Effect of Heat Treatment Schedules on Creep Performance of Ni-Based Superalloy Mar-M247 at 871 °C and 250 Mpa

Abstract

The effects of heat treatment (H1 and H2) on the creep behavior and microstructures of Mar-M247 at 871 °C/250 MPa are studied. The results show that the as-cast microstructure is composed of eutectics, γ phase, γ’ phase, MC and M₂₃C₆ carbides, while new M₆C appears in heat-treated microstructures, indicating the transformation of carbides after heat treatments. The Mar-M247 is excellent, with over 1500 h of creep life, and H1 is 48% higher than H2. The addition of post-brazing and diffusion heat treatment in H2 is detrimental to creep resistance; the two steps promoted the transformation of MC into M₂₃C₆ in advance. The increase and coarsening of M₂₃C₆ would consume more γ-phase-forming elements, weakening the solution strengthening at grain boundaries. As a result, the resistance of the grain boundary and γ/γ’ interface to dislocation motion is significantly reduced, leading to the cracks’ initiation and propagation along the grain boundaries.

1. Introduction

With excellent mechanical properties, corrosion resistance and oxidation resistance at high temperature (~800 to 1300 K), Ni-based superalloys are applied in hot end parts of aeroengines and industrial gas turbines, such as integral wheels, turbine blades and disc rotors [1,2,3,4]. The service time of superalloys in turbine blades of gas turbines is longer compared with aeroengines, up to tens of thousands of hours, and the service environment is in high temperature and stress for higher efficiency, which will test the creep resistance and microstructure stability of the alloy [5,6]. Mar-M247, a typical polycrystalline casting superalloy, strengthened by about a 60% volume fraction of L1₂ γ’ phase precipitates, has been widely applied in high-pressure movable blades and stationary blades of F-class gas turbines due to its excellent comprehensive high-temperature performance and castability [7,8].

Mar-M247 contains many refractory elements, such as W, Mo, Ta and Hf, resulting in severe composition segregation between the dendrite core (DC) and interdendritic region (ID) due to the different solidification rate of each component, which is challenging for subsequent heat treatment. Meanwhile, γ/γ’ eutectics with a low melting point are easily precipitated in ID, which leads to a decrease of strength of the local area and the origin of microcracks [9]. By adjusting the morphology and size of γ’ and the distribution of carbides, reducing the composition segregation and eliminating the eutectic, an appropriate heat treatment schedule is the key to improving high-temperature properties of superalloys [9,10,11]. Generally, the heat treatment process of superalloys includes two parts, solution treatment and aging treatment, which have gained a lot of attention. The purpose of the solution heat treatment is to remelt the coarsening γ’ phase and eliminate low melting point eutectics completely or partially, and then the γ’ phase homogeneously precipitates with smaller particles in the subsequent cooling process. A higher solution heat treatment temperature is beneficial to eliminate eutectics; a previous study showed that the area fractions of eutectics had decreased from 8.4% to 4.3% after the temperature increased from 1120 °C to 1205 °C [12]. Moreover, the massive MC carbides also decompose into fine MC during solution heat treatment, or transform into smaller secondary carbides, such as M₂₃C₆ and M₆C, thus improving the strength of grain boundaries (GBs) [13,14]. The aging heat treatment is used to adjust the content and morphology of γ’ phase, which is an important parameter to determine service life. Hot isostatic pressing (HIP), post-brazing heat treatment and diffusion heat treatment are also commonly applied in superalloys, in addition to solution and aging heat treatment. Many research studies have shown that HIP can not only eliminate the internal porosity and shrinkage caused by casting, but also further homogenize the composition, thus improving the high-temperature performance of superalloys, especially for creep resistance [15,16,17]. For example, the research results of Bor et al. [7] showed that the area fraction of microporosity decreased from 3.16% to 0.25%, thus increasing the tensile strength of MAR-M247 alloy. The aim of post-brazing heat treatment is to remove the internal stress caused by welding, avoiding premature crack formation in the welding seam of the alloy, to prolong service life [18]. Previous papers on the Mar-M247 mainly focus on effects of heat treatment schedules on microstructures, hardness evolution and tensile properties [19]. Other studies were interested in the effects of elements additions on creep properties of the alloy, such as Re, Nb or Mg [20,21,22]. However, the effects of heat treatment on creep properties under actual working environment is rarely studied. The dorsal part of the first-stage rotor blade is directly subjected to 871 °C/250 MPa, which is a crucial property indicator for Mar-M247 superalloy in F-class gas turbines. Considering working conditions, an appropriate heat treatment schedule is key to the high-temperature performance of the superalloy. Therefore, the effects of two heat treatment schedules on creep performance and microstructures of Mar-M247 superalloy at 871 °C/250 MPa were studied in this paper.

