Enhancement of tensile properties of cast Al-Si-Fe alloys by annealing and cold rolling

Abstract

The recycling of aluminum alloy is limited by ductility reduction due to the accumulation of iron. In the literature, recycling of aluminum alloys up to one percent iron is mostly reported. In this work, annealed and cold rolled cast Al-9Si-2Fe alloy (wt.%), prepared from commercial aluminium, Al-50Si master alloy, and ferrosilicon, is reported. The as-cast microstructure comprising α-Al, elongated eutectic Si, and β-Al₉Fe₂Si₂ was refined by simultaneous annealing (535 °C) and step-wise cold rolling (55–80%). The decrease in the largest length, average length, and aspect ratio of eutectic Si and β-Al₉Fe₂Si₂ after 80% cold rolling resulted in a tensile strength of 170 MPa and percentage strain of 7.3% despite high iron content, and is comparable to existing literature.

Keywords

cold rolled ferrosilicon aspect ratio tensile strength

Introduction

Aluminium-Silicon-based alloys’ widespread application in automotive and aerospace industries is owing to their high specific strength, excellent castability, and high wear resistance. The above properties are adversely affected when these alloys are produced from recycled aluminium scraps due to iron accumulation during multiple recycling cycles.^1–3 The iron concentrations in Al-Si alloys increase with an increase in the number of recycling cycles, which could be greater than 1.5 wt.% in recycled Al-Si alloy.^2,4 Iron forms equilibrium phases, such as β-Al₉Fe₂Si₂ (hereafter β, monoclinic) and α-Al₈Fe₂Si (hexagonal) in Al-Si alloys. Out of them, the hard, brittle, and needle-like β adversely affects the mechanical properties, castability, and fluidity of the Al-Si alloys.^1–3 Dongfu Song et al.⁵ thermodynamically calculated that an increase in iron content increased the formation temperature of the Fe-rich phase (β). The length and aspect ratio of the β phase increased with an increase in iron content in Al-Si alloys.⁶ In industry, the reduction of iron concentration is achieved by diluting the recycled Al-Si alloy with pure aluminium, which reduces the economic benefits of the recycling process. Recycled Al-Si alloys can be used undiluted, i.e., without adding pure aluminium, if the harmful effect of the β phase can be eliminated.² Few approaches used to negate the harmful effect of β phase are: (i) modification by adding Mn to promote formation of some other phases,^7,8 (ii) reducing β size by a suitable heat treatment,⁹ (iii) reducing the size of β by plastic deformation.^2,10–13

