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Article

The Effect of Ce on the Microstructure, Superplasticity, and Mechanical Properties of Al-Mg-Si-Cu Alloy

1
Department of Physical Metallurgy of Non-Ferrous Metals, National University of Sciences and Technology “MISIS”, Leninsky Prospekt, 4, 119049 Moscow, Russia
2
Prokhorov General Physics Institute of the Russian Academy of Sciences, 8 Vavilov Str., 119991 Moscow, Russia
*
Author to whom correspondence should be addressed.
Metals 2022, 12(3), 512; https://doi.org/10.3390/met12030512
Submission received: 24 February 2022 / Revised: 11 March 2022 / Accepted: 15 March 2022 / Published: 17 March 2022

Abstract

:
The current study focuses on the influence of Ce on the superplastic behavior, microstructure, and mechanical properties of the Al-Mg-Si-Cu-Zr-Sc alloy. The multilevel microstructural analysis including light, scanning electron, and transmission electron microscopies was carried out. The simple thermomechanical treatment including the hot and cold rolling resulted in fragmentation of the eutectic originated particles of the Ce-bearing phases. The two-step annealing of the ingots provided the precipitation of the L12-structured Al3(Sc,Zr) phase dispersoids with 10 nm mean size and a high number density. Due to the particle stimulated nucleation (PSN) effect caused by the particles of eutectic origin, and Zener pinning effect provided by nanoscale dispersoids of L12-structured phases, the studied alloy demonstrated good superplastic properties.

1. Introduction

The Al-Mg-Si based (6xxx type) alloys are widely used in aircraft and machine building due to low density, good mechanical properties, and increased corrosion resistance [1,2,3]. The Al-Mg-Si alloys are heat treatable aluminum alloys, and the precipitation of a metastable modification of Mg2Si phase at ageing significantly improves their strength [4,5,6,7]. Copper is an additional alloying element for commercial 6xxx type alloys that improves strengthening effect at ageing, owing to formation of the complex Q-AlCuSiMgCu phase [1,4,8,9]. The sheets of 6xxx-type alloys are promising for manufacturing of thin-walled complex-shaped parts by a superplastic forming (SPF) technology. Due to low critical cooling rate for this group of alloys, the supersaturated solid solution is formed by cooling from a solid solution treatment temperature or deformation temperature. However, the low-alloyed solid solution typically formed in 6xxx-type alloys [10] complicates grain refinement and simplifies dynamic grain growth that negatively affect the superplasticity [11]. The severe plastic deformation (SPD) results in the ultrafine structure and superplasticity in Al-Mg-Si based alloys [12,13,14], whereas the conventional thermomechanical treatments were not effective. Several attempts to refine grain in commercial 6061, 6013, and 6066 alloys were performed by Troeger and Stark [15,16]. The authors applied a similar approach for the Rockwell technology [17] including over-aging with a formation of coarse Mg2Si particles providing a particle stimulated nucleation (PSN) effect at recrystallization. The obtained materials exhibited the elongation of about 200% even at low strain rates of 0.0004 s−1.
The attractive approach [18] to refine grain and provide superplasticity in aluminum alloys combines the PSN effect [17,19,20] provided by the coarse particle of eutectic origin with 0.5–2 µm size and Zener pinning effect [21,22,23,24] due to nano-scaled dispersoids. This approach was previously successfully applied for 7xxx-type [25,26,27,28], 5xxx-type [26,29,30,31,32], 2xxx-type [33] aluminum alloys dopped with Ni, Fe, Er, Y. In our previous studies the effectiveness of combined PSN and Zenner pinning effects for grain refinement and superplasticity was demonstrated in 6xxx-type alloy doped with Fe, Ni, and Y [34,35]. Among rare-earth metals, Cerium (Ce) is considered as a promising alloying element in aluminum alloys. The recent study demonstrated that the Ce-enriched phases fragmentize during the heat- and thermomechanical treatments and become homogeneously distributed in the Al matrix [36,37,38]. The positive effect of Ce on the superplasticity was demonstrated in 5xxx-type alloy [31,39]. Due to the evenly distributed particles of the Ce-bearing phases the studied alloy demonstrated a superplasticity in a strain rate range of 10−1–10−2 s−1 and a temperature range of 500–540 °C. Moreover, several studies found that Ce positively affects the castability [40,41,42,43,44,45], refines of the as-cast grains and phases [46], improves wear resistance [40], creep resistance and mechanical strength of cast Al-based alloys [41,42,47,48].
The nanoscale dispersoids enhance the structure thermal stability and inhibit dynamic grain growth that favors the fine-grained structure and superplasticity [49]. Scandium (Sc) and (Zr) had proved to be the most effective dispersoid forming elements in aluminum alloys. The complex addition of Sc and Zr initiates the nucleation of core-shell Al3(Sc,Zr) L12-structured thermally stable dispersoids with 5–10 nm size [50,51,52]. These dispersoids effectively inhibit grain growth, provide superplasticity [53,54,55,56,57], and also increase tensile strength [58,59,60,61,62,63,64,65].
The present research focuses on the influence of eutectic-forming Ce and dispersoid forming Sc and Zr on the microstructure, superplasticity, and mechanical properties of the 6xxx-type aluminum alloy.

