Effect of Dispersoids and Intermetallics on Hardening the Al-Si-Cu-Mg Cast Alloys

Hernandez-Sandoval, J.; Abdelaziz, M. H.; Samuel, A. M.; Doty, H. W.; Samuel, F. H.

doi:https://doi.org/10.1155/2021/9933168

Advances in Materials Science and Engineering

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Research Article | Open Access

Volume 2021 | Article ID 9933168 | https://doi.org/10.1155/2021/9933168

Effect of Dispersoids and Intermetallics on Hardening the Al-Si-Cu-Mg Cast Alloys

J. Hernandez-Sandoval,¹M. H. Abdelaziz,²A. M. Samuel,³H. W. Doty,⁴and F. H. Samuel³

Academic Editor: Jörg M. K. Wiezorek

Received14 Mar 2021

Accepted28 Apr 2021

Published11 May 2021

Abstract

The principal aim of the present research work was to compare the role of dispersoids Al₂O₃ (∼15 μm average particle size) and SiC (∼15 μm average particle size) with that offered by Zr- and Ni-based intermetallics (∼35–70 μm average particle size) on the hardening of cast aluminum 354 alloy (9.1% Si, 0.12% Fe, 1.8% Cu, 0.008% Mn, 0.6% Mg, and 87.6% Al) at ambient temperature. There is no observable poisoning effect on the refinement of grain size after the addition of Zr to the alloys investigated in this study. The tensile test results were examined in light of the microstructural features of the corresponding alloy samples. The contribution of the added dispersoids or Ni and Zr alloying elements on the tensile properties of the 354 alloys was determined employing ∆P plots (where P = Property, UTS, YS, or %El), using the base alloy (in the as-cast condition) as a reference point. The tensile results were supported by investigating the precipitation-hardening phases using scanning and transmission electron microscopy as well as examining the fracture surfaces of selected conditions applying field emission scanning electron microscopy. The results show that, in all cases, Al₂Cu phase is the main hardening agent. The contribution of about 1.5 vol% of SiC or Al₂O₃ to the strength of the base alloy is higher than that offered by Zr- and Ni-based intermetallics, under the same aging treatment.

1. Introduction

Dispersoids are defined as small and numerous finely divided particles of one substance dispersed in another. Dispersoid particles are normally hard and brittle, resulting in excellent dispersion strengthening of the softer aluminum matrix [1]. According to Srinivasan and Imam [2], dispersoids are thermodynamically stable second-phase particles. Their size is in the range of 0.1 to 1.0 μm, which is made use of to reduce the alloy grain size. Fracture of these particles under loading can lead to the nucleation of cracks ahead of the main advancing crack. The effect of Cr and Mn addition and heat treatment on AlSi3Mg casting alloy was investigated by Tocci et al. [3] who found that Cr- and Mn-based dispersoid particles improve Vickers microhardness of the aluminum matrix compared to the base alloy after solution treatment and quenching, which is in accordance with the dispersion-hardening mechanism. In a similar study, Remøe et al. [4] recommended decreasing heating rates which would increase the dispersoid density.

Dispersoids were also described as particles that form during ingot preheating caused by precipitation of the transition elements Cr, Mn, or Zr in the form of Al₁₂Mg₂Cr, Al₂₀Cu₂Mn₃, or Al₁₂Mn₃Si and Al₃Zr particles, respectively [5]. A significant increase in mechanical properties of the tested hypoeutectic alloys was obtained by supersaturation of the Al phase with high melting point elements during HPDC processing of AlSi9Cu3(Fe) alloy as reported by Szymczak et al. [6]. A study of dispersoid particles in two Al–Mg–Si 6xxx-type aluminum alloys and their effects on the recrystallization was conducted by Hichem and Rebai [7]. Their main conclusion was that the incorporation of Si in the metastable Al₃Zr type of dispersoid particles could increase the nucleation rate by reducing the volume-free energy of formation of the dispersoid particles. Kenyon et al. [8] added that changing the dispersoid volume fraction, size, and morphology has important implications for the pinning effectiveness of the dispersoids.

It is necessary to produce a microstructure containing thermally stable and coarsening-resistant dispersoids in order to enhance the mechanical properties of an aluminum alloy, in particular at high temperatures. The dispersoids or particles play an important role in resisting coarsening if the energy of their interface with the matrix, diffusivity, and solubility is low. The Al₃Zr particles are resistant to dissolution and coarsening; they can also control the evolution of the grain and subgrain structure, thereby making it possible to increase strength and ductility in the precipitation-hardened T6 condition [9–12]. The effectiveness of the dispersoids depends on their size, spacing, and distribution. In direct-chill cast alloys, the alloying elements are highly segregated following solidification [13–15].

