Next Article in Journal
Corrosion Behavior upon Laser Surface Texturing AISI 430 Stainless Steel
Previous Article in Journal
Research Progress on Application of Machine Learning in Continuous Casting
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
Journals
Metals
Volume 15
Issue 12
10.3390/met15121385
metals-logo
Submit to this Journal Review for this Journal Propose a Special Issue
► Article Menu

Article Menu

share Share announcement Help format_quote Cite question_answer Discuss in SciProfiles
first_page
settings
Order Article Reprints
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Effect of Solidification Conditions on High-Cycle Fatigue Behavior in DD6 Single-Crystal Superalloy

by
Hongji Xie
*,
Yushi Luo
,
Yunsong Zhao
and
Zhenyu Yang
Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, China
*
Author to whom correspondence should be addressed.
Metals 2025, 15(12), 1385; https://doi.org/10.3390/met15121385
Submission received: 13 November 2025 / Revised: 3 December 2025 / Accepted: 8 December 2025 / Published: 17 December 2025
(This article belongs to the Section Metal Failure Analysis)
Browse Figures
Versions Notes

Abstract

This study investigates the influence of solidification conditions on the high-cycle fatigue (HCF) behavior of a second-generation DD6 single-crystal superalloy. Single-crystal bars with a [001] orientation were prepared using the high-rate solidification (HRS) and liquid-metal cooling (LMC) techniques under various pouring temperatures. The HCF performance of the heat-treated alloy was subsequently evaluated at 800 °C using rotary bending fatigue tests. The results demonstrate that increasing the pouring temperature effectively reduced the content and size of microporosity in the HRS alloys. At an identical pouring temperature, the LMC alloy exhibited a significant reduction in microporosity, with its content and maximum pore size being only 44.4% and 45.8% of those in the HRS alloy, respectively. Consequently, the HCF performance was enhanced with increasing pouring temperature for the HRS alloys. The LMC alloy outperformed its HRS counterpart processed at the same temperature, showing a 9.4% increase in the conditional fatigue limit (at 107 cycles). Microporosity was identified as the dominant site for HCF crack initiation at 800 °C. The role of γ/γ′ eutectic in crack initiation diminished or even vanished as the solidification conditions were optimized. Fractographic analysis revealed that the HCF fracture mechanism was quasi-cleavage, independent of the solidification conditions. Under a typical stress amplitude of 550 MPa, the deformation mechanism was characterized by the slip of a/2<011> dislocations within the γ matrix channels, which was also unaffected by the solidification conditions. In conclusion, optimizing solidification conditions, such as by increasing the pouring temperature or employing the LMC process, enhances the HCF performance of the DD6 alloy primarily by refining microporosity, which in turn prolongs the fatigue crack initiation life.

