Monotonic and cyclic deformation behaviour of ultrafine-grained aluminium

Materials Science and Engineering A 483–484 (2008) 481–484

Monotonic and cyclic deformation behaviour of ultrafine-grained aluminium J. May ∗ , D. Amberger, M. Dinkel, H.W. H¨oppel, M. G¨oken Department of Materials Science and Engineering, Institute I: General Materials Properties WW I, University Erlangen-N¨urnberg, Martensstr. 5, 91058 Erlangen, Germany Received 6 June 2006; received in revised form 29 September 2006; accepted 5 December 2006

Abstract The effect of the enhanced strain-rate sensitivity (SRS) of ultrafine-grained commercially pure aluminium Al 99.5 on the mechanical properties under monotonic as well as under cyclic loading was investigated. Compared with the conventional grain-sized counterpart, for the monotonic tests, a strongly enhanced strength combined with a high ductility was obtained, depending on the strain rate. The enhanced SRS also affects the cyclic deformation behaviour and the fatigue lives. At a lower strain rate shorter fatigue lives and a different cyclic hardening behaviour are observed. Microstructural changes during cyclic deformation are investigated by X-ray diffraction profile analysis. Based on the fatigue behaviour and the X-ray diffraction profile analysis, thermally activated annihilation processes of dislocations are regarded to be the relevant deformation mechanism leading to enhanced SRS and ductility. © 2007 Elsevier B.V. All rights reserved. Keywords: Ultrafine-grained metals (UFG); Equal-channel angular pressing (ECAP); Strain-rate sensitivity; Fatigue; X-ray peak profile analysis

1. Introduction Ultrafine-grained (UFG) metals are known to have excellent mechanical properties compared with the conventional grainsized (CG) materials. A high monotonic strength coupled with a high ductility has been reported for several UFG and nanocrystalline materials in the meanwhile [1–5]. In addition, for UFG materials an enhanced strain-rate sensitivity (SRS) has been found [6–9], which is regarded to be one of the dominating features governing the high ductility of these materials [10]. Not only under monotonic loading, but also during cyclic loading the UFG materials show extraordinary properties. For example, in a W¨ohler SN plot UFG materials show superior fatigue lives compared to the CG counterparts. For more information about the fatigue properties of UFG materials, the reader is exemplarily referred to Refs. [11–13]. In the present work, the deformation behaviour of commercially pure Aluminium Al 99.5 has been investigated with an ultrafine grained and a conventionally grain-sized microstructure, each under monotonic and under cyclic loading conditions.

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The SRS has been investigated by varying the strain rate between the different experiments. The development of grain size and dislocation density during fatigue has been investigated by X-ray profile analysis. 2. Experimental Commercially pure aluminium Al 99.5 (99.5% purity, comparable to AA1050) has been investigated by tensile tests and by total strain controlled fatigue experiments. The material has been used in a CG and in an UFG state. The UFG microstructure was introduced by equal-channel angular pressing (ECAP). Eight ECAP passes using route BC have been applied. For details on the chemical composition and on the ECAP processing of the material, the reader is referred to Ref. [9]. As shown in this reference, the grain size was determined to be 430 nm in the elongated direction and 270 nm perpendicular to it. The CG Al was in a recrystallised condition in case of monotonic tensile tests, with a grain size of about 80 ␮m. In case of the fatigue experiments, the CG Al was recrystallised and afterwards pre-deformed by swaging at room temperature up to a strain of ε = 0.67. The specimens for the tensile tests had a total length of 40 mm with a gauge length of 12.5 mm and a diameter in the gauge section of 6 mm. Strain-rate controlled tensile tests were performed

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on an Instron 4505 testing machine at room temperature. The strain rate varied from 1 × 10−2 to 1 × 10−4 s−1 . The specimens for the fatigue experiments had a total length of 85 mm with a gauge length of 11 mm and a diameter of 5 mm. Total strain controlled fatigue experiments have been performed on a servo-hydraulic testing system MTS 810. The strain rate has been varied from 5 × 10−3 to 5 × 10−5 s−1 . X-ray diffraction (XRD) experiments were performed with a double crystal diffractometer setup, according to Wilkens and Eckert [14], corrected for instrumental peak broadening. The multiple whole profile (MWP) fitting procedure developed by Ung´ar et al. [15] and implemented into a software package by Rib´arik et al. [16] was used for profile analysis. During the last several years, this method became well established for investigating UFG and nanocrystalline materials, see, e.g. [17].

