Effect of Al-5Ti-2B on the Microstructure and Mechanical Properties of Recycled Al-7Si-0.3Mg-1Fe Alloy

Abstract

This study systematically investigates the influence of grain refinement on the microstructural evolution and mechanical properties of recycled Al-7Si-0.3Mg-1Fe alloy through the addition of varying concentrations (0–1.25 wt.%) of Al-5Ti-2B master alloy. The synergistic effects of Al-5Ti-2B on the α-Al phase, eutectic Si, and Fe-rich intermetallics were characterized using metallographic analysis, XRD, SEM-BSE imaging, and EDS. In the unrefined alloy, the microstructure consisted of an α-Al solid solution with coarse plate-like eutectic Si, while Fe primarily formed needle-like β-Al₅FeSi phases that either surrounded or penetrated the eutectic Si. Increasing the Al-5Ti-2B addition refined both the α-Al dendrites and eutectic Si, while the β-Al₅FeSi phase transitioned from coarse to fine needles. The optimal refinement was achieved at a 1% Al-5Ti-2B addition, yielding a tensile strength of 149.4 MPa and elongation of 4.3%. However, excessive addition (1.25%) led to eutectic Si aggregation and β-Al₅FeSi coarsening, resulting in mechanical property deterioration and brittle fracture behavior. These findings provide insights into optimizing grain refinement for enhancing the performance of recycled Al-Si-Mg-Fe alloys.

1. Introduction

As global carbon neutrality strategies advance, the recycled aluminum industry has become a key pathway for achieving energy conservation and emissions reduction in the non-ferrous metals sector. Statistical data show that the energy consumption of recycled aluminum production is only 5% of that of primary aluminum, while approximately 40% of global aluminum alloy demand is met through recycled sources. Among various recycled aluminum alloys, Al-Si-Mg alloys are widely used in automotive components, electronic housings, and other fields due to their excellent casting performance and mechanical properties. However, during the recycling process, recycled aluminum inevitably becomes contaminated with iron-based impurities, leading to iron enrichment. Research data indicate that the iron content in scrap aluminum can reach 3–5 times that of primary aluminum, with the iron content in recycled Al-7Si-0.3Mg alloys typically exceeding 0.5 wt.% [1]. Al-Si-Mg alloys with higher iron content exhibit larger eutectic silicon areas and more severe eutectic silicon aggregation, making eutectic silicon prone to peeling, reducing the casting performance. Iron-rich intermetallic compounds may also serve as nucleation sites for porosity, increasing porosity rates and causing casting defects [2]. Iron primarily exists in aluminum alloys as needle-like β-Al₅FeSi phases. These hard and brittle phases disrupt the continuity of the matrix, becoming the primary crack sources in tensile fractures, promoting crack initiation and propagation [3], as well as significantly reducing the alloy’s strength and elongation. When the Fe content exceeds 0.8%, the alloy’s elongation may decrease by over 50% [4]. More critically, traditional purification processes struggle to economically and effectively remove iron, rendering the mechanical properties of high-iron recycled aluminum alloys (Fe > 0.8%) unable to meet industrial standards. This has become the core bottleneck constraining the high-value utilization of recycled aluminum.