Renato Baldan et al.⁷ studied the effect of Mn on microstructure and mechanical properties of the cast Al-9Si-0.8Fe alloy and concluded that the addition of 0.4 wt.% Mn improved the ultimate tensile strength from 179 to 180 MPa (0.55% improvement) and elongation (%) from 4.7% to 5.8% (23% improvement) due to conversion of β to α-Al₁₅(Fe, Mn)₃Si₂ having polygonal and Chinese script morphology. They also found that the addition of 0.7 wt.% Mn to Al-9Si-0.8Fe alloy improved ultimate tensile strength and elongation (%) to 191 MPa and 5.6% respectively, but the β appeared in the microstructure along with α-Al₁₅(Fe, Mn)₃Si₂ polyhedral phases. Dongfu Song et al.⁸ also correlated the microstructure-properties relationship of cast Al-7Si-1.2Fe-xMn (x = 0.24, 0.48, 0.72, 0.96, and 1.20). They reported slight changes in strength and significant improvement in elongation with an increase in Mn content (up to 0.73%), but a gradual decrease of tensile properties occurred when Mn content was more than 0.73% due to an increase in the number and size of polyhedral α-Al₁₅(Fe, Mn)₃Si₂ phases. L. Narayanan et al.⁹ reported in their study that tensile strength and % elongation of cast Al-6Si-3Cu-1Fe-0.3Mg improved from 178 MPa to 263 MPa (i.e., 48% improvement) and 0.6 to 0.78% (30%), respectively, after non-equilibrium heat treatment due to reduction of size and volume fraction of β. DaeHan Kim et al.¹⁰ performed deformation semisolid extrusion (D SSE) at 555 °C on Al-4.47Si-1Cu-0.35Mg-1Fe alloy to study the microstructure-properties relationship. Tensile properties of cast alloy improved significantly after semisolid extrusion due to uniform fragmentation and dispersion of β. They concluded that fragmented β withstands the load until a certain strain, and smaller particles are resistant to cracking. The β phase, which acted as nucleation sites for microcracks in cast alloy, contributed to void formation after D-SSE and improved tensile properties. JaeHwang Kim and DaeHwan Kim¹¹ and DaeHan Kim et al.¹² studied the effect of rolling (40%, 60%, and 80%) on tensile properties of Al-7.4Si-1Fe-0.27Mg alloy. They concluded that reducing the average size and largest size of Fe intermetallic compounds below a critical value after 80% rolling helped to improve yield strength, ultimate tensile strength, and ductility by 238%, 155%, and 144% respectively, as compared to the cast alloy. Shino Sakow et al.² studied the effect of hot extrusion (450 °C) on cast Al-12 mass% Si-2 mass% Fe alloy and reported that tensile strength improved from 60 MPa to 160 MPa (i.e., 166% improvement) and % strain improved from 2 to 43% (i.e., 2050% improvement) after hot extrusion due to microstructure refinement. Similarly, Behzad Asharsheikh et al.¹³ performed friction stir processing (FSP) on A356-1 wt.% Fe alloy and reported that the tensile strength and ductility of cast A356-1 wt.% Fe alloy increased by 23% and 147%, respectively, after FSP, due to the presence of small and fragmented eutectic Si and β phases.

As mentioned, the iron content in recycled Al-Si alloy could be more than 1.5 wt.% and the size of β increased with an increase in iron content.^4,6 But the effect of Mn is reduced when the iron concentration exceeds 1.2 wt.% due to the formation of sludge in the Al-Si-Fe-based alloy.^14,15 To summarize, heat treatment did not improve the tensile properties significantly. Although plastic deformation improved the tensile properties of cast Al-Si-Fe based alloy significantly, there is limited literature on the effect of plastic deformation on cast Al-Si alloy containing more than 1.5 wt.% Fe. Shino Sakow et al.² also attempted 77% thickness reduction of Al-12mass% Si -2mass% Fe by 2-pass cold-rolling, but the sample was torn into two parts due to the presence of coarse and brittle β phases. Therefore, the primary objectives of the work are (a) to prepare a higher iron containing i.e., (2 wt.% Fe) Al-9Si-Fe alloy using ferrosilicon as one of the starting raw materials, and (b) obtain enhanced tensile properties via refinement of phases by simultaneous annealing and cold rolling to improve recyclability of this alloy despite having high iron content (2 wt. %).

Experimental details

Gravity-cast Al-9Si-2Fe alloys were prepared by melting a proportionate amount of commercially pure Al (99.9% purity), 0.85–1.7 mm-sized ferrosilicon (69 wt.% Si–29.2 wt.% Fe), and Al-50Si master alloy using an induction furnace. The melt was held at 960 °C for the complete dissolution of ferrosilicon and Al-50Si master alloy. After 5 min of degassing with hexachloroethane (C₂Cl₆), the pouring of the melt was done at 880 °C into a preheated (200 °C) split-type square mold (120 mm × 120 mm × 7 mm). Optical emission spectroscopy (OES) (SPECTRO) analysis (Table 1) revealed that the alloy contains 8.63 wt.% silicon and 2.22 wt.% iron. The as-cast samples (S1) were subjected to 55% and 80% thickness reduction, respectively, by a combination of annealing (at 535 ± 3 °C for 2 h) and multi-pass cold rolling of 2.7% reduction in each pass in a two-high rolling mill at 12 revolutions per minute. The annealing was done to homogenize the as-cast samples as well as to remove the accumulated strain during multi-pass rolling. The reduction in each pass was intentionally kept low to prevent any localized strain accumulation and subsequent crack formation. Initially, S1 samples were annealed (S2) and their thickness was reduced to 30%. Subsequently, one of the 30% rolled samples was annealed and rolled to 55% thickness (S3) of the original. An S3 sample was annealed (S4) and was decreased to 80% thickness (S5).

