Microstructure Evolution, Crystallographic Orientation Regulation and Strength-Ductility Synergy Mechanism of Al-Si-Mg Alloy Synergistically Modified by Rare Earth Y and In Situ ZrB₂ Nanoparticles

Abstract

To address the demand for lightweight, high-performance Al-Si-Mg alloys in aerospace and automotive industries, this work proposes a novel synergistic strengthening strategy by combining rare-earth Y microalloying and in situ synthesized ZrB₂ nanoparticles to construct a hybrid reinforcement architecture. The effects of Y-ZrB₂ additions on the microstructure, crystallographic orientation evolution, and mechanical properties of Al-Si-Mg alloys were systematically investigated via XRD, SEM, EBSD, and tensile/hardness tests. Results show that compared with the base alloy and single-modified alloys, the co-addition of Y and ZrB₂ simultaneously enhances mechanical properties and optimizes grain structure. The optimal comprehensive performance is achieved at 0.3 wt.% Y + 2 wt.% ZrB₂ after T6 heat treatment, with ultimate tensile strength of 332.87 MPa, yield strength of 271.35 MPa, elongation of 16.24%, and Vickers hardness of 153.9 HV. Phase analysis and SEM-EDS confirm a synergistic coupling relationship between Y-rich phases and ZrB₂ nanoparticles. EBSD characterization reveals that Y-ZrB₂ modification has negligible effect on the morphology and crystallographic orientation stability of primary α-Al grains, but effectively regulates the lattice rotation, texture redistribution, and growth behavior of eutectic Si. At the optimal composition, the fraction of high-angle grain boundaries (HAGBs) reaches a maximum of 34.3%. Furthermore, the synergistic effect significantly increases the geometrically necessary dislocation (GND) density and reduces the Schmid factor of the dominant {111}⟨110⟩ slip system, thus enhancing dislocation strengthening and plastic deformation resistance. This work clarifies the intrinsic strength-ductility synergy mechanism of Y-ZrB₂ co-modified Al-Si-Mg alloys, paving a new pathway for the development of advanced lightweight aluminum alloys.

1. Introduction

The rapid expansion of high-end strategic manufacturing sectors-including aerospace, automotive, and rail transit-has led to a growing requirement for structural materials characterized by low weight and high performance in modern industrial engineering [1,2]. Aluminum alloys are widely considered among the most attractive lightweight structural candidates because of their low density and favorable specific strength [2,3,4,5,6]. Within the various aluminum alloy families, Al-Si-Mg alloys offer excellent overall service characteristics and have thus been extensively employed in industrial structural components [7,8,9,10]. Nevertheless, conventional commercially available Al-Si-Mg alloys fail to achieve a satisfactory strength-toughness balance under extreme service conditions; their limited mechanical properties remain a major barrier to large-scale industrial application [8].

Two predominant modification strategies have been explored to address these performance challenges: particle reinforcement and rare-earth microalloying. Regarding particle reinforcement, the incorporation of ceramic nanoparticles into the aluminum matrix acts as obstacles impeding dislocation glide and restricting grain boundary mobility, thereby contributing to marked enhancements in both strength and structural rigidity [11,12,13,14,15]. As for rare-earth microalloying, the addition of trace amounts of rare-earth elements can tailor microstructural features, promote the formation of stable strengthening precipitates, and consequently improve overall mechanical behavior. Previous studies have demonstrated that combined modification using rare-earth elements together with metallic alloying elements facilitates heterogeneous nucleation, refines grain size, and suppresses casting defects, thus enabling a simultaneous improvement in strength and ductility of aluminum alloys [3,5,6,16,17,18,19,20,21,22,23,24,25,26,27,28].

In recent years, the microalloying of Al-Si-Mg alloys with Y has drawn increasing attention. For instance, as reported by Xu et al., Y addition to an Al-0.6Mg-0.5Si alloy results in progressive grain refinement. The most pronounced effect, corresponding to a reduction in average grain size to 168 μm and an increase in tensile strength to 154 MPa, is observed upon the addition of 0.4 wt.% Y, representing an enhancement of 12.4% [29]. Bi et al. further showed that combining T6 heat treatment with 0.3 wt.% Y raises the tensile strength to 206.2 MPa-a 36.6% improvement relative to the as-cast condition-although the elongation decreases substantially after heat treatment [30]. Moreover, rare-earth oxides can also effectively regulate the alloy microstructure; the work of Chen et al. revealed how La₂O₃ refines grains and improves corrosion resistance [31]. Nonetheless, the strengthening effect attainable through single rare-earth microalloying alone remains modest and insufficient for increasingly stringent engineering requirements.

With respect to particle reinforcement, in situ synthesized nanoparticles have become a research focus in aluminum matrix composites owing to their high melting point, high hardness, and favorable lattice matching with the aluminum matrix [11,12,13,14,15]. However, single-method particle reinforcement suffers from inherent drawbacks, including poor interfacial wettability and pronounced clustering of nanoscale ceramic phases. When the ZrB₂ content exceeds 3 wt.%, such agglomeration severely compromises the ductility and long-term service reliability of the alloy [11,12]. Motamedi Yegane and Alizadeh reported that introducing 10 wt.% ZrB₂ into an Al5083 matrix via in situ stir casting increases the ultimate tensile strength by 18.2%, yet the strain (elongation) is reduced by 19.5% compared with the unreinforced alloy [32]. This trade-off between strength and ductility highlights the plasticity deficiency commonly associated with single-phase nanoparticle reinforcement.

