- Email: [email protected]
Point defect diffusion and clustering in ion implanted c-Si
Nuclear Instruments and Methods in Physics Research B 178 (2001) 25±32
www.elsevier.nl/locate/nimb
Point defect diusion and clustering in ion implanted c-Si Sebania Libertino *, Salvatore Coa, Corrado Spinella, Antonino La Magna, Vittorio Privitera CNR-IMETEM, Stradale Primosole 50, I-95121 Catania, Italy
Abstract This paper reviews some fundamental aspects of point defect migration and agglomeration in crystalline Si. Both insitu and ex-situ measurements were used to reach this target. Room temperature (RT) diusivities of 1:5 10 15 and 3:0 10 13 cm2 =s for I and V, respectively, were obtained using in-situ leakage current measurements, performed during and just after ion implantation. To follow the defect evolution and clustering upon annealing, ex-situ optical and electrical measurements were used. Low temperature (300±500°C) annealing causes the formation of point-like defects, while higher temperatures (500±800°C) are necessary to have defect clustering. Finally, a well-de®ned dose 1 1013 Si=cm2 in pure Si) temperature (650°C) and time thresholds exist for the transition from I-clusters to extended {3 1 1} defects. When the transition takes place, both the optical and electrical defect properties undergo a dramatic change, suggesting an abrupt structural transition in the evolution from I-cluster to {3 1 1} defects. Kinetic lattice Monte-Carlo simulations used to model the defect agglomeration and growth con®rm these results. Ó 2001 Elsevier Science B.V. All rights reserved. Keywords: Defects; Interstitial clusters; Silicon; Migration; Diusivity; Photoluminescence
1. Introduction Defect formation and agglomeration in crystalline Si (c-Si) have been the subject of many studies in the last 30 years [1±3]. A strong driving force to the research is provided by technological interest since Si is the most used semiconductor in electronic devices fabrication. The method of choice to dope Si is ion implantation that causes the generation of point defects: self-interstitials (I) and vacancies (V). Due to their low migration energy, 1 eV [4], I and
*
Corresponding author. Tel.: +39-95-591212; fax: +39-957139154. E-mail address: [email protected] (S. Libertino).
V can freely migrate even at room temperature (RT). We showed that, after MeV He implantation in Si only a small percentage of the generated defects succeeds in escaping direct recombination according to the relation I V ;. The ®nal defect concentration strongly depends on the Si impurity content, varying from 16% (high content) to below 3% (low content) [5]. The I and V that escape recombination interact with impurities (C and O), dopant atoms (e.g. P or B) or among themselves and form RT stable defects [1,2]. V-type defects like the divacancy (VV), oxygen-vacancy (OV) and phosphorous-vacancy (PV) and I-type defects like carbon interstitial-oxygen Ci O, phosphorous± carbon interstitial PCi , boron interstitial-oxygen Bi O, boron interstitial-carbon Bi C are left in the
0168-583X/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 0 ) 0 0 5 0 4 - 8
26
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
sample. It should be mentioned that I-type defect formation is a two-step mechanism: I interact with C and B forming interstitial-carbon Ci or interstitial-boron Bi and these species migrate to form the C- or B-related defects previously listed. All these defects introduce deep levels in the Si band gap [1] and can be detected with electrical measurements such as deep-level transient spectroscopy (DLTS) [5] or spreading resistance pro®ling (SRP) [6]. Despite eorts, up to now the I and V RT diffusivity values are not known. During annealing, RT stable complexes dissociate releasing the I and V they stored. Most of them will recombine until only I remain in the sample. In fact, due to the extra-implanted ion the number of I generated per implanted ion is higher (+1) than the number of V. These extra I agglomerate in more complex structures and, eventually, extended defects (e.g. {3 1 1} defects) form. They dissociate upon further annealing causing a large I supersaturation which drives several phenomena such as B transient enhanced diusion (TED) [7]. Since the TED enhancement and duration depend on the thermal stability of defects storing the I, it is fundamental to characterize them. While there are detailed experimental and theoretical studies on extended defects, e.g. {3 1 1} [7], little is known on the nanometer-sized I-clusters which will eventually grow into {3 1 1} defects. Recently, Cowern et al. [8] studied the TED at various annealing temperatures and extracted the I supersaturation values as a function of time. Assuming an Ostwald ripening mechanism for I agglomeration, they found that the experimental data can be ®tted only assuming that two stable Icluster sizes of roughly 4 and 8 I exist, but nothing can be said on their structure. On the other hand, tight-binding (TB) calculations [9] showed that the basic {3 1 1} structure, an in®nitely extended (1 1 0) I chain, cannot be the lowest energy con®guration for small clusters. The formation energy per I, very low if the chain is in®nite (1.7 eV), becomes very large (4.7 eV in a di-interstitial) when the number n of I in the chain is low (<10). Recent TB calculations of I agglomeration [10] suggest that alternative structures storing few I n < 10 could be a more favorable pathway for I agglomeration and growth. This implies that small I-clusters are not
the direct precursors of {3 1 1} defects and that a structural transformation has to occur during the growth process. Kinetic lattice Monte-Carlo simulations used to model defect agglomeration and {3 1 1} formation [11] con®rm this scenario. The aim of this paper is to assess the basic I and V properties, follow the I agglomeration into clusters and understand the growth mechanism from I-clusters to {3 1 1} defects. 2. Experimental We used both n- and p-type Czochralski grown (CZ, 7 1017 O cm 3 , 1 1016 C cm 3 ) and epitaxial (epi, O and C concentrations 5 1015 cm 3 ) Si with resistivities in the range 0.7± 4 X cm. Schotthy barriers and p ±n junctions were used for DLTS and leakage current IL measurements. The samples were implanted with He at an energy of 2.5 MeV to doses of 1 109 cm 2 to 1 1012 cm 2 and dose rates of 1 108 cm 2 to 1 1010 cm 2 s 1 or with Si at energies of 0.04±1.2 MeV to doses of 1 1012 cm 2 to 5 1013 cm 2 . DLTS measurements were performed using 1 ms ®lling pulse and reverse biases such as to include the entire damage distribution within the depletion layer. Defect depth pro®les were obtained changing the ®lling pulse voltage. Photoluminescence (PL) and SR measurements were performed on non-structured samples. PL measurements were carried out at 17 K using the 488 nm line of an Argon laser at a pump power of 50 mW. Light emitted from the sample was dispersed by a monochromator and detected by a liquid nitrogen cooled Ge detector. Standard lock in technique was used to improve the signal-to-noise ratio. Sample morphology was analyzed by transmission electron microscopy (TEM) in both plan view and cross-sectional con®gurations. 3. Results and discussion 3.1. Point defect RT migration Precious information on I and V migration in c-Si can be obtained by SRP. Defects are generated
