Microstructure Quantification in Nickel-Based Superalloy Udimet 720Li
2019 年 60 巻 4 号 p. 593-601
詳細
2019 年 60 巻 4 号 p. 593-601
The morphology of γ′ precipitates for the wrought Ni-based superalloy Udimet 720Li was quantitatively evaluated by applying the absolute moment invariant technique. The diameter of the secondary γ′ precipitates, d, increased continuously with the decrease of the cooling rate, v, after solution treatment at 1473 K for 1 h along the following equation; d ∝ v−0.4. In addition, with decreasing cooling rate, the shape of the secondary γ′ precipitates changed from spherical to octodendritic. For the oil-quenched alloy after solution treatment, d continuously increased with the increase of the aging time at 1173 K, and the spherical shape remained unchanged during the process. On the contrary, for the furnace-cooled alloy, the shape evolved from octodendritic to spherical with the aging time, exhibiting an almost constant d.
This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 82 (2018) 375–383.
Fig. 6 Field-emission scanning electron microscopy images of Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by oil quenching (a), air cooling (b), and furnace cooling (c).
Wrought Ni-based superalloys are promising candidates as high-temperature structural materials with superior strength and are also widely applied to fabricate the turbine components of aircraft jet engines. In recent years, to improve combustion efficiency in thermal power generation, the use of superalloys in advanced ultra-supercritical pressure thermal power plants, where the combustion temperature is required to increase with time, is being examined.1) The superior high-temperature strength of wrought Ni-based superalloys is generally attributed to solid solution strengthening in the γ matrix and precipitation strengthening by the high-temperature stable intermetallic phases such as γ′-Ni3(Al, Ti), γ′′-Ni3Nb and η-Ni3Ti.2) Among them, the γ′ phase has an advantage as a precipitation strengthening phase because of the following reasons.3–5) First, it is a Berthollide compound with an ordered L12 crystal structure up to the melting point, and its strength has an inverse temperature dependence. Second, the solubility limit of the alloying elements in the γ′ phase is large, and solid solution strengthening can be fully used. Third, it can precipitate coherently within the γ matrix phase.
To improve the high-temperature strength of wrought Ni-based superalloys with a γ/γ′ two-phase microstructure, it is extremely important to properly control the precipitation microstructure of the γ′ phase. When performing solution treatment below the subsolvus temperature, the primary γ′ particles remain at the γ grain boundaries, and the increase of the grain diameter in the γ matrix during this treatment can be prevented. Continuous cooling from the high temperatures of the solution treatment also results in the precipitation of submicron-sized secondary γ′ particles within the γ grains in the high-temperature region and of finer tertiary γ′ particles in the low-temperature one, eventually leading to multimodal γ′ microstructures. To date, researches have been actively conducted to obtain optimal microstructures promoting the high-temperature strength by quantitatively evaluating microstructure parameters such as the volume fraction and the particle size and precipitation density of primary, secondary, and tertiary γ′ phases.6–8)
In the evaluation of precipitation microstructures, attempts to quantify the morphology of γ′ particles are currently very limited. The secondary γ′ particles precipitated in the γ grains during cooling after the solution treatment and/or during the subsequent aging treatment exhibit a spherical, cuboidal, or intermediate shapes in conventional wrought Ni-based superalloys. In such cases, the quantification of these shapes has been attempted using the following methods: i) defining the intrinsic shape factors, such as the area of the squares inscribed in the precipitated particles,9,10) and ii) mathematically expressing the shape of the precipitated particles.11) On the contrary, in wrought Ni-based superalloys with a higher volume fraction of the γ′ phase (fv > 30%), the secondary γ′ particles precipitated during continuous cooling after the solution treatment exhibit a complex morphology, called octodendritic, in addition to the three shape types mentioned above.12,13)
MacSleyne et al. stated that the mathematical technique using the absolute moment invariants is useful for quantitatively evaluating a variety of precipitated γ′ morphologies, including octodendritic shape.14) In this study, we first quantitatively evaluated the cooling rate effect after the solute treatment on the morphology of the precipitated γ′ particles using the absolute moment invariants by focusing on the Udimet 720Li alloy, where fv is very high (fv > 45%) and the γ/γ′ misfit is close to zero among wrought Ni-based superalloys with γ/γ′ two-phase microstructures.15,16) Next, the continuously cooled alloys were subjected to isothermal aging to observe the effect of this treatment on the morphology of the secondary γ′ precipitates. To simplify the multimodal γ′ microstructures as much as possible, the solution treatment was carried out above the supersolvus temperature to exclude the primary γ′ particles in the microstructure.
