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Extensive disordering in long-range-ordered Cu3Au induced by severe plastic deformation studied by transmission electron microscopy
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Acta Materialia 56 (2008) 2526–2530 www.elsevier.com/locate/actamat
Extensive disordering in long-range-ordered Cu3Au induced by severe plastic deformation studied by transmission electron microscopy C. Rentenberger *, H.P. Karnthaler Physics of Nanostructured Materials, Faculty of Physics, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria Received 8 May 2007; received in revised form 14 January 2008; accepted 26 January 2008 Available online 5 March 2008
Abstract Bulk nanocrystalline materials can be made by severe plastic deformation. In L12 long-range-ordered alloys, this leads to extensive disordering which influences the highly improved properties of these nanocrystalline alloys. Transmission electron microscopy methods were applied to Cu3Au; both diffraction contrast images and diffraction patterns reveal that disordering takes place locally. It is concluded that in addition to disordering by the refinement of the grown-in antiphase boundary domains, the formation of antiphase boundary tubes is a prominent process of disordering. The latter is facilitated by the fact that, unlike dislocations, antiphase boundary tubes can be stored at a very high density without causing long-range stresses. The local disordering indicates that the nanocrystalline structure nucleates inhomogeneously in the highly strained disordered regions. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Transmission electron microscopy; Nanocrystalline alloys; Order–disorder phenomena; Antiphase boundary tubes
1. Introduction The extensive disordering of L12 long-range ordered (LRO) alloys by severe plastic deformation (SPD) is of special interest, as the highly improved mechanical properties of these alloys after SPD are strongly influenced by their disorder. This was shown, for example, for Ni3Al, which is nanocrystalline and disordered after SPD by high-pressure torsion (HPT) [1]. The process of disordering of LRO alloys can be studied by X-ray and neutron diffraction monitoring of the superlattice reflections. These methods yield a global average of the LRO parameter, e.g., as carried out for Ni3Al [2,3]. For the analysis of the local variations in order and the mechanism causing the disordering, transmission electron microscopy (TEM) methods based on both electron diffraction and direct imaging are appropriate and were used in the present study. In the literature, only a few TEM reports dealing with deforma*
Corresponding author. Tel.: +43 1427751314; fax: +43 1427751316. E-mail address: [email protected] (C. Rentenberger).
tion-induced localized destruction of order are found, e.g., for Ni3Al [2,4,5] and Cu3Au [6]. It is the aim of this study to analyse the process of disordering after SPD. Various methods of SPD receive currently widespread interest, as they have been applied to make bulk materials with nanocrystalline structures [7]. In the present case, L12LRO Cu3Au was deformed by HPT, as different strains (up to 5000%) can be achieved in one sample by this method. 2. Experimental procedure Polycrystalline Cu3Au samples with a grain size of 0.2 mm were annealed at 350 °C for 140 h followed by a step cooling treatment (10 K day–1 down to 160 °C) to get a fully LRO alloy with an order parameter of 1 and with domains of diameter 0.5 lm. A sample (8 mm in diameter and 0.8 mm thick) was torsion strained under a quasi-hydrostatic pressure of 4 GPa by applying 2.5 turns at room temperature (RT). The design of the HPT tool was reported in Ref. [8].
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.01.035
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For the TEM investigation, discs were cut out by spark erosion from different parts of the HPT sample, representing different levels of strain. Sample C from the central area and sample P near the periphery correspond to shear strains c = 15 ± 5 and 50 ± 10, respectively. The discs were used to prepare TEM foils by a two-step procedure: first, they were dimpled, and afterwards they were electropolished [9]. It should be mentioned that, during electropolishing of Cu3Au, redeposition of Au on the surface of the TEM foils occurs. The TEM study was carried out with an acceleration voltage of 200 kV. Diffraction contrast images (using both fundamental and superlattice reflections) and selected area diffraction (SAD) patterns were used to analyse the microstructure and the local disordering. 3. Experimental results Fig. 1 shows a SAD pattern taken from specimen P and the corresponding intensity profile obtained by integrating rings of constant magnitude of the diffraction vector; an illuminated area 30 lm2 was used. The ring-like SAD pattern indicates a large variety of grain orientations and shows some texture. Segments of a diffuse ring observed inside the h1 1 1i ring are caused by gold redeposited during preparation; this leads to a little hump (marked Au) near the h1 1 1i-peak in the intensity profile. The comparison of the positions of the peaks with calculated ones reveals that only peaks corresponding to the fcc structure are observed, they are labelled by their indices in the intensity profile. Superlattice reflections of the L12 structure (indicated by dashed lines) are missing, indicating a LRO parameter <0.1. Therefore, the HPT deformation (c = 50 ± 10) leads to extensive disordering of Cu3Au. Fig. 2 shows dark-field images of specimen C using the superlattice reflection g ¼ ½1 1 0 (beam direction BD [1 1 0]). Two sets of faint dark lines aligned along the projections of ½0 1 1 and ½ 1 0 1 (indicated by dashed lines) are imaged in Fig. 2a. The width of the individual contrast lines is only 1–2 nm. It should be mentioned that the lines are not observed using a fundamental reflection, indicating that their contrast is caused by antiphase boundary (APB) faults. In addition, wavy contrast features are observed on the nanometre scale; they are attributed to the boundaries of the grown-in APB domains fragmented by the glide of superlattice dislocations. Fig. 2b shows an area with a high density of lines; their contrast sums up to dark bands aligned along the projection of glide planes (as indicated by the dashed line). To analyse the straight lines, a SAD pattern was taken of a small area containing a large number of them (BD [1 1 1]). The SAD pattern (cf., Fig. 3a) shows fundamental and superlattice reflections of the L12 ordered structure. In addition, streaks aligned along a direction perpendicular to V ¼ ½ 1 0 1 occur near superlattice spots only. Diffuse reflections (some are marked by arrows) are caused by gold redeposited on the surfaces. To analyse the SAD pattern, a
Fig. 1. SAD pattern and its intensity profile obtained from L12 ordered Cu3Au after SPD (shear strain 50 ± 10). The positions of the observed intensity maxima and the lack of maxima corresponding to superlattice reflections (indicated by dashed vertical lines) show that no detectable LRO is present. The hump near the {1 1 1} intensity peak is caused by a gold layer redeposited on the surface during the preparation of the TEM foil.