According to Szajewski et al. [16], three different types of dislocation-precipitate interactions could occur depending on the nature of the precipitates, i.e., whether they are coherent or noncoherent. Strengthening by coherent precipitates is caused by distortion in the crystal lattice, representing barriers to dislocation movement [17–20]. Several research articles [6, 8, 21–23] emphasized that binary dispersoids such as Al₃Zr can be less effective due to limitations encountered during conventional manufacturing processes. Improvements can be achieved through the transformation of these binary trialuminides into ternary or quaternary forms by alloying with appropriate elements.

Previously, the present authors discussed the effect of addition of SiC and Al₂O₃ particulates on the microstructure and tensile properties (mainly quality index values as a function of thermal treatment) of Al–Si–Cu–Mg cast alloys [23]. In addition, the authors analyzed the effect of different alloying elements on the performance of this family of alloys over several publications [24–28]. The subject was readdressed in the present work to highlight the difference in the effectiveness of dispersoids, mainly Al₂O_3, SiC and Al₃Zr, compared to intermetallics offered by the addition of Zr and Ni in the amounts of 0.25 wt.% and 0.5 wt.%.

2. Experimental Procedure

Alloy 354 modified with 200 ppm strontium and grain-refined using 0.25 wt% Ti (added using Al-5%Ti-1%B master alloy) was used as the base alloy (coded alloy A). The chemical composition in wt. % of the as-received 354 alloy ingots is listed in Table 1.

2.1. Additions

(i)The addition of Al₂O₃ particles was carried out using a metal matrix composite (6061 alloy + 20 vol % Al₂O₃) having a 2-to-30 μm particle size.(ii)The addition of SiC particles was done using a (359 alloy + 20 vol % SiC) metal matrix composite having a similar particle size.(iii)Zirconium and nickel additions were made using Al-20wt% Zr and Al-20wt% Ni master alloys, respectively.

Alloy A, with a chemical composition as shown in Table 1, was melted in an electrical furnace at 780°C, using a 60 kg SiC crucible. The molten alloy was degassed using pure, dry argon injected into the melt for 20 min by means of a graphite impeller rotating at 135 rpm. Grain refining and modification of the melt were carried out using Al-5% Ti-1% B and Al-10% Sr master alloys, respectively, to obtain levels of 0.25 wt.% Ti and 200 ppm Sr in the melt. Table 2 lists the chemical composition of the alloys which were investigated. The melt was poured into a preheated ASTM B-108 permanent mold (preheated to 460°C), to prepare test bars; the dimensions of which are shown in Figure 1. Three samples for chemical analysis were also taken at the time of the casting; this was done at the beginning, in the middle, and at the end of the casting process to ascertain the exact chemical composition of each alloy.

As will be mentioned in a later subsection, heat treatment of the test bars used for tensile testing involved solution heat treating them at 495°C for 8 hours, followed by quenching in warm water at 60°C, after which artificial aging was applied. In the present work, three aging temperatures (155°, 190°, and 350°C) and different aging times, up to 1000 h were used. After aging, the test bars were allowed to cool naturally at room temperature (25°C). All heat treatments were carried out in an air-forced Lindberg Blue M electric resistance furnace. All of the samples, whether as-cast, solution heat-treated, or aged, were tested to the point of fracture using an MTS servo-hydraulic mechanical testing machine at a strain rate of 4 × 10⁻⁴ s⁻¹.

Ten samples from each condition were tested, for a total of 80 conditions for each alloy. Three other conditions were tested for each alloy corresponding to 200 h, 600 h, and 1000 h aging times, from which would be selected one specific aging temperature (190°C); this will be explained in detail further on in the text. The values for the mechanical properties, namely, ultimate tensile strength (UTS), yield strength (YS), and percent elongation (% El), were gathered from the computerized system of the MTS machine.

Samples for metallography were sectioned from the tensile-tested bars of all the alloys studied, about 10 mm below the fracture surface; they were then individually mounted in Bakelite and subsequently polished to a fine finish using 1 μm diamond suspension. The instrument used in this study was a Hitachi-SU-8000 FESEM, equipped with a standard secondary electron detector (SE), a backscatter electron detector (BSD), and an energy dispersive X-ray spectrometer (EDS). The FESEM was operated at a voltage of 15 kV, with a maximum filament current of 3 amperes.