The alloy used in this study is the wrought Ni-based superalloy Udimet 720Li, and its composition is shown in Table 1. This superalloy contains Ti and Al, which are the constituent elements of the γ′ phase, in amounts of 5.07 and 2.70 mass%, respectively, and fv is about 45%, which is the largest value among wrought Ni-based superalloys. Superalloy ingots with a diameter of 500 mm were produced via the triple melt process and then hot forged to form cylindrical billets with a diameter of 250 mm. During the triple melt process, the alloy composition was adjusted by vacuum induction melting (VIM), the impurities were removed by electroslag remelting (ESR), and finally, a homogeneous microstructure was obtained by vacuum arc remelting (VAR). This method is commonly applied in the manufacturing of wrought Ni-based superalloys for turbine disks.3,4) A 60 mm portion was cut out from the outer periphery of the billets, and this was used as a received material.
Cubic specimens with an 8 × 8 × 8 mm3 size were cut out from the received billets, and after solution treatment at 1473 K for 1 h, cooling was performed via the following four methods: water quenching (WQ; 220 K/s), oil quenching (OQ; 87 K/s), air cooling (AC; 10 K/s), and furnace cooling (FC; 0.13 K/s). After cooling, the cubic specimens were cut into two pieces, mechanically polished with emery paper and alumina thriller, and then subjected to electrolytic etching in a solution of chromic acid saturated phosphoric acid. Field-emission scanning electron microscopy (FE-SEM) was used for microstructure observation, which was performed on a cross section parallel to the forging direction of the cylindrical billet. The hardness measurements were conducted using a micro-Vickers hardness tester. The aging treatment was carried out at 1173 K, with an aging time ranging from 3.6 × 103 s (1 h) to 3.6 × 105 s (100 h).
For the microstructure observation via transmission electron microscopy (TEM), 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 JEOL transmission electron microscope equipped with a double-tilt goniometer stage operating at 200 kV. In the TEM analysis, the incident vector of the electron beam was fixed as B = [001] so that the morphology of the γ′ particles could be specified most evidently.
2.2 Quantitative evaluation method of the γ′ morphologyFor quantitatively evaluating the morphology of the γ′ precipitates in a γ/γ′ two-phase microstructure, a mathematical method based on the absolute moment invariants has been recently proposed by MacSleyne et al.,14) which reported that the application of second-order moment invariants is useful in quantitatively describing this morphology in the two-dimensional (2D) cross-sectional microstructure of γ/γ′ two-phase superalloys. Here, the absolute moment invariants ω1 and ω2 are defined as
| \begin{equation} \omega_{1} = \frac{2A^{2}}{\bar{\mu}_{20} + \bar{\mu}_{02}}(0 < \omega_{1} \leq 4\pi) \end{equation} | (1) |
| \begin{equation} \omega_{2} = \frac{A^{4}}{\bar{\mu}_{20}\bar{\mu}_{02} - \bar{\mu}_{11}^{2}}(\omega_{1}^{2} \leq\omega_{2} \leq 16\pi^{2}). \end{equation} | (2) |
| \begin{equation} \bar{\mu}_{pq} = \iint\left(x - \frac{\mu_{10}}{A}\right)^{p}\left(y - \frac{\mu_{01}}{A}\right)^{q}d^{2}\textbf{r} \end{equation} | (3) |
The morphological characteristics described by the absolute moment invariants can be summarized as follows. The ω1 value is invariant to the similarity transformations, such as translation, rotation, expansion, and contraction of the precipitated particles. By contrast, the ω2 value is invariant not only to such similarity transformations but also to the affine ones, such as uniform shear and non-uniform expansion. The domain of the absolute moment invariants and the particle morphology corresponding to each (ω1, ω2) value are summarized in Fig. 1, where the former can be taken within the shaded area having (4π, 16π2) as the maximum value.
