Materials Transactions, Vol. 44, No. 1 (2003) pp. 181 to 187 #2003 The Japan Institute of Metals EXPRESS REGULAR ARTICLE

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

Seong Taek Lim1;*, Il Sang Eun2 and Soo Woo Nam1

1Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1Guseong-dong, Yuseong-gu, Daejeon, 305-701,Korea 2Agency for Defence Development, P.. Box 35-5, Yuseong-gu, Daejeon, 305-600, 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 (Al2CuMg) is effectively inhibited by higher Zn:Mg ratio rather than higher solutionizing temperature. The formation of M-phase (MgZn2) and T-phase (Al2Mg3Zn3)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 system owing to the extended solubility of copper into the precipitates. Besides the minor phases ( =Al2Cu, = Aluminum alloy AA7175, one of the newer variants of the Mg2Zn11), the resulting microstructure is generally charac- baseline alloy AA7075, has primarily been designed for the terized by equilibrium phases of soluble M (MgZn2 = ), T 1–3) forging requiring higher strength and damage tolerance. (Al2Mg3Zn3) and their isomorphous phases, S (Al2CuMg), 4,5,11–17) Because of the same alloy composition with AA7075 and insoluble Al7Cu2Fe particles. These phases (5.6Zn–2.5Mg–1.6Cu–0.23Cr) except the lowered impuri- constitute most of the microstructure and govern the general ties, AA7175 is well characterized in the casting, hot working properties. M-phase and its precursors GP zone and inter- and heat treatment areas in the case of small section mediate precipitate (0) are responsible for high strength. thickness. Low melting point T- and S-phases affect hot workability and As a thick forging, however, AA7175 with conventional fracture properties.4,8,10,12,17,18) S-phase has been a main alloy composition is not a good choice due to the massive target to be prevented in the thick 7X49 and 70X0 heavy evolution of coarse equilibrium phases and high quench plates due to its massive appearance and brittleness.8,10,12,18) sensitivity, and therefore usually replaced by 7X50 type The decompositions of these phases are strongly dependent alloys.1–6) In this respect, the alloy should entail the special on the constitution of alloying elements such as controls on the alloy composition and processing for (Zn+Mg+Cu) and Zn:Mg ratio, and on the cooling rate microstructural homogeneity. following solutionizing.3,4,8,10–13,17) Furthemore, the amount There has been a strong tendency for 7xxx alloys to be of equilibrium phases (M,T,S) is essentially determined at the used as large and thick semiproducts.7–10) These thick primary processing stages such as casting and homogeniza- materials, however, often have poor mechanical and fracture tion, eapecially in the less deformed thick materials.4,8–10) properties since they receive less deformation during hot From the backgound mentioned above, the modification of working and slower quench cooling.8–10) This is especially alloy composition for the specific purpose of thick 7175 due to that thick (open die) forgings are usually made of large forging is attempted to control the equilibrium phases. diameter ingots, reaching several hundred milimeters, which Recent studies on this area are concentrated on the 70X0 accompany large dendrite size and concurrent heavy segre- ( ¼ 1; 4; 5) type plate alloys8–10,18) and mostly regarded as gations owing to slow casting speed.4) Through the many proprietaries.9,10) The constitutional study on the 7X75 alloy slow processes, there are frequent evolutions of the coarse forging is not reported yet. In this respect, the effects of alloy equilibrium phases and constituent particles which have an redesign, especially Zn:Mg ratio, and the cooling rate on the adverse effect on strength, damage tolerance and hot work- microstructural evolutions are investigated using real scale ability.8–10,17,18) big ingots and thick forgings of 7175 alloys. From the earlier studies by Strawbridge et al.11) to the recent studies by Li and Starink,12,13) complex phases are 2. Experimental Procedure evolved in the commecial quaternary (Al–Zn–Mg–Cu) The materials used in this study are large forging ingots of *Graduate Student, Korea Advanced Institute of Science and Technology. 7175 type alloys with 700 mm in diameter prepared by DC Present address: Agency for Defence Development, P.O. Box 35-5, semicontinuous casting at Dooray Air Metal in Korea (now Yuseong-gu, Daejeon, 305-600, Korea. 182 S. T. Lim, I. S. Eun and S. . 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 AA7175 0.35 max 1.2–2.0 2.1–2.9 5.1–6.1 3.3–4.9 8.4–11.0 1.8–2.9 Alloy A 0.21 1.74 2.80 5.47 4.54 10.01 1.9 Alloy 0.16 1.50 2.65 5.63 4.15 9.78 2.1 Alloy 0.13 1.63 2.43 5.78 4.06 9.84 2.4 Alloy 0.14 1.54 2.05 5.80 3.59 9.39 2.8 Alloy 0.14 1.32 2.01 6.04 3.33 9.37 3.0

