A New Insight in the Plasticity of Ni3Al Intermetallic Compounds Using AFM Observations

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A new insight in the plasticity of Ni3Al intermetallic compounds using AFM observations Jol Bonnevillea, Christophe Coupeaub, Dimitri Charrierc Universite de Poitiers - CNRS UPR 3346 ENSMA Dpartement Physique et Mcanique des Matriaux Bat. SP2MI- Bd Marie et Pierre Curie F-86962 Futuroscope-Chasseneuil Cedex FRANCE ajoel.bonneville@univ-poitiers.fr, bchristophe.coupeau@univ-poitiers.fr, cdimitri.charrier@univ-poitiers.fr Keywords: intermetallic, atomic force microscopy, dislocation, slip trace. Abstract. The positive temperature dependence of flow stress in Ni3Al is examined through fine slip trace analysis performed by atomic force microscopy. Slip traces, which result from dislocation movements, constitute outstanding markers for investigating the elementary dislocation mechanisms that control plasticity. The experiments were performed on two Ni3Al-base alloys, with Ta or Hf as additional elements. The results give evidence that the anomaly domain is accompanied by a drastic exhaustion of mobile dislocations and very short cross-slip distances on the cube cross-slip plane. Introduction Plastic deformation of crystalline materials yields noticeable surface effects, such as for instance the occurrence of slip traces, which were used as first evidences for identifying the basic deformation mechanisms of bulk materials. However, plastic deformation induces surface characteristics that are usually observed at large length scale and only provide rough information on the plastic features at the microscopic level. Recent developments of scanning probe microscopes [1], such as atomic force and scanning tunnelling microscopy (AFM and STM), enable nano-scale surface features to be routinely imaged, measured and analysed [2-3], even under in situ deformation conditions [4]. The atypical mechanical properties of some L12 intermetallic compounds (for concise review see [5]) have attracted great interest in the community of materials scientists, since now more than half a century. In particular, the yield stress of Ni3Al intermetallic compounds exhibits positive temperature dependence in a given temperature range, commonly referred to as yield stress anomaly (YSA), which violates the principles of thermal activation theory. Extensive mechanical testing and microstructural investigations have been performed, which have yielded a variety of theoretical models. Detailed reviews are given in [6]. The models can be roughly classified in three categories: forest hardening, lattice friction or cross-slip. The latter models are more successful in predicting the YSA and it is now considered that the YSA results from specific features of dislocation core configuration and from cross-slip mechanisms. In particular, a milestone was reached when the complex effects of stress orientation on the YSA were explained [7]. However, the latter model has been strongly criticized on the basis of electron microscopy observations and the origin of the YSA has evolved gradually over the years from a point-like pinning towards a microstructure-based description [5]. In this context, the atomic-scale study by AFM of screw dislocation motion through the fine structure of slip traces constitutes a unique tool for providing unambiguous description of their propagation mechanisms. In this paper, we report AFM observations obtained on Ni3(Al,Ta) and Ni3(Al,Hf) single crystalline alloys. The slip traces were post mortem examined after compression tests performed at various temperatures in the YSA of Ni3(Al,Ta) and in situ examined during compression tests at room temperature for Ni3(Al,Hf). Particular attention was here devoted to establish whether the YSA arises from an exhaustion of mobile dislocations or from a decrease in dislocation velocity [8]. Key Engineering Materials Vol. 465 (2011) pp 403-406Online available since 2011/Jan/20 at www.scientific.net (2011) Trans Tech Publications, Switzerlanddoi:10.4028/www.scientific.net/KEM.465.403All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP,www.ttp.net. (ID: 128.206.9.138, University of Missouri-Columbia, Columbia, United States of America-26/11/13,10:47:28) Experimental details For our purpose, two monocrystalline alloys with Ni74Al24Ta1 and Ni74.8Al21.9Hf3.3 compositions were used. In the following, these alloys will be referred to as Ni3(Al,Ta) and Ni3(Al,Hf). All samples were deformed in compression along the [123] crystallographic orientation. Prior to deformation, the specimens were mechanically polished for suitable AFM observations. The AFM observation face was parallel to the )154( plane. The primary slip system is a )111](110[ , a = 0.36 nm, which corresponds to slip traces lying at 62 to the compression axis on the observation face. The unit step height h, corresponding to the emergence of one superdislocation, is equal to the Burgers vector component projected onto the surface normal, i.e. n.bh= = 0.234 nm. The Ni75Al24Ta1 samples