Annealing behavior of AlSi10Mg alloy fabricated by laser powder bed fusion

The AlSi10Mg alloy fabricated by laser powder bed fusion was subjected to isochronal annealing for 2h at temperatures in the range of 100–520 °C. Recovery and grain growth of the annealed samples were investigated by combining EBSD and in situ TEM techniques. A bimodal grain size distribution with grain sizes of 0.9 μm and 2.3 μm was developed in the initial sample, where large grains were found in the interiors of melt pools and most of small grains were distributed along melt pool boundaries. In-situ TEM investigations revealed that nucleation that occurs at the melt pool boundaries is as a consequence of migration of high angle boundaries during annealing. A slight increase in grain sizes was discovered during high-temperature heat treatment, indicating the remarkable thermal stability of the alloy compared to conventional cast Al-Si alloys. This is explained by the fact that the stable eutectic Si structure, which was originally formed during rapid solidification, can effectively retard grain boundary migration during recrystallization.


Introduction
Aluminum alloys are particularly essential lightweight materials due to their high specific strength, low density, superior electrical and thermal conductivity, and corrosion resistance [1].In recent years, due to the rapid development of automotive and aerospace fields, aluminum alloy structural parts have been widely used in large quantities in various industries.Casting and rolling are the classic forming procedures for aluminum alloys, and their products cannot fulfill the increasingly high-performance standards [2].One of the most common additive manufacturing methods [3] is laser powder bed fusion (L-PBF), which has the benefits of high product precision, high molding speed, and great product design flexibility [4].As a result, L-PBF is now one of the most promising new technologies for producing aluminum alloys.
However, due to the low absorption and high reflectivity of aluminum alloys to laser light, the aluminum alloy components prepared by L-PBF often have defects such as residual stresses and spherification phenomena [5].AlSi10Mg has the benefits of good flowability, excellent air tightness, and less hot cracking during the manufacturing process when compared to other kinds of aluminum alloys [6].At present, the research on L-PBF developing of AlSi10Mg is mainly focused on the characterization of the alloy microstructure, enhancement of the mechanical properties of the components via modifying the processing parameters, and the impact of heat treatment on the alloy structure and properties, etc. Buchbinder et al. [7] succeeded in producing parts with densities close to 100% and found that parts with relative densities close to 100% could only be produced at laser powers higher than 150 W. It was also discovered that increasing the laser power and decreasing the scanning speed will raise the equipment parameters and production costs.The high laser power also tends to cause spherification of the melt pool, which leads to uneven powder distribution.Rosenthal et al. [8] found that the AlSi10Mg alloy components prepared by the L-PBF technique consisted of the coarse cell zone (CG), the fine cell zone (FG), and the heat-affected zone (HAZ).Erhard Brandl and Damien Buchbinder et al. [9] have jointly found that heat treatment has the greatest influence on fatigue properties and print orientation has the least influence on fatigue properties.The differences in the fatigue properties between the different orientations (0°, 45°, 90°) disappeared after the T6 treatment.It has also been observed that abrupt cracks always originate from the surface or subsurface (pores, unmelted areas).H.J. Klauss et al. [10] found that the heat treatment process can modulate the strength and plasticity of L-PBF AlSi12 specimens over a wide range.The L-PBF specimens were isothermally annealed at 473-723 K.The results showed that the higher the annealing temperature, the coarser the structure.X.P. Li et al. [4] investigated the different effects of subsequent solution heat treatment on AlSi12 alloys produced by L-PBF and conventional AlSi12 alloys and found that solution heat treatment can obtain eutectic AlSi12 alloys with controlled ultra-fine structure and excellent mechanical properties, which can also be regulated by changing the solution heat treatment time to control the precipitation and aggregation of Si particles.Wei Li et al. [11] sudied the effects of solid solution and artificial aging heat treatments on the AlSi10Mg alloy prepared by L-PBF and found that the solid solubility of the Si element in the alloy was much higher than the equilibrium solid solubility.As the temperature increases, the solid solubility decreases rapidly and the Si atoms precipitate into fine Si particles.This gave rise to the idea of using various heat treatment processes to control the microstructure and mechanical properties of alloys.Shakil et al. found that the microstructure of the AlSi10Mg alloy consisted of the Al matrix, the eutectic Si, and the diffusely distributed Si particles in the Al matrix.As the annealing temperature increased, the eutectic Si fractured, coarsened, and became spherical, while the diffusely distributed Si particles gradually decreased and finally disappeared, which is a typical process of Ostwald ripening [12].
It can be seen that the heat treatment of the L-PBF AlSi10Mg alloy is mainly based on solid solution treatment and artificial aging treatment, and the microstructure and strength of the alloy after heat treatment are significantly different from those in the as-built and as-cast states.Nevertheless, there has been limited research conducted on the investigation of the evolution of the microstructure of L-PBF aluminum alloys during annealing and its effect on the thermal stability properties.Based on this, this work provides a theoretical basis for further optimization of the heat treatment process of the material by observing the evolution of the substructure at different annealing temperatures and studying the annealing behavior of the substructure in situ TEM based on the microstructure of L-PBF AlSi10Mg to understand the mechanism of alloy recovery and nucleation.