2. Experimental Section

Cast into the dumbbell bar specimen by a vacuum induction melting furnace, the actual composition of Mar-M247 superalloy by inductively coupled plasma atomic emission spectra (ICP-AES) is listed in

Table 1. The actual composition of Mar-M247 superalloy (wt. %).

Elements	Cr	Co	W	Al	Ta	Hf	Ti	Mo	C	Zr	B	Ni
Content	8.35	9.96	10.02	5.56	3.15	1.48	1.00	0.65	0.14	0.03	0.02	Bal.

. The as-cast alloy is subjected to different heat treatment schedules, which are demonstrated in Figure 1. H1, the first heat treatment schedule, is HIP (1185 ± 5 °C + 4 h + air cooling, Ar atmosphere, 200 MPa) + solution heat treatment (1195 ± 5 °C + 2 h + air cooling, He atmosphere) + aging heat treatment (870 ± 5 °C + 20 h + air cooling). The other heat treatment schedule H2 is HIP (1185 ± 5 °C + 4 h + air cooling, Ar atmosphere, 200 MPa) + solution heat treatment (1195 ± 5 °C + 2 h + air cooling, He atmosphere) + post-brazing heat treatment (1120 ± 5 °C + 0.5 h + air cooling, He atmosphere) + diffusion heat treatment (1080 ± 5 °C + 3 h + air cooling, He atmosphere) + aging heat treatment (870 ± 5 °C + 20 h + air cooling). Obviously, post-brazing heat treatment and diffusion heat treatment are added in H2 (clearly shown in the yellow part of Figure 1b), compared to H1. Mar-M247 superalloys, heat-treated by H1 or H2, were defined as HS1 and HS2 alloys, respectively, for the convenience of expression. The formulations of heat treatment schedules are based on the previous work, working conditions of alloy, and CALPHAD method. An important characteristic temperature, the precipitation of γ’ phase is about 1110 °C [19]. The temperatures of HIP and solution heat treatment are higher to eliminate eutectics and obtain more homogeneous elemental distribution.

The as-cast Mar-M247, heat-treated HS1 and HS2 alloys were cut, inserted, ground and polished, subsequently, and then they were etched by HF:HNO₃:C₃H₈O = 1:2:3. The grain size was measured by determining the mean linear intercept of the grains using a Nikon SMZ800N optical microscope (OM). A Quanta Feg650 scanning electron microscope (SEM) was used to observe the microstructures at 15 kV, and the elements distribution was analyzed by an Oxford x-max energy dispersive spectrometer (EDS) on the SEM. According to ASTME21, HS1 and HS2 were mechanically processed into standard creep samples, with a gauge diameter 5 mm and gauge length 25 mm, as shown in Figure 2. The direction of the red arrows represents the direction of loading. The creep tests were performed along the axial direction of the specimen. The RD-50 high-temperature creep testing machine was used to test the creep life of HS1 and HS2 at 871 °C/250 MPa, and two creep samples in the two heat treatment schedules were used to avoid accidental error. For studying the creep mechanism of the alloy further, thin sections with a thickness of 500 μm were cut in the near creep fracture area of the test rod perpendicular to the applied load direction. The thin foils of HS1 and HS2 were mechanically thinned to about 50 μm and then etched by 5% HClO₄ + 95% C₂H₆O to the final perforation at −30 °C by a twin-jet electropolishing device. An FEI Tecnai G2 F20 transmission electron microscope (TEM) was operated at 200 kV to observe microstructures of HS1 and HS2 after creep rupture.

3. Results and Discussion

3.1. As-Cast and Heat-Treated Microstructures

The corresponding macrostructures of the as-cast rod of Mar-M247 superalloy are shown in Figure 3. Both ends of the samples are clamping ends with a diameter of 15 mm, and the middle is a working section with a diameter of 8 mm. The ordinary superalloys castings were produced by traditional precision casting technology [23]. The grain morphology is a coarse equiaxed crystal with an average size of 0.6 ± 0.2 mm and a largest size of about 2 mm, as shown in Figure 4. Figure 4a,b illustrate as-cast microstructures of the alloy by the backscattered electron mode and second electron mode of SEM, respectively. There are obvious dendrite structures and composition segregation in the as-cast microstructure of the alloy, as shown in Figure 4a. The brighter region is ID due to the segregation of heavy elements, especially the enrichment of Ta and Hf. The darker is DC, segregated into Co, Cr, W and so on. From Figure 4b, matrix γ phase and precipitated strength γ’ phase constitute the main microstructures; rosette-like γ-γ’ eutectic and a large content of carbides with different shapes are distributed in the ingrain and GBs. The eutectics at GBs with a low melting point and severe elements segregation damage local strength. One of the purposes of HIP and solution heat treatment exceeding the dissolution temperature is to eliminate eutectics partly or completely. The morphologies of carbides are mainly bulk and script-like. The coarse bulk carbides are MC, enriched in Ta and Hf, while the small Cr-rich carbides precipitated at the GBs are identified as M₂₃C₆ by EDS.

