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TEM observations in Si: An attempt to link deformation microstructures and electrical activity
ARTICLE IN PRESS Physica B 404 (2009) 4634–4636
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
TEM observations in Si: An attempt to link deformation microstructures and electrical activity D. Eyidi a, V. Eremenko b, J.L. Demenet a, J. Rabier a, a b
PHYMAT, UMR 6630 CNRS-Universite de Poitiers, SP2MI, F-86960 Chasseneuil-Futuroscope, France Institute of Microelectronics Technology, RAS, 142432 Chernogolovka, Moscow District, Russia
a r t i c l e in f o
a b s t r a c t
Keywords: Silicon Dislocations TEM Deformation microstructures Extended defects
The deformation microstructures of Si single crystals have been studied by transmission electron microscopy (TEM) in the temperature range of deformation where ‘‘dislocation trails’’ are evidenced. The main features of the deformation microstructures are characterized by dislocations dipoles and their debris. These microscopic observations are discussed in relation with the mesoscopic properties of dislocation trails as well as the possible transformation between two different core structures of dislocations in silicon. & 2009 Elsevier B.V. All rights reserved.
1. Introduction Although silicon is considered as a model material and has been extensively studied, numerous questions still remain as far as the core structure of dislocations and their movement is concerned. As an example experimental measurements of dislocation velocities put forward the existence of point obstacles controlling the dislocation movement as well as the possible nucleation of defects of atomic dimension associated to dislocation glide [1]. In addition, dislocations which have been produced by plastic deformation in a temperature range of 480–750 1C close to the brittle to ductile transition have been shown to be associated with trails of extended defects (TEDs) located behind moving dislocations in the slip plane. Those TEDs, revealed by chemical etching [2], have been studied by a number of spectroscopic techniques giving puzzling results. Indeed they exhibit an electrical activity using electron beam induced current (EBIC), but this EBIC contrast of the defects in the dislocation trails has to be confronted to the absence of any EBIC contrast from the dislocations themselves [3]. The defects formed in the TEDs were assumed to be associated with agglomerates of intrinsic point defects formed during the dislocation motion. Dislocation-related DLTS signals have been also found when dislocations travelled a significant distance which strongly depends on dislocation velocity [4]. These DLTS signals were interpreted as resulting from some core defects and atomic impurities accumulated on the dislocations during their slow motion. The interpretations of these defect signatures interest several scales i.e. microscopic (dislocation core structure defects) and Corresponding author. Tel.: +33 5 49 49 67 32; fax: +33 5 49 49 66 92.
E-mail address: [email protected] (J. Rabier). 0921-4526/$ - see front matter & 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2009.08.134
mesoscopic (the collection of microscopic defects in trails in the wake of moving dislocations). It is of importance trying to relate those signatures to the actual dislocation microstructures of these deformed materials. As a matter of fact the deformation microstructures of silicon deformed in those conditions have not been investigated so much up to now. In an effort to shed some light on the relation between the spectroscopic measurements and the actual dislocation microstructure, TEM observations have been conducted on Si samples deformed in this range of temperature. It is also expected to obtain more insights into the fundamentals of plasticity of diamond structure materials close to the brittle to ductile transition (BDT).
2. Experimental Experiments were carried out on CZ grown Si single crystals: n- and p-type with doping levels (phosphorus and boron, respectively) ranging between 1013–1020 cm3 and FZ p-type crystals 1800 O cm. The concentrations of oxygen and carbon as the dominant electrically inactive impurities were of 1017– 1018 cm3 and of 1015–1016 cm3 respectively. Plastic deformation was performed by standard four-point bending of rectangular shaped chemically polished specimens 36 4 2 mm3 with corresponding edge orientations /11 0S, /1 0 0S and /11 0S at medium temperatures (480–750 1C) and stresses 20–60 MPa. Scratches as a source of dislocations were made on {1 0 0} surface for bending around /11 0S. These deformation conditions resulted in dislocation velocities being in the range of 102– 107 cm/s. After deformation, samples were cooled without stress at a rate of 10–20 1C/s. Then, the slices were cut after deformation parallel to the {111} glide planes and then both sides were
ARTICLE IN PRESS D. Eyidi et al. / Physica B 404 (2009) 4634–4636
4635
mechanically and chemically polished in the solution 1 HF:7HNO3 for preparing of TEM thin foils. TEM has been performed using a JEOL 200CX microscope as well as a FEG-TEM JEOL 2200-FS equipped with a O filter. Weak beam dark (WB) field imaging was used to obtain defect fine structure as well as to analyse the deformation microstructure at larger scales.
