The effect of homogenization practice on the ...

Journal of Materials Processing Technology 148 (2004) 250–258

The effect of homogenization practice on the microstructure of AA6063 billets Yücel Birol∗ Materials and Chemical Technologies Research Institute, Marmara Research Center, TUBITAK, Gebze, Kocaeli, Turkey Received 20 March 2003; accepted 16 January 2004

Abstract It is well established that homogenized billets extrude easier and faster and give better surface finish and higher tensile properties than as-cast billets. Hence, the production of Al–Mg–Si extrusions from DC-cast billets almost always starts with a homogenization cycle which typically consists of a soaking treatment followed by cooling at a predetermined rate. While soaking is performed to produce a homogeneous solid solution and to transform the ␤-AlFeSi particles to the more acceptable ␣ variety, the motivation behind controlled cooling is heterogenization. The cooling cycle has to be designed to reprecipitate as much Mg2 Si as possible, in a form and size easily redissolvable during subsequent processing for high extrudability and surface quality. The effect of various soaking and cooling cycles on the formation and transformation of insoluble and soluble constituents was investigated in order to identify the optimum homogenization practice for a semicontinuous DC-cast AA6063 billet. A homogenization practice with a 6 h soak at 580 ◦ C followed by step-cooling at 250–300 ◦ C was found to be optimum. Step-cooling gave a more complete depletion of the aluminum solid solution, i.e. lower flow stress, without forming coarse and stable Mg2 Si particles which are very difficult to solutionize during reheating and thus survive the extrusion process impairing both surface quality and mechanical properties. © 2004 Elsevier B.V. All rights reserved. Keywords: Homogenization; Extrusion; AA6063; Microstructure

1. Introduction It is well established that homogenized billets extrude easier and faster and give better surface finish and higher tensile properties than as-cast billets [1]. However, there are limits to how much a homogenization treatment can improve the performance of the billet on the press. There are several features that must be imparted to the billet before homogenization, in the cast house. The billet chemistry has to be optimized in accordance with the needs and expectations of the extruder. A cast billet should also have fine grain and cell structure uniformly distributed over its cross-section, no coarse script intermetallics, minimum surface segregation and shell zone and an acceptable surface finish [2]. When properly homogenized, such a billet will give high throughput rates, low breakout pressures, desired shape within dimensional tolerances, uniform surface finish free from streaking, pick-up, die-lines as well as high tensile properties. ∗ Present address: The Scientific and Technical Research Council of Turkey, Marmara Research Centre, PO Box 21, 41470 Gebze, Kocaeli, Turkey.

0924-0136/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2004.01.056

An industrial homogenization practice typically consists of a soaking treatment which is followed by cooling at a predetermined rate [3]. Several processes take place during soaking [4,5]. The interdendritic network of the plate-like ␤-AlFeSi intermetallics are replaced by the more rounded discrete ␣-AlFeSi particles. The ␤ → ␣ transformation is important as the ␤-AlFeSi particles are often held responsible for a number of surface defects, limiting extrudability [5–9]. The Mg2 Si particles and coarse eutectics are also solutionized in a soaked billet [5,10,11] while the coring inherited from the as-cast microstructure is leveled out in favor of a more homogeneous distribution of the solute Mg and Si. Unfortunately, a fully solutionized billet as obtained at the end of soaking, is difficult to extrude. So, in contrast to soaking which is undertaken to achieve homogenization, the motivation behind controlled cooling is heterogenization [12]. Slow cooling tends to produce coarse Mg2 Si particles while rapid cooling traps the Mg and Si in solution with little or no Mg2 Si precipitation [13–15]. Coarse ␤-Mg2 Si particles are difficult to solutionize at common preheating temperatures, resist solutionizing at high heating rates and may even survive the extrusion process. Such particles lead to incipient melting and surface tearing during extrusion,

Y. Birol / Journal of Materials Processing Technology 148 (2004) 250–258

giving poor surface quality and are also responsible for poor tensile properties. Mg and Si in solution, on the other hand, have an adverse effect on the flow stress of the alloy and thus increase its resistance to deformation. So, the cooling cycle has to be designed to reprecipitate as much Mg2 Si as possible, in a form and size easily redissolvable during subsequent processing. The present work was undertaken to investigate the effect of various soaking and cooling cycles on the formation and transformation of insoluble and soluble constituents and to identify the optimum homogenization practice for a semicontinuous DC-cast AA6063 billet.