Geometrical locations in the two-dimensional (2D) moment invariant plane (ω1, ω2), where all 2D shapes must fall inside the gray region. The rightmost parabola indicates the high-symmetry shapes.
When the precipitate morphology is perfectly spherical, the ω1 and ω2 values are 4π and 16π2, respectively, corresponding to the domain maxima. As the square degree increases, the position of the (ω1, ω2) plot moves to the left and down on the parabola at the right end of the domain. When the morphology is perfectly cuboidal, ω1 and ω2 become 12.0 and 144, respectively. As the position of the (ω1, ω2) plot moves further downward and leftward on the parabola at the right end of the domain, the morphology evolves from cubic to octodendritic. Both the ω1 and ω2 values approach zero when the dendrite arm is lengthened. On the contrary, when the aspect ratio of the precipitates increases from a perfectly spherical morphology to an elliptical one with an axis ratio of 1.5, ω1 and ω2 become 11.6 and 16π2, respectively; that is, ω1 decreases compared to that for a perfectly spherical morphology, whereas ω2 remains unchanged. As described above, ω1 and ω2 reach the maximum values of 4π and 16π2, respectively, in the case of a perfectly spherical morphology and continuously decrease with its transition to a cuboidal and/or octodendritic one. However, only the ω1 value decreases also with increasing aspect ratio of the precipitates, whereas ω2 does not depend on it.
In the hot-forged billets of Ni-based superalloys, the microstructure near the surface generally differs from the inner one.4,18) To shed light on the microstructure of continuously cooled Udimet 720Li-based samples, that of the as-received billets was firstly examined. The hardness measurements were performed as a function of the distance from the billet surface, x, to confirm the non-uniformity of their microstructure, and the results are shown in Fig. 2, as open symbols, as a function of x. The Vickers hardness was Hv 570 near the billet surface and decreased dramatically with increasing x, becoming Hv 410 at x = 15 mm; this decrease was less significant when further increasing x, and the hardness reached Hv 376 at x = 60 mm. The microstructure observation was performed to understand the reason for the marked increase in hardness detected near the billet surface, typically at x < 15 mm.
Hardness of the Udimet 720Li billets as a function of the distance from the surface, measured as received (open) and after the solution treatment at 1473 K for 1 h followed by water quenching (solid).
The FE-SEM images taken near the billet surface (x = 6 mm) and at x = 36 mm, where hardness was sufficiently reduced, are shown in Fig. 3. At x = 6 mm (Fig. 3(a)), primary γ′ particles with a 1–3 µm size and fine secondary ones precipitated at a high density among them were observed; the estimated volume fraction of the primary γ′ particles was about 20%. At x = 36 mm (Fig. 3(b)), the size of the primary γ′ particles was 2–4 µm, and that of the secondary ones was also obviously larger than that observed near the billet surface.
Field-emission scanning electron microscopy images of the as-received billets of Udimet 720Li taken at (a) 6 and (b) 36 mm from the surface.
To clarify the size and morphology of the secondary γ′ particles for both samples shown in Fig. 3, the regions where they precipitated at a high density were observed also at a higher magnification. The results are shown in Fig. 4. At x = 6 mm (Fig. 4(a)), where the marked increase in hardness was detected, their mean size was about 150 nm, and many particles exhibited a spherical or cuboidal morphologies; furthermore, tertiary γ′ particles with a spherical shape and a size of about 30 nm were partially detected among the secondary ones. By contrast, at x = 36 mm (Fig. 4(b)), where the hardness fully decreased, the secondary γ′ particles with a size of about 300 nm precipitated with an octodendritic morphology,13) and fewer tertiary γ′ particles were detected. These results suggest that the dramatic increase in hardness near the billet surface, as observed in Fig. 2, could be attributed to the decrease in the size of the secondary γ′ particles and the precipitation of the fine tertiary γ′ ones.