Alcoa Korea) as a 7xxx big ingot feasibility study program. 6.6 KÁminÀ1 from 733–743 to 523 K are obtained. Casting speed and other variables are carefully controlled to Microstructural characterizations are performed by using avoid ingot cracking and ensure internal structure. Five kinds optical microscopy (OPTIPHOT 100S, NIKON), scanning of alloys given in Table 1 are designed to have different electron microscopy (XL30FEG, PHILIPS) with EDS, and Zn:Mg ratio ranging from 1.9 (alloy A) to 3.0 (alloy E) within transmission electron microscopy (JEOL 4000FX). Thin foils or around the specification limit of commercial AA 7175. for TEM are prepared by mechanical polishing to 150 mm and Contents of chromium are kept minimum (0.18%), and final twin-jet electropolishing in a solution of Fe+Si as low as 0.15 mass% except alloy A. The richer alloy 25%HNO3+75%CH3OH at À20 K and 10 . group (B,C) and the more dilute alloy group (D,E) have Calorimetric studies are performed in the 2910 DSC respectively the similar total alloying contents. The dilute module in TA Instrument with pure aluminum reference and alloy E has relatively lower Cu and higher Zn:Mg ratio. the sample discs (about 80 mg) at the rate of 10 KÁminÀ1 from The 25 mm square samples are taken at the mid-radius 303 to 803 K. The results are average of three runs. 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 3. Results and Discussion open die forging into the intermediate forgings (380 mm diameter, "  1:2) at between 673 and 623 K using 2000 ton 3.1 General Microstructures hydraulic press. The 25 mm square samples are taken at the Figure 1 shows the microstructures of 700 mm diameter mid-radius location of the forgings. Homogenizing treat- ingot (alloy B) in as-cast and as-homogenized conditions at ments and subsequent coolings are also performed in the mid-radius, indicating the change of the microstructure with DSC module where precise coolings at the rate of 0.1– 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 , and d) fully homogenized at 733 K for 24 h and 743 K for 16 h. 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 -0.275 Alloy E boundary thickness 25 mm, which is much larger than those -0.300 of small ingots typically 30 mm and 10 mm, respectively. D The large cell is difficult to be homogenized due to the large -0.325 B interdendritic segregations. The main constituents of inter- A dendritic network are identified by EDS mostly as Al–Cu -0.350 T-phase lamella eutectic with very limited amount of magnesium and -0.375 zinc, and Al–Cu–Fe compounds. The former is generally E (Zn / Mg=3.0)

