Improved Microstructure and Enhanced Tensile Properties of Hypoeutectic AlMg5Si2Mn Alloy Modified by Yttrium

Abstract

AlMg5Si2Mn alloys are widely used in the field of automotive castings. Since the morphology, size, and distribution of the primary Al dendrite and eutectic Mg₂Si have a decisive influence on the mechanical properties of the alloy, a comprehensive analysis of AlMg5Si2Mn alloys with varying Y contents was conducted using optical microscopy (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The influence of Y on the microstructural evolution and mechanical behavior of the cast hypoeutectic AlMg5Si2Mn alloy was studied. Experimental findings indicate that the addition of Y significantly refines and alters the morphology of both primary Al and eutectic Mg₂Si in AlMg5Si2Mn alloys. Specifically, in the alloy containing 0.45 wt.% Y, the primary Al undergoes a structural transformation from a coarse dendritic morphology to finer ellipsoidal grains, with a minimum secondary dendritic arm spacing (SDAS) of 18.6 ± 1.6 μm. Simultaneously, the eutectic Mg₂Si morphology transitions from a coarse lamellar structure to finer worm-like, coral-like, and fibrous forms, exhibiting a reduced average length and aspect ratio (AR) of 3.1 ± 0.4. Furthermore, the AlMg5Si2Mn alloy leads to significant improvements in mechanical performance, particularly in tensile strength. The measured average ultimate tensile strength, yield strength, and elongation are 243.3 MPa, 199.0 MPa, and 8.5%, respectively, representing increases of 19.16%, 24.6%, and 203.6% compared to the Y-free alloy. The fracture mode of the alloy fracture transitioned from brittle fracture in its unmodified condition to ductile fracture characteristics.

1. Introduction

AlMg5Si2Mn alloy (i.e., hypoeutectic Al-6Mg₂Si alloy) has more significant specific stiffness, specific strength, and better comprehensive mechanical properties [1,2,3,4,5]. Meanwhile, it has excellent formability, such as die casting, sand casting, extrusion, and forging and is widely used in the automotive industry and for other key parts, such as gear pump housings, automobile wheel hubs, shock absorber covers, cylinder heads, vehicle frames, engine blocks, and other automobile castings [1,2,3,4,5,6,7]. Primary Al dendrites and eutectic (Al + Mg₂Si) structures are formed through pseudo-eutectic reactions during solidification in AlMg5Si2Mn alloys. Generally, coarse primary Al dendrites and interconnecting plate-like, Chinese-characteristic-morphology eutectic Mg₂Si particles have been observed in hypoeutectic Al-Mg₂Si (AlMg5Si2Mn) or eutectic Al-Mg₂Si alloys [1,2,3,4,5,8,9,10,11]. However, the coarse-flake brittle eutectic Mg₂Si particles tend to act as crack initiation sites, leading to the fracture of the alloy matrix under applied working loads, thereby significantly deteriorating both the ductility and tensile strength of the material. As a result, the morphological features, size, and distribution of the primary Al and eutectic Mg₂Si are critical determinants that profoundly affect the mechanical performance of these alloys. Typically, rare-earth (RE) elements, such as Sc [1], Gd [2], and Er [4], along with grain refiners like AlTiB [12,13,14,15], AlTiC [14,15], and AlTiBC [16], are frequently incorporated into aluminum alloys to refine the microstructure of the primary Al grain. Additionally, modifiers like Gd [2], Bi [3], Er [4], Sr [13], Li [17], Ca [18], and Sb [18,19] have been utilized to alter the morphology and reduce the size of eutectic Mg₂Si. In general, rare-earth (RE) elements, such as Sc [1], Gd [2], and Er [4], and refiners, such as AlTiB [12,13,14,15], AlTiC [14,15], and AlTiBC [16], were added to Al alloys to refine the primary α-Al dendrites. Meanwhile, modifiers such as Sc [1], Gd [2], Bi [3], Er [4], Sr [13], Li [17], Ca [18], and Sb [18,19] were used to change the morphology and reduce the dimension of eutectic Mg₂Si. Several studies have demonstrated that the added Y can effectively modify the morphology of eutectic Si and significantly refine the primary Al dendrite in cast aluminum–silicon alloys [20,21,22]. For example, Li et al. [20] studied the incorporation of 0.3 wt.% Y into an Al-7.5Si-0.5Mg alloy, which resulted in a marked transformation of the eutectic silicon morphology from coarse, needle-like structures to finer fibrous forms. Likewise, Wan et al. [22] observed that the introduction of trace amounts of Y into cast AlSi7Mg0.3Fe alloys substantially altered the coarse, needle-like eutectic silicon particles into finer fibrous and spherical morphologies and simultaneously significantly refined the dendritic structure of primary α-Al. Mao et al. [23] proposed that adding Y to alloy A357 altered the morphology of eutectic silicon by introducing additional nucleation sites due to the presence of Y atoms during solidification. Additionally, the compositional supercooling effect, which results from the accumulation of Y atoms at the solidification interface of the liquid alloy during the crystallization process, further promoted the generation of an increased number of eutectic silicon nuclei. Li et al. [24] investigated the addition of Y to hypoeutectic Al-7%Si alloys and found that it effectively reduces SDAS and promotes the partial transformation of plate-like eutectic silicon into fibrous or granular morphologies, thereby significantly enhancing both the tensile strength and ductility of the Al-Si alloy.