1. Introduction

Nickel-based single-crystal superalloys have become the preferred material for critical hot-section components in advanced aero-engines, such as turbine blades, owing to their excellent high-temperature strength, microstructural stability, and oxidation resistance. The performance of these alloys directly determines the thrust-to-weight ratio, efficiency, and service reliability of the engines. The service environment for modern turbine blades is exceptionally severe, involving simultaneous exposure to temperatures exceeding 1000 °C, complex aerodynamic loads, and high-frequency vibratory stresses. Under these conditions, high-cycle fatigue (HCF) has emerged as a predominant cause of early blade failure [1,2]. For instance, the in-service HCF failure of a DD10 single-crystal turbine blade, initiated by a film-cooling hole defect, underscores the critical influence of material HCF performance on component safety [3]. Therefore, systematically investigating the mechanisms for controlling the HCF performance of single-crystal superalloys is of urgent engineering necessity.
Solidification conditions constitute a critical process in the fabrication of single-crystal superalloy blades, as they govern the temperature gradient and thereby determine the microstructural characteristics of both the as-cast and heat-treated alloy [4]. Established research indicates that the dendrite arm spacing in superalloys is inversely related to solidification parameters such as temperature gradient (G) and solidification rate (R). Specifically, the primary dendrite arm spacing (λ1) is determined by G−1/2 × R−1/4 [5], while the secondary dendrite arm spacing (λ2) is determined by G−1/2 × R−1/2 [6]. Furthermore, studies [7] have shown that under a very high and nearly constant temperature gradient, increasing the solidification rate (V) refines the dendritic structure, with λ1 and λ2 exhibiting linear relationships with V−1/4 and V−1/2, respectively. For the widely used second-generation single-crystal superalloy DD6 in China, the LMC process has been demonstrated to significantly reduce dendrite arm spacing and γ/γ′ eutectic pool size, alleviate elemental micro-segregation, and improve the uniformity of γ′ precipitates after heat treatment compared to the HRS process [4]. These solidification-condition-dependent microstructural features, including dendrite morphology and γ/γ′ eutectic, have been proven decisive for the mechanical properties of the alloy.
Extensive research has been conducted worldwide on the relationship between microstructure and conventional mechanical properties in single-crystal superalloys [8,9,10,11,12,13]. However, in the realm of fatigue performance, existing studies have primarily focused on the effects of in-service parameters like temperature and stress [14]. For example, it has been found that the fatigue strength of a nickel-based single-crystal superalloy is 443 MPa at 870 °C, with deformation dominated by the motion of dislocations in the γ matrix channels (i.e., the continuous γ phase surrounding the γ′ precipitates) [15]. In the temperature range of 700–800 °C, the HCF limit decreases from 410 MPa to 380 MPa as temperature increases, with the fracture mechanism being quasi-cleavage in all cases [16]. Although studies [17,18] have demonstrated that hot isostatic pressing can significantly reduce the number and size of microporosity in nickel-based single-crystal high-temperature alloys, thereby improving the alloy’s high-cycle fatigue performance, systematic