2. Experimental The AA6063 alloy used in the present investigation was cast industrially with a vertical DC caster in the form of 6 m long billets with a diameter of 152 mm. Its chemical composition is given in Table 1. Due to the very low Mn content, the present alloy relied solely on the soaking practice for the ␤ → ␣ transformation. The Mg2 Si content of the present alloy suggests that it was produced for high extrudability rather than high strength. Homogenization trials were performed first by changing the soaking parameters but employing a constant cooling rate and then by changing the cooling practice but keeping the soak time and temperature constant (Fig. 1). Samples for both soaking and cooling trials, 65 mm long,

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Table 1 The chemical composition of the AA6063 alloy used in the present investigation (wt.%) Si Mg Fe Mn Cu Ti Al Siex Mg2 Si

0.429 0.446 0.218 0.004 0.003 0.017 98.80 0.097 0.704

35 mm wide and 10 mm thick, were sectioned from a transverse slice of the billet, at least 10 mm away from the surface, in each case, to avoid possible microstructural and compositional variations from one sample to the other. Soaking experiments were performed in the temperature range 540–580 ◦ C with a 2–6 h soak (Fig. 1a). One of the two sets of soaked samples were air-quenched as the focus was on the evolution of the insoluble constituents. The second set of samples were cooled to approximately 200 ◦ C at a rate of 200 ◦ C/h to facilitate precipitation of the Mg2 Si phase in order to evaluate the precipitation pattern and judge the homogenization of the solid solution matrix. Soaked samples were prepared with standard metallographic techniques: ground with SiC paper, polished with 3 ␮m diamond paste and finished with colloidal silica. They were examined after etching with a 0.5% HF solu-

Fig. 1. Soaking and cooling cycles employed in the present work.

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tion using a Olympus BX51M model optical microscope. The XRD patterns were recorded with a Shimadzu XRD 6000 Diffractometer equipped with Cu K␣ radiation. The diffractometer was operated at very low scanning rates to increase the counting frequency. A Voyager 2110 model energy dispersive X-ray analyzer (EDS) was also used for further identification of the intermetallic phases. Soaked samples were investigated for the extent of ␤-AlFeSi → ␣-AlFeSi transformation, solutionizing of the coarse Mg2 Si particles and for the homogenization of the solid solution matrix. The effect of cooling practice on the precipitation of Mg2 Si particles was investigated by employing two different types of cooling cycles following a 6 h soak at 580 ◦ C (Fig. 1b). The first of these involved continuous cooling at rates ranging from air cooling to 12 ◦ C/h. For the size of the specimens used in the present study, air cooling corresponded to a cooling rate of 2000 ◦ C/h. Temperatures were recorded directly from the samples as a function of time and measures were taken to establish a constant cooling rate down to 200 ◦ C at which point all samples were air-quenched. The second set of cooling experiments involved a 2 h isothermal step at various temperatures between 450 and 200 ◦ C. Samples soaked at 580 ◦ C for 6 h were air-quenched before and after the 2 h isothermal step. These samples were investigated for the extent of Mg2 Si precipitation and for the size of the Mg2 Si particles. They were examined after etching with a 0.5% HF solution using an Olympus BX51M model optical microscope. Samples were slightly overetched to facilitate direct observation of the fine Mg2 Si particles with the light microscope. Dark-field imaging was also employed to improve resolution. The electrical conductivity of the samples was measured with a Sigma Test Unit for the evaluation of the extent of the precipitation activities. Differential scanning calorimetry (DSC) was employed to investigate how the homogenized samples cooled at different rates and with an isothermal step at different temperatures would respond to a reheating treatment. A Polymer Laboratories DSC unit was used and the samples were heated under flowing argon at a rate of 20 ◦ C/min until before melting.