3. Experimental results Experimental results presented in this paper are related to CZ Si (r ¼ 3 O cm, p type, (3–5) 1017 cm3 oxygen). The analysis of the deformation microstructures has been performed for two deformation temperatures i.e. 570 and 620 1C. Fig. 1 shows deformation microstructure in a {111} glide plane after deformation at 570 1C. The microstructure is built with some individual dislocations but mainly with dislocation dipoles. Those dipoles exhibit different habit planes, likely the two variants for lowest energy configuration of dislocation dipoles. Dislocation loops are also seen which possess the habit planes of dipoles. Point-like contrasts are also seen along dislocation dipoles. At a
Fig. 3. A long dislocation dipole whose orientations are close to /11 0S orientations. The habit plane of the dipole depends on the dislocation line orientations (deformation temperature: 620 1C).
smaller scale faint contrasts are also evidenced which look like either small dislocation segments or very small dislocation loops. Fig. 2 shows a detail of the faint contrasts found in the vicinity of a dislocation dipole. The small loop-like objects can be seen in a quite regular aligned way or waving around. The structural objects associated to these faint contrasts have not been determined. It is thought that they correspond to dipole annihilation by-products in agreement with their aligned configurations. However their meandering configurations are not understood. For a deformation temperature of 620 1C, although the density of individual dislocations seems to be more important in the deformation substructure than after a deformation at 570 1C, dislocation dipoles remain the principal deformation microstructure characteristic. Long dislocation dipoles lie along the /11 0S orientations of the Peierls valleys. Fig. 3 shows such an example of a long dislocation dipole.
Fig. 1. Deformation microstructure in a {111} glide plane built with dislocation dipoles (D) and dislocation loops (DL). Faint contrasts are also seen (FC) (deformation temperature: 570 1C).
Fig. 2. WB (g, 3 g). Dislocation dipoles and faint contrasts (deformation temperature: 570 1C).
4. Discussion The striking feature of the deformation substructures shown in the preceding section is the presence of numerous dipoles and dipole debris. Some of the dipoles are imaged as faulted dipoles as already previously reported (see for example [5,6]). Dipole debris appear as long dipolar loops aligned along specific crystallographic directions and also in the form of small objects (loops or voidsy), keeping record of a dipole location or waving around dislocations. It comes out from these observations that dipoles or debris are candidates for being associated or related to the properties of trails. However it has to be emphasised that the scales at which properties are collected (etching and EBIC) are very different from those of TEM-EBIC and etching images exhibit trails which are distant of a few micrometers-so that dipoles and debris can only be associated with the fine structure of TEDs. In this hypothesis and to relate with the mesoscopic properties of TEDs, it is interesting to sort out what could be the etching and electrical properties of dipoles and their debris. Dislocation dipoles can be revealed by etching when the constituent dislocation lines intersect the etched surface which precludes any linear etched trails over large distances. As far as EBIC is concerned the signature has to be the same as those of parent dislocations which have been seen to have no EBIC signature.
