Hot deformation behavior and rate controlling mechanism in DILUTE AL-FE-SI ALLOYS with minor additions of MNand CU


The 1xxx series of aluminum alloys are strain hardenable and offer excellent formability, corrosion resistance, and thermal and electrical conductivity. Typical applications include foil and strips for packaging, heat exchanger tubing and finstock, coaxial cable sheathing and electrical conductors [1, 2]. In general, these products are hot deformed by extrusion or rolling from direct chill (DC) cast billets or ingots. The demand for high productivity during processing leads to the requirement for an increase in hot workability to provide low flow stress with desirable final mechanical properties.
Therefore, many variants of the 1xxx series alloys are in commercial use today, balancing mechanical and physical properties along with processability. One example is the use of 1xxx series alloys in the form of hot extruded tubing for heat transfer and sheathing applications. To reduce production costs, high extrusion speeds, which are primarily controlled by the flow stress during hot deformation, are desirable. The hot working characteristics of the cast billet or ingot have a significant influence on the hot forming process, which in turn are determined by its microstructure and hot deformation regimes [3, 4]. Commercially, DC cast billets or ingots are typically homogenized prior to extrusion or rolling to improve hot workability and mechanical properties [5, 6]. In many cases, the hot workability can be optimized by proper selection of the homogenization method applied to the cast materials [7]. Prasad et al.[6] reported that a suitable homogenization treatment, which dissolved the intermetallic particles present at the grain boundaries, improves hot workability of as-cast AZ31 alloy by reducing the intercrystalline cracking and flow instability regimes. Totik et al.[7] showed that homogenization at 510 oC significantly improved the hot workability of AA 2014 alloy and resulted in the uniform workability of the tested ingots. Zhang et al.[3] studied the influence of the initial microstructure on thehot workability of a AA 2219 aluminum alloy and found that the as-aged microstructure produced the lowest hot deformation resistance and the as-solutionized microstructure showed the highest deformation resistance under the same deformation conditions. Liu et al.[8] investigated the effect of cooling rate after homogenization on the elevated temperature flow behavior of AA 7050 aluminum alloy and found that furnace-cooled alloy exhibited higher hot workability than water-quenched alloy. However, there is very limited information in the literature on the effectiveness of the homogenization treatment in 1xxx alloy production. The 1xxx series alloys contain at least 99.0% aluminum with iron and silicon as the main alloying elements or impurities (all alloy compositions are given in wt% unless otherwise indicated). Because the solid solubility of iron in aluminum is very low (i.e., max. 0.05% at 650°C), most of the iron combines with both aluminum and silicon to form secondary intermetallic phases [9, 10]. The equilibrium intermetallic Al3Fe phase can form at slow solidification rates. However, depending on the alloy composition, cooling rate and presence of trace elements, a wide range of intermediate intermetallic phases, such as AlmFe, Al9Fe, Al6Fe, AlxFe and α-AlFeSi (Al8Fe2Si), can form and are considered metastable Fe-rich phases in the literature [9-11]. The metallurgical performance of 1xxx series alloys during hot deformation can be influenced by secondary intermetallic phases in terms of their type, size and distribution. Although there are abundant studies on the formation of iron-rich intermetallic phases during solidification of 1xxx series aluminum alloys [12], studies of their evolution during the homogenization treatment are limited. Kosuge et al. [13] and Griger et al. [14] reported a solid-state phase transformation from AlmFe, Al6Fe and α-AlFeSi to Al3Fe in dilute Al-Fe-Si alloys. Li et al. [15] also observed the phase transformation from Alm(FeMn) to Al3(FeMn) in AA5182 after homogenization at 520°C.
In the current study, the microstructural evolution of two typical dilute Al-Fe-Si alloys during homogenization was investigated using optical and electron microscopy as well as electric conductivity measurements. The effect of the homogenization parameters on the intermetallic phase transformation and solid solution levels was studied. The hot deformation behavior under different homogenization conditions was evaluated using the uniaxial hot compression test at various deformation temperatures and strain rates.