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
at the surface by low-energy ion implantation and their migration monitored through the dopant deactivation tail [6]. We performed a shallow Si implant, 40 keV, in n-type epi Si to produce a damaged region which extends from the surface to 0.4 lm. Some of the V and I generated succeed in escaping recombination or clustering in this region and are injected into the bulk. The probability of recombination and/or clustering in the bulk is small, due to the very low defect concentration in this region, hence, they undergo long-range migration until trapping occurs. In highly pure epi Si, migration up to 2 lm is reported [6]. The SRP measurements plotted in Fig. 1 refer to a sample as implanted with 5 1012 Si=cm2 . A tilt angle of 7° during implantation was used to avoid any channeling eect that could aect the measurement. The vertical solid line in the ®gure shows the region directly damaged by the beam. Within this region, a strong decrease in the free carrier concentration, due to the damage presence, is observed. A deactivation tail extending up to 0.9 lm is observed: deeper, the substrate-free carrier concentration 7 1015 cm 3 is measured. After 150°C for 30 min annealing, the deactivation tail is fully recovered H, suggesting that the defects
Fig. 1. SRP measurements on n-type epi Si implanted with 40 keV Si to a dose of 5 1012 cm 2 and a 7° tilt angle: as-implanted ; after annealing at 150°C for 30 min H. The vertical solid line indicates the region directly damaged by the beam.
27
formed are point-like. Part of the damage in the ®rst 0:4 lm recovers too, showing that those defects are also present in the region directly interested by the beam. The deactivation tail is visible only in epi Si: high contents of C or O trap the I and V as soon as they are formed preventing their migration. SRP is not a spectroscopic technique; hence it cannot identify the origin of the deactivation tail. The spectroscopy of the defects in the tail was studied by measuring the depth pro®les of the VV , PV H and PCi j complexes for the same sample, see Fig. 2. These defects are the only ones that can aect the SRP measurement since they lie close to the center of the gap. The VV and PV peaks lie at the same temperature position in DLTS: hence, when both defects are present, the two DLTS signals overlap [1]. The activation energy of the resulting level is that of the dominating one, usually the VV . We separated the two contributions by measuring the peak activation energy for each point plotted in the ®gure. In the region where the VV is present (0.4±0.6 lm), the PV presence was monitored by measuring the VV pro®le in the as-implanted state and after 150°C 30 min annealing. Since the PV complex anneals at 150°C [1], the dierence between the two VV depth pro®les provides the PV concentration in
Fig. 2. DLTS depth pro®les of VV (), PV H and PCi j defect complexes for an n-type Si sample as-implanted with 40 keV Si to a dose of 5 1012 cm 2 and a 7° tilt angle. The vertical solid line shows the region directly interested by the beam.
28
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
that region. The same procedure was used for the PV points lying between 0.6 and 0:9 lm: the activation energy in each point of the depth pro®le and the defect annealing kinetics were monitored. While the VV concentration is very high in the directly damaged region and decreases abruptly going deeper into the sample, the PV behavior is quite dierent. It is still detectable up to 0:9 lm con®rming that V migrate into the bulk until trapping at impurities occurs. On the other hand, the PCi presence shows that I migrate even to longer distances (1.3 lm), probably due to the two-step mechanism of I-type defect formation. The deactivation tail is given by both P atoms enclosed in the complexes, that become inactive, and P atoms compensated by the complexes deep levels. When the defects break at 150°C, the deactivation tail recovers since the P in complexes returns electrically active and the levels are removed from the gap (see Fig. 1). Both the PV and the PCi dissociate at this temperature. Ex-situ measurements cannot provide absolute values for I and V diusivities since a time delay exists between implantation and measure, hence in-situ measurements are necessary. We monitored in-situ the IL evolution during and just after implantation of p ±n junctions with 2.5 MeV He. A schematic of the measurement setup is shown in the inset of Fig. 3: the sample was reverse biased during and after implantation to a voltage of )30 V. The reverse bias was chosen in order to embody the entire defect pro®le in the depletion region. The results of an implant with a dose rate of 1 108 cm 2 s 1 are shown in Fig. 3. At the beam turn-on, IL (30 pA at )30 V in the unimplanted diode) increases monotonically as a result of the formation of defect complexes, hence of deep levels in the Si band gap. At the beam turn-o, when a total dose of 1 1012 cm 2 is reached (after 104 s), IL decreases suggesting that a partial annihilation of the defects responsible for carrier generation in the depletion layer is occurring. After 1 day, IL is reduced by a factor of 2 and the transient almost saturates, with small variations lasting for several days. To identify the defects responsible of IL , DLTS measurements were performed. Since IL strongly depends on levels located close to mid gap, the main contribution arises
Fig. 3. In-situ IL measurements during and after implantation with 2.5 MeV He to a dose of 1 1012 cm 2 and a dose rate of 1 108 cm 2 s 1 . The vertical dashed line indicates the beam turn-o. The solid line is a ®t of the data. In the inset schematic of the experimental setup is plotted.
