Khalil Ganjehfard, Reza Taghiabadi, Mohammad Talafi Noghani, and Mohammad Hossein Ghoncheh, Tensile properties and hot tearing susceptibility of cast Al–Cu alloys containing excess Fe and Si, Int. J. Miner. Metall. Mater., 28(2021), No. 4, pp.718-728. https://dx.doi.org/10.1007/s12613-020-2039-7
Cite this article as:
Khalil Ganjehfard, Reza Taghiabadi, Mohammad Talafi Noghani, and Mohammad Hossein Ghoncheh, Tensile properties and hot tearing susceptibility of cast Al–Cu alloys containing excess Fe and Si, Int. J. Miner. Metall. Mater., 28(2021), No. 4, pp.718-728. https://dx.doi.org/10.1007/s12613-020-2039-7
Khalil Ganjehfard, Reza Taghiabadi, Mohammad Talafi Noghani, and Mohammad Hossein Ghoncheh, Tensile properties and hot tearing susceptibility of cast Al–Cu alloys containing excess Fe and Si, Int. J. Miner. Metall. Mater., 28(2021), No. 4, pp.718-728. https://dx.doi.org/10.1007/s12613-020-2039-7
Cite this article as:
Khalil Ganjehfard, Reza Taghiabadi, Mohammad Talafi Noghani, and Mohammad Hossein Ghoncheh, Tensile properties and hot tearing susceptibility of cast Al–Cu alloys containing excess Fe and Si, Int. J. Miner. Metall. Mater., 28(2021), No. 4, pp.718-728. https://dx.doi.org/10.1007/s12613-020-2039-7
This study was undertaken to investigate the tensile properties and hot tearing susceptibility of cast Al–Cu alloys containing excess Fe (up to 1.5wt%) and Si (up to 2.5wt%). According to the results, the optimum tensile properties and hot tearing resistance were achieved at Fe/Si mass ratio of 1, where the α-Fe phase was the dominant Fe compound. Increasing the Fe/Si mass ratio above unity increased the amounts of detrimental β-CuFe platelets in the microstructure, deteriorating the tensile properties and hot tearing resistance. Decreasing the mass ratio below unity increased the size and fraction of Si needles and micropores in the microstructure, also impairing the tensile properties and hot tearing resistance. The investigation of hot-torn surfaces revealed that the β-CuFe platelets disrupted the tear healing phenomenon by blocking interdendritic feeding channels, while the α-Fe intermetallics improved the hot tearing resistivity due to their compact morphology and high melting point.
Over the last decades, demands on using A206 Al–Cu alloys in automotive and aerospace industries have grown significantly. This growth is mostly due to the excellent strength and ductility and the good fracture toughness of these alloys, which are much higher than those of conventional 3xx casting alloys and close to those of some grades of ductile iron [1–4]. However, the poor casting characteristics, especially the long freezing range (505–643°C), poor fluidity, and high hot tearing susceptibility (HTS), significantly restrict the application of these alloys in the production of complex-shape castings such as wheels, cylinder heads, and engine blocks as well as castings that need to be poured in gravity-fed permanent molds [3–5]. The other issue with A206 Al–Cu alloys is the high variability of their mechanical properties, arising from their casting-related defects such as micropores and entrained double oxides [2].
The formation of harmful Fe-bearing compounds can also limit the application of Al–Cu casting alloys. Iron is the most common impurity found in Al–Cu alloys [6]. It has been shown that due to the sudden solid-solubility drop in Al (from about 1.8wt% at 655°C in liquid to 0.005wt% at 450°C in solid Al) [7–8], the presence of Fe promotes the formation of brittle intermetallics in the microstructure of Al–Cu alloys, where the most important intermetallics are plate-like β-Al7Cu2Fe (referred to as β-CuFe) and α-Al15(FeMn)3(SiCu)2 Chinese scripts (referred to as α-Fe) [6–11]. These intermetallics, especially the β-CuFe, serve as stress-risers/micro-crack initiators, deteriorating the mechanical properties. The presence of Cu further decreases (by up to fivefold) the Fe solid-solubility in Al. Therefore, very low upper limits of Fe, 0.15wt%, 0.1wt%, and 0.07wt%, are permitted in 206.0, 206.2, and A206.2 Al–Cu alloys, respectively, above which post-dendritic Fe-rich intermetallics are likely to form, due to the delay in the chemical reaction between rejected Fe and other solute elements [8,12].
Iron-rich intermetallics also negatively affect the fluidity and HTS, the two most important indexes indicating the alloys castability. According to references [13–14], the formation of Fe compounds, especially those with plate-like morphology, substantially decreases the fluidity by increasing the melt viscosity and/or blocking the interdendritic flow channels. The β-CuFe platelets also deteriorate the hot tearing resistance of the Al–Cu alloys [15–16]. Hot tearing is the nucleation and propagation of cracks at the last stage of solidification, where the small pockets of liquid isolated along the grain boundaries or triple junctions are torn up due to the low shear strength of the mushy-state alloy against thermal and solidification contraction. It is a serious defect that occurs during the solidification of alloys, especially for alloys with a wide freezing range [15–18]. The β-CuFe platelets have been shown to act as barriers within interdendritic channels and retard the liquid permeability through the dendrite arms associated with increasing the HTSs of A206-type Al–Cu alloys [15].
