Control of Equilibrium Phases (M,T,S) in the Modified Aluminum Alloy 7175 for Thick Forging Applications

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1 Materials Transactions, Vol. 44, No. 1 (2003) pp. 181 to 187 #2003 The Japan Institute of Metals EXPRESS REGULAR ARTICLE Control of Equilibrium Phases (M,T,S) in the Modified Aluminum Alloy 7175 for Thick Forging Applications Seong Taek Lim 1; *, Il Sang Eun 2 and Soo Woo Nam 1 1 Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1Guseong-dong, Yuseong-gu, Daejeon, , Korea 2 Agency for Defence Development, P.O. Box 35-5, Yuseong-gu, Daejeon, , Korea Microstructural evolutions, especially for the coarse equilibrium phases, M-, T- and S-phase, are investigated in the modified aluminum alloy 7175 during the primary processing of large ingot for thick forging applications. These phases are evolved depending on the constitutional effect, primarily the change of Zn:Mg ratio, and cooling rate following solutionizing. The formation of the S-phase (Al 2 CuMg) is effectively inhibited by higher Zn:Mg ratio rather than higher solutionizing temperature. The formation of M-phase (MgZn 2 ) and T-phase (Al 2 Mg 3 Zn 3 )is closely related with both constitution of alloying elements and cooling rate. Slow cooling after homogenization promotes the coarse precipitation of the M- and T-phases, but becomes less effective as the Zn:Mg ratio increases. In any case, the alloy with higher Zn:Mg ratio is basically free of both T and S-phases. The stability of these phases is discussed in terms of ternary and quaternary phase diagrams. In addition, the modified alloy, Al 6Zn 2Mg 1.3%Cu, has greatly reduced quench sensitivity through homogeneous precipitation, which is uniquely applicable in 7175 thick forgings. (Received October 7, 2002; Accepted November 12, 2002) Keywords: aluminum alloy 7175, large ingot, Zn:Mg ratio, S-phase, T-phase, homogenization, thick forging, quench sensitivity 1. Introduction Aluminum alloy AA7175, one of the newer variants of the baseline alloy AA7075, has primarily been designed for the forging requiring higher strength and damage tolerance. 1 3) Because of the same alloy composition with AA7075 (5.6Zn 2.5Mg 1.6Cu 0.23Cr) except the lowered impurities, AA7175 is well characterized in the casting, hot working and heat treatment areas in the case of small section thickness. As a thick forging, however, AA7175 with conventional alloy composition is not a good choice due to the massive evolution of coarse equilibrium phases and high quench sensitivity, and therefore usually replaced by 7X50 type alloys. 1 6) In this respect, the alloy should entail the special controls on the alloy composition and processing for microstructural homogeneity. There has been a strong tendency for 7xxx alloys to be used as large and thick semiproducts. 7 10) These thick materials, however, often have poor mechanical and fracture properties since they receive less deformation during hot working and slower quench cooling. 8 10) This is especially due to that thick (open die) forgings are usually made of large diameter ingots, reaching several hundred milimeters, which accompany large dendrite size and concurrent heavy segregations owing to slow casting speed. 4) Through the many slow processes, there are frequent evolutions of the coarse equilibrium phases and constituent particles which have an adverse effect on strength, damage tolerance and hot workability. 8 10,17,18) From the earlier studies by Strawbridge et al. 11) to the recent studies by Li and Starink, 12,13) complex phases are evolved in the commecial quaternary (Al Zn Mg Cu) *Graduate Student, Korea Advanced Institute of Science and Technology. Present address: Agency for Defence Development, P.O. Box 35-5, Yuseong-gu, Daejeon, , Korea. system owing to the extended solubility of copper into the precipitates. Besides the minor phases ( =Al 2 Cu, Z = Mg 2 Zn 11 ), the resulting microstructure is generally characterized by equilibrium phases of soluble M (MgZn 2 = ), T (Al 2 Mg 3 Zn 3 ) and their isomorphous phases, S (Al 2 CuMg), and insoluble Al 7 Cu 2 Fe particles. 