from VV and PV. DLTS and IL data were compared assuming a linear dependence of IL on the defect content. During implantation, free V agglomerate into V-type defects, producing the IL increase. After implantation, a residual concentration of free I still exists and can migrate to Vtype defects dissociating them. The IL time evolution is directly linked to I and V migrations. The data were ®tted (solid line) solving a set of coupled dierential equations describing diusion-limited defect annihilation, agglomeration and trapping and having the I and V diusivity as the only parameters [12]. Values of 1:5 10 15 cm2 =s and 3 10 13 cm2 =s, respectively, were found. These low values suggest that, when monitoring the defect migration in the depletion layer the enhancement due to ionization eects is limited and values close to those expected for neutral V and I are achieved. Since V-type defects might be annihilated by the mobile Ci , we checked the C eect assuming a Ci diusivity of 10 15 cm2 =s. However, simulations performed with 5 1015 CS cm 3 (the C content in our samples) or without C give the same results. Finally, so far we assumed that the electric ®eld presence, typical of these measurements, does not aect the ®nal defect distribution. This is not fully correct since recent measurements have shown that
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
the ®nal defect pro®le is slightly dierent in concentration and spatial distribution if an electric ®eld is applied. It is shown in Fig. 4 where the VV depth pro®les of samples implanted to 1010 He cm 2 and a dose rate of 108 cm 2 s 1 reverse biased or unbiased M during implantation are compared. The unbiased sample shows a higher defect concentration and a slightly dierent depth position. The same results hold for all the measured samples and also when the OV is monitored. We believe it occurs because, under our experimental condition, there is a fraction of electrically charged defects. These defects escape from the depletion region, due to the electric ®eld presence, thus avoiding the complexes formation and causing a shift of the ®nal defect depth pro®le. Further measurements are in progress to understand and quantify this phenomenon. 3.2. Defect evolution: clustering and extended defect formation Annealing at temperatures P 350°C dissociates RT stable defects causing the release of the I and V they stored. Most of them undergo a direct recombination and the survivors are, again, stored in complexes. These ``new'' defect characteristics strongly depend on the annealing temperature and
Fig. 4. VV depth pro®le of an n-type Si sample implanted with 2.5 MeV He to a dose of 1 1010 cm 2 and a dose rate of 1 108 cm 2 s 1 reverse biased to 30 V or unbiased (M) during implantation.
29
dier from those of RT stable defects. For annealing at temperatures 6 500°C, I-type defects, storing just one or two I, remain in the sample [13]. On the other hand, samples implanted with 1.2 MeV Si to a dose of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line) and annealed at 600°C for 4 h exhibit the PL spectra shown in Fig. 5. Although the PL intensity increases with ion dose, the major features remain unchanged. They consist of two broad peaks centered at 1300 and 1400 nm, associated to I-clusters [15]. Broad peaks in PL spectra are usually observed in samples with extended defects. They were associated to the quantum con®nement of carriers in regions with high strain surrounding the defects. Since no extended defects are detected in our samples, we associated this feature to the carrier recombination in the strained region surrounding the I-clusters embedded in the Si matrix [15]. Several sharp lines, associated to point-like defect complexes, are superimposed to the I-clusters signatures in the range 1200±1280 nm. Among them, the W0 line at 1233 nm (1.0048 eV) is present. It is a perturbed form of the W line (associated to I-rich defects) observed in regions with high strain [14], con®rming the strain presence in our samples. The spectra also show the oxygen thermal donors (at 1620 nm, 0.765 eV) and the Si band-edge phonon assisted recombination line at 1121 nm (1.1056 eV). The PL spectra of I-clusters are totally dierent from those of point-like defects suggesting dierent defect structures in the two cases. Moreover, the PL peak broadening of the clusters spectra suggests a complex structure for these defects causing a measurable strain in the lattice, as con®rmed by the TEM cross-section shown in the inset of Fig. 5. No extended defects are visible even at the highest implanted dose. Only a weak contrast at the projected range is observed, demonstrating that a heavily damaged region probably consisting of small defect aggre is present in the sample. Annealing gates (<50 A) for times up to 15 h at 600°C, only produces a decrease in the PL signal intensity suggesting that the I-clusters annealate and no extended defect form. These results suggest that there is a threshold temperature for the {3 1 1} defect formation. At T 6 600°C the I are stored into clusters (for
30
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
Fig. 5. PL spectra on n-type Si samples implanted with 1.2 MeV Si to doses of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line). The samples were annealed at 600°C for 4 h. In the inset: TEM cross-section made in weak beam g 4g using the (1 1 1) re¯ection of the sample implanted at 5 1013 Si cm 2 .
doses P 1012 Si cm 2 ), eventually released and annealed, probably at the surface. Annealing at 680°C, on the same samples shown in Fig. 5, produces major modi®cations in the defect optical and structural properties. The PL spectra of the same samples annealed for 1 h are plotted in Fig. 6. The sample implanted at 1012 Si cm 2 (dotted line) shows only the Si bandedge line, while at doses P 1013 cm 2 a sharp peak at 1376 nm (0.9007 eV) dominates the spectra. The study of the PL spectra and the TEM analysis of damaged samples allowed us to associate it to optical transitions occurring at or close to {3 1 1} extended defects [15]. Its width slightly increases with the dose and a broad band centered at 1550 nm (0.799 eV) is developed after 5 1013 cm 2 implantation. These measurements clearly indicate that to have the transition from I-clusters to {3 1 1} extended defects the implantation dose must be above a threshold that, for our CZ Si samples, is 1 1013 cm 2 . Previous studies [7] showed that the minimum dose necessary to form {3 1 1} defects increases with the sample C content, since it is an ecient trap for I and prevents their agglomeration. Hence, in samples with low C content, as our
CZ Si, the implantation dose threshold should be lower. Finally, we ®xed the implantation dose at 1 1013 Si=cm2 and the temperature at 700°C, both above thresholds for {3 1 1} formation, and monitored the defect evolution as a function of the
Fig. 6. PL spectra on n-type Si implanted with 1.2 MeV Si to doses of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line), and annealed at 680°C for 1 h.