Experimental Procedure

Two dilute Al-Fe-Si alloys with chemical compositions shown in Table 3.1 were selected for investigation. Both alloys were prepared from commercially pure aluminum (99.7%), Al-20%Fe and Al-50%Si master alloys. For each alloy, approximately 10 kg of material was melted in an electrical resistance furnace. The melting temperature was kept at about 750 oC. The molten metals were grain-refined by addition of 0.015% Ti in the form of an Al-5Ti-1B master alloy. The melts then cast into a rectangular permanent steel mold measuring 30×40×80 mm3. The chemical compositions of each alloy were measured using optical emission spectroscopy (OES) technique.
Cast ingots were homogenized at 550, 590 and 630°C for various soaking times (2, 6 and 12 h) followed by water quenching. All homogenization treatments were conducted in an air circulating electrical furnace. As-cast and homogenized samples were taken from the same ingot position for metallographic examination to eliminate the influence of the cooling rate.
The microstructures of the as-cast and homogenized samples were examined using optical microscopy (OM) and scanning electron microscopy (SEM, JEOL-6480LV) equipped with an electron backscatter diffraction (EBSD) system. Because iron-rich intermetallic phases have different morphologies [16, 17] and the stoichiometry of iron-rich phases are not always fixed, the EBSD technique was employed to provide unambiguous characterization of the iron-rich phases. The crystallographic data for all of the binary Al- Fe and ternary Al-Fe-Si compounds listed in the Pearson’s Handbook [18] were entered into a customized Channel 5 software database for phase identification. A total of 12 fields with dimensions of 150 μm x 150 μm were observed in each sample. All of the Fe-rich intermetallic particles in each field were identified by the EBSD analysis and more than 100 particles in each sample were characterized. After identifying all of the intermetallic particles appearing in each field, the volume fraction of each phase was quantified using the CLEMEX image analyzing system.
The electrical conductivity was measured using an AC electrical conductivity instrument (SIGMASCOPE SMP 10) on the polished rectangular ingot blocks measuring 50x35x25 mm3. This machine measures the electrical conductivity using the eddy current method. All tests were conducted at frequency of 60 kHz. In order to consider the influence of temperature on the electrical conductivity of the samples the temperature of the specimens was directly measured by one sensor while other sensor measured the electrical conductivity. SIGMASCOPE SMP 10 gives electrical conductivity of different metals and alloys in comparison with international annealed copper standard (%IACS). IACS is 100% for annealed copper standard. For each condition, six electrical conductivity measurements were performed to provide an average value. Cylindrical specimens of 10 mm diameter and 15 mm in height were machined for hot compression testing. The uniaxial hot compression tests were conducted using a Gleeble 3800 thermo-mechanical testing unit at strain rates of 0.01 and 1 s-1 and deformation temperatures of 400 and 500°C. During the tests on the Gleeble 3800 unit, specimens were heated at a rate of 2 °C/s and maintained for 120 s at the desired temperature to ensure a homogeneous temperature distribution. The specimens were deformed to a total true strain of 0.8 and then immediately water-quenched to room temperature.

Results and Discussion

As-cast and homogenized microstructures

Figure 3.1 shows the optical microstructures of two Al-Fe-Si alloys in the as-cast condition, which consists of α-Al dendrites (matrix) and different Fe-rich intermetallic phases distributed along the aluminum dendrite boundaries. The mean aluminum grain sizes were 110 and 120 μm for the low Si alloy (Al-0.3Fe-0.1Si) and high Si alloy (Al- 0.3Fe-0.25Si), respectively. In general, four different types of Fe-rich intermetallics including equilibrium Al3Fe and metastable AlmFe, α-AlFeSi and Al6Fe phases were observed in the cast microstructure. In the low Si alloy, the dominant phases were the feather-like AlmFe and acicular Al3Fe (Fig. 3.1a), and in the high Si alloy, the major phases were Chinese script α-AlFeSi and acicular Al3Fe (Fig. 3.1b).
During homogenization at the temperature range studied (550 to 630 °C), the metastable AlmFe and α-AlFeSi phases became unstable and transformed to the equilibrium Al3Fe phase. Figure 3.2 shows an example of the microstructure after homogenization at 630 °C for 12 h for both alloys. The feather-like AlmFe was completely changed to acicular and plate-like Al3Fe in the low Si alloy (Fig. 3.2a), while most of the Chinese script α- AlFeSi was transformed to plate-like Al3Fe and only a few spheroidized α-AlFeSi remained in the high Si alloy (Fig. 3.2b).

