Origin of the influence of Cu or Ag micro-additions on the age hardening behavior of ultrafine-grained Al-Mg-Si alloys

Journal of Alloys and Compounds 710 (2017) 199e204

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Origin of the influence of Cu or Ag micro-additions on the age hardening behavior of ultrafine-grained Al-Mg-Si alloys Xavier Sauvage a, *, Seungwon Lee b, c, d, **, Kenji Matsuda d, Zenji Horita b, c Normandie Univ, UNIROUEN, INSA Rouen, CNRS, Groupe de Physique des Mat eriaux, 76000 Rouen, France Department of Materials Science and Engineering, Faculty of Engineering, Kyushu University, Fukuoka 819-0395, Japan c WPI, International Institute for Carbon-Neutral Energy Research (I2CNER), Kyushu University, Fukuoka 819-0395, Japan d Graduate School of Science and Engineering for Research, University of Toyama, Toyama 930-8555, Japan a

b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 9 January 2017 Received in revised form 20 March 2017 Accepted 21 March 2017 Available online 22 March 2017

The age hardening behavior of ultrafine-grained Al-Mg-Si alloys with micro-additions of Cu and Ag has been investigated with a special emphasis on the relationship between microstructural features and hardness evolutions. Using transmission electron microscopy, it is shown that the higher hardness and thermal stability induced by a small amount of Cu is due to clustering during grain refinement by severe plastic deformation. It is also demonstrated that the addition of Ag is beneficial for the thermal stability of the ultrafine grain structure as it segregates along grain boundaries and significantly reduces their mobility. © 2017 Elsevier B.V. All rights reserved.

Keywords: Aluminum alloys Ultrafine grains Precipitation hardening Severe plastic deformation Grain boundary segregations

1. Introduction Aluminum-based Al-Mg-Si alloys are typical age-hardenable alloys having medium strength and good corrosion resistance with excellent formability. A smart application of the alloys with the age-hardenability may be for automotive body panels when they are subjected to bake-hardening as a concurrent process with the panel painting [1]. Nevertheless, higher strength is always desired as it can allow reducing the weight of automotive parts. Because solid-solution hardening is not much applicable to alloys of the Al-Mg-Si system due to the given compositions, grain refinement can be a promising means for strengthening of the alloys. Recently, there is an approach making use of simultaneous strengthening due to grain refinement and fine dispersion of precipitates [2e5]. Although grain refinement is well achieved to submicrometer sizes using severe plastic deformation (SPD)

* Corresponding author. ** Corresponding author. Current address: Graduate School of Science and Engineering for Research, University of Toyama, Toyama 930-8555, Japan. E-mail addresses: [email protected] (X. Sauvage), [email protected] (S. Lee). http://dx.doi.org/10.1016/j.jallcom.2017.03.250 0925-8388/© 2017 Elsevier B.V. All rights reserved.

processes [6], it has been demonstrated that obtaining homogeneous fine precipitation in such an ultrafine-grained structure is challenging because of heterogeneous precipitation along structural defects such as grain boundaries [7]. Hirosawa et al. [8] proposed three strategies to overcome such a difficulty, which include (1) lowering aging temperature, (2) alloying of the elements such as Cu and Ag and (3) making use of spinodal decomposition. For the present alloys (Al-Mg-Si), the third strategy cannot apply since typical compositions are out of the spinodal domain. The first strategy has been experienced with somewhat limited success [3,4,7,9], while Akama et al. [10] and Watanabe et al. [11] reported that additions of Cu and Ag (second strategy) are effective to increase not only the simultaneous strengthening due to grain refinement and fine precipitation but also the thermal stability with aging time. They also reported that the Cu addition is more effective for the enhancement of age hardening than the Ag addition despite both additions show similar behavior for the thermal stability with aging. It should be noted that Cu addition improves the thermal stability (evaluated by hardness measurements) even in a coarse grained structure [12]. In this study, we have carefully examined the role of Cu or Ag micro-additions on the age hardening behavior of ultrafine-grained Al-Mg-Si alloys. Thanks to high resolution microscopy, a special

