Engineering Materials and Their Applications
Microstructure Evolution during Isothermal Aging for Wrought Nickel-Based Superalloy Udimet 520
2020 Volume 61 Issue 8 Pages 1689-1697
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2020 Volume 61 Issue 8 Pages 1689-1697
This paper investigates the evolution of microstructure during isothermal aging at 1173 K in the wrought nickel-based superalloy Udimet 520, after being solution-treated at 1393 K for 4 h, followed by various cooling rates. Age-hardening behavior was observed during isothermal aging for water-quenched (WQ) and air-cooled (AC) specimens after the solution treatment, whereas it could not be detected for a furnace-cooled (FC) specimen. No primary γ′ particles were observed in any continuously cooled samples. For the WQ and AC specimens, the size of the secondary γ′ precipitates increased during the isothermal aging along the Ostwald ripening and their morphology evolved from spherical to an intermediate shape between spherical and cuboidal. Conversely, the secondary γ′ particles exhibited an octodendritic shape for the FC specimen, and the octodendritic character of the secondary γ′ particles was emphasized during isothermal aging, resulting in the splitting of the secondary γ′ particles. It was found that the splitting of γ′ particles occurred during the isothermal aging for Alloy 80A with a lower volume fraction of γ′ phase around 20%.
This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 84 (2020) 11–18.
Fig. 8 FE-SEM images of Udimet 520 solution-treated at 1393 K/4 h/FC followed by the aging treatment at 1173 K/1 h (a), 10 h (b), 100 h (c), and 1000 h (d).
Wrought Ni-based superalloys are widely used for high-temperature structural components in aircraft jet engines and power generation turbines as turbine blade and disc materials with superior high-temperature strength.1) In the last decade, an additional substantial increase in the use of the superalloys has been achieved by utilizing them for boiler piping in an advanced ultra-supercritical thermal power plant and for turbine wheels in an automotive turbocharger to reduce the environmental burden and thus decrease carbon dioxide emissions.2) The superior high-temperature strength of wrought Ni-based superalloys is generally considered to be achieved through solid solution strengthening in the γ matrix and precipitation strengthening by the high-temperature stable intermetallic phases.3) The γ′-Ni3(Al,Ti) phase has an advantage as a precipitation strengthening phase for the following two reasons: first, it is a thermodynamically stable compound with an ordered L12 crystal structure up to the melting point, and its strength has an inverse temperature dependence,4–6) and second, the solubility limit of the alloying elements in the γ′ phase is large,7) and solid solution strengthening occurs.
To improve the high-temperature strength for wrought Ni-based superalloys with a γ–γ′ two-phase microstructure, it is essential to properly control the precipitation microstructure of the γ′ phase. The following three-step process is generally applied to the alloy ingot to produce the practical components of wrought Ni-based superalloys, i) an ingot-to-billet conversion process, such as gyratory forging and radial forging at high temperatures, ii) a solution treatment to control the grain diameter of the γ matrix phase of the billet, and iii) continuous cooling after the solution treatment and the subsequent isothermal aging to encourage γ′ precipitation microstructure.1,3) The solution treatment at a temperature slightly lower than the solvus temperature permits primary γ′ particles to remain at the γ grain boundaries, which prevents excessive coarsening of the γ grains.8) When these alloys are subjected to continuous cooling after solution treatment and subsequent aging treatment, they typically possess microstructures in a face-centered cubic matrix (γ). They can also contain multimodal size distribution of L12 precipitates (γ′)—the primary γ′, which is a few microns in size and located on the γ grain boundaries; two populations of intragranular particles of secondary γ′ of a few hundred nanometers radius; and tertiary γ′ of a few tens nanometers radius.9) To date, research to evaluate the correlation between high-temperature strength and microstructure parameters has been actively conducted by quantitatively evaluating the microstructure parameters such as the particle size, precipitation density, and particle morphology of primary, secondary, and tertiary γ′ phases.10,11)
In engineering components where the metal structures are thick such as turbine disks, the cooling rate after the solution treatment reduces at the inner part of the components relative to their surface. It has been reported that for wrought Ni-based superalloys, the precipitation microstructure of the γ′ phase after continuous cooling followed by solution treatment and the microstructure evolution of the γ′ phase during the subsequent aging show differences between their component surfaces and inner parts.12–16) The thickness of the superalloy components applied to turbine blades and discs has been enhanced with the recent increase in the size of gas turbines for thermal power generation. In this study, we evaluate microstructure evolution during isothermal aging followed by continuous cooling after a solution treatment in Udimet 520,17,18) in which the volume fraction of the γ′ phase is 32%.3) The following four cooling methods were adopted after the solution treatment: water quenching (WQ), oil quenching (OQ), air cooling (AC), and furnace cooling (FC). The effect of the cooling rate on the microstructure evolution during continuous cooling after the solution treatment and during the subsequent aging treatment was clarified for Udimet 520. To simplify as much as possible the multimodal γ′ microstructures obtained in this study, the solution treatment was conducted under supersolvus conditions to obtain a γ single-phase microstructure by completely dissolving the primary γ′ particles.