Therefore, to improve the mechanical properties and enhance the casting characteristics of Fe-containing Al–Cu alloys, the formation of plate-like β-CuFe intermetallics must be partially or completely suppressed via appropriate methods. The Fe content in alloys can be decreased below its maximum allowable content. However, this approach is cost-intensive and significantly limits the alloy application [19]. The other approach is to convert undesirable β-CuFe platelets to less-detrimental α-Fe particles using suitable modifier elements such as Si. Kamga et al. [9] showed that converting β-CuFe platelets to the Si-bearing Chinese-script α-Fe particles substantially improved the hot tearing resistance of B206 alloy. Liu et al. [8] found that the addition of Mn or Si to the composition of A206 alloy promoted the formation of α-Fe at the expense of β-Fe particles. Elgallad and Chen [4] observed that adding Si promoted the formation of the Al15(FeMnCu)3Si2 phase rather than the needlelike β-CuFe phase and decreased the HTS of 206-type alloys. Increasing the Si content can also substantially ameliorate the fluidity and reduce the solidification shrinkage of Al–Cu alloys, making the alloys less susceptible to hot tearing [10,20]. However, the addition of Si is likely to impair the mechanical properties of the A206 family of casting alloys. According to Han et al. [20], adding Si to the binary Al–4.5Cu alloy does not change the alloy strength but impairs its ductility. Zhang et al. [10] showed that the addition of Si up to 1.1wt% decreased the tensile properties of Al–5.0Cu–0.6Mn–1.2Fe alloy. They attributed this to an increase in the volume fraction of porosities and FeSi-rich compounds. Liu et al. [11], Kamga et al. [9], and Zhao et al. [6] showed that Si addition promoted the formation of less harmful α-Fe instead of detrimental β-Fe phase, enhancing the alloy tensile properties. According to Kamga et al. [9], a Si/Fe mass ratio of close to unity accompanied with lower Fe and Si contents ensures the best mechanical properties of the cast A206 alloy.
Therefore, to take advantage of Si addition on castability while reducing its negative impacts on the tensile properties of Fe-bearing Al–Cu alloys, controlling the contents of Si and Fe elements within an optimum range is important. However, to the best of our knowledge, few studies have been conducted so far to investigate the combined effect of Si and Fe on the mechanical properties and castability of Al–Cu alloys, most of which deal with low contents of these elements. This is the reason that the Si content in some Al–Cu alloys is over 2wt% and the Fe impurity content can exceed 0.5wt% [21]. Therefore, this study was undertaken to investigate the effect of Si (0.5wt% to 2.5wt%) on the interaction of the microstructures, tensile properties, and HTS of Al–Cu alloys (A206) containing different amounts of Fe impurity (0.5wt% to 1.5wt%) and explore the optimum Si/Fe mass ratio in terms of tensile properties and HTS.
2.
Experimental
The primary A206 alloy with various contents of constitutive elements was fabricated using high-pure Al (99.9wt%) and Al–50wt%Cu, Al–25wt%Mg, and Al–25wt%Mn master alloys. An optical emission spectrometer, Foundry-Master (WAS), was employed for the chemical analysis. The average compositions of base-sample ingots are presented in Table 1. Almost 1.5 kg of prepared ingots was melted in a fireclay–graphite crucible using an electrical resistance furnace (AZAR-VM2L1200) at (750 ± 2)°C. The surface of the melt was covered by Coverall 11TM flux to avoid the elements loss via oxidation. To investigate the effects of Fe and Si on the mechanical and casting characteristics of the target alloy, desired amounts of them were added to the melt through the flux-75Si and flux-75Fe capsulated master alloys. To effectively disperse elements within the melt, stirring was followed by a degassing process via N2-based degasser tablets. The melt was then held at (750 ± 2)°C to be homogenized and poured into a preheated (250°C) cast-iron tensile mold (ASTM B 557M-02a) (Fig. 1(a)). The chemical compositions of the fabricated specimens are given in Table 1.
Table
1.
Alloy codes and chemical compositions of experimental Al–Cu alloys
To measure the HTSs of the prepared samples, the constrained-rod casting (CRC) mold developed by Alcan Kingston Research and Development Center (AKRDC), Canada [22] was used (Fig. 1(b)). Cylindrical rods with different lengths were poured via a conventional sprue. Four 9.5 mm-diameter rods with nominal lengths of 51, 89, 127, and 165 mm (rods A, B, C, and D in Fig. 1(b)) were embedded as mold cavities and constrained between the sprue and four spherical cavities with 19 mm diameter. To minimize the effect of friction between the mold wall and the solidifying melt, the mold cavity was cleaned, covered with graphite, and preheated up to 200°C before each test.
The HTS was analyzed using visual inspection of the CRCs. As presented in Table 2, numerical quantities were attributed to the severity degrees of tears and the lengths of cylindrical rods in the CRC mold. The HTSs of alloys were quantified using the correlation given in Eq. (1).
where Ci is a numerical value referred to as the hot tear severity of rod i (i: A–D), while Li is another numerical value showing the corresponding rod length of rod i.
The results obtained by multiplying the mean value of tear severity by the numerical value of the corresponding rod were depicted in footprint charts, in which each axis was assigned a specific rod in the CRC mold. The dash region in each chart illustrates the susceptibility of the sample to hot tearing [23–24]. The footprint chart can easily show if hot tearing has taken place on a particular rod. The severity of the tear is also evident. The average of four measurements was reported as the final HTS value.
As-cast tensile samples following the dimensions given in ASTM B557M standard (Fig. 1(c)) were tested by a Zwick Roell-Z100 tensile-test machine at a constant load cell of 10 kN and a crosshead speed of 0.5 mm·min−1. The ultimate strength of each specimen was taken as a mean value of four test sets at each condition.
To evaluate the fluidity, a vacuum fluidity test setup was used according to Fig. 2. Fully skimmed and stirred melt was isothermally held and sucked into the glass tube under a predetermined pressure of about 26.7×103 Pa (200 mmHg). For each test, the distance traversed by the solidifying melt was measured to compare the data of the melt fluidity. The mean value of four data sets was reported as the ultimate fluidity length.
Fig.
2.
Layout of the apparatus used for the fluidity measurement (D—Inside diameter of glass tube).
