E ﬀ ects of Homogenization Conditions on the Microstructure Evolution of Aluminium Alloy EN AW 8006

: The industrial production of products, such as foil and aluminium alloy strips, begins with the production of semi-ﬁnished products in the form of slabs. These are produced by the continuous casting process, which is quick and does not allow the equilibrium conditions of solidiﬁcation. Non-homogeneity—such as micro and macro segregation, non-equilibrium phases and microstructural constituents, as well as stresses arising during non-equilibrium solidiﬁcation—are eliminated by means of homogenization annealing. In this way, a number of technological di ﬃ culties in the further processing of semi-ﬁnished products can be avoided. The aim of this research was the optimization of the homogenization annealing of the EN AW 8006 alloy. With the Thermo-Calc software, a thermodynamic simulation of equilibrium and non-equilibrium solidiﬁcation was performed. Di ﬀ erential scanning calorimetry (DSC) was performed on selected samples in as-cast state and after various regimes of homogenization annealing and was used for the simulation of homogenization annealing. Using an optical microscope (OM), a scanning electron microscope (SEM) and an energy dispersion spectrometer (EDS), the microstructure of the samples was examined. Based on the results, it was concluded that homogenization annealing has already taken place after 8 h at 580 ◦ C to the extent, that the material is then suitable for further processing.


Introduction
Due to their relatively high strength, combined with their sufficient ductility, the non-hardenable alloys Al-Fe-Mn-Si are largely used in the automotive industry, cooling systems, civil engineering, etc. Al-foil can be produced either by starting from a conventional direct chill (DC) cast ingot which is then hot rolled to a strip of about 2-5 mm, or from continuous casting a 6-7-mm-thick sheet (twin roll casting or belt casting) and then cold rolling it to an intermediate (foil stock) gauge of about 0.4-1 mm. This is followed by annealing in the temperature range 350-400 • C. [1] Then, the material is cold rolled to final foil thickness. Afterwards, closed gap rolling must be used to achieve thinner foil gauges below 100 µm, which is performed in a dedicated foil mill. Most Al-foil products will be used in the O-temper conditions (foil must be soft annealed). In addition to the recrystallization, the final gauge annealing is required to remove lubricant from the foil [2].
Thin aluminium sheets which exhibit good formability and sufficient rigidity are advisable for use in applications such as fin-stocks in heat exchangers. Thus, a fine grain size structure and a texture warranting low anisotropy must be achieved by alloy processing. The form in which alloying phase is isomorphous to the orthorhombic Al6Mn phase usually encountered in Mn-containing alloys (3xxx or 5xxx series alloys) [12]. Furthermore, the literature describes other metastable phases of AlmFe where m varies from 4 to 5 [14]. A transformation of the metastable Fe-bearing phases towards stable Al3Fe can be achieved by homogenization at high temperatures [7,16,17]. If Mn is present, the α-AlFeSi phase is replaced by an isostructural quaternary α-Al(Mn,Fe)Si phase from the system presented in Figure 1c, which may show large variation in exact chemical composition [6,16]. , and a polythermal projection of the solidification surface of an Al-Fe-Mn-Si system (c) reproduces from reference [12].
The industrial production of products, such as strips and foils, from aluminium alloys starts with a semi-finished product in the form of a slab, which does not allow solidification at equilibrium conditions. Inhomogeneities, micro-segregations, residual stresses and other defects that occur in the material during un-equilibrium solidification are eliminated with homogenization annealing. [1] The more present internal defects are the large AlFeMn and/or AlFeMnSi intermetallic particles with different sizes, which are aligned on the deformation direction. Some zones have a coarse and fragile compound. The agglomeration of intermetallic particles leads to the rupture of the foil. Heat treatment (the homogenization of the alloy at 570 °C/180 min.) can prevent the structure defects observed in sheets and foils produced from EN AW 8006 alloy. [18,19] Aluminium oxide and/or spinel (MgO·Al2O3) on the surface of the sheets can be observed. It is more likely for these inclusions to become sources of pinholes in the sheets. Moreover, spherical particles of aluminium boride can be seen on the surface of the sheets. They are small in size and non-deformable particles compared to the ductile aluminium matrix, but they can affect the performance of the sheet rolling. [19,20] Homogenization processing has a critical contribution by allowing the precipitation of excessive alloying elements in the solid solution before subsequent thermomechanical processing. However, the high solidification rate encountered in TRC (twin-roll casting) not only favours the supersaturation and metastable condition of the alloying elements, but also hinders their microsegregation. [21,22] In the present study, the microstructural evolution of the typical dilute Al-Fe-Si (EN AW 8006) foil alloy during homogenization was investigated using optical and electron microscopy as well as differential scanning calorimetry (DSC). The effect of different homogenization practices on the formation of intermetallic phases was analysed. Especially, the impact of homogenization on changes in the morphology of plate-like constituent particles which are known [23] to adversely affect the formability of 8xxx series alloys was examined. The conditions for equilibrium solidification are not fulfilled in practice due to quite high cooling rates prevailing during industrial DC-casting [9]. Therefore, the as-cast ingot will be supersaturated with Fe and by other slow-diffusing species, such as Mn and Cr, if they are present. Under these conditions, a metastable phase Al 6 Fe can form in addition to the stable monoclinic Al 3 Fe phase. This phase is isomorphous to the orthorhombic Al 6 Mn phase usually encountered in Mn-containing alloys (3xxx or 5xxx series alloys) [12]. Furthermore, the literature describes other metastable phases of Al m Fe where m varies from 4 to 5 [14]. A transformation of the metastable Fe-bearing phases towards stable Al 3 Fe can be achieved by homogenization at high temperatures [7,16,17]. If Mn is present, the α-AlFeSi phase is replaced by an isostructural quaternary α-Al(Mn,Fe)Si phase from the system presented in Figure 1c, which may show large variation in exact chemical composition [6,16].