research on the influence of solidification conditions on the high-cycle fatigue performance of single-crystal superalloys remains scarce. Furthermore, few studies have investigated the effects of solidification conditions on the high-cycle fatigue fracture mechanism and deformation mechanism. This study systematically investigates the effect of solidification conditions on the HCF behavior of the second-generation DD6 single-crystal superalloy [19]. The aim is to provide a theoretical basis for optimizing the directional solidification process and enhancing the HCF performance of DD6 alloy.

2. Materials and Methods

The second-generation single-crystal superalloy DD6 [20], with the nominal chemical composition listed in Table 1, was investigated in this study. The standard heat treatment was conducted as follows: 1290 °C/1 h + 1300 °C/2 h + 1315 °C/4 h/AC, followed by 1120 °C/4 h/AC, and finally 870 °C/32 h/AC.
Single-crystal bars with a [001] orientation were prepared using both the HRS and LMC processes. The bars had a diameter of 15 mm and a length of 160 mm. For the LMC process, a molten Al bath served as the cooling medium. Experiments were conducted at a constant withdrawal rate of 3.5 mm/min, with the pouring temperature varied to achieve different solidification conditions, as detailed in Table 2. The alloys prepared by HRS and LMC are hereafter referred to as the HRS and LMC alloys, respectively. The dendrite arm spacing of single-crystal superalloys prepared using the processes shown in Table 2 exhibits significant differences [4,5,6]. This indicates that the HRS and LMC processes enable precise control of solidification parameters such as temperature gradients and solidification rates by adjusting the pouring temperature.
The crystal orientation of the single-crystal bars was determined using the X-ray diffraction Laue back-reflection technique. Bars with their [001] orientation deviating less than 8° from the primary stress axis were selected for subsequent processing. After heat treatment, the single-crystal bars were machined into smooth HCF specimens (stress concentration factor, Kt = 1). The dimensions and quantity of the specimens are illustrated in Figure 1, which conforms to the requirements of the HB 5153 standard.
The HCF tests were performed on a rotating-beam fatigue testing machine under ambient atmosphere. The tests were conducted at a temperature of 800 °C with a stress ratio of R = −1 and a rotational speed of 5000 rpm. The test temperature was accurately monitored and controlled via a thermocouple. Throughout the entire process, the temperature fluctuation was maintained within ±0.5 °C of the nominal temperature.
The macroscopic morphology of the fatigue fracture surfaces was observed using an LEO 1450 scanning electron microscope (SEM, LEO Electron Microscopy Ltd., Cambridge, UK). For transmission electron microscopy (TEM, TECNAI-20, FEI Company, Hillsboro, OR, USA) analysis, thin slices approximately 0.4 mm thick were sectioned parallel to the stress axis at a location about 1 mm away from the fracture surface. These slices were then mechanically ground down to a thickness of 50 μm, punched into 3 mm discs, and finally thinned via twin-jet electro-phishing. The electrolyte was a solution of 10% perchloric acid in alcohol (by volume), and the process was conducted at a voltage of 30 V and a temperature of −25 °C. The dislocation configurations after fatigue failure were characterized using a TECNAI-20 transmission electron microscope.

