Deformation Mechanism during Dynamic Loading of ...

Deformation Mechanism during Dynamic Loading of an Additively Manufactured AlSi10Mg_200C Amir Hadadzadeh1, 2,*, Babak Shalchi Amirkhiz2, Jian Li2, Akindele Odeshi3, Mohsen Mohammadi1

1- Marine Additive Manufacturing Centre of Excellence (MAMCE), University of New Brunswick, Fredericton, NB, E3B 5A1, Canada 2- CanmetMATERIALS, Natural Resources Canada, 183 Longwood Road South, Hamilton, ON, L8P 0A5, Canada 3- Department of Mechanical Engineering, University of Saskatchewan, Saskatoon, SK, S7N 5A9, Canada *Corresponding author: Amir Hadadzadeh, [email protected]

Abstract A selective laser melted AlSi10Mg_200C was subjected to dynamic loading at two strain rates. At 1680 s-1, an ultrafine substructure formed inside the α-Al dendrites due to dynamic recovery (DRV). At 4300 s-1, continuous dynamic recrystallization (CDRX) occurred, followed by DRV inside the DRX grains with development of nanoscale subgrains. Keywords: Selective laser melting; Transmission electron microscopy (TEM); Electron backscattering diffraction (EBSD); Dynamic recovery (DRV); Dynamic recrystallization (DRX).

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1. Introduction Additive manufacturing (AM) is an incremental layer-on-layer joining process where a threedimensional part is fabricated from a digital CAD model directly from a powder or wire material feedstock [1]. The feedstock is selectively melted using an energy source which is usually laser, electron beam, electric arc or ultrasonic vibration to create the layers [2, 3, 4]. Amongst the available AM processes, selective laser melting (SLM) has shown a promising trend for fabricating metallic parts due to its design freedom, shorter production cycle (in comparison to the conventional manufacturing routes), and significant cost savings [5, 6, 7, 8]. Moreover, ultrafine, metastable and gradient microstructures formed due to fast solidification rates (103-106 °C/s) during SLM process [9]. SLM has been used widely for processing of aluminum alloy powders [10, 11, 12, 13]. In particular, SLM of AlSi10Mg has received increasing attention due to its light weight, high specific strength, and excellent corrosion resistance [14, 15, 16] for applications in aerospace, automotive and marine industries. It has been shown that SLMAlSi10Mg parts possess higher strength than the as-cast [17, 18, 19] or conventional powder metallurgy [20] counterparts under quasi-static loading conditions. Similarly, under high strain rate loading conditions, SLM-AlSi10Mg exhibited higher strength than sand cast AlSi10Mg [21]. It is important to investigate the high strain rate deformation of AlSi10Mg alloy since unexpected impact situations in service conditions (e.g. vehicle crash) can expose this alloy to dynamic loading. While dynamic impact loading of conventional metallic alloys leads to formation of adiabatic shear bands (ASBs) and strain localization [22, 23, 24], the deformation mechanism during high strain rate deformation of SLM-AlSi10Mg is not clear yet. Therefore, dynamic loading of an SLM-AlSi10Mg_200C alloy was investigated in the current study and the deformed microstructure was analyzed using electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM). This study sheds light on high strain rate deformation mechanism, which is used to explain the correlation between microstructure and mechanical behavior of the SLM-AlSi10Mg_200C alloy.