3. Results and discussion 3.1. As-cast billet

Fig. 2. (a) Microstructure of the as-cast billet, (b) the shell zone at the surface and (c) macrostructure showing the grains and the segregation pattern.

The microstructure of the as-cast billet consisted of primary dendrites of aluminum rich solid solution with a well-defined interdendritic network of ␤-AlFeSi platelets (Fig. 2a). There was hardly any evidence for the Mg2 Si phase, owing to a high enough cooling rate during solidification which apparently retained the Si remaining from the insoluble constituents and Mg in solid solution. The shell zone in the cast billet was less than 100 ␮m deep and was

generally free of coarse intermetallics (Fig. 2b), while the average grain size was estimated to be 150 ␮m (Fig. 2c). Both the cell size and the grain size were uniform across the cross-section of the billet. It was thus concluded, from the metallographic analysis that the as-cast billet was of adequate quality.


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Fig. 3. The microstructural features of samples soaked between 540 and 580 ◦ C for 2–6 h and subsequently air-quenched.

3.2. Soaking experiments Fig. 3 illustrates the microstructures of those samples that were soaked between 540 and 580 ◦ C for 2–6 h and subsequently air-cooled. Soaking at 540 ◦ C for 2 h seemed to have produced little change in the type and distribution of AlFeSi phases. The majority of the interdendritic particles were still plate-like and were identified by XRD to be of the ␤-AlFeSi variety suggesting that the ␤ → ␣ transformation at this temperature was very sluggish. ␤-AlFeSi was still the major phase in those samples soaked at 540 ◦ C for 4 and 6 h in spite of an increasing number of particles that were identified by EDS to be of the hexagonal ␣-AlFeSi type (Table 2). Both structural and morphological features of the cast billet have changed considerably starting at 560 ◦ C. The ␤

platelets broke up into relatively small discrete particles after soaking at this temperature. The latter were responsible for the new reflections in the XRD spectrum that were readily indexed by the hexagonal ␣-AlFeSi phase. The ␤-AlFeSi particles that have survived soaking implied that the ␤ → ␣ transformation was still not complete at 560 ◦ C. Further evidence for the hexagonal ␣ phase becoming the major phase was provided by the EDS work that gave a Fe/Si ratio between 1.5 and 2 for the majority of the particles analyzed. The change in the microstructural features outlined above has become more prominent after soaking at 580 ◦ C. The ␤-AlFeSi phase was clearly the minor phase only after 2 h at this temperature (Table 2). The interdendritic platelets were replaced almost entirely by discrete round particles, giving a “necklace” type configuration. The ␤-AlFeSi phase has

Table 2 The insoluble constituents in soaked billets Temperature (◦ C)

540 560 580

Time (h) 2

4

6

␤-AlFeSi, ␣-AlFeSi ␣-AlFeSi (␤-AlFeSi) ␣-AlFeSi (␤-AlFeSi)

␤-AlFeSi, ␣-AlFeSi ␣-AlFeSi (␤-AlFeSi) ␣-AlFeSi (␤-AlFeSi)

␤-AlFeSi, ␣-AlFeSi ␣-AlFeSi (␤-AlFeSi) ␣-AlFeSi

The major constituent is underlined; the minor ones are put into parentheses. The constituents are listed in the order of their volume fractions.