2. Materials and Methods

For the alloys’ preparation, the following pure metals and master alloys were used; (UC Rusal, Moscow, Russia) 99.99 wt.% Al, 99.95 wt.% Mg, Al-20 wt.% Ce, Al-2.5 wt.% Sc, Al-5.0 wt.% Zr, Al-59.5 wt.% Cu, Al-12.0 wt.% Si. The master alloys were added to the melt first after the aluminum melt reached the required temperature. The alloys’ compositions are presented in Table 1. The melting process was carried out in an inductive furnace (Interselt, Saint-Petersburg, Russia) using graphite-fireclay crucibles (Lugaabrasiv, Luga, Russia). The temperature of the melt before casting was 790 ± 20 °C. The chromel-alumel thermocouple was used to control the temperature. The casting was performed to a water-cooled copper mold with internal dimensions of 100 × 40 × 20 mm. The cooling rate during the casting procedure was about 15 K/s. The as cast samples were subjected to a two-step homogenization with the first step at 350 °C for 8 h and the second step at 480 °C for 3 h. The heat treatment was performed in a Nabertherm N30/65A furnace with air-forced convection (Nabertherm, Lilienthal, Germany). The thermomechanical treatment included hot and subsequent cold rolling (Rolling mill V-3P, GMT, Saint-Petersburg, Russia) with 75 and 80% thickness reduction, respectively. The hot rolling temperature was 420 ± 20 °C. The rolls diameter was 250 mm, the rolling speed was 10 m/s. The strain rate was 15 s−1.
The X-ray diffraction (XRD) analysis was carried out to analyze the phase composition of the studied materials. The powder of the studied alloys was exposed under Cu-Kα radiation using a Bruker D8 Advanced diffractometer (Bruker Corporation, Billerica, MA, USA). The peaks identification was carried out using the Powder Diffraction Data (PDFMaintEx library v. 9.0.133) database using a DIFFRAC.EVA software (v.12.0 rev.0, Bruker AXS GmbH, Karlsruhe, Germany).
The microstructural evolution and chemical composition of the studied alloys were studied using a Tescan-VEGA3 LMH scanning electron microscope (SEM) (Tescan Brno s.r.o., Kohoutovice, Czech Republic) equipped with an energy dispersive X-ray spectrometer (EDS) X-MAX80, and an electron backscatter diffraction (EBSD) detector HKL NordlysMax (Oxford Instruments plc, Abingdon, UK). The samples were grinded with SiC papers and further polished with OP-S silica colloidal suspension using a Struers LaboPoll-5 polishing machine (Struers APS, Ballerup, Denmark). The EBSD maps were exposed from an area of 150 × 150 μm2 with a step size of 0.3 μm.
The JEOL JEM 2100 microscope (JEOL, Tokyo, Japan) equipped with an EDS analyzer (Oxford Instruments plc, Abingdon, UK) was used for transmission electron microscopy (TEM) analysis. The microscope operating voltage was 200 kV. The disc-shaped specimens with 3 mm diameter and 0.2 mm thickness were used for probe preparation. The specimens were subjected to a twin-jet polishing at a voltage of 21 V and a temperature of −(20–25) °C using a 30% HNO3 solution in methanol and a Struers TenuPol-5 machine (Struers APS, Ballerup, Denmark).
The grain structure of samples before the start of the superplastic deformation was analyzed with an optical microscopy (OM) in a polarized light mode. The Zeiss Axiovert 200 M optical microscope (Carl Zeiss, Oberkochen, Germany) was used. The samples for OM were polished with OP-S suspension (Struers APS, Ballerup, Denmark) and further anodized with Barker’s reagent at a temperature of 0 °C and a voltage of 16 V.
The Walter Bai LFM-100 machine (Walter + Bai AG, Löhningen, Switzerland) was used to study the elevated temperature deformation behavior of the studied alloys following ASTM-E2448-11 standard. The uniaxial tensile tests were performed in a temperature range of 440–520 °C and a constant strain rate range of 2 × 10−3–5 × 10−2 s−1. The machine crosshead velocity is increased according to the Equation (1) to an accuracy of ±1% to maintain a constant true strain rate until a predetermined strain value is reached or until fracture.
V = ε ˙ [ L 0 ( 1 + e ) ]
where V is crosshead velocity,
L0 is initial gauge length,
e is true strain.
The Dion-Pro+ software (v.4.8, Walter + bai company, Löhningen, Switzerlan) was used to control the crosshead movement during the test. The standard samples with the gauge length L 0 = 5.65 F 0 , where F0 is the cross sectional area, and the initial gauge section dimensions of 6 × 1 × 14 mm3 were used for the tensile tests. Three samples were analyzed per each testing regime.
The derived “m” value was determined according to the ASTM-E2448–11 (American Society for Testing and Materials) standard using a step test, in which the true strain rate was periodically stepped to 20% above nominal, then back to nominal, starting at a true strain of 0.15 and stepping up and down every 0.1 strain.
The room temperature mechanical properties were measured according to ASTM E8/E8M standard using a Zwick Z250 universal testing machine (Zwick/Roell, Ulm, Germany). Three samples per state were analyzed.

