A High-Strain-Rate Superplasticity of the Al-Mg-Si-Zr-Sc Alloy with Ni Addition

materials

Article A High-Strain-Rate Superplasticity of the Al-Mg-Si-Zr-Sc Alloy with Ni Addition

Andrey Mochugovskiy, Anton Kotov , Majid Esmaeili Ghayoumabadi, Olga Yakovtseva and Anastasia Mikhaylovskaya *

Department of Physical Metallurgy of Non-Ferrous Metals, National University of Sciences and Technology “MISIS”, 4 Leninskiy ave. 4, 119049 Moscow, Russia; [email protected] (A.M.); [email protected] (A.K.); [email protected] (M.E.G.); [email protected] (O.Y.) * Correspondence: [email protected]; Tel.: +7-(495)-6384480

Abstract: The current study analyzed the effect of Ni content on the microstructure and superplastic deformation behavior of the Al-Mg-Si-Cu-based alloy doped with small additions of Sc and Zr. The superplasticity was observed in the studied alloys due to a bimodal particle size distribution. The

coarse particles of eutectic origin Al3Ni and Mg2Si phases with a total volume fraction of 4.0–8.0% and a mean size of 1.4–1.6 µm provided the particles with a stimulated nucleation effect. The L12– structured nanoscale dispersoids of Sc- and Zr-bearing phase inhibited recrystallization and grain growth due to a strong Zener pinning effect. The positive effect of Ni on the superplasticity was revealed and conﬁrmed by a high-temperature tensile test in a wide strain rate and temperature limits. In the alloy with 4 wt.% Ni, the elongation-to-failure of 350–520% was observed at 460 ◦C, in a strain rate range of 2 × 10−3–5 × 10−2 s−1. Keywords: superplasticity; aluminum alloys; transition metals; dispersoids; particle-stimulated nucleation

Citation: Mochugovskiy, A.; Kotov, A.; Esmaeili Ghayoumabadi, M.; Yakovtseva, O.; Mikhaylovskaya, A. A High-Strain-Rate Superplasticity of 1. Introduction the Al-Mg-Si-Zr-Sc Alloy with Ni AA6XXX-type (Al-Mg-Si-based) alloys are widely applied for engineering applica- Addition. Materials 2021, 14, 2028. tions [1,2], owing to their low density, high corrosion resistance, and increased mechanical https://doi.org/10.3390/ma14082028 properties at room temperature. The Al-Mg-Si-based alloys belong to the heat treatable group due to a significant strengthening effect, provided by β0, β”-phases (metastable Academic Editor: Hideki Hosoda modifications of the Mg2Si-phase), usually form during the ageing process [3–6]. Nu- merous commercial 6XXX-type alloys contain Cu that improves strength owing to the Received: 3 March 2021 formation of the metastable Q-(Al,Mg,Si,Cu)-type phase precipitates [3,7–9]. Due to the Accepted: 14 April 2021 Published: 17 April 2021 promising mechanical characteristics at room temperature and a low critical cooling rate for solid solution treatment, the Al-Mg-Si-based alloys are attractive for production of the complex-shaped parts by superplastic forming (SPF) technology. The superplasticity of Publisher’s Note: MDPI stays neutral with regard to jurisdictional claims in Al-Mg-Si based alloys is poorly studied. The main problems of Al-Mg-Si based alloys published maps and institutional affil- are the difficulty of the grain refinement and the significant grain growth that caused the iations. low strain rate sensitivity and weak superplasticity [10]. A comparatively low alloyed solid solution [11,12] is prone to the grain growth during heating up to the superplastic deformation temperature and during superplastic deformation that decrease an elongation- to-failure [10], and led to the “orange peel” effect on the post-forming surface [13]. Some progresses in the field of superplasticity for the Al-Mg-Si based alloys have been made Copyright: © 2021 by the authors. by Troeger and Stark [14,15]. The authors applied Rockwell-type technology, including Licensee MDPI, Basel, Switzerland. This article is an open access article heterogenization annealing [16] for the AA6061 alloy, to involve the particle stimulated distributed under the terms and nucleation (PSN) effect. The thermomechanical treatment provides fine-grained structure conditions of the Creative Commons with a grain size of about 9 µm and superplasticity with the elongation-to-failure of 375% at −4 −1 Attribution (CC BY) license (https:// a strain rate of 5 × 10 s . Moreover, a low strain rate is a cause of low SPF productivity. creativecommons.org/licenses/by/ A significant progress in the grain refinement of Al-Mg-Si bases alloys was attained due 4.0/). to severe plastic deformation (SPD) techniques. A mean grain size of 0.24–4.5 µm, and

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superplasticity in a wide strain rate range of 1 × 10−4–5 × 10−2 s−1 were achieved for the 6061 and 6063 alloys processed by equal-channel angular pressing and accumulative roll bonding, while the elongation-to-failure was 200−270% [17–19]. Nevertheless, low elongation, which is the result of a significant grain growth, limits complex shaped parts production by SPF method. In addition, most of the SPD techniques cannot be conveniently applied for mass production, due to financial and technical aspects. The modification of the alloys’ composition is an effective way to provide a thermally stable fine-grained structure and a good superplasticity for various alloys [20]. In this case, the grain-refinement approach is based on the bimodal particle size distribution that can be arranged after a simple thermomechanical treatment [20]. Coarse particles provide grain refinement during recrystallization via particle-stimulated nucleation (PSN) effect [16,21,22]. Such coarse particles may precipitate during heterogenization treatment from a solid solution supersaturated by Mg, Zn, Si, Cu [14,16,23], and precipitate during solidification due to alloying with eutectic forming elements, such as Ni, Fe, Mn, Ce [24–27]. Nickel is a promising alloying element for increasing the superplasticity of Al-based alloys [24,26–30]. The homogeneous distribution of the coarse Al3Ni particles with a near-spherical morphology is formed in the alloys’ sheets during a simple thermomechanical treatment [26,27,29]. Fine dispersoids are formed due to alloying with rare-earth (RE) and transition metals (TM), such as Zr, Sc, Mn, Cr, etc. Nanoscale dispersoids are precipitated during thermomechanical treatment from a supersaturated by TM/RE aluminum based solid solution [31–34]. Dispersoids cause Zener pinning effect and stabilize grain size [35–38]. The joint addition of Sc and Zr is the most effective dispersoids-forming combination. Sc and Zr provide the formation of a high density core-shell L12-strucured Al3(Sc,Zr) dispersoids [39–41]. These dispersoids result in strong Zener pining effect and inhibit grain growth, and, in addition, increase tensile properties at room temperature [42–48]. In novel alloys, the additions of Sc, as an expensive element, are reduced to 0.1 wt.%, while that of Zr is retained in a higher amount, 0.2 wt.% [32,49,50]. The bimodal distribution of the particles in the alloy’s structure provides a high strain rate superplasticity with strain rate sensitivity coefficient m above 0.3 for the AA5XXX- [24,30,51], AA7XXX- [29,30,42,52] and AA2XXX-types [27,34] alloys containing eutectic- and dispersoid-forming Fe [24], Ni [24,27,29,52], Ce [25,27,53], Y [34], Er [51], Mn [54,55], Cr [55], Zr [30,56,57], Sc [58–61]. The influence of the eutectic-forming and dispersoid-forming elements on the grain structure and superplastic deformation behavior of Al-Mg-Si type alloys is poorly studied. The current work focuses on the analysis of the microstructural evolution and superplastic behavior of a novel Al-Mg-Si-based alloy with different content of eutectic forming Ni and small additives of dispersoid forming Sc and Zr.

