates
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
0.0000 0.2
(a)
0.4 0.6 0.8 depth (µm)
1.0
10
SiC C concentration (at%)
C+Si SiC+Si 5
0 0 (b)
200
400
600
800
1000
1200
depth (nm)
Fig. 1. Number of type II precipitates (bars) observed in sample SiC + Si versus depth compared with the range profile of C (right scale) calculated by TRIM (a). In the range of maximum C concentration it was not possible to count the type II precipitates as they overlap there with type I precipitates and dislocations. The C depth distributions measured by AES for all implantation modes (see Table 1) are given in (b).
the right side of Fig. 3(b) reveals that the precipitate is crystalline and strongly misoriented to the Si matrix. The spacing determinated of the Fourier transform of HREM images is about 0.25 nm a typical value for all the SiC polytypes. The depth distribution of these precipitates in the sample SiC + Si is indicated in Fig. 1(a). Their distribution correlates well with the C depth distribution. Table 2 summarizes the content of SiC and C of the samples. Both, the SiC and the C content of the post-implanted sample C + Si are found to be significantly lower than the content of the two other samples in the as-implanted state and also after annealing. The ‘‘missing C’’ is assumed to
Fig. 2. XTEM micrographs for the implantation modes after annealing (see Table 1). Interstitial-type defects (dislocation loops) are visible at a depth of about 0.8 lm. Dark spots are SiC precipitates. The smaller precipitates (type I) have a diameter of 4–5 nm whereas the larger ones (type II) have a diameter of about 10 nm. One example of the type II precipitates is indicated by an arrow in each image. The left margin of the graph is the sample surface.
be dissolved in the Si matrix at C concentrations below the AES detection limit of cC P 0:5 at%. Because of the low concentration of type II precipitates the main fraction of C atoms is contained in the type I precipitates. Precipitates of type I already grow during Cþ ion implantation [7]. The size of these precipitates is limited at d 5 nm and there is no significant increase of their size during the thermal treatment (Fig. 3(a)). An increase of the C fluence does not change the size of the precipitates, but their density [1,3]. The growth limit of the 3C–SiC precipitates is ascribed to the high interface energy of the
992
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
3C–SiC/c-Si interface [5,10]. Larger randomly oriented precipitates with d 6 20 nm were reported
to grow exclusively in amorphous Si substrate [5]. The interface energy of a 3C–SiC/a-Si interface is seemingly lower. The ballistic destruction of a SiC precipitate under ion irradiation can be described by a decay law according to [4]: N =N0 ¼ expðkUÞ;
k ¼ RC =ðqC dÞ;
ð1Þ
2.5
SiC
SiC+Si
C+Si
2.0
16
2
half size fluence (10 ions / cm )
where N is the number of C atoms in the precipitate and qC is the atomic density of C in SiC. RC is the number of C recoils per incoming ion ejected out of the SiC precipitate. RC can approximately be calculated by TRIM [9] as described in more detail in [4]. The half-size fluence U1=2 ¼ ln 2=k is the ion fluence for which half of the C atoms in a SiC precipitate is recoiled outside. In Fig. 4 calculated values of U1=2 for a SiC precipitate (type I) with d ¼ 4 nm are shown for the different implantation modes. The comparison of U1=2 with
1.5
1.0
0.5 0.3 Fig. 3. HREM images of SiC precipitates. Type I precipitates showing Moire pattern and one precipitate of type II coexisting at x ¼ 0:6 lm in sample SiC (a). The same precipitate of type II is shown in (b) by two images slightly tilted (see text).
0.4
0.5
0.6
0.7
Fig. 4. Half size fluence U1=2 versus depth as calculated for the different implantation modes.