ARTICLE IN PRESS 4636
D. Eyidi et al. / Physica B 404 (2009) 4634–4636
Classical dipole debris, i.e. interstitial or vacancy loops with different sizes should behave for these properties as parent dipoles. From these remarks, it is unlikely that dipoles and their classical debris can be seen as the elementary components of TEDs. However several computational works dealing with dipole annihilation have put forward that the classical dipole annihilation issues have to be revisited and that dipole debris could be in fact more complex than originally thought. Indeed dipole annihilation products, beside those relevant to screw orientations have been found not to restore the perfect crystal [7]. For example in some materials voids can result from vacancy dipole ‘‘annihilation’’ [8]. This has also to be put together with the work of Hounsome et al. [9] who attributed the brown coloration of natural diamond to disks of vacancies which are stable up to 200 vacancies as compared to dislocation loops. Such voids of vacancy clusters do not possess any deformation field and are consequently hardly visible using TEM. Annihilation of interstitial dipoles can be also associated to debris like interstitial clusters or {11 3} defects, but these defects are then associated with deformation strain field. Such defects have not been evidenced in this study. Another point that has to be addressed: is the observed dislocation dipole density relevant to a standard dislocation mechanism common to a lot of different materials or a specificity relevant to the dislocation core structures of silicon? Indeed two types of dislocation cores have been found on both sides of the BDT, i.e. a perfect ‘‘shuffle core’’ at high stress and low temperature and a dissociated ‘‘glide core’’ at medium temperature and low stress [10,11]. Those two dislocation cores have been found to behave quite differently: perfect ‘‘shuffle core’’ dislocations are impossible to move at temperatures below the BDT [12] once they have been stopped. This seems to be corroborated by recent computations [13] showing that their lower energy configurations are sessile. Then the hypothesis of perfect shuffle dislocations pre-existing but being sessile when submitted to moderate stresses can be followed. In that context, the nucleation of glide dislocations can be facilitated in the stress field of pre existing sessile dislocations in a {111} glide plane a few atomic distances from that of the parent dislocation. Dipoles and debris
will result of such a mechanism. These by-products can provide nucleation sites of jog formation on mobile dislocations, interacting with them, feeding the formation of new dipoles.
5. Conclusion Dislocation dipoles and debris have been found to build the deformation microstructure of plastically deformed silicon in the temperature and stress range where TEDs are exhibited. TEDs could result from dipole debris under the form of stable and highly re-arranged configurations. This has to be checked through atomistic computations. STEBIC (EBIC in the TEM) imaging of constitutive defects of TEDs is also desirable to allow for a better understanding of physics of dislocations in silicon.
Acknowledgements The authors acknowledge Dr P. Veyssie re (LEM CNRS-ONERA) for stimulating discussions. This work has been performed in the frame of a CNRS/ASR cooperation (PECO-NEI project 19045).
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
R. Jones, MSE B71 (2000) 24. V. Eremenko, E. Yakimov, N. Abrosimov, PSS(C) 4 (2007) 3100. V. Eremenko, E. Yakimov, Eur. Phys. J. Appl. Phys. 27 (2004) 349. V. Kveder, V. Orlov, M. Khorosheva, M. Seibt, Solid State Phenom. 131–133 (2008) 175. C.B. Carter, Phil. Mag. 36 (1977) 147. A.T. Winter, S. Mahajan, D. Brasen, Acta Metall. 11 (1977) 1157. H. Wang, D.S. Xu, R. Yang, P. Veyssie re, Acta Materialia 57 (2009) 3725. J. Rabier, M.P. Puls, Phil. Mag. A 59 (1989) 821. L.S. Hounsome, R. Jones, P.M. Martineau, D. Fisher, M.J. Shaw, P.R. Briddon, ¨ berg, PSS(C) 4 (2007) 2950. S. O J. Rabier, P. Cordier, J.L. Demenet, H. Garem, MSE A309–310 (2001) 74. J. Rabier, J. Pizzagalli, J.L. Demenet, in: J.P. Hirth, L. Kubin (Eds.), Dislocations in Solids, Elsevier, 2010, to appear (Chapter 93). J. Rabier, J.L. Demenet, PSS(A) 202 (2005) 944. L. Pizzagalli, J. Godet, S. Brochard, PRL 103 (2009) 065505.
