In machining the hyper eutectic Al-Si alloy, the severe flank wear and the considerable deposit on the flank wear land of the tool are observed. This report deals with the deposit produced under several cutting conditions by EPMA to study the relation between the flank wear and the deposit.
The following results were obtained.
(1) The composition of the deposit on cutting edge side of the flank wear land was nearly the same as the work material, and silicon was principally adhered on the wear limit side.
(2) The distribution of aluminum and silicon, which are the main elements of the deposit, was related to the width of the flank wear almost independently of cutting tool materials, cutting oils, cutting speeds, etc.
(3) The primary silicon crystal in the deposit is crushed by abrasive action when it passes through the flank wear surface.
(4) It seems that the influence of silicon deposit on the growth of the flank wear was nearly the same as that of aluminum deposit in dry cutting, but the former was less influencial than the latter in wet cutting.

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The relationship between the microstructure and the stress corrosion susceptibility, which was tested in 3.5% aqueous salt solution by loading 80% of 0.2% proof stress on 5-9%Mg alloys, was investigated.
The stress corrosion was suppressed by the cathodic current, on the other hand, the anodic current increased the susceptibility, and a good qualitative agreement was observed between the accelarated and natural tests. Considering the dependence of the stress corrosion susceptibility on the precipitation state of Al-Mg alloys aged isothermally or isochronally, it is shown that the susceptibility depended mainly on the precipitation behavior at grain boundaries, and that the maximum susceptibility corresponded to the minimum spacing of precipitated β phase particles at grain boundaries.
The stress corrosion behavior of Al-Mg alloys can be explained by the precipitation rate, which is deduced from the degree of supersaturation and the diffusion rate of solute atoms, and by the sensitizing timetemperature curve. The main controlling factor of the stress corrosion cracking of Al-Mg alloys would be the localizedanodic dissolution of precipitates at grain boundaries.

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Effects on the stress corrosion cracking in Al-Mg alloys of the accelerating anodic current, the temperature of solid solution treatment, the anodizing treatment and the addition of some elements were investigated.
Between the logarithm of anodic current density and the logarithm of time to failure, a linear relationship was observed, but this was altered by the applied stress. As a result of considering this relationship, it is shown to be necessary to take the mechanochemical reaction into cosideration in the natural stress corrosion, and thus the one-step process mechanism in stress corrosion would not seem to hold in this case.
The stress corrosion susceptibility varied with the temperature of solid solution treatment due to changes of the grain size of alloys. The anodizing was not always effective, since the susceptibility of the alloys rather in creased when sharp notchs along grain boundaries were formed during anodizing due to the preferred dissolution of precipitates. Finally, marked effect of Cr addition on the stress corrosion of Al-Mg alloys was observed. A small addition Zr was very effective and Mn was also effective at lower applied stress. Y and Bi were little effective.

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In notched fracture tests of 5083 alloy at low temperatures such as notched tensile, Charpy impact and tear tests, numerous sharp cracks are produced perpendicular to the fracture surface and result in laminated fracture. Behaviors of sub-cracking at -196°C were investigated by V-notch Charpy tests and the nature of the laminated fracture was discussed.
It has been shown that sub-cracks are formed always in the plane of the plate and propagates along grain boundaries which are elongated in the rolling direction. For equi-axed grains, no cracking is observed. Precipitation of β-phase at grain boundaries markedly accelerates such cracking as the unit propagation energy (UPE) in the tear test in the short transverse direction decreases. Intermetallic constituents are also associated with cracking but are not the main reason for cracking.
Based upon the observations, it is concluded that cracking is caused by the intergranular fracture in the plane of the plate as a result of the precipitation of β-phase. It has also been found that the total length of sub-cracks linearly decreases as increasing the product of the tensile strength and UPE in the short transverse direction.

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For the purpose of grain boundary diffUsion experiments, 4mmφ rods with coarse bamboo-structure were prepared by cold working and recrystallization of 99.999% high pure aluminum ingot as received. These rods were covered with about 10μ thick Zn layer by electroplating, and diffusion-annealed at 250°C for up to 2000hr. After diffusion annealing, these rods were cut into an adequate form for the determination of concentration distribution around the grain boundary by line scanning of E. P. M. A.
Experimental data were analyzed by many solutions for grain boundary diffusion such as Fisher's, Whipple's, Suzuoka's and finite-thickness source solutions. The last solution was introduced by the present authors.
The results obtained concerning with the penetration of Zn along random boundaries of pure aluminum are as follows:
(1) The grain boundary diffusion experiment of Zn in aluminum may be done most adequately at about 250°C. At lower temperatures it is very time consuming, and at higher temperatures it is very diffcult to separate boundary diffusion from bulk diffusion.
(2) The penetration depth of Zn along grain boundary increases with the orientation difference of neighboring grains.
(3) Under the conditions of 10μ thick Zn layer and 250°C×1000hr diffUsion heating, C=0.01 contour obtained by Whipple's solution is in the best fit with the experiments, from which 0.57×10^-14cm³/sec is adopted as the most probable value of aD', where aD'is the mean value of aD'in the range ofrelative orientation difference of 5∼30°.

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