To characterize the constituents and the microstructures of alloys, the specimens used were taken from as close as possible to hot tear location. Following the standard metallographic process, the samples were etched using a 2vol% HF reagent (2 mL HF, 98 mL distilled water) for 8 s to reveal the microstructures. The microstructures of the samples were observed using scanning electron microscopy (SEM, Tescan-Vega) coupled with energy dispersive spectroscopy (EDS) to analyze the desired phases. In terms of the torn surface analysis, samples were collected from the bars showing hot tear phenomenon, on the joint between bars and the sprue, illustrated in Fig. 3. For the case of incomplete breaking, specimens were intentionally broken to correctly collect the torn surface. The torn surface was analyzed using a Tescan-Vega scanning electron microscope. The volume fraction of microporosities (VP) was measured via the following equation:
Fig.
3.
Macrographs of the constrained rods (rod D) showing the cracking severities of the (a) base, (b) 1.0Si, (c) 2.5Si, (d) 1.0Si–1.5Fe, (e) 1.0Si–1.0Fe, and (f) 2.5Si–1.5Fe samples.
where Dt and Da are the theoretical and actual densities of the alloy, respectively.
The theoretical density (Dt) of the samples was calculated by the rule of mixtures from their nominal compositions. The actual densities (Da) of the alloys were measured based on Archimedes’ principle [25]. For this purpose, the weight of the sample was measured in air and after immersion in distilled water. The actual density was calculated using the following equation:
Da=Wa⋅ρwWa−Ww
(3)
where Wa and Ww are the weights of the sample in air and distilled water, respectively (measured by an electronic balance with an accuracy of ±0.0001 g); and ρw is the density of distilled water (~1 g/cm3).
3.
Results and discussion
3.1
Microstructural characterization
The effect of Si addition on the cast microstructure of the A206 alloy is shown in Fig. 4. Based on the EDS results presented in Fig. 5, the microstructures of 1.0Si and 2.5Si alloys mainly consisted of white θ-Al2Cu phase and light-gray Si particles, respectively, which are revealed in the α-Al dendritic matrix. A higher Si content led to a higher volume fraction of Si particles in the microstructure.
Fig.
4.
SEM microstructures of the (a) base, (b) 1.0Si, and (c) 2.5Si alloys. The main phases are shown in micrographs.
The cast microstructure of the A206 alloy containing 1.0wt% Si and 1.5wt% Fe is shown in Fig. 6(a). As seen, a synergy exists between the sluggish diffusion coefficient of Fe in the liquid Al (2.81×10−9 m2·s−1) and the low value of Fe partition coefficient (0.02) [7], and this promoted the segregation of the Fe atoms during solidification. Moreover, the deleterious effect of Cu on Fe solid-solubility in Al [11–12] encouraged the formation of the β-CuFe compound within the microstructure of the standard 206 alloy. The EDS analysis result of the β-CuFe compound is presented in Fig. 6(b).
Fig.
6.
(a) SEM microstructure of 1.0Si–1.5Fe alloy (Al7Cu2Fe is pointed out via arrows) and (b) EDS analysis result of the Al7Cu2Fe compound in (a)
Fig. 7 shows the as-cast microstructures of 1.0Si–1.5Fe, 1.0Si–1.0Fe, and 2.5Si–1.0Fe alloys, which are taken as representatives of alloys with Fe/Si mass ratios of more than unity, equal to unity, and less than unity, respectively. Regarding the presence of an appropriate amount of Mn and a relatively high Si content (1.0wt%), a considerable volume fraction of the β-CuFe compound in the microstructure of 1.0Si–1.5Fe alloy (Fe/Si mass ratio > 1) was substituted by Chinese-script or irregular blocky α-FeMn particles (Fig. 7(a)). When the volume fraction of the β-CuFe compound was higher, the formation of these phases at higher temperatures consumed the high fraction of Fe atoms, suppressing the precipitation of β-CuFe particles [11]. The equal weight percentages of Fe and Si (Fe/Si mass ratio = 1) resulted in the minimum volume fraction of the β-CuFe phase (Fig. 7(b)). A further increase in the Si content or decrease in the Fe/Si mass ratio below unity caused a high volume fraction of interdendritic Si-rich eutectic and micropores in the microstructure (Fig. 7(c)). The EDS analysis results of different phases marked in Fig. 7 are given in Fig. 8.
Fig.
7.
SEM microstructures of (a) 1.0Si–1.5Fe, (b) 1.0Si–1.0Fe, and (c) 2.5Si–1.0Fe alloys.
The effect of Si addition on the tensile properties of A206 alloy is presented in Table 3. As seen, adding Si up to 1.0wt% improved the base-alloy tensile properties, where the fracture strength and elongation percentage of 1.0Si alloy were approximately 38% and 20% higher than those in the base sample, respectively. The improvement can be attributed to the fine distribution of hard Si particles (> KHN 1100) [26] within the matrix (Fig. 4(b)), which act as small-sized plate-like obstacles against the dislocations mobility [27]. Furthermore, due to an enhancement in the liquid fluidity, which will be discussed later, Si addition could substantially reduce the volume fraction of shrinkage micropores (Fig. 4 and Table 3) in the microstructure. This not only increased the load-bearing capability of the tensile specimens but also retarded the initiation and propagation of microcracks. Further addition of Si (higher than 1.0wt%) increased the size and volume fraction of hard Si needles (Fig. 4(c)) in the microstructure, thereby slightly impairing the alloy strength and ductility. In agreement with Qian et al. [28], during the tensile loading, a high level of stress was developed around the coarse Si needles due to their plate-like morphology and the high plasticity mismatch with the matrix. If the amplitude of accumulated stress exceeds the strength of Si needles on their faceted interface with the matrix, microcracks are likely to initiate through the nucleation and coalescence of interfacial microvoids [29].
Table
3.
