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Some Aspects of Cross-Slip Mechanisms in Metals and Alloys Daniel Caillard, J.L

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Some Aspects of Cross-Slip Mechanisms in Metals and Alloys Daniel Caillard, J.L Some aspects of cross-slip mechanisms in metals and alloys Daniel Caillard, J.L. Martin To cite this version: Daniel Caillard, J.L. Martin. Some aspects of cross-slip mechanisms in metals and alloys. Journal de Physique, 1989, 50 (18), pp.2455-2473. 10.1051/jphys:0198900500180245500. jpa-00211074 HAL Id: jpa-00211074 https://hal.archives-ouvertes.fr/jpa-00211074 Submitted on 1 Jan 1989 HAL is a multi-disciplinary open access L’archive ouverte pluridisciplinaire HAL, est archive for the deposit and dissemination of sci- destinée au dépôt et à la diffusion de documents entific research documents, whether they are pub- scientifiques de niveau recherche, publiés ou non, lished or not. The documents may come from émanant des établissements d’enseignement et de teaching and research institutions in France or recherche français ou étrangers, des laboratoires abroad, or from public or private research centers. publics ou privés. J. Phys. France 50 (1989) 2455-2473 15 SEPTEMBRE 1989, 2455 Classification Physics Abstracts 62.20H Some aspects of cross-slip mechanisms in metals and alloys D. Caillard (1) and J. L. Martin (2) (1) Laboratoire d’Optique Electronique du CNRS, B.P. 4347, F-31055 Toulouse Cedex, France (2) Institut de Génie Atomique, Ecole Polytechnique Fédérale de Lausanne, CH-1015 Lausanne, Switzerland (Reçu le 6 mars 1989, accepté sous forme définitive le 24 mai 1989) Résumé. 2014 Des modèles ont été proposés pour décrire le glissement dévié des dislocations dans les métaux de structure CFC et HC. On montre comment ces modèles se sont développés et comment des résultats expérimentaux récents confirment leurs prédictions. Quelques aspects macroscopiques de la plasticité des cristaux tels que les énergies d’activation de fluage, les anomalies de limite élastique et les stades de durcissement sont discutés à la lumière de ces modèles microscopiques. Abstract. 2014 Dislocation cross slip in FCC and HCP metals has been described in terms of various models. The development of these models is reviewed and their predictions are compared with recent experimental observations. Some macroscopic features of plastic deformation (creep activation energies, strength anomalies and deformation stages) are examined in the light of these microscopic mechanisms. 1. Introduction. Cross slip occurs when the local stresses push a dislocation into a plane which is different from the original plane of splitting. It has been evidenced long ago by the observation of slip traces at sample surfaces. The outlines of the first realistic models were then derived by Friedel, both for face-centered cubic (FCC) metals [1] in 1956 on the basis of calculations of Stroh (1954) [2] and for hexagonal close-packed (HCP) metals [3] in 1959. In the mean time, another theory for dislocation cross slip was developed (see e.g. Refs. [4] and [5]). Later on, similar slip trace observations, after creep at intermediate temperatures, led some authors to interpret the corresponding activation energies in terms of cross-slip (Lytton et al. , 1958 [6] ; Friedel, 1959 [3] and 1964 [7]). This mechanism has also been claimed to be the key process at the transition stress ’iII1 between the hardening stages II and III of FCC metals (Schoeck and Seeger, 1955 [4] ; Friedel, 1956 [1] and Haasen, 1967 [8]). More recently, the importance of cross slip in the plasticity of metals has however been ignored or underestimated. In the present paper, we attempt to describe how these first ideas developed with time into detailed models, and how these have been confirmed or improved recently through Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphys:0198900500180245500 2456 sophisticated deformation tests, or electron microscope observations. The latter consist mainly of in situ deformation experiments, which have been recently described as a new technique to investigate old problems [9]. We first review the results related to cross slip in FCC metals and then consider Peierls type friction related to slip in body centred cubic (BCC) metals and non compact slip in metals with a close packed structure. We then summarize the results concerning a newly observed type of friction, the « locking-unlocking mechanism », both for prismatic slip of beryllium and slip in L12 ordered structures. Macroscopic aspects of plastic deformation such as strength anomalies, creep activation energies, deformation stages of the FCC metals are then discussed in terms of the above dislocation mechanisms. 