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Fig. 4. The distribution of Mg2 Si precipitates in a sample soaked at: (a) 540 ◦ C for 2 h and (b) 560 ◦ C for 4 h subsequently cooled to 200 ◦ C at 200 ◦ C/h.

completely transformed after 6 h soaking at 580 ◦ C while nearly equal number of hexagonal and cubic ␣-AlFeSi particles were encountered in EDS analysis. The microstructures of the two of the samples which were given a 6 h soak at 580 ◦ C and subsequently cooled at a rate of 200 ◦ C/h to allow for Mg2 Si precipitation are shown in Fig. 4. The distribution of Mg2 Si particles in samples soaked at 540 ◦ C revealed a heterogeneous character inherited from the as-cast microstructure suggesting that the microsegregation was yet not removed. The sample soaked at 560 ◦ C for 4 h, on the other hand, showed a more uniform distribution of Mg2 Si particles. It is thus claimed that a soaking treatment of at least 4 h at 560 ◦ C is required to eliminate segregation and to even out the Mg and Si distribution in the matrix.

3.3. Cooling experiments 3.3.1. Continuous cooling The post-homogenization cooling rate appeared to have a profound effect on the precipitation and dispersion characteristics of the Mg2 Si phase, as expected. The samples soaked at 580 ◦ C for 6 h and subsequently air-cooled (at a rate of 2000 ◦ C/h) hardly revealed any Mg2 Si precipitates (Fig. 5). It is thus concluded that the Mg2 Si particles were fully solutionized after soaking at 580 ◦ C and that the free Si and the Mg were both retained in solution during air-cooling. The population of very fine Mg2 Si particles inside the dendrites increased with decreasing cooling rates and finally, the samples cooled at 100 ◦ C/h were crowded with such particles. The precipitation of Mg2 Si particles was apparently not random and those sites with relatively higher energy were decorated first upon precipitation. The increasing precipitation activity with decreasing cooling rates was reflected also by increasing conductivity values (Table 3). Mg2 Si precipitates were either too few and/or too fine to give a detectable signal in the XRD spectrum, in all samples cooled at 100 ◦ C/h or faster. Reflections of the cubic equilibrium Mg2 Si phase were noted, however, after cooling at 30 ◦ C. This change in the XRD response was accompa-

nied by a remarkable change in the microstructural features of these samples. The population of the Mg2 Si precipitates was reduced drastically as the very fine and dense dispersion was replaced by relatively coarse, stable precipitates when the cooling rate dropped below 100 ◦ C/h (Fig. 5). The increasing conductivity values in this range suggested that the depletion of the solid solution matrix was still underway (Table 3). The coarsening of the Mg2 Si precipitates continued further with decreasing cooling rate as evidenced by the increase in the population of coarse Mg2 Si particles at the expense of finer ones in the sample cooled very slowly at 12 ◦ C/h. Fig. 6 shows the DSC scans of the samples soaked at 580 ◦ C for 6 h and cooled very fast to very slowly, in the range 2000–12 ◦ C/h. Results of the optical microscopy, electrical conductivity measurements and the XRD work were paralleled by the DSC analysis of these samples. The thermograms of the samples cooled at 400 ◦ C/h and faster revealed a total of four well-defined peaks until 600 ◦ C (Fig. 6). The first three of these peaks are exothermic and are thus associated with precipitation activities while the last one is endothermic and indicates a dissolution reaction. The first exothermic peak is rather small and occurs at approximately 140 ◦ C. It is claimed to be associated with the formation of Mg–Si clusters. The precipitation of the metastable ␤ /␤ -Mg2 Si and the stable ␤-Mg2 Si particles were responsible for the second and third exothermic peaks, respectively. The endothermic peak which is estimated to start at 450 ◦ C, Table 3 Electrical conductivity values of the samples soaked at 580 ◦ C for 6 h and subsequently cooled at different rates Cooling rate (◦ C/h)

Electrical conductivity (m/ mm2 )

2000 800 400 200 100 30 12

28.6 28.7 29.0 29.6 30.1 31.0 32.0


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Fig. 5. Dark-field optical micrographs showing the distribution of Mg2 Si precipitates in samples soaked at 580 ◦ C for 6 h and subsequently cooled to 200 ◦ C at: (a) 2000 ◦ C/h, (b) 800 ◦ C/h, (c) 400 ◦ C/h, (d) 200 ◦ C/h, (e) 100 ◦ C/h, (f) 30 ◦ C/h and (g) 12 ◦ C/h.

