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Defects induced by high energy helium implantation in 4H–SiC
Nuclear Instruments and Methods in Physics Research B 186 (2002) 218±222 www.elsevier.com/locate/nimb
Defects induced by high energy helium implantation in 4H±SiC M.F. Beaufort a
a,*
, E. Oliviero a, M.L. David a, J. Nomgaudyte b, L. Pranevicius b, A. Declemy a, J.F. Barbot a
Laboratoire de M etallurgie Physique, UMR6630, BP30179, 86960 Chasseneuil-Futuroscope, France b Vytautas Magnus University, 58 Donelaiciao St., Kaunas LT-3000, Lithuania
Abstract 1.6 MeV He ions were implanted at room temperature into (0 0 0 1) 4H±SiC at a dose of 1 1017 cm 2 . Using cross-sectional transmission electron microscopy, we have investigated the damage induced by the implantation and by a 1500 °C annealing. In the as-implanted sample, the damage region consists of three layers, including a continuous amorphous layer surrounded with crystalline zones. After annealing, recrystallization of the amorphous state occurs and large bubbles or cavities are observed. A simple model based on atomic relocation has been proposed to explain the layered structure observed after implantation. Ó 2002 Elsevier Science B.V. All rights reserved. Keywords: SiC; He; Voids; Ion implantation; Defects; TEM
1. Introduction In silicon, the implantation of helium or hydrogen ions leads to the formation of gas-vacancy complexes that evolve in gas bubbles for a sucient amount of incident ions. After subsequent annealing gas desorption occurs and large cavities or voids are observed [1]. These stable cavities have many technological applications such as impurity gettering or local lifetime control in power devices [2]. Since the wide band gap semiconductor SiC is of great interest as a material for power
* Corresponding author. Tel.: +33-5-49-49-68-34; fax: +33-549-49-66-92. E-mail address: [email protected] (M.F. Beaufort).
control and high speed communication devices, it should be also interesting to form stable cavities in these materials. A fundamental understanding of irradiation damage in SiC as well as their recovery is also needed to advance in technological applications. However, ion implantation into SiC generates strong damage up to the amorphization [3] and recrystallization will occur during sucient high temperature annealing [4]. Non-destructive and contactless techniques such as infrared re¯ectivity (IRR) and X-ray experiments (XRD) have given information on ion implantation-induced damaged layers such as width and strain [5]. In this paper XTEM have been performed to study cavity formation and defects evolution in n-type 4H±SiC implanted by room temperature MeV helium. We also discuss our results in light of a model based on atomic displacement and relocation processes.
0168-583X/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 1 ) 0 0 9 1 2 - 0
M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
219
2. Experimental details Cree research single-crystalline 4H±SiC (0 0 0 1) wafers were implanted at room temperature with 1.6 MeV He ions using a van de Graa accelerator (CERI-Orleans). The ion ¯uence used was 1 1017 cm 2 leading to a peak concentration of implanted atoms larger than the 1.6 at.% required for bubble formation in silicon [1]. Annealing was done at 1500 °C for 30 min in high vacuum. Crosssectional TEM foils were prepared by an ion milling thinning down to electron transparency. The microstructure of the implanted zone was examined using a Jeol 200CX microscope operating at 200 kV.
3. Results He implantation in 4H±SiC (1.6 MeV, 1 1017 He cm 2 , without anneal) produces a continuous damaged layer of approximately 650 nm in width and located at about 3.6 lm below the surface, as expected from SRIM calculations [6]. As observed after the implantation of much heavier ions in SiC [7,8] the helium implantationinduced microstructure is non-homogeneous along thickness and can be divided in three layers. Fig. 1 shows that the damaged layer after 1.6 MeV He implantation can be divided in three dierent regions labeled A, B and C. A select area diraction pattern of the B layer (in the inset 1) shows a halo pattern corresponding to an amorphous phase while A and C are crystalline SiC. This amorphous central layer, 400 nm wide, is made up of two regions (B1 and B2 ). The B1 region is related to the presence of small bubbles with a mean diameter about 1 nm, a closer view of these bubbles is shown in the inset 2. B2 part does not show any visible bubbles. B1 and B2 widths are about 280 and 120 nm, respectively. The A and C outside regions (170 and 55 nm wide, respectively) show a dark contrast which results from small defect clusters not resolvable by conventional TEM. The 400 nm width of the amorphous layer is in good agreement with the width of the buried layer obtained by IRR [5].
