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Article

Influence of Ageing Treatment on Microstructure and Mechanical Properties of GH4169 Alloy Prepared Using Wire Arc Additive Manufacturing

by
Xuewen You
,
Xinli Song
*,
Wei Geng
and
Zhishen Li
State Key Laboratory of Refractory and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China
*
Author to whom correspondence should be addressed.
Metals 2024, 14(10), 1111; https://doi.org/10.3390/met14101111
Submission received: 24 August 2024 / Revised: 21 September 2024 / Accepted: 22 September 2024 / Published: 29 September 2024
(This article belongs to the Section Additive Manufacturing)
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Abstract

:
The effects of solid-solution and ageing treatments on the microstructure and mechanical properties of a GH4169 alloy made by wire arc additive manufacturing, micro-casting, and forging were researched. The microstructure, along with the size and type of precipitated phase, were analysed using an optical microscope, scanning electron microscope, and transmission electron microscope. The strength and toughness were tested using a tensile testing machine. The results show that a polygonal austenite microstructure was obtained for the GH4169 alloy prepared through wire arc additive manufacturing, micro-casting, and forging, followed by solid-solution and double-ageing treatments at different times. There were a few twins in the austenite matrix. A large number of nano-sized γ″- and γ′-precipitated phases and a small number of Laves phases and MX phases were found in the matrix. The tensile strength and yield strength of the GH4169 alloy increased first and then decreased with the ageing time. After ageing for 16 h, the maximum yield strength was 1287 ± 22 MPa, the maximum tensile strength was 1447 ± 19 MPa, and the elongation was about 19.5%. The main strength mechanism is precipitated phase strength and solid-solution strength. The fracture exhibited obvious ductile fracture characteristics.

1. Introduction

Nickel-based superalloy GH4169 can maintain high strength and good plasticity above 650 °C, and has good high-temperature microstructure stability, creep resistance, and oxidation resistance. Therefore, the alloy is widely used in aerospace turbine blades, missiles, and in biological nuclear energy [1,2]. Often, their harsh service environment causes fatigue fracture or high-temperature creep fracture failure. However, the properties of the GH4169 alloy can be improved through ageing strengthening. Ageing is a process in which the alloy is heated to a certain temperature so that the alloy elements are dissolved in the matrix to form a supersaturated solid solution, then cooled to room temperature quickly, and then heated to a certain temperature to form a precipitated phase. By adding Nb, Ti, Al, and other elements to Ni to form a body-centred cubic structure γ″ (Ni3Nb) and a face-centred cubic structure γ′ [Ni3 (Al, Ti)] strengthening phase, the strength is improved through precipitation strengthening [3,4,5]. Due to the severe segregation of Nb in the GH4169 alloy, Laves phases (Fe, Cr, Ni)2 (Nb, Mo, Ti) or TiN, NbC are also formed [6,7].
Traditionally, the GH4169 alloy is prepared by vacuum induction melting, electroslag remelting, and vacuum consumable remelting. The process is long, the labour intensity is high, and the environment is not friendly. With the application of additive manufacturing technology (AM) to alloys and parts, the production process is greatly shortened. According to the process mechanism, additive manufacturing can be divided into vat photopolymerization, powder bed fusion, directed energy deposition, binder jetting, material extrusion, material jetting, and sheet lamination [8]. Directional energy deposition can be roughly divided into laser metal deposition (LMD) [9], direct metal deposition (DMD) [10], laser solid forming (LSF) [11], direct laser deposition (DLD) [12], wire arc additive manufacturing (WAAM) [13], etc. There is also significant research on GH4169 alloys fabricated through wire arc additive manufacturing, mainly including the influence of the manufacturing process and subsequent heat treatment on the microstructure and properties of the alloy. Michalis Benakis et al. [14] studied the influence of the welding process on weld shape and heat-affected zones at different current frequencies by changing the pulse current frequency and travel speed. The results show that the combination of high-frequency pulse current and low-frequency pulse current can control the weld size and heat-affected zone at the same time, improving the accuracy of additive manufacturing parts. Jia et al. [15] fabricated GH4169 superalloy thin-walled parts by 40 Hz ultra-high frequency pulsed wire arc additive manufacturing. In comparison to conventional pulsed arc deposition, it was found that the microstructure was smaller and the dendrites were arranged in a more orderly manner. After homogenization + ageing heat treatment, a large number of circular granular γ″ phase and γ′ phase were found to be distributed in the microstructure, and the δ phase basically disappeared. The residual Laves phase and carbides were only broken into small-sized blocks or granules. Honnige et al. [16] combined the wire arc additive manufacturing process with interlayer rolling to treat GH4169 alloy components. Through multiple rolling, not only was the content of the Laves phase reduced, but the matrix grains were also refined and the content of texture was reduced. The residual stress in the deposition manufacturing process was also released, which reduced the risk of deformation and cracking. Cao et al. [17] found that the faster solidification rate in the deposition preparation of the GH4169 alloy will produce Nb, Ti, and other segregation bands at the grain boundaries, and the Laves phase will appear inside. When the temperature is higher than 980 °C, the composition segregation band will re-dissolve, which is beneficial to the precipitation of γ″. Zhang et al. [18] found that the higher the ageing temperature, the shorter the transition time from γ″ to δ phase. The microhardness and tensile strength after treatment at different ageing temperatures exhibited a trend of increasing first and then decreasing with time. The stability of the γ″ phase is greatly affected by the ageing time and ageing temperature. The above-mentioned research shows that, by controlling the manufacturing process parameters and subsequent heat treatment, in addition to controlling the formation of columnar crystals, Laves phase, and Nb, Ti, and other segregation bands, γ″/γ′ precipitates favourable for performance are obtained. In this paper, a GH4169 alloy was prepared by a wire arc additive manufacturing and micro-casting–forging composite process (Figure 1). The microstructure evolution after homogenization and solid-solution and ageing treatments was evaluated, and γ″/γ′ precipitates and their influence on the mechanical properties were systematically studied.