on the other hand, represents the solutionizing of the equilibrium ␤-Mg2 Si phase, i.e. the solvus temperature. The intensity of the exothermic peaks correlates very well with the homogenization cooling rate. The higher precipitation capacity of those samples cooled at 2000 and 800 ◦ C/h, is demonstrated by their relatively larger exothermic peaks. The intensity of these peaks is reduced with decreasing cooling rate suggesting that increasing amounts of Mg2 Si precipitation has already taken place during post-homogenization cooling. Finally, the exothermic peaks were very weak at 30 ◦ C/h and were completely missing in the case of very

slow cooling at 12 ◦ C/h. There was hardly any Mg and Si left in solution to precipitate during reheating when the billet was cooled from the soaking temperature very slowly. The broad exothermic effect observed in the samples cooled slowly at 30 and 12 ◦ C/h are claimed to be due to further coarsening of the stable Mg2 Si precipitates. The temperature range of the dissolution peak is of great technological interest as it signals the potential problems that are often encountered with respect to surface quality and hardening capacity. The solutionizing of the Mg2 Si precipitates was over at approximately 500 ◦ C, well below the

256


Fig. 6. DSC scans of the samples soaked at 580 ◦ C for 6 h and cooled at various rates, in the range 12–2000 ◦ C/h. The heating rate was 20 ◦ C/min.

press exit temperatures commonly encountered in industrial practice in all samples cooled from the soaking temperature faster than 100 ◦ C/h. A small additional endothermic effect appeared during reheating of the sample cooled at 100 ◦ C/h, suggesting that coarse cubic ␤-Mg2 Si precipitates, in addition to the very fine ones, have formed at this cooling rate. The endothermic peaks were displaced to still higher temperatures (by as much as 50 ◦ C) in the samples cooled at 30 and 12 ◦ C/h. While the fine, metastable Mg2 Si particles which form at high cooling rates respond to preheating readily, coarse ␤-Mg2 Si particles are apparently stable at high temperatures and are thus more difficult to dissolve during reheating as evidenced by the displacement of the dissolution peak to higher temperatures in the DSC scans of the samples cooled at or slower than 30 ◦ C/h. It is thus concluded that post-homogenization cooling faster than 100 ◦ C/h should be employed for the present alloy in order to avoid the formation of coarse Mg2 Si precipitates which would be difficult to solutionize during

Fig. 7. Dark-field optical micrographs showing the distribution of Mg2 Si precipitates in samples soaked at 580 ◦ C for 6 h and step-cooled at: (a) 450 ◦ C, (b) 400 ◦ C, (c) 350 ◦ C, (d) 300 ◦ C, (e) 250 ◦ C and (f) 200 ◦ C.


reheating and extrusion. A cooling rate below 200 ◦ C/h is desired, on the other hand, to achieve an effective depletion of the solid solution matrix for a low flow stress and low resistance to extrusion deformation. 3.3.2. Step-cooling The temperature of the isothermal step inserted in the cooling cycle influenced Mg2 Si precipitation just as much as the cooling rate did in continuous cooling cycles. Decreasing the isothermal-step temperature seemed to produce an effect similar to that of decreasing the cooling rate. The samples soaked at 580 ◦ C for 6 h and air-quenched before and after a 2 h isothermal hold at 450 ◦ C did not reveal any Mg2 Si precipitates (Fig. 7). This was confirmed also by the electrical conductivity of this sample which was nearly equal to that of the soaked and quenched (supersaturated) sample. Hence, the solvus temperature is claimed to be slightly below 450 ◦ C for the present alloy and Mg2 Si, if any, were solutionized rather than precipitated after 2 h at this temperature. Some precipitation activity was first noted after step-cooling at 400 ◦ C with an accompanying increase in electrical conductivity. The precipitation at this high temperature has produced a very coarse dispersion of the stable cubic ␤-Mg2 Si particles. The population of Mg2 Si particles increased but their size decreased with decreasing isothermal-step temperatures. Finally, the sample step-cooled at 250 ◦ C revealed a very fine dispersion of Mg2 Si particles. The majority of these particles were of the metastable ␤ -Mg2 Si variety as evidenced by the lack of their exothermic effect in the calorimetry scan obtained during reheating. The sample step-cooled at 200 ◦ C, on the other hand, were free from any Mg2 Si particles indicating that no precipitation has taken place at this temperature. Electrical conductivity measurements were in good agreement with the microstructural observations (Table 4). The increasing overall precipitation activity with decreasing