Tensile properties and porosity contents of experimental alloys
The combined effect of Fe (0.5wt%, 1.0wt%, and 1.5wt%) and Si (1.0wt% and 2.5wt%) on the tensile properties of A206 alloy is also presented in Table 3. The maximum fracture strength in 1.0wt% Si-containing alloys belongs to the 1.0Si–1.0Fe sample, where α-FeMn is the dominant Fe-bearing phase (Fig. 7(b)). Compared with plate-like β-CuFe particles, α-compounds are less prone to strain accumulation due to their compacted morphology [11] and jogged interface with the matrix [30] that brings less vulnerability to the interfacial cracking during tensile loading. Therefore, in agreement with previous findings [30–31], the presence of hard α-FeMn compounds is likely to improve the strength in such a condition that they can be finely distributed within the matrix. Increasing the volume fraction of α-FeMn compounds, however, decreased the alloy ductility (Table 3). Increasing the Fe/Si mass ratio encouraged the formation of detrimental β-CuFe compounds along with α-FeMn in the microstructure (Fig. 7(a)). Therefore, the decreased tensile properties of the 1.0Si–1.5Fe sample (Fe/Si mass ratio > 1.0) (Table 3) can be explained by the formation of β-platelets in its microstructure.
Increasing the Fe content in 2.5wt% Si-containing alloys resulted in continuous improvement in fracture strength (Table 3). Due to the Fe/Si mass ratio being less than one, post-dendritic α-FeMn is likely to be the dominant Fe-rich phase in the microstructures of 2.5Si–xFe (x = 0.5–1.5) alloys, in which its formation improves the tensile properties. However, the segregation of excess Si, as well as Cu, into the remaining interdendritic liquid promoted the formation of low-melting-point ternary Al–Si–Cu eutectic phases at the last stage of solidification when the majority of the casting was already solidified. These regions are difficult to be fed by the remaining liquid [32]. Therefore, after the completion of solidification, considerable amounts of micropores were formed in the microstructure (Fig. 7(c), decreasing the tensile properties, where the fracture strength and fracture strain of 2.5Si–xFe alloys were lower than those of 1.0Si–xFe samples (Table 3).
3.2
Fluidity and hot tearing susceptibility
Aluminum–copper alloys are very prone to hot tear formation, since they experience a wide range of solidification in which the solidifying alloy is exposed to the vulnerable mushy state of hot tearing for a long time [4,11,15]. The effects of Si addition on the fluidity and HTS index of A206 alloy are presented in Fig. 9. The positive impact of Si on the HTS of A206 alloy is also evident from the footprint charts (Fig. 10), where the HTS of each alloy is assessed by the chart area. Moreover, according to each footprint chart, the occurrence severity of hot tearing on a particular rod is also distinguishable. Adding Si substantially suppressed the HTS due to an increase in fluidity (Fig. 9), mediated by the high latent heat of Si (4.5 times of that of Al) and the narrowing of the mushy zone [33]. Silicon also reduced the shrinkage during solidification because of its expansion upon solidification and formation of a higher volume fraction of the eutectic phase at the end of solidification [34].
Fig.
9.
Effects of Si content on HTS and fluidity length of A206 alloys.
The SEM results of the hot-torn surfaces of the base and 2.5Si samples are shown in Fig. 11. In conformity with the HTS results (Figs. 9 and 10), the severe tearing, rough dendritic nature of the fracture surface, and the formation of spikes on the tip of secondary dendrite arms suggest that interdendritic separation occurred in the presence of a liquid film at the last stage of solidification, where the amount and/or fluidity of the remaining liquid were probably not high enough to properly cover the dendrites surfaces or heal the hot tear microcracks. Spikes are generally considered as evidence of solid bridging between the growing dendrites and their further separation when they are pulled apart by temperature-induced mechanical stresses [35–36].
Fig.
11.
SEM micrographs showing hot-torn fracture surfaces of the (a, b) base and (c, d) 2.5Si alloys.
The high HTS of the base A206 alloy can be explained by its long freezing range. This property prolongs the period in which the alloy is susceptible to hot tearing; reduces the feeding ability of interdendritic liquid as a result of increased dendrites tortuosity; promotes the formation of micro-shrinkage cavities, which further decreases the alloy strength.
As seen previously, adding Si can substantially improve the hot tearing resistance of A206 alloy (Figs. 9 and 10). The hot tear morphology of the 2.5Si sample is shown in Figs. 11(c) and 11(d). In contrast to the torn surface of the base sample (Figs. 11(a) and 11(b)), the smooth and bumpy appearance of the fracture surface of the 2.5Si alloy demonstrates that sufficient amounts of interdendritic liquid were present at the initiation of hot tearing, which flowed within the interdendritic channels, and healed the interdendritic voids and cracks in the process referred to as “self-healing” [24].
Fig. 12 reveals the influence of Fe on the fluidities and HTS indexes of 1.0Si and 2.5Si alloys. As seen, irrespective of the Si content, adding Fe up to 1.5wt% slightly improved the fluidity of 2.5Si–xFe alloys, which can be attributed to the decreased alloy freezing range [37]. In 1.0Si–xFe alloys, maximum fluidity and consequently less HTS occurred at the addition of 1.0wt% Fe, where the Fe/Si mass ratio is unity. The addition of 0.5wt% Fe (Fe/Si mass ratio = 0.5) and 1.0wt% Fe (Fe/Si mass ratio = 1) to the 1.0Si alloy enhanced the hot tearing resistance; however, further addition of Fe decreased the hot tearing resistivity due to microporosity formation and coalescence. The addition of 0.5wt%, 1.0wt%, and 1.5wt% Fe corresponding to the Fe/Si mass ratios of 0.2, 0.4, and 0.6, respectively, in 2.5Si–xFe alloys, progressively decreased the HTS, where the HTS of the 2.5Si–1.5Fe alloy was 29% lower than that of 2.5Si alloy.
Fig.
12.
Effects of Fe content on fluidities and HTSs of (a) 1.0Si–xFe and (b) 2.5Si–xFe alloys.