2. Cross slip in FCC metals. The elementary process of dislocation cross slip from a compact plane onto another compact plane although frequently observed experimentally, is a very poorly known process. The theoretical aspects of cross slip have been studied by Escaig (1968 [10, 11]), following Friedel’s view. In his model, cross slip occurs through the mechanism sketched in figure 1. It starts from a primary constriction on a screw dislocation. Under the influence of the applied stress aided by thermal activation, the two halves of the constriction are separated and, as suggested by Friedel (1956 [1], 1957 [12]), the dislocation segment between them, splits at once on the cross slip plane. This model seems to be more realistic than the former one of Schoeck and Seeger (1955) [4]. In particular, Escaig has shown that the process depends mainly on the ratio of the width of splitting under stress on the primary and the cross slip planes. This causes the process to be orientation dependent, which will be discussed below. Fig. 1. - Successive stages of cross-slip starting at a constriction, in Friedel’s and Escaig’s model [1, 10, 11]. The activation energy AGcs has been derived for weak effective stresses : with i = 1 or 2 depending on how many constrictions are necessary to initiate the process ; bo is the width of a recombined screw dislocation ; 1£ the shear modulus, a (i ) an orientation dependent parameter, is the applied resolved shear stress ; f the stacking fault energy ; 2457 Do the dissociation width at zero stress ; A is a slowly variable function of IL, b, f. The activation volume of the process is : A new experimental technique has been developed to produce a burst of cross slip at the yield point [13] : a copper single crystal is deformed in compression along [110] to the end of stage II to produce a homogeneous dislocation forest. New samples are then cut out of this crystal, in a single slip orientation. The new primary dislocations are expected to cross slip at the yield point because of the numerous forest dislocations acting as obstacles. The tension compression asymmetries predicted by the model, as well as the orientation dependence of the critical stress for cross slip, are in good agreement with experimental results obtained at room temperature. As the temperature is changed, load relaxation experiments are performed along the stress strain curves to obtain activation volumes. An example is shown in figure 2 for a predeformation experiment at 473 K. The corrected activation volume exhibits a minimum as a function of strain. This corresponds to a value of about 300 atomic volumes, in fairly good agreement with equation (2) and the width of splitting of copper. The stress corresponding to this minimum is then considered to be the critical stress for cross slip, which is confirmed by slip trace observations. The variation as a function of temperature of this critical stress and of the corresponding activation volume, is shown in figure 3 [14]. The barrier energy for cross slip in copper was estimated to be : AG§s = 1.15 ± 0.37 eV, which falls between the two values predicted by equation (1) for i = 1 and 2 respectively. An important conclusion is that after the above measurements, it became obvious that the TIII stress and its variation as a function of temperature are different from the critical stress for cross slip between compact planes. The above experimental results, especially the orientation dependance of the critical stress for cross slip, are clearly in favor of Friedel’s and Escaig’s views. The comparison with the second type of theory for cross slip [4] and [5] has already been given elsewhere [13]. Similar orientation effects have also been found in Ni3Al and interpreted successfully using the Friedel-Escaig cross slip theory [15]. Fig. 2. - Resolved shear stress and corrected activation volume as a function of resolved shear strain. Cu single crystals. Deformation tests at 293 K. Predeformation tests at 473 K [14]. 2458 Fig. 3. - Critical stress for cross slip and corresponding activation volume as a function of temperature in Cu. Cross slip takes place at the macro elastic limit T between 250 and 410 K. v is the atomic volume [14]. 3. Friction forces of the Peierls type. The principle of this kind of friction force is that the energy of moving dislocations necessarily varies according to the lattice periodicity. Dislocations thus have their minimum energy in the so called « Peierls valleys », and their movement can be described as a series of jumps between adjacent valleys, through the nucleation and propagation of kink pairs. In the early model of Peierls and Nabarro, the dislocations are assumed to spread only in their glide planes. This results in very low friction forces, especially for dislocations which are largely split (case of glide on the most compact planes). More recent studies have shown that even larger frictional forces can originate when dislocations spread out of their glide plane. Under such conditions, the Peierls mechanism can play a very important role in the plasticity of materials, even at high temperatures. This is for instance the case of L12 ordered alloys, for which friction forces are present up to 0.7 Tm (Tm : melting temperature). These higher friction forces are found to concern mainly screw dislocations, which can spread more easily in different planes, because of their higher symmetry within the crystal lattice.
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