4,5,11 17) These phases constitute most of the microstructure and govern the general properties. M-phase and its precursors GP zone and intermediate precipitate ( 0 ) are responsible for high strength. Low melting point T- and S-phases affect hot workability and fracture properties. 4,8,10,12,17,18) S-phase has been a main target to be prevented in the thick 7X49 and 70X0 heavy plates due to its massive appearance and brittleness. 8,10,12,18) The decompositions of these phases are strongly dependent on the constitution of alloying elements such as (Zn+Mg+Cu) and Zn:Mg ratio, and on the cooling rate following solutionizing. 3,4,8,10 13,17) Furthemore, the amount of equilibrium phases (M,T,S) is essentially determined at the primary processing stages such as casting and homogenization, eapecially in the less deformed thick materials. 4,8 10) From the backgound mentioned above, the modification of alloy composition for the specific purpose of thick 7175 forging is attempted to control the equilibrium phases. Recent studies on this area are concentrated on the 70X0 (X ¼ 1; 4; 5) type plate alloys 8 10,18) and mostly regarded as proprietaries. 9,10) The constitutional study on the 7X75 alloy forging is not reported yet. In this respect, the effects of alloy redesign, especially Zn:Mg ratio, and the cooling rate on the microstructural evolutions are investigated using real scale big ingots and thick forgings of 7175 alloys. 2. Experimental Procedure The materials used in this study are large forging ingots of 7175 type alloys with 700 mm in diameter prepared by DC semicontinuous casting at Dooray Air Metal in Korea (now

2 182 S. T. Lim, I. S. Eun and S. W. Nam Table 1 Chemical compositions of 7175 type alloys used in this study. Alloy Si+Fe Cu Mg Zn Cu+Mg Zn+Mg+Cu Zn:Mg AA max Alloy A Alloy B Alloy C Alloy D Alloy E Alcoa Korea) as a 7xxx big ingot feasibility study program. Casting speed and other variables are carefully controlled to avoid ingot cracking and ensure internal structure. Five kinds of alloys given in Table 1 are designed to have different Zn:Mg ratio ranging from 1.9 (alloy A) to 3.0 (alloy E) within or around the specification limit of commercial AA Contents of chromium are kept minimum (0.18%), and Fe+Si as low as 0.15 mass% except alloy A. The richer alloy group (B,C) and the more dilute alloy group (D,E) have respectively the similar total alloying contents. The dilute alloy E has relatively lower Cu and higher Zn:Mg ratio. The 25 mm square samples are taken at the mid-radius location of ingots and are homogenized at 733 or 743 K for a given time. On the other hand, the 700 mm ingots subject to open die forging into the intermediate forgings (380 mm diameter, " 1:2) at between 673 and 623 K using 2000 ton hydraulic press. The 25 mm square samples are taken at the mid-radius location of the forgings. Homogenizing treatments and subsequent coolings are also performed in the DSC module where precise coolings at the rate of Kmin 1 from K to 523 K are obtained. Microstructural characterizations are performed by using optical microscopy (OPTIPHOT 100S, NIKON), scanning electron microscopy (XL30FEG, PHILIPS) with EDS, and transmission electron microscopy (JEOL 4000FX). Thin foils for TEM are prepared by mechanical polishing to 150 mm and final twin-jet electropolishing in a solution of 25%HNO 3 +75%CH 3 OH at 20 K and 10 V. Calorimetric studies are performed in the 2910 DSC module in TA Instrument with pure aluminum reference and the sample discs (about 80 mg) at the rate of 10 Kmin 1 from 303 to 803 K. The results are average of three runs. 3. Results and Discussion 3.1 General Microstructures Figure 1 shows the microstructures of 700 mm diameter ingot (alloy B) in as-cast and as-homogenized conditions at mid-radius, indicating the change of the microstructure with homogenizing procedures. In as cast state (Figs. 1a, b), due to Fig. 1 General micrographs of the large ingots in the alloy B at mid-radius: a) as-cast, b) interdendritic region in a), c) partly homogenized at 733 K for 1 h, and d) fully homogenized at 733 K for 24 h and 743 K for 16 h.