The modification research on iron-rich aluminum alloys primarily focuses on the following three aspects: (1) Control of the iron-phase morphology, with the addition of elements such as Mn and Cr, of which Mn preferentially combines with Fe to form the α-Al(Fe,Mn)Si phase, replacing the brittle β-Al₅FeSi phase (needle-like or plate-like), while Cr tends to form independent intermetallic phases, such as Al₁₅(Fe,Cr)₃Si₂ [5,6]. Studies have shown that adding 0.5–1.0 wt.% Mn can reduce the β-phase proportion from 25% to below 5%, while controlling the α-phase size within the range of 5–10 μm [7]. Cr promotes the non-uniform nucleation of primary Si and restricts its growth, thereby refining the grain size [8]. (2) Rapid solidification technology inhibits iron-phase growth. During rapid solidification (cooling rate of 10⁵–10⁷ K/s), Fe elements in the aluminum matrix are “frozen” within the α-Al lattice due to an insufficient diffusion time, significantly reducing Fe segregation at the grain boundaries [9]. (3) Melt overheating treatment promotes iron-phase spheroidization [10], but it has high energy consumption costs. Although these methods have made some progress, they still have obvious limitations, as follows: Mn addition requires strict control of the Fe/Mn ratio, as excess Mn can lead to the formation of enriched phases such as Al₁₅(Fe,Mn)₃Si₂ or Al₆(Fe,Mn) [11], whose crystal structures are predominantly cubic or hexagonal [12], resulting in poor compatibility with the aluminum matrix, and they can easily become stress concentration points [13]. Rapid solidification technology is constrained by equipment investment, and melt overheating treatment is only applicable to specific composition systems. Recent studies have found that grain refinement can indirectly improve the microstructural uniformity of iron-rich aluminum alloys. Grain refinement increases the intergranular area per unit volume, providing more heterogeneous nucleation sites for the β-Al₅FeSi phase. This spatial restriction effect significantly reduces the anisotropic growth trend, resulting in a more uniform short rod-like or granular morphology of the β phase [14]. The three-dimensional network-like grain-boundary structure (grain-boundary spacing <10 μm) formed by the fine-grained microstructure alters the diffusion kinetics of Fe elements. Molecular dynamics simulations show that the grain-boundary diffusion coefficient is 2–3 orders of magnitude higher than bulk diffusion, effectively suppressing the localized enrichment of Fe elements in grain-boundary regions. Experimental measurements indicate that the Fe segregation degree can be reduced from 15% in the original state to below 5%, thereby suppressing the preferential growth of the β phase along specific crystallographic directions [15]. Modalavalasa et al. [16] noted that while the addition of Co does not alter the morphology of the β phase, it refines its distribution. The fine-grained microstructure provides more grain boundaries as nucleation sites for phase formation, thereby disrupting its oriented growth. Simultaneously, it causes multiple deflections in crack propagation paths, consuming more fracture energy and enhancing the matrix’s resistance to stress concentration [17]. The above provides new insights into regulating the performance of iron-rich recycled aluminum alloys through grain refiners. However, the mechanism of the grain refinement behavior of traditional Al-Ti-B grain refiners in iron-rich aluminum–silicon–magnesium alloy systems remains poorly understood, particularly regarding potential interface reactions, element segregation, or chemical interactions between Fe and TiB₂ particles, which may significantly influence the heterogeneous nucleation efficiency. This critical scientific issue has not been systematically explored in existing research and urgently requires in-depth analysis to guide the microstructural regulation of high-performance aluminum alloys.

This study systematically controlled the addition of Al-5Ti-2B refiner (0–1.25%) to investigate its effect on the as-cast microstructure and mechanical properties of iron-rich (1%) recycled Al-Si-Mg alloy, with a focus on quantitatively analyzing the effect of TiB₂ particles on the efficiency of α-Al heterogenous nucleation, establishing a structure–property relationship between the addition amount and grain size, revealing the regulatory effect of Al-5Ti-2B on the morphology and spatial segregation behavior of the β-Al₅FeSi phase and determining the optimal addition amount to suppress the harmful effects of the Fe phase through a combination of microstructural characterization and performance testing. By conducting a comprehensive analysis of the “refining agent dosage–microstructure–performance” chain, this study aims to advance the development of processes for the high-value utilization of recycled aluminum.

2. Materials and Methods

The experimental materials for this study were obtained from a company in Baise City and consisted of recycled Al-Si-Mg aluminum alloy (Si: 7.51%; Mg: 0.31%; Fe: 1%; Al balance), hereinafter referred to as recycled Al-7Si-0.3Mg-1Fe alloy. By adding different proportions of Al-5Ti-2B intermediate alloy, the base material was remelted to prepare the corresponding experimental samples. The specific process is as follows: (1) Melting process—Melting was performed in a crucible melting furnace at 750 °C→Al-5Ti-2B intermediate alloy was added→the melt was maintained at temperature and allowed to settle for 15 min. (2) Refining treatment—Refining was carried out by rotating stirring with high-purity nitrogen introduced from the bottom (10 min)→after slag removal, the melt was allowed to settle for another 10 min. (3) Casting and forming—When the melt temperature dropped to 720 °C, it was poured into a preheated steel mold (300 °C) to prepare the tensile test specimens. For the Al-Si-Mg-Fe alloys, due to their high eutectic reaction temperature (approximately 577 °C), the Al-5Ti-2B grain refiner was prone to poisoning by Si in a silicon-rich environment, promoting the precipitation of the TiSi₂ phase, thereby weakening the refining effect. Additionally, the coarsening of the iron-rich phases (such as β-Al₅FeSi) further deteriorated the alloy’s mechanical properties, particularly the ductility and toughness. Therefore, to suppress the formation of harmful phases and to optimize the microstructure, a rapid cooling strategy was required to minimize the alloy’s residence time in the high-temperature zone (>500 °C). Based on this, the demolding time in this experiment was strictly controlled to within 45 s to ensure that the casting rapidly separated from the mold after basic solidification, thereby maximizing the grain refinement effect of Al-5Ti-2B and minimizing the adverse effects of the brittle phases.