4) W / Heat Flow, designated by eutectic [M] or T [Al–Cu–Mg–Zn]. These are -0.400 S-phase D (Zn / Mg=2.8) severely cored by copper-rich outer rims. In the slightly B (Zn / Mg=2.1) -0.425 A (Zn / Mg=1.9) homogenized state (Fig. 1c), the Cu-rich eutectics are Homo. Cooling Rate : 0.67K min-1 partially dissolved and become discrete particles and some -0.450 730740 750 760 770 780 790 800 T (Al2Mg3Zn3) phase are seen at the cell boundary as well as within the grain. In fully homogenized state (Fig. 1d), the Cu- Temperature, T / K rich eutectics are converted into S (Al CuMg), and Al–Cu– 2 Fig. 2 DSC thermograms at high temperature interval of the alloys in the Fe compounds into Al7Cu2Fe particles. Impurity-originated homogenized and cooled at the rate of 0.67 KÁminÀ1 following homo- Al7Cu2Fe particles are very coarse and bulky, varying 5– genization at 733 K for 12 h. 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, 15% in areal fraction. The major constituent particles are also varying 2–10 mm. The matrix is characterized by coarse Al7Cu2Fe and S-phase, aligned along the working axis. The platelets of M (MgZn2) and occational T (Al2Mg3Zn3) with large quantity of S-phase decorating the grain boundary are polygonal shape, which are high temperatute precipitates seen in the alloy A and B, which is in contrast to the clean decomposed during cooling after homogenization. alloy E. The resulting microstructure is generally characterized by The evolution of S-phase through the constitutional the equilibrium phases of soluble M-, T- and S-phase, and variations is reviewed from the two standpoints, Cu+Mg 19) insoluble impurity-originated Al7Cu2Fe particles. Platelets of and Zn:Mg ratio. In the classical work by Hume-Rothery, M-phase (often referred ) and polygonal T-phase are curve separating the and (+S) fields in Al–Cu–Mg is evolved depending on alloy composition, especially on the given by: Log[at%Cu]Á[at%Mg] ¼ 5:603 À 3975ÁTÀ1, where Zn:Mg ratio, and the cooling rate from solutionizing T is temperature in Kelvin. The solubility curves at the temperature. The evolution of the S-phase depends also on commercial solutionizing temperatures from 723 to 773 K are the alloy composition, and solutionizing time and tempera- solved and plotted in Fig. 4. Strawbridge et al.11) also ture. T- and S-phases have low melting points so that they presents the phase stability of Al–Mg–Cu–6 mass%Zn limit the heating range in hot working and solutioniz- system at 733 K, which is added in Fig. 4 for comparison. ing.4,13,17) The brittleness of the S-phase greatly impairs the Alloy D and E lie on a little outside the limit of the registered fracture properties.8,10) The large ingots are good examples AA7175 specification. The solubility of (Cu+Mg) at 733 K for the evolution of these phases because they subject to tends to extend in the quaternary system (Strawbridge’s). In severe segregation and very slow cooling after solutionizing. 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 3.2 Evolution of S-phase that at 733 K alloy A, B, and C are within the field of (+S), In the homogenized state, alloy A and B have large amount whereas alloy D and E are within -Al. The soluble S-phase of intermetallic phases, Fe-containing Al7Cu2Fe and Mg- is expected to be completely dissolved into the matrix at containing Al2CuMg (S-phase). This is in contrast to alloy E higher temperature over 763 K even in the alloy A. However, in which there are a little Al7Cu2Fe only, and is basically free this is impractical because of very slow dissolution rate of S- of S-phase. Large amount of Al7Cu2Fe in alloy A is due to phase and possible overheating at the highly segregated the relatively high iron content. Considering the alloy areas.4,5,9) The upper limits of the solutionizing temperature composition, it can be seen that as Cu+Mg content decreases in the 7xxx alloys are mostly determined by the S-phase or Zn:Mg ratio increases, the quantity of S-phase decreases. solvus.4,5) In this regards, Alloy E has a S-phase solvus about The microstructural evidences of the S-phase varying with 723 K which is about 40 K lower than that of alloy A. the alloy designs are consistent with the results of DSC tests Hyatt3) explained the phase stability of S-phase in shown in Fig. 2. It can be noted that alloy A and B show large 90 mass%Al–Zn–Mg–Cu system at 733 K in terms of Zn:Mg and sharp endothermic peaks, corresponding to the melting of ratio. Figure 5 shows that as Zn:Mg ratio (iso-ratio line) S-phase which onset at around 763 K. In the alloy D and E, varies from 2.0 to 4.0 for the constant copper level about however, S-phase melting peaks are diminished or comple- 2.0 mass%, the alloys move from a two-phase (+S) to a tely disappeared. single phase field. In addition, solubility of copper increases The evolution of the S-phase is similarly characterized in in high Zn:Mg ratio alloy which, as a design concept of the intermediate forgings as shown in Fig. 3. All alloys AA7050, has more Al2CuMg to be dissolved at the exhibit similar grain structures with less elongated grains solutionizing temperature.3–5) Although alloy A (90 mas- about 300 mm long, and with well developed subgrains of 3 to s%Al) only is exactly defined in the diagram, the appearance 8 mm diameter. Partial recrystallizations are kept less than of massive S-phase for the lower Zn:Mg ratio alloy is well 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.