To date, the effect of Y on hypoeutectic AlMg5Si2Mn alloy has been sparsely documented. This work aims to provide a comprehensive understanding of yttrium’s dual functionality in microstructural refinement and mechanical performance optimization of hypoeutectic Al-Mg₂Si alloy (AlMg5Si2Mn) systems, coupled with an in-depth investigation of the underlying strengthening mechanisms.

2. Experimental Methods and Procedures

A series of AlMg5Si2Mn-x wt.% Y alloys with Y additions were fabricated, where x represents Y concentrations of 0.0, 0.15, 0.30, 0.45, and 0.60. The detailed chemical composition of the alloy samples without Y is provided in

Table 1. The composition of AlMg5Si2Mn alloy (wt.%).

Element	Mg	Si	Mn	Fe	Ti	Cu	Al
Nominal value	5.5	2	0.6	<0.15	<0.09	<0.01	Bal.
Actual value	5.668	2.126	0.642	0.128	0.004	0.003	Bal.

. The alloy preparation employed commercially available raw materials sourced from Shen Yang Bo Yu Metal Co., Ltd. (Shenyang, China), including industrial pure Al (99.7 wt.%), industrial pure Mg (99.8 wt.%), and master alloys such as Al-20 wt.% Si (containing 19.923 wt.% Si), Al-10 wt.% Mn (with 9.952 wt.% Mn), and Al-10 wt.% Y (containing 9.926 wt.% Y). The constituent materials, consisting of Al ingots complemented by Al-10 wt.% Mn and Al-20 wt.% Si master alloys, were meticulously proportioned according to the designated alloy formulation before being transferred into a graphite crucible. This assembly was then positioned within a 5 KW electric resistance furnace, undergoing controlled heating to reach and sustain a temperature of 750 °C, ensuring complete melting of the metal. Upon reducing the temperature to approximately 700 °C, magnesium briquettes, pre-wrapped in aluminum foil to account for an estimated 10 wt.% combustion loss, were introduced into the melt and maintained at this temperature for approximately 10 min to ensure adequate reaction. The molten metal was subsequently heated to reach the target temperature of 750 °C, at which point the Al-10 wt.% Y master alloy was introduced into the melt. To ensure a homogeneous composition, the mixture was stirred and maintained at this temperature for 30 min. Then, dry C₂Cl₆, equivalent to 0.5% of the total melt mass, was added for degassing and slag removal. After maintaining the temperature for an additional 5 min, the molten alloy was carefully transferred into a pre-heated steel mold maintained at 250 °C, which had cylindrical sizes of φ25 mm in diameter and 120 mm in length, thereby producing the final cast aluminum alloy specimen.

The specimens designated for metallurgical characterization and mechanical property evaluation were precisely sectioned using wire electrical discharge machining (WEDM), with the sampling position maintained at about 20 mm from the bottom of the cast bar. The tensile test specimens were manufactured in strict accordance with the specifications prescribed by the Chinese national standard [25], with comprehensive dimensional specifications illustrated in Figure 1. Quantitative compositional analysis of the experimental alloys was performed utilizing an ICP-8100 analytical platform through inductively coupled plasma optical emission spectrometry (ICP-OES) methodology.

The metallographic specimens of the cast alloy were prepared using standardized grinding and polishing techniques. Subsequently, specimens were etched in 0.5 vol.% aqueous hydrofluoric acid solution for 5 s to reveal the metallographic microstructure, and some other specimens were subjected to chemical etching in a 20 wt.% NaOH solution for 40 min at ambient temperature to reveal the 3D structural characteristics of the eutectic Mg₂Si particle. Microstructural characterization of specimens was initially conducted using an optical microscopy system from PTI for image acquisition and preliminary analysis. Subsequently, a detailed examination of tensile fracture surfaces and 3D morphological features of the eutectic Mg₂Si particle was performed using scanning electron microscopy (SEM). Microstructural quantification was carried out using a PTI SOFT image analysis system coupled with optical microscopy manufactured by Shanghai CHAWEAR Co., Ltd. (Shanghai, China) to comprehensively evaluate the morphological features and dimensions of both the primary Al grain and eutectic Mg₂Si in cast AlMg5Si2Mn alloys. For quantitative analysis, five representative microstructural fields were selected from each specimen, with statistical parameters derived from a minimum of 10 independent measurements. These measurements enabled precise determination of key microstructural parameters, including SDAS of primary α-Al dendrites, as well as the characteristic length (L_Mg2Si) and AR of eutectic Mg₂Si particles. The phase identification in the experimental alloys was performed using a D8 Focus X-ray diffractometer operated at a scanning speed of 2°/min. The microstructural characterization and elemental distribution analysis were performed using a Zeiss SIGMA 500 field-emission scanning electron microscope (FE-SEM) integrated with energy-dispersive X-ray spectroscopy (EDS) and backscattered electron (BSE) imaging, enabling precise identification of Y distribution within the AlMg5Si2Mn alloy matrix.