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
TEM observations in Si: An attempt to link deformation microstructures and electrical activity D. Eyidi a, V. Eremenko b, J.L. Demenet a, J. Rabier a, a b
PHYMAT, UMR 6630 CNRS-Universite de Poitiers, SP2MI, F-86960 Chasseneuil-Futuroscope, France Institute of Microelectronics Technology, RAS, 142432 Chernogolovka, Moscow District, Russia
a r t i c l e in f o
a b s t r a c t
Keywords: Silicon Dislocations TEM Deformation microstructures Extended defects
The deformation microstructures of Si single crystals have been studied by transmission electron microscopy (TEM) in the temperature range of deformation where ‘‘dislocation trails’’ are evidenced. The main features of the deformation microstructures are characterized by dislocations dipoles and their debris. These microscopic observations are discussed in relation with the mesoscopic properties of dislocation trails as well as the possible transformation between two different core structures of dislocations in silicon. & 2009 Elsevier B.V. All rights reserved.
1. Introduction Although silicon is considered as a model material and has been extensively studied, numerous questions still remain as far as the core structure of dislocations and their movement is concerned. As an example experimental measurements of dislocation velocities put forward the existence of point obstacles controlling the dislocation movement as well as the possible nucleation of defects of atomic dimension associated to dislocation glide [1]. In addition, dislocations which have been produced by plastic deformation in a temperature range of 480–750 1C close to the brittle to ductile transition have been shown to be associated with trails of extended defects (TEDs) located behind moving dislocations in the slip plane. Those TEDs, revealed by chemical etching [2], have been studied by a number of spectroscopic techniques giving puzzling results. Indeed they exhibit an electrical activity using electron beam induced current (EBIC), but this EBIC contrast of the defects in the dislocation trails has to be confronted to the absence of any EBIC contrast from the dislocations themselves [3]. The defects formed in the TEDs were assumed to be associated with agglomerates of intrinsic point defects formed during the dislocation motion. Dislocation-related DLTS signals have been also found when dislocations travelled a significant distance which strongly depends on dislocation velocity [4]. These DLTS signals were interpreted as resulting from some core defects and atomic impurities accumulated on the dislocations during their slow motion. The interpretations of these defect signatures interest several scales i.e. microscopic (dislocation core structure defects) and Corresponding author. Tel.: +33 5 49 49 67 32; fax: +33 5 49 49 66 92.
E-mail address: [email protected] (J. Rabier). 0921-4526/$ - see front matter & 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2009.08.134
mesoscopic (the collection of microscopic defects in trails in the wake of moving dislocations). It is of importance trying to relate those signatures to the actual dislocation microstructures of these deformed materials. As a matter of fact the deformation microstructures of silicon deformed in those conditions have not been investigated so much up to now. In an effort to shed some light on the relation between the spectroscopic measurements and the actual dislocation microstructure, TEM observations have been conducted on Si samples deformed in this range of temperature. It is also expected to obtain more insights into the fundamentals of plasticity of diamond structure materials close to the brittle to ductile transition (BDT).
2. Experimental Experiments were carried out on CZ grown Si single crystals: n- and p-type with doping levels (phosphorus and boron, respectively) ranging between 1013–1020 cm3 and FZ p-type crystals 1800 O cm. The concentrations of oxygen and carbon as the dominant electrically inactive impurities were of 1017– 1018 cm3 and of 1015–1016 cm3 respectively. Plastic deformation was performed by standard four-point bending of rectangular shaped chemically polished specimens 36 4 2 mm3 with corresponding edge orientations /11 0S, /1 0 0S and /11 0S at medium temperatures (480–750 1C) and stresses 20–60 MPa. Scratches as a source of dislocations were made on {1 0 0} surface for bending around /11 0S. These deformation conditions resulted in dislocation velocities being in the range of 102– 107 cm/s. After deformation, samples were cooled without stress at a rate of 10–20 1C/s. Then, the slices were cut after deformation parallel to the {111} glide planes and then both sides were
ARTICLE IN PRESS D. Eyidi et al. / Physica B 404 (2009) 4634–4636
4635
mechanically and chemically polished in the solution 1 HF:7HNO3 for preparing of TEM thin foils. TEM has been performed using a JEOL 200CX microscope as well as a FEG-TEM JEOL 2200-FS equipped with a O filter. Weak beam dark (WB) field imaging was used to obtain defect fine structure as well as to analyse the deformation microstructure at larger scales.