Table 2 SiC and C content (normalized to the reference sample) Sample
Analysis XRD 3C–SiC(1 1 1) peak area
SiC C + Si SiC + Si
0.8
depth (m)
AES C- KLL Auger peak area
As-implanted
Annealed
As-implanted
Annealed
100% 115% 7%
100% 150% 12%
100% 128% –
100% 100% 36%
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
the actual fluence U given in Table 1 leads to the conclusion that the SiC precipitate will be destructed under ion bombardment. The Siþ fluence of the post-implanted sample SiC + Si is U U1=2 . This value is also sufficient to completely decompose the precipitate taking into account that the diameter is just above the estimated critical diameter of dC ¼ 1–4 nm [5]. Notice that in the case of the single and simultaneous double implantation not the total ion fluence contributes to precipitate dissolution as the precipitates must previously be formed by C accumulation before they are destructed. However, as SiC precipitates decay under irradiation it is evident that the processes of precipitate nucleation and growth must simultaneously proceed. The type II precipitates were formed during thermal treatment as they were not observed in the as-implanted state [7]. Their higher growth velocity, the misorientation in regard to the Si matrix and their distribution (Fig. 2) can be explained by the assumption that these precipitates form in amorphous surrounding which reduces their interface energy. In our samples an amorphous phase was not seen. However, a careful TEM analysis in as-implanted samples revealed few nanometer-sized disturbed areas (not shown). The formation of amorphous areas can be stimulated by the high amount of dissolved C atoms enhancing the damage accumulation [11,12]. This happens especially for sample SiC + Si. The Cinduced amorphization of Si was reported for even higher C concentrations above 17 at% [3]. Fig. 2 shows that the highest density of type II precipitates just appears for the sample SiC + Si. On the other hand, amorphous areas are expected to crystallize by ion beam induced epitaxial crystallization (IBIEC) [13]. The temperature TI ¼ 450 C is well above the reverse temperature TR 6 150 C characterizing the onset of the IBIEC process [14]. The strength of the IBIEC process is given by the velocity v of a moving (flat) a-Si/c-Si interface. It can be calculated by (Eq. (2)) [15] v ðZjÞ
3=4
expð0:3 eV=kT Þ;
ð2Þ
where Z is the number of atomic displacements created by one ion per unit depth and kT has its usual meaning. The crystallization rate v results to
993
values of 43% for the SiC + Si implantation mode and to 124% for the C + Si mode as compared with SiC (100%). It means that the IBIEC effect is stronger for the simultaneous dual implant C + Si and relatively weak for the Siþ post-implantation SiC + Si. In summary, both types of precipitates consist of crystalline SiC, but they are different in size, concentration and orientation. The simultaneous dual beam implantation stimulates the formation of epitaxially oriented type I precipitates and avoids the amorphization of the substrate. Therefore, the growth of type II precipitates during annealing is suppressed.
References [1] A. Romano-Rodriguez, C. Serre, L. Calvo-Barrio, A. Perez-Rodriguez, J.R. Morante, R. K€ ogler, W. Skorupa, Mater. Sci. Eng. B 36 (1996) 282. [2] F. Eichhorn, N. Schell, W. Matz, R. K€ ogler, J. Appl. Phys. 86 (1999) 4184. [3] J.K.N. Lindner, B. Stritzker, Nucl. Instr. and Meth. B 147 (1999) 249. [4] J.K.N. Lindner, B. Stritzker, Nucl. Instr. and Meth. B 148 (1999) 528. [5] J.K.N. Lindner, Nucl. Instr. and Meth. B 178 (2001) 44. [6] F. Eichhorn, N. Schell, A. M€ ucklich, H. Metzger, W. Matz, R. K€ ogler, J. Appl. Phys. 91 (2002) 1287. [7] R. K€ ogler, F. Eichhorn, J.R. Kaschny, A. M€ ucklich, H. Reuther, V. Heera, W. Skorupa, C. Serre, A. PerezRodriguez, Appl. Phys. A 76 (2003) 827. [8] W. Matz, N. Schell, G. Bernhard, F. Prokert, T. Reich, J. Claußner, W. Oehme, R. Schlenk, S. Dienel, H. Funke, F. Eichhorn, M. Betzl, D. Pr€ ohl, U. Strauch, G. H€ uttig, H. Krug, W. Neumann, V. Brendler, P. Reichel, M.A. Denecke, H. Nitsche, J. Synchrotron Rad. 6 (1999) 1076– 1085. [9] W.J. Taylor, T.Y. Tan, U. G€ osele, Appl. Phys. Lett. 62 (1993) 3336. [10] J.F. Ziegler, J.B. Biersack, U. Littmark, The Stopping and Ranges of Ions in Solids, Pergamon, NY, 1985. [11] J.P. de Souza, H. Boudinov, P.F.P. Fichtner, Appl. Phys. Lett. 64 (1994) 3596. [12] V. Heera, W. Skorupa, B. Pecz, L. Dobos, Appl. Phys. Lett. 76 (2000) 2847. [13] T. Henkel, V. Heera, R. K€ ogler, W. Skorupa, J. Appl. Phys. 84 (1998) 3090. [14] R. K€ ogler, V. Heera, W. Skorupa, E. Glase, T. Bachmann, D. R€ uck, Nucl. Instr. and Meth. B 80–81 (1993) 556. [15] V. Heera, T. Henkel, R. K€ ogler, W. Skorupa, Phys. Rev. B 52 (1995) 15776.