The quantitative microstructural characterization was conducted on Y-modified aluminum alloys to evaluate the dimensional parameters and morphological characteristics of constituent phases. Advanced 2015 version of PTI SOFT image analysis software was employed to systematically acquire precise measurements of grain dimensions and aspect ratios across various Y concentrations. The characteristic lengths and morphological aspect ratios for each distinct phase were determined based on the methodology outlined below [2,3,5,18,19,25]:

$$SDAS={\displaystyle \frac{L}{N}}$$

(1)

Mean length size

$$={\displaystyle \frac{1}{m}}{{\displaystyle \sum _j^m\left(\frac{1}{n}{\displaystyle \sum _{i=1}^n{l}_{\mathrm{i}}}\right)}}_j$$

(2)

Mean aspect ratio

$$={\displaystyle \frac{1}{m}}{{\displaystyle \sum _j^m\left({\displaystyle \sum _{i=1}^n\frac{a_i}{b_i}}\right)}}_j$$

(3)

In this study, L represents the edge-to-edge spacing between secondary dendritic arms of the primary α-Al dendrite; N indicates the number of dendritic arms; l_i signifies the maximum size of a single eutectic Mg₂Si particle; and a_i and b_i represent the maximum and minimum sizes of individual eutectic Mg₂Si particles, respectively. The ratio of a_i to b_i is defined as the aspect ratio. n denotes the number of Mg₂Si particles measured within a single observation field, while m represents the total number of evaluation areas.

The tensile testing was conducted at ambient temperature (23 ± 2 °C) utilizing a WAW-300 computer-controlled electro-hydraulic servo testing system in displacement control mode with a speed of 1 mm/min. For each alloy composition, a minimum of three identical specimens were evaluated, and the mechanical property parameters were determined based on the arithmetic mean of the replicate measurements. This stringent testing protocol was established to ensure reproducibility of the experimental data.

3. Results

3.1. XRD Analysis

Figure 2 displays characteristic X-ray diffraction patterns obtained from AlMg5Si2Mn alloys with varied Y concentrations. The diffraction analysis reveals that all Y-containing alloy compositions consistently demonstrate distinct diffraction peaks corresponding solely to the α-Al matrix and Mg₂Si intermetallic, with no detectable evidence of Y-containing compounds.

3.2. Microstructural Characteristics of the Alloy

Based on the phase equilibrium described in the Al-Mg₂Si pseudo-binary phase diagram [26], the solidification sequence of AlMg5Si2Mn alloys follows the transformation pathway L → L + α-Al_p → α-Al_p + (α-Al + Mg₂Si)_E. The nucleation and growth of the primary α-Al initiate at approximately 620 ±1 °C through the liquid-to-solid transformation, i.e., L → L + α-Al_p. When the temperature decreases to 597 ± 0.5 °C and the Mg₂Si concentration in the residual liquid reaches approximately 13.9%, a eutectic transformation initiates, leading to the precipitation of a eutectic particle, specifically L → (α-Al + Mg₂Si)_E. During the solidification of the above alloy liquid, the subscripts P and E denote the precipitation of primary α-Al and eutectic components in the microstructure, respectively.

The low-magnification microstructures of AlMg5Si2Mn alloys with Y addition are presented in Figure 3a,c,e,g,i. It is evident from the figure that the microstructure of these alloys primarily comprises a primary α-Al dendrite and black eutectic Mg₂Si. The primary Al exhibits a more pronounced coarse dendritic morphology in the alloy samples without Y, as shown in Figure 3a,b. The quantitative microstructural analysis demonstrated that the SDAS of primary α-Al dendrites in Y-free alloy specimens is measured at 38.8 ± 1.8 μm, as seen in Figure 4. The addition of Y at 0.15 wt.% induced significant dendrite refinement, decreasing the SDAS to 32.4 ± 1.6 μm (Figure 3c,d). Further increases in Y content beyond this level resulted in microstructural modifications, as presented in Figure 3e,g,i. Notably, the average SDAS progressively decreases from 24.3 ± 1.5 μm to 18.6 ± 1.6 μm with the increase in Y from 0.3 wt.% to 0.45 wt.% (Figure 3e–h). However, at the highest Y concentration of 0.6 wt.%, a slight increase in SDAS to 22.5 ± 1.4 μm is observed (Figure 3i,j), indicating the presence of an optimal Y content for achieving maximal microstructural refinement.