For different Al-5Ti-2B addition ratios (as shown in

Table 1. Specimen ID and Al-5Ti-2B (wt.%).

Specimen ID	Al-5Ti-2B Additions (wt.%)
1	0
2	0.25%
3	0.5%
4	0.75%
5	1%
6	1.25%

), five tensile specimens were prepared in parallel for each group and tested. The specimen dimensions were as follows (Figure 1): gauge length, L₀ = 50 mm, and original diameter, d₀ = 10 mm. Concurrently, matching the metallographic specimens, the SEM-BSE analysis specimens were prepared as follows: metallographic specimens were mechanically ground and polished, then etched with Keller’s reagent for 10–20 s; the SEM-BSE specimens were only polished (no etching required). After the tensile test, samples were taken 5 mm below the fracture surface for the SEM fracture morphology observation. The specific correspondences between the specimen numbers and Al-5Ti-2B addition ratios are detailed in Table 1.

Phase identification was performed using an X-ray diffractometer (XRD, Rigaku SmartLab 9, Tokyo, Japan). Backscattered electron (BSE) images, tensile fracture morphology, elemental distribution, and composition analysis were obtained using a tungsten-filament scanning electron microscope (SEM, Hitachi SU5000, Tokyo, Japan) equipped with energy-dispersive spectroscopy (EDS). Room-temperature tensile tests were conducted using a universal testing machine (CMT-4015, Shenzhen Junrui Instrument, Co., Ltd., Shenzhen, China) to determine the ultimate tensile strength and elongation. Round specimens were employed for testing, with five replicates performed for each sample type. The maximum load (F_max, unit: N) sustained during the tensile testing was recorded for each trial, and the average value was calculated. The original cross-sectional area (A₀) of the specimens was measured, and the ultimate tensile strength (R_m, unit: MPa) was calculated using Equation (1). Elongation was determined as the ratio of permanent elongation after fracture to the original gauge length, expressed as a percentage. The elongation (A) was calculated using Equation (2), where L₀ represents the original gauge length, and L_u denotes the final gauge length after fracture and reassembly.

$$R_m={\displaystyle \frac{F_{max}}{A_0}}$$

(1)

$$A={\displaystyle \frac{L_u-L_0}{L_0}}\times 100\%$$

(2)

The microstructure of the specimens was observed using an optical microscope (OM, Axio Observer 3m, Beijing Precise Instrument Co., Ltd., Beijing, China), and micrographs were taken. Metallographic images at the same magnification ratio (×100) were selected, and ImageJ 1.52 image software was used to perform binarization processing on the metallographic images to highlight the grain boundaries. Based on the processed metallographic images, the average grain boundary length, $\overline{L}$, was measured using the straight-line intersection method. The grain boundary length refers to the length of the line segment between the intersection points of the test line and the grain boundary. By calculating the average value of a large number of grain boundary lengths, the grain size can be reflected. The specific method is as follows: (1) Overlay a set of parallel straight lines on the processed metallographic images, ensuring that the total length of the lines covers multiple grains. (2) Count the total number of intersection points P between the test lines and the grain boundaries and measure the total length, L_T, of the test lines. (3) Calculate the average intercept using the formula shown in Equation (3).

$$\overline{L}={\displaystyle \frac{L_T}{P}}$$

(3)