Fig. 1. XTEM micrograph of 4H±SiC implanted at 1.6 MeV with 1 1017 He cm 2 . In the inset: 1 ± diraction pattern of the B zone, 2 ± small bubbles in the B1 part only.
Fig. 2 shows dark ®eld cross-section image of the annealed sample (1500 °C ± 30 min). As seen, the amorphous layer is totally recrystallized. The A and C regions present in the as-implanted sample are partially recovered; only a low density of small defects (not resolvable using conventional microscopy) can be observed on both sides of the B layer.
Fig. 2. Dark ®eld XTEM image on 4H±SiC implanted with 1 1017 He cm 2 and annealed (1500 °C, 30 min).
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M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
Fig. 3. Bright ®eld XTEM image showing bubbles in the middle of the B layer.
These small defects extend up to the implanted surface and only a few hundred nanometers into the bulk. In the same way, XRD studies have shown an almost total structural recovery of the ion implantation damage crystal after 1500 °C annealing [5]. Diraction pattern shows the complete crystallization of the amorphous layer including polytypism [4]. As seen in Fig. 3, large bubbles and/or cavities with size ranging from 15 to 50 nm are present inside the B region after annealing. Two bubble shapes can be distinguished, circular and faceted shapes with facets in the [)1 1 0 0] or/and [0 0 0 1] directions. We do not observe any dislocation loops in contrary to what was expected from a previous study on hot-pressed SiC [9]. These authors mentioned that bubbles are always accompanied by faulted dislocation loops of interstitial type. The implanted doses, however, were smaller than the one used in this study.
4. Modeling of structural modi®cations The room temperature helium bombardment induces atomic displacements, which result in ballistic relocations of atoms between neighboring monolayers [10]. For the quantitative analysis the solid is divided into monolayers with thickness h0
and the frequency probability of relocation of i K;L atom from K into monolayer L, wi , is introduced (i 1 for Si and i 2 for C atoms). The determination of the relocation function needs the solution of integro-dierential equations [11]. In the present work the relocation function is constructed from the following physical considerations. The number of displaced i atoms per 1 s in K K the K-th monolayer is equal to wi Di I0 =C, K where Di is the number of displaced i atoms per one incident ion in monolayer K; I0 is the ¯ux of incident ions and C is the number of atoms in one monolayer. The assumptions are made that the number of displaced atoms in the monolayer K, K Di , follows the Gaussian distribution of the nuclear energy losses, and the relocation probability of displaced atoms from monolayer K to L decreases with the increase in relocation distance 2 h0 jK Lj as exp h0 jK Lj=li , where li is the characteristic relocation distance of i atoms [12]. K;L It gives that wi can be written as " K;L wi
0 wi
exp
h0 jK li
h0 K
1 Rpd DRpd 2 # Lj ;
2
where DRpd and Rpd are the standard deviation and projected range of distribution of energy losses. 0 The parameter wi characterizes the displacement rate of i atoms in the monolayer corresponding to the maximum of nuclear energy losses (K 1 Rpd ) and are not relocated (jK Lj 0). It is function of the ¯ux and energy of incident ions and depends on the displacement energy of atoms. Computer simulation results show that the number of displaced Si atoms in SiC exceeds the number of displaced C atoms per one incident ion 0 0 about two times (w1 2=w2 ) [13]. Parameters 0 wi , DRpd and Rpd have been evaluated by SRIM: 0 0 w1 2 10 5 s 1 , w2 1 10 5 s 1 , Rpd 3:6 lm and DRpd 0:1 lm, other parameters used in calculations are l1 l2 h0 and h0 0:26 nm. The equation describing the time variations of the partial concentration of components in dierent monolayers initiated by the simultaneous ac-
M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