The high-magnification microstructures of AlMg5Si2Mn alloy containing varying amounts of Y are demonstrated in Figure 3b,d,f,h,j. The eutectic Mg₂Si displays a coarse plate-like morphology with average L_Mg2Si and AR of 14.6 ± 1.4 µm and 6.6 ± 0.5 in the Y-free alloy, respectively (Figure 3b). The microstructural analysis revealed that the incorporation of 0.15 wt.% Y led to a refinement of the eutectic Mg₂Si particles, although localized regions retained their lamellar morphology (Figure 3d). Upon increasing the Y content to 0.30–0.45 wt.%, the eutectic Mg₂Si morphology underwent a significant transformation from flake-like and Chinese-script-like structures to refined platelet, fibrous, and vermicular morphologies, as illustrated in Figure 3f,h. The quantitative evaluation indicated a progressive reduction in both the characteristic L_Mg2Si and AR of the eutectic Mg₂Si particles, decreasing from 6.8 ± 0.6 μm and 3.9 ± 0.5 to 3.7 ± 0.5 μm and 3.1 ± 0.4, respectively. Specifically, the alloy containing 0.45 wt.% Y exhibits the most significant microstructural refinement, indicating that an optimal Y addition can effectively modify the morphology of eutectic Mg₂Si in hypoeutectic AlMg5Si2Mn alloys. However, the eutectic Mg₂Si particles begin to coarsen when their concentration exceeds 0.45 wt.%. Notably, for alloys with an Y addition of 0.6 wt.% (Figure 3j), the average length of Mg₂Si (L_Mg2Si) is 6.5 ± 0.7 µm and AR is 3.5 ± 0.3. Figure 4 illustrates the average L_Mg2Si and AR values of the eutectic Mg₂Si particles across different Y contents.

The eutectic Mg₂Si morphology plays a crucial role in determining the mechanical properties of AlMg5Si2Mn alloys, thereby requiring a comprehensive analysis. Figure 5 presents a three-dimensional visualization of the eutectic Mg₂Si particles.

In Y-free alloys, the eutectic Mg₂Si crystals display a coarse, plate-like morphology. Additionally, parallel stepped structures were observed on the eutectic Mg₂Si plates (see Figure 5a), which are characteristic of a typical eutectic structure. Three-dimensional microstructural analysis indicates that the eutectic Mg₂Si in the AlMg5Si2Mn alloy containing 0.15 wt.% Y predominantly exhibits a plate-like morphology (as depicted in Figure 5b). Upon the addition of 0.3 wt.% Y, a morphological transition from coral-like to fibrous structures is observed, accompanied by a significant reduction in particle dimensions (Figure 5c). Notably, in the alloy containing 0.45 wt.% Y, the three-dimensional architecture of the eutectic Mg₂Si undergoes substantial modification, with the plate-like morphology transforming into an array of parallel fibrous structures and vermicular particles (Figure 5d). The microstructural evolution reveals that Y addition effectively refines the eutectic Mg₂Si particle in hypoeutectic AlMg5Si2Mn alloys, demonstrating modification behavior similar to that induced by Sc, Gd, and Er additions in hypoeutectic Al-Mg₂Si alloy systems [1,2,4]. When the Y concentration reaches 0.6 wt.%, the microstructural analysis indicates a marked increase in the dimensional parameters of eutectic Mg₂Si, along with significant morphological coarsening. Specifically, its morphology transforms into a coarser worm-like and short rod morphology, as depicted in Figure 5e.

3.3. Tensile Properties of the Alloys

Figure 6 and Figure 7, respectively, illustrate the tensile behavior and corresponding mechanical properties of Y-modified AlMg5Si2Mn alloys at room temperature. The baseline alloy without Y addition exhibits mechanical properties with an ultimate tensile strength (UTS) of 195.25 MPa, yield strength (YS) of 167 MPa, and elongation (EL) of 2.8%. The mechanical properties follow a distinct trend, initially increasing with Y content until reaching an optimal value, after which they decrease. At the optimal Y concentration of 0.45 wt.%, the alloy achieves its peak mechanical performance, with UTS, YS, and EL values of 243.3 MPa, 199 MPa, and 8.5%, respectively. These values correspond to significant improvements of 19.16%, 24.6%, and 203.6% in YS, UTS, and EL compared to the alloy without Y.