The alloy used in this study was the wrought Ni-base superalloy Udimet 520, and its composition is shown in Table 1. This superalloy contains aluminum (Al) and titanium (Ti), which are the constituent elements of the γ′ phase, in amounts of 2.1 and 3.1 mass%, respectively. The superalloy ingot was produced via the double melt process of vacuum induction melting and electroslag melting and then hot forged to form a cylindrical billet with a diameter of 90 mm using a the high-speed four-sided forging machine. Cubic specimens with an 8 × 8 × 8 mm3 size were cut out from the received billet and, after solution treatment at temperatures between 1353 and 1413 K for 1.4 × 104 s (4 h), continuous cooling was performed via the following four methods: WQ, 220 K s−1; OQ, 87 K s−1; AC, 10 K s−1; and FC, 0.13 K s−1. After cooling, the aging treatment was conducted at 973–1173 K for 1.1 × 103–3.6 × 106 s (0.3–1000 h).
The microstructure of the specimens was observed by field emission scanning electron microscopy (FE-SEM), which was performed on a cross section parallel to the forging direction of the billet. The embedded samples were subjected to standard mechanical polishing with emery paper and alumina slurry, followed by electrolytic etching using a supersaturated chromic acid phosphate solution in a hot water bath; the current was approximately 10 mA, and the etching time was 30 s. In the FE-SEM observations, secondary electron imaging was conducted at an accelerating voltage of 15.0 kV by focusing on the grains located close to the {001} plane of the γ matrix.
For microstructure observations by high-resolution transmission electron microscopy (HRTEM), thin films cut out from cubic test pieces were shaped into disk-like samples with a diameter of 3 mm and a thickness set to 120 µm by mechanical polishing. These samples were electrolytically polished using a standard twin-jet polisher and a solution of methanol and perchloric acid (9:1). The polishing conditions were set to 243 K and 25 V, giving a polishing current of approximately 30 mA. The perforated foils were examined using a Cs-corrected scanning transmission electron microscope FEI Titan3 G2 60-300 operating at 300 kV with the incident vector of the electron beam B = [001].
The hardness measurements were conducted using a micro Vickers hardness tester; the load was set at 9.8 N, and the holding time was constant at 10 s. Seven measurements were conducted for each specimen, and the hardness was determined from the average of the five measurements, excluding the maximum and minimum values. Care was taken to position the indenter of the hardness tester within γ grains of the alloy specimen for the hardness measurements.
The FE-SEM image of the billet for Udimet 520 is shown in Fig. 1. For wrought Ni-based superalloys, the γ′ particles undissolved during the solution treatment are termed the primary γ′ phase; the γ′ particles precipitated at high temperatures on cooling from the solution treatment are called the secondary γ′ phase; and the extremely fine γ′ particles precipitated at low temperatures on cooling are termed the tertiary γ′ phase.3,9) Primary γ′ particles of approximately 500 nm in size and exhibiting a globular shape were observed at the grain boundaries, and a high density of secondary γ′ particles of approximately 150 nm in size were detected within the γ grains (see Fig. 1). Notably, the grains of the γ matrix were equiaxed in shape, and their diameter was approximately 140 µm throughout the billet.
