Investigation on Tensile Behavior of Solid Solution-Strengthened Ni-Co-Cr-Based Superalloy During Long-Term Aging

Abstract

This study investigated how long-term aging (750 °C and 950 °C) affects the microstructure and room-temperature tensile properties of the Ni-Co-Cr superalloy GH3617. Characterization (SEM, EDS, EBSD) showed that initial aging (750 °C, 500 h) formed discontinuous M₂₃C₆ carbides, pinning grain boundaries and improving strength. Prolonged aging (750 °C, 5000 h) caused M₂₃C₆ to coarsen into brittle chain-like structures (width up to 1.244 μm) and precipitated M₆C carbides, degrading grain boundaries. Aging at 950 °C accelerated this coarsening via LSW kinetics (rate constant: 6.83 × 10⁻² μm³/s), with Mo segregation promoting M₆C formation. Tensile properties resulted from competing γ′ precipitation strengthening (post-aging strength increased up to 23.3%) and grain boundary degradation (elongation dropped from 70.1% to 43.3%). Fracture shifted from purely intergranular (cracks along M₂₃C₆/γ interfaces at 750 °C) to mixed mode (cracks initiated by M₆C fragmentation at 950 °C). These insights support superalloy microstructure optimization and lifetime prediction.

1. Introduction

Nickel-based superalloys have become critical materials for high-temperature components in aerospace and energy systems due to their exceptional strength at elevated temperatures, oxidation resistance, corrosion resistance, and creep properties [1]. With the continuously increasing demands on operating temperatures and service lifetimes of advanced gas turbines and aeroengines, research on microstructural stability and performance degradation mechanisms of nickel-based superalloys under extreme environments has gained significant attention [2]. Among these alloys, GH3617, a representative solid solution-strengthened Ni-Co-Cr-based superalloy, is designed for long-term service at 750 °C with a maximum operating temperature of 950 °C. This capability positions it as an ideal material for high-temperature applications, including aircraft engines, power generation plants, and industrial gas turbines.

In addition to solid solution-strengthening effects, the presence of secondary phase particles such as γ′ precipitates, M₂₃C₆ carbides, M₆C carbides, and MX complex phases also contributes to the strength enhancement of Alloy IN617 [3,4,5,6]. The creep resistance and tensile performance of Alloy IN617 at elevated temperatures are critically dependent on the thermal stability of these precipitates [7,8].

However, during prolonged high-temperature aging (600–1000 °C), dynamic microstructural evolution occurs in superalloys, including grain boundary carbide precipitation, γ′ phase transformation, and dislocation substructure reorganization. These changes induce the gradual degradation of mechanical properties (creep resistance, fatigue life) and oxidation resistance, directly compromising component reliability and service lifespan [9]. Zhang et al. [10] demonstrated that the precipitation behavior of γ′ phase and carbides in a GH4175 alloy can be modulated via sub-solvus heat treatment, thereby enhancing high-temperature mechanical performance. Gao et al. [11] systematically investigated the microstructural evolution and mechanical property variations of an as-cast Alloy IN617B nickel-based superalloy subjected to long-term aging at 700 °C for up to 10,000 h following solution treatment at 1200 °C. Their findings revealed that the casting process significantly influenced carbide morphology, while the synergistic interaction mechanisms between γ′ precipitates and M₂₃C₆ carbides (involving elemental diffusion constraints and nucleation promotion) served as critical factors in maintaining microstructural stability. In the present study, a standardized solution treatment was introduced during the fabrication of a GH3617 superalloy, achieving an optimized design of the initial microstructure through precise control of phase transformation pathways.

Although existing studies have revealed the microstructural stability of IN617B during long-term aging at 700 °C, systematic experimental evidence on the evolution of grain boundary carbides in GH3617 at high service temperatures (750–950 °C)—particularly the kinetics of the M₂₃C₆-to-M₆C phase transition and its direct impact on room-temperature tensile fracture behavior—remains lacking. This study systematically investigates the long-term aging and tensile behavior of a GH3617 Ni-based alloy under varying temperatures and aging durations. By quantitatively characterizing the dynamic evolution of grain boundary carbides during aging at 750–950 °C, establishing a carbide coarsening kinetic model, and determining coarsening rate constants, we reveal the temperature-dependent evolution of fracture modes. Combined with microstructural characterization (SEM, EDS, EBSD), we elucidate the alloy’s microstructural evolution and its influence on tensile behavior. Our findings provide new insights for the interfacial engineering of superalloys, aiming to enhance the service reliability of GH3617 in high-temperature components such as aeroengines through optimized aging.