3. Results

3.1. High-Cycle Fatigue Performance

Figure 2 presents the stress amplitude versus fatigue life (S-N) curves for the DD6 alloy at 800 °C under different solidification conditions. Among these, the medium-life zone depicted in the figure was obtained using the grouping method, while the long-life zone was obtained using the staircase method. The results indicate that the fatigue life of all alloys decreases with increasing stress amplitude, a trend that is independent of the solidification conditions. Specifically, as the pouring temperature increased from 1500 °C to 1590 °C, the HCF performance of the HRS alloys continuously improved. Their conditional fatigue limits, defined as the stress amplitude at which the specimen endured 107 cycles without failure, were determined to be 400.1 MPa, 421.4 MPa, and 427.5 MPa, corresponding to increases of 5.4% and 1.4%, respectively. It is noteworthy that under the same pouring temperature (1590 °C), the LMC alloy exhibited superior performance, achieving a conditional fatigue limit of 453.3 MPa, which is 9.4% higher than that of its HRS counterpart. These findings conclusively demonstrate that both increasing the pouring temperature and employing the LMC process can effectively enhance the HCF performance of the DD6 alloy, with the LMC process offering a more pronounced improvement.

3.2. High-Cycle Fatigue Fraction

Figure 3 shows the macroscopic fracture morphologies of the DD6 alloy after HCF testing at 800 °C under different solidification conditions. As illustrated, all fractured specimens, regardless of solidification condition, exhibited no significant macroscopic plastic deformation, characteristic of typical brittle fracture. The fatigue fractures consistently comprised three distinct regions: the fatigue crack initiation zone, the crack propagation zone, and the final instantaneous fracture zone. The first two zones collectively accounted for the vast majority of the fracture surface area (exceeding 80% in most cases). The fracture surfaces of all specimens were composed of one or more specific crystallographic planes. The angles between these planes and the direction of the primary stress axis remained stable between 40° and 50°, indicating a strong crystallographic orientation dependence. X-ray orientation analysis confirmed that these crystallographic planes were the {111} planes. Notably, these fractographic features were consistent across both high and low stress amplitudes and did not vary with the applied stress. These results demonstrate a remarkable consistency in the macroscopic HCF fracture morphology of the DD6 alloy, whether prepared by HRS or LMC. This suggests that the HCF fracture mode of the DD6 alloy at 800 °C is primarily governed by the test temperature and is not directly related to the solidification conditions or the applied stress level. Furthermore, comparative analysis revealed that the HCF fracture characteristics of the DD6 alloy at 800 °C observed in this study are essentially consistent with those reported in the literature for DD6 alloy subjected to pre-treatments such as overheating [21] and shot-peening [22] in the medium-temperature regime, and also highly consistent with the HCF fracture features of DD32 alloy [23] within a similar temperature range.
Additionally, typical river patterns were observed in the crack propagation zones of all alloy specimens, as shown in Figure 4. The convergence of tributaries into mainstream rivers indicates the direction of fatigue crack propagation. Simultaneously, step-like terrace structures were observed in the final fracture zones, as shown in Figure 5. These terrace structures appeared in fractures composed of multiple intersecting crystallographic planes but were absent in fractures consisting of a single crystallographic plane.
In summary, the macroscopic HCF fracture morphologies of the DD6 alloy under different solidification conditions are fundamentally consistent: the fatigue fracture surfaces consist of one or more intersecting {111} crystallographic planes, accompanied by river patterns and step-like terraces. These characteristics indicate that the HCF fracture of the DD6 alloy at 800 °C is quasi-cleavage in nature, independent of the solidification conditions. This conclusion agrees with findings reported in the literature [21,24,25].
Figure 6 shows the morphologies of the HCF crack initiation zones for the DD6 alloy under different solidification conditions. The results indicate that under all solidification conditions, fatigue cracks in the DD6 alloy initiated at the surface or sub-surface regions of the specimens, and the fractures exhibited characteristics of a single macroscopic fatigue source. A similar result has been reported in HCF studies of CMSX-4 alloy [26]. This phenomenon is consistent with the stress distribution under rotating bending loading, where the specimen surface experiences the maximum tensile stress, making it the most susceptible region for crack initiation. After initiating at the surface or sub-surface, the fatigue cracks propagated along the {111} slip planes until final instability and fracture. Although the crack initiation sites were common across different solidification conditions, the specific initiation mechanisms showed subtle variations: at a pouring temperature of 1500 °C, fatigue cracks in the HRS alloy mainly initiated at stress concentration sites formed by the combined presence of microporosity and γ/γ′ eutectic, as shown in Figure 6a, where the smooth interface of the eutectic structure contrasts distinctly with the irregular morphology of microporosity. When the pouring temperature was increased to 1560 °C and 1590 °C, the crack initiation sites in the HRS alloy became more singular, originating primarily from microporosity, as shown in Figure 6b,c. For the LMC alloy, fatigue cracks also initiated from microporosity, as shown in Figure 6d. Based on the above discussion, microporosity represent critical locations for fatigue crack initiation. This conclusion aligns with observations from high-temperature in situ fatigue testing, which directly confirmed that under high-temperature cyclic loading, fatigue cracks preferentially originate from microporosity [27].
Figure 7 presents magnified views of the HCF crack propagation zones for the DD6 alloy under different solidification conditions. It can be seen that the crack propagation path exhibits distinct crystallographic characteristics, primarily extending along the {111} slip planes. Notably, dispersed small-sized (≤10 μm) microporosity was observed on the {111} slip planes within the propagation zones. This indicates that microporosity played a promoting role during crack propagation. The presence of these micropores disrupts the continuity of the matrix, facilitating crack advance and thereby being detrimental to the HCF life of the DD6 alloy.