FE-SEM image of the as-received billet of Udimet 520.
The relationship between the grain diameter of the γ matrix and the temperature of the solution treatment is shown in Fig. 2, where the billet was subjected to the solution treatment at 1353–1413 K for 4 h. When the temperature of the solution treatment was lower than 1373 K, the grain diameter of the γ matrix remained constant at approximately 140 µm. Conversely, the grain diameter increased continuously with rising temperature above 1373 K and reached 280 µm at 1413 K. We inferred that for Udimet 520, the coarsening of the γ grains was suppressed by the pinning of the primary γ′ particles located on the grain boundaries below 1373 K, whereas its coarsening was promoted because of the dissolution of the primary γ′ particles above 1373 K. Therefore, in this study, the solution treatment was performed at 1393 K for 4 h to obtain a γ single-phase microstructure by completely dissolving the primary γ′ particles.
Plots of grain diameter vs. temperature for Udimet 520 solution-treated for 4 h.
The hardness was measured with a micro Vickers hardness tester to clarify the age-hardening behavior of Udimet 520. The cubic specimens were firstly solution-treated at 1393 K for 4 h and continuously cooled using four types cooling method, followed by the aging treatment at 1173 K up to 1000 h. The resulting age-hardening curve is shown in Fig. 3 for every continuously cooled specimen. The hardness was Hv 222 for the specimen continuously cooled by WQ after the solution treatment. When the aging treatment at 1173 K was subjected to the WQ specimen, the hardness began to increase before 0.3 h; after reaching the maximum value of Hv 358 for 1 h, the hardness decreased continuously with increasing aging time, becoming Hv 319 at 1000 h.
Plots of Vickers hardness vs. aging time at 1173 K for Udimet 520 solution-treated at 1393 K/4 h followed by WQ, OQ, AC, and FC.
As the cooling rate decreased, the hardness after the solution treatment increased continuously, reaching Hv 327 for the AC specimen, which was as high as 105 compared with that for the WQ specimen. The increase of hardness with decreasing cooling rate after the solution treatment may be due to the precipitation of the secondary γ′ phase within the γ grains during continuous cooling. In the AC specimen, the hardness gradually increased with increasing aging time and reached the maximum value at Hv 357 for 3 h. Afterward, the hardness decreased continuously with increasing aging time and became Hv 305 for 1000 h, which was quite similar to the WQ specimen.
For the FC specimen, the hardness after cooling was Hv 325, which was very close to that for the AC specimen. However, the hardness for the FC specimen decreased with increasing aging time, and it showed a plateau at around Hv 300 in the aging time between 0.3 and 30 h. Afterward, the hardness for the FC specimen increased at 100 h, and it became close to that for the WQ and AC specimens above 300 h.
3.3 Microstructure evolution during isothermal agingFrom the results of the hardness tests for Udimet 520 presented in the previous section, it was revealed that the WQ and AC specimens after solution treatment showed age-hardening behavior during subsequent aging treatment, yet age-hardening behavior was not identified for the FC specimen. In this section, microstructure evolution during the aging treatment is investigated for both the WQ specimen, which exhibited age-hardening behavior, and the FC specimen, which did not.
3.3.1 The WQ specimenThe increase of hardness in age-hardening curves is generally accepted to correspond to the onset of the precipitation of the secondary phase.3) For the WQ specimen, the hardness abruptly increased at the aging time below 0.3 h (Fig. 3). The HRTEM image for the WQ specimen aged at 973 K for 1 h was observed to clarify the microstructure evolution in the early stage of aging for the WQ specimen. This is shown in Fig. 4(a), together with its selected-area diffraction pattern (SADP) that had an incident beam direction of B = [001] with respect to the γ matrix. The areas where every two lattice point appears dark in the γ matrix are indicated in Fig. 4(a) by the yellow arrowheads. The superlattice diffraction spots derived from the ordered crystal structure in the SADP, which are shown by the red arrows, suggest the precipitation of the γ′ phase in the γ matrix.