Therefore, increasing the Fe/Si mass ratio up to 1 progressively improved the hot tearing resistance of alloy, while its further increase (up to 1.5) adversely affected the hot tearing resistance. According to the microstructural evaluation results (Fig. 7), the positive impact of increasing Fe/Si mass ratio (up to 1.0wt%) on hot tearing resistance can be partly explained by the formation and increase in the volume fraction of thermally stable α-FeMn compound in the microstructure of A206 alloy. Due to their high melting point [7], these intermetallics can enhance the alloy resistance against temperature-induced strains arising during solidification [19] whenever they have appropriate size and morphology. Moreover, as previously mentioned, due to the suitable effect of Fe addition on decreasing the alloy freezing range, Fe addition (up to 1.0wt%) improved the fluidity and increased the healing chance of incipient tears. However, if the Fe/Si mass ratio exceeds 1, the formation of interdendritic β-CuFe platelets in the microstructure can block interdendritic feeding channels, disrupting the healing process and encouraging the nucleation and coalescence of solidification micropores; this subsequently results in hot tear formation.
Fig. 13 shows the surface analysis results of the broken rods of 1.0Si–1.5Fe (Fe/Si mass ratio > 1), 1.0Si–1.0Fe (Fe/Si mass ratio = 1), and 2.5Si–1.0Fe (Fe/Si mass ratio < 1) alloys. In agreement with the microstructural evaluation result (Fig. 7(a)), a high volume fraction of β-CuFe platelets with averages of 61.80wt% Al, 12.37wt% Fe, and 25.83wt% Cu was present on the hot tear surface of the 1.0Si–1.5Fe alloy (Fig. 13(a)). It is evident that the β-CuFe platelets blocked the interdendritic feeding channels, promoting the formation of shrinkage pores, inhibiting the microcracks healing, and therefore decreasing the alloy resistance to hot tearing.
Fig.
13.
SEM micrographs showing the hot-torn fracture surfaces of (a) 1.0Si–1.5Fe, (b) 1.0Si–1.0Fe, and (c) 2.5Si–1.5Fe alloys.
In the alloy with the lowest susceptibility to hot tearing (1.0Si–1.0Fe alloy with Fe/Si mass ratio of 1), the widespread self-healing occurred through most interdendritic regions as well as the voids entrained between hot cracks (Fig. 13(b)). Therefore, the low HTS of this alloy can be explained by its high potential of the eutectic healing of already initiated tears (Fig. 12(a)) and the absence of β-CuFe particles in its microstructure (Fig. 7(b)). The average composition of the frozen eutectic phase, obtained by EDS analysis (Fig. 13(b)), was 56.04wt% Al, 0.68wt% Si, 0.04wt% Fe, and 43.24wt% Cu. Moreover, the rather full conversion of plate-like β-CuFe particles to high-melting-point α-FeMn intermetallics [7,11] in the microstructure of 1.0Si–1.0Fe alloy is likely to improve the high-temperature strength of the alloy, making it more resistant to applied solidification strains. In the 2.5Si–1.0Fe alloy shown in Fig. 13(c), the shrinkage porosities may decrease the effective cross-sectional area of the castings, and consequently their strength, which, in turn, decreases the hot tearing resistance of the alloy.
4.
Conclusions
(1) The tensile properties of Al–4.5Cu–(0.5–2.5)Si–(0.5–1.5)Fe alloys were optimal when the Fe/Si mass ratio was unity. Increasing the volume fraction of Si platelets and shrinkage micropores deteriorated the tensile properties of the samples for which the Fe/Si mass ratios were less than unity. However, the larger size and higher volume fractions of Fe-rich platelets were the main factors responsible for weaker tensile properties of the samples with Fe/Si mass ratios of more than unity.
(2) Due to the better fluidity and formation of high-melting-point α-FeMn compounds in the microstructure, the hot tearing resistance was maximum when Fe/Si mass ratio was 1. Increasing the Si content (i.e., a lower Fe/Si mass ratio) improved the liquid fluidity. However, a higher volume fraction of interdendritic shrinkage micropores seemed to adversely affect the alloy hot tearing resistance.
(3) The SEM examination of hot-torn surfaces showed significant healing of the hot tears already formed on the fracture surface of the 1.0Si–1.0Fe A206 sample (Fe/Si mass ratio = 1). The fractography study of the hot-torn surfaces also indicated the negative impacts of β-Fe compounds and shrinkage micropores on the hot tearing resistances of the 1.0Si–1.5Fe and 2.5Si–1.5Fe alloys, respectively.
A. Lemieux, J. Langlais, D. Bouchard, and X.G. Chen, Effect of Si, Cu and Fe on mechanical properties of cast semi-solid 206 alloys, Trans. Nonferrous Met. Soc. China, 20(2010), No. 9, p. 1555. DOI: 10.1016/S1003-6326(09)60338-1
M. Tiryakioğlu, J. Campbell, and N.D. Alexopoulos, On the ductility potential of cast Al–Cu–Mg (206) alloys, Mater. Sci. Eng. A, 506(2009), No. 1-2, p. 23. DOI: 10.1016/j.msea.2008.10.044
F. Sheykh-jaberi, S.L. Cockcroft, D.M. Maijer, and A.B. Phillion, Comparison of the semi-solid constitutive behaviour of A356 and B206 aluminum foundry alloys, J. Mater. Process. Technol., 266(2019), p. 37. DOI: 10.1016/j.jmatprotec.2018.10.029
E.M. Elgallad and X.G. Chen, On the microstructure and solution treatment of hot tearing resistant semi-solid 206 alloy, Mater. Sci. Eng. A, 556(2012), p. 783. DOI: 10.1016/j.msea.2012.07.064
M.R.N. Esfahani and B. Niroumand, Study of hot tearing of A206 aluminum alloy using instrumented constrained T-shaped casting method, Mater. Charact., 61(2010), No. 3, p. 318. DOI: 10.1016/j.matchar.2009.12.015
Y.L. Zhao, W.W. Zhang, C. Yang, D.T. Zhang, and Z. Wang, Effect of Si on Fe-rich intermetallic formation and mechanical properties of heat-treated Al–Cu–Mn–Fe alloys, J. Mater. Res., 33(2018), No. 8, p. 898. DOI: 10.1557/jmr.2017.441
N.A. Belov, A.A. Aksenov, and D.G. Eskin, Iron in Aluminnium Alloys: Impurity and Alloying Elements, CRC Press, New York, 2002.