3.3. Microstructure

Microporosity is a common microstructural defect in nickel-based single-crystal superalloys. Based on their formation mechanisms, micropores can be classified into two types: solidification pores and solution pores. Solidification pores form due to insufficient liquid metal feeding to compensate for volumetric shrinkage during solidification, and they typically occur near dendrite cores and eutectic pools. Solution pores form during the heat treatment of the as-cast single-crystal superalloy, resulting from non-equilibrium elemental diffusion in inter-dendritic regions. Generally, solidification pores are larger and irregularly shaped, while solution pores are often regularly spherical. Since both types of pores coexist in the heat-treated alloy specimens and are difficult to distinguish completely in practical observation, the term “microporosity” will be used uniformly hereafter without strict differentiation between solidification and solution pores.
Figure 8 shows the typical morphology of microporosity in the heat-treated DD6 alloy under different solidification conditions. Quantitative metallographic analysis revealed that at pouring temperatures of 1500 °C, 1560 °C, and 1590 °C, the microporosity content in the HRS alloys was 0.28%, 0.22%, and 0.18%, respectively, indicating a decreasing trend with increasing pouring temperature. At the same pouring temperature of 1590 °C, the LMC alloy exhibited a microporosity content of only 0.08%, which is merely 44.4% of that in the HRS alloy. The morphology of microporosity in HRS alloys at 1500 °C and 1560 °C was mostly irregular. When the pouring temperature was increased to 1590 °C, the micropores in both the HRS and LMC alloys became predominantly regularly spherical.
Figure 9 illustrates the pore size distribution characteristics of the microporosity in the DD6 alloy under different solidification conditions. To accurately quantify the size of irregularly shaped pores, they were approximated as equivalent circular pores, and the reported pore size is the equivalent diameter. The results show that the maximum equivalent diameters of microporosity in HRS alloys at pouring temperatures of 1500 °C, 1560 °C, and 1590 °C were 48 μm, 41 μm, and 31 μm, respectively, demonstrating a decreasing trend with increasing pouring temperature. This trend is consistent with the change in microporosity content, collectively confirming that increasing the pouring temperature can effectively suppress the formation of microporosity by improving melt feeding. Notably, the LMC alloy exhibited a maximum equivalent diameter of only 22 μm, which is merely 45.8% of that in the HRS alloy, indicating a more significant improvement achieved by the LMC process. Furthermore, the pore size distributions of both HRS and LMC alloys were generally similar, characterized by a gradual decrease in the number of micropores as the pore size increased.
In summary, increasing the pouring temperature or employing the LMC process resulted in varying degrees of improvement in the content and size of microporosity in the DD6 alloy. Although the shapes of the pore size distribution curves were similar for HRS and LMC alloys, the LMC process shifted the entire curve towards smaller pore sizes. This indicates that the LMC process not only reduces the pore size but also systematically optimizes the pore size structure, resulting in superior directional solidification quality.