3. Results

3.1. Metallographic Analysis

Figure 2 shows the as-cast microstructure of the recycled Al-7Si-0.3Mg-1Fe alloy at different Al-5Ti-2B addition levels. The microstructure of the alloy without Al-5Ti-2B addition is shown in Figure 2a. The matrix is primarily composed of an α-Al solid solution, with eutectic Si forming coarse plate-like structures. The Mg elements are mainly dissolved in the matrix or form nanoscale Mg₂Si strengthening phases [1,2,18], but due to their small size, they are difficult to distinguish in the figure. Fe primarily exists in the alloy as needle-like β-Al₅FeSi phases, which grow from the eutectic silicon boundaries and penetrate the entire eutectic silicon structure [3]. Iron is a harmful impurity in aluminum–silicon alloys. When its content is too high, it forms coarse, brittle intermetallic compounds (such as β-Al₅FeSi). These compounds tend to segregate at grain boundaries, acting as stress concentration points that trigger microcracks and significantly reduce the strength and plasticity of the alloy [4,19]. For recycled aluminum alloys, the presence of Fe is very common; however, completely separating iron from the aluminum melt is quite challenging. Current research primarily focuses on regulating the morphology of iron phases [5,6,7,8], inhibiting their growth [9], and promoting spheroidization [10,11,12,13]. This experiment attempted to add Al-5Ti-2B intermediate alloy to refine the matrix structure while improving the morphology of the iron phases, aiming to enhance the performance of recycled aluminum alloys. Figure 2b–f show the micrographs of the alloy after adding Al-5Ti-2B, with addition levels of 0.25%, 0.5%, 0.75%, 1%, and 1.25%. As shown in the figures, with the increasing addition of Al-5Ti-2B, the primary α-Al dendrites gradually decreased in size, and the eutectic Si morphology transitioned from plate-like to granular and fibrous. Figure 3 shows the XRD pattern of the alloy with 1% Al-5Ti-2B added. The experimental results indicate that the XRD patterns of the base alloy remain unchanged after adding different amounts of Al-5Ti-2B. This may be due to the low addition amount of Al-5Ti-2B, whose diffraction signals are masked by the base alloy, or because the elements dissolve into the alloy without forming new phases. Combining the XRD patterns in Figure 3, as the addition of Al-5Ti-2B increases, the number of TiB₂ particles per unit volume of the aluminum melt increases. These particles serve as heterogeneous nucleation substrates for α-Al, with Al₃Ti forming a thin film on their surfaces, promoting α-Al heterogeneous nucleation and increasing the nucleation rate, thereby refining the matrix grains [20,21]. As the addition of Al-5Ti-2B increases, TiAl₃ accumulates or adsorbs onto the surface defects of the β-Fe phase, hindering its growth and causing its morphology to transition from needle-like to short, rod-like structures. When the Al-5Ti-2B addition reaches 1.25%, compared to a 1% addition, the alloy grain size increases, showing a trend toward grain coarsening. Lin et al. [22] and Ke et al. [23] found that Si in the alloy reacts with Ti to form Ti-Si compounds, consuming free Ti and reducing the grain refinement effect, known as the “silicon poisoning” phenomenon. Additionally, Chowwanonthapunya & Peeratatsuwan [24] investigated the effect of Fe on the grain refinement of Al-5Ti-1B in recycled A356 alloy, finding that as the number of cycles increased, the concentrations of Ti and B decreased, and the grain refinement effect weakened. This may be due to the presence of Fe promoting the consumption of Ti or forming complex Fe-Ti-B phases, thereby affecting the effectiveness of TiB₂.

To further investigate the effect of Al-5Ti-2B addition on grain refinement in the alloy, an Axio Observer 3m metallographic microscope was used to measure the grain size in the microstructure of the alloy using the straight-line intercept method. Figure 4 shows the relationship between the addition of Al-5Ti-2B and the grain intercept. As shown in the figure, with an increase in the addition of Al-5Ti-2B, the average grain intercept of the alloy exhibits a trend of first decreasing and then increasing. When no Al-5Ti-2B is added, the average grain intercept is 18.12 μm; when the addition is 0.25%, the average grain intercept is 16.96 μm; when the addition is 0.5%, the average grain intercept is 16.78 μm; when the addition is 0.75%, the average grain size is 16.5 μm; when the addition is 1%, the average grain size is 16.42 μm; and when the addition is 1.25%, the average grain size is 17.06 μm. The grain size did not decrease further but, instead, showed a trend toward increasing, confirming the previous discussion.