Transmission electron microscopy was used in order to observe and identify the strengthening precipitates in heat-treated samples and also to investigate the coherency of the precipitates with the matrix. The FEI Tecnai G² F20 electron microscope employed was equipped with an advanced control system which permits the integration of an EDAX™ chemical analysis system, scanning transmission electron microscopy (STEM), and electron energy loss spectroscopy (EELS). The microscope was operated at an accelerating voltage of 200 kV.

3. Results and Discussion

In the present work, we are referring to the Al₂O₃ or SiC particulates as dispersoids in composites B and C containing these particulates since these particulates have no solubility in the matrix, maintaining at the same time uniform spacing. On the contrary, intermetallics are phases/compounds made up of two or more elements, producing a new phase with its own composition, crystal structure, and properties (Alloys D-G). Alloy A is the base alloy.

3.1. Microstructural Characterization

Figure 2(a) shows the microstructure of alloy A in the as-cast condition, revealing almost no presence of porosity. The blue arrows point to the change in the morphology of dendrites from elongated ones to a more rounded (rosetta-like) form which enhances feeding of the interdendritic regions, reducing the porosity in the grain-refined alloy as clearly displayed in Figure 2(b), a high-magnification micrograph of Figure 2(a). According to Campbell and Tiryakiog˘lu [29], based on the work of Fuoco et al. [30], at an early stage of solidification, the dendrites can move as a slurry. Later, when eutectic freezing starts, the dendrites are fully solidified and fixed (about 45% remaining liquid) so that significant mobility of the eutectic is to be expected. However, after adding Sr, no mobility of the residual liquid is observed even at the high level of 45%. This explanation, in addition to the change in the dendrite shape presented in Figure 2(b), contradicts the model presented by Argo and Gruzleski [31] which suggests that, in unmodified alloys, the eutectic is characterized by its irregular solid/liquid interface, leading to entrapping of small pockets of liquid between advancing solidification fronts, causing the formation of microporosity. Figures 2(c) and 2(d), respectively, exhibit the distribution of the added Al₂O₃ and SiC particulates throughout the matrix, following solutionizing at 495°C for 8 h (SHT) revealing the complete dissolution of the Al₂Cu phase. Figures 3 and 4 reveal the adherence (no debonding) between the particulates and the surrounding matrix.

In alloy D, as the Al₃Zr precipitate must be present before the peritectic reaction, the formation of Al₃Zr from liquid is called a properitectic reaction [32, 33]. In the present work, the Zr concentration used was 0.2 wt.% and 0.4 wt.% for alloy G and alloys E-F, respectively. From Figure 5(a), the liquidus temperature should be 718°C and 780°C, respectively. The melt was maintained at 780°C, to avoid reaching the softening point of the SiC crucible. Thus, in the present case, Zr-rich phases will be precipitated in the form of intermetallics as well as dispersoids (alloys E and F), whereas in the case of alloy G, most of the Zr will be precipitated in the form of Al₃Zr dispersoids.

(a)

(b)

It should be borne in mind that the present alloys contain at least 15 elements and that the solidification rate was approximately 8°C/s. These two criteria make it difficult to strictly follow the binary equilibrium diagram shown in Figure 5(a). Figure 5(b) depicts the Zr-rich phase distribution in alloy D in the as-cast condition, solidified at a very slow rate ∼0.35°C/s. In this case, the chemical composition of the dispersoid is Al-Zr-Ti as inferred from the associated EDS spectrum displayed in Figure 6(a). In the presence of Ti in the matrix, Zr has a strong affinity to react with Ti forming Al₃(Zr, Ti)₅ [34–36] or Al₃(Zr_1−xTi_x) dispersoids [36]. Figure 7 depicts the precipitation of Al₃Zr dispersoids in alloy D in the as-cast condition, in a sample prepared from a tensile test bar solidified at about 8°C/s. According to Ikhlaq et al. [34], Zr has a preferred orientation (002) along other significant diffraction planes (100), (101), (102), and (110) of Zr confirming the hexagonal crystal structure as shown in Figure 6(b).

(a)

(b)

(a)

(b)

(c)

(d)

Figure 8 displays the progress in the size and distribution of the Al₂Cu precipitation hardening particles (main hardening agent) at 190°C (peak aging). It is evident from Figures 8(b) and 8(c) that with increase in the aging time, there is an increase in the size of the precipitated Al₂Cu particles coupled with change in their morphology from spherical to short rods, which leads to a gradual decrease in contribution of these precipitates to both UTS and YS regardless the alloy type, reaching almost nil after 1000 h [37]. Figure 8(d) shows dislocation/precipitate interactions; note the significant increase in the dislocation density compared to that after solutionizing treatment (Figure 8(e)). Figure 8(f) presents the EDS spectrum, corresponding to the white square area in Figure 8(c), indicating reflections due to Al and Cu.