2. Experimental Procedures

This study comprehensively investigates the microstructural evolution, mechanistic variations, room-temperature tensile deformation, and fracture behavior of a GH3617 nickel-based superalloy under long-term aging at varying temperatures and durations. A detailed analysis of phase transformation kinetics, dislocation substructure reorganization, and crack propagation mechanisms was conducted, with specific alloy composition parameters provided in

Table 1. Chemical composition of GH3617 (wt%).

Elements	Cr	Co	Mo	Al	Ti	C	Ni
wt%	21.9	11.8	8.71	0.98	0.41	0.058	Bal.

.

In this study, all GH3617 alloys used in the experiment were directly smelted by vacuum induction melting and cast into ingots. After ingot casting, the material was subjected to cogging and forging to obtain 35 mm × 35 mm bar billets. Subsequently, multi-pass hot rolling was performed on the forged billets to produce 16 mm-diameter hot-rolled bars, and the alloy was subjected to a standard solution heat treatment by holding at 1175 °C for 1 h in a furnace, followed by water quenching (WQ). Subsequent aging treatments were conducted at 750 °C and 950 °C, with durations of 500 h, 1000 h, 5000 h, and 10,000 h. The microstructural evolution of each specimen was characterized using scanning electron microscopy (SEM, TESCAN MAIA3, TESCAN (China) Co., Ltd., Shanghai, China) equipped with energy-dispersive spectroscopy (EDS, Ultim Max Infinity SDD, Oxford, UK) and electron backscatter diffraction (EBSD, Symmetry S3, Oxford, UK). Before SEM observation, all samples were ground with 2000-grit SiC abrasive paper and mechanically polished. To reveal carbide distributions and longitudinal sections of tensile specimens, electrochemical etching was performed at room temperature using a solution containing 5 g CuCl₂ + 100 mL HCl + 100 mL C₂H₅OH. Tensile tests were carried out at room temperature with a constant strain rate of 10⁻⁴ s⁻¹. Smooth cylindrical specimens with a total length of 70 mm were employed, featuring a gauge length of 25 mm and a diameter of 5 mm.

3. Results

3.1. Initial Microstructure After Solution Treatment

The microstructural characteristics of the GH3617 nickel-based superalloy subjected to standard solution treatment are illustrated in Figure 1. Figure 1a reveals abundant carbides precipitated along grain boundaries in the hot-rolled original microstructure. To address this, the specified heat treatment regime—solution heat treatment at 1175 °C for 1h followed by water quenching—was applied to obtain a fully austenitic structure. Optical microscopy observations (Figure 1b) revealed a typical equiaxed austenitic grain structure formed in the specimen after holding at 1175 °C for 1 h under the standard solution treatment protocol. The microstructure exhibited randomly distributed crystallographic orientations, with a high density of annealing twin boundaries prevalent within the grain boundary network (Figure 1c). Quantitative analysis via electron backscatter diffraction (EBSD) demonstrated that the average grain size of the solution-treated GH3617 was approximately 35 μm, while annealing twin boundaries accounted for 60.28% of the total grain boundary length.

As demonstrated by previous studies [12], the high proportion of annealing twins in the alloy is intrinsically linked to the relatively low stacking fault energy (SFE) characteristic of nickel-based superalloys. The reduced SFE significantly lowers the energy barrier for twin formation, facilitating the generation of abundant annealing twin structures through grain boundary migration-mediated stacking fault accumulation during thermal processing. Furthermore, no discernible elemental segregation bands were observed in the microstructure, indicating that the solution heat treatment effectively eliminated severe compositional inhomogeneity. This thermal protocol enabled sufficient diffusion of solute atoms from the original as-cast microstructure into the matrix γ phase, thereby eradicating the compositional segregation induced during casting or thermomechanical processing.

3.2. Carbide Evolution in GH3617 Alloy During Long-Term Aging Process

Studies have demonstrated that the performance of the superalloy GH3617 is predominantly governed by the microstructure of its secondary phases and carbides. During the long-term aging of the GH3617 alloy, the evolution of grain boundary carbides constitutes the primary mechanism responsible for the degradation of microstructural stability and mechanical properties. Prolonged aging induces significant coarsening of grain boundary carbides (Figure 2), driven by the progressive enrichment of carbide-forming elements at grain boundaries. To elucidate this phenomenon, electron probe microanalysis (EPMA) was conducted to characterize carbide types at grain boundaries after aging at 950 °C for 500 h and 5000 h.