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emphasis is given on atomic scale mechanisms that operate during severe plastic deformation and natural and artificial aging. 2. Experimental procedures Alloys with compositions similar to the material investigated in a previous work [10] have been adopted for the present study. Regarding the Mg2Si phase stoichiometry, they contain some excess Mg: Al-1.0 mass%Mg2Si-0.4mass%Mg (equivalent to Al-1.03mass% Mg-0.37mass%Si). Besides, part of the aluminum has been substituted with 0.5mass%Cu or 0.5mass%Ag. In the following, they are designated, respectively, as exMg, exMg þ Cu and exMg þ Ag. All alloys were subjected to solution treatment (hereafter designated as ST) at 848 K for 3.6 ks followed by quenching in ice water. Then, High Pressure Torsion (HPT) was conducted at room temperature on disk shape samples (10 mm diameter and 0.85 mm thickness) under a pressure of 6 GPa and up to 5 turns with a rotation speed of 1 rpm. Subsequently, the aging behavior was examined at two different temperatures: 373 K (hereafter called artificial aging: AA) and at room temperature (hereafter called natural aging: NA). More technical details are found in Ref. [10]. The Vickers microhardness of samples was measured for a total aging time up to 1200 ks (14 days) with a Mitsutoyo HM-102 hardness tester (load of 100 g, dwell time 15s). It should be noted that microhardness data of AA materials were simply taken from earlier work [10]. For microstructural characterizations, materials were stored at 243 K after HPT (during a maximum period of 2 months) to avoid natural aging of the as-HPT state. Then, samples remained not more than a week at room temperature to allow specimen preparation and Transmission Electron Microscopy (TEM) observations. TEM samples have been prepared from 3 mm disks that were punched out at 2.5 mm (±0.5 mm) from the center of HPT discs. Then, electron transparency was reached using a twin-jet electropolishing facility (electrolyte: 70 vol% CH3OH-30 vol% HNO3, temperature 243 K, voltage 17 V). TEM observations were carried out with JEOL ARM200F microscope operated at 200 kV in conventional mode (parallel beam) to record bright-field images and selected area electron diffraction (SAED) patterns. Bright-field (BF) and high angle annular dark field (HAADF) images were also recorded in the scanning mode (STEM) with a probe size of 0.1 nm and Energy Dispersive X-ray Spectroscopy (EDS) was performed with a JEOL JED2300 detector with a probe size of 0.2 nm. 3. Results Fig. 1 shows the hardness plotted against the aging time for all samples processed by HPT including the ST samples without HPT processing. Here, the hardness values are plotted for the averages obtained from the values measured between 1 and 4 mm from the disk center where the hardness remains constant as it is saturated to a steady state [10]. Several interesting features should be highlighted on this plot: i) For all alloys without HPT processing (coarse grains), the micro-hardness is similar in the as-ST state, then it increases significantly during natural aging but this increase is more pronounced for alloys with a micro-addition of Cu and Ag (DHV~30 in the micro-alloyed states against DHV~20 in the state without alloying after 300 h), ii) After HPT processing, the hardness of the exMg þ Cu alloy is significantly higher than that of the two other alloys (~170 H V against ~150 H V) and this higher level is maintained during natural or artificial aging, iii) After HPT processing, and for each alloy, the maximum hardness achieved by NA or AA is similar, iv) under AA condition, the hardness peak of the exMg alloy is reached for only several minutes (~0.1 h) while it is achieved in about an hour for exMg þ Cu and the exMg þ Ag alloys,

Fig. 1. Evolution of the Vickers microhardness of the exMg, ex-Mg þ Cu and exMg þ Ag alloys after solution treatment followed by natural aging (ST þ NA), after HPT processing followed by natural aging (HPT þ NA) and after HPT processing followed by artificial aging (HPT þ AA).