3.2. BSE Imaging and EDS Analysis

To investigate the regulatory effect of Al-5Ti-2B on the morphology and spatial segregation behavior of the β-Al5FeSi phase, SEM and EDS analyses were conducted on recycled Al-7Si-0.3Mg-1Fe alloys containing different amounts of Al-5Ti-2B. Figure 5a–f show the SEM morphology of the recycled Al-7Si-0.3Mg-1Fe alloy, with Al-5Ti-2B addition levels of 0%, 0.25%, 0.5%, 0.75%, 1%, and 1.25%. Figure 6 shows the EDS elemental layer map of the recycled Al-7Si-0.3Mg-1Fe alloy. As shown in the figure, Fe elements are attached to the edges of the eutectic silicon, while Mg elements are uniformly distributed throughout the matrix. Figure 6a shows the EDS elemental layer map of the alloy without Al-5Ti-2B addition, and Figure 6b shows the EDS elemental layer map of the alloy with 0.5% Al-5Ti-2B addition. Combining Figure 5, Figure 6 and Figure 7, it can be seen that the dark gray elongated strips in Figure 5 represent eutectic Si, and TiB₂ and TiAl₃ are dispersed around the eutectic silicon. As the addition of Al-5Ti-2B increases, TiAl₃ aggregates on the surface of the eutectic silicon, inhibiting its growth, thereby refining the alloy grains. Mg forms the Mg₂Si strengthening phase with Si, which is dispersed throughout the alloy matrix, enhancing the alloy’s mechanical properties. The bright white, elongated, and needle-like structures are β-Fe phases, which grow around the eutectic silicon, disrupting the alloy matrix, and severely impairing the alloy’s mechanical properties. As shown in Figure 5b–f, with an increase in the addition of Al-5Ti-2B, the number of β-Fe phases in the matrix structure decreases. When the addition of Al-5Ti-2B reaches 1%, the refinement effect of the β-Fe phase is significant, with the elongated needle-like structures becoming finer and smaller. The morphology of the eutectic silicon also changes, transitioning from strip-like to granular. When the addition of Al-5Ti-2B reaches 1.25%, both the Fe phase and the Si phase exhibit obvious coarsening phenomena. The β-Fe phase appears as coarse, elongated needles intertwined within the grains and tending to aggregate, thereby weakening the mechanical properties of the alloy. In Figure 5f, the eutectic silicon exhibits severe aggregation. Chen et al. [2] and Song et al. [25] found that alloys with a higher iron content have larger eutectic silicon areas and more severe eutectic silicon aggregation. Additionally, iron-rich intermetallic compounds may serve as nucleation sites for pores, increasing the porosity and leading to casting defects.

Figure 7 and Figure 8 show the EDS spectra of the alloy without addition and with an Al-5Ti-2B addition of 0.5%. As can be seen from the figures, the main elements in the alloy are Al and Si, with a relatively low content of Mg. Based on the metallographic microscopy, SEM morphology, and XRD analysis, the limited Mg content primarily dissolves into the Al matrix to form Mg₂Si strengthening phases, which are difficult to detect in the XRD pattern due to their weak peaks. As shown in Figure 3, the XRD pattern of the alloy without Al-5Ti-2B reveals a faint peak at 2θ = 44.7°, corresponding to the α-Fe phase. In contrast, after adding 1% Al-5Ti-2B, a distinct diffraction peak of the β-Al₅FeSi phase appears at 2θ = 32.5°. However, the XRD analysis results demonstrate that increasing the Al-5Ti-2B addition from 0.25% to 1.25% does not lead to significant changes in the diffraction peaks. Therefore, only the XRD patterns of the alloy without Al-5Ti-2B and with 1% Al-5Ti-2B addition were compared in the subsequent analysis. The EDS elemental mapping further reveals that Mg is predominantly distributed at the interface between α-Al and eutectic Si, forming Mg₂Si particles that disperse throughout the alloy matrix. These particles impede dislocation motion, thereby enhancing the alloy’s strength. The key secondary phases in the Al-5Ti-2B grain refiner are TiB₂ and Al₃Ti. According to Wang et al. [21], the (0001) plane of TiB₂ provides a low-energy nucleation substrate for α-Al, while the Al₃Ti phase forms an “Al₃Ti/TiB₂” composite structure by encapsulating TiB₂, further reducing the nucleation barrier. Even in Fe-rich alloys, this dual-phase synergistic effect remains dominant.

