Abnormally large grains have been observed in Al-Mg alloy AA5182 sheet material after forming at elevated temperature, and the reduced yield strength that results is a practical problem for commercial hot-forming operations. The process by which abnormal grains are produced is investigated through controlled hot tensile testing to reproduce the microstructures of interest. Abnormal grains are shown to develop strictly during static annealing or cooling following hot deformation; the formation of abnormal grains is suppressed during plastic straining. Abnormal grains grow by static abnormal grain growth (SAGG), which becomes a discontinuous recrystallization process when abnormal grains meet to form a fully recrystallized microstructure. Nuclei, which grow under SAGG, are produced during hot deformation by the geometric dynamic recrystallization (GDRX) process. The mechanism through which a normally continuous recrystallization process, GDRX, may be interrupted by a discontinuous process, SAGG, is discussed.

The influence of texture on the bendability in CuNiSi alloys was examined by using an age-hardened polycrystalline strip with various recrystallization textures. In multiple samples manufactured by adjusting conditions of rolling and heat treatment, the Cube orientation f100gh001i, the RD-rotated Cube orientation f012gh100i, the BR orientation f362gh853i and the R(S) orientation f231gh346i, which are the representative recrystallization textures of FCC metals, developed up to 40% in respective area fractions. The bendability was clearly dependent on texture. The sample that had a strongly developed Cube orientation showed the best bendability with respect to both the good and bad ways (GW, BW) in bending. In comparison, the samples in which the BR and the R orientations developed showed poor GW bendability. The sample having a comparably random orientation showed poor GW and BW bendability. The shape of the cracks generated by bending was linear, and these cracks developed in a direction about 40 degrees from the surface. Further, they developed along shear bands, and this result was confirmed by the EBSD measurement. Therefore, the cause of cracking resulting from bending was shear bands. The correlation between the good bendability and a low average Taylor factor was confirmed. More uniform deformation by crystalline slips through texture control was effective for restraining the shear bands, i.e., for obtaining excellent bendability. [

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The changes in the states of carbon (C) together with hardness and the tensile properties of low C steel (0.045C0.34Mn in mass%) quenched from 710°C and aged at 50°C were investigated as a function of aging time using TEM and atom probe tomography. Vickers hardness increases at about 1.1 © 10 4 s, exhibits significant increase at 5.8 © 10 4 s (16 h) and maintains peak hardness untill 8.6 © 10 5 s (10 d) followed by a decrease after further aging time. At the start of peak aging, C clusters form with an irregular shape that resembles a sphere about 10 nm in diameter. The number of C atoms is about 700, and the C content is in the range of 12 at% at 1.0 © 10 5 s (28 h), where no enrichment of elements except for C is observed. At the end of peak aging, the plate-shaped precipitates (about 1 nm wide and 12 nm long) having a C content greater than 10 at% are distributed with the {100} habit plane, thus confirming the transition from C clusters to fine carbides. Lower yield strength (LYS) is the lowest for the specimen with solute C, and significantly increases for the specimen with C clusters and fine carbides in this order. LYS is determined presumably by the cutting mechanism for the C cluster specimen and the Orowan mechanism for the fine carbide specimen. The work hardening for the solute C and C cluster specimens is high, while the carbide specimen shows less work hardening. The C cluster is assumed to be decomposed into solute C through shearing by dislocations, causing work hardening and relatively good uniform elongation. Post uniform elongation (l-El) was the lowest for the C cluster specimen followed by the fine carbide specimen with the same strength level. This is because dynamic strain aging caused by solute C promotes the strain localization leading to the deterioration in l-El.

The tensile properties of an Al-Mg-Si alloy with Mg-Si clusters were compared with those of an Al-Mg-Si alloy with β ′′ precipitates of the same strength. The elongation of the alloy with Mg-Si clusters was found to be greater than that of the alloy with β ′′ precipitates because of the high work hardening rate of the former alloy, particularly in the high-strain region. Decomposition of Mg-Si clusters into solute Mg and Si atoms during the tensile deformation was revealed by differential scanning calorimetry. Transmission electron microscopy revealed three types of dislocation characteristics in these alloys: homogeneous distribution of dislocations with β ′′ precipitates, cell structures in the alloy with solute Mg and Si, and a combination of these two types in the alloy with Mg-Si clusters. In the case of the alloy with Mg-Si clusters, the yield strength increased signi cantly owing to the dislocation cutting mechanism; simultaneously, the elongation of this alloy improved greatly because of the presence of solute Mg and Si atoms formed by decomposition via plastic deformation, which were inferred to prevent dynamic recovery in the later stage of tensile deformation. Consequently, a comparison of conventional 6000 series and 7000 series Al alloys revealed that the alloy with clusters had advantages over the alloy with precipitates and the alloy with solutes in terms of the balance between strength and elongation.