K. Liu, X. Cao, and X.G. Chen, Solidification of iron-rich intermetallic phases in Al–4.5Cu–0.3Fe cast alloy, Metall. Mater. Trans. A, 42(2011), No. 7, p. 2004. DOI: 10.1007/s11661-010-0578-7
H.K. Kamga, D. Larouche, M. Bournane, and A. Rahem, Mechanical properties of aluminium–copper B206 alloys with iron and silicon additions, Int. J. Cast Met. Res., 25(2012), No. 1, p. 15. DOI: 10.1179/1743133610Y.0000000012
W.W. Zhang, Y.L. Zhao, D.T. Zhang, Z.Q. Luo, C. Yang, and Y.Y. Li, Effect of Si addition and applied pressure on microstructure and tensile properties of as-cast Al–5.0Cu–0.6Mn–1.2Fe alloys, Trans. Nonferrous Met. Soc. China, 28(2018), No. 6, p. 1061. DOI: 10.1016/S1003-6326(18)64765-X
K. Liu, X. Cao, and X.G. Chen, Effect of Mn, Si, and cooling rate on the formation of iron-rich intermetallics in 206 Al–Cu cast alloys, Metall. Mater. Trans. B, 43(2012), No. 5, p. 1231. DOI: 10.1007/s11663-012-9694-7
K. Liu, X. Cao, and X.G. Chen, Tensile properties of Al–Cu 206 cast alloys with various iron contents, Metall. Mater. Trans. A, 45(2014), No. 5, p. 2498. DOI: 10.1007/s11661-014-2207-3
E. Taghaddos, M.M. Hejazi, R. Taghiabadi, and S.G. Shabestari, Effect of iron-intermetallics on the fluidity of 413 aluminum alloy, J. Alloys Compd., 468(2009), No. 1-2, p. 539. DOI: 10.1016/j.jallcom.2008.01.079
T.O. Mbuya, B.O. Odera, and S.P. Ng'ang'a, Influence of iron on castability and properties of aluminium silicon alloys: Literature review, Int. J. Cast Met. Res., 16(2003), No. 5, p. 451. DOI: 10.1080/13640461.2003.11819622
H.K. Kamga, D. Larouche, M. Bournane, and A. Rahem, Hot tearing of aluminum–copper B206 alloys with iron and silicon additions, Mater. Sci. Eng. A, 527(2010), No. 27-28, p. 7413. DOI: 10.1016/j.msea.2010.08.025
A. Bolouri, K. Liu, and X.G. Chen, Effects of iron-rich intermetallics and grain structure on semisolid tensile properties of Al–Cu 206 cast alloys near solidus temperature, Metall. Mater. Trans. A, 47(2016), No. 12, p. 6466. DOI: 10.1007/s11661-016-3744-8
F. D’Elia, C. Ravindran, D. Sediako, K.U. Kainer, and N. Hort, Hot tearing mechanisms of B206 aluminum–copper alloy, Mater. Des., 64(2014), p. 44. DOI: 10.1016/j.matdes.2014.07.024
H.F. Bishop, C.G. Ackerlind, and W.S. Pellini, Metallurgy and mechanics of hot tearing, AFS Trans., 60(1952), p. 818.
K. Liu, X. Cao, A. Bolouri, and X.-G. Chen, Effect of Fe-rich intermetallics on tensile behavior of Al–Cu 206 cast alloys at solid and near-solid states, [in] M. Tiryakioğlu, W. Griffiths, M. Jolly, eds., Shape Casting, The Minerals, Metals & Materials Series, Springer, Cham, 2019, p. 85.
N. Han, X.F. Bian, Z.K. Li, T. Mao, and C.D. Wang, Effect of Si on the microstructure and mechanical properties of the Al–4.5%Cu alloys, Acta Metall. Sinica, 19(2006), No. 6, p. 405. DOI: 10.1016/S1006-7191(06)62080-2
J.G. Kaufman, Introduction to Aluminum Alloys and Tempers, ASM International, Materials Park, OH, 2000, p. 15.
S. Lin, A Study of Hot Tearing in Wrought Aluminum Alloys [Dissertation], University of Québec in Chicoutimi, Chicoutimi, PQ, 1999.