HRTEM image of Udimet 520 solution-treated at 1393 K/4 h/WQ followed by the aging treatment at 973 K/1 h, taken with B = [001] (a). IFFT image of (a) is shown in (b).
To clarify the HRTEM image of Fig. 4(a), we have masked and inversely transformed the area other than the high-intensity spots derived from the periodic structures of the γ matrix in the fast-Fourier transformation image, and the resulting inverse FFT (IFFT) image is shown in Fig. 4(b). The area where the lattice points looked dark for every two atom becomes more distinct (yellow arrowheads). Moreover, the particles exhibiting the dark lattice points for every two atom become evident in the IFFT image, as indicated by the white arrowheads. It is considered that the γ′ precipitates with a high degree of order in L12 structure are mixed with those with a low degree of order in the early stage of aging for the WQ specimen. Figure 4(b) shows that the secondary γ′ phase exhibited a spherical shape with a diameter of approximately 3 nm and it precipitated coherently with respect to the γ matrix. The inter-particle distance of the γ′ precipitates was evaluated at approximately 2 nm.
The FE-SEM images of the WQ specimen aged at 1173 K for 10 and 1000 h, corresponding to its over-aged conditions, are shown in Fig. 5. The secondary γ′ particles with a size of approximately 100 nm precipitated uniformly within the γ grains for the specimen aged for 10 h (Fig. 5(a)), which corresponded to the early stage of the over-aged condition. The precipitation density of the secondary γ′ particles was evaluated at 2.4 × 1013 m−2. By contrast, the size of the secondary γ′ particles increased to approximately 400 nm, and the precipitation density decreased dramatically to 1.8 × 1012 m−2 for the specimen aged for 1000 h (Fig. 5(b)), which corresponded to the latter stage of the over-aged condition. The distribution of the secondary γ′ particles was homogeneous and the secondary γ′ phase exhibited a spherical shape for the specimen aged for 10 h. Contrarily, the secondary γ′ precipitates were aligned along one direction and their morphology evolved to an intermediate shape between spherical and cuboidal for the specimen aged for 1000 h.
FE-SEM images of Udimet 520 solution-treated at 1393 K/4 h/WQ followed by the aging treatment at 1173 K/10 h (a) and 1000 h (b).
In Fig. 6, the size of the secondary γ′ particles, d, for the WQ and AC specimens over-aged at 1173 K for 3–1000 h is summarized as a function of the aging time, t. In quantifying the size of the secondary γ′ particles exhibiting an intermediate shape between spherical and cuboidal, the diameter of the circle with the same area was used as d. The value of d was 69 nm at 3 h for the WQ specimen, which was the initial stage of the over-aged condition, and it increased continuously with increasing aging time and reached 427 nm at 1000 h. d linearly increased with increasing aging time in the whole of the over-aged condition, and the gradient was evaluated as 0.33, which suggested that it increased along the Ostwald ripening during the over-aged condition.19–23) The value of d for the AC specimen was slightly smaller than that for the WQ specimen when compared for the same aging time, whereas the slope of the d–t line during the over-aged condition was 0.33, as with the WQ specimen. From the above microstructure observation for the WQ and AC specimens, which exhibited age-hardening behavior, it was revealed that the secondary γ′ phase exhibiting a spherical morphology precipitated at the early stage of aging and that the secondary γ′ phase coarsened along the Ostwald ripening during the over-aged condition with the morphology evolution from spherical to an intermediate shape between spherical and cuboidal.
Plots of diameter of secondary γ′ precipitates vs. aging time at 1173 K for Udimet 520 solution-treated at 1393 K/4 h followed by WQ and AC.