2. Experimental Materials and Methods

The experimental material is GH4169 nickel-based superalloy wire with a diameter of 1.0 mm, and the chemical composition meets the national standard composition range. The specific components are shown in Table 1.
Wire arc additive manufacturing technology is based on a discrete/accumulation principle. The arc is used as a moving heat source to continuously melt the welding wire, and then a rolling force is applied to the semi-solid molten pool, which is then cooled and deposited on the substrate. The welding current is 150 A, the voltage is 12 V, the wire feeding rate is 1.0 m/min, the rolling force is 20 kN, and the welding rate is 120 mm/min. A welding process using high-purity argon protection is used, followed by a micro-casting and forging process during the deposition process to achieve the hot rolling of the alloy. Subsequent heat treatment and mechanical property testing were carried out on the round bar samples cut along the direction parallel to the deposition direction. The sample size was Φ13 mm × 100 mm. The time–temperature transformation curve of Inconel 718 alloy [20,21] was used to formulate the homogenization heat treatment process before ageing, which was homogenized at 1080 °C for 1 h, followed by air-cooling to room temperature. The purpose was to eliminate the Laves phase, make the element Nb return to the matrix, and inhibit the element segregation [22]. After solution treatment at 960 °C for 1 h, the alloy was cooled to room temperature, and then aged at 720 °C for 4 h, 8 h, 16 h, and 32 h; it was then cooled to 620 °C for 12 h, and then air-cooled to room temperature (Figure 2).
According to the national standard ASTME9 [23], the aged GH4169 alloy was processed into a standard tensile rod sample, whose specific size is shown in Figure 3. Then, the mechanical properties of the alloy at room temperature were tested using an INSTRON8801 testing machine from NSTRON Corporation, USA. The tensile rate was 0.075 mm/min, and the yield strength was calculated using the 0.2% offset method [24]. Three samples for each ageing time were tested to calculate the mechanical properties error.
The aged GH4169 alloy was coarsely ground and finely ground using 280 #–2000 # sandpaper, and the surface was then mechanically polished with an automatic polishing machine, washed with water and alcohol, and dried. Finally, the surface was corroded using Kallings solution [25] (5% CuCl2 + 100 mL HCl + 100 mL C2H5OH). An Axioplan 2 Imaging Zeiss metallographic microscope from Zeiss, Oberkochen, Germany was used to observe the microstructure. The fracture morphology of the tensile samples and the distribution of micron-sized precipitates and microstructure morphology in ageing samples were analysed by an FEI Nova-Nano-400 scanning electron microscope from FEI, Hillsboro, OR, USA. Transmission samples were prepared by replica extraction using an IB-29510 VET high vacuum coating instrument from JEOL, Tokyo, Japan, and transmission samples were prepared by electrolytic double spraying (electrolyte is 10% HClO4 + 90% absolute ethanol [26]) using a TenuPol-5 electrolytic double spray thinning instrument from Struers, Copenhagen, Denmark. The morphology and size of the precipitated phases in the replica samples and thin-film samples were analysed using a JEM-F200 transmission electron microscope (TEM) from JEOL, Japan, and the types of nano-sized precipitated phases were identified by bright-field, dark-field, and selected-area electron diffraction, alongside high-resolution imaging.