Figure 10 shows the cross-sectional microstructure near the HCF fracture surface of the DD6 alloy at 800 °C under different solidification conditions. The γ′ precipitates in both HRS and LMC alloy specimens maintained a complete cuboidal morphology, exhibiting a regular spatial arrangement and good cuboidal alignment, with no evidence of rafting. Comparative analysis with the microstructure of the heat-treated DD6 alloy under different solidification conditions reported in previous studies [4] revealed that the morphology and size distribution of the γ′ precipitates after HCF fracture remained consistent with those before fatigue testing. After fatigue fracture, the γ′ precipitates size in the HRS alloy at pouring temperatures of 1500 °C, 1560 °C, and 1590 °C were 0.426 μm, 0.506 μm, and 0.483 μm, respectively. At the pouring temperature of 1590 °C, the γ′ precipitates size in the LMC alloy was 0.351 μm. This indicates that the γ′ precipitates size in the HRS alloy initially increased and then decreased with increasing pouring temperature, reaching a maximum at 1560 °C. Furthermore, under the same pouring temperature condition, the γ′ precipitates size in the LMC alloy was significantly smaller than that in the HRS alloy. These findings indicate that the DD6 alloy exhibited excellent microstructural stability under 800 °C HCF loading, regardless of the solidification condition. The γ′ precipitates did not undergo significant coarsening, dissolution, or morphological transformation. The stability of the alloy’s microstructure confirms that the plastic strain experienced during HCF was relatively low, with deformation primarily confined to the γ matrix channels, which is consistent with the limited macroscopic plastic deformation observed in the specimens.
Figure 11 shows the typical dislocation configurations in the DD6 alloy after HCF fracture at 800 °C under different solidification conditions. As the pouring temperature increased from 1500 °C to 1590 °C, the dislocation density showed a significant decreasing trend. At a pouring temperature of 1500 °C, dense dislocation tangles were observed in the HRS alloy specimen, indicating relatively severe plastic deformation under this solidification condition. The cuboidal morphology of some γ′ precipitates began to degrade, transforming towards irregular shapes, accompanied by an increase in the width of the γ matrix channels. When the pouring temperature increased to 1560 °C, the dislocation tangles in the HRS alloy specimen began to diminish. At the pouring temperature of 1590 °C, the dislocation configurations in both HRS and LMC alloys transformed into sparse, wavy dislocation lines without tangles, reflecting a significant reduction in the degree of plastic deformation under these solidification conditions. Correspondingly, at pouring temperatures of 1560 °C and 1590 °C, the overall integrity of the cuboidal γ′ precipitates were well maintained, and the γ matrix channel structure remained stable, a phenomenon applicable to both HRS and LMC alloys. In summary, under a fixed stress amplitude, the pouring temperature significantly influences the evolution of dislocation density and configuration in the DD6 alloy after 800 °C HCF fracture.
Further observation revealed that dislocation shearing of the γ′ precipitates was not observed in any DD6 alloy specimens after HCF fracture, regardless of solidification condition. Similarly, the formation of dislocation networks was not observed. Dislocations exhibited only distinct interfacial configurations, moving through the γ matrix channels by bowing and cross-slip mechanisms. This characteristic indicates that the deformation mechanism for the DD6 alloy during 800 °C HCF is the slip of a/2<011> dislocations within the γ matrix channels. The solidification conditions did not alter the fundamental HCF deformation mechanism of the alloy.