The fractographic analysis is performed on both transverse and longitudinal specimen orientations after tensile testing, revealing microstructural features as shown in Figure 8 and Figure 9. This thorough investigation offers an in-depth analysis of the fracture mechanisms and mechanical property characteristics of the tested alloys. Specifically, the transverse fracture surface of the Y-free alloy in Figure 8a is covered with a large range of flat, irregular, smooth disintegrated surfaces, and river-like patterned faults appear on these disintegrated surfaces. The detachment of eutectic Mg₂Si flakes from the Al matrix on these flat regions leads to the formation of smooth surfaces. This phenomenon is attributed to the fracture of the coarse, slab-like, brittle Mg₂Si particles. The Y-free specimens exhibit predominantly brittle fracture characteristics, as indicated by the presence of cleavage facets and a lack of significant plastic deformation. This brittle fracture behavior is directly linked to the inferior mechanical performance observed in Figure 6, which is characterized by reduced strength and limited elongation capacity in the unmodified alloy.

The alloy containing 0.15 wt.% Y induces significant microstructural modifications, manifested by the gradual disappearance of planar fracture features and the emergence of dimpled regions; this indicates a shift towards a mixed fracture mode that combines quasi-cleavage and micro-dimple characteristics, as illustrated in Figure 8b. Progressive Y addition beyond this concentration promotes extensive formation of dimples, which correlates with enhanced mechanical performance, particularly in terms of ductility improvement (Figure 6). At the optimal Y concentration of 0.45 wt.%, the fracture surface is entirely characterized by fine, equiaxed dimples with significant depth, indicative of a fully developed ductile fracture (Figure 8d). These fractographic observations offer direct evidence supporting the superior tensile properties and maximum plastic deformation capacity achieved in the 0.45 wt.% Y-modified AlMg5Si2Mn alloy, as shown in Figure 6.

However, at the elevated Y concentration of 0.60 wt.%, the fracture morphology exhibits significant alterations, marked by the re-emergence of localized planar fracture features and the onset of microcrack formation, along with a marked reduction in dimple density (Figure 8e). The degradation in microstructure is closely linked to the decline in mechanical properties, evidenced by decreased tensile strength and reduced ductility. The primary mechanism responsible for the property degradation is attributed to the coarsening of eutectic Mg₂Si, which compromises the alloy’s ability to undergo plastic deformation.

For comparative analysis, Figure 9 illustrates the lateral fracture characteristics of both unmodified and 0.45 wt.% Y-containing AlMg5Si2Mn alloys. In the Y-free alloy (Figure 9a), crack initiation predominantly occurs at the interlamellar boundaries of coarse eutectic Mg₂Si particles, followed by crack propagation along the interfaces between primary α-Al dendrites. Furthermore, additional crack nucleation and extension are frequently observed along these dendritic interfaces, which aligns with the typical brittle fracture behavior noted in the longitudinal section. Figure 9b provides an enlarged view highlighting specific morphological features characteristic of the unmodified alloy’s fracture surface. The areas delineated by a dashed line and indicated by an arrow represent the regions of crack initiation and propagation. Notably, the figure also reveals that the alloy exhibits no evidence of plastic deformation. The crack source primarily originates from the interface between the eutectic Mg₂Si particles region and the primary α-Al dendrites, specifically from the coarse, lamellar eutectic Mg₂Si particles. The presence of such coarse, lamellar Mg₂Si particles not only accelerates premature fracture of the sample but also facilitates continuous crack propagation around the eutectic particles. Although the primary α-Al exhibits favorable plasticity characteristics, fracture occurs when crack propagation in the eutectic Mg₂Si particles is dominated due to stress concentration, thereby preventing the manifestation of its plastic behavior. This is similar to the case where the coarse, slab-like eutectic Mg₂Si in Al-Mg₂Si alloys acts as a significant source of stress concentration [2,4,27]. Consequently, the presence of such coarse, lamellar eutectic Mg₂Si critically influences the fracture mechanism and is primarily responsible for the reduced strength and poor ductility observed in the Y-free alloy.

Figure 9c illustrates the lateral fracture of the AlMg5Si2Mn alloy that has been modified with 0.45 wt.% Y. The eutectic Mg₂Si displays a fine fibrous morphology (Figure 5d), and the lateral fracture surfaces exhibit predominantly point cracking, suggesting that the alloy experienced substantial plastic deformation prior to failure. Figure 9d presents a localized magnified view of the Y-modified AlMg5Si2Mn alloy, where the region enclosed by the dashed line indicates an area that has undergone significant plastic deformation. Consequently, in Y-modified alloys, fracture propagation is not confined to the eutectic Mg₂Si particles but also involves both the primary and eutectic α-Al. The fracture behavior is evidenced by the presence of plastic deformation zones, where crack initiation sites are clearly identified within both the primary and eutectic α-Al. These observations suggest a more intricate fracture mechanism in comparison to the unmodified alloy. In contrast, the fine fiber-like eutectic Mg₂Si demonstrates considerable resistance to fracture due to its inherent toughness, effectively inhibiting crack propagation. This finding aligns with prior research indicating that the microstructural modification of eutectic particles in aluminum-based alloys markedly enhances their mechanical properties. Specifically, in Al-Mg₂Si systems, the incorporation of various rare-earth elements such as Sc [1], Gd [2], and Er [4], along with Bi [3,5,28], effectively alters eutectic Mg₂Si morphology. Similarly, La, Ba, Nd, Sc, and Zr induce transformations in the characteristics of eutectic silicon in Al-Si alloys. These modifications consistently result in significant improvements in the overall mechanical performance of these alloy systems [26,29,30,31].