SiC precipitates formed in Si by simultaneous dual beam implantation of Cþ and Siþ ions €gler R. Ko a
a,*
, F. Eichhorn a, A. M€ ucklich a, H. Reuther a, V. Heera a, W. Skorupa a, J.K.N. Lindner b
Forschungszentrum Rossendorf, Institut f€ur Ionenstrahlphysik und Materialforschung, P.O. Box 510119, 01314 Dresden, Germany b Universit€at Augsburg, Institut f€ur Physik, 86135 Augsburg, Germany
Abstract Nanometer-sized SiC precipitates were synthesized at 450 C in Si by simultaneous dual beam implantation of Cþ and Siþ ions and subsequent annealing. The results are compared with those of sequential dual beam implantation and of single beam implantation. Two types of SiC precipitates were found. Precipitates of type I with a diameter of d ¼ 4–5 nm consist of 3C-SiC epitaxially oriented with the Si matrix. They were formed already in the as-implanted state and do not grow further during subsequent annealing. The SiC precipitates of type II with d 10 nm are not oriented with the Si matrix and grow exclusively during the subsequent annealing. The high growth velocity, the misorientation in regard to the Si matrix and the lower concentration of type II precipitates can be explained by the assumption that these precipitates were formed in an amorphous substrate which modifies their interface energy. 2003 Elsevier Science B.V. All rights reserved. PACS: 61.72.T; 81.10.Jt Keywords: Silicon; SiC; Ion beam synthesis; Dual beam implantation
1. Introduction The ion beam synthesis (IBS) of SiC in Si has been investigated for many years because of the outstanding physical material properties of SiC making it applicable for devices working under extreme conditions. The basic processes of the formation and evolution of SiC precipitates during IBS and subsequent thermal treatment, their nu-
*
Corresponding author. Tel.: +49-351-260-3613; fax: +49351-260-3411. E-mail address: [email protected] (R. K€ ogler).
cleation, growth and ion induced destruction were widely described in the literature [1–6]. Recently the IBS of nanometer-sized SiC precipitates in Si was performed for the first time by simultaneous implantation of two beams, one Cþ ion beam and a second Siþ ion beam [7]. The effect of the simultaneous dual beam IBS was compared with the single ion beam IBS and with the sequential dual beam IBS of SiC. It was shown that the additional 1.5 MeV Siþ ion beam stimulates the SiC precipitate formation due to excess vacancy generation in the region where SiC is synthesized. The present study is directed to the investigation of the thermal evolution of such SiC precipitates
0168-583X/03/$ - see front matter 2003 Elsevier Science B.V. All rights reserved. doi:10.1016/S0168-583X(03)00908-X
990
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
which were formed by single Cþ ion implantation or by dual beam IBS using two beams of Cþ and Siþ ions.
(XTEM) and high resolution electron microscopy (HREM) were performed using a Philips CM300.
2. Experimental
3. Results and discussion
Samples of (1 0 0) CZ-Si were implanted at a temperature of TI ¼ 450 C with 360 keV Cþ ions and with Siþ ions of 1.5 MeV. The mean projected ion ranges are RP ðCþ Þ ¼ 0:75 lm and RP ðSiþ Þ ¼ 1:55 lm. Both implantations were performed under an angle of a ¼ 22:5 simultaneously as well as subsequently using two ion implanters. In the simultaneous implantation mode the ion beams are synchronized in phase and amplitude that the beam spots permanently overlap during beam scanning. The time averaged ion current densities were about jðCþ Þ ¼ 7 1012 cm2 s1 and jðSiþ Þ ¼ 1:5 1012 cm2 s1 , and they were stable in the range of ±15%. SiC is synthesized mainly in the depth range at 0:55 < x < 0:85 lm (C-range) where the highest C concentration of cC P 2 at% is located. The Siþ ions come to rest in much deeper position. The energy deposition into collisions by Siþ ions inside the C-range amounts to 3.8 · 1020 eV/cm2 a value somewhat smaller than the energy deposition of 1.3 · 1021 eV/cm2 calculated for the Cþ ions themselves [7]. The implantation parameters are summarized in Table 1. A thermal treatment at 1150 C for 2 h in Ar was performed for one part of each sample. The C concentration profiles were analyzed by Auger electron spectroscopy (AES). Depth profiling was performed by sputtering with 3 keV Arþ ions of 1 lA/mm2 . The SiC content in the samples was determined by high resolution X-ray diffraction (XRD) using synchrotron radiation (0.154 nm) of ROBL beamline at ESRF Grenoble [8]. Cross-sectional transmission electron microscopy