3 Control of Equilibrium Phases (M,T,S) in the Modified Aluminum Alloy 7175 for Thick Forging Applications 183 slow casting speed, cell size reaches about 95 mm and boundary thickness 25 mm, which is much larger than those of small ingots typically 30 mm and 10 mm, respectively. The large cell is difficult to be homogenized due to the large interdendritic segregations. The main constituents of interdendritic network are identified by EDS mostly as Al Cu lamella eutectic with very limited amount of magnesium and zinc, and Al Cu Fe compounds. The former is generally designated by eutectic [M] or T [Al Cu Mg Zn]. 4) These are severely cored by copper-rich outer rims. In the slightly homogenized state (Fig. 1c), the Cu-rich eutectics are partially dissolved and become discrete particles and some T (Al 2 Mg 3 Zn 3 ) phase are seen at the cell boundary as well as within the grain. In fully homogenized state (Fig. 1d), the Curich eutectics are converted into S (Al 2 CuMg), and Al Cu Fe compounds into Al 7 Cu 2 Fe particles. Impurity-originated Al 7 Cu 2 Fe particles are very coarse and bulky, varying 5 20 mm. S-phase is known to form rapidly during the dissolution of [M] eutectic into the solid solution. 8,18) Particles of S-phase are generally round and clustered, varying 2 10 mm. The matrix is characterized by coarse platelets of M (MgZn 2 ) and occational T (Al 2 Mg 3 Zn 3 ) with polygonal shape, which are high temperatute precipitates decomposed during cooling after homogenization. The resulting microstructure is generally characterized by the equilibrium phases of soluble M-, T- and S-phase, and insoluble impurity-originated Al 7 Cu 2 Fe particles. Platelets of M-phase (often referred ) and polygonal T-phase are evolved depending on alloy composition, especially on the Zn:Mg ratio, and the cooling rate from solutionizing temperature. The evolution of the S-phase depends also on the alloy composition, and solutionizing time and temperature. T- and S-phases have low melting points so that they limit the heating range in hot working and solutionizing. 4,13,17) The brittleness of the S-phase greatly impairs the fracture properties. 8,10) The large ingots are good examples for the evolution of these phases because they subject to severe segregation and very slow cooling after solutionizing. 3.2 Evolution of S-phase In the homogenized state, alloy A and B have large amount of intermetallic phases, Fe-containing Al 7 Cu 2 Fe and Mgcontaining Al 2 CuMg (S-phase). This is in contrast to alloy E in which there are a little Al 7 Cu 2 Fe only, and is basically free of S-phase. Large amount of Al 7 Cu 2 Fe in alloy A is due to the relatively high iron content. Considering the alloy composition, it can be seen that as Cu+Mg content decreases or Zn:Mg ratio increases, the quantity of S-phase decreases. The microstructural evidences of the S-phase varying with the alloy designs are consistent with the results of DSC tests shown in Fig. 2. It can be noted that alloy A and B show large and sharp endothermic peaks, corresponding to the melting of S-phase which onset at around 763 K. In the alloy D and E, however, S-phase melting peaks are diminished or completely disappeared. The evolution of the S-phase is similarly characterized in the intermediate forgings as shown in Fig. 3. All alloys exhibit similar grain structures with less elongated grains about 300 mm long, and with well developed subgrains of 3 to 8 mm diameter. Partial recrystallizations are kept less than Heat Flow, W / g T-phase Alloy E D B A E (Zn/ Mg=3.0) S-phase D (Zn/ Mg=2.8) B (Zn/ Mg=2.1) Homo. Cooling Rate A (Zn/ Mg=1.9) : 0.67K min Temperature, T / K Fig. 2 DSC thermograms at high temperature interval of the alloys in the homogenized and cooled at the rate of 0.67 Kmin 1 following homogenization at 733 K for 12 h. 15% in areal fraction. The major constituent particles are also Al 7 Cu 2 Fe and S-phase, aligned along the working axis. The large quantity of S-phase decorating the grain boundary are seen in the alloy A and B, which is in contrast to the clean alloy E. The evolution of S-phase through the constitutional variations is reviewed from the two standpoints, Cu+Mg and Zn:Mg ratio. In the classical work by Hume-Rothery, 19) curve separating the and (+S) fields in Al Cu Mg is given by: Log[at%Cu][at%Mg] ¼ 5: T 1, where T is temperature in Kelvin. The solubility curves at the commercial solutionizing temperatures from 723 to 773 K are solved and plotted in Fig. 4. Strawbridge et al. 11) also presents the phase stability of Al Mg Cu 6 mass%zn system at 733 K, which is added in Fig. 4 for comparison. Alloy D and E lie on a little outside the limit of the registered AA7175 specification. The solubility of (Cu+Mg) at 733 K tends to extend in the quaternary system (Strawbridge s). In any case, high Cu+Mg contents make the alloy lie in the (+S) field, e.g. increases the S-phase solvus. It can be seen that at 733 K alloy A, B, and C are within the field of (+S), whereas alloy D and E are within -Al. The soluble S-phase is expected to be completely dissolved into the matrix at higher temperature over 763 K even in the alloy A. However, this is impractical because of very slow dissolution rate of S- phase and possible overheating at the highly segregated areas. 4,5,9) The upper limits of the solutionizing temperature in the 7xxx alloys are mostly determined by the S-phase solvus. 4,5) In this regards, Alloy E has a S-phase solvus about 723 K which is about 40 K lower than that of alloy A. Hyatt 3) explained the phase stability of S-phase in 90 mass%al Zn Mg Cu system at 733 K in terms of Zn:Mg ratio. Figure 5 shows that as Zn:Mg ratio (iso-ratio line) varies from 2.0 to 4.0 for the constant copper level about 2.0 mass%, the alloys move from a two-phase (+S) to a single phase field. In addition, solubility of copper increases in high Zn:Mg ratio alloy which, as a design concept of AA7050, has more Al 2 CuMg to be dissolved at the solutionizing temperature. 3 5) Although alloy A (90 mass%al) only is exactly defined in the diagram, the appearance of massive S-phase for the lower Zn:Mg ratio alloy is well

4 184 S. T. Lim, I. S. Eun and S. W. Nam Fig. 3 General micrographs of the T7 treated intermediate forgings at mid-radius section: a) and b) alloy A, c) alloy B, d) alloy E K (Al) + S Cu, mass % (Al) D C B A E Hume-Rothery at K Strawbridge at 733K 7175 Limit Mg, mass % Fig. 4 S-phase stability of the alloys in ternary Al Cu Mg (Hume- Rothery) 19) and quaternary Al 6%Zn Cu Mg (Strawbridge) 11) systems. explained. Reflecting that alloy group (B,C) and (D,E) have essentially the same respective (Cu+Mg) contents (Table 1), it is reasonable to say that the formation of the S-phase depends first on the Zn:Mg ratio, and second on the (Cu+Mg) content around the limited range of AA7175 specification. 3.3 Evolutions of M- and T-phases Other equilibrium phases of concern are M- and T-phases. Equilibrium phases of M and T reprecipitate in a coarse Fig. 5 Phase diagram of quaternary 90 mass%al Zn Cu Mg system at 733 K, showing the effect of changing Zn:Mg ratio. 