Table 1.

OES analysis of alloy.

Al (wt.%)	Si (wt.%)	Fe (wt.%)	Mg (wt.%)	Cu (wt.%)	Ni (wt.%)
89.05 ± 0.06	8.63 ± 0.03	2.22 ± 0.02	< 0.05	< 0.05	< 0.05

The metallography and characterization of cast, annealed, and cold-rolled samples were done using an optical microscope (LEICA DMI3000 M), an X-ray diffractometer (XRD) (Bruker D8 Advance x-ray diffractometer, Cu Kα radiation with λ= 0.1540 nm, and step size of 0.02 °), a field emission scanning electron microscope (FESEM) (Zeiss, Merlin Compact) with an energy dispersive spectroscope (EDS) (OXFORD Instrument X-Max^N, 51-XMX1004), and electron backscatter diffraction (EBSD) (OXFORD NORDLYS NL04-2550-20). The EBSD analysis was carried out at a tilt angle of 70 ° and a 0.15 μm step size. The hardness test (20 measurements at each condition) was done using a Vickers microhardness tester (Zwick/Roell ZHVμ) at 500 gf load and a dwell time of 10 s. The position and the dimensions of the tensile sample are shown in Figure 1. The tensile samples were tested according to ASTM E8 subsize standard with a gage length of 25 mm at 0.3 mm/min crosshead speed and room temperature using a 50 kN universal testing machine (UTM) (SHIMADZU AGX-V). The thicknesses of the cast, 55% rolled, and 80% rolled alloys specimens were 5, 2.4, and 1.3 mm, respectively. The fracture surfaces were further characterized by FESEM and EDS. The equilibrium solidification sequence of the alloy was predicted using Thermocalc 2022b ^®. The phase quantification was done using ImageJ software (ImageJ 1.53t).

Figure 1.

Schematic of experimental details with sample positions and dimensions of tensile sample in the inset.

Results and discussion

Figure 2(a) represents the equilibrium solidification curve of Al-9Si-2Fe alloy calculated using Thermocalc software. The precipitation sequences are as follows: (1) L→ β-Al₉Fe₂Si₂ + L (at 620 °C), (2) L → β-Al₉Fe₂Si₂ + α-Al + L (at 602 °C), and (3) L→ β-Al₉Fe₂Si₂ + α-Al + eutectic Si + L (at 576 °C). Some β phase forms prior to the binary eutectic reaction and has a higher formation temperature than α-Al in Al-9Si-2Fe alloy.^1,16 The XRD pattern of S1 (Figure 2(b)) confirmed the presence of α-Al (Al), Si, and β phases. No new phase was found after annealing in S2 (Figure 2(b)). The S3, S4, and S5 alloys also contain α-Al (Al), Si, and β phases (Figure 2(c)).

Figure 2.

(a) The equilibrium solidification curve of Al-9Si-2Fe alloy calculated by Thermocalc software and (b, c) XRD pattern of Al-9Si-2Fe alloy at different conditions.

The S1 alloy microstructure reveals α-Al, β, and eutectic Si phases ( Figures 3(a) and 4(a)). The EDS analysis of all phases in the Al-Si-Fe alloy in all conditions is shown in Figure 4. Due to thin plates of eutectic Si and β, the accurate atomic fraction of both phases could not be quantified in EDS.^17–20 The S2 alloy also contains α-Al, β, and eutectic Si phases (Figure 3(b) and 4(b)). The optical micrographs (Figures 3(c)-(e)) and FESEM micrograph (Figure 4(c)-(e)) of the S3, S4, and S5 alloy show fine fragmented β and rounded eutectic Si distributed uniformly in α-Al matrix. Figure 5 shows the length distribution curve of β in each condition. The β phase size ranges from 4.8 to 350 μm (Figure 5(a)) in the S1 alloy. Around 43% of the β phases have a length greater than 100 μm (Figure 5(a)). The β phase starts forming at high temperatures and grows until a significant size of Al-Si eutectic cell is attained.²¹ Figure 5(e) shows that the length of β falls within the range of 0.67 to 18 μm in S5 alloy. The average length and aspect ratio of the β and average length, aspect ratio, and surface shape factor (ϕ) of the eutectic Si phase are given in Table 2, where ϕ is calculated as²²