Materials and Methods
The industrial production of products, such as strips and foils, from aluminium alloys starts with a semi-finished product in the form of a slab, which does not allow solidification at equilibrium conditions. Inhomogeneities, micro-segregations, residual stresses and other defects that occur in the material during un-equilibrium solidification are eliminated with homogenization annealing. [1] The more present internal defects are the large AlFeMn and/or AlFeMnSi intermetallic particles with different sizes, which are aligned on the deformation direction. Some zones have a coarse and fragile compound. The agglomeration of intermetallic particles leads to the rupture of the foil. Heat treatment (the homogenization of the alloy at 570 • C/180 min.) can prevent the structure defects observed in sheets and foils produced from EN AW 8006 alloy. [18,19] Aluminium oxide and/or spinel (MgO·Al 2 O 3 ) on the surface of the sheets can be observed. It is more likely for these inclusions to become sources of pinholes in the sheets. Moreover, spherical particles of aluminium boride can be seen on the surface of the sheets. They are small in size and non-deformable particles compared to the ductile aluminium matrix, but they can affect the performance of the sheet rolling [19,20].
Homogenization processing has a critical contribution by allowing the precipitation of excessive alloying elements in the solid solution before subsequent thermomechanical processing. However, the high solidification rate encountered in TRC (twin-roll casting) not only favours the supersaturation and metastable condition of the alloying elements, but also hinders their micro-segregation [21,22].
In the present study, the microstructural evolution of the typical dilute Al-Fe-Si (EN AW 8006) foil alloy during homogenization was investigated using optical and electron microscopy as well as differential scanning calorimetry (DSC). The effect of different homogenization practices on the formation of intermetallic phases was analysed. Especially, the impact of homogenization on changes in the morphology of plate-like constituent particles which are known [23] to adversely affect the formability of 8xxx series alloys was examined.

Materials and Methods
The aim of the study was to optimize the homogenization annealing of the EN AW 8006 alloy slab. Standard composition is listed in Table 1. A thermodynamic analysis was conducted using Thermo-Calc Software (TCW2019, Stockholm, Sweden) and TCAL6 database to obtain the phase diagram and the stability of different phases, as well as the equilibrium concentrations of these phases.
Samples for heat treatment investigations were cut from the middle of the front surface of a slab with dimensions of 510 × 1310 × 4800 mm, in the form of cubes. The samples' dimensions were approximately 15 mm (l) × 15 mm (w) × 15 mm (h). Single-step solution heat treatments (homogenization treatments) were done using an electric chamber furnace, which was previously calibrated. For the investigation of solution heat treatments, samples were held for various times-4 h, 6 h, 8 h, 10 h and 12 h-at temperatures of 580 • C and 600 • C to investigate the effect of homogenization on the dissolution of particles. Samples were, after a certain time, removed from the furnace and cooled in air in order to simulate industrial conditions. All metallographic examinations were carried out on half of the samples after homogenization annealing, whereas the remainder were grinded and polished according to the standard metallographic procedure for aluminium alloys. Differential scanning calorimetry (DSC) was performed on samples in an as-cast state and on the second half of the homogenized samples, after different regimes of homogenization annealing, using an STA Jupiter 449c apparatus (NETZSCH Group, Selb, Germany). DSC analysis was performed as follows: heating up to 710 • C, followed by cooling to room temperature using 10 K/min heating and cooling rate in a protective atmosphere of argon. Using an optic microscope Zeiss Axio Observer 7 (ZEISS International, Jena, Germany) and a scanning electron microscope SEM-Jeol JSM 6610LV (JEOL Ltd., Tokyo, Japan) with an energy dispersion spectrometer (EDS), the microstructural components were analysed.