(a)

(b)

(c)

(d)

(e)

(f)

3.2. Tensile Properties

Table 3 lists the tensile parameters of alloy A in the as-cast condition. Figure 9 presents the contribution of added particulates or alloying elements on the tensile properties of the 354 alloys B through G, using alloy A in the as-cast condition as a reference point. It should be mentioned here that, in a previous study, the authors reported on the aging behavior of alloy A in the temperature rage 155°C–350°C for times up to 100 h [38].

(a)

(b)

(c)

(d)

(e)

As can be seen in Figure 9(a), when the alloys were aged at 155°C for 2 h, composite B (containing maximum 1.5 vol% Al₂O₃) contributed significantly to the UTS of alloy A (about 120 MPa) followed by a linear decrease up to alloy E (40 MPa) beyond which other alloys (alloys F and G) showed a slight increase in UTS, going to 56 and 60 MPa, respectively. Aluminum oxide particulates are a ceramic compound with a hexagonal crystal lattice. The oxygen anions define a hexagonal close packed structure, and the aluminum cations occupy 2/3 of the octahedral sites in the hcp lattice. It is used for its hardness and strength, about 9 on the Mohs scale of mineral hardness (just below diamond) [34]. Although SiC (composed of tetrahedra of carbon and silicon atoms with strong bonds in the crystal lattice) is a very hard and strong material [39], its contribution is slightly lesser than that of Al₂O₃ which may be interpreted in terms of the actual particulate volume fraction. However, both composites B and C offered a better resistance to softening, yield strength, as well as a better improvement in the alloy ductility compared to those obtained from the other alloys containing Ni or Zr. Figures 9(b) to 9(d) demonstrate the variation in the UTS, YS, and % elongation, respectively, when both particulates and alloying elements were added to alloy A. The tensile bars were exposed to aging at 190°C for times up to 1000 h.

A comparison of Figures 9(b) and 9(c) shows that the contribution level to the alloy YS is higher than that to the UTS. Although the contribution to the ductility of alloy A in Figure 9(d) followed a similar pattern as those for UTS and YS, a sharp decrease in the %El was observed for alloys F and G (1000 h). On the contrary, low aging times (2 h–100 h) resulted in a negative change in the % elongation. Some improvement can be seen for alloy G that contains 0.25% Zr + 0.25% Ni producing same levels as composites B and C due to precipitation of Al₃Zr dispersoids.

Daoud and Reif [40] suggested that the hardening process is significantly affected by the presence of micro-oxides, carbides, or SiC. The precipitation takes place on the dislocation lines. As a result of the mismatch between the coefficient of thermal expansion of the matrix and the reinforcement, it is not expected to alter the precipitation sequence. Samuel et al. [41] proposed that the relative amount of hardening precipitates is also influenced by the added dispersoids. In order to achieve a significant improvement in the alloy strength, the authors recommended increasing the reinforcement particulates by an order of magnitude than those used in composites B and C. Apparently, the added amounts of Al₂O₃ or SiC in the present study have a marginal effect on the composite tensile properties compared to Zr and Ni addition.

As will be observed from Figure 10, the highest standard deviations with respect to UTS are exhibited by the alloys containing micro-oxides; also, the variations are more accentuated for the temperatures where the maximum strength values were observed, i.e., at 155°C and 190°C. This figure is showing the average standard deviation values for all alloys studied, where the higher standard deviation values may be observed for the properties pertaining to composites B and C. In Figure 11, the maximum standard deviations correspond to the maximum ductility values, observed at 155°C and 350°C for alloys containing micro-oxides/carbides. In both Figures 10 and 11, the tensile test values corresponding to these two temperatures reveal the highest standard deviations, particularly for composites B and C, when the seven alloys used in this study are subjected to comparison.

