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In-situ transmission electron microscopy study of glissile grain boundary dislocation relaxation in a near Σ = 3 {1 1 1} grain boundary in copper
Materials Science and Engineering A 400–401 (2005) 264–267
In-situ transmission electron microscopy study of glissile grain boundary dislocation relaxation in a near Σ = 3 {1 1 1} grain boundary in copper J.P. Couzini´e a , B. D´ecamps b,∗ , L. Boulanger c , L. Priester a a b
Centre d’Etudes de Chimie M´etallurgique, CNRS, 15 rue Georges Urbain, 94407 Vitry Sur Seine, France Laboratoire de Chimie M´etallurgique des Terres Rares, CNRS, 2-8 rue Henri Dunant, 94320 Thiais, France c Commissariat a ` l’Energie Atomique, DTA/SRMP, 91191 Gif Sur Yvette, France Received 13 September 2004; received in revised form 6 December 2004; accepted 28 March 2005
Abstract An in-situ annealing experiment has been performed on an intergranular dislocation configuration composed only of glissile grain boundary dislocations observed in a near Σ = 3 {1 1 1} grain boundary in copper. Relaxation phenomena are not obvious than those predicted by theoretical models. Upon annealing, glissile intergranular dislocations are shown to overcome dislocation obstacles by node movement leading to a decrease of the total grain boundary energy. © 2005 Elsevier B.V. All rights reserved. Keywords: In-situ transmission electron microscopy; Grain boundary; Extrinsic grain boundary dislocations; Copper
1. Introduction
2. Experimental
During plastic deformation, lattice dislocations interact with grains boundaries (GBs). It is commonly assumed that, at low temperatures (T < 0.3Tm , with Tm the melting temperature), GBs constitute strong barriers to dislocation motion. At higher temperatures or high stresses, dislocations enter GBs forming extrinsic grain boundary dislocations (EGBDs) whose stress field accommodation is necessary for the plastic deformation to go on and for the GBs to return to equilibrium. The incorporation model of EGBDs by decomposition into sessile and glissile products followed by reorganization of the sessile dislocations in GBs is now relatively well-known [1–4]. On the contrary, glissile GBDs accommodation is still unclear as most of the time they are considered to rapidly glide towards the GBs extremities. The object of the present paper is to follow by in-situ transmission electron microscopy (TEM) the evolution and the relaxation processes of an intergranular configuration (analyzed in a previous paper [5]) in a near Σ = 3 GB in copper involving only glissile GBDs.
The near Σ = 3 {1 1 1} boundary is obtained by recrystallization of an initially deformed pure copper single crystal under annealing for 45 min at 813 K (0.6Tm ). After thinning the slices by standard polishing techniques, the specimen observation procedure is described as follows. The analyses of the GB parameters and of the initial intergranular configuration is performed in a JEOL 2000EX TEM. The in-situ anneal of the thin foil is achieved in a Philips CM20 equipped with a Megaview II camera and a VHS recorder.