tion of displacement and relocation processes can be written as 1 X dci L;K L wi ci dt L1 K
K
ci
1 X L 1
K;L
wi
:
The ®rst term de®nes the ¯ux of i atoms arriving into monolayer K from any other one, and the second term de®nes the ¯ux of i atoms leaving monolayer K into any other one including vacuum (negative indexes). Fig. 4 illustrates calculated concentration pro 0 ®les of atoms for two sets of parameter wi . It is seen that the area with maximum nuclear losses is depleted by atoms (c1 c2 < 1). The asymmetry in distribution pro®les manifests non-symmetrical distribution of nuclear energy losses. The best ®t s with experiments is obtained if DRpd 4DRpd , s where DRpd is the standard deviation of the nonsymmetrical Gaussian distribution pro®le from the front side. The displaced atoms occupy interstitial positions forming areas c1 c2 > 1. The coordinates corresponding to c1 c2 1 de®ne boundaries between areas enriched by vacancies and interstitials. Processes of annihilation and formation of extended defects are not considered in the present work. The calculated results are in qualitative agreement with TEM observations. The three areas can be distinguished: the central enriched by vacancies and two neighboring areas enriched by intersti-
221
tials. It is seen that the front layer A is wider (170 nm) than the near layer C (55 nm) and the density appears to be larger in the C layer than in A one. The above conducted analysis of the displacement±relocation processes gives complementary information about distribution pro®les of vacancies and interstitials. It can be concluded that nonhomogeneous displacement of atoms in dierent monolayers calculated for identical relocation probabilities initiates ¯uxes of atoms from areas of maximum nuclear losses to periphery forming areas enriched by vacancies and interstitials.
5. Conclusion The microstructures of high energy helium implanted 4H±SiC have been studied using conventional electron microscopy. When a sucient dose of helium is implanted into crystalline SiC sample, nucleation sites and bubbles are formed but the implantation also leads to amorphization. A three layer structure is observed: an amorphous one surrounded by crystalline zones with point defects. Bubbles evolve with recrystallization during thermal annealing at 1500 °C. A model based on the analysis of displacement and relocation processes of atoms in SiC gives a three layered structure as experimentally observed, enriched by vacancies in the area of maximum nuclear losses of incident ions and interstitials on the peripheries.
References
Fig. 4. Concentration pro®le of atoms after implantation for two couples of w0i values.
[1] D.M. Follstaedt, S.M. Myers, C.A. Petersen, J.W. Medernach, J. Electron. Mater. 25 (1996) 157. [2] V. Raineri, G. Fallica, S. Libertino, J. Appl. Phys. 79 (1996) 9012. [3] J.A. Spitznagel, S. Wood, W.J. Choyke, N.J. Doyle, J. Bradshaw, G. Fishman, Nucl. Instr. and Meth. B 16 (1996) 237. [4] E. Oliviero, M.L. David, M.F. Beaufort, J. Nomgaudyte, L. Pranevicius, A. Declemy, J.F. Barbot, J. Appl. Phys., to be published. [5] A. Declemy, E. Oliviero, M.F. Beaufort, J.F. Barbot, M.L. David, C. Blanchard, Y. Tessier, E. Ntsoenzok, Nucl. Instr. and Meth B 186 (2002) 318. [6] J.F. Ziegler, J.P. Biersack, U. Littmark, The Stopping and Range of Ions in Solids, Pergamon, New York, 1985.
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[7] M. Ishimaru, S. Harada, T. Motooka, T. Nakata, T. Yoneda, M. Inoue, Nucl. Instr. and Meth. B 127±128 (1997) 195. [8] V. Heera, J. Stoemenos, R. K ogler, W. Skorupa, J. Appl. Phys. 77 (1995) 2999. [9] J. Chen, P. Chung, H. Trinkaus, Phys. Rev. B 61 (2000) 12923.
[10] H.M. Urbassek, G. Mayer, H. Gades, M. Vicanek, Nucl. Instr. and Meth. B 103 (1995) 275. [11] P. Sigmund, A. Gras-Marti, Nucl. Instr. and Meth. B 182±183 (1981) 25. [12] P. Sigmund, A. Oliva, Nucl. Instr. and Meth. B 82 (1993) 269. [13] W. Jang, Nucl. Instr. and Meth. B 148 (1999) 557.