3. Results and Discussion

3.1. Microstructural and Phase Composition Analyses

The XRD analysis was performed to reveal the phase composition for the studied alloys in the homogenized state (Figure 1). According to the obtained spectrums, the (Al), Al3CuCe, Al11Ce3, Al12Ce10Si8, and Mg2Si phases solidified in the alloys studied.
The aluminum based solid solution and eutectic originated phases at the periphery of dendrite cells were observed after the homogenization annealing (Figure 2). According to the SEM-EDS analysis, the phase with dark contrast was enriched with Mg and Si and corresponded to Mg2Si phase in the 0Ce (Ce-free) alloy (Figure 2a). The Mg2Si phase volume fraction was 3.1 ± 0.3%. A negligible fraction of the Sc and Si bearing phase was also found. The solidification of the AlSc2Si2 phase [66,67] is suggested.
The dark particles enriched with Mg and Si and attributed with the Mg2Si phase were also observed in the Ce–bearing alloys. The Mg2Si-phase with volume fraction was 0.6 ± 0.1 for the 2Ce alloy, and it was noticeably smaller at 0.2 ± 0.1 for the 4Ce alloy (Figure 2b,c).
In the Ce-bearing alloys, a high fraction of bright contrast eutectic phases enriched with the Si, Cu, and Ce were observed (Figure 2b,c). The eutectic phases particles predominately exhibited a needle-shaped morphology. The volume fraction of the particles was 5.5 ± 0.4% and 10.5 ± 0.6% for the 2Ce and 4Ce alloys, respectively.
The TEM-EDS analysis of the eutectic originated particles in the 4Ce alloy revealed the Ce, Cu, Al, Si, Zr, and Sc elements (Figure 3). The concentrations of the Sc and Zr in the particles (Spectrum 1 and 2 in Figure 3a and Table 2) were similar to that in the solid solution (Spectrum 3 in Figure 3a and Table 2). The obtained spectrums suggested two types of Ce-rich phases with low Cu and high Si (spectrum 1 in Figure 3a) and with high Cu and low Si (Spectrum 2 in Figure 3a).
The comparative analysis of the XRD, SEM, and TEM results suggested the presence of several Ce-bearing phases in the studied alloys. The Al3CuCe, Al11Ce3, Al12Ce10Si8 were observed by XRD in both 2Ce and 4Ce alloys. The increase in Ce content provide a decrease in the Mg2Si-phase volume fraction that was explained by a formation of Si-and Ce-bearing phases. It should be noted that the increase in Ce content from 2 to 4% resulted in an increase in the heights of the XRD-peaks corresponded to the Al11Ce3-phase, therefore, its fraction can increase. The SEM-EDS and TEM-EDS data demonstrated that the Ce-bearing phase was also enriched with Cu and Si. The TEM-EDS analysis allowed distinguishing two-types of Ce-bearing phases with high Cu/low Si content and low Cu/high Si content. We suggested that the Si partially dissolved in the Al3CuCe, whereas the Cu dissolved in the Al12Ce10Si8 phases and both elements dissolved in the Al11Ce3 phase without the changes of the phases’ lattice structure. As an example we can provide the following papers [67,68,69] that revealed the dissolution of Si in Al3(Sc,Zr) phase or dissolution of Cr and Fe in the Al6Mn phase.
The mechanical treatment provided the fragmentation of the eutectic originated particles that were uniformly distributed in the matrix (Figure 4). The Mg2Si particles in the 0Ce alloy exhibited a mean size of 1.8 ± 0.1 µm and a volume fraction of 3.0 ± 0.1% (Figure 4a). After sheet processing, the Mg2Si phase particles size was 1.3–1.8 µm and their volume fraction was 0.5 ± 0.1% and 0.2 ± 0.1% for the 2Ce and 4Ce alloys, respectively. The Ce-bearing particles exhibited a mean size of 1.9 ± 0.1 µm and 1.4 ± 0.1 µm for the 2Ce and 4Ce alloys, respectively (Figure 4b,c). The particles size varied in a range of 0.4 to 5.6 µm. The particles size distribution histograms demonstrated near-normal size distribution. The volume fraction of the particles was 5.5 ± 0.4% and 10.5 ± 0.6% for the 2Ce and 4Ce alloys, respectively.
The TEM study of the alloys revealed a high number density of dispersoids with a mean size of 10 ± 1 nm (Figure 5). The selected area electron diffraction (SAED) of the analyzed area demonstrated the reflection pattern of Al [110] zone axis and ordered superlattice reflections corresponded to the L12-phase (insert in Figure 5b). The high-resolution image revealed the coherency of the dispersoids and Al matrix (Figure 5c) that agrees with earlier studies performed in Sc- and Zr-bearing Al-Mg-Si-based alloys [35,70]. The fast Fourier transform (FFT) of the high-resolution image also confirmed the L12 structural type of the precipitates (insert in Figure 5c). The L12 precipitate’s size and structure were the same for the alloys studied.
To analyze the grain structure of the studied alloys before the start of superplastic deformation, the sheets were annealed in a temperature range of 460–500 °C for 20 min (Figure 6). In studied conditions the sheets exhibited non-recrystallized banded grain structure. The high number density and small size of L12 dispersoids provided a strong dislocation pinning effect and inhibited recrystallization during heating. Therefore, the thermomechanical treated samples demonstrated a non-recrystallized structure before the start of the superplastic deformation.