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
annealing time. The resulting PL spectra are compared in Fig. 7. Already after 10 min (dotted line) the I-clusters are formed, suggesting a fast agglomeration process. Annealing up to 30 min causes only a decrease in the I-clusters signal intensity, probably due to their partial dissolution. The Si band-edge line, barely visible after 10 min becomes stronger as the time proceeds, con®rming that the crystal damage is recovering. After 1h annealing (solid line), the {3 1 1} characteristic signature appears, even if a strong contribution of the I-clusters signatures is visible. Only after 2 h (dash±double dot-dashed line) the {3 1 1} signal is the only one detectable in the spectrum. The results so far reported, together with our studies using DLTS [16], indicate that major transformations in the optical, electrical and structural properties of Si occur to allow the {3 1 1} defect formation. At the early stages of nucleation small I-clusters form. We believe that their structures signi®cantly dier from of the {3 1 1} defects. As annealing proceeds, some clusters grow bigger, but there is a critical dimension at which the I-clusters need to rearrange themselves in a dierent structure, probably close to the (1 1 0) I chain that constitutes the {3 1 1} defects. Of course, a nucleation barrier must exist in order to have the transition. After annealing at low
31
temperature, or for short time, the probability to overcome this barrier is small and most of the I remain in clusters until they anneal out. On the other hand, increasing the temperature, and/or the time, for doses above threshold, the probability of transition becomes very large and {3 1 1} defects form. Our results are in agreement with those by Cowern et al. [8]. He assumes that a very stable Icluster con®guration exists and that a nucleation barrier has to be overcome to form {3 1 1} defects. Kinetic lattice Monte-Carlo simulations reported in this same number [11] con®rm this hypothesis. Further studies are in progress to determine the critical cluster size. 4. Conclusion We have shown that I and V RT migration can be monitored using ex-situ and in-situ techniques, since their migration give the ®nal defect distribution. Upon annealing, I-clusters and extended defects form. By PL measurements we followed the transitions since the defect optical properties abruptly change. In particular, we found that for MeV Si implantation a threshold dose 1013 cm 2 and temperature (650°C) exist for {3 1 1} forma-
Fig. 7. PL spectra on p-type Si implanted with 1.2 MeV Si to a dose of 1 1013 cm 2 and annealed at 700°C for 10 min (dotted line), 20 min (dash±dot-dashed line), 30 min (dashed line), 1 h (solid line) and 2 h (dash-double dot-dashed line).
32
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
tion. Once the dose 1013 cm 2 and temperature (700°C) are above threshold a certain time (1 h) is necessary to allow the {3 1 1} defect formation. We believe that a structural transformation accompanies the transition from I-cluster to {3 1 1} defects. Acknowledgements The authors acknowledge A. Spada and N. Parasole for the invaluable technical assistance, A. Marino for the implants and S. Pannitteri for the TEM samples preparation. This work was partially supported by Progetto Finalizzato MADESS II and Progetto 5% Microelettronica. References [1] L.C. Kimerling, in: N.B. Urli, J.M. Corbett (Ed.), Radiation Eects in Semiconductors, Inst. Phys. Conf. Ser. 31, London, 1977, p. 221. [2] G.D. Watkins, Mater. Res. Soc. Symp. Proc. 469 (1997) 139. [3] B.G. Svensson, C. Jagadish, A. Hallen, J. Lalita, Nucl. Instr. and Meth. B 106 (1995) 183. [4] P.M. Fahey, P.B. Grin, J.D. Plummer, Rev. Mod. Phys. 61 (1989) 289.
[5] S. Coa, V. Privitera, F. Priolo, S. Libertino, G. Mannino, J. Appl. Phys. 81 (1997) 1639. [6] K. Kyllesbech Larsen, V. Privitera, S. Coa, F. Priolo, S.U. Campisano, A. Carnera, Phys. Rev. Lett. 76 (1996) 1493. [7] P.A. Stolk, D.J. Eaglesham, H.J. Gossmann, J.M. Poate, Appl. Phys. Lett. 66 (1995) 1370. [8] N.E.B. Cowern, G. Mannino, P.A. Stolk, F. Rooseboom, H.G.A. Huizing, J.G.M. van Berkum, F. Cristiano, A. Claverie, M. Jaraiz, Phys. Rev. Lett. 82 (1999) 4460. [9] J. Kim, J.W. Wilkins, F.S. Khan, A. Canning, Phys. Rev. B 55 (1997) 16186. [10] A. Bongiorno, L. Colombo, F. Cargnoni, C. Gatti, M. Rosati, in: D. Gershoni (Ed.), Proceedings of the 24th International Conference on the Physics of Semiconductor, World Scienti®c, Singapore, 1999. [11] A. La Magna, S. Coa, S. Libertino, L. Brambilla, P. Alippi, L. Colombo, Nucl. Instr. and Meth. B 179 (2001) 154. [12] S. Coa, S. Libertino, Appl. Phys. Lett. 73 (1998) 3369. [13] S. Libertino, J.L. Benton, D.C. Jacobson, D.J. Eaglesham, J.M. Poate, S. Coa, P.G. Fuochi, M. Lavalle, Appl. Phys. Lett. 71 (1997) 389. [14] G. Davies, Phys. Reports (Rev. Sector of Phys. Lett.) 176 (1989) 83. [15] S. Coa, S. Libertino, C. Spinella, Appl. Phys. Lett. 76 (2000) 321. [16] S. Libertino, J.L. Benton, S. Coa, D.J. Eaglesham, in: J.C. Barbour, S. Roorda, D. Ila (Eds.), Materals Reserach Society Symposium Proceedings 504 (1998) 3.