3. Experimental Results

Figure 4 shows the metallographic microstructure of the GH4169 alloy aged for 4–32 h. After ageing treatment for 4 h, polygonal austenite is obtained in the alloy. A large number of twins exist in the austenite crystal, and there are precipitated phases in the matrix. Some of the precipitated phases are distributed in strips. With increases in the ageing time, the austenite grain size increases significantly. The microstructure of GH4169 is FCC austenite with low stacking fault energy, and twins are prone to appear after deformation and recrystallization annealing [27]. The internal grain morphology of the GH4169 alloy obtained by arc micro-casting and forging composite additive manufacturing reported in a previous study [28] reflects the coexistence of columnar crystals and equiaxed crystals, as well as a large number of fragmented dendrites. The existence of columnar crystals is unfavourable to the mechanical properties. In this paper, no columnar crystals were found in the matrix of the GH4169 alloy prepared by wire arc additive manufacturing, micro-casting, and forging processes after solid-solution and ageing heat treatment; equiaxed austenite with a small number of twins and precipitated phases were obtained. Due to the high content of Nb in the alloy, the solidification process was segregated between dendrites to form the Laves phase or MX (M=Ti, Nb X=C, N) phase. The alloy is not completely dissolved into the matrix during homogenization or the solid-solution process, and ageing has an effect on the γ″ phase and γ′ phase. The precipitates were identified by scanning electron microscope (SEM), energy dispersive spectroscopy (EDS) composition analysis, and transmission electron microscopy high-resolution imaging and diffraction techniques. The results are shown in Figure 5, Figure 6, Figure 7, Figure 8 and Figure 9.
Figure 5 shows the microstructure morphology of GH4169 alloy with different ageing times under a scanning electron microscope. The austenite in GH4169 alloy is polygonal, and there are two types of precipitated phases. The precipitated phases with sizes of 200 nm~900 nm are distributed in strips, mainly along the boundary of the molten pool, as shown in Figure 5a,c,e,g. Figure 5b,d,f,h shows that a large number of nano-sized precipitates are distributed on both sides of the austenite grain boundary. With increases in the ageing time, the number of nano-sized precipitates increases. The composition analysis and phase identification of the precipitated phase were carried out by energy spectrum analysis and an electron diffraction technique using a transmission electron microscope, and the results are shown in Figure 6, Figure 7, Figure 8 and Figure 9.
Figure 6a shows the morphology and composition of the complex precipitates, which are mainly MX-type carbides, such as TiN, NbC, etc. TiN is usually precipitated during the solidification of the alloy; therefore, its size is larger. Figure 6b shows the morphology of the precipitated phase in the film sample analysed by transmission electron microscope. The precipitated phase contains Ni, Cr, Mo, Nb, etc., mainly in the Laves phase. The precipitated phase is rod-shaped and about 100 nm in length. The Laves phase that precipitated between dendrites during solidification did not completely dissolve into the matrix during homogenization and solution treatment. The identification of the precipitated phase, with a size of about 50 nm in Figure 6b, is shown in Figure 7, Figure 8 and Figure 9.
Figure 7, Figure 8 and Figure 9 show the bright-field and dark-field images, diffraction patterns, and high-resolution images of nano-sized precipitates. Figure 7 shows the bright-field and dark-field imaging of the precipitated phase. Figure 7a is the bright-field image of the transmission electron microscope of GH4169 alloy aged for 16 h. The matrix phase has a higher contrast (the black area in the figure), and the precipitated phase has a lower contrast (the white area in the figure). The dark-field image of Figure 7b is opposite to the contrast in Figure 7a, and the size of the precipitated phase is about 50 nm. Figure 8a is the transmission electron diffraction pattern of the sample when the ageing time is 4 h. The blue circle represents the diffraction spot of the γ″/γ′ coated structure. The yellow quadrilateral spot is calculated to be the body-centred cubic structure, which is the γ″ precipitated phase. The red triangular spot is calculated to be the face-centred cubic structure, which is the γ′ precipitated phase. Similar to Figure 8a, the