2. Materials and Methods A novel Al-Mg-Si-Cu-Ni-Zr-Sc alloy was studied. The concentrations of the alloying elements are shown in Table1.

Table 1. Chemical composition of the studied alloys (wt.%).

Alloy Mg Si Cu Ni Sc Zr Fe Al Reference 1.2 0.7 1.0 0 0.1 0.2 <0.01 balance 1 1.2 0.7 1.0 0.5 0.1 0.2 <0.01 balance 2 1.2 0.7 1.0 2.0 0.1 0.2 <0.01 balance 3 1.2 0.7 1.0 4.0 0.1 0.2 <0.01 balance

The alloys were prepared in a 20 kW induction furnace (Interselt, Saint-Petersburg, Russia) using graphite-ﬁreclay crucibles (Lugaabrasiv, Luga, Russia). For the preparation of the alloys, the following pure metals and “master alloys” (UC Rusal, Moscow, Russia) were used: Al (99.99 wt.%), Mg (99.96 wt.%), Al-5 wt.% Zr, Al-2 wt.% Sc, Al-12 wt.% Si, Al-53.5 wt.% Cu, and Al-20 wt.% Ni. The temperature was controlled by a chromel-alumel thermocouple. The temperature of the melt before casting was 830 ± 10 ◦C. The casting was Materials 2021, 14, 2028 3 of 16

performed using water-cooled copper mold with a size of 100 × 40 × 20 mm3, providing the casting cooling rate of approximately 15 K/s. The heat treatment of as-cast material plays the most important role for the dispersoid strengthening and grain boundary pinning effect. The high density of fine L12-strucutred Sc- and Zr-bearing dispersoids is observed after low temperature annealing [43,62–65] and two-step annealing regimes [43,56,63,65–67]. The homogenization annealing for the studied alloys was performed in two steps, with the first low-temperature step at 350 ◦C for 8 h and the second high-temperature step at 480 ◦C for 3 h to provide precipitation of the L12 dispersoids and fragmentation and spheroidization of the Al3Ni particles [29,52]. The as-homogenized alloy was subjected to a hot rolling (Rolling mill V-3P, GMT, Saint- Petersburg, Russia), with 85% total reduction and subsequent cold rolling (Rolling mill V-3P, GMT, Saint-Petersburg, Russia),), with 65% total reduction to prepare sheets with a thickness of 1 mm. The hot rolling was performed at 400 ± 20 ◦C. The hot and cold rolling were performed in 6 and 10 passes, respectively. The phase composition of the alloys in the as-cast state was controlled by X-ray diffraction (XRD) analysis, using Bruker D8 Advance diffractometer with Cu-Kα radiation (Bruker Corporation, Billerica, MA, USA). The microstructural evolution was studied by electron scanning microscopy (SEM) and optical microscopy (OM). A Tescan-VEGA3 LMH electron microscope with a tung- sten filament cathode (Tescan Brno s.r.o., Kohoutovice, Czech Republic) and an Oxford Instruments Advanced AZtecEnergy energy dispersive X-ray microanalysis system (EDS) (Oxford Instruments plc, Abingdon, UK) was used in the current study. Backscattered electron imaging (BSE) was performed with a 15 mm work distance and acceleration voltage of 20 kV. The grain structure of samples before the superplastic deformation was analyzed under a Carl Zeiss Axiovert 200M optical microscope (Carl Zeiss, Oberkochen, Germany) in a polarized light. The samples for OM and SEM analysis were prepared by mechanical grinding with subsequent polishing using the colloidal silica suspension Struers OP-U (Struers APS, Ballerup, Denmark). The samples for OM were subjected to anodic oxidation in a 10% water solution of the H3BO4 saturated in the HF. The transmission electron microscopy (TEM) analysis was carried out using a JEOL JEM 2100 microscope (JEOL, Tokyo, Japan). The disc-shaped samples of 3 mm diameter and 0.18–0.25 mm thickness were used for TEM analysis. The samples were preliminar- ily subjected to a twin-jet electropolishing in a Struers TenuPol-6 system (Struers APS, ◦ Ballerup, Denmark) at −25 ± 5 C using 30% HNO3 methanol solution (Struers APS, Ballerup, Denmark). The tensile test at elevated temperatures was performed using a Walter-Bai LFM 100 universal testing machine (Walter + Bai AG, Löhningen, Switzerland) in a temperature range of 440–500 ◦C and a constant strain rate range of 2 × 10−3–5 × 10−2 s−1. Three samples with a gauge part size of 14 × 6 × 1 mm3 were used for each testing condition. The largest dimension of the gauge part was parallel to the rolling direction. The strain rate was maintained constant due to controlled increasing travers speed during the test. The strain rate sensitivity m coefficient values were calculated as a first derivative of the logarithmically transformed Backofen equation for a strain of 50%.

. m = ∂ ln σ/∂ ln ε (1)

. where σ is ﬂow stress and ε is strain rate.

3. Results 3.1. Phase Analysis Figure1 shows the polythermal section of the Al-Mg-Si-Cu-Ni equilibrium phase diagram for the Ni concentration of below 5 wt.% calculated in a ThermoCalc (ThermoCalc- database TTAl5, Thermo-Calc Software AB, Stockholm, Sweden). The studied alloys are marked in the diagram by the blue vertical lines. The following phases were sequentially Materials 2021, 14, x FOR PEER REVIEW 4 of 16

3. Results 3.1. Phase Analysis Figure 1 shows the polythermal section of the Al-Mg-Si-Cu-Ni equilibrium phase di-

Materials 2021, 14, 2028 agram for the Ni concentration of below 5 wt.% calculated in a ThermoCalc4 of 16 (ThermoCalc- database TTAl5, Thermo-Calc Software AB, Stockholm, Sweden). The studied alloys are marked in the diagram by the blue vertical lines. The following phases were sequentially solidiﬁedsolidified for thefor alloysthe alloys studied studied in the equilibrium in the equilibrium state: the aluminum-based state: the aluminum solid solution-based solid solu- (Al),tion the (Al), Al3 Nithe and Al Mg3Ni2 Siand phases. Mg2Si phases.