∗
Corresponding author. Tel.: +33 1 49781214; fax: +33 1 49781203. E-mail address: [email protected] (B. D´ecamps).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.045
3. Results and interpretation 3.1. Description of the initial configuration The intergranular configuration in the near Σ = 3 {1 1 1} GB is given in Fig. 1a with its schematic representation in Fig. 1b. The GB is characterized by a misorientation angle of θ = 59.95◦ around [1 1 1]1,2 . The GB plane is close to (1 1 1) in both grains. The upper grain with respect to the electron source is crystal 2. The configuration has been completely analyzed in a recent paper [5] using a combination of TEM
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
265
Table 1 Dislocation characteristics of the initial configuration Dislocation
Burgers vector
Plane
Lattice A B
(a/2)[1¯ 0 1]2 (a/2)[1¯ 1 0]1
¯ 2 (1¯ 1 1) (1 1 1)1 , (0 0 1)1
GBDs C, E, G D, H, I2 F I1 J K
(a/6)[2¯ 1 1]1 ¯ 1 (a/6)[1¯ 2 1] (a/6)[1¯ 1¯ 2]1 (a/2)[0 1¯ 1]1 ¯ 1 (a/6)[2 1¯ 1] (a/6)[1 2¯ 1]1
= (a/6)[1¯ 1¯ 2]2 = (a/6)[2¯ 1 1]2 = (a/6)[1 2¯ 1]2 = (a/2)[1 1¯ 0]2 ¯ 2 = (a/6)[1 1 2] ¯ 2 = (a/6)[2 1¯ 1]
(1 1 1)1,2
techniques (invisibility criterion, Marukawa and Matsubara technique [6] and contrast simulation [7]). For convenience, Table 1 gives the characteristics of the different dislocations involved in the configuration. Dislocations A and B are located in grains 2 and 1, respectively, and the other dislocations (C–H, I1 , I2 , J, K) are in the GB. The TEM study [5] shows that all the interfacial dislocations are displacement symmetry conserving (DSC [8]) glissile GBDs, which result from the interaction between dissociated lattice dislocations and the twin boundary. For example, the lattice dislocation A decomposes into two DSC products C and D according to the following reaction (Fig. 1): (a/2)[1¯ 0 1]2 ⇒ (a/6)[1¯ 1¯ 2]2 + (a/6)[2¯ 1 1]2 A
Fig. 1. General view of the initial dislocation configuration in and near the GB: (a) simultaneous two-beam micrograph; (b) schematic representation of the configuration. Enlargement of the central part may be seen in the inset.
C
D
The possible mechanisms by which the initially dissociated lattice dislocation could give rise to two glissile products are discussed by Couzini´e et al. [5,9,10]. A recombination process of the Shockley partials seems to be
Fig. 2. Micrographs in double two beams diffraction condition taken from the video sequence showing the evolution of the configuration depicted in Fig. 1 upon annealing. Temperature and time are indicated as an insert in each micrograph. Schemes describing the evolution are given under each micrograph.
266
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
Fig. 2. (Continued ).
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
the more relevant as it is supported by atomistic simulations and theoretical models [11,12]. 3.2. Evolution upon in-situ annealing The micrographs presented in Fig. 2 are taken from the video sequence of the TEM in-situ annealing. The evolution can be described as follows: (a) Two dislocations coming from the left approach the configuration at T = 423 ◦ C (0.5Tm ) (Fig. 2a). Their rapid motion at relatively low temperature suggests that they are glissile GBDs. As a result, the configuration is shifted to the right indicating repulsive interaction between the defects, especially between the first dislocation and dislocation C belonging to the configuration. According to this repulsion and to the evolution that will follow, the two dislocations are assumed to be of D- and F-type, they are noted D and F , respectively. (b) Due to the configuration shift, dislocation B still located in the crystal elongates as the segment close to the surface is locked (Fig. 2b). (c) In Fig. 2c, the contrast change along dislocations I1 and I2 may be explained by a reaction between the two attractive defects to give rise to a F-type dislocation noted F . At the bottom part of the GB, the node N moves toward the surface. Due to these reactions, a complete F-type dislocation is lying in the twin from bottom to top. (d) The next step of the evolution is the emission of the F-type dislocation out of the configuration as shown in Fig. 2d. This emission requires that the F dislocation glides along the G and H dislocations by a rapid node movement. Such reaction allows the configuration to relax by a shift back with the decrease of the B dislocation length segment. (e) In Fig. 2e, a short segment of the B dislocation incorporates the GB and reacts with the attractive F and D dislocations giving rise to the schemes depicted at the bottom of Fig. 2e. (f) As for the F dislocation, a C-type dislocation is emitted on the right of the configuration by a rapid node movement. The resulting configuration is composed of the two A and B dislocations partially decomposed into DSC products but
267
not totally incorporated in the GB. It has to be noted that upon further annealing, the configuration disappears (around 0.7Tm ) by complex reactions of the same type as described above.