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2. Experimental procedure AlSi10Mg_200C virgin powder with an average particle size of 9 ± 7 µm and chemical composition of Al-10wt.%Si-0.33wt.%Mg-0.55wt.%Fe [19] was used to fabricate vertical rodshaped samples with 12 mm diameter and 12 cm height using an EOS M290 machine. The dimensions of the building plate were 250 mm × 250 mm × 325 mm and the machine was equipped with a 400 W Yb-fiber laser and a beam spot size of 100 µm. The building platform was preheated to 200°C and kept at this temperature to minimize residual stresses. The laser power, scanning speed, hatching distance, and powder layer thickness was 370 W, 1300 mm/s, 190 µm and 30 µm, respectively. The compressive dynamic shock loading test were performed using a Split Hopkinson Pressure Bar (SHPB) on cylindrical samples with 9.5 mm diameter and 10.5 mm height along the building direction. Details of shock loading tests set-up can be found in [24]. Microstructure of the as-built sample was analyzed using optical microscopy (OM), scanning electron microscopy (SEM) and EBSD. The EBSD scans were conducted in a field emission gun scanning electron microscope (FEG-SEM) (FEI Nova NanoSEM-650) equipped with an OIM 6.2 EBSD system (EDAX), with a step size of 0.12 µm. To analyze the solidification behavior of the alloy, a computational thermodynamics database, FactSage™ with the FTlite database [25] was utilized. As-built and deformed samples were analyzed using an FEI Tecnai Osiris TEM equipped with a 200 keV X-FEG gun. The Super-EDS X-ray detection system combined with the high current density electron beam in the scanning mode (STEM) was also utilized to analyze the precipitates. Spatial resolutions in the order of 1 nm were obtained during EDS elemental mapping by using a sub-nanometer electron probe.

3. Results and discussion Fig. 1(a) and (b) show the OM and SEM images of the as-built sample from the side view, where the cross-sections of the melt pools are visible. The overlapped melt pools are almost halfcylindrical in shape [26], which is the typical so-called fish-scale structure in the SLMAlSi10Mg parts [27]. Referring to the SEM image of the as-built sample (Fig. 1(b)), the center of the melt pools consists of ultrafine cellular dendritic microstructure with an average size of ~0.6 µm. Each cell consisted of α-Al with eutectic Si boundaries. This microstructure has been 3

reported for SLM of AlSi10Mg in the previous studies [20, 28]. At the boundaries of the melt pools, partially melted heat affected zone (HAZ) is observed [29]. Above the HAZ a coarseelongated columnar structure along the building direction is observed as a result of epitaxial growth [30]. The average cell size in the HAZ and coarse columnar zone is ~1.5 µm and ~1.03 µm, respectively. Beneath the HAZ, a finer columnar structure is observed with an average cell size of ~0.82 µm. Variation of cell size in the microstructure of SLM-AlSi10Mg_200C is mainly due to the change in thermal gradient during the SLM process [31]. Fig. 1(c) shows the inverse pole figure (IPF-Z) obtained from EBSD, with detailed information on grain size and orientation. Similar to the cell structure, a heterogeneous grain structure is observed for the material; however, the grain size is at least one order of magnitude larger than the cell size, which is consistent with the previous studies [20]. In other words, each grain is consisted of several cells. The width of the elongated grains can reach ~20 µm, while the equiaxed grains at the vicinity of the melt pool boundaries are much finer in size [32]. The majority of the elongated grains developed {100} texture which is consistent with the results of previous studies [26, 32, 33]. Fig. 1(d) illustrates the STEM bright field (STEM-BF) microstructure of the as-built material along with the EDS elemental maps of Al, Si, Mg and Fe superimposed on the image (Fig. 1(e)). The bright areas in the STEM-BF image are primary α-Al phase (shown by green in the map), surrounded by mainly eutectic Si (red in the map). Other intermetallic phases consisted of Mg and Fe are also observed.

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Fig. 1- (a) OM, (b) SEM and (c) inverse pole figure (IPF-Z) images of the as-built AlSi10Mg_200C sample from side view. (d) STEM-BF microstructure of the as-built sample and (e) the corresponding TEM EDS compositional map superimposed on the STEM-BF. z-axis represents the building direction.