diffraction spots of γ″ precipitated phase, γ′ precipitated phase, and the γ″/γ′ coated structure appear in Figure 8b. It can be proven that a large number of γ′ and γ″ precipitates are distributed in the matrix. Figure 9 shows the high-resolution images of the precipitated phase in Figure 7. The Fourier transform is used to calculate the average value of 10 crystal plane spacings, and the average crystal plane spacing of the Figure 9a diagram is 0.1504 nm, which is the interplanar spacing of [1 1 4 ¯ ], and it is determined as the γ″ precipitated phase. The average interplanar spacing of the Figure 9b diagram is 0.1269 nm, which is the interplanar spacing of [2 0 2 ¯ ], and it is determined to be γ′ precipitates. The crystal plane spacing of the Figure 9c diagram is 0.3275 nm, which is the crystal plane spacing of [0 1 1 ¯ ], and it is determined to be the γ″ precipitated phase. The crystal plane spacing in Figure 9d is 0.2071 nm, which is the interplanar spacing of [1 1 ¯ 1], and it is determined to be the γ′ precipitated phase. The distribution is more uniform and coherent with the matrix. The lower the mismatch between the precipitated phase and the matrix during the deformation process, the less obvious the local stress concentration of the substructures, such as grain boundaries. This can effectively inhibit the initiation of cracks and maintain the strength and plasticity of the alloy [29].
Figure 10 shows the size of the precipitated phase at different ageing times counted using the Image-pro-plus software. The calculation method is to select the area of 3 microns × 3 microns in the transmission electron microscope photo and select at least five regions for statistics. The alloy after ageing treatment for 16 h has more precipitated phases than the alloy after ageing treatment for 4 h, and all precipitated phases are less than 50 nanometres in size. The precipitation temperature of γ′ phase is 593–816 °C, and the optimum precipitation temperature is about 620 °C [30,31]. The precipitation temperature of γ″ phase is 595–870 °C, and the optimum precipitation temperature is about 720 °C [32]. The extension of the ageing holding time at 720 °C is more conducive to the precipitation of the main strengthening phase γ″. On the one hand, small-sized precipitates continue to precipitate. However, at the same time, the original precipitated phase will inevitably coarsen; therefore, the average size of the precipitated phase does not change much. The content of γ″ precipitated phase becomes increasingly dispersed, the precipitation strengthening effect is more significant, and the mechanical properties of the material are improved significantly.
Figure 11 shows the engineering stress–strain curve of GH4169 alloy. The wire arc additive manufacturing and micro-cast forging GH4169 alloy have high strength and plasticity after ageing for 4–32 h, but the strength and plasticity are slightly different at different ageing times. When the ageing time is 4 h, the yield strength and tensile strength of GH4169 alloy are the lowest; the yield strength is 1164 ± 14 MPa, the tensile strength is 1314 ± 18 MPa, and the elongation is about 26%. The tensile strength and yield strength are higher than the room temperature performance σb > 1240 MPa specified in the forging technical standard σ0.2 > 1040 MPa [33]. When the ageing time is extended to 16 h, the maximum yield strength is 1287 ± 22 MPa, the tensile strength is 1447 ± 19 MPa, and the elongation is about 19.5%. When the ageing time is extended to 32 h, the yield strength decreases to 1191 ± 21 MPa.
Figure 12 shows the tensile fracture morphology of GH4169 alloy with different ageing times. From Figure 12a,c,e,g, it can be seen that the fracture shape is a cup–cone fracture, and the section and the tensile stress direction exhibit a 45° angle. The surface is very rough, and the degree of necking is not large. The tensile fracture is mainly composed of the fibre area, radiation area, and edge shear lip. With extensions in the ageing time, the shear lip area becomes smaller and smaller, and the internal fibre area and radiation area increases. There are a large number of dimples in the fracture, which reflects typical ductile fracture. It can be seen from Figure 5 and Figure 6 that there are a large number of TiN and NbC precipitated Laves phases in the austenite matrix. The formation of dimples is related to the precipitated phase. During the tensile process, dislocations move to the surroundings of the precipitated phase to produce pile-up, resulting in the formation of micropores [34].