FigureFigure 1. 1.The The polythermal polythermal section section of the of equilibrium the equilibrium Al-1.2%Mg-0.7%Si-1.0%Cu-var.%Ni Al-1.2%Mg-0.7%Si-1.0%Cu phase-var.%Ni phase diagram;diagram; blue blue vertical vertical lines showlines theshow studied the alloysstudied compositions. alloys compositions. 3.2. XRD Analysis 3.2.The XRD XRD Analysis analysis was performed to determine the phase composition of Alloys 1 and 3 (FigureThe2). XRD The diffractions analysis peakswas performed corresponding to todetermine the (Al), Al the3Ni, phase Al2Cu, composition and Mg2Si of Alloys 1 phases were detected in as-cast state. The (Al), Al3Ni, and Mg2Si phases were found in and 3 (Figure 2). The diffractions peaks corresponding to the (Al), Al3Ni, Al2Cu, and Mg2Si as-annealed state. phases were detected in as-cast state. The (Al), Al3Ni, and Mg2Si phases were found in as- annealed state.

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FigureFigure 2. 2.XRD XRD spectrumsspectrums for for the the alloys alloys 1 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni) (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni) and and 3 3 (Al-1.2Mg-0.7Si-1.0Cu- (Al-1.2Mg-0.7Si- 4Ni)1.0Cu in-4Ni) as-cast in andas-cast annealed and annealed states. states.

3.3.3.3. SEMSEM AnalysisAnalysis FigureFigure3 3shows shows the the BSE BSE images images for for the the studied studied alloys alloys in the in as-castthe as- state.cast state. The dominantThe domi- microstructuralnant microstructural feature feature was the was dendrites the dendrites of the Al-basedof the Al- solidbased solution. solid solution. The two The types two oftypes crystallization-origin of crystallization-origin phases phases surrounded surrounded the dendrite the dendrite cells ofcells Al of solid Al solid solution solution in the in studiedthe studied alloys; alloys; predominantly, predominantly, the Ni-richthe Ni-rich bright bright phase phase and and the darkthe dark Mg- Mg and- and Si-bearing Si-bear- phase (Mg Si). The volume fraction of Ni-bearing phase expectedly increased with in- ing phase 2(Mg2Si). The volume fraction of Ni-bearing phase expectedly increased with creasingincreasing Ni Ni content. content. The The colonies colonies of of fine fine lamellar lamellar eutectic eutectic were were observed observed in alloysin alloys 2 and 2 and 3 (Figure3 (Figure3b,c) 3b,c),, whereas whereas the the bright bright phase phase formed formed the planesthe planes on the on periphery the periphery of dendrite of dendrite cells incells the in alloy the 1alloy with 1 low with Ni low content Ni content (Figure 3(Figurea). The volume3a). The fraction volume of fraction the Ni-rich of the particles Ni-rich was 2.8 ± 0.4%, 4.8 ± 0.3%, and 8.2 ± 0.4% for the alloys 1, 2, and 3, respectively. The particles was 2.8 ± 0.4%, 4.8 ± 0.3%, and 8.2 ± 0.4% for the alloys 1, 2, and 3, respectively. Mg2Si phase appeared predominantly on the periphery of dendrite cells and its volume The Mg2Si phase appeared predominantly on the periphery of dendrite cells and its vol- fraction of 1.2 ± 0.3% was similar for the studied alloys. An incomplete superposition ume fraction of 1.2 ± 0.3% was similar for the studied alloys. An incomplete superposition of Mg-rich and Si-rich zones in the EDS-maps suggested the presence of a small fraction of Mg-rich and Si-rich zones in the EDS-maps suggested the presence of a small fraction of (Si)-phase in as-cast state. The EDS analysis demonstrated the Cu segregations on the of (Si)-phase in as-cast state. The EDS analysis demonstrated the Cu segregations on the periphery of the dendrite cells. The Sc and Zr-bearing phases of solidification origin were periphery of the dendrite cells. The Sc and Zr-bearing phases of solidification origin were unobserved and the EDS maps for Sc and Zr exhibited uniform elemental distribution for unobserved and the EDS maps for Sc and Zr exhibited uniform elemental distribution for the alloys. the alloys. SEM micrographs for as-homogenized samples are presented for the alloys 1 and 2 in Figure4 and for the alloy 3 in Figure5. The first annealing step at 350 ◦C, 8 h, did not initiate the significant changes in the as-cast microstructure for the studied alloys (Figure5a–f ). The Al3Ni phase retained lamellar morphology, and a partial fragmentation of Mg- and Si-bearing phase was observed (Figure5f). The chemical composition of the phases was studied in detail after the second step of homogenization at 480 ◦C. According to the SEM-EDS data, the bright Ni-bearing phase contained Cu (Figure4c).

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Figure 3. BSE images corresponding to SEM-EDS maps for the (a) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), (b) alloy 2 (Al- 1.2Mg-0.7Si-1.0Cu-2Ni), and (c) the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) in as-cast state.

SEM micrographs for as-homogenized samples are presented for the alloys 1 and 2 in Figure 4 and for the alloy 3 in Figure 5. The first annealing step at 350 °C, 8 h, did not initiate the significant changes in the as-cast microstructure for the studied alloys (Figure

5a–f). The Al3Ni phase retained lamellar morphology, and a partial fragmentation of Mg- Figure 3. BSE images corresponding to SEM-EDS maps for the (a) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), (b) alloy 2 (Al- Figure 3. BSE images correspondingand Si-bearing to SEM-EDS phase maps was forobserved the (a) alloy(Figure 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), 5f). The chemical composition (b) alloy of the 2 (Al- phases 1.2Mg-0.7Si-1.0Cu-2Ni), and (c) the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) in as-cast state. 1.2Mg-0.7Si-1.0Cu-2Ni), andwas (c) the studied alloy 3 in (Al-1.2Mg-0.7Si-1.0Cu-4Ni) detail after the second step in as-cast of homogenization state. at 480 °C. According to the SEM-EDS data, the bright Ni-bearing phase contained Cu (Figure 4c). SEM micrographs for as-homogenized samples are presented for the alloys 1 and 2 in Figure 4 and for the alloy 3 in Figure 5. The first annealing step at 350 °C, 8 h, did not initiate the significant changes in the as-cast microstructure for the studied alloys (Figure 5a–f). The Al3Ni phase retained lamellar morphology, and a partial fragmentation of Mg- and Si-bearing phase was observed (Figure 5f). The chemical composition of the phases was studied in detail after the second step of homogenization at 480 °C. According to the SEM-EDS data, the bright Ni-bearing phase contained Cu (Figure 4c).