The isopleth of Al-0.33wt.%Mg-0.55wt.%Fe-xSi system at Al-rich corner predicted by FactSage™ is shown in Fig. 2(a). The chemical composition of AlSi10Mg_200C used in the current study is shown by a dashed red line on the graph. The possible final phases after solidification of AlSi10Mg_200C include α-Al (primary and eutectic), Si (eutectic), and Al3FeSi and Al8FeMg3Si6 intermetallics. The amount of phases after equilibrium solidification was also predicted and shown in Fig. 2(b). Although the solidification during SLM occurs under nonequilibrium conditions, reheating of the substrate layers during deposition of the subsequent layers [11, 34] can elevate the diffusion of the alloying elements and promote the intermetallic phases towards the equilibrium conditions [35]. Referring to Fig. 2(b), the amount of Al8FeMg3Si6 and Al3FeSi phases are 2.3 wt.% and 0.9 wt.%, respectively. The dominance of 5

Al8FeMg3Si6 phase over Al3FeSi is evident from TEM images in Fig. 1(e) and Fig. 2(d)-(g). Fig. 2(c) shows detailed features of a primary α-Al tilted so that the e-beam was parallel to Al [011] zone axis. A network of entangled dislocations in the α-Al dendrites is observed that can be a result of rapid solidification during SLM and evolution of residual stresses [33].

Fig. 2- (a) Al-Si-Mg-Fe isopleth calculated by FactSage™ with 0.33wt.%Mg-0.55wt.%Fe at Alrich corner, (b) phase amount during equilibrium solidification of AlSi10Mg_200C, (c) TEM-BF image of as-built sample with the beam direction parallel to the Al [011] zone axis and (d)-(g) corresponding EDS mappings.

Fig. 3(a) shows the schematic drawing of the as-built rod and the orientation of dynamic loaded sample and direction of dynamic loading. The loading direction was along the building direction. The strain rates used for shock loading are shown in Fig. 3(b). Two strain rates were employed to impact the SLM-AlSi10Mg_200C alloy with maximum level of 2200 s-1 and 5400 s-1. Since the strain rates were not constant over the course of deformation, an average value for each curve 2

  d

was calculated using avg 

1

 2  1

. The corresponding calculated average strain rates are 1680

s-1 and 4300 s-1, respectively. The dynamic stress-strain curves of samples are shown in Fig. 3(c). 6

The sample deformed at avg =1680 s-1 exhibits strain hardening beyond the yield stress, followed by reaching the peak flow stress (PFS). By increasing the average strain rate to 4300 s1

, two distinct peak stresses appear in the flow curve. In general, the flow behavior during high

strain rate dynamic loading represents the complicated competition between strain hardening and thermal softening [22, 24]. Moreover, it has been reported that compressive dynamic loading of metallic materials results in formation of ASBs and dynamic recrystallization (DRX) inside the shear bands [36]. However, formation of shear bands was not observed for the SLMAlSi10Mg_200C alloy in the current study. Fig. 3(d) and (f) show the representative EBSD-IPFZ maps of deformed samples at avg =1680 s-1 and 4300 s-1, respectively. Grain structure and orientation of the deformed samples are similar to those of the as-built sample (Fig. 1(c)), i.e. a columnar structure with {100} texture. Therefore, impact loading of the material did not alter the texture significantly [37]. The corresponding Kernel average misorientation (KAM) maps of Fig. 3(d) and (f) are shown in Fig. 3(e) and (g). KAM values represent the local misorientation in terms of dislocation density distribution in different grains [38]. As a general rule, KAM > 1° represents deformed grains (due to high density of dislocations) and KAM < 1° is an indication of DRX grains [39]. KAM maps of the deformed samples did not change by increasing the strain rate and the majority of KAM values was between 0.5-1.5°, as shown in Fig. 3(h). Such a characteristic could represent occurrence of DRX in the sample; however, formation of DRX grains is not concluded from the EBSD-IPF maps.