As shown in Figure 3a, short-term aging (500 h) resulted in pronounced segregation of C and Cr at grain boundaries, facilitating the formation of M₂₃C₆ carbides, while Mo—a critical element for M₆C carbide formation—remained uniformly distributed within the matrix. By contrast, after 5000 h of aging (Figure 3b), substantial Mo enrichment was observed at grain boundaries, accompanied by localized substitution of Cr by Mo. These findings indicate that M₂₃C₆ carbide precipitation dominates at grain boundaries during the initial aging stage due to its lower thermodynamic barrier, whereas prolonged aging significantly enhances the driving force for M₆C carbide formation.

3.3. Room Temperature Tensile Properties Under Different Aging Conditions

To investigate the tensile property evolution of the GH3617 superalloy upon long-term aging, room-temperature tensile tests were systematically conducted on specimens subjected to aging at different temperatures and durations. The resulting tensile properties (yield strength, σ_0.2; tensile strength, σ_b; elongation at break, δ; reduction in area, ψ) after various aging regimes are summarized in Figure 4. The data on the room-temperature tensile properties (average (AVE), standard deviation (SD)) of the GH3617 alloy after aging treatment are shown in

Table 2. Room-temperature tensile properties of GH3617 alloy after aging treatment.

Aging Temperature	Aging Time	σ0.2 (MPa)			σb (MPa)			δ (%)			ψ (%)
		Data	AVG	SD	Data	AVG	SD	Data	AVG	SD	Data	AVG	SD
750 °C	500 h	458	477.67	20.53	928	940.33	11.67	41.6	41.7	0.54	40.3	40.2	0.14
		469			937			42.4			40.3
		506			956			41.1			40
	1000 h	435	460.67	25.82	916	925.67	12.28	44.7	43.03	2.08	42.5	40.77	1.68
		451			918			44.3			41.3
		496			943			40.1			38.5
	5000 h	487	457	22.20	947	929	13.37	41.8	43.13	1.20	44	43.8	0.28
		450			925			44.7			44
		434			915			42.9			43.4
	10,000 h	454	437.67	12.71	932	918.33	11.03	44.5	43.67	0.85	49	48	0.82
		436			918			44			48
		423			905			42.5			47
950 °C	500 h	305	302	2.16	776	772	2.94	51.9	51.4	0.64	38.8	39	0.14
		301			771			51.8			39.1
		300			769			50.5			39.1
	1000 h	304	303	1.41	775	775.33	0.47	52	53.3	0.93	43.4	36	1.31
		304			776			53.8			43.1
		301			775			54.1			42.5
	5000 h	295	291	2.94	750	742.33	5.56	46.2	45.3	1.42	36	36	1.31
		290			740			46.4			37.6
		288			737			43.3			34.4
	10,000 h	278	283.33	3.86	713	723	8.60	45.5	47.67	1.55	38	37.33	0.94
		285			722			48.5			36
		287			734			49			38

.

As illustrated in Figure 4a, the tensile strength during aging at 750 °C increased from 762.3 MPa (unaged) to 940 MPa after 500 h, and then to 918 MPa after 10,000 h of aging, representing a 23.3% enhancement. During aging at 950 °C (Figure 4b), the tensile strength exhibited a modest increment from 762.3 MPa (unaged) to 772 MPa after 500 h aging, and then to 723 MPa after 10,000 h, corresponding to a 1.3% improvement. Both scenarios displayed a characteristic trend of initially increasing (750–950 °C, 0–500 h), and then slightly declining (750–950 °C, 500–10,000 h) before ultimately stabilizing.

Concurrently, the elongation to failure decreased from 70.1% (unaged) to 43.7% after 10,000 h of aging at 750 °C (Figure 4c), equivalent to a 37.7% reduction. Similarly, aging at 950 °C reduced the elongation from 70.1% (unaged) to 44.5% after 10,000 h (Figure 4d), showing a 36.5% decline. These ductility variations consistently followed a two-stage pattern: an initial rapid decline was succeeded by gradual stabilization.

3.4. Fracture Behavior After 5000 h of Aging

Figure 5 illustrates the evolution of the room-temperature tensile fracture morphology of the GH3617 alloy subjected to 5000 h of aging treatment at different temperatures (750 °C and 950 °C). Microstructural analysis revealed that the fracture modes under both aging conditions exhibited intergranular–transgranular mixed characteristics, as evidenced by the complex morphological features within the region enclosed by the white dashed line in Figure 5c.