4. Conclusion On the contrary of what is theoretically proposed in the literature [1,2], the movement of intergranular glissile dislocations is not straightforward. Complex intergranular reactions (combination, annihilation, . . .), which occur have already been mentioned in the case of aluminum [13]. In our case, EGBD motion is activated starting from 0.5Tm but is impeded by dislocation configuration resulting from the interaction between lattice dislocations and the grain boundary [5]. Nevertheless, a node motion mechanism allows them to overcome the obstacles yielding a stress relaxation within the interface and the decrease of the intergranular energy.
References [1] R.Z. Valiev, V. Gertsman, O.A. Kaibyshev, Phys. Status Solidi A 78 (1) (1983) 177. [2] A.A. Nazarov, A.E. Romanov, R.Z. Valiev, Scripta Metall. Mater. 24 (10) (19901929). [3] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 79 (11) (1999) 2655. [4] M. Elkajbaji, J. Thibault-Desseaux, Philos. Mag. A 58 (2) (1988) 325. [5] J.P. Couzini´e, B. D´ecamps, L. Priester, Int. J. Plast. 21 (2005) 759. [6] K. Marukawa, Y. Matsubara, Trans. Jpn. Inst. Met. 20 (10) (1979) 560. [7] A.K. Head, P. Humble, L.M. Clarebrough, A.J. Morton, C.T. Forwood, Computed Electron Micrographs and Defect Identification, North-Holland, Amsterdam, Netherlands, 1973. [8] R.C. Pond, W. Bollmann, Philos. Trans. R. Soc. Lond. A 292 (1395) (1979) 449. [9] J.P. Couzini´e, B. D´ecamps, L. Priester, Philos. Mag. Lett. 83 (12) (2003) 721. [10] J.P. Couzini´e, B. D´ecamps, L. Priester, Z. Metallkd. 95 (2004) 223. [11] B.J. Pestman, J.T.M. de Hosson, V. Vitek, F.W. Schapink, Philos. Mag. A 64 (4) (1991) 951. [12] A.H. King, C. Fu Rong, Mater. Sci. Eng. 66 (2) (1984) 227. [13] C. Solenthaler, Mater. Sci. Eng. A A125 (1) (1990) 57.
In-situ transmission electron microscopy study of glissile grain boundary dislocation relaxation in a near Σ = 3 {1 1 1} grain boundary in copper J.P. Couzini´e a , B. D´ecamps b,∗ , L. Boulanger c , L. Priester a a b
Centre d’Etudes de Chimie M´etallurgique, CNRS, 15 rue Georges Urbain, 94407 Vitry Sur Seine, France Laboratoire de Chimie M´etallurgique des Terres Rares, CNRS, 2-8 rue Henri Dunant, 94320 Thiais, France c Commissariat a ` l’Energie Atomique, DTA/SRMP, 91191 Gif Sur Yvette, France Received 13 September 2004; received in revised form 6 December 2004; accepted 28 March 2005
Abstract An in-situ annealing experiment has been performed on an intergranular dislocation configuration composed only of glissile grain boundary dislocations observed in a near Σ = 3 {1 1 1} grain boundary in copper. Relaxation phenomena are not obvious than those predicted by theoretical models. Upon annealing, glissile intergranular dislocations are shown to overcome dislocation obstacles by node movement leading to a decrease of the total grain boundary energy. © 2005 Elsevier B.V. All rights reserved. Keywords: In-situ transmission electron microscopy; Grain boundary; Extrinsic grain boundary dislocations; Copper
1. Introduction
2. Experimental
During plastic deformation, lattice dislocations interact with grains boundaries (GBs). It is commonly assumed that, at low temperatures (T < 0.3Tm , with Tm the melting temperature), GBs constitute strong barriers to dislocation motion. At higher temperatures or high stresses, dislocations enter GBs forming extrinsic grain boundary dislocations (EGBDs) whose stress field accommodation is necessary for the plastic deformation to go on and for the GBs to return to equilibrium. The incorporation model of EGBDs by decomposition into sessile and glissile products followed by reorganization of the sessile dislocations in GBs is now relatively well-known [1–4]. On the contrary, glissile GBDs accommodation is still unclear as most of the time they are considered to rapidly glide towards the GBs extremities. The object of the present paper is to follow by in-situ transmission electron microscopy (TEM) the evolution and the relaxation processes of an intergranular configuration (analyzed in a previous paper [5]) in a near Σ = 3 GB in copper involving only glissile GBDs.