To circumvent these technical bottlenecks, the synergistic introduction of rare-earth elements together with ZrB₂ particles into the aluminum matrix to construct a hybrid reinforcement architecture has emerged as an effective strategy to overcome the limitations of conventional modification approaches. Kai et al. demonstrated that the combined introduction of in situ ZrB₂ nanoparticles and Sc into a 7N01 Al alloy mitigates particle agglomeration, modifies Fe-rich precipitates, and refines the grain structure. These microstructural changes translate into enhanced mechanical properties, with the yield strength, ultimate tensile strength, and elongation reaching 447 MPa, 481 MPa, and 9.6%, respectively-representing gains of 16.1%, 11.3%, and 29.7% compared to the unreinforced matrix [33]. Ahmad et al. microalloyed an Al-Si-Mg alloy with Sc and Zr and observed that the combined addition refines the grain structure, alters precipitation behavior, and thereby contributes to a more favorable strength-ductility combination [34]. Ye et al. employed a high-throughput computational approach to screen microalloyed Al-Mg-Si alloys and identified Yb, Sc, and Ce as efficient microalloying elements; the enhanced overall performance of the material arises primarily from the formation of rare-earth-based strengthening phases within the alloy [35]. Collectively, these investigations demonstrate that the synergistic effect between rare-earth elements and ZrB₂ particles can effectively alleviate particle agglomeration, improve interfacial bonding, refine the grain structure, and enhance precipitation strengthening, thereby significantly optimizing the comprehensive mechanical properties of the alloys.

Based on the above considerations, this study synergistically combines in situ synthesized ZrB₂ nanoparticles with rare-earth Y microalloying to create a novel hybrid strengthening system in an Al-Si-Mg alloy. The influence of Y on the microstructural characteristics and crystallographic orientation distribution of the ZrB₂-modified alloy is systematically examined, and the fundamental mechanisms governing the observed synergy between reinforcement and toughening are clarified. The work aims to uncover the microstructural evolution behavior of Al-Si-Mg alloys and to provide a theoretical basis for the development of new high-performance lightweight aluminum alloys suitable for demanding industrial environments.

2. Experimental Procedure

2.1. Material Preparation and Heat Treatment Process

TheAl-Si-Mg-Y-ZrB₂ alloy used in this study was prepared using Al-3B, Al-10Zr, and Al-5Y master alloys. Based on the as-prepared Al-Si-Mg-ZrB₂ alloy, an appropriate amount of Al-5Y master alloy was added to adjust the Y content to the optimal value of 0.3 wt.%, with the aim of further optimizing the microstructure and mechanical properties of the alloy [36]. The chemical composition of the alloy was determined by inductively coupled plasma atomic emission spectrometry (ICP-AES, Agilent Technologies Inc., Santa Clara, CA, USA), and the results of the elemental analysis are given in

Table 1. Nominal addition amounts and the corresponding experimentally determined actual contents of the Al-Si-Mg-Y-ZrB₂ alloy (wt.%).

Samples	Si	Mg	Zr	B	Y	Other Impurity	Al
Al-Si-Mg	7.06	0.44	0	0	0	<0.01	Bal.
0.3 wt.% Y + 1 wt.% ZrB2	7.02	0.41	0.8	0.19	0.31	<0.01	Bal.
0.3 wt.% Y + 2 wt.% ZrB2	6.98	0.41	1.62	0.48	0.31	<0.01	Bal.
0.3 wt.% Y + 3 wt.% ZrB2	6.88	0.39	2.43	0.72	0.30	<0.01	Bal.

.

The specific preparation procedure was as follows. A graphite crucible was preheated at 423 K for 20 min. A predetermined amount of Al-Si-Mg alloy was weighed, placed into the preheated graphite crucible, and melted at 1123 K. After the alloy was completely melted, Al-3B and Al-10Zr master alloys were added. To ensure sufficient reaction and in situ formation of ZrB₂ nanoparticles, the melt was maintained at 1123 K for 30 min with simultaneous refining. Subsequently, the melt temperature was raised to 1193 K, and Al-5Y master alloy was added to tailor the microstructure, followed by another holding period of 30 min. After removing the crucible, the high-temperature melt was mechanically stirred to ensure uniform distribution of yttrium in the aluminum liquid. When the melt temperature decreased to 993 K, the melt was cast into a mold. Finally, the as-cast alloy was subjected to T6 heat treatment, which consisted of solution treatment at 808 K for 6 h followed by aging at 443 K for 7 h.

2.2. Material Characterization and Testing Methods

To prepare specimens for EBSD observation, the Al-Si-Mg-Y-ZrB₂ alloy samples were mechanically ground using SiC abrasive papers with grit sizes progressively increased from 180 to 2000 mesh. This was followed by electrolytic polishing carried out at 253 K under an applied voltage of 12 V for 30 s in a solution of perchloric acid and ethanol (HClO₄:C₂H₅OH = 1:9 by volume), which yielded a surface free from deformation-induced strain.

Mechanical property testing: According to GB/T 40123-2021 [37], the specimen was machined into the shape shown in Figure 1. Tensile tests were performed using an Shimadzu AG-X plus electronic universal testing machine (Shimadzu Corporation, Kyoto, Japan) at a crosshead speed of 0.5 mm/min. The tensile properties were reported as the average values of three replicate measurements.

Vickers hardness testing: Vickers hardness tests were conducted in accordance with GB/T 4340.1-2009 [38], with a test force of 4.9 N and a dwell time of 15 s. Fifteen random indentations were made on the specimen, and the arithmetic mean of the measurements was calculated.