3.2. The Superplastic Deformation Behavior

The superplastic behavior for the studied alloys was analyzed in a temperature range of 460–500 °C and a constant strain rate range of 2 × 10−3–5 × 10−2 s−1. The obtained true stress vs true strain curves are presented in Figure 7. The flow stress decreased with a decrease in strain rate and increase in temperature. The increase in Ce content resulted in a decrease in stress values and increase in elongation-to-failure. The strain hardening effect was observed at low strains and a strain softening occurred at larger strain values. A similar behavior was observed for various superplastic alloys with initially non-recrystallized grain structure [71,72]. Strain hardening is associated to the grain growth effect [11,18]. In the studied alloys, the strain hardening effect was observed at low strains owing to the increased dislocation density. The strain softening is attributed with dynamic recrystallization (DRX). Increase in the dislocation density to a critical level is required to start the dynamic recrystallization (DRX) [73]. When above the critical point, the DRX starts and helps to form a fine-grained structure. Due to a simplifying of the DRX, the strain softening effect is weakened with an increase in the deformation temperature and the Ce content in the alloys.
The maximum elongation-to-failure of 400–500% was observed at 480 °C for the 4Ce alloy in a strain rate range of 2 × 10−3–1 × 10−2 s−1. A better superplastic behavior of the 4Ce alloy was the result of a higher volume fraction of the coarse particles that provided the PSN effect. As a result, the coarse particles simplified the DRX and this led to a grain refinement during the deformation. At a low deformation temperature of 460 °C and elevated temperature of 500 °C, an elongation to failure increased less significantly with an increase in the fraction of coarse particles, however the maximum value of 350 ± 10% was also observed for the 4Ce alloy.
A strain-induced change of the strain rate sensitivity m-coefficient was analyzed for the studied alloys at 480 °C that provided the maximum elongation to failure (Figure 8). For the nominal strain rate values of 5 × 10−3 s−1 and 1 × 10−2 s−1, the calculated values of the m-coefficient varied from 0.25 to 0.43 for the 0Ce alloy and from 0.28 to 0.45 for the 2Ce and 4Ce alloys. An increase in the m-value was observed at small strains, which agreed with DRX behavior. A similar strain induced change to the m-value was found for alloys with initial non-recrystallized grain structure in Ref. [11].
According to the Equation (2) [74] the flow stress at superplastic deformation (σ) depends on the deformation temperature (T), the strain rate ( ε ˙ ), and the strain (ε),
σ = f ( T , ε , ˙ ε )
The Zener–Holloman parameter (Z) is used to describe the hot deformation behavior (Equations (3) and (4)) [75,76].
Z = ε ˙ exp ( Q R T   )
ε ˙ = A f ( σ ) exp ( Q R T   ) = { A 1 σ n 1 ( exp ( Q 1 R T   ) ) A 2 exp ( β σ ) exp ( Q 2 R T   ) A 3 [ sin h ( α σ ) ] n 2 ( exp ( Q 3 R T   ) )
where A1,2,3, α (α = β/n1), β, n1, n2 are the material constants that depend on the effective strain, Q1,2,3 are the effective activation energy of superplastic deformation, kJ/mol, T is the absolute temperature, K, R is the universal gas constant, 8.314 J/(mol·K). A mean strain rate sensitivity m-coefficient is calculated as 1/n1.
The sigmoidal-type equation covers the deformation behavior well with a wide strain rate and temperature ranges, and therefore, this model is often applied to analyze the activation parameters of the superplastic deformation [77,78]. The calculated mean values of the m-coefficient for the studied alloy in a temperature range of 460–500 °C and strain rate range of 2 × 10−3–1 × 10−2 s−1 was 0.34–0.40 that excess the threshold of superplastic deformation (Table 3). The experimental m-values determined via a step test at the optimal temperature of 480 °C and strain rate of 5 × 10−3 s−1 was 0.39–0.42 at a strain of 0.41. The m = 0.5 is associated with the grain boundary sliding (GBS) as the main superplastic deformation mechanism [74]. Lower m values for the alloys studied can be explained by partially recrystallized grain structure and less developed GBS.
The effective activation energy of the superplastic deformation Q3 decreased from 100 to 69 kJ/mol with an increase in the Ce content from 0 to 4%. Considering error bars, the calculated values of the activation energy were close to the activation energy of the grain boundary self-diffusion in aluminum (Q = 84–86 kJ/mol [74,79,80]. Similar effective activation energy values of the superplastic deformation were observed for various aluminum-based alloys in Refs. [53,81].
The decrease in Q values with increasing Ce was the result of an intensification in dynamic recrystallization. The increase in Ce provides a stronger PSN effect and refines grain. As a result, the fraction of high-angle grain boundaries increases, and the grain boundary diffusion intensifies. Both the calculated effective activation energy and m-coefficient values speak for the grain boundary sliding is a predominant deformation mechanism in studied alloys.
The grain/sub-grain structure was studied with EBSD technique after 200% of the superplastic deformation at 480 °C and 1 × 10−2 s−1 (Figure 9). An increase fraction of low angle grain boundaries was observed in the alloys studied, therefore, the microstructure was partially non-recrystallized, even after the superplastic deformation. The misorientation angle distribution transformed from the low-angle grain boundaries domination to a high angle one with an increase in Ce content and the related increase fraction of the coarse particles. The Ce-bearing alloys demonstrated finer grain structure. The 0Ce alloy exhibited the grain and sub-grain sizes of 4.3 ± 0.2 µm and 3.9 ± 0.1 µm, respectively (Figure 9a). The Ce-bearing alloys demonstrated smaller grain and sub-grain sizes. The grain/sub-grain sizes were 3.8 ± 0.3 µm/3.3 ± 0.1 µm for 2Ce alloy and 3.4 ± 0.2 µm/3.2 ± 0.1 µm for 4Ce alloy, respectively (Figure 9b,c).