S. Lin, C. Aliravci, and M.O. Pekguleryuz, Hot-tear susceptibility of aluminum wrought alloys and the effect of grain refining, Metall. Mater. Trans. A, 38(2007), No. 5, p. 1056. DOI: 10.1007/s11661-007-9132-7
R. Taghiabadi, A. Fayegh, A. Pakbin, M. Nazari, and M.H. Ghoncheh, Quality index and hot tearing susceptibility of Al–7Si–0.35Mg–xCu alloys, Trans. Nonferrous Met. Soc. China, 28(2018), No. 7, p. 1275. DOI: 10.1016/S1003-6326(18)64783-1
S.G. Bai, N. Perevoshchikova, Y. Sha, and X.H. Wu, The effects of selective laser melting process parameters on relative density of the AlSi10Mg parts and suitable procedures of the Archimedes method, Appl. Sci., 9(2019), 3, p. 583. DOI: 10.3390/app9030583
N. Tenekedjiev and J.E. Gruzleski, Hypereutectic aluminium–silicon casting alloys—A review, Cast Met., 3(1990), No. 2, p. 96. DOI: 10.1080/09534962.1990.11819026
R. Jamaati, S. Amirkhanlou, M.R. Toroghinejad, and B. Niroumand, CAR process: A technique for significant enhancement of as-cast MMC properties, Mater. Charact., 62(2011), No. 12, p. 1228. DOI: 10.1016/j.matchar.2011.10.008
L. Qian, H. Toda, S. Nishido, and T. Kobayashi, Experimental and numerical investigations of the effects of the spatial distribution of α phase on fracture behavior in hypoeutectic Al–Si alloys, Acta Mater., 54(2006), No. 18, p. 4881. DOI: 10.1016/j.actamat.2006.06.036
M. Gupta and S. Ling, Microstructure and mechanical properties of hypo/hyper-eutectic Al–Si alloys synthesized using a near-net shape forming technique, J. Alloys Compd., 287(1999), No. 1-2, p. 284. DOI: 10.1016/S0925-8388(99)00062-6
F. Průša, M. Bláhová, D. Vojtěch, V. Kučera, A. Bernatiková, T.F. Kubatík, and A. Michalcová, High-strength ultra-fine-grained hypereutectic Al–Si–Fe–X (X = Cr, Mn) alloys prepared by short-term mechanical alloying and spark plasma sintering, Materials, 9(2016), No. 12, p. 973. DOI: 10.3390/ma9120973
O. Prach, O. Trudonoshyn, P. Randelzhofer, C. Körner, and K. Durst, Effect of Zr, Cr and Sc on the Al–Mg–Si–Mn high-pressure die casting alloys, Mater. Sci. Eng. A, 759(2019), p. 603. DOI: 10.1016/j.msea.2019.05.038
C.H. Cáceres, M.B. Djurdjevic, T.J. Stockwell, and J.H. Sokolowski, The effect of Cu content on the level of microporosity in Al–Si–Cu–Mg casting alloys, Scripta Mater., 40(1999), No. 5, p. 631. DOI: 10.1016/S1359-6462(98)00492-8
M. Di Sabatino, On fluidity of aluminium alloys, Metall. Ital., 100(2008), No. 3, p. 17.
T. Magnusson and L. Arnberg, Density and solidification shrinkage of hypoeutectic aluminum–silicon alloys, Metall. Mater. Trans. A, 32(2001), No. 10, p. 2605. DOI: 10.1007/s11661-001-0050-9
G. Razaz and T. Carlberg, Hot tearing susceptibility of AA3000 aluminum alloy containing Cu, Ti, and Zr, Metall. Mater. Trans. A, 50(2019), No. 8, p. 3842. DOI: 10.1007/s11661-019-05290-1
M. Rappaz, I. Farup, and J.-M. Drezet, Study and modeling of hot tearing formation, [in] Proceedings of the Merton C. Flemings Symposium on Solidification and Materials Processing, Cambridge, Massachusetts, 2000, p. 213.
M.C. Flemings, E. Niyama, and H.F. Taylor, Fluidity of aluminium alloys: An experimental and quantitative evaluation, AFS Trans., 69(1961), p. 625.
Dong Sheng Wang, Wei Wei, Kun Xia Wei, et al. Grain Refinement and Mechanical Properties of Photovoltaic Frame 6063 Al Alloy with an Addition of ECAP-ed Al-5%Ti-1%B Master Alloy. JOM, 2024, 76(4): 2027.
DOI:10.1007/s11837-024-06413-9
2.
Yoshihiro Nagata, Ryohei Nakagawa, Takumi Kumaki, et al. Novel Control Factor for Tensile Strength and Solidification Cracking in Partially Solidified Al–Mn–Cu Alloy Based on Campbell’s Model with Fe-Rich Intermetallic Compounds. Metallurgical and Materials Transactions A, 2024.
DOI:10.1007/s11661-024-07564-9
3.
Shenghua Wu, Chong Yang, Peng Zhang, et al. Review of Sc microalloying effects in Al–Cu alloys. International Journal of Minerals, Metallurgy and Materials, 2024, 31(5): 1098.
DOI:10.1007/s12613-024-2841-8
4.
Padmakumar A. Bajakke, Vinayak R. Malik, Sudhakar C. Jambagi, et al. Unidirectional flipped multiple-pass friction stir process: an innovative step in the fabrication of in situ Al-Cu composites. The International Journal of Advanced Manufacturing Technology, 2024, 134(3-4): 1767.
DOI:10.1007/s00170-024-14248-6
5.
Kang Xiang, Ling Qin, Yuliang Zhao, et al. Operando study of the dynamic evolution of multiple Fe-rich intermetallics of an Al recycled alloy in solidification by synchrotron X-ray and machine learning. Acta Materialia, 2024, 279: 120267.
DOI:10.1016/j.actamat.2024.120267
6.
Sh Khalesi, R. Omidi, R. Taghiabadi, et al. Effect of Si Addition on the Weibull Distribution of Tensile Properties of Fe-bearing Al-Cu Alloys. Silicon, 2024, 16(11): 4609.
DOI:10.1007/s12633-024-03026-4
7.
Jing-Ran Sun, Bai-Xin Dong, Hong-Yu Yang, et al. The Role of Lithium in the Aging Precipitation Process of Al-Zn-Mg-Cu Alloys and Its Effect on the Properties. Materials, 2023, 16(13): 4750.
DOI:10.3390/ma16134750
8.
Zhengfeng Zhu, Renbao Qin, Yishan Sun, et al. Improvement in Mechanical Properties of Al2024 Alloy Using Mechanical Working and Heat Treatment. Materials, 2023, 16(16): 5568.
DOI:10.3390/ma16165568
9.
Hongchen Xiang, Wenjun Liu, Qiang Wang, et al. Improvement of Hot Tearing Resistance of AZ91 Alloy with the Addition of Trace Ca. Materials, 2023, 16(10): 3886.
DOI:10.3390/ma16103886
10.