Under the implantation conditions used a buried layer of nanometer-sized SiC precipitates aligned with the Si matrix is formed already during implantation [1,3,7]. The C depth distribution calculated by the computer code Transport of Ions in Matter (TRIM) [9] is shown in Fig. 1(a) and correlates well with the measured C depth profiles in Fig. 1(b). Fig. 2 shows XTEM images of the samples after annealing. Dislocation loops are observed at a depth of around x ¼ 0:8 lm at the C-range. Numerous dark spots, the SiC precipitates, are visible. The size distribution of these precipitates is bimodal. Smaller precipitates (type I) with a diameter of d ¼ 4–5 nm are hard to see due to the low magnification in Fig. 2 and are mainly located in the C-range close to the dislocations. The depth distribution of the bigger precipitates (type II) with d 10 nm is broader and their concentration is essentially lower. The simultaneously dual implanted sample C + Si is an exception as the type II precipitates are exclusively observed at x P 0:6 lm. Fig. 3 shows crosssectional HREM images of the precipitates. Type I precipitates are 3C–SiC precipitates epitaxially oriented with the Si matrix. Such precipitates were also reported in the literature [1,2,4,6]. In Fig. 3(a) the SiÆ1 1 0æ lattice image is presented of an area where precipitates of type I and II coexist. The type II precipitate is misoriented to the Si matrix. Two images of the same type II precipitate are shown in Fig. 3(b). The bright contrast in the SiÆ1 1 0æ lattice image on the left side is originated by this precipitate. Another orientation shown on
Table 1 Experimental details Sample
Cþ fluence / (cm2 )
Cþ energy E (keV)
Siþ fluence / (cm2 )
Siþ energy E (keV)
Mode
SiC C + Si SiC + Si
8.41 · 1016 8.41 · 1016 8.41 · 1016
360 360 360
– 2.00 · 1016 1.64 · 1016
– 1500 1500
Single Simultaneous Sequential
40
precipitates II
30
C profile
0.0003 0.0002
20
0.0001
10 0
991
C atoms / ion Å
Number of precipitates
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
0.0000 0.2
(a)
0.4 0.6 0.8 depth (µm)
1.0
10
SiC C concentration (at%)
C+Si SiC+Si 5
0 0 (b)
200
400
600
800
1000
1200
depth (nm)
Fig. 1. Number of type II precipitates (bars) observed in sample SiC + Si versus depth compared with the range profile of C (right scale) calculated by TRIM (a). In the range of maximum C concentration it was not possible to count the type II precipitates as they overlap there with type I precipitates and dislocations. The C depth distributions measured by AES for all implantation modes (see Table 1) are given in (b).
the right side of Fig. 3(b) reveals that the precipitate is crystalline and strongly misoriented to the Si matrix. The spacing determinated of the Fourier transform of HREM images is about 0.25 nm a typical value for all the SiC polytypes. The depth distribution of these precipitates in the sample SiC + Si is indicated in Fig. 1(a). Their distribution correlates well with the C depth distribution. Table 2 summarizes the content of SiC and C of the samples. Both, the SiC and the C content of the post-implanted sample C + Si are found to be significantly lower than the content of the two other samples in the as-implanted state and also after annealing. The ‘‘missing C’’ is assumed to
Fig. 2. XTEM micrographs for the implantation modes after annealing (see Table 1). Interstitial-type defects (dislocation loops) are visible at a depth of about 0.8 lm. Dark spots are SiC precipitates. The smaller precipitates (type I) have a diameter of 4–5 nm whereas the larger ones (type II) have a diameter of about 10 nm. One example of the type II precipitates is indicated by an arrow in each image. The left margin of the graph is the sample surface.