3.0 753 763 773K 743 733 2.5 723 (Al) + S

2.0 A C B D 1.5 (Al) Cu, mass % Cu, E 1.0 --- Hume-Rothery at 723-773K Strawbridge at 733K 7175 Limit 0.5 0.51.0 1.5 2.0 2.53.0 3.5 Mg, mass % Fig. 5 Phase diagram of quaternary 90 mass%Al–Zn–Cu–Mg system at 3) Fig. 4 S-phase stability of the alloys in ternary Al–Cu–Mg (Hume- 733 K, showing the effect of changing Zn:Mg ratio. Rothery)19) and quaternary Al–6%Zn–Cu–Mg (Strawbridge)11) systems.

morphology during cooling from homogenizing temperature. explained. Reflecting that alloy group (B,C) and (D,E) have Figure 6 shows the micrographs of alloy B and E which are essentially the same respective (Cu+Mg) contents (Table 1), fast cooled or slowly cooled during critical temperature it is reasonable to say that the formation of the S-phase interval between 733 and 523 K at the rate of 3.3 and depends first on the Zn:Mg ratio, and second on the (Cu+Mg) 0.1 KÁminÀ1, respectively. In the fast cool (Figs. 6a, c), both content around the limited range of AA7175 specification. alloys reveal the precipitation of platelets of M-phase only. However, in the slow cooled alloy B with lower Zn:Mg ratio, 3.3 Evolutions of M- and T-phases coarse precipitates of T-phase about 5 mm are seen at both Other equilibrium phases of concern are M- and T-phases. cell boundary and grain interior, along with coarse platelets Equilibrium phases of M and T reprecipitate in a coarse of M-phase about 10 mm. White particles of T-phase are seen 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 KÁminÀ1, b) and d) 0.1 KÁminÀ1. only by the backscattered electron image in the SEM. There to the fast cool (Fig. 2), separated double endothermic peak T are only fine and uniform M-phases in the higher Zn:Mg ratio and S are clearly evident in the slowly cooled from alloy E. Generally, given a same cooling rate, M-phase homogenization temperature. The peak T at about 752 K precipitates in a fine and uniform manner in the alloy E, but (onsets at 748 K) is responsible for the melting of T-phase coarse and heterogeneous manner in the alloy B. The greater reformed during slow cooling at high temperature, and peak density of M-phase in the alloy E will be discussed below in S at about 771 K (onsets at 763 K) for the melting of S-phase terms of the Zn:Mg ratio and the quench sensitivity. already formed during homogenization.12,13,17) Peak of T- Figure 7 shows the DSC thermograms for the higher phase melting is not observed in the high Zn:Mg alloy E. temperature region showing the effect of cooling rate Figure 8 shows the cooling rate dependence on the following homogenization. Single endothermic peak at dissolution energies associated with the melting of T- and 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 2.4

Alloy E -0.04 2.0 C.W..

-0.06 1.6 Alloy D

-0.08 1.2 Alloy B S Peak S T As-cast Alloy B 0.8 -0.10 Peak T

Heat Flow, W/g Heat Flow, 0.4 -0.12 /g Dissolution Energy, C.W.Q. Homo. Cooling Rate 0.0 -0.14 -1 : 0.1K min 042 6 730 740 750 760 770 780 790 800 6000 Temperature, T / K Cooling Rate, K min-1

Fig. 7 Evolutions of T and S-phase in the very slow cooling (0.1 KÁminÀ1) Fig. 8 Homogenization cooling rate effect of the alloy B on the evolution following homogenization at 733 K. of T and S-phases from DSC tests (cooled after holding at 733 K for 12 h). 186 S. T. Lim, I. S. Eun and S. W. Nam

5.0 4.5 S-phase, 0.1K min-1 from 733K -1 4.0 T -phase, 0.1K min from 733K -1 3.5 S-phase, 1.67K min from 743K -1 3.0 T -phase, 1.67K min from 743K 2.5 2.0 1.5 1.0 0.5

