Evolution of texture and precipitates in Al–Cu–Li alloy during deformation and heat treatment and its effect on anisotropy

To analyze the influence of rolling and heat treatment on the anisotropy of the alloy, metallographic structure and texturewereinvestigated on the solid solution plates of different rolling processes, as shown in figures 2–4. The ODF diagram and α, β grain orientation line distribution of hot rolled plates after solution quenching treatment are described in both figures. The difference between A, B, and C rolling processes are not visible clearly, as shown in figure 2. Figure 3 illustratesthat the plates prepared by the A/B/C process mainly formed theBrass texture. Also, S{123}〈634〉, C{112}〈111〉 texture and a small amount of G{011}〈100〉 texture were formed in the A and B process, and a small amount of G texture and {011}〈011〉 texture were formed in C process. Figure 4 indicatesthat increasing pass deformation(A)or reversing rolling(C) changes the orientation density of the main texture components of α and β grain orientation line. Moreover, the total volume fraction of the texture (S{123}〈634〉, G{011}〈100〉, {011}〈011〉〉) was increased,which greatly reduced the anisotropy of the sheet. This is one of theleading causes for the declineof the plane anisotropy index of yield strength.

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

The alloy was hot rolled at 430 °C and immediately cooled down by water and the images were taken by TEM and arepresented in figure 5. It can be seen from figure 5 that a large amount of T₁ phase has been precipitated, and δ' and θ' phasewere not present in larger amounts. The content of brass texture in a hot rolled plate of Al–Cu–Li alloy is closely related to the T₁ phase [39, 40]. The precipitation of the T₁ phase during hot deformation may causethe formation of brass texture.

Zoom In Zoom Out Reset image size

Figures 6 and 7 are theTEM images and selected area electrondiffraction in the Al–Cu–Li alloy after solution-treated at490 °C for 1.5 h and aged at 175 °C for48 h (T6 treatment). In figure 6 all T₁ phase precipitated during hot rolling has been dissolved back to the alloy matrix, only δ' phase precipitated in the solution-treated alloy. The δ' phase and matrix have a low mismatch, about 0.08 ∼ 0.3%, and the interfacial energy between them is about 14 mJm⁻². Thus, theδ' phasewasprecipitated easily in the solution stage due to the low mismatch and interfacial energy. However, the δ' phase, which is coherent with the matrix,was sliced by dislocations and induces coplanar slip, which limits the deformation of the alloy to a few movable slip surfaces. The slip surfaces were softened and the anisotropy of the alloy plate was raised. After T6 treatment, a large number of T₁ and θ' phases precipitateas exhibited in figure 6, and thus consume a large number of Li atoms in the Al solid solution, which results in adecrease in number density of δ' precipitates. Compared with δ' precipitates, T₁ and θ' precipitates can impede the dislocation movementmore efficiently, and exhibit a stronger strengthening effect. Thus the precipitation of the T₁, θ' and δ' phase results in the decrease of the coplanar slip compared with the solid-solution-treated alloy, reducing the anisotropy index of the alloy, as shown in tables 3 and 4.

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

The results of TEM analysis of the plate prepared by a process after T8 treatment are demonstrated in figure 8. After T8 treatment, the strengthening phases were still δ', θ' and T₁ phases, but T₁ phase was the main strengthening phase. With the increase of pre-deformation, the size of the T₁ phase graduallydecreased, but the number and density were expanded, and the distribution becomesmore dispersed. The number and distribution of T₁ phase nucleation on the four {111} planes depend on the density and distribution of dislocation defects [41, 42]. According to the difference of critical shear stress, four {111} faces can be divided into two groups: 'Vertical plane' and 'Inclined plane', as shown in table 6. The two groups of {111} planes exhibit large difference in dislocation distribution due to the pre-deformation before aging, which led to the discrepancy of the T₁ phase on the four {111} planes. The texture component of the three processes was mainly brass texture. If the only brass texture is assumed in Al–Cu–Li alloy, the relationship between {111} four crystal planes and Brass texture is shown in figure 9 [43]. Two {111} planes are perpendicular to the rolling surface, and the other two {111} planes are at an angle of 35.3° with the rolling surface. This orientation relationship urged different {111} faces to obtaininconsistentshear stress during the pre-deformation. The faces with higher shear stress have higher dislocation density and provided more nucleation sites forthe T₁ phase; on the contrary, the nucleation sites were less.

Zoom In Zoom Out Reset image size

Zoom In Zoom Out Reset image size

If there is only Brass texture in the plate, the Schmid factors of the slip systems at different orientations are displayed in table 6. The {111} plane inclined to the rolling plane achievedgreater critical shear stress than the vertical plane, which causes the formation of a higher density T₁ phase in the aging process of the inclined plane after pre-deformation. The higher the pre-deformation degree, the larger the T₁ phase distribution difference on two groups of {111} planes. T₁ precipitates shown in figures 10(a), (c) (when thefoil plane is parallel to rolling plane) can be regardedas the ones perpendicular to the rolling plane (designated as 'Vertical plane' type) and those shown in figures 10(b), (d) (when the foil plane is perpendicular to rolling plane) as the ones inclined to rolling plane (designated as 'Inclined plane' type). As can be seen in figure 10 and table 7, there are large differences in the size and density of T₁ precipitates between the 'Vertical plane' and 'Inclined plane' types. The density of T₁ precipitate formed on the 'inclined plane' {111} planes is higher than that formed on the 'Vertical plane' {111} planes. It is indicated that the T1 precipitates exhibits a more uniformly distributionon the 'inclined plane' {111} planes.

Zoom In Zoom Out Reset image size

In the process of stretching along the rolling direction (0°), the deformation is mainly concentrated on the inclined plane,and the high-density T₁ phase precipitated on the inclined plane. T₁ phase can effectively hinder the dislocation movement, resulting in relatively uniform deformation on the plane, and the transverse (90°) stretching also has a similar effect. However, when the stretching occurred along the 45° direction, the shear stress on the vertical {111} plane was higher than that on the inclined plane, and the deformation was concentrated on the vertical plane. Meanwhile, the T₁ phase was the main barrier for the dislocation movement on the vertical {111} plane. But, due to the limited number of T₁ phases on the vertical {111} plane, there was a coplanar slip. This also revealed that the strengthening effect of T₁ is weaker in the 45° direction but stronger in the rolling direction (0°) and transverse direction (90°). As presented in table 5, the tensile strength was lower in the 45° direction, but higher in the 0° and 90° direction. With the rise of pre-deformation, the anisotropy index of yield strength was elevated.