Samples for homogenisation simulation on the DSC apparatus were taken from the middle of the slab. On the DSC device, the homogenization of the as-cast samples was carried out, which lasted for 12 h at 580 • C and 600 • C. By using the tangent method, the change in the slope of the DSC curves was determined, which is also attributed to changes in the course of homogenization. Figure 2 illustrates the results of the thermodynamic calculations. The results indicate that the temperature range for the dissolution is 550-620 • C. As shown in Figure 2, the equilibrium phases and the temperature ranges in which they are present are: the Al 13 Fe 4 phase up to 652 • C, the Al 15 Si 2 Mn 4 phase up to 630 • C, the Al 9 Fe 2 Si 2 phase up to 477 • C, and the Al 6 Mn phase up to 444 • C. This implies that by heating at 580 • C or 600 • C for a long holding time, Al 15 Si 2 Mn 4 will partially dissolve, and Al 9 Fe 2 Si 2 and Al 6 Mn will completely dissolve, but the iron phase Al 13 Fe 4 , as well as a fraction of the Al 15 Si 2 Mn 4 phase, will still be present. It should also be noted that the homogenization treatment for these alloys should not be conducted at temperatures over 640 • C, as the Fe-rich eutectic phase Al 13 Fe 4 will start to melt. In Figure 3, the heating and cooling DSC curves of the experimental samples are presented, whereas only samples in an as-cast state, and after 6 h and 12 h at temperatures of 580 °C and 600 °C, were analysed.

Results and discussion
From these results, characteristic temperatures at heating and cooling, as well as melting (ΔHm) and solidification (ΔHs) enthalpies were determined.
From Figure 3a it can be determined, according to Thermo-Calc calculations, that at a temperature of about 649 °C the melting of the eutectic with the Mn-phase occurs, at a temperature of about 664 °C the melting of eutectic (α-Al + Al13Fe4) occurs, and at a temperature of 685 °C the melting of the primary crystals α-Al takes place. From the cooling DSC curves, the solidification is as follows (Figure 3b): the primary solidification of α-Al occurs at about 650 °C, and the solidification of the eutectic (α-Al + Al13Fe4) occurs at a temperature of about 621 °C. The solidification of Mn-phase could not be detected due to an equilibrium solidification.  In order to improve transparency, the temperatures and enthalpies are collected in Table 2. It can be considered that the melting temperature of the Mn-eutectic is the lowest in the sample in the as-cast state (647.3 °C), the melting temperature is increased in homogenized samples, and is the highest (649.7 °C) in samples that were homogenized at 600 °C. This trend is also observed in the In Figure 3, the heating and cooling DSC curves of the experimental samples are presented, whereas only samples in an as-cast state, and after 6 h and 12 h at temperatures of 580 • C and 600 • C, were analysed. In Figure 3, the heating and cooling DSC curves of the experimental samples are presented, whereas only samples in an as-cast state, and after 6 h and 12 h at temperatures of 580 °C and 600 °C, were analysed.
From these results, characteristic temperatures at heating and cooling, as well as melting (ΔHm) and solidification (ΔHs) enthalpies were determined.
From Figure 3a it can be determined, according to Thermo-Calc calculations, that at a temperature of about 649 °C the melting of the eutectic with the Mn-phase occurs, at a temperature of about 664 °C the melting of eutectic (α-Al + Al13Fe4) occurs, and at a temperature of 685 °C the melting of the primary crystals α-Al takes place. From the cooling DSC curves, the solidification is as follows (Figure 3b  In order to improve transparency, the temperatures and enthalpies are collected in Table 2. It can be considered that the melting temperature of the Mn-eutectic is the lowest in the sample in the as-cast state (647.3 °C), the melting temperature is increased in homogenized samples, and is the highest (649.7 °C) in samples that were homogenized at 600 °C. This trend is also observed in the From these results, characteristic temperatures at heating and cooling, as well as melting (∆H m ) and solidification (∆H s ) enthalpies were determined.
From Figure 3a it can be determined, according to Thermo-Calc calculations, that at a temperature of about 649 • C the melting of the eutectic with the Mn-phase occurs, at a temperature of about 664 • C the melting of eutectic (α-Al + Al 13 Fe 4 ) occurs, and at a temperature of 685 • C the melting of the primary crystals α-Al takes place. From the cooling DSC curves, the solidification is as follows (Figure 3b): the primary solidification of α-Al occurs at about 650 • C, and the solidification of the eutectic (α-Al + Al 13 Fe 4 ) occurs at a temperature of about 621 • C. The solidification of Mn-phase could not be detected due to an equilibrium solidification.