Evolution of iron-rich intermetallic phases during homogenization

To better quantify their evolution during heat treatment, all of the Fe-rich intermetallic particles were first identified by EBSD and then the volume fraction of each phase was quantified by image analysis [19, 20]. Figure 3.3 shows typical morphologies and their EBSD solutions for the four different intermetallics observed in the two alloys in the ascast and homogenized conditions. In the EBSD analysis, the mean angular deviation (MAD) between the experimental and simulated patterns represents the accuracy of the solution. A smaller value indicates a closer match between the experimental and simulated results. Typically, a MAD value smaller than 0.7 is considered to be critical for an accurate solution [21]. The MAD values obtained for AlmFe, Al6Fe, Al3Fe and α-AlFeSi phases presented in figure 3.3, were approximately 0.2, 0.14, 0.18 and 0.23, respectively, which are much lower than 0.7 and confirm the reliable and accurate identification of all four phases. The quantitative results for phase evolution as a function of homogenization temperature are shown in Figure 4. In the low Si alloy (Fig. 3.4a), the as-cast microstructure was dominated by both AlmFe and Al3Fe with a minor amount of α-AlFeSi. After homogenization at 550 °C for 6 h, AlmFe decomposed and completely transformed to Al3Fe while the α-AlFeSi remained unchanged. After increasing the temperature to 590 °C, α-AlFeSi dissolved and transformed to Al3Fe. Therefore, at higher homogenization temperatures (590 to 630 °C), Al3Fe is the only intermetallic phase in the microstructure. For the high Si alloy (Fig. 3.4b), the as-cast microstructure consisted of a major phase of α-AlFeSi and two minor phases consisting of Al3Fe and AlmFe. During homogenization at 550 °C for 6 h, all of the AlmFe was converted to Al3Fe but the α-AlFeSi remained stable. The major α-AlFeSi phase only began to decompose at 630°C. After 6 h, most of the α-AlFeSi particles transformed to Al3Fe. Therefore, the major α-AlFeSi co-existed with Al3Fe at 550 and 590 °C, and Al3Fe become the predominant phase at 630 °C. In both alloys, Al6Fe was present as trace amounts of isolated particles (i.e., less than 3% of the constituent particles) in the as-cast microstructure. As reported by Griger et al. [14] and Tezuka et al. [22], the use of Al-Ti-B as a grain refiner can promote the formation of AlmFe over Al6Fe. During homogenization, the Al6Fe phase was stable at 550 °C and disappeared after homogenization at 590 °C in both alloys.

Solid-state transformation from AlmFe to Al3Fe

Figure 3.5 shows the solid-state transformation from AlmFe to Al3Fe in detail at 550°C for the low Si alloy, starting with the cast microstructure in Fig. 3.5a where the AlmFe phase formed during solidification. During homogenization at 550°C for 2 h (Fig. 3.5b and c), AlmFe first spheroidized and then dissolved in the aluminum matrix. Simultaneously, plate-like Al3Fe particles precipitated either in the aluminum matrix close to the original AlmFe phase or on the pre-existing Al3Fe phase formed during solidification. It is apparent that the solid-state transformation from AlmFe to Al3Fe proceeds via a dissolutionreprecipitation mechanism where iron is required to diffuse from the dissolving AlmFe to the precipitating Al3Fe particles. Therefore, the controlling factor for the phase transformation is the diffusion of iron between the two species. By increasing the homogenization time to 6 h (Fig 3.5d), all of the AlmFe transformed to Al3Fe, which continued to coarsen into large plate-like particles. Homogenization at a higher temperature or for a longer time, as shown in Figure 3.2 (a), promoted the formation of coarse blockshaped or plate-like Al3Fe, which most likely result from phase ripening and coalescence.

Solid-state transformation from α-AlFeSi to Al3Fe

Figure 3.6 illustrates the microstructural evolution during homogenization at 630°C and shows the solid-state transformation from α-AlFeSi to Al3Fe in the high Si alloy. In the as-cast condition, the Chinese script α-AlFeSi particles were located along the aluminum dendrite boundaries (Fig. 3.6a). During homogenization at 630 °C for 2 h, α-AlFeSi firstcoarsened and spheroidized (Figs. 3.6b and c) and finally dissolved releasing Fe and Si atoms into the solid solution. Simultaneously, Al3Fe precipitated either in the vicinity of the original α-AlFeSi (Fig. 3.6c) or on the pre-existing Al3Fe particles. The α-AlFeSi to Al3Fe transformation also proceeded via a dissolution-reprecipitation mechanism. The phase transformation rate is believed to be controlled by the growth rate of the Al3Fe phase, which in turn is controlled by the diffusion of Fe in the aluminum matrix from the dissolving α-AlFeSi to the precipitating Al3Fe. After homogenization for 12 h, most of α-AlFeSi had completely transformed and Al3Fe continued to coarsen into large plate-like particles. However, a small amount of spheroidal α-AlFeSi still remained in the microstructure (Fig. 3.6d). For the low Si alloy, all of the α-AlFeSi particles transformed to Al3Fe at 590 °C, whereas in the high Si alloy, α-AlFeSi was stable at this temperature and only started to transform at 630°C. Due to the higher bulk silicon content, α-AlFeSi in the high Si alloy is more stable than that in the low Si alloy.

