IN-SITU OBSERVATION OF CREEP DAMAGE IN Al-Al 2 O 3 MMCs BY SYNCHROTRON X-RAY TOMOGRAPHY

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1 73 IN-SITU OBSERVATION OF CREEP DAMAGE IN Al-Al 2 O 3 MMCs BY SYNCHROTRON X-RAY TOMOGRAPHY A.Pyzalla 1,4, B.Camin 2, B.Lehrer 2, M.Wichert 2, A.Koch 2, K.Zimnik 1,2, E.Boller 3, W.Reimers 2 1 Institute of Material Science and Technology, TU Wien, Karlsplatz , 1040 Wien, Austria 2 Institute of Material Science and Technology, TU Berlin, Ernst-Reuter-Platz 1, Berlin, Germany 3 European Synchrotron Radiation Facility ESRF, Grenoble, France 4 Now: Max-Planck-Institute for Iron Research, Düsseldorf, Germany ABSTRACT Creep damage evolution in AA %Al 2 O 3 particles is observed in-situ by synchrotron radiation tomography using a miniature creep device. Reconstructions of the gauge volume of the creep samples show the particle delamination, particle fracture, void formation and growth at 340 C at various stages of the creep process. KEYWORDS MMC, tomography, creep, damage, X-ray INTRODUCTION Compared to monolithic metals or ceramics, particle reinforced metal matrix composites (MMC p ) can be advantageous due to combining the properties of the matrix and the reinforcement. In case of aluminium matrix composites usually the strength, stiffness and wear resistance are superior to the aluminium matrix alloy. The high thermal conductivity and low thermal expansion coefficient makes Al-base MMC p attractive for applications in automotive engine technology and aircraft industry. The creep behaviour of most MMC p is significantly different from their monolithic matrices. This has been attributed to the internal stresses generated by the load transfer between matrix and particles, which may increase damage, e.g., by promoting particle cracking, delamination at the particle/matrix interface as well as by facilitating void nucleation and void growth. The presence of these internal stresses during the whole creep processes has been proven using neutron diffraction by Winand et al. [1,2] in in-situ creep experiments of Al/SiC MMCs with whisker reinforcements. The experiments performed by [1,2] also revealed that the overall load transfer is extremely sensitive to the reinforcement geometry and test conditions. The influence of the reinforcement geometry and distribution on the nucleation and growth of creep voids has also been indicated by numerical simulations, e.g., by [3-5]. Creep curves of MMC p often show a very brief secondary (steady-state) stage and a far more extensive tertiary stage, which may be attributed to the earlier damage formation compared to

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3 74 their monolithic matrices [e.g. 6]. However, quantitative studies on damage evolution during creep deformation are rare [7-9]. In order to link damage evolution in Al-Al 2 O 3 metal matrix composites to their strain rate we explored the possibilities for in-situ X-ray tomography during the creep process in an experiment at the beamline ID19 of ESRF, Grenoble, France. X-ray tomography for studying creep damage is strongly advantageous to ex-situ metallography and microscopy since it provides the chronology of damage events and their interactions. Moreover it reveals damage location in a defined volume, e.g., next to a particle corner. A particular advantage of X-ray microtomography over determining damage evolution e.g. in in-situ creep tests in a scanning electron microscope is that the data obtained in microtomography represents bulk information. Buffière et al. [10] have shown that bulk values, e.g., of volume fractions of broken particles can substantially differ from those obtained at the surface. CREEP DEVICE A creep device (fig.1) has been designed and manufactured, where the incoming and the emerging X-rays for a 180 turn of the sample are not hindered by any parts of the device except a 0.5mm thick homogeneous magnesium alloy tube. Thus absorption is homogeneous and artefacts can be avoided. Since electric motors are known to introduce vibrations into the set-up, the sample was mechanically loaded by a spring rather than by an electric motor. The external stress on the sample can be adjusted by a spring with a suitable spring constant and adjusting the deformation of the spring. The forces on the sample are diverted by the magnesium shielding tube. The sample is heated by the radiant heat exerted by an inductor, which is coiled around the lower part of the sample beneath the area where the tomograms are taken. Temperature is controlled by temperature measurements using a thermocouple sitting in a small hole near the gauge volume (above the heating device) and a Fig. 1: Creep device temperature controller. During the experiment the sample temperature variations are ±1K. Ex-situ measurements of the temperature gradient along the sample have shown that the temperature difference along the active area of the sample amounts to ΔT = 2 C at T = 200 C res. ΔT = 4 C at T = 470 C. The maximum temperature reachable in the device is about 500 C (limitation is the thermal stability of the magnesium tube).