The hardness for the FC specimen decreased continuously with increasing aging time at 1173 K, which was indicative that the age-hardening behavior was not identified for the FC specimen, unlike the WQ and AC specimens. The FE-SEM image for the FC specimen is shown in Fig. 7. The secondary γ′ phase exhibiting an octodendritic shape,24–26) one side of which was approximately 200 nm, precipitated within the whole of the γ grains. It was also observed that a high density of fine tertiary γ′ particles with a size of approximately 10 nm precipitated between the secondary γ′ particles.
FE-SEM image of Udimet 520 solution-treated at 1393 K/4 h followed by FC.
The FE-SEM images for the FC specimen subjected to the aging treatment at 1173 K for 1–1000 h are shown in Fig. 8. For the specimen aged for 1 h (Fig. 8(a)), in which the hardness decreased to approximately Hv 300, the size of the secondary γ′ particles increased to approximately 400 nm while retaining an octodendritic morphology (green arrowheads). Deep cuts were introduced from the side of a part of the secondary γ′ particles into the particle interior (red arrowheads), and a coarse secondary γ′ particle could split into several γ′ particles (yellow arrowheads).27,28) Notably, the split secondary γ′ particles exhibited an intermediate shape between spherical and cuboidal, and fine tertiary γ′ particles observed in the FC specimen were barely discernible.
FE-SEM images of Udimet 520 solution-treated at 1393 K/4 h/FC followed by the aging treatment at 1173 K/1 h (a), 10 h (b), 100 h (c), and 1000 h (d).
As the aging time increased to 10 h (Fig. 8(b)), the split secondary γ′ particles increased in size to approximately 400 nm while retaining their intermediate shape between spherical and cuboidal. When the aging time was further increased to 100 h (Fig. 8(c)), the shape of the secondary γ′ particles evolved into octodendritic form (green arrowheads) and deep cuts were introduced from the side of the secondary γ′ particles again (red arrowheads) followed by the splitting (yellow arrowheads). Finally, for the specimen aged for 1000 h (Fig. 8(d)), the secondary γ′ particles with a size of approximately 400 nm were uniformly distributed and exhibited an intermediate shape between spherical and cuboidal; the same as the specimen aged for 10 h (Fig. 8(b)). As mentioned above, a continuous increase in the size of the secondary γ′ phase could not be identified during the aging treatment for the FC specimen, and it became clear that the following two kinds of morphology evolution occur alternately: i) the secondary γ′ particles with an octodendritic shape split into particles exhibiting an intermediate shape between spherical and cuboidal and ii) the secondary γ′ particles with the intermediate shape coarsen with the morphology evolution to an octodendritic shape.
The morphology of the secondary γ′ particles for the FC specimen aged at 1173 K was quantitatively evaluated as a function of aging time. For quantitatively evaluating the morphology of the γ′ precipitates in a γ–γ′ two-phase microstructure, a mathematical method based on the absolute moment invariants; ω1 and ω2, has recently been proposed by MacSleyne et al.29) who reported that the application of second-order moment invariants is useful in quantitatively describing the morphology of the γ′ particles such as spherical, cuboidal, and octodendritic, in the two-dimensional cross-sectional microstructure of γ–γ′ two-phase superalloys. Moreover, the usefulness of the mathematical method to quantitatively evaluate the morphology of the γ′ precipitates in a γ–γ′ two-phase microstructure has been verified for the wrought Ni-based superalloys.13,14,30) The (ω1, ω2) results for the secondary γ′ particles for the FC specimen aged at 1173 K for 100 h are shown in Fig. 9 as an example. When the particle morphology was perfectly spherical, the ω1 and ω2 values were 4π and 16π2, respectively, corresponding to the domain maxima. The ω1 and ω2 values continuously decreased with the transition to a cuboidal and/or octodendritic form. Only the ω1 value decreased with the increasing aspect ratio of the precipitates, whereas ω2 did not depend on it. The domain of the absolute moment invariants and the particle morphology corresponding to each (ω1, ω2) value are summarized in Fig. 9, where the former can be taken within the shaded area having (4π, 16π2) as the maximum value.