To verify the findings described above and to understand the distribution of the phases in the alloy, an EDS (energy-dispersive spectroscopy) analysis was performed on the recycled Al-7Si-0.3Mg-1Fe alloy, as shown in Figure 9. Figure 9a shows the EDS spectrum image obtained when the addition of Al-5Ti-2B was 1%, and Figure 9b shows the SEM backscatter image when the addition of Al-5Ti-2B was 0.5%. Figure 9c–f are the energy spectrum diagrams of points A, B, C, and D in Figure 9a,b, respectively. From Figure 9a,c, it can be seen that the dark gray elongated structures in the alloy microstructure are eutectic Si. According to the Al-Si alloy phase diagram, primary α-Al precipitates first from the aluminum melt, and through the eutectic reaction, Al-Si eutectic silicon forms between the dendrite arms. After cooling, Mg₂Si and Si particles precipitate. Point B in Figure 9a contains the following four elements: Al, Si, Mg, and Fe. As shown in Figure 9d, due to the high Fe content in the matrix alloy, Mg consumes Si to form the Mg₂Si phase, indirectly affecting the composition and morphology of the Al-Fe-Si phase. Doroshenko et al. [26] found that in Al-6Mg-2Ca-2Zn alloys, the addition of Mg transforms the Fe-Al-Si phase from needle-like to Chinese-character-like, with the results similar to those of this experiment. Therefore, the microstructure at point B should be the AlFeSi phase. In this case, Mg dissolves into the Al matrix, exerting a solid solution strengthening effect and indirectly influencing the precipitation kinetics of the Fe-Si intermetallic compound [27]. Additionally, Mg enriches at the grain boundaries, inhibiting the growth of coarse Fe phases and promoting the formation of nanoscale α-AlFeSi phases [28]. According to the energy spectrum diagram at point C in Figure 9e, the bright white stripes in Figure 9a may be α-Fe(Al₁₂Fe₃Si) or β-Fe(Al₅FeSi or Al₉Si₂Fe₂), but the Si content in the α-Fe phase is relatively low. Combined with the XRD patterns in Figure 3, the structure is identified as the β-Al5FeSi phase. Its morphology is needle-like or plate-like, growing along the eutectic silicon boundaries and penetrating the entire eutectic structure. If Fe elements are locally enriched due to uneven stirring during melting or other factors, the Fe-rich phase will serve as an excellent nucleation site for pores during solidification, leading to increased porosity in the casting and a reduced casting performance of the alloy.

In summary, as shown in Figure 9a, when the addition of Al-5Ti-2B increases to 1%, the Fe-rich phase exhibits a fine needle-like or fibrous morphology, primarily in the form of a β-Fe phase. As shown in Figure 5e, the needle-like or plate-like β-Fe morphology becomes finer and shorter, and it is insufficient to penetrate the dendrites, resulting in improved mechanical properties of the alloy. From the energy spectrum at point D in Figure 9f, it can be seen that this point contains the following four elements: Al, Si, Ti, and B. This is because the addition of Al-5Ti-2B to the regenerated Al-7Si-0.3Mg-1Fe alloy forms particles such as Al₃Ti and TiB₂, which act as heterogeneous nucleation or nucleation promotion agents, refining the alloy grains. As the addition of Al-5Ti-2B increases, Al₃Ti aggregates at defects on the eutectic silicon surface, inhibiting its growth and achieving grain refinement of the eutectic silicon. Meanwhile, TiB₂ promotes the nucleation rate of Al₃Ti, thereby achieving grain refinement.

3.3. Tensile Properties Analysis

As shown in

Table 2. Tensile strength and elongation of Al-7Si-0.3Mg-1Fe alloys.

Specimen Number	Tensile Strength (MPa)	Elongation (%)
1	117.7	2.5
2	121.1	2.0
3	139.2	1.9
4	145.7	3.3
5	149.4	4.3
6	142.6	2.5

and Figure 10, the tensile strength of the recycled Al-7Si-0.3Mg-1Fe alloy first increases and then decreases with an increase in the Al-5Ti-2B addition. When the Al-5Ti-2B addition is 1%, the tensile strength and elongation of the alloy reach their maximum values at 145.7 MPa and 3.3%, respectively. Fe forms β-Fe phases in Al-Si-Mg alloys, which are relatively coarse and reduce the fluidity of the aluminum alloy melt, affecting the feeding of the melt and leading to shrinkage porosity in castings, thereby degrading the casting performance of the alloy [2]. Additionally, the hard and brittle, elongated, and needle-like Fe phase causes cracks to propagate preferentially along the Fe phase/silicon interface, severely damaging the alloy’s mechanical properties [29]. Bacaicoa et al. [3] found that the interface between the β-Al₅FeSi phase and the matrix is prone to microcracks due to differences in thermal expansion coefficients (β phase: 19.5 × 10⁻⁶/K, while the Al matrix has a coefficient of 23.6 × 10⁻⁶/K), microcracks are prone to form. Eutectic Si particles act as obstacles during crack propagation, causing cracks to deflect and branch [30]. At Al-5Ti-2B addition levels between 0 and 1%, combined with SEM images in Figure 5, eutectic Si primarily distributes at the α-Al grain boundaries, and bright white β-Al₅FeSi phases are distributed around the eutectic Si, with some penetrating the entire eutectic structure. As the Al-5Ti-2B addition increases, the distribution of the bright white β-Al₅FeSi phases with coarse, elongated, and needle-like morphology decreases per unit area, and the morphology of the eutectic Si transitions from rod-like to granular, with the. distribution becoming more uniform, reducing interfacial stress. TiB₂ and TiAl₃ are distributed around the eutectic Si, and the pinning dislocation effects are enhanced, further improving the alloy’s tensile strength and mitigating the adverse effects of β-Fe on the alloy’s mechanical properties. When the Al-5Ti-2B addition reaches 1.25%, the grain size increases again, grain refinement strengthening weakens, and the β-Fe phase forms coarse needle-like structures penetrating the entire eutectic Si structure, disrupting the aluminum alloy matrix. Concurrently, eutectic Si exhibits enrichment phenomena, leading to a decrease in the alloy’s mechanical properties. Song et al. [25] investigated the microstructure and fracture behavior of recycled Al-Si-Mg alloys, finding that the Fe-Si phase serves as an excellent nucleation site for pores, resulting in an increase in pores within the alloy castings and an increase in inter-dendritic defects. The hard and brittle β-Fe phase causes the alloy to develop a series of microcracks, with the stress concentration at the crack tips, promoting crack initiation and reducing the alloy’s tensile strength.