Solid solution levels

Solid solution levels can significantly contribute to the high temperature flow stress in dilute aluminum alloys [5, 23]. However, due to low solubility of iron in aluminum and also low amount of Si concentration in the investigated alloys it is very hard to directly measure the solute level of these elements in the aluminum matrix by EDX technique precisely. Furthermore, using XRD technique to measure the solute levels of iron and silicon based on the variation of lattice parameters is technically challenging, because it needs an advance XRD machine and also reliable standard. The increase or decrease in electrical conductivity is widely used to follow the dissolution and precipitation of solutes in aluminum alloys [24, 25] and was used as an index of solid solution levels of Fe and Si in this study.
The impact of the homogenization treatment on the electrical conductivity of the two alloys is shown in Figure 3.7. In comparison to the low Si alloy, the high Si alloy generally exhibited lower conductivity for a given condition, indicating overall higher levels of solute in solution. During the rapid solidification in the permanent mold casting, which has a similar cooling rate range to DC cast ingots, some Fe and Si atoms are retained in the supersaturated solid solution. Homogenization at 550°C significantly increased the electrical conductivity of both alloys compared to the as-cast condition due to the elimination of supersaturation as excess Fe and Si were precipitated as intermetallics. As the homogenization temperature was increased to 630 °C, the conductivity decreased to levels similar to the as-cast condition, indicating the dissolution of Fe and Si containing constituent phases.
In the low Si alloy, the phase transformation of AlmFe to Al3Fe occurred at 550 °C and reached completion with increased homogenization time. However, the conductivity changed only slightly as the homogenization time increased at this temperature (Fig. 3.7a), which indicated that this type of phase transformation did not significantly affect the solid solution levels. However, by increasing the homogenization temperature from 550 to 590 °C, the transformation of α-AlFeSi to Al3Fe was promoted resulting in a release of Si atoms into the solution. Therefore, the decrease in conductivity was more significant. An increase in the temperature to 630 °C further decreased the conductivity, which corresponded to the re-dissolution of the constituent phases and further release of Fe and Si into the solution due to increased solid solubility at a higher temperature.
As shown in Fig. 3.7b, the high Si alloy only exhibited a slight decrease in electrical conductivity from 550 to 590 °C because no phase transformation occurred in this temperature range (Fig. 4b). However, a more significant drop in conductivity was observed when the temperature was increased from 590 to 630 °C because the transformation of α-AlFeSi to Al3Fe released free Si atoms into the solution and the solid solubility increased with temperature. After soaking at 630 °C, the conductivity continued to decrease as the homogenization time increase, and after 12 h the conductivity of the high Si alloy was actually lower than that at the as-cast condition, which indicated that the solid solution levels progressively increased as the homogenization time increased.
The results for both alloys in Fig. 3.7 indicate that the solid solution levels initially decreased during homogenization at 550 °C and then increased at higher homogenization temperatures. It is also evident that the AlmFe to Al3Fe phase transformation had little influence on the solid solution levels, and the phase transformation of α-AlFeSi to Al3Fe significantly increased solute levels in the solution during homogenization.