Electron backscatter diffraction (EBSD) measurements were performed using a scanning electron microscope (SEM) equipped with an Oxford Instruments Nordlys Max3 detector (Oxford Instruments plc, Abingdon, Oxfordshire, UK) to acquire crystallographic orientation data. The acquired EBSD data were subsequently processed with TSL OIM™ analysis software(Version 7.3, EDAX, Inc., Mahwah, NJ, USA) to extract information on grain orientation, texture distribution, and grain boundary characteristics.

3. Results

3.1. XRD Phase Analysis

X-ray diffraction (XRD) was used to identify the phase composition of the investigated alloys, and the resulting patterns are shown in Figure 2. For the base Al-Si-Mg alloy, the diffraction pattern exhibits only peaks attributable to α-Al, eutectic Si, and the Mg₂Si phase, in agreement with the equilibrium phase assemblage reported for conventional cast Al-Si-Mg alloys.

Upon the addition of 0.3 wt.% Y, additional diffraction reflections corresponding to the ternary intermetallic Al₂Si₂Y phase become observable [39,40]. These new peaks, however, are found to partially overlap with those of the eutectic Si phase. This overlap arises primarily from two aspects. First, Al₂Si₂Y adopts a hexagonal crystal system (space group P6/mmm, a = b = 0.420 nm, c = 0.664 nm), whereas Si crystallizes in a diamond cubic structure (space group Fd-3m, a = 0.543 nm). The interplanar spacings of certain planes in these two phases are very close, leading to nearly identical diffraction angles under Cu-Kα radiation [6]. Second, because the Y content is only 0.3 wt.%, the volume fraction of the in situ formed Al₂Si₂Y phase is inherently low. Consequently, the diffraction signal from this intermetallic phase is weak and can be easily masked by the intense peaks of the predominant Si phase.

For the alloy synergistically modified with 0.3 wt.% Y and 2 wt.% ZrB₂, the characteristic peaks of both the in situ synthesized ZrB₂ particles and the Y-containing phase appear around 2θ ≈ 35°, where they overlap substantially. This overlapping suggests spatial proximity and likely interfacial interaction between the ZrB₂ nanoparticles and the rare-earth phase. Such an interpretation is further supported by the subsequent SEM-EDS analysis.

3.2. SEM-EDS Results

The microstructural evolution of the investigated alloys was examined by scanning electron microscopy (SEM), and the representative micrographs are compiled in Figure 3, including the unmodified base alloy, the alloys singly modified with 0.3 wt.% Y or 2 wt.% ZrB₂, and the Al-Si-Mg alloys compositely modified with different Y-ZrB₂ ratios.

Figure 3a displays the microstructure of the as-cast base alloy. After electropolishing treatment, the eutectic Si in the matrix exhibits a typical coarse flake-like morphology, a characteristic feature of conventional as-cast Al-Si-Mg alloys. Such a flake-like eutectic Si morphology is generally considered detrimental to mechanical performance, as the sharp edges of the Si plates tend to act as stress concentrators, promoting premature crack initiation under loading [41].

Compared with the base alloy, the introduction of the Y-ZrB₂ modifying phases induces a marked refinement of the eutectic Si structure. As shown in Figure 3b–f, with increasing Y-ZrB₂ composite addition, the eutectic Si undergoes progressive morphological transformation, accompanied by a continuous decrease in grain size. This refinement behavior is commonly observed in Al-Si-Mg alloys upon rare-earth microalloying, where rare-earth elements suppress the growth of eutectic Si by adsorbing at the solid-liquid interface and restricting Si atom diffusion [31,42]. At the optimal modification composition of 0.3 wt.% Y + 2 wt.%ZrB₂ (Figure 3e), the originally coarse flake—like eutectic Si is substantially refined and completely transformed into a uniform, fine rod-like morphology, achieving the most pronounced refinement and morphological modification effect. This morphological transition from flake-like to rod-like Si is known to significantly reduce stress concentration at the Si-matrix interface, thereby contributing to enhanced ductility [30].

Upon further increasing the Y-ZrB₂ addition (Figure 3f), the modification effect noticeably deteriorates: the refined, uniform rod-like eutectic Si regresses into a mixed morphology consisting of both flake-like and rod-like features, accompanied by a pronounced loss of microstructural uniformity. Such over-modification behavior is well documented in rare-earth modified Al-Si alloys, where excessive modifier addition leads to the reappearance of coarse Si phases due to the saturation of adsorption sites and the formation of coarser intermetallic compounds [42].

Figure 4 presents the EDS elemental mapping results of the Al-Si-Mg alloy compositely modified with 0.3 wt.%Y + 2 wt.%ZrB₂. The modified alloy matrix simultaneously contains Y-rich rare-earth phases, in situ ZrB₂ nanoparticles, and Mg₂Si strengthening phases. Notably, the elemental distribution maps reveal that the spatial distribution of the in situ ZrB₂ nanoparticles and the Y-containing rare-earth phases largely overlap. This observation provides direct evidence of spatial proximity and potential interfacial interaction between the two modifying phases.