3) morphology during cooling from homogenizing temperature. Figure 6 shows the micrographs of alloy B and E which are fast cooled or slowly cooled during critical temperature interval between 733 and 523 K at the rate of 3.3 and 0.1 Kmin 1, respectively. In the fast cool (Figs. 6a, c), both alloys reveal the precipitation of platelets of M-phase only. However, in the slow cooled alloy B with lower Zn:Mg ratio, coarse precipitates of T-phase about 5 mm are seen at both cell boundary and grain interior, along with coarse platelets of M-phase about 10 mm. White particles of T-phase are seen

5 Control of Equilibrium Phases (M,T,S) in the Modified Aluminum Alloy 7175 for Thick Forging Applications 185 Fig. 6 Micrographs of homogenized alloy B (a,b) and E (c,d) with different cooling rate after homogenization at 733 K for 12 h: a) and c) cooling rate 3.3 Kmin 1, b) and d) 0.1 Kmin 1. only by the backscattered electron image in the SEM. There are only fine and uniform M-phases in the higher Zn:Mg ratio alloy E. Generally, given a same cooling rate, M-phase precipitates in a fine and uniform manner in the alloy E, but coarse and heterogeneous manner in the alloy B. The greater density of M-phase in the alloy E will be discussed below in terms of the Zn:Mg ratio and the quench sensitivity. Figure 7 shows the DSC thermograms for the higher temperature region showing the effect of cooling rate following homogenization. Single endothermic peak at around 753 K in the as-cast alloy B is related with the reaction involving the M-, T-, and S-phases, usually encountered in the 7xxx quaternary alloys. 15,17) In contrast to the fast cool (Fig. 2), separated double endothermic peak T and S are clearly evident in the slowly cooled from homogenization temperature. The peak T at about 752 K (onsets at 748 K) is responsible for the melting of T-phase reformed during slow cooling at high temperature, and peak S at about 771 K (onsets at 763 K) for the melting of S-phase already formed during homogenization. 12,13,17) Peak of T- phase melting is not observed in the high Zn:Mg alloy E. Figure 8 shows the cooling rate dependence on the dissolution energies associated with the melting of T- and 2.4 Heat Flow, W/g Alloy E Alloy D Alloy B T S As-cast Alloy B Dissolution Energy, J/g Peak S Peak T C.W.Q. C.W.Q Homo. Cooling Rate : 0.1K min Temperature, T / K Cooling Rate, K min Fig. 7 Evolutions of T and S-phase in the very slow cooling (0.1 Kmin 1 ) following homogenization at 733 K. Fig. 8 Homogenization cooling rate effect of the alloy B on the evolution of T and S-phases from DSC tests (cooled after holding at 733 K for 12 h).

6 186 S. T. Lim, I. S. Eun and S. W. Nam Dissolution Energy, J/g A B S-phase, 0.1K min -1 from 733K T -phase, 0.1K min -1 from 733K S-phase, 1.67K min -1 from 743K T -phase, 1.67K min -1 from 743K S-phases in the low Zn:Mg alloy B. T-phase is not appeared when the critical cooling rate about 0.7 Kmin 1 is exceeded. The average cooling rate during critical range is measured as about 0.7 Kmin 1 at the one third of the radius in the 700 mm diameter ingot. This suggests the possible occurrence of the T-phase within the large ingots. S-phase is moderately decreased with increasing homogenization cooling rate. However, it is observed that the fast cooling alone cannot exclude S-phase effectively. Figure 9 is a master diagram showing the effects of the constitution, and temperature and cooling rate of the homogenization on the evolution of T and S-phases. It is clearly seen that slow cooing after homogenization