Flow stress behavior during hot deformation

The hot workability of two dilute Al-Fe-Si alloys was assessed by a series of flow stress curves. Figures 3.8 and 3.9 show the true stress-true strain curves for both alloys as a function of homogenization temperature for a 6 h soak time compared to the as-cast starting material. For most of the test conditions, the peak flow stress was followed by a steady state region, as shown in Figures 3.8 and 3.9. However, at the highest Z condition (=1 s-1 and T=400 °C), the flow stress continued to increase with increasing strain. The former case occurs when dynamic softening is in balance with work hardening while the latter phenomenon is indicative of work hardening being stronger than softening during deformation. In general, such flow stress behaviors are characteristic of hot working where dynamic recovery is the main softening mechanism [5]. For the low Si alloy (Fig. 3.8), the flow stress behavior was very similar for all four deformation conditions. The as-cast material always exhibited the highest flow stress values. Homogenization at 550 °C yielded the lowest flow stress values. The flow stress gradually increased when the homogenization temperature was increased to 590 and 630 °C. For the high Si alloy (Fig. 3.9), the highest overall value of flow stress was obtained after homogenization at 630°C followed by the as-cast condition. The minimum values of flow stress were still obtained after homogenization at 550 °C.
Typical flow stress values at a strain of 0.8 are shown in Figure 3.10 as a function of the homogenization temperature for deformation conditions where =1 s-1 and T=400 and 500°C. The two alloys exhibited a similar trend as a function of the homogenization temperature. Homogenization decreased the flow stress compared to the as-cast condition, and treatment at 550°C yielded the minimum flow stress for both alloys. This result is consistent with the supersaturated state of the as-cast ingot and the maximum precipitation of the solute at 550°C. However, the flow stress increased again at higher homogenization temperatures as the solute levels increased. In general, the high Si alloy exhibited higher flow stress values (4-11%) under all deformation conditions due to the higher solute content.
In the low Si alloy, an increase in the soak temperature from 550 to 630°C produced a 10 to 23% increase in the flow stress over the range of deformation conditions due to increased Fe and Si levels in the solid solution. For the same range of conditions, the flow stress of the high Si alloy increased by 15 to 45%. This larger increase in the flow stress was associated with the higher proportion of α-AlFeSi, which promoted the release of more Si solute during the transformation to Al3Fe.


The Fe and Si present in the 1xxx DC cast ingots as deliberate alloy additions or impurities are generally present in two forms: (1) constituent particles that are Fe-rich intermetallics distributed along aluminum cell and grain boundaries and (2) solute elements that are in the aluminum solid solution. Both constituent particles and solute elements can have significant impacts on the hot workability of aluminum alloys [5]. Homogenization of DC cast ingots prior to extrusion or rolling often results in significant changes in the type, size and distribution of constituent particles as well as the solute levels in the aluminum matrix, which in turn, can alter the hot workability of a given alloy.

Effect of constituent particles

Al-Fe and Al-Fe-Si intermetallic (constituent) particles in dilute Al-Fe-Si alloys were originally eutectic secondary phases that solidified around aluminum cells and grains. In general, when intermetallic particles are relatively fine (<1 μm) and uniformly distributed in the matrix, they can effectively pin dislocations, stabilize subgrains and retard dynamic recovery and recrystallization, leading to a large increase in the flow stress during the hot deformation process[5]. In the current study, the size of all of the intermetallic particles in the as-cast and homogenized conditions were quite large (i.e., from one to tens of μm, as shown in Figs. 3.1 and 3.2), and the total volume fractions of the intermetallic particles were also quite low (i.e., approximately 1.3 and 1.5 vol.% in the low Si and high Si alloys, respectively). Therefore, the low volume fraction, large size and large interparticle spacing of the intermetallic particles limit their overall pinning ability. During homogenization at different temperatures, phase transformations (AlmFe to Al3Fe and α-AlFeSi to Al3Fe) occurred, and the type and size of intermetallic particles were partially modified. However, the size of the particles after homogenization was still large (1-20 μm), and their location in the microstructure remained unchanged (Fig. 3.2).
Therefore, it is reasonable to conclude that the contribution to high temperature flow stress from the various intermetallic particles and their distributions produced by changes in the alloy composition and homogenization treatment is not significant.