Defects induced by high energy helium implantation in 4H±SiC M.F. Beaufort a
a,*
, E. Oliviero a, M.L. David a, J. Nomgaudyte b, L. Pranevicius b, A. Declemy a, J.F. Barbot a
Laboratoire de M etallurgie Physique, UMR6630, BP30179, 86960 Chasseneuil-Futuroscope, France b Vytautas Magnus University, 58 Donelaiciao St., Kaunas LT-3000, Lithuania
Abstract 1.6 MeV He ions were implanted at room temperature into (0 0 0 1) 4H±SiC at a dose of 1 1017 cm 2 . Using cross-sectional transmission electron microscopy, we have investigated the damage induced by the implantation and by a 1500 °C annealing. In the as-implanted sample, the damage region consists of three layers, including a continuous amorphous layer surrounded with crystalline zones. After annealing, recrystallization of the amorphous state occurs and large bubbles or cavities are observed. A simple model based on atomic relocation has been proposed to explain the layered structure observed after implantation. Ó 2002 Elsevier Science B.V. All rights reserved. Keywords: SiC; He; Voids; Ion implantation; Defects; TEM
1. Introduction In silicon, the implantation of helium or hydrogen ions leads to the formation of gas-vacancy complexes that evolve in gas bubbles for a sucient amount of incident ions. After subsequent annealing gas desorption occurs and large cavities or voids are observed [1]. These stable cavities have many technological applications such as impurity gettering or local lifetime control in power devices [2]. Since the wide band gap semiconductor SiC is of great interest as a material for power
* Corresponding author. Tel.: +33-5-49-49-68-34; fax: +33-549-49-66-92. E-mail address: [email protected] (M.F. Beaufort).
control and high speed communication devices, it should be also interesting to form stable cavities in these materials. A fundamental understanding of irradiation damage in SiC as well as their recovery is also needed to advance in technological applications. However, ion implantation into SiC generates strong damage up to the amorphization [3] and recrystallization will occur during sucient high temperature annealing [4]. Non-destructive and contactless techniques such as infrared re¯ectivity (IRR) and X-ray experiments (XRD) have given information on ion implantation-induced damaged layers such as width and strain [5]. In this paper XTEM have been performed to study cavity formation and defects evolution in n-type 4H±SiC implanted by room temperature MeV helium. We also discuss our results in light of a model based on atomic displacement and relocation processes.
0168-583X/02/$ - see front matter Ó 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 8 - 5 8 3 X ( 0 1 ) 0 0 9 1 2 - 0
M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
219
2. Experimental details Cree research single-crystalline 4H±SiC (0 0 0 1) wafers were implanted at room temperature with 1.6 MeV He ions using a van de Graa accelerator (CERI-Orleans). The ion ¯uence used was 1 1017 cm 2 leading to a peak concentration of implanted atoms larger than the 1.6 at.% required for bubble formation in silicon [1]. Annealing was done at 1500 °C for 30 min in high vacuum. Crosssectional TEM foils were prepared by an ion milling thinning down to electron transparency. The microstructure of the implanted zone was examined using a Jeol 200CX microscope operating at 200 kV.
3. Results He implantation in 4H±SiC (1.6 MeV, 1 1017 He cm 2 , without anneal) produces a continuous damaged layer of approximately 650 nm in width and located at about 3.6 lm below the surface, as expected from SRIM calculations [6]. As observed after the implantation of much heavier ions in SiC [7,8] the helium implantationinduced microstructure is non-homogeneous along thickness and can be divided in three layers. Fig. 1 shows that the damaged layer after 1.6 MeV He implantation can be divided in three dierent regions labeled A, B and C. A select area diraction pattern of the B layer (in the inset 1) shows a halo pattern corresponding to an amorphous phase while A and C are crystalline SiC. This amorphous central layer, 400 nm wide, is made up of two regions (B1 and B2 ). The B1 region is related to the presence of small bubbles with a mean diameter about 1 nm, a closer view of these bubbles is shown in the inset 2. B2 part does not show any visible bubbles. B1 and B2 widths are about 280 and 120 nm, respectively. The A and C outside regions (170 and 55 nm wide, respectively) show a dark contrast which results from small defect clusters not resolvable by conventional TEM. The 400 nm width of the amorphous layer is in good agreement with the width of the buried layer obtained by IRR [5].