3. Experimental results Experimental results presented in this paper are related to CZ Si (r ¼ 3 O cm, p type, (3–5) 1017 cm3 oxygen). The analysis of the deformation microstructures has been performed for two deformation temperatures i.e. 570 and 620 1C. Fig. 1 shows deformation microstructure in a {111} glide plane after deformation at 570 1C. The microstructure is built with some individual dislocations but mainly with dislocation dipoles. Those dipoles exhibit different habit planes, likely the two variants for lowest energy configuration of dislocation dipoles. Dislocation loops are also seen which possess the habit planes of dipoles. Point-like contrasts are also seen along dislocation dipoles. At a
Fig. 3. A long dislocation dipole whose orientations are close to /11 0S orientations. The habit plane of the dipole depends on the dislocation line orientations (deformation temperature: 620 1C).
smaller scale faint contrasts are also evidenced which look like either small dislocation segments or very small dislocation loops. Fig. 2 shows a detail of the faint contrasts found in the vicinity of a dislocation dipole. The small loop-like objects can be seen in a quite regular aligned way or waving around. The structural objects associated to these faint contrasts have not been determined. It is thought that they correspond to dipole annihilation by-products in agreement with their aligned configurations. However their meandering configurations are not understood. For a deformation temperature of 620 1C, although the density of individual dislocations seems to be more important in the deformation substructure than after a deformation at 570 1C, dislocation dipoles remain the principal deformation microstructure characteristic. Long dislocation dipoles lie along the /11 0S orientations of the Peierls valleys. Fig. 3 shows such an example of a long dislocation dipole.
Fig. 1. Deformation microstructure in a {111} glide plane built with dislocation dipoles (D) and dislocation loops (DL). Faint contrasts are also seen (FC) (deformation temperature: 570 1C).
Fig. 2. WB (g, 3 g). Dislocation dipoles and faint contrasts (deformation temperature: 570 1C).
4. Discussion The striking feature of the deformation substructures shown in the preceding section is the presence of numerous dipoles and dipole debris. Some of the dipoles are imaged as faulted dipoles as already previously reported (see for example [5,6]). Dipole debris appear as long dipolar loops aligned along specific crystallographic directions and also in the form of small objects (loops or voidsy), keeping record of a dipole location or waving around dislocations. It comes out from these observations that dipoles or debris are candidates for being associated or related to the properties of trails. However it has to be emphasised that the scales at which properties are collected (etching and EBIC) are very different from those of TEM-EBIC and etching images exhibit trails which are distant of a few micrometers-so that dipoles and debris can only be associated with the fine structure of TEDs. In this hypothesis and to relate with the mesoscopic properties of TEDs, it is interesting to sort out what could be the etching and electrical properties of dipoles and their debris. Dislocation dipoles can be revealed by etching when the constituent dislocation lines intersect the etched surface which precludes any linear etched trails over large distances. As far as EBIC is concerned the signature has to be the same as those of parent dislocations which have been seen to have no EBIC signature.