FigureFigure 4.4. SEMSEM imagesimages forfor thethe ((aa)) alloyalloy 11 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni)(Al-1.2Mg-0.7Si-1.0Cu-0.5Ni) andand ((bb)) alloyalloy 22 (Al-1.2Mg-0.7Si-1.0Cu-2Ni)(Al-1.2Mg-0.7Si-1.0Cu-2Ni) afterafter thethe two-steptwo-step homogenizationhomogenization andand correspondedcorresponded to to EDS EDS elemental elemental distribution distribution maps maps captured captured from from the the area area 1 in1 in (a )(a and) and area area 2 2 in (b); (c) the EDS spectrums corresponded to the area marked-up in insert in (a,b). in (b); (c) the EDS spectrums corresponded to the area marked-up in insert in (a,b).

Figure 4. SEM images for the (a) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni) and (b) alloy 2 (Al-1.2Mg-0.7Si-1.0Cu-2Ni) after the two-step homogenization and corresponded to EDS elemental distribution maps captured from the area 1 in (a) and area 2 in (b); (c) the EDS spectrums corresponded to the area marked-up in insert in (a,b).

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Figure 5. SEM images for the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) after (a–f) one step homogenization at 350 °C, 8 h and (g–l) two-step Figure 5. SEM images for the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) after (a–f) one step homogenization at 350 ◦C, 8 h and (g–l) with the first step at 350 °C, 8 h and second step at 480 °C, 3 h; (b–e) and (h–k) are the EDS maps captured from the are marked up two-step with the first step at 350 ◦C, 8 h and second step at 480 ◦C, 3 h; (b–e) and (h–k) are the EDS maps captured from with red square in (a) and green square in (g), respectively; (f) and (l) are the magnified areas marked up with blue square in (a) and yellowthe square are marked in (g), uprespectively. with red square in (a) and green square in (g), respectively; (f) and (l) are the magnified areas marked up with blue square in (a) and yellow square in (g), respectively. The EDS analysis revealed that the bright phase contained 23.7 wt.%Ni, 2.7 wt.% Cu The EDS analysis revealed that the bright phase contained 23.7 wt.%Ni, 2.7 wt.% Cu for the alloy 1 (Figure 4 spectrum 1), and 7.5 wt.% Ni, 1.9 wt.% Cu for the alloy 2 (Figure for the alloy 1 (Figure4 spectrum 1), and 7.5 wt.% Ni, 1.9 wt.% Cu for the alloy 2 (Figure4 4 spectrum 4). The dark phase contained 21.6 wt.% Mg and 19.3 wt.% Si (Figure 4 spec- spectrum 4). The dark phase contained 21.6 wt.% Mg and 19.3 wt.% Si (Figure4 spectrum trum 3), whereas the total concentration of the other elements was less than 1 wt.%. The 3), whereas the total concentration of the other elements was less than 1 wt.%. The analysis analysis of Al solid solution revealed 0.7% Mg, 0.9% Si, and 0.9% Cu (Figure 4 spectrum of Al solid solution revealed 0.7% Mg, 0.9% Si, and 0.9% Cu (Figure4 spectrum 2). For the 2).alloy For 3, the the alloy Cu-distribution 3, the Cu-distribution became more became homogeneous more homogeneous after high-temperature after high-temperature annealing. annealing.However, aHowever, slight increase a slight of increas the Cue signal of thewas Cu signal noticed was near noticed the Ni-rich near the zones. Ni-rich zones. The second step of homogenization annealing resulted in the fragmentation and spheroidization of the Mg2Si phase for the studied alloys (Figure 4 and Figure 5g,l). The spheroidization of the Mg2Si phase for the studied alloys (Figures4 and5g,l). The dominant dominantfraction of fraction the Ni-rich of the phase Ni-rich exhibited phase non-fragmentized exhibited non-fragmentized lamellae for lamellae the alloy for 1 with the alloy 0.5% 1Ni with after 0.5% the Ni second after stepthe second of homogenization step of homogenization (Figure4a). (Figure The partial 4a). fragmentationThe partial fragmen- of Ni- tationbearing of particles Ni-bearing was particles observed was for observed the alloy 2 for with the 2 a wt.%lloy 2 Ni with (Figure 2 wt.%4b). Ni The (Figure fragmentation 4b). The fragmentationof the Ni-bearing of the phase Ni- wasbearing significant phase w foras thesignificant alloy 3 withfor the 4% alloy Ni. 3 with 4% Ni. Figure 66 showsshows thethe SEMSEM micrographsmicrographs ofof thethe alloysalloys studiedstudied afterafter a simple thermomechanical treatment and and post-deformation post-deformation annealing annealing at 480 at 480◦C for°C 20for min. 20 min.A near-homogeneous A near-homoge- neousdistribution distribution of the of solidification the solidification origin origin particles particles was observed was observed in the in alloys the alloys studied studied after aftersheet sheet processing. processing.

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Figure 6. SEM images of cold rolled sheets after final 20 min annealing at 480 °C for the (a) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu- Figure 6. 6. SEMSEM images images of ofcold cold rolled rolled sheets sheets after after final ﬁnal 20 min 20 minannealing annealing at 480 at °C 480 for◦ Cthe for (a) the alloy (a) 1 alloy (Al-1.2Mg 1 (Al-1.2Mg-0.7Si--0.7Si-1.0Cu- 0.5Ni), (b) alloy 2 (Al-1.2Mg-0.7Si-1.0Cu-2Ni), and (c,d) alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) with corresponded SEM-EDS 0.5Ni)1.0Cu-0.5Ni),, (b) alloy (b 2) alloy(Al-1.2Mg 2 (Al-1.2Mg-0.7Si-1.0Cu-2Ni),-0.7Si-1.0Cu-2Ni), and (c,d) and alloy (c ,3d (Al) alloy-1.2Mg 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni)-0.7Si-1.0Cu-4Ni) with corresponded with corresponded SEM-EDS maps. mapsSEM-EDS. maps. The particles size distribution histograms (Figure 7) exhibited gamma distribution TheThe particles particles size size distribution histograms (Figure (Figure 77)) exhibitedexhibited gammagamma distributiondistribution forfor bothboth both typestypes types ofof of particles.particles. particles. TheThe The meanmean mean sizessizes sizes ofof of brightbright bright NiNi Ni-bearing--bearingbearing particlesparticles particles werewere were 1.1.77 1.7 ±± 00.1,±.1, 0.1,1.1.66 ± 0.1, and 1.4 ± 0.2, and the mean sizes of Mg2Si particles were 1.4 ± 0.1, 1.4 ± 0.1, and 1.6 ± ±1.6 0.1,± and0.1, and1.4 ± 1.4 0.2±, and0.2, the and mean the mean sizes sizes of Mg of2Si Mg particles2Si particles were were 1.4 ± 1.4 0.1,± 1.0.1,4 ± 1.40.1,± and0.1, 1. and6 ± 01.60.1.1,, ± forfor0.1, thethe for alloysalloys the alloys 1,1, 2,2, andand 1, 2, 3,3, and respectively.respectively. 3, respectively.