3.3. Room Temperature Mechanical Properties

The mechanical properties at room temperature were analyzed using uniaxial tensile test and hardness measurement (Table 4). The thermomechanical-treated samples were preliminarily subjected to a solid solution treatment (SST) at 520 °C for 20 min and subsequently aged at 180 °C for 8 h. After SST and water cooling, the strength characteristics of the studied alloys were not significantly different. Due to an increase in the volume fraction of Ce-bearing particles, the hardness changed from 65 ± 2 to 72 ± 3 HV and yield strength (YS) from 140 ± 1 to160 ± 1 MPa with an increase in Ce content from 0 to 4%. The subsequent ageing results in a noticeable increase in the strength characteristics of the Ce-free alloy. The hardness, YS and ultimate tensile strength (UTS) values for the 0Ce alloy were 126 ± 3 HV, 367 ± 1 MPa and 420 ± 4 MPa, respectively, that correlate well with data in Ref. [70]. The strengthening effect at aging of the Al-Mg-Si-(Cu) alloys is due to the precipitation of β″, β′, Q′ metastable phases [4,8,9] The strength characteristic for the 0Ce alloy were approximately 50 MPa higher than that for the conventional AA6013 T6-treated sheets [82]. The increase in strength can be explained by a strengthening effect of the L12 phase precipitates. The L12 dispersoids can increase the alloy’s strength following the dislocation shearing and Orowan dislocation bypass looping mechanisms. The Al3Zr L12-structured dispersoids smaller than 5.6 nm in size [79] are sheared by dislocation, whereas coarser precipitates induce Orowan strengthening that can be estimated following the Equation (5a):
Δ σ o r = M · 0.4 · G b π ( 1 ν ) · ln ( 2 R ¯ r 0 ) λ ,  
where ν = 0.345 is the Poisson ratio, M~3.0 is the Taylor factor, G = 26GPa is the shear modulus, b = 0.286 nm is the Burgers vector, R ¯ = π · D s / 8 and Ds are the mean precipitate radius and the diameter, r0 = 1.5b, λ is the interparticle space (nm):
λ = 0.5 D s · ( 2 π 3 φ π 4 ) ,  
where φ is the volume fraction of precipitates.
The concentrations of Zr and Sc in solid solution were about 0.15 and 0.12%, respectively. Taking into account that Zr and Sc are fully precipitated from the solid solution after annealing, the L12 dispersoid volume fraction was estimated as 0.48%. Considering a mean dispersoid size of 10 nm, the maximum Orowan strengthening is approximately 100 MPa.
The alloying with Ce provides a noticeable decrease in the alloy’s room temperature strength and an increase in ductility at 4 wt.% Ce. The strength properties degradation can be explained by a formation of the Si and Cu-enriched Ce-bearing phases. As a result, the solute Si in the Ce-bearing alloys decreased and, therefore, the further aging effect due to precipitation of Si and Cu-bearing β″, β′, Q′ phases was weak.