Impact of solute elements

Solute elements can have a strong influence on the hot deformation behavior of aluminum alloys by interacting with mobile dislocations and retarding softening processes [5, 23, 26]. The solution strengthening effect of a solute is strongly dependent on the specific solvent/solute pair [27]. The rapid solidification of the permanent mold casting used in this study resulted in the retention of Fe and Si atoms in the supersaturated solid solution. Therefore, the as-cast samples typically exhibited the highest flow stress in both alloys. Homogenization at a lower temperature (550 °C) significantly reduced the flow stress compared to the as-cast condition due to the elimination of Fe and Si supersaturation. However, an increase in the homogenization temperature from 550 to 630 °C increased the flow stresses by 10-23% and 15-45% for the low and high Si alloys, respectively (Figs. 3.8 and 3.9). The increase in flow stress is believed to be closely related to the high solid solution levels of both Fe and Si produced at a higher homogenization temperature. Figure 3.11 shows the Al-Fe phase diagram where the solubility of iron increases rapidly as the temperature increases [28]. The iron solid solubility limits for the Al-Fe binary system at 550, 590 and 630 °C are approximately 0.012, 0.024 and 0.04%, respectively (i.e., the iron solute content can be more than tripled over this temperature range). Although iron typically has low solid solubility in aluminum (i.e., several hundred ppm), it is a potent solute that has a significant influence on dislocation movement due to its low diffusivity [29]. Marshall et al. [30] reported that very small concentrations of iron in a solid solution of high purity aluminum may suppress softening processes. Shelby et al. [29] also demonstrated that a small amount of iron in solution (i.e., at the level of several hundred ppm) could significantly improve the creep resistance of pure aluminum. In addition to the effect of iron solid solubility, the phase transformation from α- AlFeSi to Al3Fe at higher temperatures released Si atoms to the matrix and increased the overall solute level, which augmented the flow stress. For the high Si alloy, α-AlFeSi was the dominant intermetallic phase, and it occupied approximately 60% of the total intermetallic volume fraction prior to the transformation (Fig. 3.4b). The phasetransformation of α-AlFeSi to Al3Fe at 630 °C produced a significant increase in the silicon content in the solid solution (Fig. 3.7b), which resulted in a significant increase in the flow stress that exceeded the as-cast value.

Industrial aspect

Based on industrial experience, the as-cast microstructure is undesirable for hot deformation processes (e.g., high speed hot extrusion of thin wall tubing) due to the inherent high solid solution levels, which promote high flow stress, and the coarse intermetallic structures, which can cause hot ductility related surface defects, such as die lines or pick-up. The selection of homogenization parameters (i.e. temperature and time) should consider both aspects. Homogenization of dilute Al-Fe-Si alloys induces a progressive transformation either from AlmFe to Al3Fe at lower temperatures or from α-AlFeSi to Al3Fe at higher temperatures. In both cases, extended homogenization times can promote coarsening of Al3Fe, which may be detrimental to the hot ductility and surface finish quality of extruded products.
The results of the current study indicate that the solute levels of Fe and Si are the key factors controlling the flow stress during hot deformation of dilute Al-Fe-Si alloys. Homogenization at a lower temperature (550 °C) eliminated the supersaturation of solute elements and therefore greatly reduced the flow stress. However, an increase in the homogenization temperature cannot further soften the materials, as traditionally believed.
On the contrary, an increase in the homogenization temperature significantly increased the flow stress up to 23 and 45% from 500 to 630 °C for the low and high Si alloys, respectively, due to the increase in the solid solution levels at higher temperatures. Increases of this magnitude in the flow stress are highly significant for the extrusion process where the productivity is typically limited by the press capacity. One of the key factors affecting extrusion productivity of dilute Al-Fe-Si alloys is the high temperature flow stress. In general, an extrusion press has a specific pressure available to push the billet through the die and this is a direct function of the flow stress. Aluminum alloys exhibit strain rate sensitivity, namely the higher the strain rate (the ram speed), the higher the flow stress is. Therefore, for a given press capacity and profile geometry, the flow stress of the material will dictate the maximum ram speed at which the press can operate. A higher flow stress also produces a larger temperature increase during extrusion, which increases the surface exit temperature for a given extrusion speed resulting in an earlier onset of surface defects.
If the minimum flow stress is the primary consideration, the use of a relatively low homogenization temperature, such as 550°C, offers significant potential benefits in terms of the extrusion speed and productivity for dilute Al-Fe-Si alloys. A reduction in the silicon level can also be beneficial. In comparison to the low Si alloy, the high Si alloy exhibited higher flow stresses (4 to 11%) for a given deformation condition. This further increase in flow stress is also significant in term of the processing speed. This effect was further compounded at high homogenization temperatures by the release of Si during the α-AlFeSi to Al3Fe transformation, which resulted in a further increase in flow stress. Commercially, an increase in Si in the alloy is often utilized to increase the room temperature strength.
Therefore, there is a potential trade-off between these two aspects. If the final mechanical property requirement dictates the use of a high silicon alloy, the proper selection of homogenization temperature and time should be considered to avoid the phase transformation from α-AlFeSi to Al3Fe.