were conventionally deformed up to approximately 10-2 plastic shear strain at a nominal applied strain-rate of 10-4 s-1, using a Schenk RMC 100 testing machine. Three deformation temperatures were selected, namely: 293 K, 505 K and 720 K. The first temperature corresponds to the onset of the YSA domain while the third one corresponds to a temperature nearly 80 K below the stress peak with temperature. To prevent surface contamination, the deformation tests at 505 K and 720 K were performed under Ar atmosphere. The resulting slip traces were then post mortem examined by AFM. The AFM device, described in details elsewhere [4], also allows for surface imaging during in situ deformation. This facility was used to deform the Ni3(Al,Hf) single crystalline samples at room temperature, which for this alloy corresponds to a temperature located within the YSA domain [9]. For avoiding spurious vibrations, the compression experiment is stopped just prior to scanning the surface The AFM images are recorded while the specimen is under stress relaxation conditions, resulting in a small stress decrease of a few percents of the total applied stress [9], so that each AFM image can be ascribed to one strain (or stress) value only. All AFM images presented below were taken in signal error contact mode, so that a single surface step feature appears as one peak. These images are not calibrated in the direction perpendicular to the observation surface, but have visual advantage of fine detail enhancement in the plane surface. Results Ni3(Al,Ta) Figure 1 shows typical images of slip traces obtained on the Ni3(Al,Ta) samples for the three deformation temperatures. At all temperatures, the slip traces that are visible in the central part of the samples belong to the primary (111) octahedral glide plane. Two striking features emerge from the images. 2 m2 m 2 m(a) (b) (c) Fig. 1, Typical AFM images of slip traces observed on Ni3(Al,Ta) samples after nearly 1% plastic strain deformed at (a) 293 K, (b) 505 K, (c) 720 K. The slip traces correspond to the primary (111) octahedral plane. 404 Materials Structure & Micromechanics of Fracture VI First, slip traces are rectilinear, without appreciable deviation from the primary octahedral glide plane. Imaging at higher magnification indicates that the intimate slip trace structure is composed by several finer slip traces that are distributed on parallel primary octahedral glide plane. At the limit of in-plane AFM resolution under atmospheric conditions, the finer slip traces are not connected by cross-slip events. On this point, it is necessary to clarify that the present AFM device allows for out-of-plane atomic resolution, but in-plane resolution is much lower and of the nanometer range. Second, majority of slip traces crosses the entire image at 293 K, but with increasing temperature it is clearly observed an increasing proportion of slip traces ending within the images. This leads to shorter slip traces with increasing temperature, which lengths are given in figure 2. Figure 2 clearly shows that the slip trace length strongly decreases with increasing temperature, even for temperature higher than the work-hardening peak temperature. It should be also noted that although the slip traces are observed after 1% plastic strain they account for all plastic strain history of the samples. In order 0100200300400500600700800200 300 400 500 600 700 800slip line length / mT / K Fig. 2, Slip trace lengths for Ni3(Al,Ta) as a function of temperature. The error bars correspond to the statistical error about the average values (open symbol). to check for possible artefacts, some samples were re-polished and re-deformed by 1% plastic strain. The slip trace length is found significantly shorter at 293 K only, but the decreasing trend is confirmed [11]. Ni3(Al,Hf) Figure 3 shows a series of AFM images recorded on a Ni3(Al,Hf) single crystalline sample deformed at increasing plastic strain. The slip traces exhibit similar characteristics to those of Ni3(Al,Ta) samples: they correspond to the primary (111) octahedral slip plane, they are rectilinear and exhibit limited lengths. Topological analysis of AFM observations indicates that slip-trace length remains essentially constant with increasing plastic strain. In addition, a tedious, but instructive, examination of the number of emerging dislocations per slip line at the different plastic strain levels indicates that they are composed by a few dislocations only, of the order of four dislocations per slip traces. Therefore, once a slip line has been created during a plastic strain increment, it does not evolve anymore. It is also remarkable from figures 3(a) to 3 (d) that new lines are continuously created after each successive plastic strain increment. [111][123](a) (b) (c) (d)2 mFig. 3, AFM images of slip traces observed on Ni3(Al,Hf) samples at increasing plastic strain. The slip traces correspond to the primary (111) octahedral plane: (a) p 0.2%, (b) p 0.3%, (c) p 0.4%, (d) p 0.5%. Key Engineering Materials Vol. 465 405 Discussion and conclusion