Figure 7. The particle size distribution histograms of the (a–c) Al3Ni and (d–f) Mg2Si phases for the thermomechanical treated (a,d) Figure 7. The particle size distribution histograms of the (a–c) Al3Ni and (d–f) Mg2Si phases for the thermomechanical treated (a,d) alloy Figure1 (Al-1.2Mg 7. The-0.7Si particle-1.0Cu size-0.5Ni) distribution, (b,e) alloy histograms 2 (Al-1.2Mg of the-0.7Si (a–-1.0Cuc) Al3-Ni2Ni) and, and (d– (fc),f Mg) alloy2Si phases3 (Al-1.2Mg for the-0.7Si thermomechanical-1.0Cu-4Ni). alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), (b,e) alloy 2 (Al-1.2Mg-0.7Si-1.0Cu-2Ni), and (c,f) alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni). treated (a,d) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), (b,e) alloy 2 (Al-1.2Mg-0.7Si-1.0Cu-2Ni), and (c,f) alloy 3 (Al-1.2Mg-0.7Si- 1.0Cu-4Ni). 3.4.3.4. DispersoidsDispersoids ParametersParameters

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3.4. Dispersoids Parameters The TEM analysis of the alloy 3 after a two-step annealing revealed a high-density of The TEM analysis of the alloy 3 after a two-step annealing revealed a high-density of uniformlyThe TEM distributed analysis nanoscale of the alloy dispersoids 3 after a two with-step a meanannealing size ofrevealed 10 ± 1 nma high (Figure-density8a,b). of uniformly distributed nanoscale dispersoids with a mean size of 10 ± 1 nm (Figure 8a,b). uniformlyThe corresponding distributed selected nanoscale area dispersoids electron diffraction with a mean (SAED) size in of [111] 10 ±l zone1 nm axis(Fig exhibitedure 8a,b). The corresponding selected area electron diffraction (SAED) in [111]l zone axis exhibited Thetypical corresponding for the L12-phase selected ordered area electron superlattice diffraction reﬂections (SAED (Figure) in8 [111]c). Thel zone [111] axis zone exhibited axis of typical for the L12-phase ordered superlattice reflections (Figure 8c). The [111] zone axis typicalL12 phase for isthe parallel L12-phase to [111] ordered zone superlattice axis of Al. reflections (Figure 8c). The [111] zone axis of L12 phase is parallel to [111] zone axis of Al. of L12 phase is parallel to [111] zone axis of Al.

FigureFigure 8. 8. TEMTEM images images for the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-3Ni)(Al-1.2Mg-0.7Si-1.0Cu-3Ni) afterafter two-steptwo-stepannealing; annealing; ( a(a)) bright bright ﬁeld; field; ( b(b)) dark dark ﬁeld; field; (c ) (Figurec) corresponded 8. TEM images SAED. for the alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-3Ni) after two-step annealing; (a) bright field; (b) dark field; corresponded(c) corresponded SAED. SAED. 3.5.3.5. Grain Grain Structure Structure for for the the Thermomechanically Thermomechanically Treated Treated Samples 3.5. Grain Structure for the Thermomechanically Treated Samples HighHigh density of of L1 L12 2precipitates precipitates provided provided a partially a partially non-recrystallized non-recrystallized band-like band grain-like High density of L12 precipitates provided a partially non-recrystallized band-like grainstructure structure after postafter deformation post deformation annealing annealing in a temperature in a temperature range ofrange 440–500 of 44◦0C– for50020 °C min for grain structure after post deformation annealing in a temperature range of 440–500 °C for 20(Figure min (Figure9). The 9). fraction The fraction of the of recrystallized the recrystallized grains grains decreased decreased with with increasing increasing volume vol- 20 min (Figure 9). The fraction of the recrystallized grains decreased with increasing vol- umefraction fraction of Al of3Ni Al particles.3Ni particles. The The mean mean thickness thickness of non-recrystallized of non-recrystallized grains grains decreased decreased with ume fraction of Al3Ni particles. The mean thickness of non-recrystallized grains decreased withincreasing increasing Ni content Ni content and wereand were 8.4 ± 8.1.3,4 ± 6.81.3,± 6.0.5,8 ± 5.50.5,± 5.0.55 ± and0.5 and 4.0 ±4.00.3 ± 0µ.3m µm for Ni-freefor Ni- with increasing Ni content and were 8.4 ± 1.3, 6.8 ± 0.5, 5.5 ± 0.5 and 4.0 ± 0.3 µm for Ni- freereference reference alloy, alloy, alloy alloy 1, alloy 1, alloy 2, and 2, and alloy alloy 3, respectively 3, respectively (Figure (Figure9a–d). 9a–d). free reference alloy, alloy 1, alloy 2, and alloy 3, respectively (Figure 9a–d).

Figure 9. The grain structure for the (a,e) reference Ni-free alloy, (b,f) alloy 1 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni), (c,g) alloy 2 (AlFigure-1.2Mg 9. The-0.7Si grain-1.0Cu structure-2Ni), and for ( d the,h) (alloy,e) referencereference 3 (Al-1.2Mg Ni-freeNi--free0.7Si alloy,-1.0Cu ((bb-,4Ni),ff)) alloyalloy after 11 (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni),the(Al -thermomechanical1.2Mg-0.7Si-1.0Cu -treatment0.5Ni), (c, gand) alloyalloy final 22 (Al-1.2Mg-0.7Si-1.0Cu-2Ni), and (d,h) alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) after the thermomechanical treatment and final annealing(Al-1.2Mg-0.7Si-1.0Cu-2Ni), at (a–d) 500 °C for and20 min (d, hand) alloy (e–h 3) (Al-1.2Mg-0.7Si-1.0Cu-4Ni) 300% of strain at 460 °C and after constant the thermomechanical strain rate of 5 × 10 treatment−3 s−1. and ﬁnal annealing at (a–d) 500 °C for 20 min and (e–h) 300% of strain at 460 °C and constant strain rate of 5 × 10−3 s−1. annealing at (a–d) 500 ◦C for 20 min and (e–h) 300% of strain at 460 ◦C and constant strain rate of 5 × 10−3 s−1. 3.6. The Tensile Test at Elevated Temperatures 3.6. The Tensile Test at Elevated Temperatures 3.6. TheTheTensile superplastic Test at Elevateddeformation Temperatures behavior was analyzed in a temperature range of 440– The superplastic deformation behavior was analyzed in a temperature range of 440– 500 °CThe and superplastic a constant strain deformation rate range behavior of 2 × 10 was−3–5analyzed × 10−2 s−1 (Figure in a temperature 10). The flow range stress of 500 °C and a constant strain rate range of 2 × 10−3–5 × 10−2 s−1 (Figure 10). The flow stress decreased440–500 ◦C withand increasing a constant temperature strain rate range and ofdecreasing 2 × 10−3–5 strain× 10 rate.−2 s −In1 (Figurethe studied 10). Thetemper- ﬂow decreased with increasing temperature and decreasing strain rate. In the studied temper- aturestress limits decreased, the maximum with increasing stress was temperature observed andat the decreasing maximum strain studied rate. strain In the rate studied of 5 × ature limits, the maximum stress was observed at the maximum studied strain rate of 5 × 10−2 s−1. In a temperature range of 440–500 °C , the maximum stress values varied in a range 10−2 s−1. In a temperature range of 440–500 °C , the maximum stress values varied in a range