1) The as-cast microstructures of two dilute Al-Fe-Si alloys (Al-0.3Fe-0.1Si and Al- 0.3Fe-0.25Si) consisted of α-Al dendrites and metastable AlmFe and α-AlFeSi, as well as equilibrium Al3Fe intermetallic particles. The proportion of α-AlFeSi intermetallic increased with a higher silicon content.
2) Homogenization promoted the phase transformation from the metastable AlmFe or α-AlFeSi phase to the equilibrium Al3Fe phase via a dissolution-reprecipitation mechanism. The AlmFe dissolved and transformed completely at 550°C in both alloys. The α-AlFeSi was transformed at 590°C in the low Si alloy (Al-0.3Fe-0.1Si), whereas it began to decompose and transform to Al3Fe at 630°C in the high Si alloy (Al-0.3Fe-0.25Si).
3) Homogenization at 550°C significantly reduced the solid solution levels in both alloys due to the elimination of the supersaturation originating from the cast ingot. Above 550°C, the solid solution levels progressively increased.
4) The flow stress behavior of dilute Al-Fe-Si alloys was primarily controlled by the solute levels of Fe and Si. Homogenization at 550°C produced the lowest flow stress for all of the deformation conditions studied. An increase in the homogenization temperature from 550 to 630°C increased the flow stress by 10 to 23% and 15 to 45% for the low Si and high Si alloys, respectively, which is commercially significant in terms of the productivity during hot forming processes.
5) An increase in the silicon content from 0.10 to 0.25% in dilute Al-0.3Fe-Si alloys increased the solid solution levels for all of the homogenized conditions studied, which resulted in an increase in the overall flow stress by 4 to 11%.



The hot deformation behavior of dilute Al-Fe-Si alloys (1xxx) containing various amounts of Fe (0.1 to 0.7 wt%) and Si (0.1 to 0.25 wt%) was studied by uniaxial compression tests conducted at various temperatures (350-550 °C) and strain rates (0.01-10 s-1). The flow stress of the 1xxx alloys increased with increasing Fe and Si content. Increasing the Fe content from 0.1 to 0.7% raised the flow stress by 11-32% in Al-Fe-0.1Si alloys, whereas the flow stress increased 5-14% when the Si content increased from 0.1 to 0.25% in Al-0.1Fe-Si alloys. The influence of the temperature and the strain rate on the hot deformation behavior was analyzed using the Zener-Holloman parameter, and the effect of the chemical composition was considered in terms of the materials constants in the constitutive analysis. The proposed constitutive equations yielded an excellent prediction of the flow stress over wide ranges of temperature and strain rate with various Fe and Si contents. The microstructural analysis results revealed that the dynamic recovery is the sole softening mechanism of the 1xxx alloys during hot deformation. Increasing the Fe and Si content retarded the dynamic recovery and resulted in a decrease in the subgrain size and
mean misorientation angle of the boundaries.