It is notable that a low temperature homogenization provided an increased fraction of Mg2Si phase of solidification origin and a high number density of the L12 Al3(Sc,Zr) precipitates for Ce-free alloy with a low Sc content. Such a bimodal particle size distribution provided a fine-grained structure and the superplastic behavior with m ≈ 0.4 and elongation above 300%. At the same time, the Ce-free alloy demonstrated an increased room temperature strength. Alloying with Ce provided a high fraction of the coarse particles that improved the PSN effect, provided grain refinement, and superplasticity. Due to a collaboration of the Zener pinning effect caused by fine L12-dispersoids and PSN effect provided by coarse Ce-bearing particles, the alloys with Ce demonstrated a good superplasticity with m = 0.42 and 400–500% of elongation at (0.2–1) × 10−2 s−1. The disadvantage of the Ce alloying of the studied Al-Mg-Si-Cu based alloys is the dissolution of the precipitation strengthening Si and Cu elements into Ce bearing phase. As a result, Si and Cu content in a solid solution decreased and the aging effect decreased due to the Ce. To improve the strength properties of the Ce bearing Al-Mg-Si-alloys, further chemical composition optimization is required.

4. Conclusions

In this paper, the microstructural evolution, superplastic deformation behavior, and mechanical properties of the Al-Mg-Si-Cu-based alloys containing eutectic forming Ce and dispersoid forming Sc and Zr elements were studied. The main conclusions were drawn as follows:
According to the XRD, SEM, and TEM results, the Al3CuCe, Al11Ce3, and Al12Ce10Si8 phases were solidified in the Ce bearing alloys, and Mg2Si was formed in the Ce-free and Ce-bearing alloys. The EDS analyses showed that the alloying elements were dissolved into Ce bearing phases; the dissolution of the Si in the Al3CuCe, Cu in the Al12Ce10Si8 and both Si and Cu in the Al11Ce3 phase without the changes of the phases’ lattice structure were observed. The increase in the Ce content resulted in an increase in the volume fraction of the Ce-bearing phase particles and a decrease in the volume fraction of the Mg2Si phase of eutectic origin.
The thermomechanical treatment, including two-step homogenization, with hot and cold rolling, provided a bimodal particle size distribution of the nanoscale L12-structured Al3(Sc,Zr) phase precipitates of 10 ± 1 nm size and the coarse particles of eutectic origin. In the Ce-free alloy, the volume fraction of the Mg2Si particles was 3% and the size was 1.8 ± 0.1 µm. In the Ce-bearing alloys the volume fraction of the Mg2Si phase was below 0.01. The mean size of the Ce-bearing particles was 1.9 ± 0.1 µm and 1.4 ± 0.1 µm, the volume fraction was 5.5 and 10.5% for alloys with 2% and 4% Ce, respectively.
Due to the pronounced particle stimulated nucleation effect, an increase in the Ce-content improved the superplastic properties, increased the value of elongation-to-failure and decreased flow stress values of the alloys studied. The maximal elongation-to-failure of 500% was achieved at 480 °C and 5 × 10−3 s−1 for the alloy with 4% Ce.
The maximal values of the YS and UTS of 367 ± 1 and 420 ± 4 MPa, respectively, were achieved for the Al-Mg-Si-Cu based Ce-free alloy in the T6 state. The Si and Cu dissolution in the Ce-bearing phases weakened the aging effect and led to a significant decrease in strength properties.