3.3. Fracture Behavior

Figure 12(a) represents the fracture surface of alloy A in the as-cast condition (secondary electron image) where severe cracks (arrowed blue) can be seen propagating through the fine-dimpled structure (1.3% elongation). Solutionizing at 495°C for 8 h resulted in a significant increase in the alloy ductility (6.3%) which reflected on the alloy surface fracture as displayed in Figure 12(b) where a deep dimple structure is seen predominating the fracture surface-blue arrows. Slip bands are observed covering the wall of the dimples shown in Figure 12(b). Figure 12(c) exhibits the fracture surface of alloy A aged at 190°C for 2 h revealing severe cracks (white arrows), whereas similar features as those shown in Figure 12(b) were reported covering the fracture surface when the alloy was aged at 350°C for 100 h (overaging), see arrows in Figure 12(d). In composite C, Figure 13, the presence of micro-SiC particulates were seen at the bottom of the coarse dimples (white arrows) when the alloy was solutionized (reaching 5.3% ductility). In all solutionized alloys, clear slip bands were commonly observed as shown by the blue arrows in the figure. The fracture surface of alloy D (containing 0.4%Zr) after solutionizing treatment resulted in enhancing the alloy ductility, reaching about 4.5%. While the SHT did not change the shape of the star-like particles, Figure 14(a) shows several slip lines clearly visible on the edges (black arrows) of the particle, as well as a long crack reaching the star center (white dashed arrow). The associated EDS spectrum, Figure 14(b), exhibits reflections due to Al, Si, Zr, and Ti elements suggesting that the composition of this phase is (Al,Si)₂(Ti,Zr) [28].

(a)

(b)

(c)

(d)

(a)

(b)

Figure 15(a) demonstrates the fracture of alloy E in the as-cast condition. It is evident that the presence of such a significant amount of insoluble Ni-rich phases would cause marked brittleness (1.44%El). The associated EDS spectrum in Figure 15(b), corresponding to the marked spot in (a), shows strong peaks of Al, Ni, Cu, and Si with a small peak of Mg indicating a mixture of Al₃CuNi and Q-Al₅Cu₂Mg₈Si₆ phases. The fracture surface of alloy F (containing about 0.4%Ni + 0.45Zr) is illustrated in Figure 16(a) showing a fractured massive Zr-rich phase particle. The black arrow points to parallel markings on the surface of the particle. The associated EDS spectrum corresponding to the area marked X in (a) reveals that this phase is possibly (Al,Si)₃(Ti,Zr), depending on the Si/Zr intensity ratio [42].

(a)

(b)

(a)

(b)

Based on the EDS spectra presented in Figures 7(c), 14(b), and 16(b), the Si/Zr intensity ratios are 0, 0.5, and 1, respectively, depending on the Zr-rich phase particle size. According to Garza-Elizondo et al. [43], the (Al,Si)₂(Ti,Zr) phase contains 36 wt.% Si and 38 wt.% Zr, whereas the (Al,Si)₃(Ti,Zr) phase contains 11 wt.% Si and 47 wt.% Zr, reaching Al₃Zr when Si is practically nil. The atomic radii of Al, Si, Ti, and Zr are listed in Table 4. The Al₃Zr phase is a closed packed structure (Figure 17(a)) with an interatomic spacing of 0.441 nm. However, there is controversy regarding the particle structure which has been reported as being cubic Ll₂ or tetragonal DO23, as shown in Figure 17(b) [42, 43]. It is also reported that Al₃Zr has the stable DO23 structure, but for small supersaturations of the solid solution, Al₃Zr precipitates with the metastable L1₂ structure and precipitates with the DO23 structure only appear for prolonged heat treatment and high enough supersaturations [44, 45].

(a)

(b)

It is suggested that since the Si atom has a small radius (Table 4) compared to the interatomic spacing in Figure 17(a), the Si atoms can easily diffuse into the Al₃Zr cells during the solidification process. The concentration of Si in the Al₃Zr phase particle is a function of the temperature and the size of the formed Al₃Zr phase particle. In other words, the larger the size of the Al₃Zr phase particle, the lesser is the Si concentration, as shown in Figure 16(b). Considering the atomic radius of Ti (0.145 nm) is very close to that of the Al atom (0.143 nm), this would facilitate the substitution of a certain number of Al atoms by Ti atoms.

4. Conclusions

Based on the results obtained in this study, the following conclusions may be drawn:(1)There is no observable poisoning effect on the refinement of grain size after the addition of Zr to the alloys investigated in this study.(2)The Si/Zr ratio varies depending on the size of Zr-rich phase particles.(3)The standard deviation (SD) of the UTS and %El values is higher in SiC and Al₂O₃ particulate-containing alloys due to their low volume fraction and tendency for segregation.(4)In order to achieve a lower SD, the volume fraction of Al₂O₃ or SiC particulates should be increased by an order of magnitude.(5)The results show that, in all cases, Al₂Cu phase is the main hardening agent. The contribution of about 1.5 vol.% of SiC or Al₂O₃ particulates to the strength of the base alloy is higher than that those offered by Zr- and Ni-based intermetallics.(6)The reduction in mechanical properties due to addition of different elements is attributed principally to the increase in the percentage of intermetallic phases formed during solidification; such phase particles would act as stress concentrators, decreasing the alloy ductility.