4. Analysis and Discussion

The GH4169 alloy prepared by wire arc additive manufacturing has excellent strength and plasticity after ageing for 4–32 h. The yield strength increases from 1164 MPa to 1287 MPa, the tensile strength increases from 1314 MPa to 1447 MPa, and the elongation is 19.5–26%. Its excellent mechanical properties are closely related to the microstructure evolution. The GH4169 alloy is prepared by wire arc additive manufacturing and micro-casting and forging. Furthermore, after solid-solution and double ageing treatment, the columnar crystals are eliminated and equiaxed crystals are obtained. The strengthening mechanism includes second-phase precipitation strengthening, solid-solution strengthening, and grain-boundary strengthening [35]. The shape, quantity, size, and distribution of the precipitated phase affect the strength and toughness of the alloy. Figure 13 shows the schematic diagram of the γ″/γ′ composite sandwich structure precipitated phase.
According to the morphology and diffraction spots of the nano-precipitates shown in Figure 7, Figure 8 and Figure 9, it can be seen that γ″/γ′ is a composite precipitate, and it is a γ″/γ′/γ″ sandwich structure, as shown in Figure 13 [36]. When ageing at 720 °C, the precipitation rate of the γ″ phase is lower than that of the γ′ phase. With increases in the ageing time, although the γ′ secondary strengthening phase is continuously precipitated, the coarsening rate of γ′ phase is faster than the precipitation rate [37], resulting in an increase in the number of precipitated phases; however, the first formed precipitated phase is coarsened, and the average size of the precipitated phase increases. When the furnace cooling is reduced to 620 °C, the precipitation rate of the γ″ phase is higher than that of the γ′ phase, and a large number of γ″ phases are precipitated. At this time, the Nb element continues to diffuse from the γ′ phase to the γ″/γ′ interface, and the γ″/γ′ can easily form flat discs and ellipsoids during the ageing process. This structure reduces the overall elastic energy and lattice distortion of the precipitated phase, which is a benefit for the mechanical properties of the alloy [38].
When the ageing time is 4 h (which is relatively short), the number of precipitated phases is small, and the precipitated phases are mainly distributed along the weld line and their distribution is not uniform, so the precipitation strengthening effect is weak, the strength is low, and the plasticity is high. As the ageing time extends to 16 h, the number of precipitated phases in the matrix increases, and the distribution becomes more dispersed. The nano-sized γ′ and γ′ precipitated phases are coherent with the matrix. The relative strength contribution mechanism of coherent precipitation is mainly through dislocation cutting. The precipitated phase particles destroy the atomic health, increase the P-N force, and increase the resistance of the dislocation motion. The dislocation cuts the precipitated phase particles to form a slip step, increases the interface energy, and increases the energy consumed by the dislocation motion. The coherent strain field of the precipitated phase particles interacts with the dislocation strain field to produce a strengthening effect through the coherent strain field zone. Therefore, the strengthening effect on the matrix is high, and the strength reaches a maximum [39]. When the ageing time is further extended to 32 h, the precipitated phase is coarsened and the transformation from γ″ phase to δ phase occurs in part, which makes it lose the coherent relationship with the matrix, and the precipitation strengthening effect is weakened; consequently, the strength is reduced.
In addition, grain-boundary strengthening and solid-solution strengthening also contribute to the strength of GH4169. Both austenite grain boundaries and twin boundaries hinder the movement of dislocations during deformation, which is beneficial to improving the strength. According to the Hall–Petch formula, the smaller the original austenite grain size, the higher the material strength. As the ageing time is extended to 32 h, the average grain size of austenite increases slightly, so its strength decreases. The solid-solution strengthening elements Fe, Cr, Mo, Co, and so on [40,41] are added to the alloy, which causes lattice distortion and increases the resistance of dislocation slip, thus playing a good solid-solution-strengthening role.
In summary, it can be seen that a large number of nano-coherent precipitated phases are obtained after the solid-solution and double-ageing treatment of the GH4169 alloy by arc micro-casting and forging composite additive manufacturing. Precipitation strengthening, solid-solution strengthening, and grain boundary strengthening increase the strength and plasticity of GH4169 alloy. A large number of dimples in the tensile fracture are closely related to a large number of micron-sized precipitates in the matrix. During the tensile deformation process of the alloy, dislocations move to the interface between the precipitated phase particles and the matrix and accumulate, initiating cracks. With the plastic deformation, a large number of holes are produced; therefore, a large number of dimples can be observed in the fracture of the tensile sample.