The microstructures after heat treatment of HS1 and HS2 are shown in Figure 5. The obvious dendrite structure is alleviated, and most of the eutectic is eliminated by heat treatment schedules H1 or H2, comparing Figure 4 and Figure 5, indicating improvement of homogeneous microstructures. Nevertheless, the microstructures exhibit a discrepancy between HS1 and HS2 due to the different heat treatment schedules. The eutectic in HS1 is not completely removed, but its size and area fraction are significantly reduced, as shown in Figure 5b. Most of the eutectics are completely eliminated in HS2, as shown in Figure 5d. Apparently, post-brazing heat treatment and diffusion heat treatment have played a key role in the elimination of eutectic for Mar-M247. Moreover, dendrite microstructures of HS1 are more obvious compared with HS2, according to Figure 5a,c, indicating that the above two heat treatment steps further alleviate composition segregation.

The carbides in the microstructures of HS1 and HS2 are shown in Figure 6. A large amount of Cr-rich chain M₂₃C₆ and a small amount of W-rich block M₆C appeared at the GBs. M₆C is not found in as-cast Mar-M247, indicating that some of the primary MC has transformed during the heat treatment. There is still some massive or script-like MC carbide at the GBs, whose size significantly decreases after heat treatment. The massive MC carbide at GBs in Figure 6c significantly reduces compared with Figure 6a, indicating that schedule H2 is conducive to the transformation of MC carbide at GBs into other carbides, such as M₆C and M₂₃C₆ carbides. The compositions of MC, M₂₃C₆ and M₆C in Mar-M247 superalloy by SEM-EDS are listed in

Table 2. Composition of various carbides in MAR-M247 alloy (at. %).

Carbides	C	Al	Ti	Cr	Co	Ni	Hf	Mo	Ta	W
MC	72.26	0.14	4.98	0.40	0.53	2.56	7.8	0.31	9.33	1.69
M6C	62.64	0.48	1.05	4.92	3.83	9.88	0.52	1.08	2.58	13.22
M23C6	40.46	1.44	4.40	10.83	5.72	14.11	8.21	1.04	6.29	7.50

. The results show that M₂₃C₆ contains more solid solution strengthening elements, such as Cr, Co, Mo and W. However, MC consumed more precipitation strengthening elements (such as Hf and Ta).

The area fraction of massive carbide MC in as-cast, heat-treated and after-creeped alloys are illustrated in Figure 7 after data statistics. The area fraction of massive MC reduces significantly from 6.7% to 4.9% or 3.3% after heat treatment schedules H1 or H2, respectively. This suggests that more primary coarse MC carbide undergoes the following transformation reactions [24,25]:

$$\mathrm{M}\mathrm{C}+\mathsf{\gamma }\to {\mathrm{M}}_6\mathrm{C}\left({\mathrm{M}}_{23}{\mathrm{C}}_6\right)+{\mathsf{\gamma }}^{\prime }$$

(1)

The decomposed carbides, M₂₃C₆ and M₆C, usually form with discontinuous or granular precipitations at the grain boundaries, as show in Figure 6. These continuous distributions at the GBs will prevent grain boundary sliding during the creep of the alloy [26]. The main carbides at GBs is M₂₃C₆ rather than M₆C due to the content of W in Mar-M247. Previous research results showed that if the superalloys have a sufficiently high content of Mo + W (Mo + 1/2 W > 6%), the MC carbides in such alloys tend to be unstable and decompose into M₆C-type carbide, while Mo + 1/2 W = 5.66 wt. % in the Mar-M247 superalloy [27]. As a result, the MC mainly transformed to M₂₃C₆-type carbide after heat treatment or subsequent high-temperature creep.