4. Discussion

4.1. Refinement Mechanism of Alloys

The microstructural analysis as presented in Figure 10a,b confirms the presence of finely dispersed bright particles, which are predominantly located in the central regions of primary α-Al dendrites(shown by the yellow directional arrows) as well as at the boundaries of the eutectic Mg₂Si particles (shown by the red arrows), respectively. According to the EDS analysis results (Figure 10c–h), the white particles are identified as Y-rich compounds. Based on the surface-scanning energy spectrum analysis presented in Figure 11b,c, the results can be cross-referenced with the relevant literature, confirming that the compound is an Al₃Y particle [32]. According to the Al-Y phase diagram, the alloying solution undergoes a eutectic reaction, L → α-Al + Al₃Y, at 637 °C ± 2 °C [32]. The eutectic α-Al, which exhibits lattice parameters identical to those of the primary α-Al grain, serves as a substrate that promotes heterogeneous nucleation and growth of the primary α-Al.

Bramfitt’s theory indicates that a smaller two-dimensional (2D) lattice mismatch leads to a superior lattice match between the substrate and the nucleation particle, thereby facilitating the nucleation process. The 2D lattice mismatch can be quantified using the following equation [33]:

$${\delta }_{{\left(\mathit{hkl}\right)}_n}^{{\left(hkl\right)}_{\mathrm{s}}}={\displaystyle \frac{1}{3}}{\displaystyle \sum _{i=1}^3\frac{\left|\left(d_{{\left[uvw\right]}_s^i}\mathit{cos}\theta \right)-d_{{\left[uvw\right]}_n^i}\right|}{d_{{\left[uvw\right]}_n^i}}}\times 100\%$$

(4)

In Equation (4), (hkl)_s represents the low-index crystallographic plane of the substrate material, while [uvw]_s denotes the corresponding low-index crystallographic direction within the (hkl)_s plane. Similarly, (hkl)_n designates the low-index plane of the nucleating particle, with [uvw]_n representing the associated low-index direction within the (hkl)_n plane. The parameter d[uvw]_s quantifies the interatomic spacing along the [uvw]_s direction, and the notation d[uvw]_n represents the interatomic spacing along the [uvw]_n direction. The angular relationship between these crystallographic directions is defined by θ, which specifies the angle formed between [uvw]_s and [uvw]_n.

The Al₃Y intermetallic compound is a hexagonal close-packed (HCP) structure with lattice parameters of a = 0.6310 nm and c = 0.4590 nm [34]. This crystallographic arrangement is notably distinct from the face-centered cubic (FCC) structure of primary α-Al, which has a lattice constant of a = 0.4050 nm. According to previous studies [33,34], the 2D lattice mismatch between Al₃Y and primary Al has been precisely determined to be 10.19%, a value notably lower than the critical threshold of 15% that is essential for achieving effective heterogeneous nucleation. The relatively small lattice mismatch facilitates the role of Al₃Y as an efficient nucleation catalyst in AlMg5Si2Mn alloys, thereby promoting the formation of Al₃Y precipitates that act as favorable heterogeneous nucleation sites. Consequently, this mechanism markedly enhances the nucleation of primary α-Al, leading to the substantial microstructural refinement.