Data Availability

Data will be available upon request.

Conflicts of Interest

There are no conflicts of interest between the authors.

References

Y. Li and L. Arnberg, “Precipitation of dispersoids in DC-cast AA3103 alloy during heat treatment,” in Essential Readings in Light Metals, J. F. Grandfield and D. G. Eskin, Eds., pp. 1021–1027, Springer, Cham, Switzerland, 2016.
View at: Publisher Site | Google Scholar
R. Srinivasan and M. A. Imam, “Role of dispersoids on the fatigue behavior of aluminum alloys: a review,” in Fatigue of Materials III, T. S. Srivatsan, M. A. Imam, and R. Srinivasan, Eds., pp. 11–22, Springer, Cham, Switzerland, 2016.
View at: Publisher Site | Google Scholar
M. Tocci, R. Donnini, G. Angella, and A. Pola, “Effect of Cr and Mn addition and heat treatment on AlSi3Mg casting alloy,” Materials Characterization, vol. 123, pp. 75–82, 2017.
View at: Publisher Site | Google Scholar
M. Remøe, I. Westermann, and K. Marthinsen, “Characterization of the density and spatial distribution of dispersoids in Al-Mg-Si alloys,” Metals, vol. 9, no. 1, p. 26, 2019.
View at: Publisher Site | Google Scholar
N. C. W. Kuijpers, F. J. Vermolen, C. Vuik, P. T. G. Koenis, K. E. Nilsen, and S. v. d. Zwaag, “The dependence of the β-AlFeSi to α-Al(FeMn)Si transformation kinetics in Al-Mg-Si alloys on the alloying elements,” Materials Science and Engineering: A, vol. 394, no. 1-2, pp. 9–19, 2005.
View at: Publisher Site | Google Scholar
T. Szymczak, G. Gumienny, L. Klimek, M. Goły, J. Szymszal, and T. Pacyniak, “Characteristics of Al-Si alloys with high melting point elements for high pressure die casting,” Materials, vol. 13, no. 21, p. 4861, 2020.
View at: Publisher Site | Google Scholar
F. Hichem and G. Rebai, “Study of dispersoid particles in two Al-Mg-Si aluminium alloys and their effects on the recrystallization,” Applied Physics A, vol. 119, no. 1, pp. 285–289, 2015.
View at: Publisher Site | Google Scholar
M. Kenyon, J. Robson, J. Fellowes, and Z. Liang, “Effect of dispersoids on the microstructure evolution in Al–Mg–Si alloys,” Advanced Engineering Materials, vol. 21, pp. 1–7, 2019.
View at: Publisher Site | Google Scholar
T. Szymczak, J. Szymszal, and G. Gumienny, “Statistical methods used in the assessment of the influence of the Al-Si alloy’s chemical composition on its properties,” Archives of Foundry Engineering, vol. 18, pp. 203–211, 2018.
View at: Google Scholar
T. Szymczak, J. Szymszal, and G. Gumienny, “Evaluation of the effect of the Cr, Mo, V and W content in an Al-Si alloy used for pressure casting on its proof stress,” Archives of Foundry Engineering, vol. 18, pp. 105–111, 2018.
View at: Google Scholar
T. Szymczak, J. Szymszal, and G. Gumienny, “Evaluation of the eect of Cr, Mo, V and W on the selected properties of silumins,” Archives of Foundry Engineering, vol. 18, pp. 77–82, 2018.
View at: Google Scholar
M. Cai, J. D. Robson, and G. W. Lorimer, “Simulation and control of dispersoids and dispersoid-free zones during homogenizing an AlMgSi alloy,” Scripta Materialia, vol. 57, no. 7, pp. 603–606, 2007.
View at: Publisher Site | Google Scholar
K. Marthinsen, E. Nes, O. Daaland, and T. Furu, “The spatial distribution of nucleation sites and its effect on recrystallization kinetics in commercial aluminum alloys,” Metallurgical and Materials Transactions A, vol. 34, no. 12, pp. 2705–2715, 2003.
View at: Publisher Site | Google Scholar
C. D. Marioara, S. J. Andersen, T. N. Stene et al., “The effect of Cu on precipitation in Al-Mg-Si alloys,” Philosophical Magazine, vol. 87, no. 23, pp. 3385–3413, 2007.
View at: Publisher Site | Google Scholar
T. Marlaud, A. Deschamps, F. Bley, W. Lefebvre, and B. Baroux, “Influence of alloy composition and heat treatment on precipitate composition in Al-Zn-Mg-Cu alloys,” Acta Materialia, vol. 58, no. 1, pp. 248–260, 2010.