φ = \frac{4 π A}{P^{2}}

(1)

where A and P stand for the area and perimeter of eutectic Si phases, respectively.

Figure 3.

Optical micrographs of (a) as-cast (S1), (b) as-cast and annealed (S2), (c) 55% rolled (S3), (d) 55% rolled and annealed (S4), and (e) 80% rolled (S5) Al-9Si-2Fe alloy.

Figure 4.

FESEM micrographs and corresponding EDS spot scan of all phases of (a) as-cast (S1), (b) as-cast and annealed (S2), (c) 55% rolled (S3), (d) 55% rolled and annealed (S4), and (e) 80% rolled (S5) Al-9Si-2Fe alloy. (Where solid, round dot, and square dot arrow mark indicates α-Al, eutectic Si, and β-Al₉Fe₂Si₂ (β) phases, respectively).

Figure 5.

Length distribution of β-Al₉Fe₂Si₂ (β) phase in (a) as-cast (S1), (b) as-cast and annealed (S2), (c) 55% rolled (S3), (d) 55% rolled and annealed (S4), and (e) 80% rolled (S5) Al-9Si-2Fe alloy.

Table 2.

Characteristics of β and eutectic Si in Al-9Si-2Fe alloys.

Conditions	Largest length of β (μm)	Average length of β (μm)	Aspect ratio of β	Average length of eutectic Si (μm)	Aspect ratio of Si	Surface shape factor (ϕ) of eutectic Si
as-cast (S1)	324.1	79.0	44.0	9.1	9.2	0.41
as-cast and annealed (S2)	320.3	65.1	38.3	7.5	6.8	0.53
55% rolled (S3)	26.5	7.5	6.8	4.2	4.3	0.65
55% rolled and annealed (S4)	28.1	7.3	6.6	4.9	4.8	0.60
80% rolled (S5)	17.3	3.8	4.1	3.4	3.3	0.72

Figure 6 shows the schematic diagram of the effect of heat treatment and rolling on the microstructure of the cast Al-9Si-2Fe alloy. After annealing of the as-cast alloy, the eutectic Si was spherodized, and heat treatment had a negligible effect on the size of the β phase. After 55% rolling, both eutectic Si was spherodized and β was fragmented. Annealing of 55% rolled alloy has less effect on the size of β, but eutectic Si was slightly coarsened. After 80% rolling, the β phase was completely fragmented, and eutectic Si was spherodized, and both phases were distributed uniformly in the α-Al matrix.

Figure 6.

Schematic diagram showing the effect of annealing and rolling on the microstructure of Al 9Si-2Fe alloy ((S1) as-cast, (S2) as-cast and annealed, (S3) 55% rolled, (S4) 55% rolled and annealed, and (S5) 80% rolled alloy).