Author Contributions

Conceptualization, A.G.M. and A.V.M.; methodology, N.Y.T. and A.S.P.; software, A.G.M.; validation, A.S.P. and A.G.M.; formal analysis, A.S.P.; investigation, A.G.M. and A.V.M.; data curation, A.G.M.; writing—original draft preparation, A.G.M.; writing—review and editing, A.V.M.; visualization, A.G.M.; supervision, A.V.M. All authors have read and agreed to the published version of the manuscript.

Funding

The idea of the study, SEM, EBSD and TEM investigations, mechanical tests were funded by the Russian Science Foundation Grant No. 20-79-00269. The calculations and discussion on the theoretical contribution of the strengthening mechanisms were done in a framework of State Task to Universities of Russian Federation (No. 0718-2020-0030). For TEM studies the equipment of Collective Use Equipment Center “Materials Science and Metallurgy” that works due to financial support of the Russian Federation represented by the Ministry of Higher Education and Science (No. 075-15-2021-696) was used.

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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Figure 1. The XRD data for the studied alloys after a two-step homogenization annealing.
Figure 1. The XRD data for the studied alloys after a two-step homogenization annealing.
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Figure 2. SEM-BSE micrographs and corresponded SEN-EDS maps for (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys after a two-step homogenization.
Figure 2. SEM-BSE micrographs and corresponded SEN-EDS maps for (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys after a two-step homogenization.
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Figure 3. (a) TEM micrographs of the eutectic originated particles in the as-homogenized 4Ce alloy; (b) the TEM-EDS spectrums for the areas 1–3 marked-up with circles in (a).
Figure 3. (a) TEM micrographs of the eutectic originated particles in the as-homogenized 4Ce alloy; (b) the TEM-EDS spectrums for the areas 1–3 marked-up with circles in (a).
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Figure 4. SEM-BSE images of the studied alloys in the as-cold rolled state and corresponded particle size distribution histograms.
Figure 4. SEM-BSE images of the studied alloys in the as-cold rolled state and corresponded particle size distribution histograms.
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Figure 5. TEM micrographs for the 4C alloy; (a) bright field; (b) dark field; (c) high-resolution image of dispersoid; insert in (b) is corresponding SAED and insert in (c) is corresponding FFT.
Figure 5. TEM micrographs for the 4C alloy; (a) bright field; (b) dark field; (c) high-resolution image of dispersoid; insert in (b) is corresponding SAED and insert in (c) is corresponding FFT.
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Figure 6. Grain structure of the (ac) 0Ce, (df) 2Ce, (gi) 4Ce alloy sheets after 20 min annealing at (a,d,g) 460 °C, (b,e,h) 480 °C, (c,f,i) 500 °C.
Figure 6. Grain structure of the (ac) 0Ce, (df) 2Ce, (gi) 4Ce alloy sheets after 20 min annealing at (a,d,g) 460 °C, (b,e,h) 480 °C, (c,f,i) 500 °C.
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Figure 7. (ac) True stress vs true strain curves for 0Ce, 2Ce, and 4Ce alloys for a constant strain rates of 2 × 10−3, 5 × 10−3, and 1 × 10−2 s−1; (df) elongation-to-failure diagrams as a function of strain rate and Ce content in the alloys for the deformation temperatures of (a,d) 460, (b,e) 480, (c,f) 500 °C.
Figure 7. (ac) True stress vs true strain curves for 0Ce, 2Ce, and 4Ce alloys for a constant strain rates of 2 × 10−3, 5 × 10−3, and 1 × 10−2 s−1; (df) elongation-to-failure diagrams as a function of strain rate and Ce content in the alloys for the deformation temperatures of (a,d) 460, (b,e) 480, (c,f) 500 °C.