Magnified view of the secondary γ′ particles precipitated in the as-received billets of Udimet 720Li at (a) 6 and (b) 36 mm from the surface.
The Vickers hardness results for the water-quenched alloy after the solution treatment at 1473 K for 1 h, which is sufficiently higher than the solvus temperature of Udimet 720Li (1428 K), are shown in Fig. 2 as solid symbols; the hardness was around Hv 410 both near the billet surface and in the bulk. This result suggests that the non-uniformity of the original microstructures near the billet surface and in the bulk was sufficiently dissolved by the solution treatment.
3.2 Microstructure of the continuously cooled alloyCubic specimens with an 8 mm side were cut out from the as-received billets, and subjected to WQ, OQ, AC, and FC after the solution treatment at 1473 K for 1 h; the resulting Vickers hardness values are summarized as a function of the distance from the specimen surface, y, in Fig. 5, where y = 4 mm indicates the center position in the specimen. The grain size of the γ matrix phase of the specimens subjected to the solution treatment was 839 µm, and thereafter, indenter was taken care to put within γ grains in the hardness measurement. The hardness of the furnace-cooled alloy remained constant at Hv 365 regardless of y, whereas those of the water-quenched, oil-quenched, and air-cooled ones were almost constant at y ≧ 1.5 mm but slightly increased near the surface at y < 1.5 mm.
Hardness of the Udimet 720Li cubic specimens as a function of the distance from the surface measured after the solution treatment at 1473 K for 1 h followed by water quenching (WQ), oil quenching (OQ), air cooling (AC), and furnace cooling (FC).
The FE-SEM images taken at y ≧ 1.5 mm, which can be assumed to be unaffected by the increased hardness near the sample surface, are shown in Fig. 6. In the oil-quenched alloy, which had a relatively high cooling rate (Fig. 6(a)), secondary γ′ particles with a spherical shape and a size of about 40 nm precipitated homogeneously within the γ grains. In the air-cooled alloy (Fig. 6(b)), the size of the secondary γ′ particles was about 90 nm, and compared to the oil-quenched specimen, the frequency of the smooth interface increased, and the particle morphology tended to be more quadrangular. In the furnace-cooled alloy, which had the slowest cooling rate (Fig. 6(c)), the secondary γ′ particles exhibited a size of about 500 nm and an octodendritic morphology; in addition, fine tertiary γ′ particles with a size below 10 nm precipitated at a high density among the secondary γ′ particles were clearly recognized. No primary γ′ particles were observed in any of these continuously cooled specimens. The volume fractions of the secondary γ′ precipitates in the water-quenched, oil-quenched, and air-cooled alloys were 30%, 39%, and 42%, respectively. The larger hardness values for the oil-quenched and air-cooled alloys compared to the water-quenched one, as shown in Fig. 5, could be attributed to the corresponding higher fv.
Field-emission scanning electron microscopy images of Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by oil quenching (a), air cooling (b), and furnace cooling (c).
The results of the microstructure observation performed on the Udimet 720Li specimens subjected to the solution treatment above the solvus temperature and then cooled in four different ways can be summarized in the following four points. First, no primary γ′ particles were observed at any cooling rate from WQ to FC. Second, the size of the secondary γ′ precipitates continuously increased with decreasing cooling rate. Third, the morphology of the secondary γ′ precipitates evolved from a spherical shape to a cuboidal and/or an octodendritic one when decreasing the cooling rate. Fourth, the precipitation of tertiary γ′ particles became evident with the decrease of the cooling rate. As described above, the difference observed in the microstructure when the Udimet 720Li specimens were cooled at four different cooling rates after the solution treatment was most pronounced in the size and morphology of the secondary γ′ precipitates. Therefore, in the following sections, these two parameters are quantified and evaluated from the cooling rate viewpoint.