sketch of a sector (indicated by the two dashed lines) is shown in Fig. 3b. Circles correspond to reflections of the zero-order Laue zone (ZOLZ), whereas squares and triangles indicate reflections of the neighbouring Laue zones (+1 and –1). Full (m odd) and open (m even) symbols belong to parallel reciprocal lattice planes perpendicular to V. Only the reflections indicated by full symbols show streaks. For reflections of the ZOLZ (full circles), the streaks lie in this Laue zone. At the positions of reflections of neighbouring Laue zones, the weak contrast is caused by an intensity expanding perpendicular to the Laue zones. Reflections of neighbouring Laue zones can be moved into the ZOLZ by tilting, e.g., when BD = [2 2 1], reflection ½1 0 2 belongs to the ZOLZ and shows pronounced streaks in the plane of the projection. From the results obtained
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Fig. 2. TEM dark-field images of HPT-deformed L12 ordered Cu3Au (shear strain <20) using the superlattice reflection g ¼ ½1 0 1 (BD[1 1 0]). (a) Faint contrast lines aligned along two different h1 1 0i directions (indicated by dashed lines) are interpreted as APB tubes. The wavy contrast variations are caused by fragmented domain walls. (b) The contrast of several closely spaced APB tubes can sum up to broad bands aligned along the projection of h1 1 0i (indicated by a dashed line).
using different BD vectors, it is concluded that each superlattice reflection belonging to reciprocal lattice planes with odd m values has a disc-like shape extended in ±[1 1 1] directions. This intensity distribution of reflections with m = 1 is sketched in Fig. 3c for the ZOLZ and for the first-order Laue zone (±FOLZ). Table 1 shows the phase factor a = 2pg R for different values of g when a displacement vector R ¼ a=2½ 1 0 1 parallel to V and parallel to the defect line is assumed. R is taken to correspond to the Burgers vector of the superpartial dislocations which, together with the APB fault ribbon, form the superlattice glide dislocations. The value of a is constant and equal to pm for the different planes perpendicular to V, where m increases with increasing distance from the centre, as indicated in Fig. 3b. The occurrence of streaks agrees with the visibility criterion a = p (modulo 2p) for APB faults when R ¼ a=2½ 1 0 1 is used. In addition, streaks perpendicular
Fig. 3. Analysis of streaks observed in a SAD pattern obtained from an area showing a high density of parallel straight defects (BD [1 1 1]). (a) Experimental SAD pattern; in addition to the spot pattern, streaks aligned perpendicular to V ¼ ½1 0 1 are observed at some superlattice reflections of the L12 structure. (Some spots of the redeposited Au layer are marked.) (b) Schematic sketch of a sector of the experimental SAD pattern (indicated by dashed lines in (a)). Large circles correspond to reflections in the ZOLZ, whereas squares and triangles indicate reflections of the first-order (±FOLZ) Laue zones. Streaks (perpendicular to V) are observed only near superlattice reflections lying in planes with odd m values. (c) Schematic sketch of the intensity distribution of superlattice spots belonging to plane m = 1 and to different Laue zones.
to other h1 1 0i directions are observed, e.g., faint streaks running perpendicular to ½0 1 1 in the SAD pattern correspond to defects aligned along ±½0 1 1.
