Aged microstructure

The evolution of the microstructure during aging is further investigated through FESEM observations of SLM Inconel 718 samples submitted to different aging treatments at the characteristic temperatures obtained from the thermal analysis reported previously in this chapter (figures 4.5 and 4.26). The samples used for this study were previously submitted to a solution annealing treatment at 1065°C for 2 hours in order to remove the microstructural features of the as built state and obtain a small initial number of preexisting precipitates, making easier the identification of the new formed ones.

Table 2.5 summarized the thermal history of the aged samples used for this investigation. The number of the exposition steps at the characteristic temperatures and the respective durations were chosen in order to obtain some indication on the precipitation and coarsening kinetics of the second phases which are expected to form. At 565°C the formation of ’ phase is expected, the comparison between the microstructure after a relatively short exposition (4 hours) and a prolonged one (24 hours) allows to determine if a short time is sufficient to reach a saturation level of the ’ content in the aged alloy. The precipitation of ’’ is reported to be more sluggish with respect to ’ [1] [17] [211], therefore the observation after a very short period (2 hours) of exposition at 740°C is performed to verify that ’ formation precedes the ’’ phase. Conversely, the prevalence of ’’ on ’ is expected after 8 hours of exposition. The aging treatments at 800°C have the aim to verify an eventual transformation of the metastable ’’ into . Finally, at 870°C both ’’ and

 phases could form. The latter is expected to prevail on the former with the proceeding of the exposition. The overall aging period is 24 hours for all samples, the analysis after this period allows to verify an eventual coarsening of the phases

’, ’’ and  phases.

The collected Vickers microhardness are reported in figure 4.28, they provide a preliminary feedback of the evolution of the strengthening particles in the alloy during the aging treatment.

Figure 4.28. Mean Vickers microhardness obtained on the samples solutioned at 1065°C for 2 hours and then aged at the characteristic temperatures of the alloy. The bars indicate the 95%

confidence ranges. Data published in [186].

The x-ray patterns obtained on the as built sample and on the 1065°C for 2 hours solutioned one are shown in figure 4.29. The peaks relative to the face centered cubic (FCC)  matrix are clearly recognizable. No peaks related to second phases are detected. However, the present peaks are modified by the solution treatment, in fact they become narrower and shifted at a lower value of 2 angle. In figure 4.29 is shown, as an example, the magnification of the main peak, the Full Width at Half Maximum (FWHM) of the peaks relative to the as built and the solutioned samples are 0.307° and 0.0895°, respectively. The reduction of the FWHM values is probably due to the release of the internal stresses in the as built material during solution treatment. The shift of the peaks toward lower 2 angle, which corresponds to a greater value of the interplanar distance d according to the Bragg’s law (equation 2.10), is due to the increase of the lattice distortion arising from the increase of the solute content. In particular Nb is progressively released in solid solution and this is in agreement with the results obtained with the TMA analyses (figures 4.4 and 4.27). Therefore, the shift of the peaks, also reported for Inconel 625 alloy [212] [213], is an indirect proof of the dissolution of the second phases present in the as built material during solution annealing.

Figure 4.29. X-ray diffractograms of the as built sample and the sample solutioned at 1065°C for 2 hours. First published in [186].

The x-ray patterns obtained on the differently aged samples are shown in figure 4.30. An additional peak at 2𝜃 ≅ 46° is revealed only on samples treated at 870°C for 8 and 24 hours. This peak can be indexed as the (211) of the  phase, which is the stronger peak of this compound [44]. No one of the other diffraction peaks of the  phase was detected during the XRD analysis; similar spectra are also reported by Cao et al. on SLM Inconel 718 alloy after solution treatment at 1065°C for 1 hour and double aging at 760°C for 10 hours and 650°C for 8 hours [46]. The x-ray diffraction peaks related to the  phase are typically observed on Inconel 178 alloy after thermal treatment between 850°C and 990°C [49] [214] [215].

The presence of the ’ and ’’ phases cannot be ascertained by XRD analysis because their peaks are superimposed to the ones of the  matrix. Nevertheless, the precipitation of these new phases can be deduced from the shift of the peaks (a magnification of the main peak is shown in figure 4.30 as an example).

Figure 4.30. X-ray diffractograms of the samples aged at the characteristic temperatures of the alloy with magnification on the  (200) peak in order to observe the shift respect to as solutioned

state. First published in [186].