3.4. Fractography Analysis

Figure 11 shows the fracture surface morphology of the alloy after adding different amounts of Al-5Ti-2B. According to the tensile test results, as the addition of Al-5Ti-2B increases, the mechanical properties of the alloy first improve and then deteriorate. The fracture mechanism is analyzed through fracture surface morphology analysis. As reported in the literature [31,32,33], in unmodified iron-rich aluminum–silicon–magnesium alloys, the primary crack propagates along the β-Al₅FeSi phase/α-Al interface, forming an intergranular fracture path; eutectic silicon particles act as stress concentration points, inducing secondary cracks, and the fracture surface exhibits extensive cleavage planes. Based on the test results, the macro fracture surfaces of all samples show no obvious yield or necking phenomena, and the tensile fracture surfaces exhibit typical brittle fracture characteristics. The tensile fracture surface of the recycled Al-7Si-0.3Mg-1Fe alloy without the addition of Al-5Ti-2B is shown in Figure 11a. From the figure, it can be seen that the fracture surface morphology is primarily river-like, exhibiting a cleavage fracture surface morphology. From the fracture path analysis, the fracture surface is an intergranular fracture, belonging to brittle fracture. This is because β-Fe grows along the eutectic Si boundary, penetrating the entire eutectic structure and severing the alloy matrix. Under external force, fractures first occur, forming crack sources, and secondary cracks are prone to occur during crack propagation, accelerating material fracture. As the addition of Al-5Ti-2B increases, quasi-cleavage planes appear in Figure 11b–e, and the river-like patterns become shorter. Notably, when the Al-5Ti-2B addition reaches 1%, both the β-Al₅FeSi intermetallic phases and eutectic silicon undergo significant refinement, effectively mitigating stress concentration effects. This microstructural modification alters the fracture mechanism, causing crack propagation to circumvent the refined eutectic Si particles and leading to characteristic intergranular fracture behavior [29,30], as shown in Figure 11e. The fracture surface morphology exhibits minor tearing ridges (i.e., ductile pits (circled in the figure)), but the fracture mode remains predominantly brittle. When the Al-5Ti-2B addition reaches 1.25%, as shown in Figure 11f, the fracture surface exhibits extensive micro-porosity (indicated by arrows in the figure). Figure 11g is an enlarged view of the micro-porosity in Figure 11f, and the EDS spectrum at point A in Figure 11g is shown in Figure 11h. The Fe content at this point reaches 32.16%, and the O element is likely introduced due to improper sample handling, indicating that the microstructure at point A corresponds to the β-Fe phase. The formation of shrinkage porosity is attributed to Fe enrichment, which has a segregation effect on Al. Based on the research results in Reference [2], it is known that under external force, shrinkage porosity leads to crack initiation and propagation, resulting in a decline in the mechanical properties of the alloy.

4. Conclusions

(1)

In the recycled Al-7Si-0.3Mg-1Fe alloy, Fe primarily exists in the form of the β-Al₅SiFe phase, appearing as coarse needle-like structures distributed around the eutectic Si, with some penetrating the entire Al-Si eutectic structure and disrupting the matrix structure. As the addition of Al-5Ti-2B increases, the α-Al primary phase and eutectic Si in the matrix alloy structure become refined, and the structure tends to homogenize. When the addition of Al-5Ti-2B reaches 1%, the alloy properties reach their optimal state, with the average grain size decreasing to 16.42 μm, tensile strength increasing to 149.4 MPa, and the elongation increasing to 4.3%. When the addition of Al-5Ti-2B continues to increase, the eutectic Si aggregates, the β-Al₅SiFe phase coarsens, and micro-porosity appears in the alloy structure, leading to a decline in the alloy’s performance.