Figure 7(a) shows the microhardness of Al-9Si-2Fe alloys under different conditions. The average hardness of as-cast alloy (S1) is 72 HV_0.5 due to a higher volume fraction of hard eutectic Si and β phases.⁶ The difference in thermal expansion between aluminium and the second phases results in high interfacial dislocation density and thus contributes to increased average hardness of the as-cast alloy.²³ A decrease in stored strain energy value after annealing²³ (535 °C for 2 h) reduces the average hardness to 50 HV_0.5. Subsequent rolling (55% reduction) and annealing changed the hardness of S4 alloy to 44 HV_0.5. The hardness of S5 alloy is 63 HV_0.5. Figure 7(b) shows the engineering stress-strain curves of Al-9Si-2Fe alloys at different conditions and the corresponding average values of tensile properties are shown in Table 3. The annealed cast alloy (S2) shows slight improvement in ultimate tensile strength (UTS) and percentage strain as compared to the as-cast alloy (S1). This could be due to decreased aspect ratio and average size of eutectic Si and improved surface shape factor (ϕ) of eutectic Si. Annealing alters the morphology of eutectic Si from sharp-edged acicular morphology to rounded coral-like morphology, which reduces the stress concentration (marked in red dashed circle in Figure 3(a), (b)).²⁴ The annealing has less contribution towards the reduction of the size of β phases (Figures 3(b), 4(b), and 5(b)) in the cast alloy. So, the percentage strain of the alloy did not improve significantly due to the presence of coarser β with a maximum size of 320 μm. The improvement of tensile properties after 55% rolling (S3) and 80% rolling (S5) is due to the presence of fine and rounded phases in the rolled alloy (Figure 3(c), (e) and 4(c), (e)), which has more resistance to cracking compared to the coarser phases.¹² The decrease in average length, largest length, and aspect ratio (Table 2) of uniformly distributed β and eutectic Si, along with increased surface shape factor (ϕ) of eutectic Si (Table 2), also resulted in better tensile properties in the rolled alloys (55% and 80%). The reason for the decrease in strength and improved percentage strain in S4 alloy can be attributed to recrystallization and grain growth at higher annealing temperature, causing a significant reduction in the dislocation density.²⁵ Moreover, annealing reduces the residual stress developed during cold rolling, which further improves the ductility.²⁶ The increase in aspect ratio and average size of silicon and decrease in spherical factor after annealing (Table 2) reduced the tensile strength of S4 compared to S3.

Figure 7.

(a) The average micro-hardness of Al-9Si-2Fe alloy at different conditions (b) Tensile properties of Al-9Si-2Fe alloy at different conditions (c) EBSD map of 80% rolled Al-9Si-2Fe alloy (S5), where red, yellow, green colour represents α-Al, eutectic Si, and β-Al₉Fe₂Si₂ (β) phases, respectively and black lines represent grain boundary (d) Optical micrograph of cast Al-9Si-2Fe alloy (S1), and (e) Optical micrograph of 80% rolled Al-9Si-2Fe alloy (S5).

Table 3.

Tensile properties of Al-9Si-2Fe alloys with different conditions.

Conditions	Yield strength (YS) (MPa)	Ultimate tensile strength (UTS) (MPa)	Strain (%)
as-cast (S1)	32 ± 4.5	69 ± 6.6	1.6 ± 0.4
as-cast and annealed (S2)	29 ± 1.8	78 ± 4.1	2.7 ± 0.4
55% rolled (S3)	75 ± 7.8	148 ± 3.6	4.7 ± 1.2
55% rolled and annealed (S4)	31 ± 3.5	74 ± 7.1	25.8 ± 1.2
80% rolled (S5)	108 ± 4.1	170 ± 9.1	7.3 ± 1.8

The 80% rolled alloy (S5) has a better combination of strength and percentage strain than other conditions (Table 3) because of the presence of uniformly distributed fine and fragmented β phases and rounded eutectic Si compared to other conditions (Figures 3, 4, and Table 2). An estimate of yield strength (σ_0.2) for 80% rolled alloy (S5) is given as^27,28

σ_{0.2} = σ_{0} + σ_{G B} + Δ σ_{S S} + σ_{D I S}

(2)

Where σ₀ is lattice friction stress, i.e., 13 MPa for pure aluminium alloy.²⁷ Grain boundary strengthening (σ_GB) is calculated using Hall-Petch formula,²⁷

σ_{G B} = \frac{K}{d^{\frac{1}{2}}}

(3)