The backscattered electron (BSE) images in Figure 10a,b and Figure 12 show that the Y element is not only present in the boundary region of the eutectic Mg₂Si particles but also appears at the junction of the primary α-Al and the eutectic Mg₂Si particles, as confirmed by the bright white particles indicated by the red arrows in Figure 10. With the equilibrium partition coefficient of Y during solidification of the alloy being less than 1, Y atoms are expelled from the solid phase, instead migrating to and accumulating at the dynamic solid–liquid interface, leading to enrichment at the advancing front of the growing eutectic Mg₂Si particles. This phenomenon is similar to the effects observed in Sc-, Gd-, Bi-, and Er-containing Al-Mg₂Si alloys [1,2,3,4], as well as Y-, La-, and Eu-containing Al-Si alloys [23,26,35]. The energy line scan in Figure 13 reveals a misalignment between the peak Y concentration in the Y-rich region and the concentrations of Mg and Si, indicating the absence of Y element from the eutectic Mg₂Si particles. However, there is an observed increase in Y content surrounding the Mg₂Si particles; it suggests that the white Y-rich particle is located proximate to the eutectic Mg₂Si. The relatively large atomic radius of Y significantly limits its solid solubility in the aluminum matrix, thereby promoting the formation of an Y-enriched boundary layer at the advancing solid–liquid interface during the intricate process of alloy solidification. This phenomenon triggers compositional undercooling at the solid–liquid interface, thereby effectively suppressing the growth mechanism of eutectic Mg₂Si particles. Moreover, the magnesium and silicon atoms diffuse through the Y-rich liquid boundary layer, resulting in their accumulation at the solidification front. Consequently, the process promotes the growth of the Mg₂Si particles. Simultaneously, based on the EDS spectra acquired from the spot scan depicted in Figure 11, it is evident that the Y-rich layer can form a ternary (Al, Mg, or Y) particle. According to the existing literature [36], the Y-rich layer specifically forms as the Al₄MgY compound. The Y-rich interfacial layer partially serves as a diffusional barrier for the magnesium and the silicon atoms, thereby suppressing the growth of the Mg₂Si particles and facilitating microstructural refinement. The observation aligns with prior studies on the modification of eutectic silicon through the incorporation of rare-earth elements, including Y, Ce, and Yb, in Al-Si alloy systems [37,38,39]. Conversely, it is postulated that a portion of the Y atoms are concentrated at the growth interface of the eutectic Mg₂Si, undergo dissolution within the Mg₂Si particles, and the dissolved Y atoms in the Mg₂Si particles induce severe lattice distortions due to their significantly larger atomic radius compared to Mg and Si atoms, resulting in a substantial increase in internal defects within the crystals. Consequently, the presence of numerous crystal defects leads to roughening of the faceted growth interfaces of the eutectic Mg₂Si particles, either partially or entirely, thereby promoting isotropic growth and further refinement of their microstructure.

According to the aforementioned analysis and discussion, it is evident that Al and Y form Al₃Y compounds in AlMg5Si2Mn alloys. As clearly observed in the transmission electron microscopy (TEM) image presented in Figure 14, these Al₃Y compounds are distributed within the primary α-Al dendrite and around the eutectic Mg₂Si particles. Consequently, it can be concluded that Al₃Y compounds serve as effective heterogeneous nucleation cores.

4.2. Mechanical Property Enhancement Mechanisms for Alloys

The experimental results demonstrate that adding Y into the as-cast AlMg5Si2Mn alloy enhances both the strength and plasticity. The enhancement in tensile strength of Y-containing AlMg5Si2Mn alloys can be ascribed to three predominant factors: the refinement of the primary α-Al grains, the modification of the eutectic Mg₂Si particles, and strengthening by second-phase particles. According to the Hall–Petch formula, which elucidates the relationship between the tensile strength and grain size in the metallic materials:

$${\sigma }_1={\sigma }_0+{\mathit{kd}}^{-0.5}$$

(5)

In Equation (5), σ₁ represents the tensile strength in MPa, k is a constant, usually micron-sized grains k > 0, the parameter d signifies the mean dimension of the primary α-Al dendrites, measured in micrometers (μm), whereas σ₀ signifies the intrinsic lattice friction stress. As indicated by Equation (5), a decrease in the dimension (d) of primary α-Al dendrites results in an increase in the tensile strength (σ₁). Adding Y within the concentration range of 0.15–0.45 wt.% to the AlMg5Si2Mn alloys results in a remarkable refinement of the microstructure, specifically in the primary α-Al dendrites, thereby markedly elevating the tensile properties, as vividly depicted in Figure 6. This remarkable enhancement can be ascribed to the synergistic interplay between dendrite refinement and the modified distribution of second-phase particles. A notable observation is that the introduction of 0.6 wt.% Y results in a slight increase in the number of primary α-Al dendrites, along with a decrease in the average tensile strength to 235.5 MPa. The parameter σ₀ quantitatively represents the grain boundary resistance to plastic deformation. The microstructural refinement increases the density of grain boundaries that impede dislocation motion, thereby elevating the energy barrier for the crack propagation and enhancing tensile properties [40,41]. The alloy modified with 0.45 wt.% Y exhibits optimal microstructural characteristics, featuring the finest α-Al dendrites and consequently achieving the highest tensile strength. Moreover, the strengthening mechanism is further augmented by the second-phase particles’ reinforcement from the eutectic Mg₂Si and Al₃Y particles. These second-phase particles’ constituents effectively restrict dislocation movement during plastic deformation. Consequently, these particles can circumvent the moving dislocations and form a dislocation ring around them, thereby increasing the density of dislocations. As a result, the critical resolved shear stress (CRSS) necessary for dislocation glide is significantly increased, enhancing the material’s resistance to plastic deformation through improved dislocation pinning mechanisms. Furthermore, the enhancement of the second-phase particle dispersion can be quantitatively assessed utilizing the following Orowan equation [42]:

$${\sigma }_{\mathit{Or}}={\displaystyle \frac{Gb}{2\pi \left(1-\nu \right)}}.{\displaystyle \frac{1}{\lambda }}.\ln \left({\displaystyle \frac{D}{r_0}}\right)$$