(2)

Iron (Fe) elements adhere to the eutectic silicon (Si), forming the β-Al₅FeSi phase. Magnesium (Mg) enriches at grain boundaries, inhibiting the growth of coarse Fe phases. Simultaneously, Mg reacts with silicon to form Mg₂Si, achieving solid solution strengthening and enhancing the alloy’s performance. Al-5Ti-2B serves as a grain refiner, with its effectiveness primarily dependent on TiB₂ and Al₃Ti, which act as nuclei for the heterogeneous nucleation of α-Al. When the addition of Al-5Ti-2B is 1%, the refinement effect on the primary α-Al phase and eutectic Si in the matrix structure is significant, but its effect on altering the morphology of the β-Al₅FeSi phase is not prominent. However, after the matrix structure is refined, it inhibits the growth and distribution of the β-Al₅FeSi phase, thereby mitigating the adverse effects of the β-Al₅FeSi phase on the matrix alloy. When the addition of Al-5Ti-2B increases to 1.25%, eutectic Si aggregates, and the β-Al₅FeSi phase regrows into coarse needle-like structures, severely impairing the alloy’s mechanical properties. Additionally, iron-rich intermetallic compounds may serve as nucleation sites for pores, increasing the porosity and leading to casting defects. Therefore, when improving the properties of iron-rich regenerated Al-Si-Mg alloys through the refinement effect of Al-5Ti-2B, the optimal addition amount is 1%.

(3)

Based on the relevant literature and experimental results, it is known that due to the difference in thermal expansion coefficients at the interface between the β-Al₅FeSi phase and the matrix, microcracks are prone to form. Under external force, these microcracks propagate rapidly along the interface between the β-Al₅FeSi phase and the α-Al matrix. Meanwhile, unmodified eutectic Si particles act as obstacles during crack propagation, causing crack deflection and branching, which severely impair the mechanical properties of the alloy. The presence of Fe elements affects the distribution of eutectic Si, thereby reducing the alloy’s casting performance. In tensile tests, the fracture surface shows no obvious yield or necking phenomena, with brittle fracture being the primary mode of failure.

5. Outlook

(1)

Interfacial Behavior of TiB₂/Al and Fe Segregation Mechanisms

The current understanding confirms that Fe impurities in recycled aluminum (typically forming β-Al₅FeSi intermetallics) detrimentally affect the heterogeneous nucleation efficiency of TiB₂ particles in Al-5Ti-2B refiners. To address this fundamental issue, future investigations should prioritize the following:

(i)

In situ TEM characterization of Fe segregation kinetics at TiB₂/Al interfaces, particularly focusing on the atomic-scale interaction between Fe and the (0001)TiB₂ nucleation substrate;

(ii)

Quantitative assessment of nucleation potency deterioration through interfacial energy measurements (e.g., AFM-based nanomechanical testing), establishing a predictive model correlating Fe concentration with undercooling requirements.

Such studies would provide critical insights into impurity-mediated nucleation inhibition mechanisms.

(2)

Rare-Earth-Enhanced Impurity Tolerance in Recycled Al-Si-Mg Alloys

A promising research direction involves developing RE-modified grain refinement systems, as follows:

(i)

The synergistic effects of Ce/La multi-component additions could mitigate Si poisoning by competitively binding with Ti (as verified by ab initio MD simulations of Ti-Si-RE cluster stability);

(ii)

Advanced correlative microscopy (combining APT for nanoscale Mg mapping and synchrotron μ-XRD for phase identification) is essential to quantify Mg partitioning behavior in Fe-containing systems;

(iii)

Phase-field modeling incorporating experimental inputs from high-throughput XRD-EBSD could elucidate the stabilization mechanism of α-Al(Fe,Mn)Si phases during solidification.

This integrated approach may redefine impurity management strategies in recycled alloys.

(3)

Thermomechanical Processing Optimization

Key knowledge gaps remain regarding the following:

(i)

The critical strain threshold for Fe-intermetallic fragmentation during multi-pass rolling, which could be determined through high-temperature EBSD coupled with crystal plasticity finite element modeling (CPFEM);

(ii)

The role of stored energy distribution (characterized by 3D-EBSD and TEM dislocation analyses) in governing precipitation kinetics during post-deformation aging.

Establishing these process–structure–property relationships would enable microstructure-informed manufacturing of high-performance recycled alloys.