4 75 EXPERIMENTAL DETAILS Samples were machined from AA6061+Al 2 O 3 hot extruded rods. The sample longitudinal axes coincide with the longitudinal axes of the rods. Care was taken to achieve a high surface quality during the machining process. High-resolution X-ray tomography experiments were performed at beamline ID19 at the European Synchrotron Radiation Facility ESRF, Grenoble, France using a monochromatic beam with 20.5 kev energy. The field of view of the camera was 600µm x 600µm. Scans of 900 twodimensional radiographs were taken by rotating the sample in 0.2 steps. Each radiograph took ~3s, one complete tomogram needed app. 55min. The resolution obtained was ~0.28µm. From the 2D radiographs the volume element of the sample was reconstructed using a conventional 2D filtered back projection algorithm. The creep device was mounted on the rotation stage upon an x-y-z-translation stage, necessary for sample alignment. The sample was heated to 340 C (heating-up to the set-point takes about 1 min.). The creep stress exerted on the sample was σ = 25MPa. Tomograms were taken in the initial state of the sample and continuously during the creep experiments. The damage evolution in the MMC with increasing creep time was analysed manually on the reconstructed images of the composite. This method is rather time-consuming and limits the size of the sample volume that can be analysed. However, it was preferred to automatic analyses (similar to the studies by [10]) since the phase contrast at the interface between particles and matrix was underlined by dark/bright fringes. RESULTS AND DISCUSSION The tomograms reveal a good contrast between the Al 2 O 3 -particles and the AA6061 metal matrix (fig. 2). In a few slices small (~ 3µm diameter) voids are visible in this initial state, which are due to decohesion at the particle-matrix interfaces during the hot extrusion process. Similar extrusion process induced voids appeared in the Al-SiC composites investigated by [10]. A comparison of a horizontal and a vertical slice through the sample shows a preferential orientation of the particle longitudinal axis in the extrusion direction (fig. 3). The particle shape is polygonal and the mean diameter of the particles is app. 11µm. The AA6061 matrix appears homogeneous in its absorption (fig.2, 3). 150µm Fig. 2: 3D representation of a sample during creep, view in the direction of the sample longitudinal axis, T = 340 C, t creep = 360 min

5 76 The tomograms obtained during the in-situ creep experiments at ID19 reveal that it is possible to observe the time dependent damage formation in metal matrix composites in-situ. Limits of this technique today are mainly the available resolution, which does not give access to the very early stages of damage and the time necessary for a complete tomogram. If the sample creep is fast, especially in the tertiary stage, the rapid growth of the sample leads to blurred images. The reconstructed sample volumes showed particle fracture (fig. 3, fig. 4); some particles, however, were already broken during extrusion. Particle fracture due to creep processes has Fig. 3: Longitudinal slice through a sample during creep, T = 340 C, t creep = 660min, slice taken in the center of the sample also been observed by [11] in this temperature range around 300 C. Experimental data and a model for the competition between particle fracture and particle decohesion of MMCs has recently been proposed by [12], which showed particle fracture to be dominant in composites with high yield stress and work hardening rate and particle decohesion to prevail in case of MMCs with a low yield stress and medium work hardening rate. In case of creep damage of AA6061+Al 2 O 3 at 340 C the absolute dominance of particle decohesion is presumably due to the high strength of the Al 2 O 3 particles (fracture probability 37% at σ=1990mpa [13]), the low creep stress resulting in a low stress level in the metal matrix and the possibility of stress relaxation due to diffusional Fig. 4: Voids at the particle matrix interface and voids connecting particles, part of a longitudinal slice through the sample, T = 340 C, t creep = 11h

