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deck face locations. In general, the ten- sile strength of 351 alloy is higher than that of A356+0.5%Cu alloy. Te elongation of 351 alloy is, how-


Fig. 9. This SEM fractographic image shows a crack initiated from a single large pore in the high pressure oil line sample of the 351 alloy with coarse microstructure.


Fig. 8. SEM fractographic images show crack initiation from multiple small pores and quick shearing of material between pores in the deck faces of the 351 alloy with fine microstructure. One fractured sample failed at 498,749 cycles (a), and another failed at 1,988,585 cycles (b).


refined cylinder heads may be attrib- uted to an increase in the amount of oxides generated when grain refiner was introduced into the melt. Te grain size remained large in both grain refined and non-grain refined heads, above 500µm, in both fast and slowly solidi- fied locations. Figure 4 and Table 2 show a


comparison of the tensile properties between the 351 and A356+0.5%Cu alloys at various test conditions, for specimens taken from cylinder head


ever, lower than that of A356+0.5%Cu alloy. Te improved tensile strength may be attributed to an increased amount of Q precipitates in the aluminum matrix, and zirconium- and vanadium-containing dispersoids in the aluminum matrix and at the grain boundaries, as shown in Figure 5. Figure 6 shows the creep strain as a


function of exposure time for both 351 alloy and A356+0.5%Cu alloy tested at 300C and 22MPa. At the stress and temperature tested, 351 alloy is clearly superior to A356+0.5%Cu alloy, par- ticularly with the increase of exposure time. Te drastic improvement of creep resistance of the 351 alloy is attributed to the presence of fine, semi-coherent and thermally extremely stable zirconium- and vanadium-containing dispersoids formed during solution treatment as


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