Influence of Post-Treatment Process on Microstructure and Properties of Laser Additively Manufactured Nickel-Based Superalloy

Abstract

Defects such as cracks and micropores exist in nickel-based superalloy during laser powder bed fusion (LPBF), hindering their application in various fields. Hot isostatic pressing (HIP) was combined with conventional heat treatment (HT) to obtain LPBF nickel-based superalloy parts with ideal properties and fewer defects. The results show that HIP process can improve the densification, while the conventional HT can eliminate the micro-defects to improve the mechanical properties. After HIP treatment, the defect volume fraction of LPBF specimens decreases. After HT, the defect content of HIP+HT specimens increases slightly. After post-treatment, the hardness shows a decreasing trend, and the tensile strength and post-break elongation of HIP+HT specimens increase to 1326 MPa and 21.3%, respectively, at room temperature.

Keywords

nickel-based superalloy; laser powder bed fusion; post-treatment; microstructure; properties

The rapid development of additive manufacturing technology has solved problems in rapid manufacturing of complex parts made of nickel-based superalloys^[1–3]. In particular, the application of laser powder bed fusion (LPBF) technique realizes the direct manufacturing of nickel-based superalloy parts, which greatly improves the manufacturing efficiency of important parts in the fields of aerospace, petrochemical, medical and automotive manufacturing^[4–7].

LPBF is a metal 3D printing technique based on high-energy laser beams^[7–9]. The principle is to selectively melt and bond loose powders by heating the powder bed with concentrated energy scanning from a laser heat source^[10–11]. This enables the molding of complex shaped parts^[12–14]. Compared with traditional manufacturing methods, it overcomes the problems of long molding cycle time, poor molding quality, cumbersome post-processing, etc, and has the advantages of high molding density, good surface molding quality, high precision, etc. With high molding density and dimensional accuracy up to 20‒50 μm, LPBF technique has a very high efficiency in the manufacture of metal parts^[15–16].

LPBF produces defects (pores and cracks) during the manufacturing process, and these defects degrade the mechanical properties of the superalloy. In addition, due to the extremely high solidification rate of the alloy during the manufacturing process, high residual stresses are generated in the material. In order to eliminate the alloy defects and to improve its properties, different heat treatments (HTs) were used. Xuan et al^[17] reported that single crystal superalloy was subjected to hot isostatic pressing (HIP) treatment, which resulted in an increase in mid-temperature plasticity. Microcracks disappear after HIP treatment, which increases the densification of the alloy. Baskov et al^[18] investigated the effect of heat treatments such as HIP and HIP+solid solution+aging on EO741NP alloy. The post-treatment results in precipitation of γʹ phase, dissolution of Laves phase and formation of MC carbide phase in the alloy. The mechanical properties of the alloy after HIP+solid solution+aging treatment were optimal. Roncery et al^[19] carried out a suitable hot isostatic complete HT process on ERBO/1 alloy, which not only significantly reduced the number density of micropores, but also refined the microstructure of the γ/γʹ phase in alloy, leading to an increase in the enduring creep life of the alloy at 750 °C/800 MPa. Sun et al^[20] used a two-step HT method of HIP and solid solution to reduce micro-defects from 0.96% to 0.11%, which increased the tensile strength and elongation at break of the alloys at room temperature by 3.6% and 1135%, respectively, and increased the tensile strength and plasticity at 900 °C by 11.9% and 410%, respectively.

In this study, laser additive manufacturing of ZGH451 was researched and the effect of post-treatment (HIP, solid solution+aging and HIP+solid solution+aging) on the microstructure and mechanical properties of the alloy was analyzed. The alloys before and after post-treatment were characterized and their mechanical properties (hardness and tensile property) were compared.

1 Experiment

A laser additively manufactured nickel-based superalloy material, ZGH451, independently developed by the Institute of Metal Research, Chinese Academy of Sciences, was selected as the test material. Table 1 shows the chemical composition of the alloy. The LPBF experiments were carried out using a Concept Laser M2, equipped with a 400 W fiber laser on the SLM machine. In the multilayer addition stage, a 316L stainless steel plate was preheated and kept at a constant temperature of 100 °C. The working chamber was filled with argon gas, so the oxygen content was below 10 μL/L. Specific process parameters are shown in Table 2. The size of the resulting sample was 20 mm×20 mm×80 mm.

Table 1  Chemical composition of ZGH451 superalloy (wt%)

Ta+Mo+W	Al+Ti	Cr	C	Co	B	Ni
15.8	6.0	8.0	0.05	8.0	0.015	Bal.