5. Conclusions

The GH4169 alloy was made by wire arc additive manufacturing and micro-casting forging. The effect of solid-solution and double ageing treatment processes on microstructure and mechanical properties was researched in detail, and the main conclusions are as follows.
(1) The equiaxed austenite grains were obtained and a large number of nano-scale γ′ and γ″ precipitated phases and a little of MX and Laves phase were found in the matrix.
(2) With the increase in ageing time from 4 h to 16 h, the number of nano-sized γ′ and γ″ precipitated phases increased, which is a benefit for the improvement of the strength of the GH4169 made by WAAM and ageing treatment.
(3) The best strength and toughness of GH4169 alloy prepared by WAAM plus micro-casting and forging were obtained after solution and ageing at 720 °C for 16 h.
The tensile strength is 1447 MPa, the yield strength is 1287 MPa, and the elongation is 19.5%.
(4) The strength and plasticity of GH4169 alloys prepared by WAAM and micro-casting and forging are superior to those prepared by traditional melting and forging methods.

Author Contributions

Conceptualization, X.Y.; methodology, X.Y., W.G. and Z.L.; validation, X.Y. and X.S.; investigation, X.Y.; writing—original draft preparation, X.Y.; writing—review and editing, X.Y. and X.S.; project administration, X.S. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by technology innovation project of Hubei Province of China, grant number, No. 2022BEC025.

Data Availability Statement

The datasets presented in this article are not readily available because privacy or ethical restrictions.

Acknowledgments

The authors would like to thank Li Yuanyuan, Analysis and Testing Center, Wuhan University of Science and Technology.

Conflicts of Interest

The authors declare no conflicts of interest. We solemnly state that all authors of this article have read and approved the final version submitted. The content of this manuscript has never been copyrighted or published before, and is not currently considered for publication elsewhere.