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Fig. 3- (a) Schematic representation of the as-built rod, dynamic loading sample and loading direction, (b) strain rates employed for impact loading, (c) corresponding stress-strain curves, (d) and (e) IPF-Z and KAM maps of deformed sample with avg =1680 s-1, (f) and (g) IPF-Z and KAM maps of deformed sample with avg =4300 s-1 and (h) KAM profile of the deformed samples.

Details of microstructural changes in the deformed samples were analyzed using TEM and the results are presented in Fig. 4. Fig. 4(a) shows the TEM-BF of the as-built sample as a reference for comparison with the deformed microstructures. As mentioned, a network of entangled dislocations exists in the α-Al dendrites in the as-built sample. Dynamic loading of the material with avg =1680 s-1 resulted in formation of a subgrain network, consisting of low angle grain boundaries (LAGBs) featured by a network of accumulated dislocations (Fig. 4(b)), which was formed as a result of dynamic recovery (DRV) [40]. By increasing the average strain rate to 4300 s-1, LAGBs progressively transformed to high angle grain boundaries (HAGBs) and DRX grains evolved, as seen in Fig. 4(c). This DRX mechanism, which is controlled by conversion of substructures to HAGBs, is known as continuous DRX (CDRX) [41]. Meanwhile, LAGBs developed inside the DRX grains and nanoscale subgrains formed. Details of DRV and DRX evolution in the deformed samples are shown in Fig. 4(d)-(h) for avg =1680 s-1 and Fig. 4(i)-(m) 8

avg =4300 s-1. In both cases, the microstructural evolutions occurred inside the cell structures and morphology of eutectic Si (cell boundaries) was not altered during dynamic loading. Plastic deformation of SLM-AlSi10Mg_200C at high strain rate led to generation of new dislocations as a result of dislocation glide. Meanwhile, impact loading of metallic materials results in heat generation and temperature rise in the sample [22]. Increase of temperature during high strain rate deformation promoted the dislocations movement. At avg =1680 s-1, movement of dislocations resulted in formation of accumulated dislocation networks and dislocation-free subgrains. In fact, by development of accumulated dislocation networks, LAGBs formed and the rest of dislocations entrapped in these networks [42]. From the Si EDS mapping of the deformed sample with avg =1680 s-1 (Fig. 4(f)) it appears that in addition to the eutectic Si network, Si precipitates (shown by white arrows) also act as favorable locations for entanglement of mobile dislocations and formation of accumulated dislocation networks. Therefore, the strain hardening in the flow curve of sample impacted with avg =1680 s-1 (Fig. 3(c)) was a result of dislocation generation up to the PFS, followed by softening due to substructure formation and DRV. By increasing the strain rate to avg =4300 s-1, the level of generated heat and temperature rise in the sample increased [43]. Therefore, formation of accumulated dislocation networks, LAGBs and DRV structure accelerated, followed by CDRX and development of DRX grains. It seems that the first PFS and the consequent softening (at ε≈0.04) appeared in the dynamic stress-strain curve of this sample (Fig. 3(c)) is a result of completion of CDRX. By continuing the deformation, dislocations generated in the DRX grains and strain hardening occurred in the material up to the second PFS (at ε≈0.14). In addition, beyond this strain the LAGBs and nanoscale subgrains formed in the DRX grains and the final softening was a result of DRV in the DRX grains. Similar to the sample deformed at avg =1680 s-1, both eutectic Si networks and Si precipitates (shown by the white arrows in Fig. 4(k)), acted as suitable locations for formation of the HAGBs.

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Fig. 4- TEM-BF images of (a) as-built and deformed samples with (b) avg =1680 s-1 and (c) avg =4300 s-1. Grain structure evolution in α-Al dendrites and the corresponding EDS mappings of deformed samples with (d)-(h) avg =1680 s-1 and (i)-(m) avg =4300 s-1. All images were taken with the beam direction along the Al [011] zone axis orientation.