Table 2  LPBF process parameters

Laser power/W	Speed /mm·s-1	Spot size/μm	Focus movement/mm	Offset to original contour/mm
240	680	180	0	0.125

The specimens were investigated in four states i.e. LPBF, HIP, HIP+HT (solid solution+aging) and HT. HIP, HIP+HT and HT were firstly treated by stress relief annealing (650 °C/4 h, referred to as HT₀), followed by subsequent treatments. HIP parameters are 1200 °C/180 MPa/2 h; solid solution para-meters are 1240 °C/3 h; aging parameters are: 1120 °C/4 h+850 °C/24 h.

The specimens were ground and polished by LAP-2MV metallographic grinding and polishing machine, respectively. Electrolytic corrosion was carried out on the finely polished samples, the corrosion solution was 15 g CrO₃+10 mL H₂SO₄+150 mL H₃PO₄, and the corrosion time was 15 s. The microstructures were observed by DM4M metallurgical microscope (OM) and SU8010 field emission scanning electron microscope (SEM), and the volume fractions of the micropores were calculated by the image analysis software Image-pro Plus. Electron backscattered diffraction (EBSD) analysis was also conducted. The residual stress of annealed specimens was tested with the LM-12 residual stress tester. Nanoindentation experiments were performed on the specimens by a Nano Indenter* G200 device at room temperature and the indenter was pressed down under 30 mN for 15 s. For the three-dimensional information (number, volume fraction and size) of the micropores inside the alloy, the specimens were quantitatively characterized by X-ray computed tomography (XCT, Versa XRM-500 device). The specimens used for XCT observation were cylindrical specimens with 2 mm in diameter and approximately 5 mm in height. The tensile properties of the specimens at the room temperature were tested by a TSE504D universal testing machine, and the tensile specimen is shown in Fig 1.

Fig.1  Schematic of tensile specimen

2 Results and Discussion

In order to prevent the excessive residual stress in the alloy due to deformation and cracking during post-processing, the laser additive manufacturing of nickel-based superalloy ZGH451 was firstly de-stressed and annealed, and the residual stress was reduced from 840 MPa to 744.6 MPa.

Fig.2a‒2b show the microstructures of LPBF specimens in the perpendicular and printing directions. In the LPBF specimens, overlapping “fish-scale” melt pools with a depth of 70 μm and a width of 110 μm can be observed. Microcracks of different sizes are randomly distributed at grain boundaries, which is similar to the microstructure of Hastelloy-X alloy prepared by LPBF^[21]. At large magnification, columnar dendritic structures with different orientations can be observed, and this type of structure is also typical of other nickel-based superalloys prepared by LPBF^[22–24].

Fig.2  SEM images of nickel-based superalloy ZGH451 after different treatments: (a‒b) LPBF, (c‒d) HIP, (e‒f) HIP+HT, and (g‒h) HT

Fig.2c‒2h show the microstructures of HIP, HIP+HT and HT specimens perpendicular to the printing direction. After the alloy is treated by the post-treatment process, the “fish-scale” melt pool boundary and the columnar dendritic structure with different orientations disappear, and the Ni₃Al-based γʹ phase is precipitated in the nickel-based solid solution. After the HIP process (Fig.2d), the γʹ phase has different sizes and irregular shapes, and a small part of the γʹ phase is cubic. After HIP+HT and HT (Fig.2e‒2h), the γʹ phases are uniformly precipitated and cubic in shape, and it can be seen clearly that in the specimens treated by HIP+HT (Fig.2e‒2h), the size of γʹ phases is significantly reduced, with an average size of 0.533 μm and a high volume share of 60.9%.

Fig.3 shows the EBSD results of LPBF specimens, HT₀ specimens, HIP specimens, HIP+HT specimens and HT specimens parallel to the printing direction. The columnar grain size of different specimens ranges from 5 μm to 400 μm, and the average grain size of LPBF specimen is 23.5 μm. The average grain sizes of HT₀ and HIP specimens are 25.7 and 27 μm, respectively, and the grain size, morphology and orientation are almost the same as those of LPBF specimen, indicating that there is no phenomenon such as recrystallization and grain growth during annealing and HIP. The average grain size of HIP+HT specimen is 60.4 μm, and grain morphology changes from columnar crystal to equiaxial crystal. The average grain size of HT sample is 71.3 μm. The grains of HIP+HT and HT specimens are mainly dominated by large grains, and the annealing twins can be observed, which indicates that the grains grow up and become coarse in the HT process.

Fig.3  EBSD images of nickel-based superalloy ZGH451 after different treatments: (a) LPBF, (b) HT₀, (c) HIP, (d) HIP+HT, and (e) HT

As shown in Fig.4 and Fig.5, the microscopic defect contents of LPBF, HIP, HIP+HT and HT specimens are measured by OM and μ-CT analysis. Statistically, the OM and μ-CT densities in the samples are greater than 99.6%. It should be noted that due to the limitation of the specimen size and the number of specimens used in the experiments, OM observation is a two-dimensional detection and μ-CT is a three-dimensional detection, so the statistical results will be somewhat different. As can be observed from Fig.4, HT cannot improve the densification of the specimens, so the densities of the HT specimens and the LPBF specimens are close to each other. It can be clearly observed in Fig.4c‒4f and Fig.5d‒5i that the HIP process eliminates most of the defects and increases the specimen densities. From the OM and μ-CT plots, it can be observed that most of the pore morphology is spherical and near-spherical, which is consistent with the gas pores produced by the original powder particles^[25–27]. After HIP+HT (Fig.4e‒4f), it can be found that the microcracks do not reappear, which suggests that there is no gas residue inside the microcracks. However, there is a slight increase in the micro-defects, which is attributed to the fact that the residual stress in the pores provides the impetus for the growth or regeneration of micropores during solid solution+aging, and an equivalent phenomenon was also found by Sun et al^[20].

Fig.4  Morphologies and distribution of micropores in ZGH451 specimens after various treatments: (a‒b) LPBF, (c‒d) HIP, (e‒f) HIP+HT, and (g‒h) HT

Fig.5  OM images of as-fabricated ZGH451 specimens obtained after various treatments: (a‒c) LPBF, (d‒f) HIP, (g‒i) HIP+HT, and (j‒l) HT

Fig.6a‒6d show the nanoindentation load-displacement curves of the as-deposited specimens and the specimens with different HTs, under a peak load of 30 mN with six indentations (No.1‒No.6) for each specimen. As can be seen from Fig.6a‒6d, all the load-displacement curves are smooth, indicating that no “sawtooth rheological behavior” occurs, which is inconsistent with the phenomenon of heat-treated Hastelloy X and GH3536 alloys prepared by laser powder-bed melting^[28–29]. Hastelloy X and GH3536 alloys undergo a significant “sawtooth rheological behavior” after HT. In general, the phenomenon of “sawtooth rheological behavior” (sudden displacement shifts) in the load-displacement curve of nanoindentation indicates that the mechanical behavior of the alloys changes from fully elastic deformation to plastic deformation. Each displacement shift corresponds to one or more shear zones which coordinate the plastic deformation of the material, and the larger the displacement shift, the more severe the shear plastic deformation inside the shear zone. As can be seen from Fig.6, during the loading process, the stress-strain curve is smooth and upward, and there is no obvious displacement mutation or “sawtooth rheological behavior”, indicating that no matter whether the deposited specimen is post-treated or not, the specimen is deformed by elastic deformation. This indicates that both the deposited and post-treated specimens mainly suffer elastic deformation.

Fig.6  Load-displacement curves (a‒d), hardness (e) and Young's modulus (f) of ZGH451 specimens after different treatments: (a) LPBF, (b) HIP, (c) HIP+HT, and (d) HT

Fig.6e‒6f show the average hardness and the average elastic modulus of the as-deposited specimens and the specimens after different post-treatments obtained from the load-displacement curves. The hardness and elastic modulus of post-treated specimens are smaller than those of LPBF specimens, because after high-temperature treatment, the fine grains will be partly eliminated and partly transformed into recrystallization grains, in addition to the release of residual stresses in the material, which will lead to the reduction of hardness. The average hardness and average elastic modulus of HIP+HT specimens are higher than those of other specimens, probably because of the uniform precipitation of γ' and the reduced micro-defects.

Fig.7 summarizes the stress-strain curves and tensile properties of ZGH451 alloy specimens in four different states. The tensile strength of the LPBF specimen is 1002.7 MPa and the elongation at break is 3.7%. The micro-defects in the specimen act as stress concentration areas under tensile loading, accelerating crack expansion and leading to premature fracture. The HT specimen has a tensile strength of 1276 MPa and an elongation at break of 12.1%, better than those of the LPBF specimen and the HIP specimen. The tensile strength of the HT specimen is 1326 MPa, and the elongation after break is 21.3%, significantly better than those of the LPBF, HIP and HT specimens. The HIP+HT specimen has a smaller size and a higher proportion of γ' phase. The HIP+HT specimen has a smaller size and a higher proportion of γ' phase. In addition, the reduction of grain size also increases the deformation resistance. After HIP treatment, the micro-defects are significantly reduced, and the tensile properties of the alloy are further improved.

Fig.7  Stress-strain curves (a) and tensile properties (b) of ZGH451 specimens after different treatments

3 Conclusions

1) After HIP treatment, the densification of LPBF specimens increases. After HT, the microscopic defects in HIP+HT specimens increase slightly compared with that in HIP specimens.

2) Because of higher dislocation density and larger residual stresses, LPBF specimens have the highest hardness, which decreases after post-treatment. With post-treatment, HIP+HT specimens have the highest hardness due to high volume share of γ' phase and higher densification.

3) At room temperature, the tensile strength and elongation at break of HIP+HT specimens increase to 1326 MPa and 21.3%, respectively, mainly due to the increase in densification, size reduction and increase in the volume share of γ' phase.