One of the prominent features of AFM observations is the high density of ending slip traces for both alloys, which indicates that the slip trace lengths are markedly small, especially for single crystals oriented for single slip. In addition, once the slip traces are formed, their lengths and their dislocation number remain almost constant with increasing plastic strain. Considering that the slip trace length is, in some way, related to the mean free path of dislocations, these results are interpreted as a strong exhaustion of mobile dislocations, resulting from a permanent locking of the immobilised dislocations. In this frame, (1) the strong decrease in slip trace length with increasing temperature reflects an increasing exhaustion rate of mobile dislocations when the temperature is raised and (2) the increasing number of slip traces with increasing plastic strain suggests that new dislocation sources are continuously activated to compensate for the high exhaustion of mobile dislocations. Mobile dislocation exhaustion has already been reported using indirect experimental techniques [12], but to our knowledge these AFM observations constitute the first direct evidence of mobile dislocation exhaustion, an ingredient which has not been sufficiently considered in the modelling of the YSA. Another characteristic of the slip lines is their straightness, indicating that dislocation movement is predominantly planar. Therefore, the cross-slip distance of moving dislocations should be very small, i.e. not larger than the dislocation Burgers vector. This result would be in agreement with point-like pinning models, but does not exclude cross-slip microstructure-based models that consider the cross slip distance on the cube cross-slip plane as an adjustable parameter. Such a small cross-slip distance has been proposed for instance to explain the peak in hardening rate, which is observed at a temperature below the stress peak temperature [13]. The decrease in hardening after the hardening peak is however attributed to the stress assisted re-mobilisation of the incomplete Kear-Wilsdorf lock, a feature which a priori is not observed in the present work. To conclude, AFM observations of the slip traces in the YSA of Ni3Al intermetallic alloys suggest that the YSA results from a strong exhaustion of the density of mobile dislocations. This exhaustion certainly results from a cross slip mechanism of the screw dislocation over short distances from the primary onto the cube cross slip plane, which preserves the orientation effect predictions of the so-called PPV model [7]. Further work is in progress to access the in-plane resolution in order to visualise at the atomic scale the three dimensional geometry of the slip lines. References [1] G. Binnig, H. Rohrer, Ch. Gerber, E. Wiebel, Phys. Rev. Letters 49 (1982), p. 57. [2] C. Coupeau, J.-C. Girard, J. Grilh, J. of Vacuum Science Technology B 16 (1998), p 1964. [3] C. Coupeau, J. Bonneville, Applied Physics Letters 90 (2007) 171914. [4] C. Coupeau, J. C. Girard, J. Rabier, Dislocations in Solid Volume 12, eds F.R.N. Nabarro and J.P. Hirth, Elsevier Science (2004), p. 273. [5] P. Veyssire, Intermetallics 6 (1998), p. 587. [6] F. R. N. Nabarro, M. S. Duesbery, Dislocations in Solids, vol. 10 (1996). [7] V. Paidar, D. P. Pope, V. Vitek, Acta Metall. 32 (1984), p. 435. [8] F. Louchet, Encyclopedia of Marerials: Science and Technology, Elsevier Science, Amsterdam (2001), p. 4158. [9] P. Sptig, J. Bonneville, J-L. Martin, Mat. Sci. Eng. A 167 (1993), p. 73. [10] P.H. Thornton, R.G. Davies, T.L. Johnston, Metall. Trans. 1 (1970), p 207. [11] J. Fikar, C. Coupeau, T. Kruml, J. Bonneville, Mat. Sci. Eng. A 387-389 52004), p. 926. [12] B.L. Cheng, E. Carreo-Morelli, N. Baluc, J. Bonneville, R. Schaller, Phil. Mag. A 79 (1999), p. 2227. [13] T. Kruml, E. Conforto, B. Lo Piccolo, D. Caillard, J-L. Martin, Acta Mater. 50 (2002) p. 5091. 406 Materials Structure & Micromechanics of Fracture VIMaterials Structure & Micromechanics of Fracture VI 10.4028/www.scientific.net/KEM.465 A New Insight in the Plasticity of Ni3Al Intermetallic Compounds Using AFM Observations 10.4028/www.scientific.net/KEM.465.403 DOI References[1] G. Binnig, H. Rohrer, Ch. Gerber, E. Wiebel, Phys. Rev. Letters 49 (1982), p. 57.doi:10.1103/PhysRevLett.49.57 [2] C. Coupeau, J.-C. Girard, J. Grilh, J. of Vacuum Science Technology B 16 (1998), p 1964.doi:10.1116/1.590234 [3] C. Coupeau, J. Bonneville, Applied Physics Letters 90 (2007) 171914.doi:10.1063/1.2731436 [7] V. Paidar, D. P. Pope, V. Vitek, Acta Metall. 32 (1984), p. 435.doi:10.1016/0001-6160(84)90117-2 [12] B.L. Cheng, E. Carreo-Morelli, N. Baluc, J. Bonneville, R. Schaller, Phil. Mag. A 79 (1999), . 2227.doi:10.1080/01418619908210419 [12] B.L. Cheng, E. Carreo-Morelli, N. Baluc, J. Bonneville, R. Schaller, Phil. Mag. A 79 (1999), p. 2227.doi:10.1080/01418619908210419 [13] T. Kruml, E. Conforto, B. Lo Piccolo, D. Caillard, J-L. Martin, Acta Mater. 50 (2002) p. 5091.doi:10.1016/S1359-6454(02)00364-6

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