4. Conclusions In summary, high strain rate dynamic loading of an SLM-AlSi10Mg_200C alloy was studied using SHPB at two strain rates level, i.e. avg =1680 s-1 and 4300 s-1. The microstructure of the as-built alloy consisted of heterogenous cell structures, where the boundaries were made of eutectic Si phase. The grain size was one order of magnitude higher than the cell size. In contrast to dynamic loading of the conventional alloys, no ASBs and strain localization was observed for this alloy. Dynamic loading of the sample at avg =1680 s-1 led to the development of ultrafine substructure and DRV. By increasing the strain rate to avg =4300 s-1, the substructure converted 10

to a DRXed structure through CDRX mechanism followed by the development of nanoscale subgrains in the DRX grains as a result of DRV.

Acknowledgements The authors would like to acknowledge Natural Sciences and Engineering Research Council of Canada (NSERC) project number RGPIN-2016-04221 and New Brunswick Innovation Foundation project number (NBIF)-RIF2017-071 for the financial support of this work. The authors would also like to acknowledge AMM for fabricating the SLM samples, Dr. Mark Kozdras at CanmetMATERIALS for facilitating the research and Renata Zavadil for her cooperation for OM and SEM analyzes and Catherine Bibby for TEM sample preparations.

References

[1] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals. Acta Mater. 117 (2016) 371-392. [2] T. DebRoy, H.L. Wei, J.S. Zuback, T. Mukherjee, J.W. Elmer, J.O. Milewski, A.M. Beese, A. Wilson-Heid, A. De, W. Zhang, Additive manufacturing of metallic components – Process, structure and properties, Prog. Mater. Sci. 92 (2018) 112–224. [3] H.T. Fujii, S. Shimizu, Y.S. Sato, H. Kokawa, High-strain-rate deformation in ultrasonic additive manufacturing. Scr. Mater. 135 (2017) 125–129. [4] M.N. Gussev, N. Sridharan, Z. Thompson, K.A. Terrani, S.S. Babu, Influence of hot isostatic pressing on the performance of aluminum alloy fabricated by ultrasonic additive manufacturing. Scr. Mater. 145 (2018) 33–36. [5] D.D. Gu, W. Meiners, K. Wissenbach, R. Poprawe, Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int. Mater. Rev. 57(3) (2012) 133-164. [6] J.H. Martin, B.D. Yahata, J.M. Hundley, J.A. Mayer, T.A. Schaedler, T.M. Pollock, 3D printing of highstrength aluminium alloys. Nature 549 (2017) 365-369. [7] L. Li, B. Post, V. Kunc, A.M. Elliott, M.P. Paranthaman, Additive manufacturing of near-net-shape bonded magnets: Prospects and challenges. Scr. Mater. 135 (2017) 100–104. [8] B.H. Jared, M.A. Aguilo, L.L. Beghini, B.L. Boyce, B.W. Clark, A. Cook, B.J. Kaehr, J. Robbins, Additive manufacturing: Toward holistic design. Scr. Mater. 135 (2017) 141–147. [9] X.P. Li, G. Ji, Z. Chen, A. Addad, Y. Wu, H.W. Wang, J. Vleugels, J. Van Humbeeck, J.P. Kruth, Selective laser melting of nano-TiB2 decorated AlSi10Mg alloy with high fracture strength and ductility. Acta Mater. 129 (2017) 183–19. [10] K.G. Prashanth, S. Scudino, H.J. Klauss, K.B. Surreddi, L. Löber, Z. Wang, A.K. Chaubey, U. Kühn, J. Eckert, Microstructure and mechanical properties of Al–12Si produced by selective laser melting: Effect of heat treatment, Mater. Sci. Eng. A 590 (2014) 153–160.