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temperature limits, the maximum stress was observed at the maximum studied strain rate of 5 × 10−2 s−1. In a temperature range of 440–500 ◦C, the maximum stress values varied inof a27 range–48 MPa, of 27–48 26–45 MPa, MPa, 26–45 and 25 MPa,–40 MPa and 25–40for alloys MPa with for alloys0.5, 2, withand 4 0.5, wt.% 2, andNi, respec- 4 wt.% Ni, −3 −1 respectively.tively. The minimum The minimum stress value stress was value achieved was achieved at 2 × 10− at3 s 2−1.× In10 a temperatures . In a temperature range of ◦ range440–500 of 440–500°C , the minimumC, the minimum stress values stress varied values in varied a range in of a range 15–18 of MPa,15–18 12 MPa–16 MPa,, 12–16 and MPa , and10–15 10–15 MPa MPafor 0.5, for 2, 0.5, and 2, 4 and wt.% 4 wt.%Ni, respectively. Ni, respectively. An increas An increasee in Ni content in Ni content led to a ledde- to a decreasecrease in inthe the flow ﬂow stress stress for forthe thestudied studied temperatures. temperatures.

FigureFigure 10. 10.The The ﬂow flow stress-strain stress-strain dependencesdependences for the the ( (aa––dd) )alloy alloy 1 1(Al (Al-1.2Mg-0.7Si-1.0Cu-0.5Ni),-1.2Mg-0.7Si-1.0Cu-0.5Ni), (e–h (e)– alloyh) alloy 2 (Al 2- (Al-1.2Mg-1.2Mg- 0.7Si-1.0Cu-2Ni), and (i–l) alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) in a strain rate range of 2 × 10−3 s−1–5 × 10−2 s−1 and tempera- 0.7Si-1.0Cu-2Ni), and (i–l) alloy 3 (Al-1.2Mg-0.7Si-1.0Cu-4Ni) in a strain rate range of 2 × 10−3 s−1–5 × 10−2 s−1 and tures of (a,e,i) 440 °C, (b,f,j) 460 °C, (c,g,k) 480 °C, and (d,h,l) 500 °C. temperatures of (a,e,i) 440 ◦C, (b,f,j) 460 ◦C, (c,g,k) 480 ◦C, and (d,h,l) 500 ◦C. At the studied deformation conditions, the elongation-to-failure was varied from 130 At the studied deformation conditions, the elongation-to-failure was varied from 130 to 520%. The alloy 1 with 0.5% Ni exhibited superplasticity with elongation-to-failure of to 520%. The alloy 1 with 0.5% Ni exhibited superplasticity with elongation-to-failure of 300–370% at 440–480 °C and strain rates of (2–8) × 10−3 s−1 (Figure 10a–c). At higher strain 300–370% at 440–480 ◦C and strain rates of (2–8) × 10−3 s−1 (Figure 10a–c). At higher rates, the elongation did not exceed 200%. The alloys with 2 and 4% Ni exhibited a similar strain rates, the elongation did not exceed 200%. The alloys with 2 and 4% Ni exhibited deformation behavior. For the alloy 2 with 2 % Ni, the maximum elongation-to-failure of a similar deformation behavior. For the alloy 2 with 2% Ni, the maximum elongation-to-−3 −1 370–400% was achieved at temperatures of 440–460 °C in a strain ◦rate range of 2 × 10 s – failure1 × 10−2 of s−1 370–400% (Figure 10e,f). was achievedThe maximum at temperatures elongation of of 520 440–460% was achievedC in a strain at 460 rate °C range and of × −3 −1 × −2 −1 2(2–810) × 10s−3 s−1–1 for the10 alloys 3(Figure with 4 %10 Nie,f). (Figure The maximum 10j). At the elongation higher strain of 520% rates wasof (1 achieved–5) × ◦ −3 −1 at10 460−2 s−1, Cthe and alloy (2–8) exhibited× 10 elongations for the of alloy 350– 3420% with. For 4% the Ni (Figuredecreased 10 j).temperature At the higher of 440 strain −2 −1 rates°C, at of high (1–5) strain× 10 ratess of 5, the× 10 alloy−2 s−1, the exhibited elongation elongation was below of 230%, 350–420%. whereas For a thestrain decreased rate ◦ −2 −1 temperatureof 2 × 10−3 s−1 of provided 440 C, atthe high elongation strain rates of 520 of% 5 ×(Figure10 s10i)., theThe elongation temperature was increase below to 230%, −3 −1 whereas480–500 °C a strain resulted rate in of a 2decrease× 10 ins elongationprovided in the both elongation alloys 2 and of 520%3 (Figure (Figure 10g,h,k,l). 10i). The ◦ temperatureThe elongation increase of 360% to at 480–500 500 °C wasC resultedachieved infor a the decrease alloy 3 in with elongation 4% Ni at ina low both strain alloys 2 ◦ andrate 3of (Figure 2 × 10−3 10 s−1g,h,k,l). (Figure The 10l). elongation of 360% at 500 C was achieved for the alloy 3 with − − 4% NiThe at asuperplastic low strain behavior rate of 2 ×of 10the 3Nis-free1 (Figure reference 10l). alloy was analyzed in similar testing conditions.The superplastic The alloy behavior exhibited ofthe elongation Ni-free- referenceto-failure alloyof 250 was–300 analyzed% in a strain in similar rate range testing conditions.of 2 × 10−3–1 The× 10−2 alloy s−1 and exhibited a temperature elongation-to-failure range of 440–500 of °C 250–300% (Figure 11), in aand strain non- ratesuper- range ofplastic 2 × 10flow−3 –1accompanied× 10−2 s− 1byand necking a temperature with 120% rangeof elongation of 440–500 was observed◦C (Figure at a11 constant), and non- superplasticstrain rate of ﬂow5 × 10 accompanied−2 s−1. by necking with 120% of elongation was observed at a − − constant strain rate of 5 × 10 2 s 1.