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Table 1 Values of the phase factor a = 2pgR for some reflections of different reciprocal planes m (cf., Fig. 3b) when a displacement vector R ¼ a=2½1 0 1 parallel to the defect line is assumed 1 0 0 1 g ½1 ½1 ½2 2 0 ½1 1 2 ½1 0 2 ½2 0 1 ½2 1 1
vacancies at the melting temperature [13,14]; still the reduction of the LRO by vacancies will be <1%.
m a = 2pgR
The present TEM results obtained by comparing images taken with fundamental reflections with those of superlattice reflections (cf., Fig. 2) and by the analysis of streaks occurring in the diffraction patterns (cf., Fig. 3) agree with the results expected for individual APB tubes as well as APB tubes accumulated in bands. A high density of APB tubes leads to a local and crystallographically related destruction of the LRO order and can occur without the storage of glide dislocations. Therefore, the stress field associated with APB tubes is negligible, and they can be stored at a high density without causing long-range stresses [15]. APB tubes can be formed by superlattice screw dislocations (e.g., by annihilation and by double cross-slip [16]) and by the movement of superlattice dislocations with nonaligned jogs [17]. The length of the streaks in the SAD pattern (expanding to the next Laue zone) indicates that the dimension of the tubes in ±[1 1 1] is very small (three atomic (1 1 1) planes or less). This result favours mechanisms that lead to the formation of tubes of small height; like the one by double cross-slip and that by movement of non-aligned jogs.
1 p
2 2p
2 2p
3 3p
3 3p
3 3p
3 3p
4. Discussion In several structural intermetallics, the analysis of the nanocrystals formed by SPD shows that they are disordered. This was observed in L12 ordered alloys by TEM methods (e.g., in Cu3Au and in Ni3Al deformed by HPT [5,6]) and in a B2 ordered alloy by X-ray diffraction (FeAl deformed by ball milling [10]). It is therefore proposed that the formation of a nanocrystalline structure made by SPD is correlated with disordering. Based on the present TEM results of SPD deformed Cu3Au, the following disordering mechanisms leading to a deformation-induced order–disorder transition are discussed. 4.1. Disordering by APB fault ribbons of dissociated dislocations and by point defects Deformation in L12 ordered alloys proceeds by the glide of superlattice dislocations (Burgers vector b = ah1 1 0i) [11]. They contain ribbons of APB faults reducing the local order. Intense shearing leads to storage and accumulation of dislocations and finally to a refined grain structure [12]. To estimate the contribution of APB fault ribbons of dissociated superlattice dislocations to the disordering process, calculations of the long-range order parameter were carried out. They show that, in the case of Cu3Au, a very high superlattice dislocation density of 2 1017 m–2 would be needed to decrease the order by 50%, the complete loss of order would require dislocation densities near the stability limit of the crystal (for this calculation a dissociation width of 11 nm was used, as deduced from weak-beam images of edge dislocations). It should be noted that, in contrast to the present observations, APB fault ribbons bounded by dissociated dislocations should show a contrast in the image taken, with the fundamental reflection corresponding to the superlattice reflection in Fig. 2. Also the length of the streaks in the SAD pattern caused by superlattice screw dislocations would be different from the observed one. Assuming a dissociation width of 7 nm on the (1 1 1) plane for screw dislocations lying parallel to ½ 1 0 1, the length of the streaks of the reflections in the ZOLZ would be only 1/7 of the observed one. In addition to APB faults, point defects formed during deformation can also contribute to disordering. It is reported that, after SPD, the density of deformation-induced vacancies is at RT as high as the density of thermal
4.2. Disordering by copious formation of APB tubes
4.3. Disordering by the refinement of the grown-in APB domains Gliding superlattice dislocations intersect domain walls and form new APB faults [18]. The refinement of domains can be identified by TEM methods using superlattice reflections (cf., Fig. 2a). Measurements of the reduction of the domain size with progressing deformation were carried out in L12 ordered Cu–22 at.% Pt using X-ray diffraction [19]. 4.4. Comparison with nanocrystalline Ni3Al In the case of Ni3Al that remains fully ordered up to the melting point, only the first two disordering mechanisms are possible, as no grown-in APB domains are present. Furthermore, disordering by APB fault ribbons bounded by dissociated dislocations is not, so effective as in Cu3Au, as their dissociation widths are 40% smaller in Ni3Al. Therefore, it can be concluded that, in coarse-grained Ni3Al, the local strong reduction of the long-range order [5] occurs mainly by the formation of APB tubes. The loss of LRO in extended areas takes place in Ni3Al in the nanocrystalline structure only, whereas in Cu3Au, fragmented grains are already strongly disordered [5,6]. The local correlation between disorder and nanocrystallization found by TEM methods seems to indicate that disordering is a prerequisite for nanocrystallization and that the nanocrystals nucleate inhomogeneously in the highly strained disordered regions.