257

Table 4 Electrical conductivity values of the samples soaked at 580 ◦ C for 6 h and subsequently cooled at different isothermal hold temperatures Isothermal-step temperature (◦ C)

Electrical conductivity (m/ mm2 )

450 400 350 300 250 200

28.8 30.1 31.3 31.6 31.6 29.2

isothermal-step temperatures was reflected by increasing conductivity values until 300 ◦ C. The extent of the precipitation was more or less the same in the sample step-cooled at 250 ◦ C as inferred from a very similar electrical conductivity value. Lack of precipitation at 200 ◦ C is evidenced also by the sudden drop in the electrical conductivity of this sample. Fig. 8 shows the DSC scans of the samples soaked at 580 ◦ C for 6 h and step-cooled at various temperatures between 200 and 450 ◦ C. Results of optical microscopy, electrical conductivity measurements and XRD work were paralleled by the DSC analysis of the step-cooled samples. The reheating thermogram of the sample step-cooled at 200 ◦ C reveal all three exothermic peaks until 600 ◦ C in the addition to the endothermic peak as this temperature was too low to allow any precipitation within 2 h. The two low-temperature exothermic peaks were completely missing and the exothermic effect produced by the precipitation of the equilibrium ␤-Mg2 Si phase was very much weakened in that sample cooled with a 2 h isothermal hold at 250 ◦ C. It is thus concluded that much of the available Mg and Si had already precipitated in the form of ␤ /␤ -Mg2 Si particles during the 2 h isothermal hold at this temperature and little Mg and Si were available to precipitate during reheating.

Fig. 8. DSC scans of the samples soaked at 580 ◦ C for 6 h and step-cooled at temperatures between 450 ◦ C and 200 ◦ C. The heating rate was 20 ◦ C/min.

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The thermogram of the sample step-cooled at 300 ◦ C was almost identical, as one would expect in view of the metallographic work and electrical conductivity measurements. A broad exothermic effect was noted during reheating of the sample step-cooled at 350 ◦ C, as typically observed when coarsening processes are active. Exothermic peaks for the precipitation of both metastable ␤ /␤ -Mg2 Si and stable ␤-Mg2 Si particles appeared during reheating of the sample step-cooled at 400 ◦ C. The higher solute levels of Mg and Si in the sample held at this temperature is responsible for the restoration of the precipitation capacity. A major difference produced by the isothermal hold at 400 ◦ C is the shift of the dissolution peak to higher temperatures suggesting that the Mg2 Si particles formed during the isothermal hold and coarsened further during reheating are very difficult to solutionize and are potentially harmful both to the surface quality and mechanical properties. Of the several isothermal-step temperatures tested in step-cooling experiments, 250 and 300 ◦ C appeared to yield the best results as they have maximized the precipitation of Mg2 Si without shifting the dissolution peak to higher temperatures.