www.elsevier.nl/locate/nimb
Point defect diusion and clustering in ion implanted c-Si Sebania Libertino *, Salvatore Coa, Corrado Spinella, Antonino La Magna, Vittorio Privitera CNR-IMETEM, Stradale Primosole 50, I-95121 Catania, Italy
Abstract This paper reviews some fundamental aspects of point defect migration and agglomeration in crystalline Si. Both insitu and ex-situ measurements were used to reach this target. Room temperature (RT) diusivities of 1:5 10 15 and 3:0 10 13 cm2 =s for I and V, respectively, were obtained using in-situ leakage current measurements, performed during and just after ion implantation. To follow the defect evolution and clustering upon annealing, ex-situ optical and electrical measurements were used. Low temperature (300±500°C) annealing causes the formation of point-like defects, while higher temperatures (500±800°C) are necessary to have defect clustering. Finally, a well-de®ned dose 1 1013 Si=cm2 in pure Si) temperature (650°C) and time thresholds exist for the transition from I-clusters to extended {3 1 1} defects. When the transition takes place, both the optical and electrical defect properties undergo a dramatic change, suggesting an abrupt structural transition in the evolution from I-cluster to {3 1 1} defects. Kinetic lattice Monte-Carlo simulations used to model the defect agglomeration and growth con®rm these results. Ó 2001 Elsevier Science B.V. All rights reserved. Keywords: Defects; Interstitial clusters; Silicon; Migration; Diusivity; Photoluminescence
1. Introduction Defect formation and agglomeration in crystalline Si (c-Si) have been the subject of many studies in the last 30 years [1±3]. A strong driving force to the research is provided by technological interest since Si is the most used semiconductor in electronic devices fabrication. The method of choice to dope Si is ion implantation that causes the generation of point defects: self-interstitials (I) and vacancies (V). Due to their low migration energy, 1 eV [4], I and
*
Corresponding author. Tel.: +39-95-591212; fax: +39-957139154. E-mail address: [email protected] (S. Libertino).
V can freely migrate even at room temperature (RT). We showed that, after MeV He implantation in Si only a small percentage of the generated defects succeeds in escaping direct recombination according to the relation I V ;. The ®nal defect concentration strongly depends on the Si impurity content, varying from 16% (high content) to below 3% (low content) [5]. The I and V that escape recombination interact with impurities (C and O), dopant atoms (e.g. P or B) or among themselves and form RT stable defects [1,2]. V-type defects like the divacancy (VV), oxygen-vacancy (OV) and phosphorous-vacancy (PV) and I-type defects like carbon interstitial-oxygen Ci O, phosphorous± carbon interstitial PCi , boron interstitial-oxygen Bi O, boron interstitial-carbon Bi C are left in the
0168-583X/01/$ - see front matter Ó 2001 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 0 ) 0 0 5 0 4 - 8
26
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
sample. It should be mentioned that I-type defect formation is a two-step mechanism: I interact with C and B forming interstitial-carbon Ci or interstitial-boron Bi and these species migrate to form the C- or B-related defects previously listed. All these defects introduce deep levels in the Si band gap [1] and can be detected with electrical measurements such as deep-level transient spectroscopy (DLTS) [5] or spreading resistance pro®ling (SRP) [6]. Despite eorts, up to now the I and V RT diffusivity values are not known. During annealing, RT stable complexes dissociate releasing the I and V they stored. Most of them will recombine until only I remain in the sample. In fact, due to the extra-implanted ion the number of I generated per implanted ion is higher (+1) than the number of V. These extra I agglomerate in more complex structures and, eventually, extended defects (e.g. {3 1 1} defects) form. They dissociate upon further annealing causing a large I supersaturation which drives several phenomena such as B transient enhanced diusion (TED) [7]. Since the TED enhancement and duration depend on the thermal stability of defects storing the I, it is fundamental to characterize them. While there are detailed experimental and theoretical studies on extended defects, e.g. {3 1 1} [7], little is known on the nanometer-sized I-clusters which will eventually grow into {3 1 1} defects. Recently, Cowern et al. [8] studied the TED at various annealing temperatures and extracted the I supersaturation values as a function of time. Assuming an Ostwald ripening mechanism for I agglomeration, they found that the experimental data can be ®tted only assuming that two stable Icluster sizes of roughly 4 and 8 I exist, but nothing can be said on their structure. On the other hand, tight-binding (TB) calculations [9] showed that the basic {3 1 1} structure, an in®nitely extended (1 1 0) I chain, cannot be the lowest energy con®guration for small clusters. The formation energy per I, very low if the chain is in®nite (1.7 eV), becomes very large (4.7 eV in a di-interstitial) when the number n of I in the chain is low (<10). Recent TB calculations of I agglomeration [10] suggest that alternative structures storing few I n < 10 could be a more favorable pathway for I agglomeration and growth. This implies that small I-clusters are not
the direct precursors of {3 1 1} defects and that a structural transformation has to occur during the growth process. Kinetic lattice Monte-Carlo simulations used to model defect agglomeration and {3 1 1} formation [11] con®rm this scenario. The aim of this paper is to assess the basic I and V properties, follow the I agglomeration into clusters and understand the growth mechanism from I-clusters to {3 1 1} defects. 2. Experimental We used both n- and p-type Czochralski grown (CZ, 7 1017 O cm 3 , 1 1016 C cm 3 ) and epitaxial (epi, O and C concentrations 5 1015 cm 3 ) Si with resistivities in the range 0.7± 4 X cm. Schotthy barriers and p ±n junctions were used for DLTS and leakage current IL measurements. The samples were implanted with He at an energy of 2.5 MeV to doses of 1 109 cm 2 to 1 1012 cm 2 and dose rates of 1 108 cm 2 to 1 1010 cm 2 s 1 or with Si at energies of 0.04±1.2 MeV to doses of 1 1012 cm 2 to 5 1013 cm 2 . DLTS measurements were performed using 1 ms ®lling pulse and reverse biases such as to include the entire damage distribution within the depletion layer. Defect depth pro®les were obtained changing the ®lling pulse voltage. Photoluminescence (PL) and SR measurements were performed on non-structured samples. PL measurements were carried out at 17 K using the 488 nm line of an Argon laser at a pump power of 50 mW. Light emitted from the sample was dispersed by a monochromator and detected by a liquid nitrogen cooled Ge detector. Standard lock in technique was used to improve the signal-to-noise ratio. Sample morphology was analyzed by transmission electron microscopy (TEM) in both plan view and cross-sectional con®gurations. 3. Results and discussion 3.1. Point defect RT migration Precious information on I and V migration in c-Si can be obtained by SRP. Defects are generated