Due to the high content of Mo plus W, the width of chain M₂₃C₆ in alloy HS1 is 200 to 300 nm. However, the shape of M₂₃C₆ in alloy HS2 is not only chain, but also bulk, with a size of 500–600 nm. The bulk M₂₃C₆ will be detrimental to local strength at the GBs. Grain boundary slip is hindered by M₂₃C₆ in favor of creep resistance, and the change of morphology of M₂₃C₆ from blocky to chain leads to an increase of temporary creep resistance and improvement of the plasticity of superalloys [27]. The precipitation temperature of carbide M₂₃C₆ ranges from 600 °C to 1100 °C. The post-brazing heat treatment and diffusion heat treatment in the H2 schedule are within this range, which promotes the transformation of MC to carbide M₂₃C₆.

Importantly, it should be emphasized that the casting pores are all eliminated in the two alloys with the help of HIP heat treatment. The presence of microporosity will harm the properties of Mar-M247, resulting in the scatter of creep performance [28]. The application of HIP heat treatment ensured the subsequent creep test results were not influenced by casting pores.

3.2. Creep Tests and Microstructures

Figure 8 illustrates the creep performance of HS1 and HS2 alloys under 871 °C/250 MPa. From Figure 8a, Mar-M247 superalloy displays excellent creep resistance under the working circumstances. The heat treatment schedule significantly affects creep life, and the creep life of HS1 is about 48% higher than that of HS2 alloy. The addition of post-brazing heat treatment and diffusion heat treatment completely eliminates eutectic and decreases composition segregation in HS2; meanwhile, it causes adverse effects on the alloy creep performance, the mechanism of which will be discussed in Section 3.3. In addition, there is little difference in the elongation after fracture and the reduction area of the two alloys, indicating that the two heat treatment steps have little effect on the plasticity of the alloy.

The creep fracture morphology of HS1 and HS2 alloys is shown in Figure 9. A typical plastic fracture is present in the two alloys, as shown in Figure 9a,c, that is, an intergranular creep fracture, which is the main creep fracture pattern for superalloys. Some fracture features, such as dimples and cleavage, are covered by oxides, which are still found in Figure 9c. As shown in Figure 9b,d, dense microcracks are distributed on the carbides. It can be considered that carbides are related to the initiation and propagation of microcracks. The distribution of the main carbides-forming elements at the creep fracture is shown in Figure 9e; the enrichment of elements, especially for C, Ta, Hf and W (see the red ellipse), indicates that the carbides are distributed along the microcracks.

The creep microstructures of the longitudinal section near the fracture of HS1 and HS2 alloys are shown in Figure 10. The weak area in the microstructures is the GBs after a long time of service at 871 °C/250 MPa, where many microcracks, originated from the coarse carbides, appear and expand. It has been found that there is massive MC in heat-treated microstructures in Section 3.1, while the size of MC (shown by the purple arrow) carbide decreases significantly after creep, as shown in Figure 10a,c; that is, decomposition and transformation of carbides occur in the creep of the alloys. The red ellipse in Figure 10a shows the decomposition of large MC into small MC, and MC transforms into M₂₃C₆ in the orange ellipse. The studies of Xie et al. [29] showed that two adjacent MCs tend to transform to M₂₃C₆ or M₆C under the joint action of strain energy and interfacial energy, which increased the creep resistance of the alloy.

The creep microcracks appear in the surroundings of M₂₃C₆ carbide, and M₂₃C₆ seriously coarsens, and the content increases after creep, its morphology changed from chain to bulk, and its size (width) increases from nanoscale to micron-sized. The width of M₂₃C₆ in alloy HS1 increases from 200~300 nm in heat-treated to 1.0~1.1 μm in post-creep microstructures. The width of M₂₃C₆ in HS2 widens from 500~600 nm in the heat treatment state to 1.2~1.3 μm after creep fracture. Equation (1), the transformation of MC to M₂₃C₆, occurs in the H1 and H2 heat treatment schedules, as shown in Section 3.1. The transformation process is significantly intensified in the creep at 871 °C/250 MPa. Moreover, HS2 alloy undergoes MC→ M₂₃C₆ before HS1 due to the addition of post-brazing heat treatment and diffusion heat treatment. Therefore, the size of M₂₃C₆ is larger in the heat-treated or creep state HS2 alloy than in HS1. The fine carbides pined, preventing GBs from sliding and migration at the GBs. However, M₂₃C₆ is unstable and prone to coarsening during creep at high temperature, comparing with M₆C, which creates disorder in the GBs and allows easy dislocation movement thought the GBs [30]. Furthermore, M₂₃C₆ contains a higher content of γ-phase-forming elements, such as Cr, Co and W. The coarsening and increase of M₂₃C₆ will consume more solution strengthening elements, which is also demonstrated by the finding of the γ-poor phase region around M₂₃C₆ (shown by the red arrow in Figure 10b). The solution strengthening effects at the GBs is significantly weakened, resulting in the initiation and expansion of microcracks. These changes caused by the carbides greatly affect the creep properties of the alloy.