In order to improve transparency, the temperatures and enthalpies are collected in Table 2. It can be considered that the melting temperature of the Mn-eutectic is the lowest in the sample in the as-cast state (647.3 • C), the melting temperature is increased in homogenized samples, and is the highest (649.7 • C) in samples that were homogenized at 600 • C. This trend is also observed in the second eutectic (α-Al + Al 13 Fe 4 ). The melting temperature of the α-Al phase is also the lowest in the sample in the as-cast state (681.6 • C), which results from a slower and more equilibrious solidification. It is raised by a higher homogenization temperature and a longer homogenization annealing time. It is also observed that the largest melting enthalpy is in the sample in the as-cast state (−573.1 J/g). Homogenized samples have a lower melting enthalpy, where the lowest is in the sample after homogenization at 600 • C 12 h. Table 2. Listed melting and solidification temperatures, as well as corresponding enthalpies.

Sample
Heating/Melting Cooling/Solidification The liquidus temperatures of all samples are around 650 • C. At cooling, only the primary solidification and the solidification of the eutectic (α-Al + Al 13 Fe 4 ) can be observed on the cooling DSC curve, which results from a slower and more equilibrious solidification. The liquidus temperature does not change significantly with the time and temperature of the prior homogenization annealing. The highest solidification temperature of the eutectic (α-Al + Al 13 Fe 4 ) is in the sample in the cast state (625.3 • C), which decreases with increasing homogenization temperature and time. The tendency of the solidification enthalpy variation is the same as in the melting enthalpy.
These observations correlate with the fact that the effects of homogenization are greater for longer periods and higher annealing temperatures.
The optical micrographs taken from samples after various homogenization regimes are shown in Figures 4 and 5. In the as-cast sample, there is at most a finely dispersed eutectic that occurs during longer Al 13 Fe 4 needles. A slightly larger square phase of Al 15 Si 2 Mn 4 irregular forms is also observed. In homogenized samples, dispersed phases, which look like Al 13 Fe 4 , appear in the α-Al matrix. The eutectic phase Al 15 Si 2 Mn 4 gradually dissolves and disappears with longer annealing times. Those which remain become slightly larger and rounded, which is desirable due to their lower cut effect and thus their better forming properties. The longer needle phases Al 13 Fe 4 become more rounded by increasing the annealing time. Similarly, it occurs at both temperatures, but at a higher temperature of annealing ( Figure 5) the less eutectic phase is seen, which is somewhat more rounded. The content and shape of the intermetallic phases does not change significantly after eight hours of annealing.
In addition, the simulation of the homogenization was done using DSC measurement, whereas the homogenization process was analysed at two experimental temperatures, 590 • C and 610 • C, respectively. Each experiment was performed two times; the results are presented in Figure 6. The results show that when homogenization takes place at 590 • C, not all phases are dissolved in 12 h, but rather that the process is continuing, whereas the DSC curve is still dropping. At 610 • C, most of the homogenization process is finished after 160 min, and fully finished after 300 min, at 610 • C. These results are in good agreement with the microstructure results in Figures 4 and 5.   In addition, the simulation of the homogenization was done using DSC measurement, whereas the homogenization process was analysed at two experimental temperatures, 590 °C and 610 °C, respectively. Each experiment was performed two times; the results are presented in Figure 6. The results show that when homogenization takes place at 590 °C, not all phases are dissolved in 12 h, but rather that the process is continuing, whereas the DSC curve is still dropping. At 610 °C, most of the homogenization process is finished after 160 minutes, and fully finished after 300 minutes, at 610 °C. These results are in good agreement with the microstructure results in Figures 4 and 5. Furthermore, an EDS analysis of the selected samples ( Figure 7) was made in order to determine the type and phase composition of phases after certain homogenization regimes. From the EDS analysis, listed in Table 3, it can be concluded that the phases in the form of large, sharp needles contain mainly aluminium and iron, which indicates that this is the Al13Fe4 phase. Phases of smaller and rounded forms contain a slightly higher concentration of manganese and silicon, indicating that this is the Al15Si2Mn4 phase. The proportion of these two elements increases in these phases with a longer period of homogenization annealing and is in the as-cast sample up to 2 wt. %, and in the sample 580-12 as well as up to 4 mas. %.   