v) after long time artificial aging, only the exMg alloy exhibits a significant loss in hardness as compared to the as-HPT state (DHV ~ -20), indicating a lower thermal stability. To understand these particular features, microstructures were carefully examined in the as-HPT state for the exMg þ Cu (Fig. 2) and exMg þ Ag (Fig. 3). As shown on bright-field images (Figs. 2(a) and 3(a)), and as expected, the severe plastic deformation produced an ultrafine grain structure with a similar grain size for both alloys (ranging from 180 to 280 nm). Even in high resolution mode (images not shown here), no precipitate or cluster could be imaged but in the exMg þ Cu alloy, SAED patterns taken in low zone axis (Fig. 2 (b)) revealed some superstructure reflections that correspond to (21e31) reflections of the Q0 phase [13]. As shown on the Fast Fourier Transformation (FFT) set in Fig. 2(c), this feature was also confirmed by high resolution HAADF STEM imaging. Such nanoscaled clusters could not be observed in the exMg þ Ag alloy (Fig. 3(b)), however STEM HAADF images revealed a bright contrast along some GBs (arrows on Fig. 3(c)). Other contrast changes on this image are attributed to some local variation of the foil thickness (thinner area on the top provides less signal and makes the image darker). Then, bright features at GBs are linked to some local composition variation, namely segregation. To confirm this point, EDS line profile analyses were carried out across GBs. As shown on the plot in Fig. 3(d), a local Ag enrichment at GB is clearly exhibited and is attributed to strain induced segregation [14e17]. As shown on Figs. 4(a) and 5(a), after artificial aging the grain size is still in the sub-micrometer range and did not change significantly (Table 1). Besides, the ordered clusters identified in the as-HPT state for the exMg þ Cu alloy are still present inside grains (see SAED Fig. 4(b)). Unfortunately it has not been possible to estimate if the number density has changed as compared to the asHPT material. However, as shown on the STEM HAADF image (arrowed on Fig. 4(c)), few relatively large precipitates have heterogeneously nucleated along GBs and triple lines. They are about 10e30 nm in size and EDS analysis (see the map set in Fig. 4(c)) clearly revealed that they contain a significant amount of Mg and Si. Such precipitates could not be observed in the exMg þ Ag alloy where most, if not all, Ag segregations along GBs seem to be preserved during the AA (see Fig. 5(b) and (c)).

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Fig. 2. exMg þ Cu alloy processed by HPT and observed by STEM in the as-HPT state. (a) Low magnification bright-field image showing the UFG structure; (b) SAED pattern taken in a grain oriented in (001)Alfcc zone axis and showing some reflections (circled) that could be attributed to the (21e31) reflection of the Q0 phase [13]; (c) HAADF HRSTEM image and corresponding FFT (inset) giving another evidence of local ordering and clustering.

4. Discussion Experimental data clearly reveal two specific features related to the ultrafine grained structures achieved by severe plastic deformation in the considered alloys: dynamic precipitation (exMg þ Cu) and GB segregations (exMg þ Ag). Dynamic

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precipitation in Al alloys processed by SPD at room temperature has already been reported in AlMgSi alloys [7] and is often attributed to an enhanced atomic mobility resulting from a high concentration of strain induced vacancies. It is also important to note that vacancies play a major role in the early stage of clustering in binaries Al-Mg, Al-Si, Al-Cu and more complex alloys [18e20]. Besides, a higher atomic mobility combined with a relatively high proportion of GB areas may also promote the formation of GB segregations if there is a positive segregation enthalpy for the considered solute. It has also been proposed that non-equilibrium segregations may also progressively form during SPD since GBs are vacancy sinks [15,17]. Interestingly, micro-additions of Cu and Ag lead to different phenomenon, namely clustering for Cu and GB segregation for Ag. The binding energies of Cu and Ag with vacancies are relatively close [21], they both exhibit a positive mismatch in the fcc lattice of Al but the diffusivity of Ag in Al is significantly higher than Cu [22]. This difference in atomic mobility together with a different enthalpy of segregation along non-equilibrium grain boundaries [17] is probably the reason of SPD induced GB segregation of Ag. One should also note the specificity of Q0 clusters that nucleated in the exMg þ Cu alloy. They are so small that they could not be imaged while they usually appear as lath shaped precipitates during artificial aging [13]. It is believed that such lath shaped Q0 precipitates cannot grow in the present material because of the large strain imposed during HPT. In any case, both GB segregations and nanoscaled clusters affect the properties and the thermal stability of the considered alloys. They all exhibit a similar hardness in the solution treated state (non-deformed), but their natural aging response (at room temperature) is somewhat different as highlighted on Fig. 1. Indeed, it clearly appears that micro-additions of Cu and Ag lead to a significantly more pronounced hardening (see Fig. 1). This is not surprising and it could be attributed to Cu/Mg clustering in the exMg þ Cu alloy [23,24] and to Ag GP zone formation in the exMg þ Ag alloy [25]. In the as-HPT state, the hardness of the exMg þ Cu alloy is significantly higher than the exMg or exMg þ Ag (HV~170 against ~150). Our TEM observations revealed that the grain size is relatively similar for all alloys (see Figs. 2(a) and 3(a)), thus this peculiarity can be directly attributed to nanoscaled clusters that nucleated during SPD and that have been identified only in the exMg þ Cu alloy (see Figs. 2(b) and 3(b)). During natural aging of HPT-processed alloys, the hardness increases moderately, at least 50% less than in the coarse grain state (Fig. 1). Since the grain size does not significantly change during natural aging, this higher stability is probably the result of solute elements redistribution that occurred during SPD. In the exMg þ Cu alloy, most of clusters are already formed after HPT processing (Fig. 2(b)). In the exMg þ Ag alloy most of Ag has segregated along GBs (Fig. 3(c)) and consequently a lower fraction of GP zones should nucleate. This is also consistent with the stable gap in hardness between exMg þ Cu and exMg þ Ag HPT alloys during natural aging. During artificial aging at 373 K, the HV peak is achieved in less than an hour for any alloy and then a progressive and constant decrease is observed (Fig. 1). This decrease is however significantly more pronounced for the ex-Mg alloy which can be attributed to a more pronounced grain growth as revealed by our observations (Table 1). This feature points out the benefit of micro-addition of Cu or Ag in UFG AlMgSi alloys. In the exMg þ Ag alloy, grain boundary segregation of Ag stabilizes the UFG structure [26e28], while in the exMg þ Cu alloy, GBs are pinned by Q0 phase particles reducing their mobility. On a practical point view, it is clear that micro-additions of Cu is interesting to achieve a higher strength and a better thermal