4. Discussion

4.1. Evolution Behavior of Grain Boundary Carbides During Long-Term Aging

Figure 2 illustrates the grain boundary microstructure evolution of the GH3617 alloy under different aging temperatures and durations. As revealed in the figure, after 500 h of aging at 750 °C, a small quantity of serrated and discontinuous M₂₃C₆ carbides precipitated along grain boundaries, while fine granular M₂₃C₆ carbides dispersed within grains. Upon extending the aging time to 1000 h, an observable increase in intragranular carbide density was detected, with the distribution density of grain boundary M₂₃C₆ carbides remaining stable but exhibiting localized coarsening. Following prolonged aging for 5000 h, significant microstructural evolution occurred: blocky M₆C carbides formed and grew at grain boundaries, displaying an inhomogeneous distribution both at boundaries and within grains. Partial M₂₃C₆ carbides at these locations underwent dissolution, while the remaining undissolved carbides experienced substantial coarsening, transforming from discrete serrated morphologies to continuous chain-like configurations. No discernible variations in carbide type or size were observed between aging treatments of 5000 h and 10,000 h.

Notably, during short-term aging (500–1000 h) at 750 °C, the discontinuous distribution of both grain boundary carbides and intragranular particles effectively enhanced the material’s comprehensive properties, particularly creep resistance, through grain boundary migration pinning and grain boundary sliding inhibition. This mechanism contributed to excellent microstructural stability. Under 950 °C aging conditions, the carbide evolution in the GH3617 alloy demonstrated distinct characteristics. As shown in the figure, semi-continuous chain-like M₂₃C₆ carbides formed along grain boundaries after 500 h of aging. Significant coarsening occurred at grain boundaries during aging, while the partial dissolution of slender rod-shaped and blocky M₂₃C₆ carbides within grains was observed. The residual rod-shaped carbides subsequently underwent further coarsening. Notably, the continuous distribution of coarse short-rod structures was found to detrimentally affect material performance. Comparative analysis with Figure 3 revealed positive correlations between aging time and both volume fraction and average diameter of the M₆C carbides. Previous studies [13,14,15] have demonstrated that compared with the M₂₃C₆ phase, M₆C carbides exhibit superior lattice stability during long-term high-temperature service. The continuous grain boundary distribution of this characteristic phase enhances creep resistance near service temperatures through the Orowan strengthening mechanism. This interface engineering strategy provides crucial theoretical guidance for microstructure optimization in superalloy systems.

Figure 6 illustrates the dependence of grain boundary carbide width in the GH3617 alloy on aging temperature and duration during long-term aging. As shown in Figure 6a, the average carbide width initially increases with extended aging time and elevated temperature, eventually stabilizing with a maximum value of 1.244 μm. Studies [16] have demonstrated that the growth kinetics of M₂₃C₆ carbides exhibit accelerated progression during the early aging stages, followed by gradual deceleration as aging proceeds. A positive correlation between aging temperature and carbide coarsening rate was observed. This behavior originates from the significant compositional gradient between M₂₃C₆ carbides and the matrix γ during the initial non-equilibrium stage. Cr progressively segregates toward grain boundaries through bulk diffusion and grain boundary diffusion mechanisms, driving carbide growth. As the system approaches thermodynamic equilibrium, the Cr concentration gradient between the γ matrix and the M₂₃C₆ carbides diminishes, resulting in reduced interfacial elemental exchange rates that effectively suppress further carbide coarsening. As revealed in Figure 6b, a linear relationship exists between the cube of the average carbide width (d³) and aging time (t), consistent with the LSW theory governing coarsening kinetics (d³ ∝ t). The slope of this linear correlation corresponds to the coarsening rate constant, calculated as 6.83 × 10⁻² μm³/s. This quantitative agreement with LSW theory confirms the dominance of interfacial reaction-controlled Ostwald ripening in carbide evolution under the given aging conditions.

4.2. Room-Temperature Tensile Properties of GH3617 Alloy After Long-Term Aging

With increasing aging temperature and prolonged aging duration, the material’s strength demonstrates an initial enhancement followed by stabilization. Conversely, the alloy’s elongation first decreases before reaching a steady state. This behavior originates from precipitation evolution mechanisms: During the initial aging stages, supersaturated Ti and Al atoms undergo diffusional clustering to form fine γ′ precipitates. These nanoscale particles distribute uniformly within the matrix, strengthening the material through dislocation obstruction mechanisms (Orowan bypass or dislocation shearing). The homogeneous precipitation distribution equilibrates dislocation slip resistance, mitigates localized stress concentration, and delays necking initiation during tensile deformation, collectively contributing to relatively high ductility.