4. Discussion

The enhancement in the high-cycle fatigue performance of the DD6 alloy can be achieved by increasing the pouring temperature or employing the LMC process, which effectively improves the characteristics of microporosity. Under identical pouring temperature conditions, the LMC process demonstrates superior effectiveness in microstructural refinement and associated performance enhancement. Notably, the HCF life of the LMC alloy can be more than doubled compared to that of the HRS alloy, particularly under low stress amplitude conditions.
In high-cycle fatigue, the crack initiation stage typically constitutes the majority of the total fatigue life [28,29]. This characteristic makes the crack initiation behavior highly sensitive to the material’s microstructural features. Therefore, investigating the influence of solidification conditions on HCF crack initiation behavior is crucial for understanding the differences in HCF performance. It is generally accepted that HCF cracks in single-crystal superalloys predominantly initiate at microstructural discontinuities such as microporosity, γ/γ′ eutectic, and carbides [30,31,32,33,34,35]. Numerous studies [22,36,37] have reported that HCF cracks in DD6 alloy primarily initiate at microporosity, and no crack initiation at carbides was observed in the present study. Therefore, the effect of carbides on crack initiation will not be discussed further.
The present study reveals that, except for the HRS alloy at the 1500 °C pouring temperature where cracks initiated at both microporosity and γ/γ′ eutectic, the fatigue cracks in HRS alloys at other pouring temperatures and in the LMC alloy initiated solely at microporosity. This indicates that while variations in solidification conditions lead to subtle differences in crack initiation sites, all fatigue cracks originate from metallurgical defects at or near the specimen surface. These results confirm that changes in solidification conditions do not alter the fundamental mechanism of HCF crack initiation in the DD6 alloy but merely influence the specific type of defect responsible for initiation. Previous research has indicated the presence of significant residual γ/γ′ eutectic in the HRS alloy cast at 1500 °C. This is likely the underlying reason for the concurrent crack initiation at both microporosity and γ/γ′ eutectic under this specific condition. The synergistic effect of microporosity and γ/γ′ eutectic promotes fatigue crack initiation in the DD6 alloy, which is detrimental to its HCF performance. As the pouring temperature increases or the LMC process is employed, the γ/γ′ eutectic content decreases or even vanishes, consequently diminishing its role as a fatigue crack initiator. This is corroborated by the observation that fatigue cracks in alloys under other solidification conditions initiated solely from microporosity.
A deeper analysis reveals a clear causal relationship between the improvement in microporosity characteristics and the enhancement of HCF performance in the DD6 alloy with optimized solidification conditions. As the pouring temperature increases or the LMC process is used, the content and size of microporosity are improved to varying degrees, corresponding to a continuous improvement in HCF performance.
From the perspective of the fatigue damage mechanism, the size of microporosity, acting as stress concentrators, significantly influences HCF failure [25]. According to the literature [38], the stress intensity factor generated by a pore can be expressed by Equation (1):
K 1 = Y σ π A p o r e
where K1 is the stress intensity factor, Y is a constant (Y = 0.5 for an internal pore; Y = 0.65 for a pore near the surface), σ is the applied stress amplitude, and Apore is the area of the equivalent circular pore.
Equation (1) shows that under the same stress amplitude, a larger pore area, i.e., a larger pore size, results in higher local stress concentration. This geometric strengthening effect significantly increases the effective stress intensity at the crack tip, allowing the material to reach the critical condition for fatigue crack initiation earlier, thereby shortening the crack initiation life and ultimately leading to reduced total fatigue life. As the pouring temperature increased from 1500 °C to 1590 °C, the maximum equivalent diameter of microporosity in the HRS alloy decreased from 48 μm to 31 μm, and further to 22 μm in the LMC alloy. The reduction in the maximum pore size corresponds to a decrease in stress concentration, which delays crack initiation, increases the crack initiation life, and consequently enhances the HCF performance. A clear negative correlation exists between the HCF life and the maximum microporosity size. This conclusion is also supported by the research of Brundidge et al. [39].
This study also found that small, dispersed micropores were present within the crack propagation zones of DD6 alloy specimens under all solidification conditions. These pores disrupt the continuity of the matrix, facilitating accelerated crack propagation through a “pore-connection” mechanism. This propagation mode effectively reduces the energy required for crack advance, allowing the crack to cover the same distance in fewer cycles, thereby increasing the crack propagation rate. Research by Ma et al. [40] also observed that when a propagating crack encounters a micropore, the crack length increases significantly, accelerating propagation. However, similar crack–pore interaction characteristics were observed in alloy specimens under different solidification conditions, indicating that this mechanism is relatively insensitive to changes in solidification conditions. This suggests that small, dispersed micropores are not the dominant factor causing the differences in HCF performance under different solidification conditions. This conclusion is consistent with the findings of Huang [41], who concluded through quasi-in situ fatigue experiments that small, dispersed micropores have a limited effect on fatigue performance, whereas large micropores, acting as the most significant stress concentrators, directly control fatigue crack initiation and thus dominate the HCF life of the alloy.
In summary, increasing the pouring temperature or employing the LMC process enhances the HCF performance of the alloy primarily by reducing or eliminating large micropores. The small micropores commonly present within the specimens have a comparatively minor effect on the HCF performance. Effectively suppressing the maximum size of microporosity through the optimization of solidification conditions is key to improving the HCF performance of single-crystal superalloys.