Such overlapping distribution patterns have been reported in several studies on synergistic modification in aluminum matrix composites, where the co-location of reinforcing particles and rare-earth phases facilitates load transfer across the interface and retards crack propagation [33,43]. This result corroborates the aforementioned XRD analysis (Figure 2), in which the characteristic diffraction peaks of the ZrB₂ phase and the rare-earth Al₂Si₂Y phase overlap at 2θ ≈ 35°. Collectively, the combined XRD and EDS evidence confirms the existence of a significant interfacial coupling and synergistic effect between the in situ ZrB₂ nanoparticles and the rare-earth Y phase, leading to the formation of an integrated composite modification system within the Al-Si-Mg matrix.

3.3. Mechanical Properties and Vickers Hardness

Figure 5 presents the ultimate tensile strength (UTS), yield strength (YS), elongation after fracture (δ), and Vickers hardness (HV) of the base alloy, the alloy singly modified with 0.3 wt.% Y, the alloy singly modified with 2 wt.% ZrB₂, and the Al-Si-Mg alloys compositely modified with different ratios of Y-ZrB₂.

In the as-cast condition, the base alloy exhibits a UTS of 159.11 MPa, YS of 102.34 MPa, δ of 7.75%, and HV of 56.2. With the addition of 0.3 wt.%Y, these properties increase to 181.78 MPa, 130.42 MPa, 12.54%, and 83.5 HV, respectively. For the alloy singly modified with 2 wt.% ZrB₂, the corresponding values are 178.85 MPa, 119.55 MPa, 10.21%, and 106.2 HV.

A comparison reveals that the single ZrB₂-modified alloy exhibits lower strength and elongation but higher hardness than the single Y-modified alloy. This discrepancy arises from the distinct strengthening mechanisms of each approach. ZrB₂ primarily enhances hardness through grain refinement and dispersion strengthening derived from the hard ceramic particles [32], while its influence on the morphological modification of eutectic Si is limited. Consequently, a substantial amount of coarse flake-like eutectic Si persists in the microstructure. Under external loading, these flake-like Si particles readily induce stress concentration and promote microcrack initiation, thereby degrading the load-bearing capacity and fracture resistance of the alloy [44]. In contrast, trace amounts of rare earth Y effectively refine and spheroidize the coarse eutectic Si, mitigating stress concentration risks, while simultaneously achieving multiple synergistic strengthening and toughening effects via grain refinement, morphological modification, and second-phase precipitation strengthening [30].

When 0.3 wt.%Y and 2 wt.%ZrB₂ are added in combination, the as-cast alloy achieves the optimum comprehensive mechanical properties, with UTS, YS, δ, and HV reaching 203.46 MPa, 134.78 MPa, 14.15%, and 133.5 HV, respectively. After T6 heat treatment, the properties of this hybrid modified alloy are further significantly enhanced, attaining 332.87 MPa, 271.35 MPa, 16.24%, and 153.9 HV, respectively, representing a substantial improvement over the base alloy. It should be noted that T6 treatment of Al-Si-Mg alloys typically involves solution treatment followed by artificial aging, which promotes the precipitation of fine Mg₂Si strengthening phases uniformly dispersed within the α-Al matrix, thereby achieving precipitation hardening [41]. This precipitation hardening effect, combined with the morphological refinement of eutectic Si induced by Y addition and the dispersion strengthening from in situ ZrB₂ nanoparticles, accounts for the exceptional mechanical performance of the synergistically modified alloy in the T6 condition.

Figure 6 shows the fracture morphologies of the base alloy, the alloy singly modified with 0.3 wt.% Y, the alloy singly modified with 2 wt.% ZrB₂, and the Al-Si-Mg alloys compositely modified with different ratios of Y-ZrB₂.

For the base alloy (Figure 6a), the fracture surface exhibits extensive flat cleavage facets, river patterns, and a few tear ridges, with few and unevenly distributed dimples, indicating predominantly brittle fracture. The coarse flake-like eutectic Si phase acts as crack initiation sites during fracture, which is consistent with the low elongation data of this alloy. This phenomenon aligns with well-established observations that microcracks in Al-Si-Mg cast alloys are predominantly initiated within Si particles, where shear bands precede particle fracture and dislocation pile-up serves as the dominant damage-initiating process [45].

After the single addition of Y or ZrB₂ (Figure 6b,c), the alloy performance is moderately improved, with an increased number of dimples and reduced cleavage facets, exhibiting a mixed brittle-ductile fracture characteristic. Upon the combined addition of Y and ZrB₂ (Figure 6d–f), the number of dimples further increases, their size distribution becomes more uniform, and brittle cleavage facets are further reduced, indicating the onset of a synergistic modification effect. At the addition level of 0.3 wt.% Y + 2 wt.% ZrB₂ (Figure 6e), the fracture surface consists of numerous fine and uniformly distributed dimples, displaying a typical ductile fracture mode. This indicates that crack propagation is effectively inhibited, and the plasticity reaches its optimum. With excessive addition, the dimples become uneven in size, secondary cracks increase, and the plasticity declines.

The transition from brittle cleavage to ductile dimple fracture is directly correlated with the morphological evolution of eutectic Si from coarse flakes to fine rods/spheroids, which effectively reduces stress concentration at the Si-matrix interface and retards premature crack nucleation [46]. Furthermore, the presence of uniformly dispersed ZrB₂ nanoparticles can induce crack deflection and crack bridging, contributing to enhanced ductile failure behavior [47].

The evolution of fracture morphologies is in good agreement with the mechanical property test results, confirming the regulatory role of Y-ZrB₂ synergistic modification on the fracture mode of the alloy.

