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Although in the past it was common practice to fully anneal the rod before wire drawing, continuous processing is now usually practiced, that is the as-worked rod with retained substructure goes directly to the cold forming. Drawing of the rod through 12 dies with 20% reduction per die brings the total true strain to 6.89.
Although in the past it was common practice to fully anneal the rod before wire drawing, continuous processing is now usually practiced, that is the as-worked rod with retained substructure goes directly to the cold forming. Drawing of the rod through 12 dies with 20% reduction per die brings the total true strain to 6.89. Such additional strain of 2.56, about, 60% of the hot rolling, requires only 60% of the dislocation motion, but brings about much more dislocation storage and a higher rate of strain hardening because of comparatively little dynamic recovery at ambient temperature. The strain hardening in wire drawing of EC Al and some dilute alloys was linear when there was a retained hot worked substructure, but became zero or negative when the rod had recrystallized grains or coarse particle distributions. The softening is likely the result of dynamic recovery, not recrystallization, since no new grains were observed.
The cold worked cell structure is built on the existing hot worked structure without tearing it apart. Dislocations become entangled in the existing sub-boundaries, thus making them more ragged and reducing the links of the wall networks, and also form new walls partitioning the subgrains and decreasing the cell size to 0.5-0.8µm. This behavior is similar to that found upon reloading of a cold worked specimen after cell growth in static recovery; the flow curve is lower than the initial cold work curve of recrystallized material because dislocations are accumulating on the recovered substructure in a different way from on the purely cold worked. The hot worked structure in Al-0.65 Fe is much more stable and less disturbed by the cold working than EC wire or commercial aluminum because of the stabilizing effect of the 0.2µm FeAl3 particles.
The cold working decreases the conductivity of the wire relative to the hot rolled rod because of the greatly increased dislocation density.
Primary particles of FeAl3 greater than 0.6µm diameter, are not sharable, i.e. dislocations cannot pass through them. As they do not change shape, the surrounding matrix flows around them undergoing additional complex deformation. This results in creation of cells much smaller than the average size with dense, high misorientation walls. The needle-like, eutectic particles of FeAl3, (Fe.Co)2Al9, or Fe2SiAl3 of about 0.2µm diameter are also not shearable by individual dislocations. The metal flows relatively easily around these thin rods, so that they give rise to merely additional dislocations. However, dislocations accumulating along their length exert bending stresses that fracture them into segments only a little longer than their diameter.
The potential solution hardening by Fe is high because of the large atomic size difference of 5.9%; however, the actual hardening depends on the amount dissolved. In a 0.5% Fe alloy, the strength after quenching from 640°C to retain 0.05%Fe in solution and severely deforming is about 190 MPa compared to 170 MPa for furnace cooling for precipitation. A 0.05% Fe alloy similarly treated has strengths of 170 and 110 MPa respectively. These strengths increase slightly upon low temperature aging as solute segregates to the dislocations but decline to about 40 MPa at 310°C as precipitation and recrystallization take place. However, precipitation during hot rolling returns the concentration to equilibrium which at the finishing temperature is at such a low level that it provides very little hardening, even though the Fe atoms form weak.
In the course of rolling and wire drawing, the grains become fibrous, lengthening by a factor of 100 in the direction of the wire axis and decreasing in diameter by a factor of 10 in the plane normal to it; the rows of eutectic particles are reduced in spacing to about 2 µm. The longitudinal strength is raised because the area for glide of mobile dislocations in slip planes diagonal to the axis is much restricted.
Resistance to Softening of a Hot Worked Substructure Strengthening from cold working has very low stability at elevated temperatures because the high density, high energy substructure readily gives rise to recrystallization unless some additional factor blocks it and provides an opportunity for recovery to gradually lower the strain energy and improve the stability. On the other hand, since Al is highly capable of recovery, limiting its degree is important in maintaining strength.
In the first stable of recovery, tangles diminish in density and rearrange into neat sub-boundaries redundant dislocations annihilate with retention of the substructure scale and much of the strength. The polygonization in this stage is initially speeded up by the internal stresses in cell walls and interiors. In the second stage, the strength declines severely as subgrains become non-uniformly larger through walls either disintegrating as their dislocations leave to incorporate into others, or migrating to amalgamate with others.
The fractured FeAl3 eutectic rods of about 0.2 µm diameter fairly uniformly distributed play an important role in stabilizing the substructure. The 109mm-3 density of 0.075-0.5 µm particles is the same as the number of 1 µm cells per mm3, so that there is about one particle per cell; whereas for EC there is only one for every two cells.
In Al-10% Fe atomized-powder extrusion-compacted alloy, the dendritic FeAl3 uniformly distributed in particles of 0.3-0.03µm, stabilizes the hot worked substructure and impedes recrystallization for up to 1000 hrs at 320°C. In combination with recovery annealing, θ particles in Al-Cu alloys stabilize the substructure up to 400°C. Dilute dispersion alloys show good stability to work softening but coarse dispersions (0.7 Fe. 2-6 Ni) do not. Rapid non-uniform subgrain growth was observed in EC Al during annealing.
Large particles 0.6µm in diameter have a destabilizing influence in so far as they serve as centers of nucleation because they have created around themselves local regions of fine subgrains with very high misorientations.
Such large particles 0.6-2.5µm widely spaced have been observed to accelerate recrystllization in Al-Fe alloys, as well as in several other alloy systems.
Al-Mn and Al-Mg-Si alloys with large particles of 1µm, from chill casting and in the former stabilization arose from 0.04µm precipitates of MnAl6 not from Mn in solution. This behavior as confirmed for Al-Mn alloys with additional fine precipitates of either ZrAl3 or MnAl6. Similar inhibition was found in a commercial RR58 and special alloys here the large particles were Fe-Ni intermetallics and the fine, 0.2µm spheres of MgCuSi.