View at: Publisher Site | Google Scholar
B. A. Szajewski, J. C. Crone, and J. Knap, “Analytic model for the Orowan dislocation-precipitate bypass mechanism,” Materialia, vol. 11, p. 10067, 2020.
View at: Publisher Site | Google Scholar
S. Drouzy, Jacob, and M. Richard, “Interpretation of tensile results by means of quality index and probable yield strength,” AFS International Cast Metals Journal, vol. 5, pp. 43–50, 1980.
View at: Google Scholar
C. H. Cáceres, Microstructure Design and Heat Treatment Selection for Casting Alloys Using the Quality Index, ASM-IMS, Cincinnati, OH, USA, 1999.
C. H. Cáceres, T. Din, A. K. M. B. Rashid, and J. Campbell, “Effect of aging on quality index of an Al-Cu casting alloy,” Materials Science and Technology, vol. 15, no. 6, pp. 711–716, 1999.
View at: Publisher Site | Google Scholar
C. H. Ca’ceres and J. A. Taylor, “Enhanced ductility in Al–Si–Cu–Mg casting alloys with high Si content,” in Shape Casting: The John Campbell Symposium, M. Tiryakiouglu and P. Crepeau, Eds., pp. 245–254, Los Angeles, CA, USA, 2005.
View at: Google Scholar
I. El-Mahallawi, A. Shash, and A. Amer, “Nanoreinforced cast Al-Si alloys with Al2O3, TiO2 and ZrO2 nanoparticles,” Metals, vol. 5, no. 2, pp. 802–821, 2015.
View at: Publisher Site | Google Scholar
T. Dorin1, M. Ramajayam, and T. J. Langan, “Effects of Mg, Si, And Cu on the formation of the Al3Sc/Al3Zr dispersoids,” in Proceedings of the 16th International Aluminum Alloys Conference (ICAA16), Canadian Institute of Mining, Metallurgy & Petroleum, Montreal, Canada, June 2018, ISBN: 978-1-926872-41-4.
View at: Google Scholar
J. Hernandez-Sandoval, F. H. Samuel, and S. Valtierra, “Effect of additions of SiC and Al₂O₃ particulates on the microstructure and tensile properties of Al-Si-Cu-Mg cast alloys,” International Journal of Metalcasting, vol. 10, no. 3, pp. 253–263, 2016.
View at: Publisher Site | Google Scholar
G. H. Garza-Elizondo, A. M. Samuel, S. Valtierra, and F. H. Samuel, “Effect of transition metals on the tensile properties of 354 alloy: role of precipitation hardening,” International Journal of Metalcasting, vol. 11, no. 3, pp. 413–427, 2017.
View at: Publisher Site | Google Scholar
G. H. Garza-Eiizondo, S. A. Alkahtani, A. M. Samuel, and F. H. Samuel, “Role of Ni and Zr in preserving the strength of 354 aluminum alloy at high temperature,” in Light Metals 2014, pp. 305–314, Minerals, Metals and Materials Society, San Diego, CA, USA, 2014.
View at: Google Scholar
H. R. Ammar, A. M. Samuel, F. H. Samuel, E. Simielli, G. K. Sigworth, and J. C. Lin, “Influence of aging parameters on the tensile properties and quality index of Al-9 pct Si-1.8 pct Cu-0.5 pct Mg 354-type casting alloys,” Metallurgical and Materials Transactions A, vol. 43, no. 1, pp. 61–73, 2012.
View at: Publisher Site | Google Scholar
A. M. Samuel, F. H. Samuel, and H. W. Doty, “Influence of melt treatment and solidification parameters on the quality of 319.2 endchill aluminum castings,” in 4th International Conference on Molten Aluminium Processing, pp. 261–293, Orlando, FL, USA, November 1995.
View at: Google Scholar
J. Hernandez-Sandoval, A. M. Samuel, S. Valtierra, and F. H. Samuel, “Thermal analysis for detection of Zr-rich phases in Al-Si-Cu-Mg 354-type Alloys,” International Journal of Metalcasting, vol. 11, no. 3, pp. 428–439, 2017.
View at: Publisher Site | Google Scholar
J. Campbell and M. Tiryakioğlu, “Review of effect of P and Sr on modification and porosity development in Al-Si alloys,” Materials Science and Technology, vol. 26, no. 3, pp. 262–268, 2010.
View at: Publisher Site | Google Scholar
