4. Discussion
The EN AW 2618 Al alloy is considered as a material with low L-PBF processability and highly susceptible to solidification cracking [
2]. During solidification of melt pools, thermal shrinkage can promote accumulation of tensile stresses between adjacent grains. This phenomenon can lead to nucleation and subsequent propagation of hot cracks along grain boundaries. Coarse columnar grains are more susceptible to hot cracking than small equiaxed grains since they provide straight paths for cracks and also promote a larger concentration of solutes due to lower grain boundary surface area compared to equiaxed grains [
5]. This work confirmed that the modification of the 2618 with Ti and B leads to a fine crack-free equiaxed grain structure, which is much more refined than that typically found in 2618 alloy processed by L-PBF [
2].
The solidification curves based on Scheil equation reported in
Figure 1 demonstrate that TiB
2 and Al
3Ti precipitate from the liquid phase. These primary phases are believed to be able to stimulate heterogeneous nucleation of α-Al grains, leading to a refined microstructure, as shown by the micrograph of
Figure 2b and by the EBSD orientation map of
Figure 3a. The XRD diffractograms (
Figure 4) and the EBSD orientation map (
Figure 3b) confirm the existence of the TiB
2 phase in the as-built material.
As-built and solution treated samples showed rather similar DSC curves (
Figure 5). The exothermic peak A and the endothermic peak B are likely associated to precipitation and dissolution of GP zones, respectively, whereas the exothermic peak C and the endothermic peak D are due to the formation and dissolution of S’/S precipitates, respectively, as for the conventional 2618 alloy [
13,
14]. At temperatures higher than 470 °C, the as-built and solution annealed samples behave differently. In particular, the alloy in the as-built condition shows a large endothermic peak E. This peak is believed to be related to the dissolution of Mg
2Si (β) and Al
2Cu (θ) phases, in agreement with the results of XRD analyses (
Figure 4). Indeed, the solution treatment radically changes the microstructure, leading to the dissolution of the inter-cellular segregation network (
Figure 7), where β and θ generate at the end of solidification (
Figure 1) [
12].
SEM analysis showed that the as-built alloy is characterised by fine cells, typical of rapidly solidified materials, that are surrounded by zones with pronounced solute segregation (
Figure 2c). The direct aging performed from the as-built condition does not modify the morphology of second phases formed on solidification (
Figure 7a). On the contrary, solution treatment leads to a drastic change in microstructure as shown by the SEM image of the T6 alloy reported in
Figure 7b. The solute-rich inter-cellular network made by fine second phases is indeed replaced by coarser grains and particles. These latter identified as TiB
2, Al
3Ti, Al
7Cu
4Ni and Al
9FeNi by XRD analysis (
Figure 4c).
The high cooling rates produced by L-PBF process promote the formation of a supersaturated solid solution in the as-built material [
7]. Thus, the material can be directly aged from the as-built condition, as confirmed from data given in
Figure 5 and
Figure 6. The solution treatment causes a decrease in material hardness, likely due to microstructural coarsening and reduction of the dislocation density. Both as-built and solution treated alloys reached the maximum hardness after 1 h and 3 h at 200 °C and 180 °C, respectively. Additionally, aging curves shown in
Figure 6 highlight the faster precipitation kinetics of the novel alloy with respect to a conventional wrought EN AW 2618, which exhibits, according to literature, the hardness peak after 20 h of exposure at 200 °C [
15].
Results of tensile tests show that the T6 alloy modified with Ti and B shows values of UYS and UTS similar to those attained in wrought 2618 alloy in T6 temper. Indeed, R. Nunes et al. reported a UYS and UTS of 372 and 440 MPa, respectively, for the conventional T6-2618 alloy processed by forging [
16].
Figure 8 and
Table 5 show UYS values of 461, 495 and 392 MPa and UTS values of 447, 460 and 470 MPa for the for the as-built, T5 and T6 XY-specimens, respectively. The high solidification and cooling rates generated by L-PBF are supposed to be responsible for the formation of extended solid solution and the refined microstructure in the as-built material, resulting in high strength already after solidification and cooling. The T5 treated alloy showed higher strength values with respect to the as-built material, due to the capability of the alloy to respond to direct aging. The T6 treated alloy, which is characterised by coarser microstructure, showed higher elongation at fracture but lower UYS as compared to the T5 and as-built conditions. As shown by the results of tensile tests performed along the Z- and XY-directions, the alloy exhibits an almost-isotropic behaviour [
4]. The epitaxial and competitive growth of coarse columnar grains is indeed suppressed by heterogeneous nucleation of grains stimulated by the addition of nucleants.
Finally, it is to consider that the alloy with Ti and B printed by L-PBF shows higher mechanical strength at elevated temperatures with respect to the conventional 2618 alloy produced by forging.
Figure 10 shows the UYS vs. temperature (a) and UTS vs. temperature (b) plots of conventional wrought T6-treated 2618 [
16] and T7-treated Ti-B modified 2618 alloy of the present study.
It can be supposed that TiB2 and Al3Ti particles stabilise the microstructure and enhance the mechanical behaviour of the Ti-B modified 2618 alloy at high temperatures. Although the elongation at fracture of the novel alloy at high temperatures shows lower values with respect to that of the conventional 2618 alloy, their absolute values are considered as appreciable for structural applications, ranging between 14.3% and 22.0%.