Fig. 1. XTEM micrograph of 4H±SiC implanted at 1.6 MeV with 1 1017 He cm 2 . In the inset: 1 ± diraction pattern of the B zone, 2 ± small bubbles in the B1 part only.
Fig. 2 shows dark ®eld cross-section image of the annealed sample (1500 °C ± 30 min). As seen, the amorphous layer is totally recrystallized. The A and C regions present in the as-implanted sample are partially recovered; only a low density of small defects (not resolvable using conventional microscopy) can be observed on both sides of the B layer.
Fig. 2. Dark ®eld XTEM image on 4H±SiC implanted with 1 1017 He cm 2 and annealed (1500 °C, 30 min).
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M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
Fig. 3. Bright ®eld XTEM image showing bubbles in the middle of the B layer.
These small defects extend up to the implanted surface and only a few hundred nanometers into the bulk. In the same way, XRD studies have shown an almost total structural recovery of the ion implantation damage crystal after 1500 °C annealing [5]. Diraction pattern shows the complete crystallization of the amorphous layer including polytypism [4]. As seen in Fig. 3, large bubbles and/or cavities with size ranging from 15 to 50 nm are present inside the B region after annealing. Two bubble shapes can be distinguished, circular and faceted shapes with facets in the [)1 1 0 0] or/and [0 0 0 1] directions. We do not observe any dislocation loops in contrary to what was expected from a previous study on hot-pressed SiC [9]. These authors mentioned that bubbles are always accompanied by faulted dislocation loops of interstitial type. The implanted doses, however, were smaller than the one used in this study.
4. Modeling of structural modi®cations The room temperature helium bombardment induces atomic displacements, which result in ballistic relocations of atoms between neighboring monolayers [10]. For the quantitative analysis the solid is divided into monolayers with thickness h0
and the frequency probability of relocation of i K;L atom from K into monolayer L, wi , is introduced (i 1 for Si and i 2 for C atoms). The determination of the relocation function needs the solution of integro-dierential equations [11]. In the present work the relocation function is constructed from the following physical considerations. The number of displaced i atoms per 1 s in K K the K-th monolayer is equal to wi Di I0 =C, K where Di is the number of displaced i atoms per one incident ion in monolayer K; I0 is the ¯ux of incident ions and C is the number of atoms in one monolayer. The assumptions are made that the number of displaced atoms in the monolayer K, K Di , follows the Gaussian distribution of the nuclear energy losses, and the relocation probability of displaced atoms from monolayer K to L decreases with the increase in relocation distance 2 h0 jK Lj as exp h0 jK Lj=li , where li is the characteristic relocation distance of i atoms [12]. K;L It gives that wi can be written as " K;L wi
0 wi
exp
h0 jK li
h0 K
1 Rpd DRpd 2 # Lj ;
2
where DRpd and Rpd are the standard deviation and projected range of distribution of energy losses. 0 The parameter wi characterizes the displacement rate of i atoms in the monolayer corresponding to the maximum of nuclear energy losses (K 1 Rpd ) and are not relocated (jK Lj 0). It is function of the ¯ux and energy of incident ions and depends on the displacement energy of atoms. Computer simulation results show that the number of displaced Si atoms in SiC exceeds the number of displaced C atoms per one incident ion 0 0 about two times (w1 2=w2 ) [13]. Parameters 0 wi , DRpd and Rpd have been evaluated by SRIM: 0 0 w1 2 10 5 s 1 , w2 1 10 5 s 1 , Rpd 3:6 lm and DRpd 0:1 lm, other parameters used in calculations are l1 l2 h0 and h0 0:26 nm. The equation describing the time variations of the partial concentration of components in dierent monolayers initiated by the simultaneous ac-
M.F. Beaufort et al. / Nucl. Instr. and Meth. in Phys. Res. B 186 (2002) 218±222