The values of lattice parameter determined on the samples evaluated by XRD analysis are shown in figure 4.31. The lattice parameter decreases with the time and the temperature of the aging treatment due to the progressive transfer of the solute from the solid solution to the precipitates. Zhang et al. also report a reduction of the lattice parameter of the FCC matrix of an Inconel 625 alloy, produced by powder bed additive manufacturing technique, during the post heat treatment caused by the precipitation of  precipitates and the consequent reduction of the Nb content from the solid solution [216].

Figure 4.31. Values of the lattice parameter of  phase with different thermal histories obtained from the XRD analyses. Data published in [186].

Most of the precipitates present after aging at 565°C are residuals of the eutectic phases, in particular small carbides that were not dissolved during the previous solution annealing, and intergranular carbides formed during the solution annealing. Also ’ is expected to form at such temperature, but it is probably too fine to be clearly defined through FESEM images, although a population of very small particles can be observed after 24 hours of aging (figure 4.32). However, the increase of the Vickers microhardness during aging at 565°C is an indirect proof of the formation of ’ strengthening phases at this temperature and it also demonstrates that the saturation level is not reached after 4 hours of exposure at this temperature, since the microhardness can still be increased by prolonging the heat treatment.

Figure 4.32. FESEM micrographs of the samples aged at 565°C. The micrographs acquired on the sample aged for 24 hours have been already published in [186].

After aging at 740°C the microstructural modification is more pronounced and an important formation of strengthening phases occurs that lead to a strong increase of the Vickers microhardness. It can be observed that new precipitates are formed already after 2 hours of treatment. The density of newly formed precipitates is not homogeneous: the interdendritic boundary are decorated with a larger number of precipitates respect to the core of each dendrite. After 2 hours the formed precipitates are extremely small, most of them are probably ’ that form with a higher kinetics with respect to ’’. Conversely, discoidal ’’ precipitates are clearly identifiable after longer aging time. Comparing the FESEM images after 8 and 24 hours of aging, it is possible to observe that coarsening of ’’ occurs that leads to a progressive reduction of the Vickers microhardness.

The ’’ precipitates have a clear crystallographic relationship with the matrix, how it can be seen in figure 4.33. In each grain the ’’ precipitates form along one of the two orthogonal directions. The ’’ disks form parallel to the {100} planes of

 with the following reciprocal crystallographic orientation [46] [47]:

(001)_𝛾′′ ∥ {100}_𝛾 , [100]_𝛾′′ ∥ 〈001〉_𝛾.

Figure 4.33. FESEM micrographs of the samples aged at 740°C. Formation of discoidal ’’ after longer aging time.

During aging at 800°C,  precipitates start to form at the grain boundaries (figure 4.34). Inside the grain, discoidal ’’ undergo further coarsening. A precipitates depleted zone forms along grain boundaries in which ’’ precipitates are dissolved due to the formation of intergranular  compounds that reduce the local availability of Nb. In this depleted zone, ’ particles are still present because their formation does not require Nb availability. Similar observations (absence of

’’ precipitates near the intergranular  plates) are also reported by Wlodek and Field [18].

After 24 hours of exposition, the intergranular  phase is coarser and precipitates with plate-like morphology are detectable also inside the grains.

Intragranular  precipitates can nucleate on the stacking fault of preexisting precipitates of ’’ [23]. Therefore at this temperature the population of ’’ is progressively substituted by  phases leading to a decrease of the Vickers microhardness.

Figure 4.34. FESEM micrographs of the samples aged at 800°C. Formation of intergranular  and further coarsening of discoidal ’’.

After the aging temperature of 870°C the strengthening phases ’ and ’’ are completely absent and a lot of plate-like  precipitates form (figure 4.35). At the grain boundaries the formation and growth of  precipitates are faster, therefore coarser plates can be observed in these zones with respect to the intragranular areas.

Inside each grain, the plates form a regular pattern consisting of arrays disposed in a parallelepiped grid. This regular disposition of the intragranular  phase is due to its crystallographic relationship with the matrix (as reported at paragraph 1.3.1.3) [19] [44] [46] [47]: (010)_𝛿 ∥ {111}_𝛾 , [100]_𝛿 ∥ 〈11̅0〉_𝛾.

The plates of each array continue to growth with time and merge leading to the formation of very long precipitates of about 6-8 m (bottom right panel of figure 4.35).