The near Σ = 3 {1 1 1} boundary is obtained by recrystallization of an initially deformed pure copper single crystal under annealing for 45 min at 813 K (0.6Tm ). After thinning the slices by standard polishing techniques, the specimen observation procedure is described as follows. The analyses of the GB parameters and of the initial intergranular configuration is performed in a JEOL 2000EX TEM. The in-situ anneal of the thin foil is achieved in a Philips CM20 equipped with a Megaview II camera and a VHS recorder.
∗
Corresponding author. Tel.: +33 1 49781214; fax: +33 1 49781203. E-mail address: [email protected] (B. D´ecamps).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.045
3. Results and interpretation 3.1. Description of the initial configuration The intergranular configuration in the near Σ = 3 {1 1 1} GB is given in Fig. 1a with its schematic representation in Fig. 1b. The GB is characterized by a misorientation angle of θ = 59.95◦ around [1 1 1]1,2 . The GB plane is close to (1 1 1) in both grains. The upper grain with respect to the electron source is crystal 2. The configuration has been completely analyzed in a recent paper [5] using a combination of TEM
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
265
Table 1 Dislocation characteristics of the initial configuration Dislocation
Burgers vector
Plane
Lattice A B
(a/2)[1¯ 0 1]2 (a/2)[1¯ 1 0]1
¯ 2 (1¯ 1 1) (1 1 1)1 , (0 0 1)1
GBDs C, E, G D, H, I2 F I1 J K
(a/6)[2¯ 1 1]1 ¯ 1 (a/6)[1¯ 2 1] (a/6)[1¯ 1¯ 2]1 (a/2)[0 1¯ 1]1 ¯ 1 (a/6)[2 1¯ 1] (a/6)[1 2¯ 1]1
= (a/6)[1¯ 1¯ 2]2 = (a/6)[2¯ 1 1]2 = (a/6)[1 2¯ 1]2 = (a/2)[1 1¯ 0]2 ¯ 2 = (a/6)[1 1 2] ¯ 2 = (a/6)[2 1¯ 1]
(1 1 1)1,2
techniques (invisibility criterion, Marukawa and Matsubara technique [6] and contrast simulation [7]). For convenience, Table 1 gives the characteristics of the different dislocations involved in the configuration. Dislocations A and B are located in grains 2 and 1, respectively, and the other dislocations (C–H, I1 , I2 , J, K) are in the GB. The TEM study [5] shows that all the interfacial dislocations are displacement symmetry conserving (DSC [8]) glissile GBDs, which result from the interaction between dissociated lattice dislocations and the twin boundary. For example, the lattice dislocation A decomposes into two DSC products C and D according to the following reaction (Fig. 1): (a/2)[1¯ 0 1]2 ⇒ (a/6)[1¯ 1¯ 2]2 + (a/6)[2¯ 1 1]2 A
Fig. 1. General view of the initial dislocation configuration in and near the GB: (a) simultaneous two-beam micrograph; (b) schematic representation of the configuration. Enlargement of the central part may be seen in the inset.
C
D
The possible mechanisms by which the initially dissociated lattice dislocation could give rise to two glissile products are discussed by Couzini´e et al. [5,9,10]. A recombination process of the Shockley partials seems to be
Fig. 2. Micrographs in double two beams diffraction condition taken from the video sequence showing the evolution of the configuration depicted in Fig. 1 upon annealing. Temperature and time are indicated as an insert in each micrograph. Schemes describing the evolution are given under each micrograph.
266
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
Fig. 2. (Continued ).