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The highly improved ductility of nanocrystalline Ni3Al [1] is supported by the presence of disordering in the nanocrystalline structures of alloys that are L12-LRO in coarse grains. The reason is that the mobility of the screws is much higher in the disordered structure than in the ordered one, where they show locked configurations due to their nonplanar dislocation core structure [20]. In addition, the reduction in the length of the Burgers vector in the disordered structure (b = a/2h1 1 0i) as compared with the ordered one (b = ah1 1 0i) by a factor of 2 reduces the stress to activate glide dislocations accordingly. Furthermore, in the disordered nanocrystalline structure, the occurrence of twinning by the emission of Shockley partials (b = a/ 6h1 1 2i) from grain boundaries can contribute to ductility [21,22]. The increase in the strength in nanocrystalline alloys is caused by the reduction in grain size. In straininduced disordered nanocrystalline alloys (e.g., Ni3Al), the strength is increased even more when the deformation is carried out at an elevated temperature at which reordering takes place during deformation [1,23]. 5. Conclusions
(1) SPD by HPT of L12-LRO Cu3Au leads to extensive disordering and finally to the complete loss of order. (2) TEM shows that disordering starts locally along glide planes and is associated with the deformation-induced generation of APB. From dark-field images using superlattice reflections and diffraction patterns, it is concluded that extensive disordering is caused predominantly by the storage of APB tubes and of APB formed by intersecting grownin domains with dislocations. Contrary to APB of dissociated dislocations, APB tubes can be stored at a very high density without causing long-range stresses. For L12 ordered alloys that remain fully ordered up to the melting point (e.g., Ni3Al), it is concluded that the formation of APB tubes is the decisive process. (3) Based on the TEM results, it is proposed that in structural intermetallics (Cu3Au, Ni3Al) disordering occurring locally is a prerequisite for the formation of a nanocrystalline structure made by SPD and that the nanocrystals nucleate in the highly strained disordered regions. (4) The highly improved mechanical properties of bulk nanocrystalline alloys made by SPD are influenced by extensive disordering, leading to a higher mobility
of screw dislocations, as in the ordered structure screw dislocations show locked configurations due to their non-planar dislocation core structure.
Acknowledgments The authors thank Dr. A. Vorhauer (Erich Schmid Institute Leoben, Austria) for his kind help with the HPT deformation. The authors acknowledge the support by the ‘Bulk Nanostructured Materials’ research project within the ‘Materials Science’ research focus of the University of Vienna and the Austrian FWF. References [1] Valiev RZ, Mukherjee AK. Scripta Mater 2001;44:1747. [2] Ball RW, Gottstein G. Intermetallics 1994;2:205. [3] Korznikov AV, Tram G, Dimitrov O, Korznikova GF, Idrisova SR, Pakiela Z. Acta Mater 2001;49:663. [4] Horton JA, Baker I, Yoo MH. Phil Mag 1991;A63:319. [5] Rentenberger C, Karnthaler HP. Acta Mater 2005;53:3031. [6] Rentenberger C, Mangler C, Karnthaler HP. Mater Sci Eng 2004;A387–389:795. [7] Valiev RZ, Islamgaliev RK, Alexandrov IV. Progr Mat Sci 2000;45:103. [8] Wetscher F, Vorhauer A, Pippan R. Mater Sci Eng 2005;A410– 411:213. [9] Fisher RM, Marcinkowski MJ. Phil Mag 1961;6:1385. [10] Gialanella S. Intermetallics 1995;3:73. [11] Sun YQ, Hazzledine PM. Geometry of dislocation glide in L12 c’phase: TEM observations. In: Nabarro FRN, Duesbery MS, editors. Dislocation in solids, vol. 10. Amsterdam: Elsevier; 1996. p. 29. [12] Hughes DA, Hansen N. Phil Mag 2003;83:3871. [13] Schafler E, Steiner G, Korznikova E, Kerber M, Zehetbauer MJ. Mater Sci Eng 2005;A410–411:169. [14] Carling K, Wahnstrom G, Mattsson TR, Mattsson AE, Sandberg N, Grimvall G. Phys Rev Lett 2000;85:3862. [15] Rentenberger C, Waitz T, Karnthaler HP. Phys. Rev. B 2003;67:094109. [16] Hazzledine PM, Hirsch P. Antiphase domain boundary tubes in ordered alloys. In: Stoloff NS, Koch CC, Liu CT, Izumi O, editors. High-temperature ordered intermetallic alloys II, vol. 81. Pittsburgh, PA: Materials Research Society; 1987. p. 75. [17] Vidoz AE, Brown LM. Philos Mag 1962;A7:1167. [18] Ardley GW. Acta Metall 1955;3:525. [19] Starenchenko SV, Zamyatina IP, Starenchenko VA, Kozlov EV. Phys Met Metallogr 1998;85:201. [20] Yoo MH, Sass SL, Fu CL, Mills MJ, Dimiduk DM, George EP. Acta Metall Mater 1993;41:987. [21] Van Swygenhoven H, Derlet PM, Froseth AG. Nat Mater 2004;3:399. [22] Chen M, Ma E, Hemker KJ, Sheng H, Wang Y, Cheng X. Science 2003;300:1275. [23] Karnthaler HP, Waitz T, Rentenberger C, Mingler B. Mater Sci Eng 2004;A387:777.