Danyang Wang, Liang Bai, Wu Chen, et al. A Review of Experimental Research on Hot Tearing of Aluminium Alloy Casting. Transactions of the Indian Institute of Metals, 2023.
DOI:10.1007/s12666-023-03145-3
11.
Min Ao, Yucheng Ji, Pan Yi, et al. Relationship between elements migration of α-AlFeMnSi phase and micro-galvanic corrosion sensitivity of Al—Zn—Mg alloy. International Journal of Minerals, Metallurgy and Materials, 2023, 30(1): 112.
DOI:10.1007/s12613-022-2428-1
12.
Fuqing Jiang, Lei Tang, Sheng Li, et al. Achieving strength-ductility balance in a laser powder bed fusion fabricated TiB2/Al–Cu-Mg-Ag alloy. Journal of Alloys and Compounds, 2023, 945: 169311.
DOI:10.1016/j.jallcom.2023.169311
13.
R. Younes, M. Bournane, A. Idir, et al. Influence of the iron additive on the microstructural behavior of an aluminum-copper foundry alloy B206. Industrial laboratory. Diagnostics of materials, 2023, 89(9): 48.
DOI:10.26896/1028-6861-2023-89-9-48-52
14.
Kai-Yu Liang, Hao-Chuan Huang, Ching-Yao Tseng, et al. Effects of Fe, Si, Cu, and TiB2 Grain Refiner Amounts on the Hot Tearing Susceptibility of 5083, 6061, and 7075 Aluminum Ingots. Metals, 2023, 14(1): 15.
DOI:10.3390/met14010015
15.
Yuliang Zhao, Weixiang He, Dongfu Song, et al. In-situ synchrotron X-ray radiography study of primary Fe-rich phases growth in Al-Fe( Cu) alloys. Materials Characterization, 2023, 195: 112539.
DOI:10.1016/j.matchar.2022.112539
16.
Congcong Wang, Mingxing Guo, Kangcai Yu, et al. Evading the low formability dilemma in Al-Zn-Mg-Cu alloys via heterogeneous structure. Materials Characterization, 2022, 193: 112295.
DOI:10.1016/j.matchar.2022.112295
17.
Yuliang Zhao, Weixiang He, Dongfu Song, et al. Effect of ultrasonic melt processing and Al-Ti-B on the microstructural refinement of recycled Al alloys. Ultrasonics Sonochemistry, 2022, 89: 106139.
DOI:10.1016/j.ultsonch.2022.106139
18.
M. Mostafavi, R. Taghiabadi, M. Jafarzadegan. Optimizing the mechanical properties of Al-4.5Cu-xSi alloys through multi-pass friction stir processing and post-process aging. Archives of Civil and Mechanical Engineering, 2022, 22(1)
DOI:10.1007/s43452-021-00333-0
19.
Hongmin Guo, Wenjie Hu, Kunyi Yang, et al. Effects of Si and Fe additions on mechanical properties of as-cast Al–5Cu alloys. Materials Science and Technology, 2022, 38(6): 352.
DOI:10.1080/02670836.2022.2043076
20.
Chunyu Yue, Xiaoguang Yuan, Ming Su, et al. Effect of adding Pr on the microstructure and hot tearing sensitivity of as-cast Al-Cu-Mg alloys. Materials Characterization, 2022, 191: 112141.
DOI:10.1016/j.matchar.2022.112141
21.
Hengbin Liao, Liling Mo, Xiong Zhou, et al. Revealing the nucleation event of Mg-Al alloy induced by Fe impurity. International Journal of Minerals, Metallurgy and Materials, 2022, 29(7): 1317.
DOI:10.1007/s12613-021-2406-z
22.
Xin Zhang, Jia-Yu Hu, Bai-Xin Dong, et al. Effect of Cu and Zn elements on morphology of ceramic particles and interfacial bonding in TiB2/Al composites. Ceramics International, 2022, 48(18): 25894.
DOI:10.1016/j.ceramint.2022.05.266
23.
Hamed Kamali, Erfan Kamali, Massoud Emamy. Effects of Zr additions on structure and tensile properties of an Al-4.5Cu-0.3Mg-0.05Ti (wt.%) alloy. China Foundry, 2022, 19(1): 9.
DOI:10.1007/s41230-022-1094-2
24.
Ge Yi, Heng Li, Chenyang Zang, et al. Remarkable improvement in strength and ductility of Al-Cu foundry alloy by submicron-sized TiC particles. Materials Science and Engineering: A, 2022, 855: 143903.
DOI:10.1016/j.msea.2022.143903
25.
Xiangyang Chen, Mingxing Guo, Dexian Qiao, et al. Coordinated deformation and high formability mechanisms of Al–Zn–Mg–Si–Fe–Mn–Cu alloys with heterogeneous cellular structure. Materials Science and Engineering: A, 2022, 848: 143407.
DOI:10.1016/j.msea.2022.143407
26.
Xiang-Yi Song, Ya-Jie Wang, Jing-Xin Zhang, et al. Solidification microstructure manipulation mechanism of hypoeutectic and hypereutectic Al–Si alloys controlled by trace in-situ nano-Fe2B and nano-Fe3Si. Journal of Materials Research and Technology, 2022, 21: 3856.
DOI:10.1016/j.jmrt.2022.10.142
27.
Ming Su, Xiaoguang Yuan, Chunyu Yue, et al. Effect of liquid film constitutional supercooling at the end of solidification on the hot cracking behavior of Al-4.0 wt.%Cu alloy. Materials Letters, 2022, 328: 133116.
DOI:10.1016/j.matlet.2022.133116
28.
Chunyu Yue, Xiaoguang Yuan, Ming Su, et al. Effect of Rare Earth Pr on Hot Tearing Sensitivity of As-Cast Al-Cu-Mg Alloys. SSRN Electronic Journal, 2021.
DOI:10.2139/ssrn.3980043
Search
DownLoad:
