Effect of Al addition on the microstructure and mechanical properties of IN718 alloy fabricated by laser deposition manufacturing

The IN718 alloys with different Al contents were fabricated by laser deposition manufacturing (LDM) technology. The microstructure and tensile properties of LDMed IN718 alloy with different Al contents before and after homogenization-solution-double aging (HSA) treatment were analyzed. The as-deposited samples show dendrite structure which develops with the addition of Al element. The microstructure of the as-deposited samples mainly consists of the γ matrix, Laves phases and a small amount of carbides. The HSA treatment improves the precipitation of γ″, γ′ and δ strengthening phases, and with the Al content increasing, the γ′ phase increases, but the γ″ phase decreases. With the increasing of Al content, the tensile strength of as-deposited samples increases, but the elongation decreases. For HSA-treated samples, the IN718+0.5Al alloy shows the highest tensile strength and elongation. The tensile fracture surfaces of LDMed IN718 alloys show dimpled surfaces, indicating the transgranular ductile failure mode.


Introduction
Laser deposition manufacturing (LDM) technology is one of the additive manufacturing technologies [1][2][3].Compared with the traditional casting and forging processing, LDM technology has many advantages including low material consumption, low processing cycle, high flexibility and quick response etc [4,5].It can realize the demands of complex structure, lightweight and high performance of aerospace components, which has a broad application prospect in superalloy complex structure [6,7].
IN718 alloy is the key material applying for the aero engine due to its good microstructure stability, fatigue resistance, oxidation resistance and processing performance [8,9].With the rapid development of aviation science and technology, the new generation aircraft needs to meet the requirements of super high speed, high altitude and long voyage, which puts forward higher requirements for the machining technology, material structure and mechanical properties of IN718 alloy [10].IN718 alloy exhibits excellent precipitation strengthening by precipitating the γ″ (Ni 3 Nb) phase and γ′ (Ni 3 (Al,Ti)) phase.Normally, the γ″ phase is the main strengthening phase with better strengthening effect.However, γ″ phase is metastable, which tends to aggregate and grow under the action of prolonged high temperature.Moreover, the γ″ phase will be converted to a thermodynamically stable orth ogonal δ phase when the temperature exceeded the limit temperature (650 °C).The reduction of γ″ phase and the formation of excessive δ phase will cause a rapid decline in mechanical properties of IN718 alloy [11].Therefore, for the IN718 alloy applying in aero engine, it is of great significance to improve the high-temperature microstructure stability inhibiting the transformation from γ″ phase to δ phase.
Al is an important constituent element of the γ′ phase in IN718 alloy, and it can facilitate the precipitation of γ′ phase improving the high-temperature microstructure stability.Du et al [12,13] found that with the increasing of Al content in IN718 alloy, abundant of γ′ phase was precipitated forming a coated structure of γ′+γ ″ increasing the operating temperature.Xie et al [14] pointed out that the content variation of Al+Ti+Nb, Al/Ti and (A1+Ti)/Nb of IN718 alloy can change the precipitation behavior of γ′ and γ″ phases and then improve the service temperature and mechanical properties.Fu et al [15,16] found that the Al content cannot be increased indefinitely (the maximum Al content is 1.5 wt%) in IN718 alloy, otherwise it will lead to the rapid increase of δ phase and the reduction of microstructure stability of the alloy.Hence, it is feasible to control the amount and morphology of the precipitated phases and then optimize the high-temperature microstructure stability and mechanical properties by changing the Al content in IN718 alloy.However, previous works were all focused on the traditional as-cast and as-forged IN718 alloy lacking of the relevant research of the LDMed samples which exhibits completely different metallurgical mechanisms from the casting and forging processes.
In this work, the IN718 alloys with different Al content were fabricated by LDM technology, and the effect of Al addition on the microstructure and tensile properties of LDMed samples were studied.The homogenizationsolution-double aging (HSA) heat treatment processes were carried out and the microstructure evolution and properties transform of the heat-treated samples were deeply analyzed.It provides a method of improving the microstructure and mechanical properties of IN718 alloy fabricated by LDM processing by adjusting the Al content.

Materials
In this work, the carbon steel (C45E4) plates were chosen as the substrate, and was mechanical ground to remove the oxide layer, followed by being degreased in alcohol and dried in air before laser deposition test.The deposited metal powder was IN718 alloy spherical powder (the particle size is 53-150 μm) mixed with different content of pure Al powder.The chemical composition of the IN718 powder was listed in table 1.The mixed powder was fabricated by using a V-8 shaker-mixer with the rotation speed of 18 r/min for 8.5 h.Referring to the known research results [15], the mixing ratio (wt%) between the IN718 powder and the added pure Al powder was set as 100 : 0, 99.5 : 0.5, 99.15 : 0.85, and the three alloy compositions were respectively denoted as IN718, IN718 + 0.5Al, IN718 + 0.85Al for convenience.Before the laser deposition test, the three mixed powders were dried in a vacuum for 6 h at 120 °C to eliminate the moisture.

Methods
The LDM test was carried out by using the LDM-800 laser deposition manufacturing system under the protection of high purity argon.The LDMed samples with dimensions of 90 mm × 45 mm × 40 mm were fabricated in accordance with the scanning strategies of short-side unidirectional reciprocation (shown in figure 1(a)).The scanning speed was 11 mm s −1 , the laser power was 2 kW, the powder feeding rate was 8 g min −1 , and the overlapping ratio was 40%.The LDMed samples were treated with the following heat treatment process: homogenization treatment (1100 °C for 1.5 h, air cooling)-solution treatment (980 °C for 1 h, air  cooling)-double aging (720 °C for 8 h, furnace cooling to 620 °C for 8 h, air cooling).The above heat treated samples was abbreviated as HSA.
The metallographic samples were sectioned perpendicularly to the laser scanning direction and prepared by conventional inlaying, grinding, polishing and then chemically etching by using the mixture of 80 ml HCl + 7 ml HNO 3 + 13 ml HF for about 60 s.The microstructure analysis was carried out by the optical microscopy (OLYMPUS-GX51) and the scanning electron microscope (ZEISS Sigma 300).The phase identification was conducted by using the x-ray diffraction (MPDDY2094) with a Cu radiation.The step size was 0.02°, the scanning rate was 2°/min, and the scanning range was 20°−100°.
The room temperature tensile tests were performed by using a MTS universal testing machine, and the testing process were performed according to the standard of BS EN ISO 6892-1: 2019 [17].The standard dogbone shape tensile bar were taken from the XOY planes, perpendicular to the building direction (Z direction) as shown in figure 1(a).The dimension schematic of tensile samples was shown in figure 1(b).The fracture morphology of the tensile samples was observed by the scanning electron microscope.

Microstructure analysis
Figure 2 shows the optical micrograph of as-deposited IN718 samples with different Al content.As shown in figures 2(a)-(c), the microstructure of as-deposited samples is composed of non-uniform columnar dendrites and some equiaxed grains.During laser deposition process, since the heat is mainly dissipated along the building direction, the columnar dendrites grow along the direction paralleling to the building direction and passes through multiple deposited layers.The formation of equiaxed grains derives from the columnar to equiaxed transition which depends on the temperature gradient and solidification rate.During laser deposition process, the growth velocity of solideliquid interface depends on the isotherms.Once the actual temperature is below the equilibrium liquid temperature, the columnar to equiaxed transition will occur [18].For conventional IN718 sample, the dendrite morphology is basically coarse primary dendrite structure as shown in figure 2(d).With the addition of Al element, the primary dendrite structure still exists and the secondary dendrite structure develops gradually as shown in figures 2(e) and (f), which indicates a more uniform microstructure.The crystal morphology of the alloy mainly depends on the synthetic action of solute concentration, solidification rate and temperature gradient in the liquid phase [19].During the solidification process of LDMed IN718 samples, the solute concentration increases continuously with the addition of Al element and then the solute redistributes in the liquid phase, which leads to the decrease in liquidus temperature of the alloy and the increase in the supercooling degree in front of the boundary.When the super-cooling degree is higher than that required for heterogeneous nucleation in the melt, the microstructure may change from the continuous lath distribution to the continuous reticular distribution causing the developed secondary dendrite structure.
Figure 3 shows the microscopic morphology of as-deposited IN718 alloy with different Al content.The microstructure mainly consists of the γ matrix, Laves phases and a small amount of carbides [20].For Table 2 shows the element distribution of major constituent phases.It is found that Laves phase is mainly rich in Nb and Mo and lack of Cr, Fe, and Ni compared with the γ phase, which is consistent with composition characteristics of Laves phase reported in other studies [20].The segregation of Nb and Mo in IN718 alloy will affects the precipitation of strengthened phase leading to the weakening of strengthening effect.Previous study [22] pointed out that the increase in Al content for IN718 alloy can significantly decrease segregation coefficient of Nb, Mo, Ti element indicating a lower the degree of segregation of alloying elements.Also, the lower segregation coefficient means the decreasing of Nb, Mo and Ti content in liquid phase during the late-stage of solidification resulting in the morphological transformation of Laves phase.Figure 4 presents the microstructure of the HSA-treated samples.Compared with the as-deposited samples, the weld pool trace of HSA-treated samples weakens.Also, the columnar dendrite structure transforms into equiaxed recrystallized structure [23], and the microstructure is finer and more uniform with the addition of Al element due to the finer structure for as-deposited sample (shown in figure 2) [24].
IN718 alloy has the complex phase composition which gives the alloy excellent mechanical properties.The phase precipitation sequence is closely related to the heat treating temperature.After homogenization treatment (1100 °C), alloying elements are fully diffused eliminating the Laves phase and composition segregation.And, the epitaxial columnar crystals are replaced by equiaxed grains.During solution treatment (980 °C), δ phase can be precipitated along the grain boundary, which inhibits the growth of grains and insures the uniform grain.During two-stage aging treatment, a large number of strengthening phases (γ′ phase and γ″ phase) will precipitate in the matrix.And, the addition of Al element further facilitates the precipitation of γ′ phase [12][13][14].The mechanism of microstructure evolution is shown in figure 5.
The microscopic morphology of the HSA-treated IN718 alloy with different Al content is shown in figure 6.For LDMed sample after HSA treatment, short rod-like δ phase precipitates along the grain boundary (figure 6 For conventional IN718 alloy, the prime strengthening phase is γ″ phase and γ′ phase.Al and Ti elements are the constituent elements of γ′ strengthening phase, and Ti element exists by replacing the Al atom in the facecentered cubic vertex.Therefore, the content of Al element directly affects the precipitation number of γ′ phase.Normally, the addition of Al element can accelerate the precipitation rate and increase the precipitation amount  of γ′ phase.The precipitation of γ′ phase will consume partial Nb element, which leads to the decreasing of Nb content in γ matrix.And, because Nb is the constituent elements of γ″ phase, the amount of γ″ phase decreases with the Nb content decreasing, which inhibits the excessive formation of δ phase under high temperature conditions improving the thermal stability of microstructure.With the substantial increasing of Al content, the redundant Al atom excludes the Nb atom from the crystal lattice of γ′ phase and then Nb element enriches around the γ′ phase.When the concentration reaches a certain level, the γ″ phase is formed again resulting in the increasing of δ phase [16].