In addition, γ’ phase severely coalesces and rafts after creep at 871 °C/250 MPa over 1500 h, as shown by the blue arrows in Figure 10b, which significantly decreases the hindering ability of dislocation movement through the γ/γ’ interface [31]. The coarsening of M₂₃C₆ consumes solid solution strengthening elements at the GBs, which further promotes the dislocation movement. The joint weakness of the γ/γ’ interface and the GBs results in the fracture of the alloys after service for a long time.

3.3. Creep Fracture Mechanism

The microstructures of HS1 and HS2 are studied by scanning transmission microscope (STEM), and TEM-EDS is used for carbides mapping; the results are shown in Figure 11. It shows that the microcracks appeared at the surroundings of M₂₃C₆ carbide, indicating that massive M₂₃C₆ facilitates the initiation and propagation of microcracks. Importantly, the coexistence of M₂₃C₆ and MC is found in Figure 11a, which certified the transformation of MC to M₂₃C₆. However, only M₂₃C₆ is found in Figure 11b, due to the ahead of the transformation process in HS2 compared with HS1 alloy. The content of M₂₃C₆ is greater and the size is larger, resulting in the creep fracture of HS2 alloys before HS1.

Figure 12 shows the relationship between carbides and dislocation in the two alloys. Abundant dislocation networks are gathered in the γ/γ’ interface, as shown in Figure 12a, which is vital to resisting dislocation movement [32]. Nevertheless, many super-dislocations cut into the γ’ interface, indicating the weakness of the γ/γ’ interface after creep microstructures evolution. In addition, M₂₃C₆ carbide with the size of about 1.5 μm is precipitated at the γ/γ’ interface in two alloys. Lager-sized M₂₃C₆ consumes γ-forming elements, like Cr, Mo and Co, resulting in the impoverishment of γ around M₂₃C₆ carbide. It weakens solid solution strengthening and accelerates the movement of dislocation thought the GBs, leading to creep from alloys finally. And more and lager-sized M₂₃C₆ precipitates in HS2 alloy, as shown in Figure 12b, which is more harmful to the creep performance. Bor et al. [19] also showed that the carbides, like M₂₃C₆ carbide, piled up dislocations and generated stress concentration at the GBs, thus the initiation and propagation of microcracks along the GBs. In addition, small amounts of M₆C carbide are precipitated, transformed from MC carbide, which has little effect on the properties of the alloy.

Figure 13 illustrates the creep fracture mechanism schematic diagram. The microstructures after creep fracture has three important changes compared with the heat-treated microstructure. (a) Massive MC decomposes to fine MC and transforms to M₂₃C₆, and chain M₂₃C₆ is coarsened significantly and bulk MC is precipitated. (b) The γ-poor regions are formed around bulk M₂₃C₆, especially at grain boundaries, which significantly reduced the strength of the GBs. (c) The spherical γ’ coalesces and rafts, which facilitates dislocation movement, especially at grain boundaries. These promote the initiation of cracks occurring at grain boundaries. Compared with HS1, M₂₃C₆ is coarser, and γ-poor regions are larger in HS2 alloy, which greatly reduces the grain boundary strength and, thus, shortens the creep life of the alloy.

4. Conclusions

The effects of heat treatment schedules on the creep performance and its microstructures of Ni-based superalloy Mar-M247 at 871 °C/250 MPa were studied. Conclusions could be summarized as follows:

(1) The as-cast microstructure of Mar-M247 alloy is composed of eutectic, γ phase, γ’ phase and carbides, in which carbide is mainly MC and a small amount of M₂₃C₆.

(2) Both heat treatment schedules H1 and H2 alleviate composition segregation, obvious dendrite structure and most eutectic. The addition of post-brazing heat treatment and diffusion heat treatment in H2 completely eliminates eutectic, while it also promotes the transformation of MC carbide to M₂₃C₆ or M₆C.

(3) The creep life of Mar-M247 with the H1 heat treatment schedule is about 48% higher than that with the H2 schedule.

(4) The precipitation and coarsening of M₂₃C₆ at the GBs reduced the solid solution strengthening effect of γ-phase-forming elements, weakening the strength of the GBs of Mar-M247. The addition of post-brazing heat treatment and diffusion heat treatment promote the MC to M₂₃C₆ early, resulting in the decrease of the creep life of HS2 alloy.