Acta Materialia 56 (2008) 2526–2530 www.elsevier.com/locate/actamat
Extensive disordering in long-range-ordered Cu3Au induced by severe plastic deformation studied by transmission electron microscopy C. Rentenberger *, H.P. Karnthaler Physics of Nanostructured Materials, Faculty of Physics, University of Vienna, Boltzmanngasse 5, 1090 Vienna, Austria Received 8 May 2007; received in revised form 14 January 2008; accepted 26 January 2008 Available online 5 March 2008
Abstract Bulk nanocrystalline materials can be made by severe plastic deformation. In L12 long-range-ordered alloys, this leads to extensive disordering which influences the highly improved properties of these nanocrystalline alloys. Transmission electron microscopy methods were applied to Cu3Au; both diffraction contrast images and diffraction patterns reveal that disordering takes place locally. It is concluded that in addition to disordering by the refinement of the grown-in antiphase boundary domains, the formation of antiphase boundary tubes is a prominent process of disordering. The latter is facilitated by the fact that, unlike dislocations, antiphase boundary tubes can be stored at a very high density without causing long-range stresses. The local disordering indicates that the nanocrystalline structure nucleates inhomogeneously in the highly strained disordered regions. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; Transmission electron microscopy; Nanocrystalline alloys; Order–disorder phenomena; Antiphase boundary tubes
1. Introduction The extensive disordering of L12 long-range ordered (LRO) alloys by severe plastic deformation (SPD) is of special interest, as the highly improved mechanical properties of these alloys after SPD are strongly influenced by their disorder. This was shown, for example, for Ni3Al, which is nanocrystalline and disordered after SPD by high-pressure torsion (HPT) [1]. The process of disordering of LRO alloys can be studied by X-ray and neutron diffraction monitoring of the superlattice reflections. These methods yield a global average of the LRO parameter, e.g., as carried out for Ni3Al [2,3]. For the analysis of the local variations in order and the mechanism causing the disordering, transmission electron microscopy (TEM) methods based on both electron diffraction and direct imaging are appropriate and were used in the present study. In the literature, only a few TEM reports dealing with deforma*
Corresponding author. Tel.: +43 1427751314; fax: +43 1427751316. E-mail address: [email protected] (C. Rentenberger).
tion-induced localized destruction of order are found, e.g., for Ni3Al [2,4,5] and Cu3Au [6]. It is the aim of this study to analyse the process of disordering after SPD. Various methods of SPD receive currently widespread interest, as they have been applied to make bulk materials with nanocrystalline structures [7]. In the present case, L12LRO Cu3Au was deformed by HPT, as different strains (up to 5000%) can be achieved in one sample by this method. 2. Experimental procedure Polycrystalline Cu3Au samples with a grain size of 0.2 mm were annealed at 350 °C for 140 h followed by a step cooling treatment (10 K day–1 down to 160 °C) to get a fully LRO alloy with an order parameter of 1 and with domains of diameter 0.5 lm. A sample (8 mm in diameter and 0.8 mm thick) was torsion strained under a quasi-hydrostatic pressure of 4 GPa by applying 2.5 turns at room temperature (RT). The design of the HPT tool was reported in Ref. [8].
1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.01.035
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For the TEM investigation, discs were cut out by spark erosion from different parts of the HPT sample, representing different levels of strain. Sample C from the central area and sample P near the periphery correspond to shear strains c = 15 ± 5 and 50 ± 10, respectively. The discs were used to prepare TEM foils by a two-step procedure: first, they were dimpled, and afterwards they were electropolished [9]. It should be mentioned that, during electropolishing of Cu3Au, redeposition of Au on the surface of the TEM foils occurs. The TEM study was carried out with an acceleration voltage of 200 kV. Diffraction contrast images (using both fundamental and superlattice reflections) and selected area diffraction (SAD) patterns were used to analyse the microstructure and the local disordering. 3. Experimental results Fig. 1 shows a SAD pattern taken from specimen P and the corresponding intensity profile obtained by integrating rings of constant magnitude of the diffraction vector; an illuminated area 30 lm2 was used. The ring-like SAD pattern indicates a large variety of grain orientations and shows some texture. Segments of a diffuse ring observed inside the h1 1 1i ring are caused by gold redeposited during preparation; this leads to a little hump (marked Au) near the h1 1 1i-peak in the intensity profile. The comparison of the positions of the peaks with calculated ones reveals that only peaks corresponding to the fcc structure are observed, they are labelled by their indices in the intensity profile. Superlattice reflections of the L12 structure (indicated by dashed lines) are missing, indicating a LRO parameter <0.1. Therefore, the HPT deformation (c = 50 ± 10) leads to extensive disordering of Cu3Au. Fig. 2 shows dark-field images of specimen C using the superlattice reflection g ¼ ½1 1 0 (beam direction BD [1 1 0]). Two sets of faint dark lines aligned along the projections of ½0 1 1 and ½ 1 0 1 (indicated by dashed lines) are imaged in Fig. 2a. The width of the individual contrast lines is only 1–2 nm. It should be mentioned that the lines are not observed using a fundamental reflection, indicating that their contrast is caused by antiphase boundary (APB) faults. In addition, wavy contrast features are observed on the nanometre scale; they are attributed to the boundaries of the grown-in APB domains fragmented by the glide of superlattice dislocations. Fig. 2b shows an area with a high density of lines; their contrast sums up to dark bands aligned along the projection of glide planes (as indicated by the dashed line). To analyse the straight lines, a SAD pattern was taken of a small area containing a large number of them (BD [1 1 1]). The SAD pattern (cf., Fig. 3a) shows fundamental and superlattice reflections of the L12 ordered structure. In addition, streaks aligned along a direction perpendicular to V ¼ ½ 1 0 1 occur near superlattice spots only. Diffuse reflections (some are marked by arrows) are caused by gold redeposited on the surfaces. To analyse the SAD pattern, a
Fig. 1. SAD pattern and its intensity profile obtained from L12 ordered Cu3Au after SPD (shear strain 50 ± 10). The positions of the observed intensity maxima and the lack of maxima corresponding to superlattice reflections (indicated by dashed vertical lines) show that no detectable LRO is present. The hump near the {1 1 1} intensity peak is caused by a gold layer redeposited on the surface during the preparation of the TEM foil.