be dissolved in the Si matrix at C concentrations below the AES detection limit of cC P 0:5 at%. Because of the low concentration of type II precipitates the main fraction of C atoms is contained in the type I precipitates. Precipitates of type I already grow during Cþ ion implantation [7]. The size of these precipitates is limited at d 5 nm and there is no significant increase of their size during the thermal treatment (Fig. 3(a)). An increase of the C fluence does not change the size of the precipitates, but their density [1,3]. The growth limit of the 3C–SiC precipitates is ascribed to the high interface energy of the
992
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
3C–SiC/c-Si interface [5,10]. Larger randomly oriented precipitates with d 6 20 nm were reported
to grow exclusively in amorphous Si substrate [5]. The interface energy of a 3C–SiC/a-Si interface is seemingly lower. The ballistic destruction of a SiC precipitate under ion irradiation can be described by a decay law according to [4]: N =N0 ¼ expðkUÞ;
k ¼ RC =ðqC dÞ;
ð1Þ
2.5
SiC
SiC+Si
C+Si
2.0
16
2
half size fluence (10 ions / cm )
where N is the number of C atoms in the precipitate and qC is the atomic density of C in SiC. RC is the number of C recoils per incoming ion ejected out of the SiC precipitate. RC can approximately be calculated by TRIM [9] as described in more detail in [4]. The half-size fluence U1=2 ¼ ln 2=k is the ion fluence for which half of the C atoms in a SiC precipitate is recoiled outside. In Fig. 4 calculated values of U1=2 for a SiC precipitate (type I) with d ¼ 4 nm are shown for the different implantation modes. The comparison of U1=2 with
1.5
1.0
0.5 0.3 Fig. 3. HREM images of SiC precipitates. Type I precipitates showing Moire pattern and one precipitate of type II coexisting at x ¼ 0:6 lm in sample SiC (a). The same precipitate of type II is shown in (b) by two images slightly tilted (see text).
0.4
0.5
0.6
0.7
Fig. 4. Half size fluence U1=2 versus depth as calculated for the different implantation modes.
Table 2 SiC and C content (normalized to the reference sample) Sample
Analysis XRD 3C–SiC(1 1 1) peak area
SiC C + Si SiC + Si
0.8
depth (m)
AES C- KLL Auger peak area
As-implanted
Annealed
As-implanted
Annealed
100% 115% 7%
100% 150% 12%
100% 128% –
100% 100% 36%
R. K€ogler et al. / Nucl. Instr. and Meth. in Phys. Res. B 206 (2003) 989–993
the actual fluence U given in Table 1 leads to the conclusion that the SiC precipitate will be destructed under ion bombardment. The Siþ fluence of the post-implanted sample SiC + Si is U U1=2 . This value is also sufficient to completely decompose the precipitate taking into account that the diameter is just above the estimated critical diameter of dC ¼ 1–4 nm [5]. Notice that in the case of the single and simultaneous double implantation not the total ion fluence contributes to precipitate dissolution as the precipitates must previously be formed by C accumulation before they are destructed. However, as SiC precipitates decay under irradiation it is evident that the processes of precipitate nucleation and growth must simultaneously proceed. The type II precipitates were formed during thermal treatment as they were not observed in the as-implanted state [7]. Their higher growth velocity, the misorientation in regard to the Si matrix and their distribution (Fig. 2) can be explained by the assumption that these precipitates form in amorphous surrounding which reduces their interface energy. In our samples an amorphous phase was not seen. However, a careful TEM analysis in as-implanted samples revealed few nanometer-sized disturbed areas (not shown). The formation of amorphous areas can be stimulated by the high amount of dissolved C atoms enhancing the damage accumulation [11,12]. This happens especially for sample SiC + Si. The Cinduced amorphization of Si was reported for even higher C concentrations above 17 at% [3]. Fig. 2 shows that the highest density of type II precipitates just appears for the sample SiC + Si. On the other hand, amorphous areas are expected to crystallize by ion beam induced epitaxial crystallization (IBIEC) [13]. The temperature TI ¼ 450 C is well above the reverse temperature TR 6 150 C characterizing the onset of the IBIEC process [14]. The strength of the IBIEC process is given by the velocity v of a moving (flat) a-Si/c-Si interface. It can be calculated by (Eq. (2)) [15] v ðZjÞ
3=4
expð0:3 eV=kT Þ;
ð2Þ
where Z is the number of atomic displacements created by one ion per unit depth and kT has its usual meaning. The crystallization rate v results to
993
values of 43% for the SiC + Si implantation mode and to 124% for the C + Si mode as compared with SiC (100%). It means that the IBIEC effect is stronger for the simultaneous dual implant C + Si and relatively weak for the Siþ post-implantation SiC + Si. In summary, both types of precipitates consist of crystalline SiC, but they are different in size, concentration and orientation. The simultaneous dual beam implantation stimulates the formation of epitaxially oriented type I precipitates and avoids the amorphization of the substrate. Therefore, the growth of type II precipitates during annealing is suppressed.
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