Where K is Hall-Petch constant, i.e., 74 MPa μ^1/2 for pure aluminium²⁷ and d is the grain size.²⁷ The calculated average grain size (d) for 80% rolled alloy (S5) using the EBSD map (Figure 7(c)) is 3.3 μm. The calculated maximum solid solution strengthening (Δσ_SS) due to silicon and iron is 16 and 0 MPa, respectively.²⁹ The strength contribution due to dislocation strengthening (σ_DIS) is calculated by the Taylor formula,²⁷

σ_{d i s} = α M G b ρ^{\frac{1}{2}}

(4)

Where α is 0.24 for FCC materials, M is Taylor factor (3.06), b is Burger's vector (0.256 nm), G is the shear modulus of the aluminium matrix (26 GPa), and ρ is the dislocation density.²⁷ Dislocation density (ρ) is calculated as,³⁰

ρ = \frac{2 \sqrt{3} ϵ}{D b}

(5)

Where є is the heterogeneous strain and D is the crystallite size. The heterogeneous strain (є) and crystallite size (D) are calculated from the XRD pattern using the Williamson-Hall equation³⁰

Δ 2 θ c o s θ = \frac{0.9 λ}{D} + 2 ϵ s i n θ

(6)

Where Δ2θ, θ, and λ are full width at half maximum (FWHM) in radian, diffraction angle (radian), and the wavelength of X-rays (0.154 nm), respectively.³⁰ The crystallite size of α-Al in 80% rolled alloy and the heterogeneous strain of 80% rolled alloy were calculated as 44.56 nm and 0.00039, respectively. It is to be noted that XRD determines the size of coherent diffraction domains, which include both sub-grains and dislocation cells in plastically deformed materials. So, the crystallite size of α-Al measured by XRD is significantly less than the grain size of α-Al measured using EBSD.³¹ Combining equations (2) to (6), the calculated yield strength of 80% rolled alloys (S5) is 123 MPa. The difference in experimentally measured and calculated yield strength (Table 4) is due to the presence of porosity in the 80% rolled alloy (S5) (Figure 7(e)). As the elongated β phase is the first phase to solidify in this alloy, it caused difficulty filling the interdendritic region, resulting in the formation of pores and shrinkage defects^7,32 as shown in Figure 7(d). Due to rolling, the size of the pores can be minimized (Figure 7(e)). Another reason for the difference in experimentally measured and calculated yield strength in the rolled alloy is due to the consideration of maximum solid solubility of Si in the α-Al matrix, because some will form eutectic Si and some will form β.

Table 4.

The summary of the contribution of grain boundary strengthening (σ_GB), dislocation strengthening (σ_DIS), and solid solution strengthening (Δσ_SS) to calculated yield strength (YS).

Alloy	ρ (nm⁻²)	d (μm)	σ₀ +σ_GB (K d ^−1/2) (MPa)	Δσ_SS (MPa)	σ_DIS (MPa)	Calculated YS (MPa)	Experimentally Measured YS (MPa)
80% Rolled	0.00012	3.3	53.5	16	53.5	123	108 ± 4.1

The fractography studies were done to describe the tensile properties. The fracture surface of the cast alloy (S1) (Figure 8(a)) shows smooth surfaces of β, which appear as massive platelets (marked by solid yellow arrows), suggesting that the cast alloy has low ductility. The initiation and propagation of the crack⁸ (marked by the round dot yellow arrow) in Figure 8(a), through coarse β (marked by the solid yellow arrow), led to the failure of the cast alloy. The fracture surface of S2 alloy also shows a smooth surface of β (Figure 8(b)). The uniform distribution of small dimples on the fracture surface of S3, S4, and S5 alloys indicates ductile failure (Figure 8(c)-(e)) with fragmented β and rounded Si phases within the voids. The reasons for low ductility in cast alloys (S1) have been explained below. To find out the phases that act as crack initiation sites, researchers measured the hardness and elastic modulus of β, Si, and Al matrix phases by nanoindentation. The hardness of β, Si, and Al matrix is reported as 11.5, 13.3, and 2.1 GPa, respectively, whereas the elastic modulus of β, Si, and Al matrix is 196, 185, and 92 GPa, respectively.^33,34 So, the β and Si phases can be considered significantly harder compared to the matrix. The strain mismatch between the aluminium matrix and hard phases and the stress concentration of the hard phase and its nearby area causes crack initiation at the harder phase, depending on the size and morphology of the hard phases.³⁴ In this cast alloy, β controls fracture because the length (size) of β is greater than that of eutectic Si. Zaidao Li et al.³⁵ have also concluded that initiation and growth of cracks occur along with the hard phase, i.e., mainly with iron intermetallics in Al-7Si-3.4Cu-0.8Fe-0.5Mn-0.28Mg-0.01Sr alloys. Shota Kariya et al.³⁶ used equation (7) to study the effect of crack length on the tensile properties of the alloy. They found that the crack length was the same as the grain size. So, they considered the grain size as crack length and found that grain size (i.e., crack length) had an inverse relationship with ductility.