sketch of a sector (indicated by the two dashed lines) is shown in Fig. 3b. Circles correspond to reflections of the zero-order Laue zone (ZOLZ), whereas squares and triangles indicate reflections of the neighbouring Laue zones (+1 and –1). Full (m odd) and open (m even) symbols belong to parallel reciprocal lattice planes perpendicular to V. Only the reflections indicated by full symbols show streaks. For reflections of the ZOLZ (full circles), the streaks lie in this Laue zone. At the positions of reflections of neighbouring Laue zones, the weak contrast is caused by an intensity expanding perpendicular to the Laue zones. Reflections of neighbouring Laue zones can be moved into the ZOLZ by tilting, e.g., when BD = [2 2 1], reflection ½1 0 2 belongs to the ZOLZ and shows pronounced streaks in the plane of the projection. From the results obtained
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Fig. 2. TEM dark-field images of HPT-deformed L12 ordered Cu3Au (shear strain <20) using the superlattice reflection g ¼ ½1 0 1 (BD[1 1 0]). (a) Faint contrast lines aligned along two different h1 1 0i directions (indicated by dashed lines) are interpreted as APB tubes. The wavy contrast variations are caused by fragmented domain walls. (b) The contrast of several closely spaced APB tubes can sum up to broad bands aligned along the projection of h1 1 0i (indicated by a dashed line).
using different BD vectors, it is concluded that each superlattice reflection belonging to reciprocal lattice planes with odd m values has a disc-like shape extended in ±[1 1 1] directions. This intensity distribution of reflections with m = 1 is sketched in Fig. 3c for the ZOLZ and for the first-order Laue zone (±FOLZ). Table 1 shows the phase factor a = 2pg R for different values of g when a displacement vector R ¼ a=2½ 1 0 1 parallel to V and parallel to the defect line is assumed. R is taken to correspond to the Burgers vector of the superpartial dislocations which, together with the APB fault ribbon, form the superlattice glide dislocations. The value of a is constant and equal to pm for the different planes perpendicular to V, where m increases with increasing distance from the centre, as indicated in Fig. 3b. The occurrence of streaks agrees with the visibility criterion a = p (modulo 2p) for APB faults when R ¼ a=2½ 1 0 1 is used. In addition, streaks perpendicular
Fig. 3. Analysis of streaks observed in a SAD pattern obtained from an area showing a high density of parallel straight defects (BD [1 1 1]). (a) Experimental SAD pattern; in addition to the spot pattern, streaks aligned perpendicular to V ¼ ½1 0 1 are observed at some superlattice reflections of the L12 structure. (Some spots of the redeposited Au layer are marked.) (b) Schematic sketch of a sector of the experimental SAD pattern (indicated by dashed lines in (a)). Large circles correspond to reflections in the ZOLZ, whereas squares and triangles indicate reflections of the first-order (±FOLZ) Laue zones. Streaks (perpendicular to V) are observed only near superlattice reflections lying in planes with odd m values. (c) Schematic sketch of the intensity distribution of superlattice spots belonging to plane m = 1 and to different Laue zones.
to other h1 1 0i directions are observed, e.g., faint streaks running perpendicular to ½0 1 1 in the SAD pattern correspond to defects aligned along ±½0 1 1.