The 1xxx series aluminum alloys are primarily used for applications in which superior formability and excellent thermal and electrical conductivity are required. Typical applications include foil and strips for packaging, heat-exchanger tubing and fin stock, coaxial cable sheathing and electrical conductors [1, 2]. These products are generally subjected to hot-forming processes such as extrusion and rolling. Therefore, the development of a method to analyze and predict their hot deformation behavior under various thermomechanical conditions is the primary goal. The high-temperature flow behavior of various materials in hot-forming processes is very complex. The work hardening and dynamic softening are both significantly affected by many factors, such as the chemical composition, the forming temperature, the strain rate and the strain [3, 4]. The flow behavior of materials is very important for the design of hotforming processes due to its substantial impact on the required deformation load as well as the kinetics of metallurgical transformations. Traditionally, the trial and error method has been employed to optimize the thermomechanical processes. To overcome the huge number of tests required to achieve a reliable conclusion in the trial and error practice, various modeling techniques have been developed that permit a significant reduction in the production cost.
The modeling of materials flow behavior is often conducted by proper constitutive equations, which correlates the dynamic material properties such as the flow stress to the process parameters such as the deformation temperature and strain rate [5-7]. Normally, uniaxial hot compression tests are employed to provide the necessary data to extract the constitutive equations. Several analytical [8], phenomenological [9], and empirical [10] models have been proposed to describe the high temperature flow behavior for a wide range of metals and alloys. Johnson and Cook [11] proposed a phenomenological model to develop a cumulative-damage fracture model. Sellars et al. [12] proposed a hyperbolic-sine constitutive law to describe the elevated temperature flow behavior of various materials. Sloof et al. [13] introduced a strain-dependent parameter into the hyperbolic sine constitutive equation to improve its accuracy. Lin et al. [14] proposed a revised hyperbolic sine constitutive equation to describe the flow behavior of 42CrMo steel by considering the compensation of the strain and strain rate. Ashtiani et al. [6] established strain-compensated constitutive equations to predict the flow behavior of commercially pure aluminum. Among various constitutive equations available, the hyperbolic sine constitutive equation, proposed by Sellars et al. [12], has proven to be applicable over a wide range of materials and alloys [5-7, 15, 16].
Commercial 1xxx aluminum alloys exhibit higher strength and work hardening than high purity aluminum. The main alloying additions, or controlled impurities, in these alloys are Fe and Si. Zhao et al. [17] reported that Fe and Si play a major role in the strength and work hardening of commercial 1xxx alloys and that the contribution from the other impurities is negligible. McQueen et al. [4, 18, 19] concluded that dynamic recrystallization (DRX) could not occur in commercially pure aluminum (1xxx alloys) and that dynamic recovery (DRV) was the sole restoration mechanism during hot deformation. Although a few researchers [4, 6, 19] have studied the hot deformation behavior of commercially pure aluminum, no systematic investigation of the influence of different Fe and Si contents on the hot deformation behavior of dilute Al-Fe-Si alloys is available in the literature. In the present study, the hot deformation behavior of dilute Al-Fe-Si alloys with a systematic variation of the Fe and Si contents was investigated by hot compression tests conducted at various deformation temperatures and strain rates. The experimental stressstrain data were employed to drive constitutive equations correlating flow stress, deformation temperature and strain rate considering the influence of the chemical composition. Moreover, the effects of the deformation conditions and the chemical composition on the microstructural evolution associated with the dynamic softening were investigated.

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Table des matières

1.1 Background and problem definition
1.2 Objectives
2.1 Hot working
2.1.1 Work Hardening
2.1.2 Solid solution hardening
2.1.3 Precipitation hardening
2.1.4 Softening Mechanisms Dynamic Recovery Dynamic Recrystallization
2.1.5 Dynamic recovery flow curves
2.1.6 Dynamic recrystallization flow curves
2.2 AA1xxx series aluminum alloys
2.2.1 Microstructure of AA1xxx series aluminum alloys Binary Al-Fe phases Ternary Al-Fe-Si phases Transformation of metastable phases
2.2.2 Hot deformation behavior of AA1xxx aluminum alloys Substructure evolution during hot deformation Effect of homogenization Effect of Mn Effect of Cu
2.3 Modeling the high temperature flow behavior of metals and alloys
2.3.1 Phenomenological constitutive equations
2.3.2 Artificial neural network (ANN)
3.1 Introduction
3.2 Experimental Procedure
3.3 Results and Discussion
3.3.1 As-cast and homogenized microstructures
3.3.2. Evolution of iron-rich intermetallic phases during homogenization Solid-state transformation from AlmFe to Al3Fe Solid-state transformation from α-AlFeSi to Al3Fe
3.3.3. Solid solution levels
3.3.4. Flow stress behavior during hot deformation
3.3.5. Discussion Effect of constituent particles Impact of solute elements Industrial aspect
3.4 Conclusions
4.1 Introduction
4.2 Experimental
4.3 Result and discussion
4.3.1 Flow stress behavior
4.3.2 Constitutive analysis
4.3.3 Effect of chemical composition
4.3.4 Microstructural evolution during hot deformation
4.4. Conclusions
5.1 Introduction
5.2 Experimental procedure
5.3 Results and discussion
5.3.1 Flow stress behavior
5.3.2 Constitutive analyses
5.3.3 Microstructural evolution during hot deformation
5.3.4 Discussion
5.4 Conclusions
6.1 Introduction
6.2 Experimental procedures
6.3 Results and discussion
6.3.1 Effect of Cu content on flow stress behavior
6.3.2 Development of artificial neural network model Effect of Cu addition Effect of temperature Effect of strain rate Assessment of the proposed model Sensitivity analysis
6.4 Conclusions
Appendix 6.A: Example illustrating Garson’s algorithm
7.1 Conclusions
7.2 Recommendations

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