Furthermore, an EDS analysis of the selected samples ( Figure 7) was made in order to determine the type and phase composition of phases after certain homogenization regimes. From the EDS analysis, listed in Table 3, it can be concluded that the phases in the form of large, sharp needles contain mainly aluminium and iron, which indicates that this is the Al 13 Fe 4 phase. Phases of smaller and rounded forms contain a slightly higher concentration of manganese and silicon, indicating that this is the Al 15 Si 2 Mn 4 phase. The proportion of these two elements increases in these phases with a longer period of homogenization annealing and is in the as-cast sample up to 2 wt. %, and in the sample 580-12 as well as up to 4 mas. %. Furthermore, an EDS analysis of the selected samples ( Figure 7) was made in order to determine the type and phase composition of phases after certain homogenization regimes. From the EDS analysis, listed in Table 3, it can be concluded that the phases in the form of large, sharp needles contain mainly aluminium and iron, which indicates that this is the Al13Fe4 phase. Phases of smaller and rounded forms contain a slightly higher concentration of manganese and silicon, indicating that this is the Al15Si2Mn4 phase. The proportion of these two elements increases in these phases with a longer period of homogenization annealing and is in the as-cast sample up to 2 wt. %, and in the sample 580-12 as well as up to 4 mas. %.     For samples that were homogenized at 600 • C (Figure 8), the effects of homogenization are even clearer. The needles of the eutectic phase Al 13 Fe 4 grow and become even more pronounced. Mn-phases are slightly less noticeable with a longer homogenization time. They form more roundish, their particles are smaller and accumulate in larger aggregates, whereas the concentration of manganese and silicon increases (Table 4). For samples that were homogenized at 600 °C (Figure 8), the effects of homogenization are even clearer. The needles of the eutectic phase Al13Fe4 grow and become even more pronounced. Mnphases are slightly less noticeable with a longer homogenization time. They form more roundish, their particles are smaller and accumulate in larger aggregates, whereas the concentration of manganese and silicon increases (Table 4).

Conclusions
In this research, the optimization of the homogenization annealing of EN AW 8006 alloy was performed, and the following conclusions were drawn.
In samples of EN AW 8006 alloy, the following phases can occur: α-Al; long needles from the Al13Fe4 phase; finely dispersed (α-Al + Al13Fe4) eutectic; and larger, more square forms of the Al15Si2Mn4 phase. By prolonging the time of the homogenization annealing, the finely dispersed eutectic (α-Al + Al13Fe4) is partially dissolved in α-Al, and partly forms longer, slightly more rounded needles of the Al13Fe4 phase, which is desirable due to a lower notching effect. The phase Al15Si2Mn4 also dissolves during homogenization and forms larger globules. It should be noted that the content and shape of the phases after eight hours of annealing does not change significantly at each homogenization temperature.
From the results of EDS phase analysis, it can be concluded that, in the microstructure, there is a solid solution of α-Al and the phase Al13Fe4 and Al15Si2Mn4, whereas the concentration of silicon and manganese is increased with a longer homogenization time in the Al15Si2Mn4 phase.
The homogenization annealing time could have been reduced from its current 580 °C for 12 h, wherein preheating takes 5 h, to 6 h at 600 °C, according to DSC and SEM analysis, or to 8 h at 580 °C, according to SEM analysis.

Conclusions
In this research, the optimization of the homogenization annealing of EN AW 8006 alloy was performed, and the following conclusions were drawn.
In samples of EN AW 8006 alloy, the following phases can occur: α-Al; long needles from the Al 13 Fe 4 phase; finely dispersed (α-Al + Al 13 Fe 4 ) eutectic; and larger, more square forms of the Al 15 Si 2 Mn 4 phase. By prolonging the time of the homogenization annealing, the finely dispersed eutectic (α-Al + Al 13 Fe 4 ) is partially dissolved in α-Al, and partly forms longer, slightly more rounded needles of the Al 13 Fe 4 phase, which is desirable due to a lower notching effect. The phase Al 15 Si 2 Mn 4 also dissolves during homogenization and forms larger globules. It should be noted that the content and shape of the phases after eight hours of annealing does not change significantly at each homogenization temperature.
From the results of EDS phase analysis, it can be concluded that, in the microstructure, there is a solid solution of α-Al and the phase Al 13 Fe 4 and Al 15 Si 2 Mn 4 , whereas the concentration of silicon and manganese is increased with a longer homogenization time in the Al 15 Si 2 Mn 4 phase.
The homogenization annealing time could have been reduced from its current 580 • C for 12 h, wherein preheating takes 5 h, to 6 h at 600 • C, according to DSC and SEM analysis, or to 8 h at 580 • C, according to SEM analysis.