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Table 1 Values of the phase factor a = 2pgR for some reflections of different reciprocal planes m (cf., Fig. 3b) when a displacement vector R ¼ a=2½1 0 1 parallel to the defect line is assumed 1 0 0 1 g ½1 ½1 ½2 2 0 ½1 1 2 ½1 0 2 ½2 0 1 ½2 1 1
vacancies at the melting temperature [13,14]; still the reduction of the LRO by vacancies will be <1%.
m a = 2pgR
The present TEM results obtained by comparing images taken with fundamental reflections with those of superlattice reflections (cf., Fig. 2) and by the analysis of streaks occurring in the diffraction patterns (cf., Fig. 3) agree with the results expected for individual APB tubes as well as APB tubes accumulated in bands. A high density of APB tubes leads to a local and crystallographically related destruction of the LRO order and can occur without the storage of glide dislocations. Therefore, the stress field associated with APB tubes is negligible, and they can be stored at a high density without causing long-range stresses [15]. APB tubes can be formed by superlattice screw dislocations (e.g., by annihilation and by double cross-slip [16]) and by the movement of superlattice dislocations with nonaligned jogs [17]. The length of the streaks in the SAD pattern (expanding to the next Laue zone) indicates that the dimension of the tubes in ±[1 1 1] is very small (three atomic (1 1 1) planes or less). This result favours mechanisms that lead to the formation of tubes of small height; like the one by double cross-slip and that by movement of non-aligned jogs.
1 p
2 2p
2 2p
3 3p
3 3p
3 3p
3 3p
4. Discussion In several structural intermetallics, the analysis of the nanocrystals formed by SPD shows that they are disordered. This was observed in L12 ordered alloys by TEM methods (e.g., in Cu3Au and in Ni3Al deformed by HPT [5,6]) and in a B2 ordered alloy by X-ray diffraction (FeAl deformed by ball milling [10]). It is therefore proposed that the formation of a nanocrystalline structure made by SPD is correlated with disordering. Based on the present TEM results of SPD deformed Cu3Au, the following disordering mechanisms leading to a deformation-induced order–disorder transition are discussed. 4.1. Disordering by APB fault ribbons of dissociated dislocations and by point defects Deformation in L12 ordered alloys proceeds by the glide of superlattice dislocations (Burgers vector b = ah1 1 0i) [11]. They contain ribbons of APB faults reducing the local order. Intense shearing leads to storage and accumulation of dislocations and finally to a refined grain structure [12]. To estimate the contribution of APB fault ribbons of dissociated superlattice dislocations to the disordering process, calculations of the long-range order parameter were carried out. They show that, in the case of Cu3Au, a very high superlattice dislocation density of 2 1017 m–2 would be needed to decrease the order by 50%, the complete loss of order would require dislocation densities near the stability limit of the crystal (for this calculation a dissociation width of 11 nm was used, as deduced from weak-beam images of edge dislocations). It should be noted that, in contrast to the present observations, APB fault ribbons bounded by dissociated dislocations should show a contrast in the image taken, with the fundamental reflection corresponding to the superlattice reflection in Fig. 2. Also the length of the streaks in the SAD pattern caused by superlattice screw dislocations would be different from the observed one. Assuming a dissociation width of 7 nm on the (1 1 1) plane for screw dislocations lying parallel to ½ 1 0 1, the length of the streaks of the reflections in the ZOLZ would be only 1/7 of the observed one. In addition to APB faults, point defects formed during deformation can also contribute to disordering. It is reported that, after SPD, the density of deformation-induced vacancies is at RT as high as the density of thermal
4.2. Disordering by copious formation of APB tubes
4.3. Disordering by the refinement of the grown-in APB domains Gliding superlattice dislocations intersect domain walls and form new APB faults [18]. The refinement of domains can be identified by TEM methods using superlattice reflections (cf., Fig. 2a). Measurements of the reduction of the domain size with progressing deformation were carried out in L12 ordered Cu–22 at.% Pt using X-ray diffraction [19]. 4.4. Comparison with nanocrystalline Ni3Al In the case of Ni3Al that remains fully ordered up to the melting point, only the first two disordering mechanisms are possible, as no grown-in APB domains are present. Furthermore, disordering by APB fault ribbons bounded by dissociated dislocations is not, so effective as in Cu3Au, as their dissociation widths are 40% smaller in Ni3Al. Therefore, it can be concluded that, in coarse-grained Ni3Al, the local strong reduction of the long-range order [5] occurs mainly by the formation of APB tubes. The loss of LRO in extended areas takes place in Ni3Al in the nanocrystalline structure only, whereas in Cu3Au, fragmented grains are already strongly disordered [5,6]. The local correlation between disorder and nanocrystallization found by TEM methods seems to indicate that disordering is a prerequisite for nanocrystallization and that the nanocrystals nucleate inhomogeneously in the highly strained disordered regions.
