The flow of matrix materials in the gap of fibers was investigated. A multilayer aluminum alloy as a model material was used for strain measurement. An evaluation method for large plastic strain after Wolak-Parks was applied to investigate the consolidation process of mono-sheets. The flow of matrix material is considerably small in the vicinity of pressure-directional side and the normal side of fibers, until the matrix material fills up the gap of fibers. A logarithmic relation is found between the fill-up load and the combination of diameter and spacing of the fibers. The pressure required for full consolidation shall surpass the fill-up load. An effective application of multilayer materials and the Wolak-Parks method to metal forming is established.

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Polycrystalline high purity aluminum was epitaxially regrown from seed crystals having various combinations of crystallographic orientations. Stable growth of a grain potentially depends on a prefered orientation during solidification. Grains preferentially grow in <100> direction under the solidification condition in which the solidliquid interface forms hexagonal cells. The preference of <100> orientation is emphasized under the conditions in which the interface forms broken cells and dendrites. These facts lead to succesful preparation of polycrystalline aluminum having crystallographic orientations three dimensionally controlled.

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The composites were unidirectionally solidified at rates 0.5 to 20cm/hr and at temperature gradients 40 to 50°C/cm. The structure is arranged at slow freezing rates. The fibers change in shape from rod to blade as lowering the freezing rate. The shape and interfacial spacing of the fibers Al₃Fe in which manganese dissolves depend on the solidification rate. The arranged structure leads to high tensile strength. Thick fibers solidified at slow freezing rates cause lowering the tensile strength. The composites have tensile strength improved by about 2 times at 400°C than the binary Al-Fe eutectic composite.

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A method for thermal analysis of casting dies has been developed using an analogue analyzer. It is more convenient and rapid than the Acurad analysis which is stepwise and time-consuming. Electric potential distribution corresponding to the temperature distribution on the inner and outer surfaces and cooling channels is applied, in this method, at a time on electro-conductive paper. The simulated temperatures agree with real temperatures in an aluminum casting die within ± 10°C.

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The tests on 1100-0 and 5056-F alloys concern influences of contact conditions of the tool chip on the stress state in the vicinity of tool tip and the inner part of built-up edge. A little effect of the nose radius of a single rake tool on the cutting force and the deformed layer is found, because of adhering the built-up edge. The cutting force and thickness of deformed layer decrease with decreasing the 1st rake angle in a double rake tool, because the equivalent rake angle in which the built-up edge is formed on the 1st rake face increases. This fact is also photoelastically ascertained. The compressive stresses in the built-up edge are photoelastically observed both parallel and normal to the cutting direction. These stresses seem to increase with approaching to the rake face. One normal to the cutting direction decreases with approaching to the tool flank.

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Tensile tests were carried out using a large numbers of JIS No. 5 and 13A types specimens of 5052-0 plates 3mm thickness and 7N01-T6 5mm thickness. The specimens in Barba type materials such as 5052-0 alloy frequently fracture near the end of the gage length, but ones in Oliver type materials such as 7N01-T6 rupture near the center of the gage length. The elongation δ₀ is almost independent of the fracture position in 5052-0, but decreases in 7N01-T6 as fracture occurs near one end of the gage length. The presumed elongation δ₁ were calculated in the method JIS Z 2241 (Method of Tensile Test for Metallic Materials) for the specimens in which fracture occurred near the end of the gage length. δ₁ is neary equal to δ₀ for the Barba type materials independent of the fracture position, but is larger than δ₀ for the Oliver type materials as fracture occurs near one end of the gage length.

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Is is well known that a small amount of iron in aluminum markedly retards its recrystallization ¹⁾ and brings about R-orientation in annealing textures of cold rolled alloys ²⁾, which is generally attributed to the fact that iron precipitates during annealing and interacts with progress of recrystallization ^{1), 3), 4)}.
Iron in aluminum supersaturated by solid quenching ^{4), 5)} to such an extent as about 0.04wt% is possible to precipitate even below 300°C in a cold worked matrix, while precipitation kinetics in an undeformed matrix follows a C-curve with a nose at about 470°C and iron does not precipitate practically below 300°C. Electron microscopical observations have revealed that iron in an undeformed matrix precipitates between 300°C and 500°C as flat needles ^{4), 6)} in the <100>_Al directions with the {110}_Al habit planes and that iron in a deformed matrix precipitates on subboundaries as irregular spheres 1), 3), 4). The spherical precipitates ³⁾ and the born-like precipitates ⁷⁾ formed by longer annealing above 500°C have been tentatively identified by electron diffraction as a stable compound, Al₃Fe, though some ambiguity was left. The flat needle precipitates are, however, not yet reported to be identified directly either by electron or X-ray diffraction. A measurement ⁵⁾ of Mössbauer effect in an Al-0.05wt%Fe alloy showed that only Al₃Fe was present in an undeformed matrix aged between 370°C and 520°C, while a metastable compound, Al₆Fe, was prominant in a cold rolled matrix annealed between 270° and 420°C. The correspondence of the electron microscopical observations to the results of the Mössbauer effect measurement seems to have not been discussed yet.
In the case of aging of highly supersaturated alloys (several wt%Fe) formed by splat quenching ^{8), 9), 16)}, precipitation of Al₆Fe and Al₃Fe is reported to occur below 400°C and above 450°C respectively, and formation of the former phase at an earlier stage followed by growth of the latter phase is observed at intermediate temperatures.
The Al₆Fe phase is also formed during unequilibrium solidification of Al-Fe alloys, where two more Al-Fe compounds, a tetragonal phase in an Al-0.5wt%Fe alloy ¹⁰⁾ and a monoclinic phase in an Al-0.5wt%Fe-0.2wt%Si alloy ¹¹⁾, were found together with Al₃Fe and Al₆Fe. Wheather these two metastable phases concern with the aging process of solid quenched Al-Fe alloys is yet unknown.
Recent investigations ^{12), 13)} of recrystallization textures of cold rolled and annealed Al-Fe alloys have shown that relative amounts of R-orientation to cube-orientation can be remarkably varied by final annealing temperatures. One of the reasons is certainly dependence of precipitating quantity in a period of recrystallization on annealing temperatures ⁴⁾, but the possibility that precipitating phases may be also changed by annealing temperatures must be pointed out when the results of the Mössbauer effect measurement and of the splat quenched alloys are taken into consideration ¹³⁾.
An electron-microscopical work is consequently being done to identify precipitates formed during recrystallization annealing in a cold rolled Al-Fe alloy.
The 99.998% purity Al-0.03wt%Fe alloy was melted and cast in vacuo, solution treated at 630°C, cold worked, and intermediately annealed for 30s at 550°C in a salt bath. A specimen was further cold rolled to 0.2mm thickness by the reduction of 95% and finally annealed for 16h at 280°C in the salt bath. Photo 1 shows precipitates which are very likely to have been formed initially on subboundaries and left in a recrystallized grain after migration of recrystallization front.

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