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◦ −2 −3 −1 FigureFigureFigure 11. 11. The The 11. flow flowThe stress stress ﬂow-strain- stress-strainstrain dependences dependences dependences for for the the reference reference for the referenceNi Ni-free-free alloy alloy Ni-free at at ( a(a alloy) )460 460 °C at°C in( ina )a a 460strain strainC rate rate in arange range strain of of rate1 1 × × 10 10 range−2−−55 × × of10 10−3 s s−1 −2 −1 −3 −1 andand at 1at ×constant constant10−2− 5strain strain× 10 rates− rates3 s− of1 ofand ( b(b) )1 at 1 × constant× 10 10−2 s s−1 and strainand ( c()c rates)5 5 × × 10 10 of−3 s ( bs−1) for 1for× 440, 440,10− 4602 460s− and1 andand 500 500 (c )°C. °C. 5 × 10−3 s−1 for 440, 460 and 500 ◦C.

◦ TheThe strain strainstrain rate rate sensitivity sensitivity coefficient coefficient coefficient m m mwas waswas determined determined determined at at 460 at460 460°C °C for forC 50% for50% 50%strains strains strains (Fig- (Fig- −3 −1 (Figureureure 12a). 12a). 12 The a).The Them m-values-valuesm-values were were were above above above 0.3 0.3 0.3in in a ina strain strain a strain rate rate rate range range range of of ( of2(2– (2–5)–55) )× × 10 10×−3−310 s −1s−1 for fors the thefor alloy alloy the withalloywith 0.5%Ni. with0.5%Ni. 0.5%Ni. The The m m The-values-valuesm-values of of 0.3 0.35 of5––0 0.35–0.450.45.45 were were observed were observed observed in in the the inalloys alloys the alloys with with 2% with2% and and 2% 4 4% and% Ni Ni 4% in in −3 −2 −1 aNia strain strain in a strainrate rate range range rate rangeof of 5 5 × × 10 of 10−3 5−3–×–55 ×10 × 10 10−2−2–5 s −1s−1×. .For For10 the thes Ni Ni.- Forfree-free thealloy alloy Ni-free, ,the the coefficient coefficient alloy, the coefficientm m was was above abovem −3 −1 was0.30.3 in in above a a strain strain 0.3 rate inrate a strainrange range rateof of ( 2( range2––88) )× × 10 of10−3 (2–8)−3 s −1s−1 and and× 10 below belows 0.3 0.3and at at belowhigher higher 0.3strain strain at higher rates. rates. strainAn An m m rates.coef- coef- ◦ Anficientficientm coefficient value value higher higher value than than higher 0.35 0.35 thanresulted resulted 0.35 inresulted in increased increased in increased elongation elongation elongation to to failure failure to at failureat 460 460 °C °C at (Figure 460(FigureC (Figure12b).12b). 12b).

FigureFigureFigure 12. 12. The The 12. dependences Thedependences dependences of of ( a(a) of)strain strain (a) strain rate rate sensitivity sensitivity rate sensitivity m m-coefficient-coefficientm-coefﬁcient and and ( b( andb) )elongation elongation (b) elongation to to failure failure to failure vs. vs. Ni Ni vs.content content Ni content (x (x-axis)-axis) (x-axis) and and strain strain raterate (y and(y-axis)-axis) strain at at 460 rate460 °C. °C. (y-axis) at 460 ◦C.

TheThe dynamic dynamicdynamic recrystallization recrystallization occurred occurred occurred during during during the the the superplastic superplastic superplastic deformation deformation deformation and and pro- and pro- providedvidedvided a a fine fine a ﬁne-grained-grained-grained structure structure structure formation formation formation for for for Ni Ni Ni-bearing-bearing-bearing alloys alloys alloys containing containingcontaining coarsecoarse AlAl Al33Ni3Ni ◦ −3 −1− particles.particles. After After deformation deformation of of 300%300 300%% atat at 460460 460 °C C°C withwith with aa a constantconstant constant strainstrain strain raterate rate ofof of5 5 5× × × 10 10−3 3s s−,1 1,a ,a meanamean mean grain grain grain size size size was was was 9. 9.99 9.9± ± 0 0.6,±.6, 0.6,7. 7.22 ± 7.2± 0 0.6,.6,± and and0.6, 5. and5.99 ± ± 0 5.90.6.6 µm ±µm0.6 for forµ the mthe foralloys alloys the 1, alloys1, 2, 2, and and 1, 3, 2,3, respec- andrespec- 3, tivelyrespectivelytively (Figure (Figure (Figure 9f 9f––h).h). 9The f–h).The reference reference The reference alloy alloy that that alloy contained contained that contained a a low low fraction fraction a low fractionof of coarse coarse of particles particles coarse particlesdemonstrateddemonstrated demonstrated partially partially partiallynon non-recrystallized-recrystallized non-recrystallized grain grain structure structure grain structure,, ,even even after after even the the after failure failure the failure at at 300% 300% at and300%and a a andmean mean a meangrain grain size grain size in in size the the inrecrystallized recrystallized the recrystallized areas areas areaswas was 1 1 was66 ± ± 1 1 16µm µm± (Figure 1(Figureµm (Figure 9e). 9e). 9e).

4.4. Discussio Discussion Discussionn TheThe thermodynamic thermodynamic calculation, calculation, XRD XRD analysis analysis and andand SEM SEMSEM studies studiesstudies confirmed conﬁrmedconfirmed the thethe pres- pres-presence of Mg2Si phase, and the Al3Ni phase of crystallization origin. For an as-cast state, a enceence of of Mg MgSi2Si phase, phase, and and the the Al AlNi3Ni phase phase of of c crystallizationrystallization origin. origin. For For an an as as-cast-cast state, state, a a small amount of non-equilibrium (Si) phase also was found by SEM-EDS. The (Si) phase smallsmall amount amount of of non non-equilibrium-equilibrium (Si) (Si) phase phase also also was was found found by by SEM SEM-EDS.-EDS. The The (Si) (Si) phase phase was dissolved during homogenization annealing. The bright particles of crystallization waswas dissolved dissolved during during homogenization homogenization annealing. annealing. The The bright bright particles particles of of crystallization crystallization origin belonged to Al3Ni phase were found in as-cast state. The non-equilibrium Al2Cu originorigin belonged belonged to to Al Al3Ni3Ni phase phase were were f oundfound in in as as-cast-cast state state. .T Thehe non non-equilibrium-equilibrium Al Al2Cu2Cu

