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




    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.


    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.465.403

    All 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:, 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.



    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 m

    2 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










    200 300 400 500 600 700 800



    e len




    T / 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].


    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.


    (a) (b) (c) (d)

    2 m

    Fig. 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.


    [1] G. Binnig, H. Rohrer, Ch. Gerber, E. Wiebel, Phys. Rev. Letters 49 (1982), p. 57.

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    [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.

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    [7] V. Paidar, D. P. Pope, V. Vitek, Acta Metall. 32 (1984), p. 435.

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    [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),

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    [13] T. Kruml, E. Conforto, B. Lo Piccolo, D. Caillard, J-L. Martin, Acta Mater. 50 (2002) p.


    406 Materials Structure & Micromechanics of Fracture VI

  • Materials 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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