3.4. Orientation Distribution of Microstructure

To elucidate the synergistic role of in situ ZrB₂ nanoparticles and rare-earth Y in governing the crystallographic orientation evolution of the Al-Si-Mg matrix, systematic electron backscatter diffraction (EBSD) characterization was performed. Figure 7 presents the EBSD inverse pole figure (IPF) maps of Al-Si-Mg alloys incorporating different amounts of the Y-ZrB₂ composite modifier.

Irrespective of the Y-ZrB₂ addition level, the primary α-Al grains consistently exhibit an equiaxed morphology. With increasing ZrB₂ nanoparticle content, secondary dendrites progressively appear surrounding the α-Al grains. However, the morphological characteristics of the α-Al grains remain largely unaffected by the compositional variations. This observation aligns with the EBSD findings reported by Liu et al. [36], in which the addition of Y-containing nanoscale precipitates was shown to have a negligible influence on the crystallographic orientation stability of α-Al grains. Furthermore, the crystallographic orientations of the α-Al grains are randomly distributed across the {100}, {110}, and {111} planes, without the formation of any pronounced preferred texture, indicating that the Y-ZrB₂ composite modification barely alters the intrinsic crystallographic configuration of the α-Al matrix, which thus retains its disordered growth characteristic.

By contrast, the crystallographic orientation and texture evolution of eutectic Si are significantly regulated by the Y-ZrB₂ addition. In the unmodified Al-Si-Mg alloy, eutectic Si exhibits a distinct preferred growth along the {111} Si plane, whereas no notable texture features appear on the {110} or {112} planes. This preferential orientation is consistent with the well-established impurity-induced twinning (IIT) mechanism, where modifier atoms preferentially adsorb at the {111} Si twin re-entrant edges, thereby facilitating growth along specific crystallographic directions [48]. As the Y-ZrB₂ content increases, prominent texture spots emerge on the {110} plane, implying substantial orientation deviation and lattice rotation of the eutectic Si grains. This evolution substantiates the mechanism that Y atoms preferentially segregate to the growth steps of the Si phase, increasing its lattice constant and thereby suppressing the preferential growth along {111} Si [15]. This is further evidenced by the possible dissolution of Y into ZrB₂ particles during processing, leading to the formation of a nano-modified interfacial layer that effectively constrains the directional solidification of eutectic Si [49].

At the optimal addition level of 0.3 wt.% Y + 2 wt.% ZrB₂, the texture intensity of eutectic Si increases from 8.69 (base alloy) to 10.94, reflecting the maximal regulatory effect on texture refinement. Excessive ZrB₂ addition results in a progressive decline in texture intensity, which is primarily ascribed to the orientation randomization and redistribution of partial eutectic Si grains. This behavior is reminiscent of findings in textured porous Si₃N₄-ZrB₂ composites, where increasing ZrB₂ content was observed to inhibit the degree of grain orientation [50]. The deterioration of texture intensity at excessive modifier levels suggests a saturation of the adsorption sites for Y atoms on the ZrB₂ nanoparticle surfaces, accompanied by the reappearance of coarser eutectic Si structures, which aligns with the microstructural observations presented in Figure 3.

3.5. Grain Boundary Character and Grain Size Distribution

Figure 8 presents the grain boundary character distribution (GBCD) of Al-Si-Mg alloys with various Y-ZrB₂ additions. In the figure, red lines denote low-angle grain boundaries (LAGBs, 2–5°), green lines represent medium-angle grain boundaries (MAGBs, 5–15°), and blue lines indicate high-angle grain boundaries (HAGBs, >15°). Statistical analysis reveals that the unmodified Al-Si-Mg alloy is dominated by LAGBs with a proportion of 81.6%, whereas the fractions of MAGBs and HAGBs are only 9.3% and 9.1%, respectively, exhibiting the typical LAGB-dominated microstructure characteristic of as-cast aluminum alloys.

The introduction of Y-ZrB₂ composite phases enables preferential adsorption and enrichment of nanoparticles at grain boundaries, generating an obvious grain boundary pinning effect via the Zener mechanism and promoting the progressive transition from LAGBs to MAGBs and HAGBs. This phenomenon is consistent with the Zener pinning effect reported by Gutta et al., where nano-sized ceramic particles at grain boundaries effectively suppress grain growth and promote the evolution of boundary misorientation [51]. At the optimal alloy composition of 0.3 wt.% Y + 2 wt.% ZrB₂, the HAGB fraction reaches a peak value of 34.3%, achieving the most favorable reconstruction of grain boundary structure. A high HAGB fraction is known to be beneficial for mechanical properties, as HAGBs can more effectively impede dislocation motion and hinder crack propagation compared to LAGBs [52].

Further increasing the ZrB₂ content leads to severe nanoparticle agglomeration, which weakens the interfacial modification efficiency of rare earth Y and reverses the grain boundary evolution trend, inducing a backward transformation from HAGBs to MAGBs and LAGBs. Statistical data reveal that the HAGB fraction decreases from 34.3% to 27.7%, while the LAGB fraction increases from 61.5% to 65.9%. This phenomenon originates from excessive ZrB₂ accumulation at grain boundaries, which greatly compromises the grain refinement and interfacial optimization effect of rare earth Y. It should be noted that this behavior is analogous to the over-modification phenomenon observed in Al-Si-Mg alloys with excessive Sr addition, where the precipitation of hard Sr-rich particles leads to a deterioration in mechanical properties [53].