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
at the surface by low-energy ion implantation and their migration monitored through the dopant deactivation tail [6]. We performed a shallow Si implant, 40 keV, in n-type epi Si to produce a damaged region which extends from the surface to 0.4 lm. Some of the V and I generated succeed in escaping recombination or clustering in this region and are injected into the bulk. The probability of recombination and/or clustering in the bulk is small, due to the very low defect concentration in this region, hence, they undergo long-range migration until trapping occurs. In highly pure epi Si, migration up to 2 lm is reported [6]. The SRP measurements plotted in Fig. 1 refer to a sample as implanted with 5 1012 Si=cm2 . A tilt angle of 7° during implantation was used to avoid any channeling eect that could aect the measurement. The vertical solid line in the ®gure shows the region directly damaged by the beam. Within this region, a strong decrease in the free carrier concentration, due to the damage presence, is observed. A deactivation tail extending up to 0.9 lm is observed: deeper, the substrate-free carrier concentration 7 1015 cm 3 is measured. After 150°C for 30 min annealing, the deactivation tail is fully recovered H, suggesting that the defects
Fig. 1. SRP measurements on n-type epi Si implanted with 40 keV Si to a dose of 5 1012 cm 2 and a 7° tilt angle: as-implanted ; after annealing at 150°C for 30 min H. The vertical solid line indicates the region directly damaged by the beam.
27
formed are point-like. Part of the damage in the ®rst 0:4 lm recovers too, showing that those defects are also present in the region directly interested by the beam. The deactivation tail is visible only in epi Si: high contents of C or O trap the I and V as soon as they are formed preventing their migration. SRP is not a spectroscopic technique; hence it cannot identify the origin of the deactivation tail. The spectroscopy of the defects in the tail was studied by measuring the depth pro®les of the VV , PV H and PCi j complexes for the same sample, see Fig. 2. These defects are the only ones that can aect the SRP measurement since they lie close to the center of the gap. The VV and PV peaks lie at the same temperature position in DLTS: hence, when both defects are present, the two DLTS signals overlap [1]. The activation energy of the resulting level is that of the dominating one, usually the VV . We separated the two contributions by measuring the peak activation energy for each point plotted in the ®gure. In the region where the VV is present (0.4±0.6 lm), the PV presence was monitored by measuring the VV pro®le in the as-implanted state and after 150°C 30 min annealing. Since the PV complex anneals at 150°C [1], the dierence between the two VV depth pro®les provides the PV concentration in
Fig. 2. DLTS depth pro®les of VV (), PV H and PCi j defect complexes for an n-type Si sample as-implanted with 40 keV Si to a dose of 5 1012 cm 2 and a 7° tilt angle. The vertical solid line shows the region directly interested by the beam.
28
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
that region. The same procedure was used for the PV points lying between 0.6 and 0:9 lm: the activation energy in each point of the depth pro®le and the defect annealing kinetics were monitored. While the VV concentration is very high in the directly damaged region and decreases abruptly going deeper into the sample, the PV behavior is quite dierent. It is still detectable up to 0:9 lm con®rming that V migrate into the bulk until trapping at impurities occurs. On the other hand, the PCi presence shows that I migrate even to longer distances (1.3 lm), probably due to the two-step mechanism of I-type defect formation. The deactivation tail is given by both P atoms enclosed in the complexes, that become inactive, and P atoms compensated by the complexes deep levels. When the defects break at 150°C, the deactivation tail recovers since the P in complexes returns electrically active and the levels are removed from the gap (see Fig. 1). Both the PV and the PCi dissociate at this temperature. Ex-situ measurements cannot provide absolute values for I and V diusivities since a time delay exists between implantation and measure, hence in-situ measurements are necessary. We monitored in-situ the IL evolution during and just after implantation of p ±n junctions with 2.5 MeV He. A schematic of the measurement setup is shown in the inset of Fig. 3: the sample was reverse biased during and after implantation to a voltage of )30 V. The reverse bias was chosen in order to embody the entire defect pro®le in the depletion region. The results of an implant with a dose rate of 1 108 cm 2 s 1 are shown in Fig. 3. At the beam turn-on, IL (30 pA at )30 V in the unimplanted diode) increases monotonically as a result of the formation of defect complexes, hence of deep levels in the Si band gap. At the beam turn-o, when a total dose of 1 1012 cm 2 is reached (after 104 s), IL decreases suggesting that a partial annihilation of the defects responsible for carrier generation in the depletion layer is occurring. After 1 day, IL is reduced by a factor of 2 and the transient almost saturates, with small variations lasting for several days. To identify the defects responsible of IL , DLTS measurements were performed. Since IL strongly depends on levels located close to mid gap, the main contribution arises
Fig. 3. In-situ IL measurements during and after implantation with 2.5 MeV He to a dose of 1 1012 cm 2 and a dose rate of 1 108 cm 2 s 1 . The vertical dashed line indicates the beam turn-o. The solid line is a ®t of the data. In the inset schematic of the experimental setup is plotted.
