Iron is known to be the most common and at the same time the most detrimental impurity in Al-Si alloys due to its role in the formation of brittle Fe-rich intermetallic phases that precipitate during solidification process. In this way, commercial Al-Si alloys always contain Fe impurities; this undesirable contamination is inevitable and cannot be economically removed from the molten aluminum, it comes either from the use of steel tools for melting and casting or from the use of scrap as the initial material.

The Fe impurities have a detrimental effect on the mechanical properties and extrusion characteristics of the Al-Si alloys. This is attributed to the formation of different Fe-containing intermetallic phases, such as α-AlFeSi and β-AlFeSi [1] -[5] . The α phase is usually described as having a cubic, α_c-Al₁₅(Fe,Mn)₃Si₂, [6] [7] or hexagonal, α_h-Al₈Fe₂Si, [8] - [10] crystal structure, and the β phase is generally considered to be monoclinic Al₅FeSi [11] - [15] . Techniques have to be developed to neutralize the detrimental effect of Fe impurities: 1) rapid solidification; 2) manganese addition; 3) melt superheating; and 4) inclusion particles addition. The three first methods basically convert the crystallization of the β phase to the less harmful Chinese script form, α_c phase [3] [8] . Each one of these methods has its own advantages and limitations.

The inclusion particles addition technique is based on the work performed by Campbell and co-workers [16] - [18] . They have reported that the Fe-rich intermetallic phases nucleate and grow on the wetted surfaces of the oxide films, which have long been identified as cracks in the castings. Campbell et al. have further applied this nucleation mechanism to explain the cracks observed along the central axis of some Fe-rich intermetallics and for remove oxide films from the liquid aluminum. According to Cao and Campbell [17] [18] , the wet side of the oxide films acts as a preferential nucleation site for Fe-rich intermetallic phases. As these particles have an increased sedimentation tendency, they sink to the bottom, bringing the oxide films with them.

The nucleation of Fe-intermetallic phases, such as α and β, was reported to be enhanced in the presence of particular inclusions added such as α-Al₂O₃ and MgAl₂O₄ [19] . Khalifa et al. [20] have investigated the nucleation of Fe-intermetallic phases on the surface of different inclusions. They have reported that some inclusions exhibit higher potency for the nucleation of particular Fe-intermetallic phases under certain conditions and poor potency under other conditions. In general, these results showed limited effect of the inclusion particles added on the nucleation of Fe-intermetallic phases. For example, γ-Al₂O₃, which is a potent nucleant for the primary phase α-Al, exhibit poor potency for the nucleation of the Fe-intermetallic phases. On the other hand, Lu and Dahle [21] did not find direct physical contact between Fe-rich intermetallic phases and the oxide films, thus the role of the oxide films as a heterogeneous substrate of Fe-intermetallics is uncertain.

The present work was performed to provide additional experimental information about the interaction between the Fe-intermetallic phases and α-Al₂O₃ oxide films, commonly present in Al-Si alloys, in order to clarify the role of this oxide as heterogeneous substrate for the nucleation of Fe-intermetallic phases. This aspect would aid in controlling intermetallic formation, allowing control of the microstructure of the castings and hence improvement of the mechanical properties of the commercial Al-Si alloys. It is worth mentioning that any intermetallic phase formation on γ-Al₂O₃ films was not investigated in this work.

Figure 2 shows the cooling curve obtained for the alloy 1 cooled at a rate of ~1˚C/s; its first derivative is also included. Each thermal arrest in the cooling curve is associated with the formation temperature of a microstructural phase; the corresponding T_p was estimated with the aid of an associated peak appearing in the first derivative curve (dotted line in Figure 2). According to microstructural observation of the cross area of the alloy 1, the formation temperature of α_sc-AlFeMnSi phase (hereinafter denoted as α_sc) is associated with the first thermal arrest occurring in the curve, the second is related with the formation of primary silicon particles and the last represents the heat release during eutectic reaction. Figure 3 shows SEM images of the α_sc phase formed in alloy 1, the irregular morphology observed in the as-cast samples is shown in Figure 3(a), and the polyhedral morphology observed in the isothermally-treated samples is shown in Figure 3(b) and Figure 3(c). In Figure 3(d) is shown the whole sample obtained by DTA, in this sample can be seen the irregular morphology of the α_sc phase, also can be confirmed that this phase not presents direct physical contact with the oxide film of α-Al₂O₃ formed in the sample’s surface by natural oxidation. The presence of α-Al₂O₃ in the isothermally-treated samples was confirmed by SEM and XRD, as is shown in Figure 6(c), Figure 8 and Figure 10(c), shown in the following paragraphs. Additionally, some research suggest that oxides such as α-Al₂O₃ may be present within oxide films depending on the superheating temperature, in this work 900˚C, which favors the transformation of oxide films from γ-Al₂O₃ to α-Al₂O₃ [22] .

