The forming mechanism of chromium bearing compounds (E phase) in a chromium containing Al-Zn-Mg-Cu alloy during the homogenizing treatment was studied. The structures of such specimens as 1) homogenized at 470°C (T_H) after heating at different rates to T_H, 2) cooled from the temperatures below T_H during heating to T_H, 3) preaged at 300-400°C, and 4) pre-aged and then up-quenched to T_H were examined. E phases precipitate heterogeneously on the periphery of η phases formed during heating to T_H or during pre-aging. The precipitates of E phase are retained after disolution of η phases, and they Ostwald-ripen during homogenization. The achievement of uniform distribution of η phases is the most important for that of E phases.

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The molten metal was injected in a disk shaped die cavity through a gate in the tangential direction to the disk cavity under various conditions. The molten metal flows separately into the upper and lower sides of the cavity and meets at a point 225° from the gate independently of the injection condition. The velocity of molten metal lowers down to 2m/s at a final filling portion where a number of surface defects appear, as the molten metal fills the die cavity. The defects can be reduced by accelerating molten metal velocity near this portion. Water analogue was performed using the same shape of a plastic die. The mold filling behavior and effect of velocity decrease of molten metal can be simulated.

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Effects of enviromental temperature on the failure time in stress corrosion cracking of Al-6.12wt%Zn-2.05wt% Mg alloy were investigated by the Weibull distribution method. Test pieces were solution treated at 460°C for 1h, quenched into ice-water, aged at 120°C for 24h and tested at temperatures from 13° to 53°C under a constant applied stress of 80% of the yield stress in 3.5%NaCl+0.2%H₂O₂ aqueous solution of pH 5.2. Both failure time and crack initiation time obey the single Weibull distribution with location parameters. The former time depends on the latter time. Their shape parameters are independent of temperatures and equal to 1.5, which corresponds to the parameter of wear out failure. Stress corrosion cracking of this alloy considerably depends on temperature. Apparent activation energies evaluated from Arrhenius plots of reciprocals of mean values are 16.4kcal/mol in failure time, 17.4kcal/mol in crack initiation time, and 13.1kcal/mol in crack propagation period.

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Two sorts of Al-5.6%Zn-2.4%Mg-1.3%Cu alloys with 0.12%Zr and 0.26%Cr, respectively, were hot-rolled by 86% at 400°C and then solution-treated at 470°C. Transmission electron microstructures of both alloys in hot-rolled state are nearly the same. Subgrain structures of both alloys developed after disappearance of the η-particles are quite different. Zr-containing alloy shows an almost perfect subgrain structure, reflecting that the particle size of Al₃Zr is much smaller than that of E phase. In Zr-containing alloy, recrystallized grains have naturally appeared faster than in Cr-containing alloy and the grain growth is retained. Elongated grains in the rolling direction are commonly observed in both recrystallized alloys, and are attributed to the layered distribution of Al₃Zr or E particles in the rolled sheet.

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Tensile and Vickers and Rockwell hardness tests were carried out on a number of aluminum wrought alloys. The alloys are classified into the following three groups according as the relation between tensile strength and Vickers hardness Hv: (1) S-group; O, H112, F, T1 and T4 tempered alloys for which if Hv>45, σ_B=0.40Hv and if Hv<45, σ_B=0.35Hv, (2) H-group; work-hardened alloys for which σ_B=0.35Hv-3.0, and (3) T-group; heat-treated alloys except for those in S-group for which σ_B=0.32Hv and for 6xxx alloys σ_B=0.30Hv. In Rockwell hardness test, B-scale is appropriate for T-group alloys and F-scale for S- and H-group alloys. A definite chart for converting hardness to tensile strength is also shown.

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7475 alloy plates were cold rolled into sheets 1.6mm in thickness at a rolling reduction 73%, recrystallized and solutionized at 480°C at different heating and cooling rates, and were finally cold rolled to 1.2mm thickness. Tensile tests under an increasing temperature condition or hot tensile tests after warm prestraining were performed. The sheets rapidly heated and rapidly cooled in the intermediate heat-treatment show superplastic elongation 500% or more under an increasing tempuature condition and at a strain rate 2.8×10^-3/s which is one order of magnitude higher than that informed for intermediate thermomechanically treated 7475 sheets up to date. The same sheets prestrained by 15 to 30% at about 360°C also show superplasticity at the hot tensile tests. Fairly fine but elongated grains containing considerably fine subgrains 2 to 10μm in diameter are developed by ITMT and the warm prestraining. These subgrains would play a vital role during superplastic deformation.

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