Experimental
Table 1 shows the chemical composition of the raw AlSi10Mg alloy powder measured by inductively coupled plasma emission spectrometry (provided by SLM solutions Group AG) and the relative density measured by Archimedes method with a theoretical density of 2.67 g/cm 3 for the as-built AlSi10Mg alloy.At least three samples were tested to ensure the reproducibility of the results.The powder was prepared by aerosolization and the morphology is shown in figure 1a.The powder has good sphericity with agglomeration and its average particle size is 32 µm.
The production equipment is L-PBF Solution SLM125HL with the following parameters: maximum laser power 400 W, scanning speed 1.65 m/s, scan line spacing 130 μm, powder thickness 30 μm per layer, interlayer rotation angle 67°, zigzag scanning mode.Figure 1b shows the physical image of the as-built sample.The samples were heat treated in the LT15/11/B410 Nabertherm muffle furnace.When the temperature reached the set temperature, the specimens were placed inside and cooled in the air after a certain time shown in table 2.
Specimens were cut to an appropriate size, polished, and then etched with Keller's reagent for 15 s-30 s.Microstructures were observed with a JEOL JSM-7800F field emission scanning electron microscope equipped with a NordlysMax2 EBSD detector, and the composition was analyzed by scanning electron microscope with energy dispersive X-ray spectroscopy (EDX/EDS).The HKL CHANNEL5 system and analysis software were used to calculate information such as grain size.A JEOL-2100 transmission electron microscope equipped with an EM-31670SHTH double-tilt heating rod was used for in-situ heat treatment tests.This is due to solute element redistribution at the solid-liquid interface front during fast solidification (solidification rates of up to 10 6 ~10 7 K/s), resulting in some microscopic segregation between dendrite and dendrite branches, with silicon concentrated in the dendrite.And solute atoms in the coarse cell zone are subjected to more severe thermal cycling [13].When focusing on the fine grain region, it is possible to see that the fine cellular α-Al eutectic structure is distributed in a continuous network and that no secondary dendrite arms have fully developed.However, there are a small number of dispersed particles inside the eutectic cellular structure as shown in figure 2b.The size of the coarse cell zone grows dramatically, with a diameter of around 2-3 μm, whereas the diameter of cellular α-Al is approximately 0.4-1 μm in the fine cell zone.The continuous Al-Si eutectic structure also breaks in the tiny heat-affected zone that exists between the fine and coarse cell zones, resulting in discontinuous staple-like fibers.
After being annealed at 200 C, 300 C, and 520 C for 2 h, the microstructure morphology of the fine cell zone of the L-PBF AlSi10Mg alloy is shown in figure 2c-e.The sample's microstructure changes after annealing at various temperatures.The eutectic structure was also dispersed in a cell-like style after annealing at 200 °C, and the microstructure mostly retained its original look.However, some of the cells had started to fuse and produce slight spheroidization.At the same time, the number of nano-Si particles in the matrix increased, and the size was about 20-30 nm.This is due to the precipitation of the supersaturated Si elements inside the matrix by annealing.
When the annealing temperature was increased to 300 °C, both the precipitated Si and the eutectic spheroidized Si were in the form of fine particles, and the eutectic network structure was nearly destroyed, as shown in figure 2d.Further increase the annealing temperature to 520 °C, this temperature reaches the solid solution temperature of aluminum alloy, the microstructural morphology is shown in figure 2e.The eutectic network structure has completely disappeared, the Si particles grow rapidly and the particle size is large, and some of the coarsened Si grows into regular polyhedra with a more uniform shape and independent distribution.The surface of coarsened Si particles has a growth step and a size of 1-3 μm, which is about 50 times greater compared with that after annealing at 200 ℃.This indicates that with the increase of annealing temperature, the eutectic Si particles keep coarsening and the network-like structure disappears, while the Si particles precipitated in the original matrix also coarsen and grow up, which is indistinguishable from the coarsened eutectic structure.