Dissolution Energy, J/g Dissolution Energy, 0.0 -0.5 A B C D E -1.0 1.82.0 2.2 2.4 2.6 2.8 3.0 Zn:Mg ratio Fig. 10 Quaternary Al–Zn–Mg–Cu system at 733 K for 6 mass%Zn.11) Fig. 9 Master diagram showing the effects of homogenization conditions and alloy chemistries on the evolution of T- and S-phases from DSC tests. Therefore, T-phase would be more stable at high temperature S-phases in the low Zn:Mg alloy B. T-phase is not appeared in the Mg-rich alloy B (Zn:Mg = 2.1) than in the alloy E when the critical cooling rate about 0.7 KÁminÀ1 is exceeded. (Zn:Mg = 3.0), which is consistent with the Ref. 13). The average cooling rate during critical range is measured as Furthermore, when cooled to around 673 K, the region of T- about 0.7 KÁminÀ1 at the one third of the radius in the 700 mm phase in Fig. 10 is expected to extend toward the lower diameter ingot. This suggests the possible occurrence of the magnesium side, stimulating the precipitation of T-phase. At T-phase within the large ingots. S-phase is moderately the still lower temperature around 623 K, high temperature decreased with increasing homogenization cooling rate. M-phase starts to precipitate and competes with the T- 15,20) However, it is observed that the fast cooling alone cannot phase. In this respect, T-phase in the low Zn:Mg ratio exclude S-phase effectively. alloy in Fig. 6 is a quaternary T-phase (Al2Mg3Zn3 to Figure 9 is a master diagram showing the effects of the Al6CuMg4) formed at higher temperature. It is ascertained by constitution, and temperature and cooling rate of the the noticeable trace of copper in the EDS analysis. The homogenization on the evolution of T and S-phases. It is minimum temperature below which T-phase disappears is clearly seen that slow cooing after homogenization promotes around 673 K. It is determined by the interrupted quenching the precipitation of the T-phase, but becomes ineffective as to a given temperature from 523 to 723 K and holding for a Zn:Mg ratio increases. Fast cooling in any alloy definitely prolonged time (5 h). In this experiment, T-phase melting retards the T-phase precipitation, but partly effective for the peak is pronounced at around 693 K. Therefore, the appear- S-phase. Higher homogenization temperature reduces the ance of the T-phase depends on the constitution (Zn:Mg amount of the S-phase somewhat, but higher Zn:Mg ratio is ratio) and duration at higher temperature (cooling rate). found to be more effective inhibitor. In any case, alloy E having the highest Zn:Mg ratio is basically free of both T- 3.4 Quench sensitivity and S-phases. Figure 11 shows TEM micrographs from the intermediate In the most of commercial high strength 7xxx alloys, the forgings which are quenched with very slow cooling rates quaternary compositions are not single phase even at the about 4 KÁsÀ1. The alloy B (Fig. 11a) exhibits a large number solutionizing temperatures.4) Through the participation of of very coarse  platelets about 200 nm long which are copper into the precipitates, M-phase ranges from MgZn2 to heterogeneously nucleated on the E dispersoids AlCuMg described as isomorphous Mg(Al,Cu,Zn)2, and T- (Al18Mg3Cr2). Contrarily, the alloy E (Fig. 11b) is char- phase ranges from Al2Mg3Zn3 to Al6CuMg4 described as acterized by the precipitation of the dispersoid (E-phase) 5,11,14,16) isomorphous Mg3(Al,Cu,Zn)2. S-phase has a rela- 5,11) tively fixed composition as Al2CuMg. The aluminum corner of the Al–Mg–Zn diagram shows that higher Zn:Mg ratio exceeding 2.2 favors MgZn2 (), and lower Zn:Mg ratio 11,20) less than 2.2 favors Al2Mg3Zn3 (T). 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 20) Al2Mg3Zn3 is known to be stable below about 623 K, 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. Fig. 11 TEM micrographs of intermediate forgings from alloy B (a) and E Alloy A and B are still in the (+S) field, but is sided toward (b): The forgings are quenched from 743 K at the rate of 4 KÁsÀ1 and aged the (+S+T) fields due to the higher magnesium contents. at 380 K for 7 h and 450 K for 7 h. 