J.P. Couzini´e et al. / Materials Science and Engineering A 400–401 (2005) 264–267
the more relevant as it is supported by atomistic simulations and theoretical models [11,12]. 3.2. Evolution upon in-situ annealing The micrographs presented in Fig. 2 are taken from the video sequence of the TEM in-situ annealing. The evolution can be described as follows: (a) Two dislocations coming from the left approach the configuration at T = 423 ◦ C (0.5Tm ) (Fig. 2a). Their rapid motion at relatively low temperature suggests that they are glissile GBDs. As a result, the configuration is shifted to the right indicating repulsive interaction between the defects, especially between the first dislocation and dislocation C belonging to the configuration. According to this repulsion and to the evolution that will follow, the two dislocations are assumed to be of D- and F-type, they are noted D and F , respectively. (b) Due to the configuration shift, dislocation B still located in the crystal elongates as the segment close to the surface is locked (Fig. 2b). (c) In Fig. 2c, the contrast change along dislocations I1 and I2 may be explained by a reaction between the two attractive defects to give rise to a F-type dislocation noted F . At the bottom part of the GB, the node N moves toward the surface. Due to these reactions, a complete F-type dislocation is lying in the twin from bottom to top. (d) The next step of the evolution is the emission of the F-type dislocation out of the configuration as shown in Fig. 2d. This emission requires that the F dislocation glides along the G and H dislocations by a rapid node movement. Such reaction allows the configuration to relax by a shift back with the decrease of the B dislocation length segment. (e) In Fig. 2e, a short segment of the B dislocation incorporates the GB and reacts with the attractive F and D dislocations giving rise to the schemes depicted at the bottom of Fig. 2e. (f) As for the F dislocation, a C-type dislocation is emitted on the right of the configuration by a rapid node movement. The resulting configuration is composed of the two A and B dislocations partially decomposed into DSC products but
267
not totally incorporated in the GB. It has to be noted that upon further annealing, the configuration disappears (around 0.7Tm ) by complex reactions of the same type as described above.
4. Conclusion On the contrary of what is theoretically proposed in the literature [1,2], the movement of intergranular glissile dislocations is not straightforward. Complex intergranular reactions (combination, annihilation, . . .), which occur have already been mentioned in the case of aluminum [13]. In our case, EGBD motion is activated starting from 0.5Tm but is impeded by dislocation configuration resulting from the interaction between lattice dislocations and the grain boundary [5]. Nevertheless, a node motion mechanism allows them to overcome the obstacles yielding a stress relaxation within the interface and the decrease of the intergranular energy.
References [1] R.Z. Valiev, V. Gertsman, O.A. Kaibyshev, Phys. Status Solidi A 78 (1) (1983) 177. [2] A.A. Nazarov, A.E. Romanov, R.Z. Valiev, Scripta Metall. Mater. 24 (10) (19901929). [3] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 79 (11) (1999) 2655. [4] M. Elkajbaji, J. Thibault-Desseaux, Philos. Mag. A 58 (2) (1988) 325. [5] J.P. Couzini´e, B. D´ecamps, L. Priester, Int. J. Plast. 21 (2005) 759. [6] K. Marukawa, Y. Matsubara, Trans. Jpn. Inst. Met. 20 (10) (1979) 560. [7] A.K. Head, P. Humble, L.M. Clarebrough, A.J. Morton, C.T. Forwood, Computed Electron Micrographs and Defect Identification, North-Holland, Amsterdam, Netherlands, 1973. [8] R.C. Pond, W. Bollmann, Philos. Trans. R. Soc. Lond. A 292 (1395) (1979) 449. [9] J.P. Couzini´e, B. D´ecamps, L. Priester, Philos. Mag. Lett. 83 (12) (2003) 721. [10] J.P. Couzini´e, B. D´ecamps, L. Priester, Z. Metallkd. 95 (2004) 223. [11] B.J. Pestman, J.T.M. de Hosson, V. Vitek, F.W. Schapink, Philos. Mag. A 64 (4) (1991) 951. [12] A.H. King, C. Fu Rong, Mater. Sci. Eng. 66 (2) (1984) 227. [13] C. Solenthaler, Mater. Sci. Eng. A A125 (1) (1990) 57.