Fig. 2. Cyclic deformation curves of UFG Al 99.5. A total strain amplitude of εtot /2 = 3.8 × 10−3 and three different strain rates have been used.

3. Results and discussion In Fig. 1, stress–strain curves from tensile tests on UFG Al 99.5 and on recrystallised CG Al 99.5 are plotted. The tensile tests have been performed at different strain rates. CG Al shows a high elongation to failure, but the very low strength in the recrystallised condition makes it unsuitable for most engineering applications. For UFG Al a strongly increased strength, and also a relatively high ductility are obtained. As shown in Ref. [18], after an initial drop of the ductility when performing only one or two ECAP passes, the ductility is enhanced as the number of ECAP passes is increased up to eight passes. Only UFG Al exhibits an appreciable strain-rate sensitivity (see Fig. 1). This elevated SRS correlates with the good ductility of the material. For CG Al the applied strain rate has almost no influence on the stress–strain behaviour. The strain-rate sensitivity of UFG Al may play an important role for the rather high ductility of the material, because it will delay the onset of necking and hence prolongs the elongation to failure. It is also worth to note that in compression tests performed in Ref. [9] it was shown that the strain-rate sensitivity of UFG Al is higher, when the strain rate is lowered. This could explain why the ductility of UFG Al

Fig. 1. Stress–strain curves from tensile tests on UFG Al 99.5 and CG Al 99.5. Different strain rates have been applied during the tensile tests; data from Ref. [18].

is higher when the tensile test is performed with a lower strain rate. A deformation mechanism based on thermally activated annihilation of dislocations at the grain boundaries, as suggested in Refs. [2,8,9], might be able to explain the observed SRS of UFG Al. The cyclic deformation behaviour of UFG Al 99.5 is plotted in Fig. 2, exemplarily for tests performed at a constant total strain amplitude of εtot /2 = 3.8 × 10−3 and at three different strain rates of 5 × 10−3 to 5 × 10−5 s−1 . Please note, that the x-axis is plotted in a logarithmic scale to emphasise the beginning of cyclic deformation. All three cyclic deformation curves show cyclic hardening in the early stage of the fatigue tests. This is remarkable, because the material is already severely prestrained by the ECAP process. On the one hand, this observation corresponds to the behaviour found on UFG Al in Ref. [19], and on nanocrystalline nickel [20]. On the other hand, for very pure UFG copper (99.99%), intense cyclic softening due to localized grain coarsening has been reported, starting immediately at the beginning of the test [21]. One possible explanation for this behaviour might be the difference in impurity content of the investigated materials. In the present case, commercially pure Al was used, where the main impurities Fe (0.34%) and Si (0.19%) stabilize the UFG microstructure, as it is known from thermal stability tests. As shown in Fig. 2, the cyclic deformation behaviour of UFG Al is strongly influenced by the strain rate. The obtained stress level at the beginning of the fatigue test is lowest for the smallest strain rate. After about 100 cycles the obtained stress level of all three experiments at the different strain rates becomes nearly equal. Based on these results the following scenario is proposed: starting from a severely pre-deformed material with a rather high dislocation density and due to the change in the deformation path from simple shear deformation to symmetric push–pull deformation spontaneous annihilation of gliding dislocations with dislocations mainly stored at the grain boundaries takes place. This change in the deformation mode will lead to a redistribution of the dislocation structure in the vicinity of a grain boundary and will also lower the dislocation density at the beginning. Dislocations, which are not within the distance of spontaneous annihilation, will only contribute to the annihilation process

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Fig. 3. Coffin–Manson plot of UFG Al 99.5 and of CG Al 99.5 (swaged to a strain of ε = 0.67).