Excessive Y-ZrB₂ addition introduces abundant microstructural defects into the matrix, remarkably increasing dislocation density and enabling LAGBs to store high strain energy. Although the stored energy provides a driving force for grain boundary misorientation transition, high-density defects reduce the intergranular fracture energy, weaken the coordinated deformation capacity of grain boundaries, and eventually lead to synchronous deterioration of alloy strength and ductility.

Figure 9 presents the grain size distribution histograms of the Al-Si-Mg alloys with different Y-ZrB₂ addition levels. Combined with the grain boundary misorientation analysis results in Figure 8, it can be observed that the base alloy exhibits a wide grain size distribution range, corresponding to the relatively low proportion of high-angle grain boundaries (HAGBs, >15°) shown in Figure 8a. The coarse-grained structure leads to a limited total grain boundary length and a relatively high proportion of low-angle grain boundaries.

After the introduction of Y-ZrB₂ composite modification, the grain size distribution range narrows significantly, and the grain refinement effect is evident, accompanied by a substantial increase in the total grain boundary length, which is consistent with the increasing trend of the HAGB proportion in Figure 8. Among all the compositions, the addition of 0.3 wt.% Y + 2 wt.% ZrB₂ yields a grain size predominantly concentrated in the range of 100–130 µm, exhibiting the narrowest distribution and the optimal refinement effect. This corresponds to the highest HAGB proportion shown in Figure 8c, indicating that the grain structure uniformity is optimal under this condition. The abundant HAGBs effectively impede crack propagation, which is in good agreement with the observed improvement in mechanical properties.

With a further increase in the Y-ZrB₂ addition level, the grain refinement effect weakens, the size distribution broadens, and the HAGB proportion decreases, ultimately leading to a decline in the mechanical properties of the alloy. This behavior is consistent with the over-modification phenomenon documented in rare-earth modified Al-Si alloys, where excessive modifier addition leads to the reappearance of coarse Si phases and a deterioration of grain uniformity [54].

3.6. Grain Boundary Misorientation Evolution

The grain boundary misorientation distribution of the Al-Si-Mg alloys modified with different Y-ZrB₂ additions is presented in Figure 10. For the unmodified base alloy (Figure 10a), the misorientation between adjacent grains is predominantly concentrated in the range of 0–5°, with only a minor fraction distributed in the theoretical random misorientation range of 40–50°.

The evolution of grain boundary misorientation with Y-ZrB₂ addition is governed by the competition between Zener pinning and nanoparticle agglomeration. At moderate addition levels (Figure 10b,c), in situ synthesized ZrB₂ nanoparticles are well dispersed at grain boundaries, exerting a strong pinning force that forces adjacent grains to rotate toward larger misorientation angles [51,54]. According to the classical Zener pinning model, particles situated at grain boundaries impede boundary migration; importantly, the pinning effect is significantly stronger at low-angle grain boundaries (LAGBs) than at high-angle grain boundaries (HAGBs), which promotes the progressive transition from LAGBs to HAGBs [55]. The dominant misorientation peak gradually shifts from 0–5° toward 30–50°, accompanied by a marked increase in the proportion of random misorientation in the 50–60° range.

This pinning-induced lattice rotation is further assisted by Y-rich precipitates, which segregate at grain boundaries and promote dislocation accumulation. According to Jiang et al., such segregation induces high-density dislocation pile-ups and increases the stored strain energy at the boundaries, which provides an additional thermodynamic driving force for the LAGB-to-HAGB transformation [56]. The synergistic effect of Zener pinning and grain boundary segregation thus accounts for the observed increase in overall grain misorientation.

However, when the Y-ZrB₂ addition exceeds the optimal level (e.g., 0.3 wt.% Y + 3 wt.% ZrB₂ in Figure 10d), severe agglomeration of nanoparticles occurs. This reduces their effective pinning number density, weakens the Zener force, and compromises the grain boundary modification efficiency of the rare earth Y. Consequently, the overall grain misorientation gradually decreases: the dominant misorientation peak shifts back toward lower angles, the HAGB fraction declines, and LAGBs reappear. These observations underscore the critical importance of optimizing the Y-ZrB₂ addition level to maximize the grain boundary structural benefit.

3.7. Geometrically Necessary Dislocation Density

The geometrically necessary dislocation (GND) density distribution of Al-Si-Mg alloys with different Y-ZrB₂ addition levels is presented in Figure 11. GND density was calculated from the kernel average misorientation (KAM) values derived from the EBSD data, with quantitative analysis performed using the TSL OIM™ software. As shown in Figure 11a, the base Al-Si-Mg alloy exhibits relatively low GND density concentrated near grain boundary regions. In contrast, the Y-ZrB₂ modified alloys display distinctly higher GND density in the grain boundary areas, as evidenced in Figure 11b–d. As the Y-ZrB₂ addition increases, GND density progressively rises across the modified alloys, with the highest GND density observed for the 0.3 wt.% Y + 3 wt.% ZrB₂ alloy. Quantitative analysis reveals that the average GND density increases from 0.44 × 10¹⁴ m⁻² for the base alloy to 1.08 × 10¹⁴ m⁻² for the 0.3 wt.% Y + 3 wt.% ZrB₂ alloy. This elevation in GND density is consistent with the trend reported in nanoparticle-reinforced aluminum matrix composites, where increasing reinforcing particle content leads to a higher density of geometrically necessary dislocations due to thermal expansion mismatch and plastic incompatibility between the matrix and reinforcing phases [32,57].