11

[11] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, A review on selective laser sintering/melting (SLS/SLM) of aluminium alloy powders: Processing, microstructure, and properties. Prog. Mater. Sci. 74 (2015) 401–477. [12] K.G. Prashanth, S. Scudino, J. Eckert, Defining the tensile properties of Al-12Si parts produced by selective laser melting. Acta Mater. 126 (2017) 25-35. [13] A.B. Spierings, K. Dawson, T. Heeling, P.J. Uggowitzer, R. Schäublin, F. Palm, K.Wegener, Microstructural features of Sc- and Zr-modified Al-Mg alloys processed by selective laser melting. Mater. Des. 115 (2017) 5263. [14] N. Read, W. Wang, K. Essa, M.M. Attallah, Selective laser melting of AlSi10Mg alloy: Process optimisation and mechanical properties development. Mater. Des. 65 (2015) 417–424. [15] W. Li, S. Li, J. Liu, A. Zhang, Y. Zhou, Q. Wei, C. Yan, Y. Shi, Effect of heat treatment on AlSi10Mg alloy fabricated by selective laser melting: Microstructure evolution, mechanical properties and fracture mechanism, Mater. Sci. Eng. A 663 (2016) 116–125. [16] P. Delroisse, P.J. Jacques, E. Maire, O. Rigo, A. Simar, Effect of strut orientation on the microstructure heterogeneities in AlSi10Mg lattices processed by selective laser melting. Scr. Mater. 141 (2017) 32–35. [17] N.T. Aboulkhair, I. Maskery, C. Tuck, I. Ashcroft, N.M. Everitt, The microstructure and mechanical properties of selectively laser melted AlSi10Mg: The effect of a conventional T6-like heat treatment. Mater. Sci. Eng. A 667 (2016) 139–146. [18] R. Chou, A. Ghosh, S.C. Chou, M. Paliwal, M. Brochu, Microstructure and mechanical properties of Al10SiMg fabricated by pulsed laser powder bed fusion. Mater. Sci. Eng. A 689 (2017) 53–62. [19] H. Asgari, C. Baxter, K. Hosseinkhani, M. Mohammadi, On microstructure and mechanical properties of additively manufactured AlSi10Mg_200C using recycled powder. Mater. Sci. Eng. A 707 (2017) 148–158. [20] B. Chen, S.K. Moon, X. Yao, G. Bi, J. Shen, J. Umeda, K. Kondoh, Strength and strain hardening of a selective laser melted AlSi10Mg alloy. Scr. Mater. 141 (2017) 45–49. [21] E. Zaretsky, A. Stern, N. Frage, Dynamic response of AlSi10Mg alloy fabricated by selective laser melting. Mater. Sci. Eng. A 688 (2017) 364–370. [22] A.G. Odeshi, S. Al-ameeri, M.N. Bassim, Effect of high strain rate on plastic deformation of a low alloy steel subjected to ballistic impact. J. Mater. Process. Technol. 162-163 (2005) 385-391. [23] G.M. Owolabi, A.G. Odeshi, M.N.K. Singh, M.N. Bassim, Dynamic shear band formation in Aluminum 6061T6 and Aluminum 6061-T6/Al2O3 composites. Mater. Sci. Eng. A 457 (2007) 114–119. [24] A.A. Tiamiyu, R. Basu, A.G. Odeshi, J.A. Szpunar, Plastic deformation in relation to microstructure and texture evolution in AA 2017-T451 and AA 2624-T351 aluminum alloys under dynamic impact loading. Mater. Sci. Eng. A 636 (2015) 379-388. [25] FactSage™, Ver. 6.2. Thermfact (Montreal, Canada) and GTT-Technologies (Aachen, Germany) (1976–2010). [26] L. Thijs, K. Kempen, J.P. Kruth, J.V. Humbeeck, Fine-structured aluminium products with controllable texture by selective laser melting of pre-alloyed AlSi10Mg powder. Acta Mater. 61 (2013) 1809-1819. [27] Y. Bai, Y. Yang, Z. Xiao, M. Zhang, D. Wang, Process optimization and mechanical property evolution of AlSiMg0.75 by selective laser melting. Mater. Des. 140 (2018) 257-266. [28] N.T. Aboulkhair, N.M. Everitt, I. Ashcroft, C. Tuck, Reducing porosity in AlSi10Mg parts processed by selective laser melting. Additive Manuf. 1–4 (2014) 77–86. [29] M. Mohammadi, H. Asgari, Achieving low surface roughness AlSi10Mg 200C parts using direct metal laser sintering. Additive Manuf. 20 (2018) 23–32. [30] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, Densification mechanism and microstructural evolution in selective laser sintering of Al–12Si powders. J. Mater. Process. Technol. 211 (2011) 113-121. [31] F. Trevisan, F. Calignano, M. Lorusso, J. Pakkanen, A. Aversa, E.P. Ambrosio, M. Lombardi, P. Fino, D. Manfredi, On the Selective Laser Melting (SLM) of the AlSi10Mg Alloy: Process, Microstructure, and Mechanical Properties. Materials 10 (2017) 76. [32] N. Takata, H. Kodaira, A. Suzuki, M. Kobashi, Size dependence of microstructure of AlSi10Mg alloy fabricated by selective laser melting. Mater. Charact. (2017) In Press , https://doi.org/10.1016/j.matchar.2017.11.052. [33] N. Takata, H. Kodaira, K. Sekizawa, A. Suzuki, M. Kobashi, Change in microstructure of selectively laser