from VV and PV. DLTS and IL data were compared assuming a linear dependence of IL on the defect content. During implantation, free V agglomerate into V-type defects, producing the IL increase. After implantation, a residual concentration of free I still exists and can migrate to Vtype defects dissociating them. The IL time evolution is directly linked to I and V migrations. The data were ®tted (solid line) solving a set of coupled dierential equations describing diusion-limited defect annihilation, agglomeration and trapping and having the I and V diusivity as the only parameters [12]. Values of 1:5 10 15 cm2 =s and 3 10 13 cm2 =s, respectively, were found. These low values suggest that, when monitoring the defect migration in the depletion layer the enhancement due to ionization eects is limited and values close to those expected for neutral V and I are achieved. Since V-type defects might be annihilated by the mobile Ci , we checked the C eect assuming a Ci diusivity of 10 15 cm2 =s. However, simulations performed with 5 1015 CS cm 3 (the C content in our samples) or without C give the same results. Finally, so far we assumed that the electric ®eld presence, typical of these measurements, does not aect the ®nal defect distribution. This is not fully correct since recent measurements have shown that
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
the ®nal defect pro®le is slightly dierent in concentration and spatial distribution if an electric ®eld is applied. It is shown in Fig. 4 where the VV depth pro®les of samples implanted to 1010 He cm 2 and a dose rate of 108 cm 2 s 1 reverse biased or unbiased M during implantation are compared. The unbiased sample shows a higher defect concentration and a slightly dierent depth position. The same results hold for all the measured samples and also when the OV is monitored. We believe it occurs because, under our experimental condition, there is a fraction of electrically charged defects. These defects escape from the depletion region, due to the electric ®eld presence, thus avoiding the complexes formation and causing a shift of the ®nal defect depth pro®le. Further measurements are in progress to understand and quantify this phenomenon. 3.2. Defect evolution: clustering and extended defect formation Annealing at temperatures P 350°C dissociates RT stable defects causing the release of the I and V they stored. Most of them undergo a direct recombination and the survivors are, again, stored in complexes. These ``new'' defect characteristics strongly depend on the annealing temperature and
Fig. 4. VV depth pro®le of an n-type Si sample implanted with 2.5 MeV He to a dose of 1 1010 cm 2 and a dose rate of 1 108 cm 2 s 1 reverse biased to 30 V or unbiased (M) during implantation.
29
dier from those of RT stable defects. For annealing at temperatures 6 500°C, I-type defects, storing just one or two I, remain in the sample [13]. On the other hand, samples implanted with 1.2 MeV Si to a dose of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line) and annealed at 600°C for 4 h exhibit the PL spectra shown in Fig. 5. Although the PL intensity increases with ion dose, the major features remain unchanged. They consist of two broad peaks centered at 1300 and 1400 nm, associated to I-clusters [15]. Broad peaks in PL spectra are usually observed in samples with extended defects. They were associated to the quantum con®nement of carriers in regions with high strain surrounding the defects. Since no extended defects are detected in our samples, we associated this feature to the carrier recombination in the strained region surrounding the I-clusters embedded in the Si matrix [15]. Several sharp lines, associated to point-like defect complexes, are superimposed to the I-clusters signatures in the range 1200±1280 nm. Among them, the W0 line at 1233 nm (1.0048 eV) is present. It is a perturbed form of the W line (associated to I-rich defects) observed in regions with high strain [14], con®rming the strain presence in our samples. The spectra also show the oxygen thermal donors (at 1620 nm, 0.765 eV) and the Si band-edge phonon assisted recombination line at 1121 nm (1.1056 eV). The PL spectra of I-clusters are totally dierent from those of point-like defects suggesting dierent defect structures in the two cases. Moreover, the PL peak broadening of the clusters spectra suggests a complex structure for these defects causing a measurable strain in the lattice, as con®rmed by the TEM cross-section shown in the inset of Fig. 5. No extended defects are visible even at the highest implanted dose. Only a weak contrast at the projected range is observed, demonstrating that a heavily damaged region probably consisting of small defect aggre is present in the sample. Annealing gates (<50 A) for times up to 15 h at 600°C, only produces a decrease in the PL signal intensity suggesting that the I-clusters annealate and no extended defect form. These results suggest that there is a threshold temperature for the {3 1 1} defect formation. At T 6 600°C the I are stored into clusters (for
30
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
Fig. 5. PL spectra on n-type Si samples implanted with 1.2 MeV Si to doses of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line). The samples were annealed at 600°C for 4 h. In the inset: TEM cross-section made in weak beam g 4g using the (1 1 1) re¯ection of the sample implanted at 5 1013 Si cm 2 .
doses P 1012 Si cm 2 ), eventually released and annealed, probably at the surface. Annealing at 680°C, on the same samples shown in Fig. 5, produces major modi®cations in the defect optical and structural properties. The PL spectra of the same samples annealed for 1 h are plotted in Fig. 6. The sample implanted at 1012 Si cm 2 (dotted line) shows only the Si bandedge line, while at doses P 1013 cm 2 a sharp peak at 1376 nm (0.9007 eV) dominates the spectra. The study of the PL spectra and the TEM analysis of damaged samples allowed us to associate it to optical transitions occurring at or close to {3 1 1} extended defects [15]. Its width slightly increases with the dose and a broad band centered at 1550 nm (0.799 eV) is developed after 5 1013 cm 2 implantation. These measurements clearly indicate that to have the transition from I-clusters to {3 1 1} extended defects the implantation dose must be above a threshold that, for our CZ Si samples, is 1 1013 cm 2 . Previous studies [7] showed that the minimum dose necessary to form {3 1 1} defects increases with the sample C content, since it is an ecient trap for I and prevents their agglomeration. Hence, in samples with low C content, as our
CZ Si, the implantation dose threshold should be lower. Finally, we ®xed the implantation dose at 1 1013 Si=cm2 and the temperature at 700°C, both above thresholds for {3 1 1} formation, and monitored the defect evolution as a function of the
Fig. 6. PL spectra on n-type Si implanted with 1.2 MeV Si to doses of 1 1012 cm 2 (dotted line), 1 1013 cm 2 (solid line) and 5 1013 cm 2 (dashed line), and annealed at 680°C for 1 h.