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Figure 8. Strain dependence of the stress and m-value during the step test with periodically stepped true strain rate to 20% above nominal of 5 × 10−3 s−1 and 1 × 10−2 s−1, then back to nominal for (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys at 480 °C.
Figure 8. Strain dependence of the stress and m-value during the step test with periodically stepped true strain rate to 20% above nominal of 5 × 10−3 s−1 and 1 × 10−2 s−1, then back to nominal for (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys at 480 °C.
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Figure 9. EBSD-IPF grain boundary maps and corresponded misorientation angle distribution histograms for the (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys after 200% of superplastic deformation at 480 °C and 1 × 10−2 s−1 strain rate.
Figure 9. EBSD-IPF grain boundary maps and corresponded misorientation angle distribution histograms for the (a) 0Ce, (b) 2Ce, and (c) 4Ce alloys after 200% of superplastic deformation at 480 °C and 1 × 10−2 s−1 strain rate.
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Table 1. Chemical compositions of the studied alloys (wt.%).
Table 1. Chemical compositions of the studied alloys (wt.%).
AlloyMgSiCuCeScZrAl
0Ce1.20.71.0-0.100.2bal.
2Ce1.20.71.02.00.100.2bal.
4Ce1.20.71.04.00.100.2bal.
Table 2. Elemental composition (at.%) corresponded to the spectrums 1–3 in Figure 3b.
Table 2. Elemental composition (at.%) corresponded to the spectrums 1–3 in Figure 3b.
AlMgSiCuCeScZr
Spectrum 143.400.547.9026.4021.400.100.02
Spectrum 243.500.1724.202.5029.400.180.11
Spectrum 397.721.600.030.50-0.120.15
Table 3. The material parameters and the effective activation energy of the superplastic deformation values for the Equation (4) calculated for a true strain of 0.41 (50%).
Table 3. The material parameters and the effective activation energy of the superplastic deformation values for the Equation (4) calculated for a true strain of 0.41 (50%).
Strain
[%]
Q1
[kJ/mol]
n1/mlnA1Q2
[kJ/mol]
β
[MPa−1]
lnA2Q3
[kJ/mol]
n2lnA3
0Ce98 ± 42.89/0.342.1100 ± 50.167.8100 ± 72.2110.3
2Ce92 ± 32.72/0.371.998 ± 70.177.496 ± 72.069.5
4Ce67 ± 82.46/0.401.170 ± 100.183.169 ± 81.875.2
Table 4. Mechanical properties of the thermomechanical-treated samples after solid solution treatment and artificial aging (standard deviation are presented as an error bar).
Table 4. Mechanical properties of the thermomechanical-treated samples after solid solution treatment and artificial aging (standard deviation are presented as an error bar).
SST at 520 °C (20 min) SST and Aging at 180 °C for 8 h
AlloyHVYS [MPa]UTS [MPa]El [%]HVYS [MPa]UTS [MPa]El [%]
0Ce65 ± 2140 ± 1256 ± 320 ± 2126 ± 3367 ± 1420 ± 49 ± 1
2Ce68 ± 3150 ± 2255 ± 515 ± 1102 ± 2281 ± 2324 ± 210 ± 1
4Ce72 ± 3160 ± 1250 ± 511 ± 274 ± 2192 ± 7260 ± 1013 ± 1
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Mochugovskiy, A.G.; Prosviryakov, A.S.; Tabachkova, N.Y.; Mikhaylovskaya, A.V. The Effect of Ce on the Microstructure, Superplasticity, and Mechanical Properties of Al-Mg-Si-Cu Alloy. Metals 2022, 12, 512. https://doi.org/10.3390/met12030512

AMA Style

Mochugovskiy AG, Prosviryakov AS, Tabachkova NY, Mikhaylovskaya AV. The Effect of Ce on the Microstructure, Superplasticity, and Mechanical Properties of Al-Mg-Si-Cu Alloy. Metals. 2022; 12(3):512. https://doi.org/10.3390/met12030512

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Mochugovskiy, Andrey G., Alexey S. Prosviryakov, Nataliya Yu. Tabachkova, and Anastasia V. Mikhaylovskaya. 2022. "The Effect of Ce on the Microstructure, Superplasticity, and Mechanical Properties of Al-Mg-Si-Cu Alloy" Metals 12, no. 3: 512. https://doi.org/10.3390/met12030512

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