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Figure 1. Schematic diagram of the experimental material preparation.
Figure 1. Schematic diagram of the experimental material preparation.
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Figure 2. Heat treatment process parameters.
Figure 2. Heat treatment process parameters.
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Figure 3. Diagram of standard tensile specimen.
Figure 3. Diagram of standard tensile specimen.
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Figure 4. Optical microscope microstructure of GH4169 alloy at different ageing times: (a) 4 h; (b) 8 h; (c) 16 h; and (d) 32 h.
Figure 4. Optical microscope microstructure of GH4169 alloy at different ageing times: (a) 4 h; (b) 8 h; (c) 16 h; and (d) 32 h.
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Figure 5. The microstructure of GH4169 nickel-based superalloy after ageing for different times under scanning electron microscope: (a,b) 4 h, (c,d) 8 h, (e,f) 16 h, and (g,h) 32 h.
Figure 5. The microstructure of GH4169 nickel-based superalloy after ageing for different times under scanning electron microscope: (a,b) 4 h, (c,d) 8 h, (e,f) 16 h, and (g,h) 32 h.
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Figure 6. Morphology and energy spectrum of GH4169 alloy aged for 4 h and 16 h under transmission electron microscope: (a) the precipitated phase in the replica sample, A and B are precipitated phases; (b) the precipitated phase in the flake sample, 1, 2, 3 and 4 are precipitated phases; (c) the energy spectrum diagram of the precipitated phase A in (a); and (d) the energy spectrum diagram of the precipitated phase 2 in (b).
Figure 6. Morphology and energy spectrum of GH4169 alloy aged for 4 h and 16 h under transmission electron microscope: (a) the precipitated phase in the replica sample, A and B are precipitated phases; (b) the precipitated phase in the flake sample, 1, 2, 3 and 4 are precipitated phases; (c) the energy spectrum diagram of the precipitated phase A in (a); and (d) the energy spectrum diagram of the precipitated phase 2 in (b).
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Figure 7. Bright- and dark-field images of precipitated phases in GH4169 alloy ageing for 16 h under transmission electron microscope: (a) bright-field images; and (b) dark-field images.
Figure 7. Bright- and dark-field images of precipitated phases in GH4169 alloy ageing for 16 h under transmission electron microscope: (a) bright-field images; and (b) dark-field images.
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Figure 8. Diffraction spots of precipitated phase under transmission electron microscope: (a) ageing for 4 h; and (b) ageing for 16 h.
Figure 8. Diffraction spots of precipitated phase under transmission electron microscope: (a) ageing for 4 h; and (b) ageing for 16 h.
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Figure 9. High-resolution imaging, the lower right corner is the crystal plane spacing of the selected area after fast Fourier transform and inverse fast Fourier transform, (a,b) sample ageing for 4 h; and (c,d) sample ageing for 16 h.
Figure 9. High-resolution imaging, the lower right corner is the crystal plane spacing of the selected area after fast Fourier transform and inverse fast Fourier transform, (a,b) sample ageing for 4 h; and (c,d) sample ageing for 16 h.
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Figure 10. Size and distribution of precipitates at different ageing time: (a) 4 h; and (b) 16 h.
Figure 10. Size and distribution of precipitates at different ageing time: (a) 4 h; and (b) 16 h.
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Figure 11. The tensile stress–strain curves of GH4169 alloy ageing for different times.
Figure 11. The tensile stress–strain curves of GH4169 alloy ageing for different times.
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Figure 12. The fracture morphology of GH4169 alloy after ageing for different times obtained by scanning electron microscope: (a,b) 4 h; (c,d) 8 h; (e,f) 16 h; and (g,h) 32 h.
Figure 12. The fracture morphology of GH4169 alloy after ageing for different times obtained by scanning electron microscope: (a,b) 4 h; (c,d) 8 h; (e,f) 16 h; and (g,h) 32 h.
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Figure 13. Simplified structure of γ″/γ′ composite sandwich.
Figure 13. Simplified structure of γ″/γ′ composite sandwich.
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Table 1. Chemical composition of GH4169 (wt%).
Table 1. Chemical composition of GH4169 (wt%).
ElementExperimental SamplesASTMB670-07 [19]
Ni53.7550.0~55.0
Cr18.1017.0~21.0
FeBalBal.
C0.04≤0.08
Si0.24≤0.35
P<0.01≤0.015
S<0.01≤0.015
Cu0.08≤0.30
Mo3.042.80~3.30
Ti0.850.65~1.15
Mn0.32≤0.35
Nb+Ta5.134.75~5.50
Co0.82≤1.0
Al0.460.20~0.80
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You, X.; Song, X.; Geng, W.; Li, Z. Influence of Ageing Treatment on Microstructure and Mechanical Properties of GH4169 Alloy Prepared Using Wire Arc Additive Manufacturing. Metals 2024, 14, 1111. https://doi.org/10.3390/met14101111

AMA Style

You X, Song X, Geng W, Li Z. Influence of Ageing Treatment on Microstructure and Mechanical Properties of GH4169 Alloy Prepared Using Wire Arc Additive Manufacturing. Metals. 2024; 14(10):1111. https://doi.org/10.3390/met14101111

Chicago/Turabian Style

You, Xuewen, Xinli Song, Wei Geng, and Zhishen Li. 2024. "Influence of Ageing Treatment on Microstructure and Mechanical Properties of GH4169 Alloy Prepared Using Wire Arc Additive Manufacturing" Metals 14, no. 10: 1111. https://doi.org/10.3390/met14101111

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