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
annealing time. The resulting PL spectra are compared in Fig. 7. Already after 10 min (dotted line) the I-clusters are formed, suggesting a fast agglomeration process. Annealing up to 30 min causes only a decrease in the I-clusters signal intensity, probably due to their partial dissolution. The Si band-edge line, barely visible after 10 min becomes stronger as the time proceeds, con®rming that the crystal damage is recovering. After 1h annealing (solid line), the {3 1 1} characteristic signature appears, even if a strong contribution of the I-clusters signatures is visible. Only after 2 h (dash±double dot-dashed line) the {3 1 1} signal is the only one detectable in the spectrum. The results so far reported, together with our studies using DLTS [16], indicate that major transformations in the optical, electrical and structural properties of Si occur to allow the {3 1 1} defect formation. At the early stages of nucleation small I-clusters form. We believe that their structures signi®cantly dier from of the {3 1 1} defects. As annealing proceeds, some clusters grow bigger, but there is a critical dimension at which the I-clusters need to rearrange themselves in a dierent structure, probably close to the (1 1 0) I chain that constitutes the {3 1 1} defects. Of course, a nucleation barrier must exist in order to have the transition. After annealing at low
31
temperature, or for short time, the probability to overcome this barrier is small and most of the I remain in clusters until they anneal out. On the other hand, increasing the temperature, and/or the time, for doses above threshold, the probability of transition becomes very large and {3 1 1} defects form. Our results are in agreement with those by Cowern et al. [8]. He assumes that a very stable Icluster con®guration exists and that a nucleation barrier has to be overcome to form {3 1 1} defects. Kinetic lattice Monte-Carlo simulations reported in this same number [11] con®rm this hypothesis. Further studies are in progress to determine the critical cluster size. 4. Conclusion We have shown that I and V RT migration can be monitored using ex-situ and in-situ techniques, since their migration give the ®nal defect distribution. Upon annealing, I-clusters and extended defects form. By PL measurements we followed the transitions since the defect optical properties abruptly change. In particular, we found that for MeV Si implantation a threshold dose 1013 cm 2 and temperature (650°C) exist for {3 1 1} forma-
Fig. 7. PL spectra on p-type Si implanted with 1.2 MeV Si to a dose of 1 1013 cm 2 and annealed at 700°C for 10 min (dotted line), 20 min (dash±dot-dashed line), 30 min (dashed line), 1 h (solid line) and 2 h (dash-double dot-dashed line).
32
S. Libertino et al. / Nucl. Instr. and Meth. in Phys. Res. B 178 (2001) 25±32
tion. Once the dose 1013 cm 2 and temperature (700°C) are above threshold a certain time (1 h) is necessary to allow the {3 1 1} defect formation. We believe that a structural transformation accompanies the transition from I-cluster to {3 1 1} defects. Acknowledgements The authors acknowledge A. Spada and N. Parasole for the invaluable technical assistance, A. Marino for the implants and S. Pannitteri for the TEM samples preparation. This work was partially supported by Progetto Finalizzato MADESS II and Progetto 5% Microelettronica. References [1] L.C. Kimerling, in: N.B. Urli, J.M. Corbett (Ed.), Radiation Eects in Semiconductors, Inst. Phys. Conf. Ser. 31, London, 1977, p. 221. [2] G.D. Watkins, Mater. Res. Soc. Symp. Proc. 469 (1997) 139. [3] B.G. Svensson, C. Jagadish, A. Hallen, J. Lalita, Nucl. Instr. and Meth. B 106 (1995) 183. [4] P.M. Fahey, P.B. Grin, J.D. Plummer, Rev. Mod. Phys. 61 (1989) 289.
[5] S. Coa, V. Privitera, F. Priolo, S. Libertino, G. Mannino, J. Appl. Phys. 81 (1997) 1639. [6] K. Kyllesbech Larsen, V. Privitera, S. Coa, F. Priolo, S.U. Campisano, A. Carnera, Phys. Rev. Lett. 76 (1996) 1493. [7] P.A. Stolk, D.J. Eaglesham, H.J. Gossmann, J.M. Poate, Appl. Phys. Lett. 66 (1995) 1370. [8] N.E.B. Cowern, G. Mannino, P.A. Stolk, F. Rooseboom, H.G.A. Huizing, J.G.M. van Berkum, F. Cristiano, A. Claverie, M. Jaraiz, Phys. Rev. Lett. 82 (1999) 4460. [9] J. Kim, J.W. Wilkins, F.S. Khan, A. Canning, Phys. Rev. B 55 (1997) 16186. [10] A. Bongiorno, L. Colombo, F. Cargnoni, C. Gatti, M. Rosati, in: D. Gershoni (Ed.), Proceedings of the 24th International Conference on the Physics of Semiconductor, World Scienti®c, Singapore, 1999. [11] A. La Magna, S. Coa, S. Libertino, L. Brambilla, P. Alippi, L. Colombo, Nucl. Instr. and Meth. B 179 (2001) 154. [12] S. Coa, S. Libertino, Appl. Phys. Lett. 73 (1998) 3369. [13] S. Libertino, J.L. Benton, D.C. Jacobson, D.J. Eaglesham, J.M. Poate, S. Coa, P.G. Fuochi, M. Lavalle, Appl. Phys. Lett. 71 (1997) 389. [14] G. Davies, Phys. Reports (Rev. Sector of Phys. Lett.) 176 (1989) 83. [15] S. Coa, S. Libertino, C. Spinella, Appl. Phys. Lett. 76 (2000) 321. [16] S. Libertino, J.L. Benton, S. Coa, D.J. Eaglesham, in: J.C. Barbour, S. Roorda, D. Ila (Eds.), Materals Reserach Society Symposium Proceedings 504 (1998) 3.