EBSD characterization of AlSi10Mg as-built and annealed specimens
The grain size distribution of the as-built samples together with the EBSD-IPF orientation imaging of the AlSi10Mg samples both before and after annealing is shown in figure 3. Heat treatment hardly changed the topography of the EBSD maps of L-PBF AlSi10Mg alloy until 300 °C.Therefore, only EBSD-IPF orientation maps after annealing at 200 °C are shown here.It can be seen that for L-PBF AlSi10Mg alloys, heat treatment at low and medium temperatures (200 °C -300 °C) has little effect on the EBSD-IPF orientation maps, although the SEM images at these temperatures are different.By increasing the annealing temperature to 520 °C, the difference in EBSD-IPF orientation maps caused by the presence of the melt pool boundary is reduced, and abnormally large grains in the size of 20-50 μm appear in the sample.
Further statistical results of the average grain size, grain boundary orientation misorientations, and recrystallization ratio of the samples before and after annealing are plotted in figure 4. The size and fraction of the recrystallized grains in different states are determined by EBSD post-processing software.It is found that when the heat treatment temperature increases, the alloy's average grain size, grain boundary orientation misorientations, and recrystallization ratio are all continuing to increase.In particular, the grain size and HAGB only rose by around 22% and 14.2%, respectively, when the samples were heated at 520 °C for 2 h.The HAGBs curves and recrystallization ratio evolution patterns from the samples revealed that recrystallization behaviour occurred when the samples were heat-treated at 300 °C.However, the grain size of L-PBF AlSi10Mg alloys increased gradually and its thermal stability was satisfactory compared with traditional Al-Si alloys.