The intermetallics extracted from the isothermally-treated samples of alloy 1 (see Figure 3(c)), were analyzed by XRD. As can be observed in Figure 4, all peaks correspond to the Al₅₀Mn₁₂Si₇ phase (JCPDS card number 71-2182); this phase exhibits a simple cubic (sc) crystal structure with a Pm3 space group [23] . In this work, however, according to the EDS data given in Table 2, the extracted intermetallic actually corresponds to the α_sc-Al_34.1FeMn_8.5Si_6.16 formula, with a Fe/Mn weight ratio of 0.12, and a total concentration of transition metals (Fe + Mn) of 32.4%. In comparison, in Figure 4 are shown the XRD pattern of the Al₅₀Mn₁₂Si₇ phase and the experimental XRD pattern of the α_sc-Al_34.1FeMn_8.5Si_6.16 phase. It is observed that both patterns show similar peaks, in some cases the peaks intensity and the 2θ angles change in a little extent. This behavior has been attributable to the substitution of manganese atoms by iron atoms [4] [24] in the lattice sites of the crystal structure. Similar results were observed for the α_bcc and α_h-AlFeMnSi phases, shown in the following paragraphs.

The cooling curve obtained for the alloy 3 clearly reveals three slope changes during solidification process, as is shown in Figure 5. Based on microstructural analysis, the first slope change observed is related with the formation temperature of α_bcc-AlFeMnSi phase (hereinafter abbreviated as α_bcc). For the second slope change, the possible reaction indicates the formation of β-AlFeSi and silicon particles. The β phase was not analyzed by XRD due to the small amount formed in the isothermally-treated samples and the fact that it was mixed with the previously formed α_bcc phase; however, it was identified as β phase based on its typical morphology reported in literature data [25] . Finally, the third slope change is related to the eutectic reaction that is the last to solidify.

SEM micrographs of the α_bcc phase are shown in Figure 6. The irregular morphology of the α_bcc phase can be observed in the as-cast sample (Figure 6(a)), and the polyhedral morphology in some intermetallics extracted from the isothermally-treated samples shown in Figure 6(b) and Figure 6(c). Figure 6(c), also shows α_bcc particles formed on the surface of a α-Al₂O₃ oxide film, this observation provided the first direct evidence that the α_bcc phase can nucleate and grow on the surface of α-Al₂O₃ oxide films. Figure 6d can be considered as another evidence for the association between α_bcc phase and α-Al₂O₃ oxide films, in this Figure is shown a sample of the alloy 3 obtained by DTA, in the sample can be seen polyhedral α_bcc particles revealing that this phase presents direct physical contact with the α-Al₂O₃ oxide film formed in the sample’s surface by natural oxidation.

The XRD pattern of the intermetallic particles extracted from alloy 3 is shown in Figure 7. As can be observed, all peaks could be indexed according to the JCPDS card number 71-4015, which corresponds to the Al₁₇Fe_3.2Mn_0.8Si₂ phase. This phase exhibits a body-centered cubic (bcc) crystal structure and belongs to the Im3 space group [26] . For alloy 3, according to the EDS analysis given in Table 2, the extracted intermetallic par-

ticles corresponded to the α_bcc-Al_8.8Fe_1.36MnSi_1.47 formula. Also, the EDS data given show that the iron contained in the α_bcc phase increased, while the manganese contained decreased with increasing iron content in the alloy. The total concentration of transition metals (Fe + Mn) in the α_bcc phase fell in the range of 31% - 32%, when iron was added to the alloys. The Fe/Mn weight ratio was in the range of 0.45 - 1.39 for the extracted α_bcc phase, therefore the Fe/Mn weight ratio depended on the alloy’s chemical composition. It is noteworthy that similar results were obtained for the alloy 2 shown in Table 2.

XRD results also shown the presence of α-Al₂O₃ in the powder produced by grinding the extracted particles from alloys 3 and 4, the last described in the following paragraphs. As observed in Figure 8, all peaks of the experimental XRD pattern of the α-Al₂O₃ were indexed according to the JCPDS card number 74-1081, which corresponds to the α-Al₂O₃ with a hexagonal crystal structure and R-3c space group [27] . It is worth mentioning that the alumina sample was obtained by separating α phases and alumina particles from particles like those observed in Figure 6(c) and Figure 10(c). Such α-Al₂O₃ particles have been thought to originate during casting of the melt into the graphite crucible and subsequent melt superheating at 900˚C.