(6)

In the formula, ${\sigma }_{\mathit{Or}}$ is the yield strength, b signifies the Boltzmann vector, G denotes the shear modulus, ν is Poisson’s ratio, λ and D represent the effective inter-particle distance and the mean particle diameter of the second phases, respectively, while r₀ corresponds to the core radius of the dislocation. Equation (6) establishes an inverse correlation between the yield strength and the characteristic dimension of the second-phase particles. This relationship reveals that the reduction in particle size, coupled with an increase in particle density, profoundly amplifies the yield strength of the alloy, thereby resulting in a remarkable enhancement in its mechanical properties.

In addition, according to the critical stress equation for the occurrence of fracture within eutectic particle particles [43]:

$${\sigma }_s=\sqrt{{\displaystyle \frac{\pi E{\gamma }_s}{D\left(1-{\nu }^2\right)}}}$$

(7)

The critical stress in Equation (7) is denoted as σ_s and measured in MPa, while E represents the Young’s modulus. D refers to the size of eutectic Mg₂Si particles, measured in μm, and ν denotes the Poisson’s ratio. γ_s denotes the effective surface energy of the Al matrix. The critical stress σ_s for the separation of the eutectic Mg₂Si from the aluminum matrix demonstrates an increase with D, as indicated by Equation (7), which exhibits a direct correlation with the fineness of the eutectic Mg₂Si particles. Consequently, a higher level of stress is required for refined microstructure to induce fracture, resulting in an enhancement of alloy tensile strength. In this experiment, the eutectic microstructure of the alloy containing 0.45 wt.% Y demonstrates the finest characteristics, thereby enabling the alloy to achieve the highest critical fracture stress and consequently exhibit maximum tensile strength.

Beyond the notable improvement in tensile strength, Y-modified AlMg5Si2Mn alloys demonstrate exceptional ductility enhancement. The alloy containing 0.45 wt.% Y shows a remarkable 203.6% increase in elongation relative to its unmodified counterparts, as evidenced in Figure 6. This substantial plasticity improvement is primarily attributed to the microstructural modification of eutectic Mg₂Si particles, which are inherently more prone to crack nucleation and propagation compared to the more ductile α-Al matrix due to their inherent brittleness. The fracture behavior of these alloys is critically influenced by the dimensional and morphological characteristics of eutectic Mg₂Si particles, which play a pivotal role in dictating both the initiation and propagation mechanisms of cracks.

The critical stress σ_c for crack extension within the material and the corresponding crack length a can be determined according to Griffith fracture theory [44]:

$${\sigma }_c=\sqrt{{\displaystyle \frac{2E\gamma }{\pi a}}}$$

(8)

In the equation, E and γ represent Young’s modulus and the fracture surface energy, respectively. Coarse eutectic Mg₂Si particles exhibit a greater propensity for initiating the formation of internal defects compared to fine eutectic Mg₂Si particles, which results in a substantial reduction in the intrinsic fracture stress of the AlMg5Si2Mn alloy. The matrix alloy containing coarse Mg₂Si particles shows a higher density and larger size of microcracks during fracture, in contrast to the Y-containing alloy with fine eutectic Mg₂Si particles. Additionally, the elongated lamellar eutectic Mg₂Si is uniformly distributed along the α-Al dendrite boundaries, leading to localized stress concentration around the Mg₂Si particles. This phenomenon serves as a precursor for crack formation and ultimately initiates microcracks. The microcracks subsequently propagate locally along the boundaries of primary α-Al, leading to the emergence of extensively fragmented surfaces and fractures within the matrix alloy (Figure 7). The outcome is a notable decrease in its ductility (Figure 6).

5. Conclusions

The influence of adding Y on microstructural evolution and mechanical behavior in hypoeutectic AlMg5Si2Mn alloys is systematically studied, and the following significant findings are obtained:

(1)

The incorporation of Y elements in cast AlMg5Si2Mn alloys results in microstructural refinement, as indicated by a substantial reduction in the SDAS of primary α-Al dendrites, as well as significant modification of both the morphology and dimensions of the eutectic Mg₂Si.

(2)

In Y-added AlMg5Si2Mn alloys, Al₃Y particles act as effective heterogeneous nucleation sites and exhibit significant multiple strengthening effects.

(3)

The comprehensive mechanical properties of AlMg5Si2Mn alloys are significantly enhanced by the addition of Y.