Fig. 4. Development of the crystallite size and dislocation density determined by XRD. Initial as-ECAP state and after fatigue data are shown.

after thermal activation of climb. Hence, the stress level slowly starts to increase, depending on the strain rate. During further cycling the distance to possible annihilation partners increases. Hence, the density of gliding dislocations increases and consequently the material cyclically hardens. Thermal activation at room temperature should be considered, as the experiments were performed at a homologous temperature of about 0.32 and as the grain boundary diffusion coefficient in UFG materials is significantly increased [22]. After about 100 cycles, the potential annihilation partners are too far apart for a further contribution to this process, which results in more or less identical stress levels for strain rates applied. This explanation is contradictory to the often-discussed deformation mechanisms of grain boundary sliding and Coble creep [1]. The obtained results can easily be understood by the above-described dislocation-based deformation mechanism, while the grain boundary sliding mechanism cannot explain the hardening and the reduction of the difference in stress level between the high and low strain rate experiments. Fig. 2 also shows that even the fatigue life of UFG Al is dependent on the strain rate: a higher strain rate leads to a higher fatigue life. By comparison with fatigue experiments on UFG Al under vacuum atmosphere, crack oxidation effects could be excluded to make an important contribution to the fatigue life of the investigated material [23]. However, further investigations on the damage mechanisms have to be done in order to explain the strain rate dependent fatigue lives. The SRS effect including increased fatigue life with increased cyclic strain rate has been found for a large range of total strain amplitudes [23]. In Fig. 3, the data are shown in a Coffin–Manson plot. For comparison, data for a swaged CG Al 99.5 are also shown. The values for the plastic strain amplitude have been evaluated at Nf /2, where Nf is the number of cycles to failure. At a given plastic strain amplitude, swaged CG Al has a strongly increased fatigue life compared to the UFG counterpart. This fits very well to the behaviour predicted in Ref. [24]. Please note that UFG Al reaches much higher stress amplitudes at a given plastic strain amplitude than the CG counterparts. Consequently, in a W¨ohler SN plot, throughout the complete fatigue life regimes

UFG Al is better than the CG counterpart [23]. With respect to the applied strain rate, no SRS effect can be seen in case of swaged CG Al. In case of UFG Al, in the whole investigated range of fatigue lives the higher strain rate shows a higher fatigue life at the same plastic strain amplitude. To investigate the microstructural changes during fatigue of the UFG Al, X-ray peak broadening analysis has been performed. For all investigated total strain amplitudes and strain rates, the results are plotted in Fig. 4. The development of grain size and dislocation density is shown with respect to the total strain amplitude. The crystallite size increases and the dislocation density decreases during fatigue, compared to the as-ECAP state. Please note, that the X-ray profile analysis is also sensitive to low angle grain boundaries and dipole arrangements. Therefore, no conclusion can be drawn, whether coarsening of the crystallite size is due to homogeneous coarsening of all crystallites, or only subgrain boundaries and dipole walls are dissolved. The apparent discrepancy of the grain size obtained by transmission electron microscopy in Ref. [9] and X-ray profile analysis is due to the higher sensitivity of the XRD method to small misorientations, as mentioned above. Regarding the grain coarsening and the development of the dislocation density no clear differences with respect to the strain amplitude and strain rate can be depicted. However, the significant reduction of dislocation density during cyclic deformation by approximately one order of magnitude fits well to the above-described scenario of thermally activated annihilation of dislocations at the grain boundaries. 4. Conclusions Tensile tests and total strain controlled fatigue experiments show, that the deformation behaviour of UFG Al is dependent on the applied strain rate. This is not the case for CG Al with conventional grain size. In tensile tests on UFG Al, strength decreases and ductility increases with decreasing strain rate. Fatigue experiments have shown that the fatigue life of UFG Al increases with increasing strain rate. Based on the strain-rate sensitivity, the hardening in the beginning of fatigue experiments and on the reduction of dislocation density during fatigue,

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dislocation movement combined with thermally activated annihilation of dislocations at the grain boundaries has been suggested as the dominating deformation mechanism for UFG Al at room temperature. Acknowledgement The authors acknowledge the financial support by the Deutsche Forschungsgemeinschaft DFG, Contract no. GO 741/10-1, within the Research Unit Program “Mechanische Eigenschaften und Grenzfl¨achen ultrafeink¨orniger Werkstoffe”. References [1] [2] [3] [4] [5] [6] [7] [8] [9]

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