In-situ observation of the behaviour of the cellular structure during annealing
To further investigate the evolution of the precipitated phases and interfaces during annealing, in-situ heating experiments (evolution of the cellular structure during the heating phase from 280 °C to 450 °C) were conducted for the L-PBF AlSi10Mg alloy, and its microstructure at different times is shown in figure 5. From the observations in figure 5a-d, it can be seen that there is no obvious sign of the growth of Si particles at labeled point 1 at the early stage of heating (from 0 s to 3 min); the change of Si particles at labeled point 2 is not obvious, and the interface at labeled cell boundary 3 remains stable.At 4 min to 5 min, the Si particles at point 1 became smaller and the Si particles grew back; at point 2, the Si particles grew coarser; at boundary 3, the Si particles at the interface also grew back and the interface boundary became blurred; and at labeled grain 4, the recrystallized grains appeared.When the temperature was continued to 450 °C (5 min to 8 min), the recrystallized grains grew rapidly within a short time; the Si particles at the interface at cell boundary 3 basically disappeared, leaving only three small particles with less obvious lining and the interface also disappeared; the size of recrystallized grains at grain 4 grew to about 0.75 μm; the size of particle 5 grew to about 0.5 μm.
To better observe the changes in microstructure, the evolution of several sites focusing on the cellular structure in figure 5 is converted into a sketch, as shown in figure 6 (the other parts are omitted in the simplified figure).Overall, two phenomena appear in the Si particles during the annealing process.One is the smaller size particles that slowly disappear during the annealing process, such as the Si particles at labeled point 1; the other type of larger size and more stable structure of Si particles maintain a stable structure for 1-2 minutes during the annealing process and then grow rapidly.The comparison of labeled point 5 with labeled points 1 and 3 demonstrates that the Si particles at the interface of three cellular structures are more stable than the Si particles on the boundary of two adjacent cellular structures.Furthermore, it is noteworthy that, after about 4 minutes of warming, the small-sized Si particles at the point 3 interface vanished along with the melting back.According to this, the Si particles at the interface are directly connected to the breakdown of the cellular structure, and when the Si at the interface turns back-soluble, the pinned particles at the subgrain boundary vanish and the subgrain border migrates quickly.At the same time, the recrystallized grains at label point 4 nucleate at the cellular structure boundary after 4 min annealing, followed by rapid growth.In general, the morphological evolution of eutectic structures during annealing can be analyzed from two perspectives: Precipitates resolution or growth; Interface migration.Firstly, there are two variations of the precipitated Si particles, one is the nanoscale Si particles precipitated from the matrix, and the other is the Si particles obtained by spheroidization of the eutectic structures.When annealed at medium and low temperatures, the supersaturated Si in the solid solution in the matrix began to precipitate, and the eutectic Si began to spheroidize.This is because the interfacial energy γAlSi of Al-Si is large, spheroidization and granularization occur to reduce the interfacial energy, and the continuously growing eutectic structure begins to dissolve and break, and the network structure begins to be destroyed [14].By continuing to increase the annealing temperature, Si particles with sizes smaller than the average size gradually dissolve and disappear, and the phenomenon of "large particles engulfing small particles" occurs, resulting in a decreasing number of Si particles within the field of view and increasing average size.At higher annealing temperatures, abnormal growth of Si particles also occurs, i.e. some Si particles are much larger than others.Additionally, grain size increases and migration occur at the interface as a result of decreased stability brought on by the Si particles' disappearance from the cellular structure.The Si particles on the trigonal points are more stable and they increase the resistance to grain growth, so the alloy still maintains good thermal stability at high temperatures and the grain size grows slowly.In contrast, eutectic Si in the middle of the cellular boundary is found to be more prone to back-solvation.In addition, since the structure of the alloy prepared by the L-PBF technology is in a sub-stable state, recrystallization occurs on the boundary when annealed at high temperature again, and the recrystallized grains grow rapidly so that abnormally grown grains appear.

Figure 2 .
Figure 2. SEM images of the AlSi10Mg alloy with and without heat treatment: (a-b) without heat treatment; (c) at 200 C for 2h; (d) at 300 C for 2h; (e) at 520 C for 2h and their EDS results are shown in (f).

Figure 3 .
Figure 3. EBSD characterization of the AlSi10Mg alloy with and without heat treatment: (a) As-built; (b) Grain size distribution diagram; (c) At 200 C for 2h; (d) At 520 C for 2h.

Figure 4 .
Figure 4.The calculation results of microstructural features.

Figure 5 .
Figure 5.In situ TEM observations of Si particles at small-angle cell boundaries in AlSi10Mg samples (280 °C H) during heat treatment at temperatures changing from 280 °C to 450 °C.

Figure 6 .
Figure 6.Schematic diagram of in situ TEM observations of Si particles at small-angle cell boundaries in AlSi10Mg samples (280 °C H) during heat treatment at temperatures changing from 280 °C to 450 °C.

Table 2 .
Experimental parameters of isothermal heat treatment of AlSi10Mg samples.