ARTICLE IN PRESS 4636
D. Eyidi et al. / Physica B 404 (2009) 4634–4636
Classical dipole debris, i.e. interstitial or vacancy loops with different sizes should behave for these properties as parent dipoles. From these remarks, it is unlikely that dipoles and their classical debris can be seen as the elementary components of TEDs. However several computational works dealing with dipole annihilation have put forward that the classical dipole annihilation issues have to be revisited and that dipole debris could be in fact more complex than originally thought. Indeed dipole annihilation products, beside those relevant to screw orientations have been found not to restore the perfect crystal [7]. For example in some materials voids can result from vacancy dipole ‘‘annihilation’’ [8]. This has also to be put together with the work of Hounsome et al. [9] who attributed the brown coloration of natural diamond to disks of vacancies which are stable up to 200 vacancies as compared to dislocation loops. Such voids of vacancy clusters do not possess any deformation field and are consequently hardly visible using TEM. Annihilation of interstitial dipoles can be also associated to debris like interstitial clusters or {11 3} defects, but these defects are then associated with deformation strain field. Such defects have not been evidenced in this study. Another point that has to be addressed: is the observed dislocation dipole density relevant to a standard dislocation mechanism common to a lot of different materials or a specificity relevant to the dislocation core structures of silicon? Indeed two types of dislocation cores have been found on both sides of the BDT, i.e. a perfect ‘‘shuffle core’’ at high stress and low temperature and a dissociated ‘‘glide core’’ at medium temperature and low stress [10,11]. Those two dislocation cores have been found to behave quite differently: perfect ‘‘shuffle core’’ dislocations are impossible to move at temperatures below the BDT [12] once they have been stopped. This seems to be corroborated by recent computations [13] showing that their lower energy configurations are sessile. Then the hypothesis of perfect shuffle dislocations pre-existing but being sessile when submitted to moderate stresses can be followed. In that context, the nucleation of glide dislocations can be facilitated in the stress field of pre existing sessile dislocations in a {111} glide plane a few atomic distances from that of the parent dislocation. Dipoles and debris
will result of such a mechanism. These by-products can provide nucleation sites of jog formation on mobile dislocations, interacting with them, feeding the formation of new dipoles.
5. Conclusion Dislocation dipoles and debris have been found to build the deformation microstructure of plastically deformed silicon in the temperature and stress range where TEDs are exhibited. TEDs could result from dipole debris under the form of stable and highly re-arranged configurations. This has to be checked through atomistic computations. STEBIC (EBIC in the TEM) imaging of constitutive defects of TEDs is also desirable to allow for a better understanding of physics of dislocations in silicon.
Acknowledgements The authors acknowledge Dr P. Veyssie re (LEM CNRS-ONERA) for stimulating discussions. This work has been performed in the frame of a CNRS/ASR cooperation (PECO-NEI project 19045).
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
R. Jones, MSE B71 (2000) 24. V. Eremenko, E. Yakimov, N. Abrosimov, PSS(C) 4 (2007) 3100. V. Eremenko, E. Yakimov, Eur. Phys. J. Appl. Phys. 27 (2004) 349. V. Kveder, V. Orlov, M. Khorosheva, M. Seibt, Solid State Phenom. 131–133 (2008) 175. C.B. Carter, Phil. Mag. 36 (1977) 147. A.T. Winter, S. Mahajan, D. Brasen, Acta Metall. 11 (1977) 1157. H. Wang, D.S. Xu, R. Yang, P. Veyssie re, Acta Materialia 57 (2009) 3725. J. Rabier, M.P. Puls, Phil. Mag. A 59 (1989) 821. L.S. Hounsome, R. Jones, P.M. Martineau, D. Fisher, M.J. Shaw, P.R. Briddon, ¨ berg, PSS(C) 4 (2007) 2950. S. O J. Rabier, P. Cordier, J.L. Demenet, H. Garem, MSE A309–310 (2001) 74. J. Rabier, J. Pizzagalli, J.L. Demenet, in: J.P. Hirth, L. Kubin (Eds.), Dislocations in Solids, Elsevier, 2010, to appear (Chapter 93). J. Rabier, J.L. Demenet, PSS(A) 202 (2005) 944. L. Pizzagalli, J. Godet, S. Brochard, PRL 103 (2009) 065505.