K = σ_{0} (π a)^{\frac{1}{2}}

(7)

Figure 8.

Fracture surfaces and corresponding EDS spot scan of phases of (a) as-cast (S1), (b) as-cast and annealed (S2), (c) 55% rolled (S3), (d) 55% rolled and annealed (S4), and (e) 80% rolled (S5) Al-9Si-2Fe alloy.

K is the stress intensity factor, σ₀ is the uniform stress around the crack, and a is the crack length. For brittle fracture, UTS is interpreted as critical fracture stress.³⁷ This relationship can also be applied to the present alloy system. Anton Bjurenstedt et al.³ concluded that the largest intermetallics initiated cracks in Al-8Si-0.6Fe-0.4Mg alloys. DaeHan Kim et al.¹² also concluded that the largest particles have a more significant role in controlling fracture. So, assuming the as-cast Al-9Si-2Fe (S1) to be perfectly brittle, and the largest β phase having a length of 324 μm is the crack initiator, the stress intensity factor is found to be 2.2 MPa m^−1/2 for the cast alloy (S1) by using equation 7. After 80% rolling, the improved ductility is due to a lower β phase length. Assuming equation 7 is still applicable for rolled alloy, the calculated critical crack length for 80% rolled alloy (S5) is 53 μm. But in S5 alloy, the largest β length is 17.3 μm, which is below the critical crack length calculated for this alloy. Due to the above-mentioned reasons, 80% rolled alloys (S5) have higher strength and ductility than cast alloys (S1). Figure 9 shows the comparison of the tensile properties achieved in this study with a few Al-Si-based alloys processed through different routes.^2,38,39 The result (Figure 9) shows that 80% cold rolling helped in a significant improvement of both the strength and ductility of cast Al-9Si-2Fe alloys (S1).

Figure 9.

Comparisons of tensile properties of Al-Si-based alloys processed through different routes^2,38,39

Conclusions

A high-iron containing gravity-cast Al-9Si-2Fe alloy (wt.%) was prepared using ferrosilicon as one of the starting raw materials. The enhanced tensile properties in the alloy were achieved through refinement of eutectic Si and β-Al₉Fe₂Si₂ (β) phases by simultaneous annealing and stepwise cold rolling. The 80% cold rolled alloy (S5) demonstrated the best combination of tensile strength (170 MPa), yield strength (108 MPa), and percentage strain (7.3%) owing to important factors like reduction of β phase below critical crack length, spheroidization of phases, and reduction of porosity.

Footnotes

Acknowledgements

The authors are thankful to Professor Gandham Phanikumar and Miss Arpita Priyadarshani Samal from the Department of Metallurgical and Materials Engineering, Indian Institute of Technology Madras, India for providing Thermocalc software data.

Author contribution(s)

Saswati Nanda: Conceptualization; Formal analysis; Investigation; Methodology; Validation; Writing – original draft; Writing – review & editing.

Partha Sarathi De: Conceptualization; Investigation; Supervision; Writing – review & editing.

Animesh Mandal: Conceptualization; Investigation; Supervision; Writing – review & editing.

Funding

The authors received no financial support for the research, authorship, and/or publication of this article.

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

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