tion of displacement and relocation processes can be written as 1 X dci L;K L wi ci dt L1 K
K
ci
1 X L 1
K;L
wi
:
The ®rst term de®nes the ¯ux of i atoms arriving into monolayer K from any other one, and the second term de®nes the ¯ux of i atoms leaving monolayer K into any other one including vacuum (negative indexes). Fig. 4 illustrates calculated concentration pro 0 ®les of atoms for two sets of parameter wi . It is seen that the area with maximum nuclear losses is depleted by atoms (c1 c2 < 1). The asymmetry in distribution pro®les manifests non-symmetrical distribution of nuclear energy losses. The best ®t s with experiments is obtained if DRpd 4DRpd , s where DRpd is the standard deviation of the nonsymmetrical Gaussian distribution pro®le from the front side. The displaced atoms occupy interstitial positions forming areas c1 c2 > 1. The coordinates corresponding to c1 c2 1 de®ne boundaries between areas enriched by vacancies and interstitials. Processes of annihilation and formation of extended defects are not considered in the present work. The calculated results are in qualitative agreement with TEM observations. The three areas can be distinguished: the central enriched by vacancies and two neighboring areas enriched by intersti-
221
tials. It is seen that the front layer A is wider (170 nm) than the near layer C (55 nm) and the density appears to be larger in the C layer than in A one. The above conducted analysis of the displacement±relocation processes gives complementary information about distribution pro®les of vacancies and interstitials. It can be concluded that nonhomogeneous displacement of atoms in dierent monolayers calculated for identical relocation probabilities initiates ¯uxes of atoms from areas of maximum nuclear losses to periphery forming areas enriched by vacancies and interstitials.
5. Conclusion The microstructures of high energy helium implanted 4H±SiC have been studied using conventional electron microscopy. When a sucient dose of helium is implanted into crystalline SiC sample, nucleation sites and bubbles are formed but the implantation also leads to amorphization. A three layer structure is observed: an amorphous one surrounded by crystalline zones with point defects. Bubbles evolve with recrystallization during thermal annealing at 1500 °C. A model based on the analysis of displacement and relocation processes of atoms in SiC gives a three layered structure as experimentally observed, enriched by vacancies in the area of maximum nuclear losses of incident ions and interstitials on the peripheries.
References
Fig. 4. Concentration pro®le of atoms after implantation for two couples of w0i values.
[1] D.M. Follstaedt, S.M. Myers, C.A. Petersen, J.W. Medernach, J. Electron. Mater. 25 (1996) 157. [2] V. Raineri, G. Fallica, S. Libertino, J. Appl. Phys. 79 (1996) 9012. [3] J.A. Spitznagel, S. Wood, W.J. Choyke, N.J. Doyle, J. Bradshaw, G. Fishman, Nucl. Instr. and Meth. B 16 (1996) 237. [4] E. Oliviero, M.L. David, M.F. Beaufort, J. Nomgaudyte, L. Pranevicius, A. Declemy, J.F. Barbot, J. Appl. Phys., to be published. [5] A. Declemy, E. Oliviero, M.F. Beaufort, J.F. Barbot, M.L. David, C. Blanchard, Y. Tessier, E. Ntsoenzok, Nucl. Instr. and Meth B 186 (2002) 318. [6] J.F. Ziegler, J.P. Biersack, U. Littmark, The Stopping and Range of Ions in Solids, Pergamon, New York, 1985.
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[7] M. Ishimaru, S. Harada, T. Motooka, T. Nakata, T. Yoneda, M. Inoue, Nucl. Instr. and Meth. B 127±128 (1997) 195. [8] V. Heera, J. Stoemenos, R. K ogler, W. Skorupa, J. Appl. Phys. 77 (1995) 2999. [9] J. Chen, P. Chung, H. Trinkaus, Phys. Rev. B 61 (2000) 12923.
[10] H.M. Urbassek, G. Mayer, H. Gades, M. Vicanek, Nucl. Instr. and Meth. B 103 (1995) 275. [11] P. Sigmund, A. Gras-Marti, Nucl. Instr. and Meth. B 182±183 (1981) 25. [12] P. Sigmund, A. Oliva, Nucl. Instr. and Meth. B 82 (1993) 269. [13] W. Jang, Nucl. Instr. and Meth. B 148 (1999) 557.