XRD results
Figure 7 shows the x-ray diffraction pattern of the LDMed samples.It is found in figure 7(a) that the as-deposited samples mainly consist of γ matrix phase and Laves phase.As the Al content increases, the 2θ angle gradually shifts to the right as shown in figure 7(b).It is mainly because that the addition of Al element with larger atomic radius leads to the lattice expansion decreasing the lattice spacing (d) between crystal planes.According to the Bragg's Law ( d 2 sinq l = ) [25], the incident wavelength λ remains constant under the same processing parameters, so the decrease in d value may result in the increasing of θ value.After HSA treatment (figure 7(c)), the γ phase, γ′ phase and γ″ phase precipitate, but the main peaks of γ phase, γ′ phase and γ″ phase are overlapped.The peak of Laves phase and δ phase for HSA-treated samples is not visible from the x-ray diffraction patterns due to the lower content.

Tensile property analysis
Figure 8 shows the engineering stress-strain curves of the as-deposited and HSA-treated IN718 alloys with different Al content, and the corresponding ultimate tensile strength (UTS), yield strength (YS) and elongation (EI) values are listed in table 3.For the as-deposited sample (figure 8(a)), the IN718 sample shows the lowest  24%.The improving in strength can also be credited to precipitation of strengthening phase.The precipitated γ′ and γ″ phases may impede the dislocation movement, which leads to the dislocation stacking increasing the driving force for dislocation movement [26].Finally, the dislocations will move again and cut through the precipitated γ′ and γ″ phases once the driving force is high enough [20].Therefore, the HSA-treated samples exhibit a higher strength than that of the as-deposited samples due to the precipitation of strengthening phase.Besides, the precipitated δ phases, which are incoherent with the matrix phase (γ phase), can withstand little plastic deformation leading to the initiation and propagation of crack [20].With the addition of Al element, the amount of δ phases in grain boundary for HSA-treated samples first decrease and then increased (figures 6(a)-(c)) resulting that the elongation first increases and then decreases.Figure 9 shows the fracture surfaces of as-deposited IN718 alloy with different Al contents.As shown in figure 9(a), the fracture surface of as-deposited IN718 alloy contains a fibrous area and a shear lip zone.The wavy boundary appears at the edge of the fracture, indicating a heavy plastic deformation has been experienced during the tensile process.With the increasing of Al content, the surface of the tensile fracture is gradually smooth showing the brittle characteristics as shown in figures 9(b) and (c).The fiber area on the fracture surface is enlarged as shown in figures 9(d)-(f).All tested as-deposited samples exhibit mostly dimpled surfaces, which indicate a transgranular ductile failure mode [27].The dimple size and depth of the as-deposited IN718 alloy are much larger and deeper than the Al-added samples.This indicates that the as-deposited IN718 alloy experienced a more homogeneous deformation compared with the Al-added samples [28,29], which was consistent with the higher elongation of the as-deposited IN718 alloy.
Figure 10 shows the fracture surfaces of HSA-treated IN718 alloy with different Al contents.As with the asdeposited tensile samples, the fracture surface of HSA-treated IN718 alloy also exhibits the fibrous area and the shear lip zone as shown in figures 10(a)-(c).Also, the dimpled surfaces appear on the surface of HSA-treated samples indicating the transgranular ductile failure mode as shown in figure 10(d