Geometrical locations in the 2-D moment invariant plane (ω1, ω2). In the (ω1, ω2) plane, all 2-D shapes must fall inside the gray region. The rightmost parabola indicates the high symmetry shapes. The open circles indicate the ω1–ω2 plots of secondary γ′ precipitates observed in Udimet 520 solution-treated at 1393 K/4 h/FC followed by the aging treatment at 1173 K/100 h, in which the number of data is 206. The open square means the median of all the ω1–ω2 data, where ω1 = 11.2 and ω2 = 134.2.
The number of the (ω1, ω2) data plots for the secondary γ′ precipitates obtained in the FC specimen aged at 1173 K for 100 h was 206 and is shown using open symbols in Fig. 9. By focusing on the (ω1, ω2) distribution, it was found that most plots appeared to be concentrated in the higher ω1 and ω2 regions in the vicinity of the parabola at the right end of the domain; that is, the value of ω1 ranged from 10.5 to the upper limit of the domain at 4π (∼12.5) and the ω2 value ranged from 120 to the upper limit of the domain at 16π2 (∼157.9). To identify the (ω1, ω2) values representing the morphology of the secondary γ′ precipitates, the median of the 206 plots was determined, and it is shown in Fig. 9 as an open square; it was (ω1, ω2) = (11.2, 134.2).
To evaluate the morphology evolution of the secondary γ′ particles, the median for the FC specimens aged at 1173 K is summarized in Fig. 10, in which the higher region of the domain of the absolute moment invariants ω1 and ω2 is enlarged. The median (ω1, ω2) = (11.3, 142.6) for the FC specimen (plot A). The plots for the FC specimens aged below 3 h (B, C, D) are located in the vicinity of the position of the plot A. By contrast, the median (ω1, ω2) = (12.1, 153.8) for the specimen aged for 10 h (E), which is positioned in the far upper-right direction, along the parabola at the right end of the domain from the positions of plots A–D. When the aging time was further increased to 30 h (F) and 100 h (G), the plots of the median moved substantially toward the lower-left direction and returned to the vicinity of plots A–D. Then, the plots for the specimens aged for 300 h (H) and 1000 h (I) again moved substantially toward the upper-right direction. From the above results of the microstructure analysis, it was clarified for the FC specimen, which did not exhibit the age-hardening behavior, that i) the secondary γ′ particles exhibiting an octodendritic shape precipitate within the γ grains during the continuous cooling followed by the solution treatment and ii) the shape of the secondary γ′ precipitates reciprocates between two types of morphology; such as an octodendritic shape and an intermediate shape between spherical and cuboidal while retaining its small aspect ratio during the subsequent aging treatment.
Plots of the median of ω1–ω2 data for Udimet 520 solution-treated at 1393 K/4 h/FC followed by the aging treatment at 1173 K. The plot A is the median for the as-cooled specimen, and the plots B (0.3 h), C (1 h), D (3 h), E (10 h), F (30 h), G (100 h), H (300 h), and I (1000 h) are the data for the specimens aged at 1173 K.
The schematic illustration showing the morphology evolution of the secondary γ′ particles during isothermal aging at 1173 K for the WQ and FC specimens after solution treatment at 1393 K for 4 h is summarized in Fig. 11. In the cases with relatively higher cooling rates after solution treatment, such as WQ, OQ, and FC for Udimet 520, the secondary γ′ particles exhibiting a spherical morphology precipitated during continuous cooling or at the early stage of isothermal aging, and the secondary γ′ phase coarsened along the Ostwald ripening during the over-aged condition with the morphology evolution from spherical to an intermediate shape between spherical and cuboidal (Fig. 11(i)). However, for the FC specimen with the slowest cooling rate, the secondary γ′ particles exhibited an octodendritic morphology (ii). Deep cuts were introduced from the side of the secondary γ′ particles during the aging treatment (iii), and finally, the secondary γ′ particle split into several particles (iv). The split secondary γ′ particles exhibited an intermediate shape between spherical and cuboidal, the size of which increased during the subsequent aging treatment while retaining the morphology (v). For the FC specimen, the coarsening and splitting shown in (ii)–(v) are repeated during the aging treatment.