As aging progresses with extended time and elevated temperatures, Ostwald ripening dominates, inducing γ′ precipitate coarsening. This microstructural evolution increases both particle size and inter-precipitate spacing. Larger precipitates require elevated stress levels for dislocation bypass while suppressing shear mechanisms, consequently increasing material brittleness and reducing elongation. Concurrently, high-temperature aging promotes the grain boundary segregation of Cr and Mo elements, forming continuous or chain-like M₂₃C₆ carbides. These grain boundary phases weaken intergranular cohesion and provide preferential crack propagation paths, further exacerbating the reduction in ductility.

4.3. Analysis of Room-Temperature Tensile Fracture Morphology of GH3617 Alloy After Long-Term Aging

To systematically investigate post-service fracture behavior, a metallographic examination of fracture longitudinal sections becomes imperative. Figure 7 illustrates the microstructural evolution characteristics in longitudinal profiles of the GH3617 alloy subjected to 5000 h of aging treatment followed by room-temperature tensile fracture. The damage nucleation predominantly originates at secondary phase carbides (Mo-, Cr-, and Ti-rich) and grain boundaries, with carbide-induced damage demonstrating preferential elongation along the tensile axis. During aging at 750 °C, weakened grain boundary cohesion drives crack initiation and propagation through a dominant intergranular fracture mechanism. At elevated temperatures (950 °C), intensified microcrack density accompanies fragmented carbides along grain boundaries. EDS mapping analysis (Figure 8) confirms coexisting M₂₃C₆ and M₆C carbide variants, where carbide clustering induces localized stress concentration. This stress amplification exceeds the critical load-bearing capacity of boundaries under applied loads, triggering preferential microvoid nucleation at both grain boundaries and M₂₃C₆/γ phase interfaces. Subsequent plastic instability facilitates microvoid coalescence, ultimately propagating to form macroscopic cracks.

Figure 9 schematically illustrates the grain boundary carbide evolution and tensile crack propagation mechanisms at 750 °C and 950 °C. During low-temperature short-term aging, discontinuous granular M₂₃C₆ carbides nucleate at grain boundaries due to C/Cr segregation, while Mo remains homogeneously distributed in the matrix. Prolonged aging of 5000 h induces M₂₃C₆ coarsening with partial dissolution, accompanied by Mo diffusion-driven segregation at grain boundaries that promotes limited M₆C carbide formation. During high-temperature short-term aging (500 h), accelerated thermal activation enables continuous chain-like M₂₃C₆ precipitation along grain boundaries. Extended aging duration triggers a three-stage evolution: (1) initial chain-like carbides transform into coarsened rod-shaped morphologies, (2) partial M₂₃C₆ dissolution releases Cr into the matrix, and (3) Mo-dominated solute redistribution facilitates progressive replacement by M₆C carbides through coupled dissolution–precipitation kinetics. Both 750 °C and 950 °C aging treatments at a duration of 5000 h exhibit analogous mixed intergranular–transgranular fracture modes. At 750 °C, weakened grain boundary strength promotes secondary intergranular cracks as dominant failure origins. These cracks nucleate preferentially at grain boundaries and propagate along boundaries under increasing applied stress. By contrast, aging at 950 °C induces M₂₃C₆ carbide coarsening, generating localized stress concentrations that trigger carbide fragmentation. This bears similarity to the research of S. Gao et al. and Sai Rajeshwari K. et al. [11,17]. This process initiates microcracks at grain boundary carbides, which subsequently propagate through carbide/matrix interfacial decohesion and coalesce into macrocracks under sustained loading.