4. Conclusions Soaking at 540 ◦ C falls far too short of expectations as it fails to remove microsegregation and to achieve the ␤-AlFeSi → ␣-AlFeSi transformation both of which are desirable to improve the surface quality and extrudability of the material. The interdendritic network of the ␤-AlFeSi platelets is replaced by a dispersion of equiaxed, discrete ␣-AlFeSi particles and a homogeneous Al(Mg,Si) solid solution is obtained starting at 560 ◦ C. It takes a 6 h soak at 580 ◦ C to fully transform the ␤-AlFeSi phase to the ␣-AlFeSi phase in the present alloy with a very low Mn level. The extent of precipitation increased with decreasing post-homogenization cooling rates. However, cooling rates below 100 ◦ C/h produced coarse ␤-Mg2 Si particles that were shown by DSC to be difficult to solutionize during reheating. A cooling rate below 200 ◦ C/h is desired, on the other hand, to achieve an effective depletion of the solid solution matrix for a low flow stress and low resistance to extrusion deformation. It is thus concluded that a cooling rate between 200 and 100 ◦ C/h is optimum for the present alloy for high extrudability and surface quality should continuous cooling be employed after soaking. Step-cooling at 250–300 ◦ C for 2 h was found to give extensive precipitation of the fine, metastable ␤ -Mg2 Si

particles which were readily solutionized during reheating. Step-cooling under the conditions described above appears to be a better practice than continuous cooling at a rate between 100 and 200 ◦ C/h for optimum billet performance. The former gives a more complete depletion of the aluminum solid solution, i.e. lower flow stresses, avoiding the formation of coarse and stable Mg2 Si particles which are very difficult to solutionize during reheating and thus survive the extrusion process impairing both surface quality and mechanical properties.

Acknowledgements It is a pleasure to thank Miss Selda Üçüncüoˇglu and Mr. Fahri Alageyik for their help in the metallographic work, Mr. Osman Çakir for performing the heat treatments, Mrs. Feriha Birol of ARÇELI˙ K for running the DSC tests.

References [1] E.C. Beatty, in: Proceedings of the First International Aluminum Extrusion Technology Seminar, Paper no. 10, Aluminum Association, Washington, DC, 1969. [2] A.J. Bryant, G.E. Macey, R.A.P. Fielding, Light Met. Age (2002) 6. [3] Y. Birol, G. Kara, S. Üçüncüoˇglu, M. Usta, in: Z. Jin, A. Beaudoin, T.A. Bieler, B. Radhakrishnan (Eds.), Hot Deformation of Aluminum Alloys III, TMS, 2003, p. 463. [4] S.R. Claves, D.L. Elias, W.Z. Misiolek, Materials Science Forum, vols. 396–402, Trans Tech Publications, Switzerland, 2002, p. 667. [5] S. Zajac, et al., Mater. Sci. Technol. 10 (1994) 323. [6] E.C. Beatty, in: Proceedings of the Second International Aluminum Extrusion Technology Seminar, vol. 1, pp. 225–228. [7] J. Langerweger, in: Proceedings of the Conference on Aluminium Technology, London, March 1986, p. 216. [8] K. Karhausen, A.L. Dons, T. Aukrust, Mater. Sci. Forum 217–222 (1996) 403. [9] T. Oka, M. Yanagawa, M. Hirano, in: Proceedings of the Sixth International Aluminum Extrusion Technology Seminar, vol. 1, Aluminum Association, Washington, DC, 1996, p. 157. [10] N.C.W. Kuijpers, W.H. Kool, S. van der Zwaag, Materials Science Forum, vols. 396–402, Trans Tech Publications, Switzerland, 2002, p. 675. [11] H. Tanihata, et al., J. Mater. Sci. 34 (1999) 1205–1210. [12] S. Zajac, B. Bengtsson, C. Jonsson, Materials Science Forum, vols. 396–402, Trans Tech Publications, Switzerland, 2002, p. 399. [13] E.B. Bjornbakk, et al., Materials Science Forum, vols. 396–402, Trans Tech Publications, Switzerland, 2002, p. 405. [14] O. Reiso, et al., in: Proceedings of the Sixth International Aluminum Extrusion Technology Seminar, vol. 1, Aluminum Association, Washington, DC, 1996, p. 1. [15] N. Dahl, et al., in: Proceedings of the Sixth International Aluminum Extrusion Technology Seminar, vol. 1, Aluminum Association, Washington, DC, 1996, p. 529.