Schematic illustration showing the γ′ morphology during the aging treatment for Udimet 520. (i) is the morphology evolution for the WQ specimen, whereas (ii)–(v) denote that for the FC specimen.
Splitting of γ′ particles for Ni-based superalloys27,28) has been reported to occur during continuous cooling after the solution treatment (Nimonic 115,31) IN738LC32)) and during isothermal aging (RR100033)). Although the volume fraction of the γ′ phase is as high as 60%, 50%, and 46% for Nimonic 115, IN738LC, and RR1000, respectively,3) splitting was first identified for Udimet 520 with the γ′ fraction of 32% in this study. To examine whether splitting occurs in wrought Ni-based superalloys with a further reduced γ′ fraction or not, Alloy 80A3) (Ni–19.2Cr–1.4Al–2.2Ti, mass%) with a γ′ fraction of 20% was continuously cooled after the solution treatment, and the microstructural evolution during the subsequent isothermal aging was investigated.
The FE-SEM image of the Alloy 80A specimen continuously cooled at a rate of 0.013 K s−1 after the solution treatment at 1423 K for 1 h, is shown in Fig. 12(a). Note that the cooling rate of 0.013 K s−1 was extremely slow, one-tenth that of the FC (0.13 K s−1). The size of the secondary γ′ particles for the as-cooled specimen was approximately 200 nm, and some of the γ′ particles exhibited an octodendritic shape (green arrowheads), whereas the fine tertiary γ′ particles could not be observed between the secondary γ′ particles. In the specimen subjected to the aging treatment at 1173 K for 0.3 h with respect to the as-cooled specimen, the size of the secondary γ′ precipitates increased to approximately 300 nm while retaining the octodendritic morphology (green arrowheads) and deep cuts were introduced from the side of the secondary γ′ particles (red arrowheads). Furthermore, for the specimen aged at 1173 K for 1 h (Fig. 12(c)), a coarse secondary γ′ particle splits into small cuboidal particles with a size of approximately 150 nm (yellow arrows), which is smaller than the size of the secondary γ′ particles for the as-cooled specimen. From the above results, it is clarified that splitting can occur not only in the wrought Ni-based superalloys with a higher γ′ volume fraction but also in the wrought Ni-based superalloy with a low γ′ volume fraction of 20%.
FE-SEM images of Alloy 80A solution-treated at 1423 K/1 h/SC (slow cooling; 0.013 K s−1) (a) followed by the aging treatment at 1173 K/0.3 h (b) and 1 h (c).
The secondary γ′ particles with an octodendritic morphology can be obtained by continuous cooling with a slow cooling rate followed by the solution treatment for both Udimet 520 with the γ′ volume fraction of 32% and Alloy 80A with the γ′ volume fraction of 20%. In such cases, it was clarified that the splitting of γ′ particles occurs during the subsequent isothermal aging treatment. By contrast, when the cooling rate after the solution treatment is relatively higher and the γ′ phase exhibiting an octodendritic shape cannot be obtained after the continuous cooling, the splitting does not occur during the subsequent isothermal aging and the γ′ phase coarsens along the Ostwald ripening. A future project might systematically investigate the critical cooling rate to obtain the γ′ particles exhibiting an octodendritic morphology after the solution treatment for the wrought Ni-based superalloy from the viewpoints of γ′ fraction and γ–γ′ lattice misfit.
In this study, the wrought Ni-based superalloy Udimet 520 was subjected to continuous cooling at various cooling rates after solution treatment, and the age-hardening behavior and microstructure evolution during the subsequent isothermal aging for the superalloy were investigated. The following results were obtained.
The alloy samples of Alloy 80A used in this study were provided by Daido Steel Co., Ltd. The authors would like to thank Prof. Susumu Onaka of Tokyo Institute of Technology, and Mr. Kenji Okubo and Mr. Ryo Ota of Hokkaido University for their kind assistance with the microstructure observation using electron microscopy. A part of this work was conducted at Hokkaido University, supported by the Nanotechnology Platform Program of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan.