12

melted AlSi10Mg alloy with heat treatments. Mater. Sci. Eng. A 704 (2017) 218–228. [34] M.M. Kirka, P. Nandwana, Y. Lee, R.R. Dehoff, Solidification and solid-state transformation sciences in metals additive manufacturing. Scr. Mater. 135 (2017) 130–134. [35] S.M. Thompson, L. Bian, N. Shamsaei, A. Yadollahi, An overview of Direct Laser Deposition for additive manufacturing; Part I: Transport phenomena, modeling and diagnostics. Additive Manuf. 8 (2015) 36-62. [36] N.P. Gurao, A.O. Adesola, A.G. Odeshi, J.A. Szpunar, On the evolution of heterogeneous microstructure and microtexture in impacted aluminum–lithium alloy. J. Alloys Compd. 578 (2013) 183–187. [37] H. Asgari, A. Odeshi, K. Hosseinkhani, M. Mohammadi, On dynamic mechanical behavior of additively manufactured AlSi10Mg_200C. Mat. Lett. 211 (2018) 187-190. [38] J. Su, M. Sanjari, A.S.H. Kabir, I.H. Jung, S. Yue, Dynamic recrystallization mechanisms during high speed rolling of Mg–3Al–1Zn alloy sheets. Scr. Mater. 113 (2016) 198–201. [39] H. Li, E. Hsu, J. Szpunar, H. Utsunomiya, T. Sakai, Deformation mechanism and texture and microstructure evolution during high-speed rolling of AZ31B Mg sheets. J. Mater. Sci. 43 (2008) 7148–7156. [40] H.E. Hu, L. Zhen, B.Y. Zhang, L. Yang, J.Z. Chen, Microstructure characterization of 7050 aluminum alloy during dynamic recrystallization and dynamic recovery. Mater. Charact. 59 (2008) 1185-1189. [41] R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juul Jensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen, A.D. Rollett, Current issues in recrystallization: a review, Mater. Sci. Eng. A 238 (1997) 219–274. [42] E.U. Lee, H.H. Krazlein, E.E. Underwood, Dynamic Recovery in Aluminum. Mater. Sci. Eng. 7 (1971) 348356. [43] G.M. Owolabi, D.T. Bolling, A.G. Odeshi, H.A. Whitworth, N. Yilmaz, A. Zeytinci, The Effects of Specimen Geometry on the Plastic Deformation of AA 2219-T8 Aluminum Alloy Under Dynamic Impact Loading. J. Mater. Eng. Perform. 26 (2017) 5837–5846.

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