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 sensitivity by promoting homogeneous precipitation, provid- (0 þ ) phases, suggesting the lower quench sensitivity ing the usefulness in the 7175 thick forgings. through the homogeneous nucleation of precipitates. This is consistent with the Fig. 6bwhere coarse and heterogeneous Acknowledgments precipitation of M-phase is observed in alloy B. The quench sensitivity generally depends first on the The authors wish to acknowledge ADD for the financial nature of dispersoids, and secondly on the total alloying support, and formerly Dooray Air Metal (now Alcoa Korea) content (Cu+Mg+Zn).6,21) Copper addition in Al–Zn–Mg for the primary processing of the materials. alloys increases the quench sensitivity by reducing the solubility of zinc and magnesium and thereby increases the REFERENCES supersaturation. At the same time, as in the Cu-free 7xxx alloys, the lower Zn:Mg ratio leads to higher strength but 1) J. T. Staley: Treatise on Materials Science and Technology, Vol. 31, higher quench sensitivity.4,15) Excess magnesium in the Alunimum Alloys-Contemporary Research and Application, ed. by A. K. Vasudevan and . D. Doherty, (Academic Press, 1989) pp. 3–31. lower Zn:Mg ratio alloy enhance the kinetics of precipitation 2) J. T. Staley: Metals . Quarterly, May (1976) 137–142. of coarse (precipitates during slow quench cooling which are 3) M. V. Hyatt: Aluminio 46 (1977) 81–99. heterogeneously nucleated on grain boundaries or Cr- 4) T. Sheppard: Extrusion of Aluminum Alloys, (Kluwer Academic containing dispersoids.4) This would decrease the homoge- Publishers, 1999) 81–86, 232–236. neous nucleation temperature for the GP zones to grow to a 5) J. E. Hatch: Aluminum, Properties and Physical Metallurgy, (ASM, 22) Metals Park, 1984) 79–104. stable size. Therefore, the minimized quench sensitivity in 6) I. J. Polmear: Light Alloys, Metallurgy of the Light Metals, 2nd ed., ed. the alloy E is connected with the lower level of (Cu+Mg) by P. W. K. Honeycombe and P. Hancock, (Edward Arnold, 1989) content, and with the higher Zn:Mg ratio. pp. 54–104. Consequently, the 7175 alloy redesign proposed in this 7) J. Liu and M. Kulak: Mater. Sci. Forum 331–337 (2000) 127–140. study is effective for the minimizing the low melting point T 8) P. Sainfort, C. Sigli, G. M. Raynaud and Ph. Gomiero: Mater. Sci. Forum 242 (1997) 25–32. and S-phases, providing a flexible heating range, T  20 K 9) R. Shahani, T. Warner, C. Sigli, P. Lassince and P. Lequeu: Proc. Sixth or more. Together with low quench sensitivity, the rede- Int. Conf. on Aluminum Alloys, Toyohashi, Japan, July 5–10, (1998) signed alloy with microstructural homogeneity would ensure 1105–1110. higher mechanical and fracture properties, and better hot 10) A. J. Morris, R. . Robey, P. D. Couch and E. De los Rios: Mater. Sci. workability, which is expected to be uniquely applicable in Forum 242 (1997) 181–186. 11) D. J. Strawbridge, W. Hume-Rothery and A. T. Little: J. Inst. Metals 74 7175 thick forgings. (1948) 191–225. 12) X.-M. Li and M. J. Starink: Mater. Sci. Forum 331–337 (2000) 1071– 4. Conclusions 1076. 13) X.-M. Li and M. J. Starink: Mater. Sci. Tech. 17 (2001) 1324–1328. (1) The evolution of equilibrium phases (M,T,S) in the 14) D. Godard, P. Archambault, E. Aeby-Gautier and G. Lapasset: Acta Mater. 50 (2002) 2319–2329. 7175 alloys are greatly influenced by the limited change of 15) . F. Mondolfo: Aluminum Alloys, Structure and Properties, (1976) alloy design, notably Zn:Mg ratio during the primary 844–864. processing of large ingots. 16) J. A. Wert: Scr. Metall. 15 (1981) 445-47. (2) The modified alloy design with lower (Cu+Mg) 17) P. D. Couch, A. Burer and E. W. Sunam: Proc. 4th Int. Conf. on contents and higher Zn:Mg ratio (3.0) shows the lower S- Aluminum Alloys, (Gorgia Institute of Technology, Atlanta, USA, 1994) 10–17. phase solvus compared to the conventional 7175 alloy with 18) . Kamp, I. Sinclair and M. J. Starink: Metall. Mater. Trans. A 33A higher (Cu+Mg) contents and lower Zn:Mg ratio (2.0), (2002) 1125–1136. leading to the controlled amount of the coarse S-phase. Slow 19) Hume-Rothery: J. Inst. Metals 70 (1944) 491. cooing after homogenization promotes the precipitation of 20) H. Loffler, I. Kovacs and J. Lendvai: J. Mater. Sci. 18 (1983) 2215– the T-phase, but become ineffective as the Zn:Mg ratio 2240. 21) A. Deschamps and . Brechet: Mater. Sci. Eng. A251 (1998) 100–107. increases. 22) A. K. Mukhopadhyay: Metall. Mater. Trans. A 28A (1997) 2429–2433. (3) High Zn:Mg ratio alloy has remarkably reduced quench