Correlating the GND distribution with the grain boundary evolution presented in Figure 8, the regions of elevated GND density gradually expand from grain boundaries toward grain interiors as the Y-ZrB₂ content increases. This spatial propagation of dislocation-rich zones from interfaces into the grain matrix is consistent with bright-field TEM microstructural observations [36]. The underlying mechanism is primarily ascribed to the grain boundary segregation of Y-ZrB₂ nanoparticles, which induces interfacial dislocation proliferation and promotes the continuous outward propagation of dislocations from boundary-adjacent zones into the matrix interior. This phenomenon aligns with the established understanding that nanoparticle clustering at grain boundaries acts as dislocation sources under applied stress, thereby enhancing the overall GND density within the alloy [51]. The observed GND evolution trend also correlates well with the variation in adjacent and random grain misorientation presented in Figure 10, where higher misorientation angles are generally accompanied by increased GND densities, reflecting the intrinsic coupling between crystallographic rotation and dislocation accumulation during plastic deformation [56,58].

3.8. Schmid Factor Statistics

Figure 12 presents the statistical distribution of Schmid factors for the dominant {111}⟨110⟩ slip system in Al–Si–Mg alloys with different Y-ZrB₂ additions. The Schmid factor describes the ratio of the resolved shear stress on a slip system to the applied uniaxial tensile stress. A smaller Schmid factor implies a lower probability of dislocation slip activation when the same external load is applied. As a result, materials possessing a lower average Schmid factor tend to offer greater resistance to plastic deformation and consequently display enhanced mechanical strength. In contrast, a higher average Schmid factor facilitates the activation of slip systems and promotes extensive crystallographic slip, which is beneficial for ductility but may reduce the yield strength of the material.

When compared with the unmodified base alloy, all Y-ZrB₂ modified alloys exhibit a reduced average Schmid factor. This reduction is attributed to the combined influence of grain refinement and the distributed nanoparticle phases, which act as obstacles hindering dislocation motion and raising the stress required for slip initiation. It is also noteworthy that heterogeneous local stress fields, originating from the non-uniform dispersion of reinforcing phases or from strain mismatches at grain boundaries, can cause appreciable Schmid factor variations even among grains sharing the same crystallographic orientation.

4. Discussion

The combined addition of rare earth Y and in situ ZrB₂ nanoparticles induces a pronounced multi-scale modulation of the Al-Si-Mg alloy microstructure, especially affecting the crystallographic orientation relationship and texture evolution between primary α-Al and eutectic Si grains. Systematic EBSD pole figure analysis (Figure 7, Figure 8, Figure 9, Figure 10, Figure 11, Figure 12 and Figure 13) demonstrates the unique advantage of this composite modification strategy in optimizing both microstructure and mechanical properties.

• Multi-scale modulation of α-Al orientation and eutectic Si texture

The Y-ZrB₂ composite modification has only a weak influence on the spatial orientation distribution of α-Al matrix grains, with no noticeable preferred orientation shift before and after modification. This observation confirms the inherent stability of the crystallographic growth configuration of the Al matrix [59,60]. Neither second-phase particles nor rare earth doping can alter the intrinsic lattice growth characteristics of α-Al grains.

In contrast, rare earth Y exerts a decisive and targeted regulation on the anisotropic growth, lattice orientation evolution, and preferred growth mode of eutectic Si [6,61]. Through interfacial segregation and atomic adsorption, Y suppresses the continuous growth of eutectic Si along its intrinsic crystallographic direction and breaks its coarse lamellar directional growth habit [6,62].

Local orientation analysis of eutectic Si grains (Figure 13a,b) reveals essential differences in grain misorientation. In the unmodified alloy, adjacent eutectic Si grains exhibit misorientation angles of 0.75°/⟨1 2 3⟩, 4.22°/⟨0 2 3⟩, and 2.01°/⟨1 1 2⟩, all below 5°, which are typical low-angle grain boundary (LAGB) characteristics [62]. The atomic arrangement inside Si grains remains highly ordered with low lattice distortion, and eutectic Si grows in a single-orientation directional mode with strong anisotropy.

At the optimal composition (0.3 wt.% Y + 2 wt.% ZrB₂), adjacent eutectic Si grains undergo significant lattice rotation and orientation reconstruction. Their misorientation angles increase to 42.95°/⟨1 3 3⟩, 29.38°/⟨0 2 3⟩, and 54.93°/⟨0 3 4⟩, forming abundant high-angle grain boundaries (HAGBs) [60]. The refinement and structural reconstruction of eutectic Si at grain boundaries break the original directional growth mode and induce a systematic texture transformation. The orientation evolution shows a positive correlation with grain deflection and increased texture intensity.

{111} pole figure analysis further verifies the texture reconstruction mechanism induced by Y-ZrB₂ synergy. The unmodified alloy displays four regular concentrated bright spots corresponding to the directional growth of eutectic Si. After modification, the pole figure spots appear mostly in pairs with clear angular deviation, increasing texture dispersion [59,60].

• Formation of a Y-ZrB₂ nano-transition layer and interfacial energy modulation

Combined with grain boundary evolution (Figure 8) and GND variation (Figure 11), Y-ZrB₂ composite phases preferentially segregate and accumulate at grain boundaries during solidification, raising the GND density from 0.44 × 10¹⁴ m⁻² to 1.08 × 10¹⁴ m⁻² [59,60]. The dislocation pinning and strain accumulation induced by Y-modified ZrB₂ nanoparticles elevate interfacial energy and create high-density local dislocations and strain-concentrated zones at the Si/Al interface [59].