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Fig. 4. exMg þ Cu alloy processed by HPT and observed by STEM in the A.A. condition. a) Low magnification bright-field image showing the UFG structure; (b) SAED pattern taken in a grain oriented in (001)Alfcc zone axis and showing superstructure reflections that could be attributed to the (21e31) reflection of the Q0 phase [13]; (c) HAADF STEM image showing few bright nanoscaled particles (arrowed). The EDS map (inset, Al blue, Si green, Mg red) indicates that these particles contain a significant amount of Mg and Si. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Fig. 3. exMg þ Ag alloy processed by HPT and observed by STEM in the as-HPT state. (a) Low magnification bright-field image showing the UFG structure; (b) SAED pattern taken in a grain oriented in (110)Alfcc zone axis. Contrary to the exMg þ Cu alloy, no reflections that could be attributed to another phase are detected; (c) HAADF STEM

image showing bright grain boundaries indicated by white arrows; (d) EDS line profile recorded across the boundary indicated by the black arrow on (c), it confirms that the bright contrast is related to Ag segregation.

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Table 1 Mean grain size evolution during aging at 373 K for the exMg and exMg þ Cu alloys (measured from TEM images). Aging time at 373 K

exMg

exMg þ Cu

<10 min 103 min 104 min

233 ± 54 nm 252 ± 64 nm 263 ± 73 nm

223 ± 57 nm 213 ± 57 nm 236 ± 62 nm

5. Conclusions i) A higher hardness in the as-processed state was achieved with micro-addition of Cu due to cluster formation during grain refinement by HPT processing. This higher hardness, as compared to the material without Cu or with micro-addition of Ag, is maintained during artificial aging at 373 K because these clusters do not dissolve and inhibit the grain growth. ii) The micro-addition of Ag improves the thermal stability since this element segregates at grain boundaries. It reduces their mobility and this leads a better thermal stability of the ultrafinegrained structure. However, grain interiors being depleted in Ag, the natural aging response is significantly reduced as compared to the coarse-grained material. Acknowledgments The authors would like to thank the Agence Nationale de la Recherche (ANR) for financial support (PRASA project - ANR-15CE08-0029). This work was also supported in part by Japan Science and Technology Agency (JST) under Collaborative Research Based on Industrial Demand “Heterogeneous Structure Control: Towards Innovative Development of Metallic Structural Materials”, in part by the Light Metals Educational Foundation of Japan, and in part by a Grant-in-Aid from the MEXT, Japan, for Scientific Research (S) (No.26220909). The HPT process was carried out in the International Research Center on Giant Straining for Advanced Materials (IRC-GSAM) at Kyushu University. References

Fig. 5. a) Low magnification STEM bright-field image of the exMg þ Ag alloy processed by HPT þ AA showing the UFG structure; (b) Low magnification HAADF STEM image showing Ag segregation along GBs (arrowed); (c) HAADF high resolution STEM image showing a grain in (001) zone axis (bottom left) and a GB with a bright contrast indicating some Ag segregation (arrowed).

stability for UFG AlMgSi alloys. However careful attention must be paid as it could reduce the corrosion resistance [1,29]. Microaddition of Ag is also beneficial on the thermal stability thanks to GB segregation of Ag, but the peak hardness achievable is significantly lower. In any case, the temperature and/or time for bake hardening has to be reduced as far as such an UFG structure are considered since we can still take advantage of simultaneous strengthening due to grain refinement and fine precipitation.

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