The most interesting observation in the as-received T351 materials was the change in volume fraction of recrystallisation as a function of the dispersoid content. This provided the initial evidence of a possible interaction between Zr and Mn when these two elements are added together. There was a rather strange phenomenon regarding the GB pinning behaviour, since the combined Zr-Mn addition was effective for reducing grain size, but not for inhibiting recrystallisation. Fig. 4.1 clearly shows that the recrystallised volume fraction in the 0.1Zr-0.3Mn alloy at mid-thickness was fairly high (13%), while in the 0.1Zr alloy it was minor (1.5%). The value for the 0.1Zr-D alloy shows that even a coarser Al3Zr distribution could pose a more successful barrier to recrystallisation (Vf=5%). A Zr content of 0.05 wt% was not sufficient to prevent recrystallisation and a large volume of the material ended up with a recrystallised microstructure. At the other extreme of adding only Mn, the microstructure became completely recrystallised in the presence of 0.4 wt% Mn. In general, recrystallisation was more pronounced near the sheet surface.
0 20 40 60 80 100
surface mid-thickness
Recrystallised V f (%)
0.1Zr 0.05Zr-0.4Mn 0.3Mn
0.1Zr-0.3Mn 0.1Zr-D
Fig. 4.1: Recrystallised volume fraction at the surface and mid-thickness plane of the sheet for all alloys in the T351 condition.
The microstructures in this condition depend strongly on the presence of dispersoids which provide the means to control grain growth during annealing treatments, as well as the extent of recrystallisation [14]. The pinning ability of dispersoids is determined by their size, volume fraction, distribution and coherency [14]. Hence, it was expected that the finer and coherent Al3Zr phase would be a more potent recrystallisation inhibitor compared to the much larger and incoherent Al20Cu2Mn3 phase [16]. Indeed, from Fig. 4.2 it is clear that the addition of 0.1 wt% Zr, with or without the presence of Mn, ensured that the fibrous grain structure resulting from hot rolling was maintained even after solution treatment, with insignificant levels of recrystallisation. Even a coarser Al3Zr distribution in the 0.1Zr-D alloy, produced by a more extended homogenisation treatment with a second stage at a higher temperature (§3.1), had the same effect. Lowering the Zr content to 0.05 wt% led to a significant decrease in the GB pinning ability in the 0.05Zr-0.3Mn alloy, which was 62.5% recrystallised. In the case when no Zr was added at all, the sole presence of Mn was not sufficient to restrict GB motion during solution treatment so the outcome was a coarse and fully recrystallised grain structure, as shown in Figs. 4.2i,j for the 0.4Mn alloy. In all cases the grains remained highly elongated along RD due to the high strain imposed during rolling. For the slightly recrystallised alloys, the effects of the various dispersoid contents on the microstructures and the extent of recrystallisation are much clearer in the EBSD maps near the surface of the sheet than at mid-thickness.
(a) (b)
Fig. 4.2: EBSD maps showing the grain structure morphology at the surface (left) and mid-thickness (right) planes of all alloys in the T351 condition (orientations are represented by Euler colours): a,b) 0.1Zr-D, c,d) 0.1Zr, e,f) 0.1Zr-0.3Mn, g,h) 0.05Zr-0.3Mn, and i,j) 0.4Mn.
Regarding the substructure present in the unrecrystallised regions, the 0.1Zr and 0.1Zr-0.3Mn alloys exhibited typical hot-rolled microstructures that consisted of well-defined subgrains (Fig. 4.3a,b) with a low remnant density of free dislocations in their interiors (Fig.
(e)
(c) (d)
(f)
(h) (g)
(i) (j)
4.3c). Most of the dislocations had formed subgrain boundaries. Subgrains in the 0.1Zr alloy were longer and more elongated along the RD, compared to the 0.1Zr-0.3Mn alloy. The subgrain aspect ratio values at the mid-thickness plane of the sheet were measured to be 5.1 and 2.4 respectively for the two alloys. This difference was caused by a difference in the pinning pressures depending on the dispersoid family present and their respective distribution, which defines how they interact with the recovered substructure. More details will be given on this topic in Chapter 5 and once the dispersoid volume fractions are discussed in §4.4.
Fig. 4.3: STEM-HAADF images showing subgrain structures: a) well-defined subgrains in the 0.1Zr alloy, highly elongated along RD, b) well-defined subgrains in the 0.1Zr-0.3Mn alloy with smaller aspect ratio, and c) two subgrains in a two-beam condition in the 0.1Zr-0.3Mn alloy, exhibiting strong diffraction contrast from free dislocations in their interior.
Quantified data from EBSD measurements aided in the comparison of all grain structures in the T351 temper. Only maps taken at the sheet’s mid-thickness were considered for these results. The HAGB spacings in all alloys clearly illustrated the previously observed trends. The average HAGB spacing and LAGB spacing along ND appeared to change in a similar fashion as a function of the dispersoid content (Fig. 4.3a). The 0.1Zr-0.3Mn alloy had the finest grain spacing of all, equal to 5.8 µm, followed by the 0.1Zr alloy with an average HAGB spacing of 7.4 µm. This difference becomes more notable if one considers that the recrystallised volume fraction in the former alloy was higher than in the latter (Fig. 4.4). Thus, even with a significant number of coarse recrystallised grains being included in the calculations, the final HAGB spacing value was still smaller. The coarser Al3Zr distribution in the 0.1Zr-D alloy was not as effective as the previous two alloys, so the resulting HAGB spacing was slightly larger (8.2 µm). A much higher increase was observed when the amount of Zr was lowered to 0.05 wt% (18.4 µm) and when no Zr was added, the HAGB spacing increased even further, as was the case for the 0.4Mn alloy (50.6 µm). It has to be noted here that the LAGB spacing for the 0.4Mn alloy does not appear in the graph in Fig. 4.3a because there was no defined substructure present in this alloy.
(a) (b) (c)
1 10 100
HAGB spacing LAGB spacing
GB spacing (µm)
0.1Zr-D 0.05Zr-0.4Mn
0.3Mn
0.1Zr-0.3Mn 0.1Zr
Fig. 4.4: HAGB and LAGB spacing along ND measured from EBSD maps in the T351 condition for all five alloys.
The apparent contradiction that the Mn-dispersoids are able to restrict grain growth, but also less effective in limiting recrystallisation has been reported before, as mentioned in §2.2.4 [16], but no explanations were given. At this stage, no safe explanation can be given for this observation. Details will be given later in this thesis after the dispersoid distributions have been thoroughly examined.