From the perspective of crystal defect theory, a high density of interfacial dislocations aggravates lattice distortion at phase boundaries, forcing lattice deflection and orientation adjustment of eutectic Si grains [60]. The texture intensity of the Y-ZrB₂ modified alloy is lower than that of the alloy modified only with Y [36]. This arises from interfacial adsorption competition: active Y atoms are preferentially adsorbed onto ZrB₂ surfaces, forming a nano-transition layer [6,59], which consumes some Y atoms that would otherwise passivate the Si growth interface and regulate growth habits. Consequently, the texture-modifying effect of Y alone is weakened, confirming that Y dominates the growth direction and crystallographic orientation of eutectic Si [6,61].

• Thermodynamic-kinetic correlation and eutectic Si refinement mechanism

The layered structure of the Y-ZrB₂ composite does not directly alter the crystallographic orientation of eutectic Si. Instead, it improves the interfacial wettability between the reinforcing particles and the Al matrix and increases the heterogeneous nucleation activation energy (ΔG*↑) [59,60]. This thermodynamic-kinetic coupling promotes uniform dispersion of in situ ZrB₂ nanoparticles, improves the heterogeneous nucleation rate, and suppresses particle agglomeration, achieving size refinement and homogeneous distribution of the ZrB₂ nanoparticles [59,60].

The synergistic effect between Y and ZrB₂ achieves comprehensive optimization of the microstructure and crystallographic orientation. The interfacial adsorption behavior of Y on ZrB₂ can be described by the Langmuir-McLean model. Y adsorption reduces the Si/Al interfacial energy and suppresses the preferential {111}Si growth [6,63].

• Essential mechanism for strength-ductility synergy

The increase in GND density is closely correlated with the transition of grain boundary character distribution (GBCD); specifically, the LAGB fraction decreases from 81.6% while the HAGB fraction rises to 34.3% [60,61]. High-density GNDs promote strength enhancement through the Taylor-Orowan strengthening mechanism. The transition from LAGBs to HAGBs improves the grain boundary resistance to dislocation motion and increases the work-hardening capacity. An increased HAGB fraction also benefits ductility by absorbing deformation energy and hindering crack propagation [59,60].

Excessive Y-ZrB₂ addition (beyond 0.3 wt.% Y + 2 wt.% ZrB₂) causes severe ZrB₂ agglomeration and a decrease in Y adsorption saturation, inducing abnormal dislocation accumulation and reducing intergranular fracture energy. This restricts further improvement of mechanical properties and may even cause simultaneous deterioration of strength and ductility. Hence, there exists an optimal composition window for Y-ZrB₂ synergistic modification, and precise control of the composition is key to achieving strength-ductility synergy. The coupling of grain boundary character distribution and grain boundary engineering (GBE) further clarifies the intrinsic relationship between grain boundary structure and macroscopic mechanical performance.

In summary, the synergistic addition of rare earth Y and in situ ZrB₂ nanoparticles constructs a Y-modified ZrB₂ nano-transition layer, precisely modulates the Si/Al interfacial energy, suppresses preferential {111} Si growth, induces local lattice rotation, and promotes interfacial dislocation proliferation. This multi-scale thermodynamic–kinetic coupling mechanism realizes eutectic Si refinement, texture redistribution, interfacial dislocation proliferation, and eventual strength-ductility synergy of the alloy.

5. Conclusions

Based on the systematic investigation of grain orientation, grain boundary character distribution, geometrically necessary dislocation (GND) density, and Schmid factor analysis, the influence of Y-ZrB₂ addition on the microstructural evolution and mechanical performance of Al-Si-Mg alloys is clarified. The main findings are drawn as follows:

The Y-ZrB₂ composite modification has a minor influence on the crystallographic orientation and morphological stability of primary α-Al grains, but exerts a strong regulation on the lattice rotation, texture redistribution, and growth behavior of eutectic Si. In the unmodified alloy, eutectic Si grains exhibit low misorientation angles and a dominant {111} preferred orientation. At the optimal addition (0.3 wt.% Y + 2 wt.% ZrB₂), significant lattice rotation and orientation randomization of eutectic Si occur, leading to a much more uniform texture.

The introduction of Y-ZrB₂ promotes a progressive transition from low-angle grain boundaries (LAGBs) to high-angle grain boundaries (HAGBs). The HAGB fraction reaches its maximum of 34.3% at the optimal composition. Excessive addition causes severe nanoparticle agglomeration, weakens the grain boundary pinning effect, and reverses the grain boundary evolution, resulting in a decrease in HAGBs and an increase in LAGBs.

The synergistic effect of Y and ZrB₂ significantly increases the GND density (from 0.44 × 10¹⁴ m⁻² to 1.08 × 10¹⁴ m⁻²) while simultaneously reducing the Schmid factor of the dominant {111}⟨110⟩ slip system. This dual action strengthens the alloy through the Taylor-Orowan dislocation mechanism and restricts the initiation of plastic slip, thereby enhancing both deformation resistance and work-hardening capacity.

There exists an optimal composition window for Y-ZrB₂ synergistic modification (0.3 wt.% Y + 2 wt.% ZrB₂). Within this window, the alloy achieves the best combination of microstructural refinement, texture homogenization, and mechanical properties: ultimate tensile strength 332.87 MPa, yield strength 271.35 MPa, elongation 16.24%, and Vickers hardness 153.9 HV after T6 heat treatment. Either too low or too high addition deteriorates the comprehensive performance, confirming that precise composition control is the key to realizing strength–ductility synergy.