As depicted in Figure 10, the austenitic grains in the GH3617 alloy exhibit pronounced plastic deformation during room-temperature tensile testing compared to the as-solutionized microstructure, showing elongation along the tensile axis. A significant fraction of grains undergo crystallographic reorientation, transitioning from random orientations to a predominant blue coloration in inverse pole figure (IPF) maps, indicative of lattice rotation toward the <111> direction. Quantitative texture analysis reveals maximum orientation distribution function intensities of 7.89 and 6.62 for specimens aged at 750 °C and 950 °C, respectively, confirming the development of a characteristic tensile fiber texture [18,19,20]. This texture evolution aligns with dominant {111}<110> slip system activation, consistent with typical deformation mechanisms in nickel-based superalloys, as corroborated by Qu et al. [21]. Furthermore, deformation twins are observed within the γ matrix, contributing to enhanced strength relative to non-aged conditions. Figure 10a,b demonstrate that the majority of grains exhibit <111>∥RD orientations. With increasing aging temperature, progressive replacement by <001>-oriented grains occurs, a phenomenon attributed to strain energy minimization in FCC materials, as validated by Zhang et al. [22]. This grain rotation mechanism accommodates severe plastic strain near grain boundary (GB) regions, thereby stabilizing ductility at elevated temperatures. Concurrently, extensive intergranular cracking is evident at GBs, as manifested by microcrack networks along carbide-depleted zones.

Figure 10c,d present kernel average misorientation (KAM) maps of post-tensile fracture surfaces, where the KAM magnitude quantitatively correlates with localized strain distribution and geometrically necessary dislocation (GND) density. Both specimen groups exhibit analogous KAM patterns: elevated misorientation values concentrate near GBs and twin boundaries, while lower values prevail within grain interiors. This spatial heterogeneity confirms significant stress localization at GB regions, thereby rationalizing the dominance of intergranular crack initiation and propagation mechanisms in both conditions. The comparable localized strain profiles further indicate that aging temperature (750 °C vs. 950 °C) does not substantially alter dislocation density or lattice distortion magnitudes, preserving consistent fracture modality across thermal exposure conditions.

Notably, Wright et al. [23] established that KAM distributions in deformed materials exhibit direct correspondence with GND configurations, which serve as microstructural proxies for strain partitioning. The observed KAM homogeneity between specimens demonstrates that aging temperature exerts a negligible influence on strain distribution uniformity during room-temperature tensile deformation, maintaining equivalent work-hardening capacity regardless of thermal history.

5. Conclusions

This study systematically elucidates the microstructural evolution mechanisms and their impact on the room-temperature tensile properties of solution-strengthened GH3617 nickel-based superalloys during long-term aging at 750–950 °C. The principal conclusions are summarized as follows:

• Grain boundary carbides exhibit temperature–time coupled dynamic evolution characteristics, where carbide morphology and phase selection are synergistically governed by aging parameters. Serrated M₂₃C₆ carbides formed during early-stage aging at 750 °C (500 h) enhance creep resistance via effective grain boundary pinning. Prolonged aging (5000 h) induces M₆C carbide coarsening into continuous chain-like structures (max. width of 1.244 μm), progressively deteriorating boundary cohesion. Accelerated coarsening kinetics at 950 °C are driven by Mo boundary segregation and obey LSW theory (d³ ∝ t), exhibiting a coarsening rate constant of 6.83 × 10⁻² μm³/s.
• The evolution of room-temperature tensile properties arises from the competitive mechanism between γ′ precipitate strengthening and grain boundary weakening. Post aging, the alloy achieves a maximum strength enhancement of 23.3% (940.3 MPa at 750 °C/500 h), which is attributed to nano-sized γ′ precipitates reinforcing the matrix through Orowan looping. Conversely, the elongation reduction (70.1% → 43.3%) correlates strongly with deteriorated plasticity caused by GB carbide coarsening and γ′ spacing enlargement (via Ostwald ripening, quantified by the LSW coarsening exponent n = 3). Fracture mode transitions exhibit thermal history dependence.
• The fracture mode demonstrates temperature-dependent evolutionary characteristics. Following aging at 750 °C, intergranular fractures dominate, with cracks nucleating preferentially at M₂₃C₆/γ interfaces. During aging at 950 °C, mixed intergranular–transgranular fractures develop, where microcrack propagation initiated by fragmented M₆C carbides serves as the primary failure mechanism. Tensile deformation-induced <111>-oriented texture formation, coupled with elevated KAM values at grain boundary regions, confirms that stress concentration at the boundaries constitutes the preferential crack nucleation sites.

The focus of subsequent work is to extend the evolution laws of grain boundary carbides and tensile fracture cracks during long-term aging, originally established for Ni-Cr-Mo-based superalloys, to all superalloys. This aims to unify the grain boundary carbide evolution laws across superalloys during long-term aging. Concurrently, the effects of trace element additions (e.g., Hf/Ta) on the interfacial bonding strength between carbides and the γ matrix will be investigated to suppress the formation of chain-like carbides.