promotes the precipitation of the T-phase, but becomes ineffective as Zn:Mg ratio increases. Fast cooling in any alloy definitely retards the T-phase precipitation, but partly effective for the S-phase. Higher homogenization temperature reduces the amount of the S-phase somewhat, but higher Zn:Mg ratio is found to be more effective inhibitor. In any case, alloy E having the highest Zn:Mg ratio is basically free of both T- and S-phases. In the most of commercial high strength 7xxx alloys, the quaternary compositions are not single phase even at the solutionizing temperatures. 4) Through the participation of copper into the precipitates, M-phase ranges from MgZn 2 to AlCuMg described as isomorphous Mg(Al,Cu,Zn) 2, and T- phase ranges from Al 2 Mg 3 Zn 3 to Al 6 CuMg 4 described as isomorphous Mg 3 (Al,Cu,Zn) 2. 5,11,14,16) S-phase has a relatively fixed composition as Al 2 CuMg. 5,11) The aluminum corner of the Al Mg Zn diagram shows that higher Zn:Mg ratio exceeding 2.2 favors MgZn 2 (), and lower Zn:Mg ratio less than 2.2 favors Al 2 Mg 3 Zn 3 (T). 11,20) Higher density of M-phase in the high Zn:Mg ratio alloy in Fig. 6 would be explained in this regards. However, Cu-free ternary Al 2 Mg 3 Zn 3 is known to be stable below about 623 K, 20) suggesting that the T-phase in the low Zn:Mg ratio alloy in Fig. 6bis not a ternary phase. As a real alloy system, Fig. 10 shows the quaternary Al Zn Mg Cu system at 733 K for 6 mass%zn. 11) M-phase is completely taken into solid solution at this temperature. Alloy A and B are still in the (+S) field, but is sided toward the (+S+T) fields due to the higher magnesium contents. C Zn:Mg ratio Fig. 9 Master diagram showing the effects of homogenization conditions and alloy chemistries on the evolution of T- and S-phases from DSC tests. D E Fig. 10 Quaternary Al Zn Mg Cu system at 733 K for 6 mass%zn. 11) Therefore, T-phase would be more stable at high temperature in the Mg-rich alloy B (Zn:Mg = 2.1) than in the alloy E (Zn:Mg = 3.0), which is consistent with the Ref. 13). Furthermore, when cooled to around 673 K, the region of T- phase in Fig. 10 is expected to extend toward the lower magnesium side, stimulating the precipitation of T-phase. At the still lower temperature around 623 K, high temperature M-phase starts to precipitate and competes with the T- phase. 15,20) In this respect, T-phase in the low Zn:Mg ratio alloy in Fig. 6 is a quaternary T-phase (Al 2 Mg 3 Zn 3 to Al 6 CuMg 4 ) formed at higher temperature. It is ascertained by the noticeable trace of copper in the EDS analysis. The minimum temperature below which T-phase disappears is around 673 K. It is determined by the interrupted quenching to a given temperature from 523 to 723 K and holding for a prolonged time (5 h). In this experiment, T-phase melting peak is pronounced at around 693 K. Therefore, the appearance of the T-phase depends on the constitution (Zn:Mg ratio) and duration at higher temperature (cooling rate). 3.4 Quench sensitivity Figure 11 shows TEM micrographs from the intermediate forgings which are quenched with very slow cooling rates about 4 Ks 1. The alloy B (Fig. 11a) exhibits a large number of very coarse platelets about 200 nm long which are heterogeneously nucleated on the E dispersoids (Al 18 Mg 3 Cr 2 ). Contrarily, the alloy E (Fig. 11b) is characterized by the precipitation of the dispersoid (E-phase) Fig. 11 TEM micrographs of intermediate forgings from alloy B (a) and E (b): The forgings are quenched from 743 K at the rate of 4 Ks 1 and aged at 380 K for 7 h and 450 K for 7 h.