Conclusion
In this work, the IN718 alloys with different Al content were fabricated by LDM technology, and the HSA heat treatment was carried out.The effect of Al addition on microstructure and mechanical properties of LDMed IN718 alloy before and after heat treatment were analyzed.The obtained conclusions were as follows: (1) The microstructure of as-deposited samples is consisted of non-uniform columnar dendrites and some equiaxed grains, and the dendrite structure develops with the Al addition.After HSA treatment, the equiaxed grains increase and the microstructure is more uniform with the addition of Al element.
(2) The as-deposited samples mainly contain the γ phases and Laves phases.After HSA treatment, the γ′ and γ″ strengthening phase precipitates.With the Al content increasing, the γ′ phase increases, but the γ″ phase decreases.(3) The tensile strength of as-deposited IN718 alloy increases with the increasing of Al content, but the elongation decreases.After HSA treatment, the tensile strength and elongation first increase and then decrease.The fracture surfaces of the tensile LDMed IN718 alloys show dimpled surfaces, indicating the transgranular ductile failure mode.

Figure 1 .
Figure 1.(a) Schematic diagram of scanning strategies and sampling location; (b) dimension schematic of tensile samples.
Figure4presents the microstructure of the HSA-treated samples.Compared with the as-deposited samples, the weld pool trace of HSA-treated samples weakens.Also, the columnar dendrite structure transforms into equiaxed recrystallized structure[23], and the microstructure is finer and more uniform with the addition of Al element due to the finer structure for as-deposited sample (shown in figure2)[24].IN718 alloy has the complex phase composition which gives the alloy excellent mechanical properties.The phase precipitation sequence is closely related to the heat treating temperature.After homogenization treatment (1100 °C), alloying elements are fully diffused eliminating the Laves phase and composition segregation.And, the epitaxial columnar crystals are replaced by equiaxed grains.During solution treatment (980 °C), δ phase can be precipitated along the grain boundary, which inhibits the growth of grains and insures the uniform grain.During two-stage aging treatment, a large number of strengthening phases (γ′ phase and γ″ phase) will precipitate in the matrix.And, the addition of Al element further facilitates the precipitation of γ′ phase[12][13][14].The mechanism of microstructure evolution is shown in figure5.The microscopic morphology of the HSA-treated IN718 alloy with different Al content is shown in figure6.For LDMed sample after HSA treatment, short rod-like δ phase precipitates along the grain boundary (figure6(a)), and circular γ′ phase and strip-like γ″ phase are uniformly distributed within the grain (figure 6(d)).As shown in figure 6(b), the amount of δ phase precipitated at the grain boundary markedly decreases for IN718 +0.5Al alloy, and the amount of lath-like γ″ phase decreases significantly in the grain, but the γ′ phase increases (figure 6(e)).For HSA-treated IN718+0.85 alloy, more coarser δ phase precipitates at grain boundaries (figure 6(c)) and coarser γ′ phase are uniformly distributed within the grain (figure 6(f)).For conventional IN718 alloy, the prime strengthening phase is γ″ phase and γ′ phase.Al and Ti elements are the constituent elements of γ′ strengthening phase, and Ti element exists by replacing the Al atom in the facecentered cubic vertex.Therefore, the content of Al element directly affects the precipitation number of γ′ phase.Normally, the addition of Al element can accelerate the precipitation rate and increase the precipitation amount

Figure 5 .
Figure 5. Schematic of microstructure evolution for LDMed IN718 alloy during HSA heat treatment processes.
)-(f).By contrast, the HSAtreated IN718 + 0.5Al alloy shows the larger and deeper dimples indicating the best plasticity for the HSAtreated IN718 alloy.The fracture morphology of HSA-treated samples is consistent with the results of tensile properties shown in figure 8(b).

Table 2 .
Composition statistics of the γ phase and laves phase in the as-deposited samples.

Table 3 .
Tensile strength and elongation of the as-deposited and HSA-treated samples.