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phase observed in as-cast state was dissolved during homogenization. Meanwhile, the SEM-EDS analysis for the bright phase after homogenization revealed an increased concentration of Cu, comparing it to that of for the aluminum-based solid solution (Al). This can be a consequence of partial substitution of Ni atoms by Cu atoms in Al3Ni phase or the result of Al2Cu phase that was incompletely dissolved during the homogenization. It can also be suggested that Cu-rich Al7Cu4Ni-phase can be formed in the alloy 1 with low Ni/Cu ratio [67]. The Mg2Si phase was fragmentized and spheroidized during homogenization of the studied alloys. The microstructural studies of the as-annealed samples revealed differences in the evolution of the Al3Ni lamellae. The fragmentation was observed for the alloys 2 and 3 with increased Ni content, whereas Al3Ni phase retained its elongated shape in the alloy 1 with 0.5% Ni. It can be suggested that Cu atoms were dissolved in the Al3Ni phase and increased its thermal stability and inhibited fragmentation. The increasing of Ni content results in a higher fraction of Al3Ni eutectic phase and, therefore, the effect of Cu becomes increasingly insignificant. After the thermomechanical treatment, a near-uniform distribution of the Al3Ni and Mg2Si particles with a mean size of 1.4–1.6 µm was observed. Nickel has an extremely low solubility in aluminum at conventional solidification conditions [68], and the dissolution of Ni in the Sc and Zr bearing precipitates was not observed by atomic probe tomography [69,70], therefore, Ni cannot significantly influence the precipitation kinetics and parameters of dispersoids. It has been suggested that the precipitation behavior of the alloys with different Ni content was similar. It should be mentioned that the applied homogenization annealing provided L12-Al3(Sc,Zr) dispersoids with a size of about 10 nm that is similar to the size of precipitates formed at low temperature annealing [43]. We also suggest that Si atoms partially substitute Al in L12 phase providing a more complex dispersoids of (Al,Si)3(Sc,Zr) phase as it was demonstrated in Refs. [71,72]. The zone axis [111] of the dispersoids agglomeration was parallel to the corresponding zone axis of the Al matrix that suggested coherency/semi-coherency for the dispersoids and Al matrix. The alloys exhibited partially unrecrystallized grain structure before the start of the superplastic deformation up to a temperature of 500 ◦C. This emphasizes a strong dislocation pinning effect, due to the L12 dispersoids, which inhibited recrystallization during heating to the superplastic deformation temperature. The role of the secondary particles on the superplastic behavior is follows. The well- known superplastic deformation mechanisms are grain boundary sliding, dislocation and diffusional creep [10,73–78]. Due to the diffusion-controlled nature of these mechanisms, the superplasticity is usually observed at elevated temperatures and low strain rates. Grain refinement simplifies grain boundary sliding and its accommodation by dislocation and diffusional creep, therefore, superplastic behavior is observed at higher strain rates and lower temperatures [20,79]. The bimodal particles size distribution with Al3Ni and Mg2Si particles of about 1–2 µm in size, and L12 precipitates with size of 10 nm contributed significantly to the grain refinement and provided high strain rate superplasticity for the studied alloys. The Al3Ni phase particles accumulate dislocations near interphase boundaries and provide the pronounced PSN effect during the superplastic flow [20,42,58]. Due to PSN mechanism, Al3Ni particles increased the dynamic recrystallization rate during the deformation and provided grain refinement. The reference alloy without Ni and Al3Ni particles demonstrated weak superplasticity at higher strain rates due to suppressed dynamic recrystallization. In addition, Al3Ni particles form interphase boundaries with aluminum matrix. An increased fraction of the interphase boundaries as well as grain boundaries can accelerate the alloy diffusivity and support the acting of diffusion-controlled superplastic deformation mechanisms. As a result, an increase in Al3Ni phase fraction increased the elongation-to-failure and decreased the flow stress values. The fine L12 precipitates provided a strong Zener pinning effect, inhibiting dynamic grain growth and providing a stable flow during superplastic deformation. Materials 2021, 14, 2028 13 of 16

5. Conclusions The microstructure and the superplastic deformation behavior were studied for the Al-1.2Mg-0.7Si-1Cu-xNi-0.1Sc-0.2Zr alloys, where x = 0, 0.5, 2, and 4 wt.%. The sheets were processed by a simple thermomechanical treatment included homogenization annealing, hot and cold rolling. The main conclusions are drawn as follows: 1. The bimodal particles size distribution with the coarse crystallization-origin particles of the Mg2Si and Al3Ni phases with 1.4–1.6 µm mean size and L12-structured dispersoids of the Al3(Sc,Zr) phase with a mean size of 11 ± 1 nm were formed in the studied alloys. The volume fraction of the Mg2Si phase was 1.2% and the volume fraction of the Al3Ni phase increased from 2.8 to 8%, with increasing Ni-content from 0.5 to 4%. 2. The alloys studied exhibited superplasticity in a strain rate range of 2 × 10−3–5 × 10−2 s−1 and a temperature range of 440–500 ◦C. Due to a particle stimulated nucleation effect, an increase in Al3Ni phase fraction provided grain reﬁnement during the superplastic deformation, increased the value of elongation-to-failure and decreased ﬂow stress values. An elongation increased from 250–300% for the alloys, with 0–0.5% Ni to 400–500% for the alloy with 4% Ni, which exhibited superplasticity even at a strain rate of 5 × 10−2 s−1.

Author Contributions: Conceptualization, A.M. (Andrey Mochugovskiy) and A.M. (Anastasia Mikhaylovskaya); methodology, M.E.G. and A.K.; software, A.K.; validation, M.E.G. and O.Y.; formal analysis, A.K.; investigation, A.M. (Andrey Mochugovskiy), M.E.G. and A.M. (Anastasia Mikhaylovskaya); data curation, O.Y.; writing—original draft preparation, A.M. (Andrey Mochugovskiy) and A.M. (Anastasia Mikhaylovskaya); writing—review and editing, A.M. (Anastasia Mikhaylovskaya) and M.E.G.; visualization, O.Y.; supervision, A.M. (Anastasia Mikhaylovskaya). All authors have read and agreed to the published version of the manuscript.

Funding: This work was funded by the Russian Science Foundation in a framework of [Grant#20-79-00269]. Institutional Review Board Statement: Not applicable. Informed Consent Statement: Not applicable. Data Availability Statement: The raw and processed data required to reproduce these results are available by contacting the authors. Conﬂicts of Interest: The authors declare no conﬂict of interest.

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