3.3 Secondary γ′ particle size and cooling rateThe size of the secondary γ′ particles is plotted against cooling rate in Fig. 7 for the Udimet 720Li specimens that were subjected to solution treatment at 1473 K for 1 h followed by continuous cooling via the four methods mentioned above; the sizes of the secondary γ′ particles reported for the same superalloy by Mao et al.19) and Radis et al.20) are also included in the figure. When evaluating the size of the precipitates having cuboidal and/or octodendritic morphologies, the diameter of a circle having the same area as the particles was adopted as d. The size of the secondary γ′ particles continuously increased with decreasing cooling rate, and when displayed in logarithmic scales, both parameters appeared as the straight line with a gradient of −0.4. This means that the following equation, stating the relation between the size of the secondary γ′ precipitates and the cooling rate, was satisfied for the solution-treated and continuously cooled Udimet 720Li specimens:
| \begin{equation} d \propto v^{-n} \end{equation} | (4) |
In René 88DT (fv = 42%)21) and AD 730 (fv = 40%),22) which are practical wrought Ni-based superalloys with higher fv, the relationship between d and v after the solution treatment is reported and included in Fig. 7. For both superalloys, the size of the secondary γ′ particles continuously increases with the decrease of the cooling rate, as for the Udimet 720Li case. In addition, the four plots for René 88DT and AD 730 are approximately parallel to the straight line corresponding to n = 0.4 obtained for Udimet 720Li.
In Ni-based superalloys, the size of the secondary γ′ particles does not necessarily increase when the cooling rate is reduced after the solution treatment, as reported for IN-738 LC, which has an equilibrium fraction of the γ′ phase above 50%.23) This is because a phenomenon called splitting occurs during the superalloy cooling, in which a coarsened secondary γ′ precipitate divides into eight smaller γ′ particles. This splitting occurs markedly when Ni-based superalloys with high γ′ fraction (>50%) and large γ/γ′ lattice misfit are subjected to continuous cooling24) or isothermal aging.25) It is evident that this phenomenon does not occur in the Udimet 720Li case, even during FC; therefore, in this study, d continuously increases with the decrease of the cooling rate, and the universal relationship shown in eq. (4) is satisfied.
3.4 Secondary γ′ particle morphology and cooling rateThe TEM images of the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by OQ and FC are shown in Fig. 8 with the incident beam direction of B = [001]. The microstructure of the oil-quenched alloy was recorded as a dark-field image, whereas that of the furnace-cooled one was a bright-field image. In the oil-quenched alloy, most precipitates exhibited a spherical shape, and d was 37 nm, as estimated from the FE-SEM image (Fig. 6(a)); in some cases, precipitates close to each other had a smooth interface, indicated by the arrowheads in Fig. 8(a). By contrast, the octodendritic shape was clearly detected with B = [001] for the furnace-cooled alloy. On the basis of the TEM analysis, the morphology of the secondary γ′ particles was quantitatively evaluated by using absolute moment invariants.
Transmission electron microscopy images of Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by (a) oil quenching (dark-field image) and (b) furnace cooling (bright-field image), taken with incident beam direction B = [001]. The parallel interface between neighboring γ′ precipitates is indicated with arrowheads.
The (ω1, ω2) plots for the oil-quenched alloy with B = [001] are shown in Fig. 9. The domains were 0 < ω1 ≦ 4π and 0 < ω2 ≦ 16π2, as shown in Fig. 1, and the plots obtained in this study were located on their higher region; therefore, the region with ω1 ≧ 9 and ω2 ≧ 100 is enlarged in Fig. 9. The data plots (38) for the secondary γ′ precipitates obtained in this study are shown using open symbols in the figure. By focusing on the (ω1, ω2) distribution, it is found that 38 plots appeared to be concentrated in the higher ω1 and ω2 regions, that is, the ω2 value was hardly reduced from the upper limit of the domain at 16π2 (∼157.9). In addition, although the ω1 values were concentrated near the upper limit of the domain at 4π (∼12.5), some data were positioned in the ω1 < 12 region. The plot with the smallest ω1 value among the 38 data obtained for the oil-quenched alloy was (ω1, ω2) = (11.9, 157.4), which indicates an elliptical morphology with an axis ratio of 1.4. Statistically, there is no tendency for the precipitate morphology to move away from a spherical shape and approach a cuboidal one, whereas extended secondary particles with an increased aspect ratio are found sporadically. To identify the (ω1, ω2) values representing the morphology of the secondary γ′ precipitates for the oil-quenched alloy, the median of the 38 plots was determined, and it is shown in Fig. 9 as a solid square; it was (ω1, ω2) = (12.4, 156.9), which is very close to the domain maxima (ω1, ω2) = (4π, 16π2), revealing a true circular shape. On the average, the secondary γ′ particles in the oil-quenched alloy remained a spherical morphology during continuous cooling.