6 77 accomodation. Decohesions occurred at the particle/matrix interface (fig. 4, 5); they are easily seen at the front face of particles elongated in the direction of the load axis. The shape of the voids produced by decohesion adapts to the shape of the particle-matrix-interface. Voids form at the poles (in load direction) of the particles. These poles according to [14,15] are regions with maximum hydrostatic stresses and strain concentrations. Particularly often voids seem to be formed at sharp particle corners. In the vicinity of these sharp corners Christman et al. [3] have shown by finite element modelling that the hydrostatic stress level is elevated, thus these regions are prone to void formation and void growth. Analyses of sequences of tomograms obtained with increasing creep time showed that new decohesions often appear at the same particles or at a neighbouring particle. The decohesions tend to grow and to join along the particle/matrix interface. In regions where particles agglomerate and thus appear close to each other, often voids are visible between particles which are only separated by a thin bridge of matrix material (fig. 5). These voids tend to be comparatively large. Regions with high particle density usually coincide with a high void volume in the metal matrix. FE modelling by [15] has shown that this damage accumulation due to void nucleation and growth within regions with Fig. 5: Voids between agglomerated particles particle clusters can be attributed to high stress multiaxality and constrained plastic flow within the metal matrix. CONCLUSIONS The tomography of in-situ creep experiments at ID19 at ESRF revealed that it is possible to observe the time dependent damage formation in metal matrix composites in-situ. Limits of this technique today are mainly the available resolution, which does not reveal the very early stages of void formation, and the time necessary for a complete tomogram. If the sample creep is fast, especially in the tertiary stage, the rapid change of the sample may lead to blurred images.

7 78 The observation of creep damage in AA vol.-%Al 2 O 3 particles revealed that o under the experimental conditions chosen, damage by interface delamination dominates and particle fracture is rare, o voids are formed at the particle/matrix interfaces, especially at sharp particle corners o once a void nucleates at a particle soon other voids appear at the same particle o damage appears to proceed by void agglomeration at the particle/matrix interfaces rather than by void growth o high void concentrations are visible in regions with high particle density and in-between close neighbour particles, where multi-axial strain states exist o void shape is not related to external stress orientation, but rather to the orientation of the particle/matrix interface. More insight into the time-dependence of the damage process is expected from faster tomography experiments and the combination of diffraction with tomography, since this will give access also to the load transfer between matrix and particles. ACKNOWLEDGEMENTS The authors gratefully acknowledge the use of beamline ID19 at ESRF and the support of the beamline staff. REFERENCES [1] Winand, H.M.A., Withers, P.J., Physica B (1997) [2] Winand, H.M.A., Whitehouse, A.F., Withers, P.J.: Mat. Sci. Eng. A284 (2000) [3] Christman, T., Needleman, A., Suresh, S.: Acta Met. 37 (1989) [4] Pandorf R., Broeckmann, C. Metal and Materials International 4 (1998) [5] Goda, K., Int. J. Plasticity 18 (2002) [6] Lin, Z., Li, Y, Mohamed, F., Mat. Sci. Eng. A332 (2002) [7] Carroll, H.E., Whitehouse, A.F., Scripta mater. 42 (2000) [8] Schürhoff, J., Yawny, A., Skrotzki, B., Eggeler, G, Mat. Sci. Eng. A (2004) [9] Humphreys, F.J.: Proc. EMAG-MICRO 89, IOP, Bristol (1990) 465 [10] Buffière, J.-Y., Maire, E., Cloetens, P., Lormand, G., Fougères, R., Acta. mater. 47 (1999) [11] Ma, Y., Xia, K., Mihlich, J.L., Langdon, T.G.: Creep behaviour in an aluminium composite containing alumina microspheres, in: T. Chandra, A.K. Dhingra: Proc. Conf. Advanced Composites, Warrendale, USA (1993) [12] Babout, L., Maire, E.,Fougère, R., Acta Mat. 52 (2004) [13] Mochida, T., Taya, M., Lloyd, D.J, Materials Transactions JIM 32 (1991) [14] Xu, X.Q., Watt, D.F., Acta metall. mater. 42 (1994) [15] Watt, D.F., Xu, X.Q., Lloyd, D.J, Acta mater. 44 (1996)