R. Fuoco, E. R. Correa, and A. V. O. Correa, “Effect of modification treatment on microporosity formation in 356 Al alloy,” AFS Transactions, vol. 103, pp. 379–387, 1995.
View at: Google Scholar
D. Argo and J. E. Gruzleski, “Porosity in modified aluminum alloy castings,” AFS Transactions, vol. 96, pp. 65–74, 1988.
View at: Google Scholar
H. W. Kerr, J. Cisse, and G. F. Bolling, “On equilibrium and non-equilibrium peritectic transformations,” Acta Metallurgica, vol. 22, pp. 677–686, 1974.
View at: Publisher Site | Google Scholar
J. Murray, A. Peruzzi, and J. P. Abriata, “The Al-Zr (aluminum-zirconium) system,” Journal of Phase Equilibria, vol. 13, no. 3, pp. 277–291, 1992.
View at: Publisher Site | Google Scholar
U. Ikhlaq, A. Hirose, R. Ahmad et al., “The role of Ar flow rates on synthesis of nanostructured zirconium nitride layer growth using plasma enhanced hot filament nitriding (PEHFN) technique,” The European Physical Journal Applied Physics, vol. 69, no. 1, p. 10801, 2015.
View at: Publisher Site | Google Scholar
K. E. Knipling, D. C. Dunand, and D. N. Seidman, “Precipitation evolution in Al-Zr and Al-Zr-Ti alloys during isothermal aging at 375–425°C,” Acta Materialia, vol. 56, no. 1, pp. 114–127, 2008.
View at: Publisher Site | Google Scholar
A. De Luca, S. Shua, and D. N. Seidman, “Effect of microadditions of Mn and Mo on dual L12- and α-precipitation in a dilute Al-Zr-Sc-Er-Si alloy,” Materials Characterization, vol. 169, pp. 114–127, 2020.
View at: Publisher Site | Google Scholar
H. Basoalto and M. Anderson, “An extension of mean-field coarsening theory to include particle coalescence using nearest-neighbour functions,” Acta Materialia, vol. 117, pp. 122–134, 2016.
View at: Google Scholar
J. Hernandez-Sandoval, G. H. Garza-Elizondo, A. M. Samuel, S. Valtiierra, and F. H. Samuel, “The ambient and high temperature deformation behavior of Al-Si-Cu-Mg alloy with minor Ti, Zr, Ni additions,” Materials & Design, vol. 58, pp. 89–101, 2014.
View at: Publisher Site | Google Scholar
A. Rendtel, B. Moessner, and K. A. Schwetz, “Hardness and hardness determination in silicon carbide materials, book editor,” Advances in Ceramic Armor, vol. 26, 2005.
View at: Publisher Site | Google Scholar
A. Daoud and W. Reif, “Influence of Al2O3 particulate on the aging response of A356 Al-based composites,” Journal of Materials Processing Technology, vol. 123, no. 2, pp. 313–318, 2002.
View at: Publisher Site | Google Scholar
A. M. Samuel, H. Liu, and F. H. Samuel, “Effect of melt, solidification and heat-treatment processing parameters on the properties of Al-Si-Mg/SiC(p) composites,” Journal of Materials Science, vol. 28, no. 24, pp. 6785–6798, 1993.
View at: Publisher Site | Google Scholar
O. Prach, O. Trudonoshyn, P. Randelzhofer, С. Körner, and K. Durst, “Effect of Zr, Cr and Sc on the Al-Mg-Si-Mn high-pressure die casting alloys,” Materials Science and Engineering: A, vol. 759, pp. 603–612, 2019.
View at: Publisher Site | Google Scholar
G. H. Garza-Elizondo, A. M. Samuel, S. Valtierra, and F. H. Samuel, “Effect of Ni, Mn, Sc, and Zr addition on the tensile properties of 354-type Alloys at ambient temperature,” International Journal of Metalcasting, vol. 11, no. 3, pp. 396–412, 2017.
View at: Publisher Site | Google Scholar
R. L. Zhang, Experiential Electronic Theory of Solid and Molecules Changchun, Jilin Science and Technology Press, Changchun, China, 1993.
J. D. Robson and P. B. Prangnell, “Dispersoid precipitation and process modelling in zirconium containing commercial aluminium alloys,” Acta Materialia, vol. 49, no. 4, pp. 599–613, 2001.
View at: Publisher Site | Google Scholar

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Copyright © 2021 J. Hernandez-Sandoval et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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