Absolute moment invariant (ω1, ω2) plots of the secondary γ′ particles observed in the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by oil quenching, built from 38 data. The solid square indicates the median of all data, with ω1 = 12.4 and ω2 = 156.9.
The (ω1, ω2) results for the secondary γ′ particles in the furnace-cooled alloy, with a slow cooling rate, are shown in Fig. 10. The number of data plots was 19, which is smaller than that for the oil-quenched alloy because the size of each secondary γ′ precipitate was comparatively large. The wide distribution of the plots is evident for the furnace-cooled alloy, compared with the oil-quenched one. In this case, d was about 500 nm, as shown in Fig. 6(c), which is sufficiently larger compared to a typical TEM film thickness. In the octodendritic shape case, the shape of the secondary γ′ particles depends on the cut surface, resulting in the widespread distribution of the plots observed in the figure. Each plot lies near the right end of the domain, which means that the aspect ratio of each precipitate is close to unity. Furthermore, more than half of the plots were located in the ω2 < 144 region, indicating that a high portion of the secondary γ′ particles exhibited an octodendritic shape. The median obtained from the 19 plots was (ω1, ω2) = (11.3, 133.5). In the furnace-cooled alloy, the secondary γ′ precipitates are considered to evolve from a spherical to a cuboidal and, further, to an octodendritic shape during continuous cooling, and the aspect ratio of the precipitates remains around unity.
Absolute moment invariant (ω1, ω2) plots of the secondary γ′ particles observed in the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by furnace cooling, built from 19 data. The solid square indicates the median of all data, with ω1 = 11.3 and ω2 = 133.5.
The ω1 value was firstly summarized as a function of the cooling rate to evaluate the morphology of the secondary γ′ particles for Udimet 720Li specimens continuously cooled with four cooling rates after the solution treatment at 1473 K for 1 h, and the results are shown in Fig. 11. The ω1 value was 12.4, which is close to the upper limit of the domain (4π), when the cooling rate was as high as that for WQ and OQ and decreased continuously with the decrease of the cooling rate, reaching 11.4 for the furnace-cooled alloy. As discussed in Section 2.2, the ω1 value decreased not only with the aspect ratio increase but also with the transition of the particle morphology from a spherical to a cuboidal and/or an octodendritic shape. In this study, almost no increase in the aspect ratio of the secondary γ′ precipitates was observed either in the oil-quenched alloy with a relatively high cooling rate or in the furnace-cooled one with a low cooling rate. The decrease in the ω1 value with the cooling rate decrease, as observed in Fig. 11, could be due to the transition of the precipitate morphology from a spherical to a cuboidal and/or an octodendritic shape.
Absolute moment invariant ω1 values for the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h as a function of the cooling rate.
The ω2 values, which represent the morphology of the precipitates and do not depend on the aspect ratio, are shown as a function of the cooling rate in Fig. 12 as open symbols; the morphological illustration of the secondary γ′ particles corresponding to each ω2 value with the aspect ratio of unity is also included. The ω2 values were 16π2 and 144 in the cases of a true circular and a cuboidal shape, respectively, and the trend of an octodendritic morphology was emphasized as it decreased below 144. The ω2 value was close to the upper limit of the domain (16π2) when the cooling rate was high, as in the case of the water-quenched and oil-quenched alloys. It dramatically decreased with the decrease of the cooling rate, typically below 5 × 100 K/s, becoming 134 in the furnace-cooled alloy. As the cooling rate decreased from WQ to OQ and AC, the morphology of the secondary γ′ particles evolved from a spherical to a cuboidal shape with an increased frequency of a smooth interface, and the octodendrite shape became visible for the furnace-cooled alloy, which had the lowest cooling rate.