7 Control of Equilibrium Phases (M,T,S) in the Modified Aluminum Alloy 7175 for Thick Forging Applications 187 particles about 50 nm and the fine and uniform matrix ( 0 þ ) phases, suggesting the lower quench sensitivity through the homogeneous nucleation of precipitates. This is consistent with the Fig. 6bwhere coarse and heterogeneous precipitation of M-phase is observed in alloy B. The quench sensitivity generally depends first on the nature of dispersoids, and secondly on the total alloying content (Cu+Mg+Zn). 6,21) Copper addition in Al Zn Mg alloys increases the quench sensitivity by reducing the solubility of zinc and magnesium and thereby increases the supersaturation. At the same time, as in the Cu-free 7xxx alloys, the lower Zn:Mg ratio leads to higher strength but higher quench sensitivity. 4,15) Excess magnesium in the lower Zn:Mg ratio alloy enhance the kinetics of precipitation of coarse (precipitates during slow quench cooling which are heterogeneously nucleated on grain boundaries or Crcontaining dispersoids. 4) This would decrease the homogeneous nucleation temperature for the GP zones to grow to a stable size. 22) Therefore, the minimized quench sensitivity in the alloy E is connected with the lower level of (Cu+Mg) content, and with the higher Zn:Mg ratio. Consequently, the 7175 alloy redesign proposed in this study is effective for the minimizing the low melting point T and S-phases, providing a flexible heating range, T 20 K or more. Together with low quench sensitivity, the redesigned alloy with microstructural homogeneity would ensure higher mechanical and fracture properties, and better hot workability, which is expected to be uniquely applicable in 7175 thick forgings. 4. Conclusions (1) The evolution of equilibrium phases (M,T,S) in the 7175 alloys are greatly influenced by the limited change of alloy design, notably Zn:Mg ratio during the primary processing of large ingots. (2) The modified alloy design with lower (Cu+Mg) contents and higher Zn:Mg ratio (3.0) shows the lower S- phase solvus compared to the conventional 7175 alloy with higher (Cu+Mg) contents and lower Zn:Mg ratio (2.0), leading to the controlled amount of the coarse S-phase. Slow cooing after homogenization promotes the precipitation of the T-phase, but become ineffective as the Zn:Mg ratio increases. (3) High Zn:Mg ratio alloy has remarkably reduced quench sensitivity by promoting homogeneous precipitation, providing the usefulness in the 7175 thick forgings. Acknowledgments The authors wish to acknowledge ADD for the financial support, and formerly Dooray Air Metal (now Alcoa Korea) for the primary processing of the materials. REFERENCES 1) J. T. Staley: Treatise on Materials Science and Technology, Vol. 31, Alunimum Alloys-Contemporary Research and Application, ed. by A. K. Vasudevan and R. D. Doherty, (Academic Press, 1989) pp ) J. T. Staley: Metals Eng. Quarterly, May (1976) ) M. V. Hyatt: Aluminio 46 (1977) ) T. Sheppard: Extrusion of Aluminum Alloys, (Kluwer Academic Publishers, 1999) 81 86, ) J. E. Hatch: Aluminum, Properties and Physical Metallurgy, (ASM, Metals Park, 1984) ) I. J. Polmear: Light Alloys, Metallurgy of the Light Metals, 2nd ed., ed. by P. W. K. Honeycombe and P. Hancock, (Edward Arnold, 1989) pp ) J. Liu and M. Kulak: Mater. Sci. Forum (2000) ) P. Sainfort, C. Sigli, G. M. Raynaud and Ph. Gomiero: Mater. Sci. Forum 242 (1997) ) R. Shahani, T. Warner, C. Sigli, P. Lassince and P. Lequeu: Proc. Sixth Int. Conf. on Aluminum Alloys, Toyohashi, Japan, July 5 10, (1998) ) A. J. Morris, R. F. Robey, P. D. Couch and E. De los Rios: Mater. Sci. Forum 242 (1997) ) D. J. Strawbridge, W. Hume-Rothery and A. T. Little: J. Inst. Metals 74 (1948) ) X.-M. Li and M. J. Starink: Mater. Sci. Forum (2000) ) X.-M. Li and M. J. Starink: Mater. Sci. Tech. 17 (2001) ) D. Godard, P. Archambault, E. Aeby-Gautier and G. Lapasset: Acta Mater. 50 (2002) ) L. F. Mondolfo: Aluminum Alloys, Structure and Properties, (1976) ) J. A. Wert: Scr. Metall. 15 (1981) ) P. D. Couch, A. Burer and E. W. Sunam: Proc. 4th Int. Conf. on Aluminum Alloys, (Gorgia Institute of Technology, Atlanta, USA, 1994) ) N. Kamp, I. Sinclair and M. J. Starink: Metall. Mater. Trans. A 33A (2002) ) Hume-Rothery: J. Inst. Metals 70 (1944) ) H. Loffler, I. Kovacs and J. Lendvai: J. Mater. Sci. 18 (1983) ) A. Deschamps and Y. Brechet: Mater. Sci. Eng. A251 (1998) ) A. K. Mukhopadhyay: Metall. Mater. Trans. A 28A (1997)