The results of the alloy 4 (Figure 9) indicate that there are two Fe-rich intermetallic phases prior to the eutectic reaction. The primary phase for this alloy was α_h-AlFeMnSi (hereinafter denoted as α_h), which is associated to the first slope change. The second slope change represents the formation temperature of β-AlFeSi phase and some silicon particles. Finally, the last slope change represents the heat release during the eutectic reaction. SEM images of the α_h phase revealed that this phase exhibits irregular morphology in the as-cast samples (Figure 10(a)), and polyhedral morphology in the isothermally-treated samples, as can be seen in the intermetallics extracted shown in Figure 10(c). In this Figure, also can be observed that α_h particles are formed on the surface of a α-Al₂O₃ oxide film. This microstructural observation provided direct evidence that confirms that the α_h phase

can nucleate and grow on the wetted surface of α-Al₂O₃ oxide films. Figure 10(d) shown a sample of the alloy 4 obtained by DTA, this sample has revealed that α_h particles with polyhedral morphology are located on the wetted surface of the oxide film formed in the sample’s surface by natural oxidation, so can be established that the α_h phase presents direct physical contact with the α-Al₂O₃ oxide film. This observation can support the association between α_h phase and α-Al₂O₃ oxide film.

The intermetallic particles extracted from isothermally-treated samples of alloy 4, were analyzed by XRD. As observed in Figure 11, all peaks of the experimental XRD pattern were indexed according to the JCPDS card number 71-0238, which corresponds to the Al_167.8Fe_44.9Si_23.9 phase. This phase exhibits a hexagonal crystal struc- ture, and belongs to the P63/mmc space group [28] . In this work, however, the chemical composition of the extracted intermetallic corresponded to the α_h-Al₁₄Fe_3.36MnSi_3.15 formula according to Table 2. EDS data of the α_h phase revealed that the total concentration of transition metals (Fe + Mn) in the α_h phase fell in the range of 33.3% - 34.2%. The Fe/Mn weight ratio was in the range of 3.41 - 4.22 for the extracted α_h phase so can be es-

tablished that the Fe/Mn weight ratio depended on the alloy’s chemical composition. It is worth mentioning that similar results were observed for the alloy 5 shown in Table 2.

Based on microstructural observation there is good experimental evidence that led us to assume that the α_bcc and α_h phases present direct physical contact with the α-Al₂O₃ oxide films, as is shown in Figure 6(c) and Fig- ure 10(c). If α_bcc and α_h phases nucleate and grow on the wetted surface of α-Al₂O₃ oxide films, it is expected that there will be good atomic contact and good lattice match between the wetted surface of α-Al₂O₃ oxide films and the α_bcc and α_h particles that have grown on it. This assumption led us to suggesting that α-Al₂O₃ oxide films act as a preferential nucleation site for α_bcc and α_h phases due to the similarity in their crystal structure. Also, if α-Al₂O₃ oxide film is a good nucleant for α_bcc and α_h phases, it is expected that it will determine the spatial distribution of these phases, as can be observed in Figure 6(d) and Figure 10(d), where the α_bcc and α_h particles are located on the oxidized melt surface in DTA samples of alloys 3 and 4. This observation tends to support the assumption that α-Al₂O₃ oxide films act as a preferential nucleation site for α_bcc and α_h phases.

Careful studies of crystallographic orientation relationships may be another way to confirm that α-Al₂O₃ oxide film is a good nucleant for α_bcc and α_h phases. In general, it is accepted that the mismatch of crystal structures, the electronic nature of the crystal bonds of the two phases, superheating, and the presence of alloying elements are factors that may influence the favorability of nucleation.

On the other hand, for the β phase formed in the experimental alloys, was not found direct physical contact with the α-Al₂O₃ oxide film formed in the sample’s surface by natural oxidation, as can be seen in Figure 6(d) and Figure 9(d). This observation led us to assume that between the β phase and α-Al₂O₃ oxide films there is not good atomic contact and good lattice match, thus, the β phase did not appear to nucleate and grow on the surface of α-Al₂O₃ oxide films.

According to the Mn content of the studied alloys could be expected the formation of MnO·Al₂O₃, Mn₂O₃, or MnO·SiO₂ particles, however, the presence of these oxides was not detected by X-ray diffraction, as can be seen in Figure 8. Another effect that could be expected is that Mn content had an effect on the alumina transformation from γ-Al₂O₃ to α-Al₂O₃, nonetheless, according to the experimental XRD pattern of the α-Al₂O₃, shown in Figure 8, such effect was discarded for the studied temperature range. Thus, it was thought that Mn content just become embedded in the solidification of the α_sc, α_bcc and α_h phases. This observation is in accordance with some research [4] [24] , that suggest that many ternary AlFeSi intermetallics can accept Mn atoms into solid solution, which reflects the fact that Mn atoms may replace the Fe atoms in their crystal structure. Finally, it is important study whether the Fe-rich intermetallic phases can nucleate at all without the presence of oxide films in aluminum melts. This important question is not easily answered because obtaining aluminum melts free from oxide films remains difficult.

According to the experimental results, can be established that the α_bcc and α_h phases nucleate and grow on the wetted surface of α-Al₂O₃ oxide films formed by natural oxidation in aluminum melts. On the other hand, the β particles formed in the experimental alloys not presents direct physical contact with the α-Al₂O₃ oxide films, thus, the β phase did not appear to nucleate on α-Al₂O₃ oxide films.

The authors are grateful to the National Council for Science and Technology of Mexico (Conacyt) for the financial support provided to carry out this work (Project No. 149030).