Absolute moment invariant ω2 values for the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h (open) and for those aged at 1173 K for 100 h (solid) after the solution treatment.
In the previous sections, the following two points were identified with respect to the secondary γ′ particle size and morphology after the solution treatment at 1473 K for 1 h of the Ni-based superalloy Udimet 720Li followed by continuously cooling by four different methods. First, the size continuously increases with the decrease of the cooling rate so that the d ∝ v−0.4 relationship is established between them. Second, the morphology evolves from a spherical to a cuboidal and, further, to an octodendritic shape with the decrease of the cooling rate, and the aspect ratio does not increase. On the basis of the results described above, in this section, the aging treatment is performed at 1173 K to understand the associated evolution of the secondary γ′ particles for the solution-treated alloys that have been continuously cooled by each cooling method.
The FE-SEM images of the oil-quenched and furnace-cooled alloys subjected to the solution treatment at 1473 K for 1 h and then to aging treatment at 1173 K for 100 h are shown in Fig. 13. For the aged oil-quenched alloy, the d value was about 200 nm, and most precipitates exhibited a spherical morphology (Fig. 13(a)); in addition, particles with an increased aspect ratio resulting from the aggregation of adjacent secondary γ′ precipitates were partially detected. On the contrary, for the aged furnace-cooled alloy, which had a slow cooling rate, the fine tertiary γ′ particles detected before the aging treatment disappeared, the shape of the secondary γ′ particles was not octodendritic but spherical or cuboidal (Fig. 13(b)), and no aggregation of the secondary γ′ precipitates was observed.
Field-emission scanning electron microscopy images of Udimet 720Li specimens aged at 1173 K for 100 h after the solution treatment at 1473 K for 1 h followed by (a) oil quenching and (b) furnace cooling.
The d values of the oil-quenched and furnace-cooled alloys after the solution and aging treatments are summarized as a function of the aging time in Fig. 14. In the oil-quenched alloy, d linearly increased when increasing the aging time above 1 h, and the gradient was 0.3, which suggests that it increased along the Ostwald ripening during the aging treatment.26) On the contrary, d hardly increased during the first 30 h of the aging treatment for the furnace-cooled alloy and slightly increased afterward.
Change in the diameter of the secondary γ′ precipitates during isothermal aging at 1173 K for the Udimet 720Li specimens after the solution treatment at 1473 K for 1 h followed by oil quenching (OQ) and furnace cooling (FC).
The median of the ω2 value, which directly reflects the precipitate morphology, was obtained from the (ω1, ω2) distributions for the oil-quenched and furnace-cooled alloys subjected to the solution treatment followed by the aging one. The results are shown in Fig. 12 as solid symbols. The ω2 value was 155.6 for the oil-quenched alloy, with a relatively high cooling rate, which is almost the same as that of the same alloy after the solution treatment and before aging. For the oil-quenched alloy, d increased continuously along the Ostwald ripening with remaining the spherical morphology during the aging treatment. By contrast, the ω2 value was 149.8 for the furnace-cooled alloy after the aging treatment, which is apparently larger than that observed before aging. In this case, the secondary γ′ particles exhibited an octodendritic shape after the solution treatment (Fig. 6(c)), and a size reduction due to their splitting and a size increase due to the coarsening of the split particles may occur in parallel. Hence, for the furnace-cooled alloy, the morphology of the secondary γ′ precipitates is considered to evolve to a spherical shape from an octodendritic one, and d is hardly increased by the aging treatment at 1173 K.
In this study, we first investigated the effect of the cooling rate after the solution treatment on the size and morphology of γ′ precipitates for the γ/γ′ two-phase wrought Ni-based superalloy Udimet 720Li having a high γ′ volume fraction (>45%) and a low γ/γ′ phase misfit. Then, the effect of the aging treatment on their size and morphology was also investigated. The following results were obtained.
The authors would like to thank Prof. Susumu Onaka and Prof. Yoshisato Kimura of Tokyo Institute of Technology and Prof. Yoji Miyajima of Kanazawa University for kind help in microstructure observation using electron microscopy.