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The highly improved ductility of nanocrystalline Ni3Al [1] is supported by the presence of disordering in the nanocrystalline structures of alloys that are L12-LRO in coarse grains. The reason is that the mobility of the screws is much higher in the disordered structure than in the ordered one, where they show locked configurations due to their nonplanar dislocation core structure [20]. In addition, the reduction in the length of the Burgers vector in the disordered structure (b = a/2h1 1 0i) as compared with the ordered one (b = ah1 1 0i) by a factor of 2 reduces the stress to activate glide dislocations accordingly. Furthermore, in the disordered nanocrystalline structure, the occurrence of twinning by the emission of Shockley partials (b = a/ 6h1 1 2i) from grain boundaries can contribute to ductility [21,22]. The increase in the strength in nanocrystalline alloys is caused by the reduction in grain size. In straininduced disordered nanocrystalline alloys (e.g., Ni3Al), the strength is increased even more when the deformation is carried out at an elevated temperature at which reordering takes place during deformation [1,23]. 5. Conclusions
(1) SPD by HPT of L12-LRO Cu3Au leads to extensive disordering and finally to the complete loss of order. (2) TEM shows that disordering starts locally along glide planes and is associated with the deformation-induced generation of APB. From dark-field images using superlattice reflections and diffraction patterns, it is concluded that extensive disordering is caused predominantly by the storage of APB tubes and of APB formed by intersecting grownin domains with dislocations. Contrary to APB of dissociated dislocations, APB tubes can be stored at a very high density without causing long-range stresses. For L12 ordered alloys that remain fully ordered up to the melting point (e.g., Ni3Al), it is concluded that the formation of APB tubes is the decisive process. (3) Based on the TEM results, it is proposed that in structural intermetallics (Cu3Au, Ni3Al) disordering occurring locally is a prerequisite for the formation of a nanocrystalline structure made by SPD and that the nanocrystals nucleate in the highly strained disordered regions. (4) The highly improved mechanical properties of bulk nanocrystalline alloys made by SPD are influenced by extensive disordering, leading to a higher mobility
of screw dislocations, as in the ordered structure screw dislocations show locked configurations due to their non-planar dislocation core structure.
Acknowledgments The authors thank Dr. A. Vorhauer (Erich Schmid Institute Leoben, Austria) for his kind help with the HPT deformation. The authors acknowledge the support by the ‘Bulk Nanostructured Materials’ research project within the ‘Materials Science’ research focus of the University of Vienna and the Austrian FWF. References [1] Valiev RZ, Mukherjee AK. Scripta Mater 2001;44:1747. [2] Ball RW, Gottstein G. Intermetallics 1994;2:205. [3] Korznikov AV, Tram G, Dimitrov O, Korznikova GF, Idrisova SR, Pakiela Z. Acta Mater 2001;49:663. [4] Horton JA, Baker I, Yoo MH. Phil Mag 1991;A63:319. [5] Rentenberger C, Karnthaler HP. Acta Mater 2005;53:3031. [6] Rentenberger C, Mangler C, Karnthaler HP. Mater Sci Eng 2004;A387–389:795. [7] Valiev RZ, Islamgaliev RK, Alexandrov IV. Progr Mat Sci 2000;45:103. [8] Wetscher F, Vorhauer A, Pippan R. Mater Sci Eng 2005;A410– 411:213. [9] Fisher RM, Marcinkowski MJ. Phil Mag 1961;6:1385. [10] Gialanella S. Intermetallics 1995;3:73. [11] Sun YQ, Hazzledine PM. Geometry of dislocation glide in L12 c’phase: TEM observations. In: Nabarro FRN, Duesbery MS, editors. Dislocation in solids, vol. 10. Amsterdam: Elsevier; 1996. p. 29. [12] Hughes DA, Hansen N. Phil Mag 2003;83:3871. [13] Schafler E, Steiner G, Korznikova E, Kerber M, Zehetbauer MJ. Mater Sci Eng 2005;A410–411:169. [14] Carling K, Wahnstrom G, Mattsson TR, Mattsson AE, Sandberg N, Grimvall G. Phys Rev Lett 2000;85:3862. [15] Rentenberger C, Waitz T, Karnthaler HP. Phys. Rev. B 2003;67:094109. [16] Hazzledine PM, Hirsch P. Antiphase domain boundary tubes in ordered alloys. In: Stoloff NS, Koch CC, Liu CT, Izumi O, editors. High-temperature ordered intermetallic alloys II, vol. 81. Pittsburgh, PA: Materials Research Society; 1987. p. 75. [17] Vidoz AE, Brown LM. Philos Mag 1962;A7:1167. [18] Ardley GW. Acta Metall 1955;3:525. [19] Starenchenko SV, Zamyatina IP, Starenchenko VA, Kozlov EV. Phys Met Metallogr 1998;85:201. [20] Yoo MH, Sass SL, Fu CL, Mills MJ, Dimiduk DM, George EP. Acta Metall Mater 1993;41:987. [21] Van Swygenhoven H, Derlet PM, Froseth AG. Nat Mater 2004;3:399. [22] Chen M, Ma E, Hemker KJ, Sheng H, Wang Y, Cheng X. Science 2003;300:1275. [23] Karnthaler HP, Waitz T, Rentenberger C, Mingler B. Mater Sci Eng 2004;A387:777.