5. Conclusions

Based on a systematic investigation into the effects of solidification conditions on the microstructure and high-cycle fatigue performance of the DD6 single-crystal superalloy, the main conclusions are summarized as follows:
(1)
Increasing the pouring temperature effectively reduced both the content and size of microporosity in the heat-treated HRS alloys. Under the same pouring temperature, the LMC alloy exhibited significantly less microporosity content (only 44.4% of the HRS alloy) and smaller pore sizes (only 45.8% of the HRS alloy).
(2)
The high-cycle fatigue performance of the HRS alloys was enhanced with increasing pouring temperature. Under the same pouring temperature, the LMC alloy demonstrated superior HCF performance compared to the HRS alloy, with a 9.4% increase in the conditional fatigue limit.
(3)
For the HRS alloy at a pouring temperature of 1500 °C, HCF cracks initiated at surface/subsurface microporosity and γ/γ′ eutectic. When the pouring temperature was increased to 1560 °C and 1590 °C, or when the LMC process was employed, cracks initiated solely from surface/subsurface microporosity. Microporosity is the dominant factor for HCF crack initiation in DD6 alloy at 800 °C, while the role of γ/γ′ eutectic in crack initiation weakens or even vanishes as solidification conditions are optimized.
(4)
The HCF fracture of the DD6 alloy at 800 °C exhibited characteristics of quasi-cleavage, which was independent of the solidification conditions.
(5)
At a stress amplitude of 550 MPa, dislocation shearing of the γ′ precipitates was not observed in the DD6 alloy after HCF fracture at 800 °C. The deformation was characterized by the slip of a/2<011> dislocations within the γ matrix channels via bowing and cross-slip mechanisms. The solidification conditions did not alter the fundamental HCF deformation mechanism of the alloy.
(6)
Approaches that increase the temperature gradient, such as raising the pouring temperature or employing the LMC process, are effective in improving the content and size of microporosity in the DD6 alloy. This improvement subsequently enhances the overall HCF performance primarily by prolonging the fatigue crack initiation life.

Author Contributions

Conceptualization, H.X. and Y.Z.; methodology, H.X. and Y.Z.; software, H.X. and Y.L.; validation, Y.Z.; formal analysis, H.X. and Z.Y.; investigation, H.X. and Y.L.; resources, Y.Z.; data curation, H.X. and Y.L.; writing—original draft preparation, H.X.; writing—review and editing, H.X. and Y.L.; visualization, Y.Z.; supervision, Y.Z. and Z.Y.; project administration, Y.Z. and Z.Y.; funding acquisition, H.X. and Y.L. All authors have read and agreed to the published version of the manuscript.

Funding

This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

The author declares that he has no known competing financial interests or personal relationships that could have appeared to influence the work reported in this article.

References

  1. Bu, J.; Gao, Z.; Niu, J.; Cao, Y. Crack failure analysis of a fan stator vane. Aeroengine 2021, 47, 91–95. [Google Scholar]
  2. Wang, H. Failure analysis of turbine blade fracture in R20 gas turbine. Gas Turbine Technol. 2023, 36, 62–67. [Google Scholar]
  3. Gao, Z.; Bu, J.; Zhang, K.; Tong, W.; Li, M.; Wang, W. Fracture Failure Analysis of DD10 single crystal turbine blade under service environment. Aeroengine 2025, 51, 149–154. [Google Scholar]
  4. Xie, H.; Li, J.; Luo, Y.; Zheng, S.; Luo, K. Effect of Solidification Conditions on Microstructure Evolution in DD6 single-crystal superalloy. Acta Metall. Sin. 2025. Available online: https://link.cnki.net (accessed on 14 August 2025).
  5. Kurz, W.; Fisher, D.J. Dendrite growth at the limit of stability: Tip radius and spacing. Acta Metall. 1981, 29, 11–20. [Google Scholar] [CrossRef]
  6. Wagner, A.; Shollock, B.A.; McLean, M. Grain structure development in directional solidification of nickel-base superalloys. Mater. Sci. Eng. A 2004, 374, 270–279. [Google Scholar] [CrossRef]
  7. Tang, X.; Zhang, Y.; Li, J. Directional solidification of a Ni-based single crystal superalloy under high temperature gradient. Rare Met. Mater. Eng. 2012, 41, 738–742. [Google Scholar]
  8. Wang, R.; Li, J.; Yue, X.; Zhao, J.; Yang, W. Deformation and fracture mechanism of third-generation single crystal superalloy during in-situ tension at room temperature. Rare Met. Mater. Eng. 2025, 54, 1410–1416. [Google Scholar]
  9. Han, M.; Yu, J.; Li, J.; Xie, H.; Dong, J.; Yang, Y. Influence of shot peening on tensile properties of DD6 single crystal superalloy. J. Mater. Eng. 2019, 47, 169–175. [Google Scholar]
  10. Sun, J.; Liu, J.; Chen, C.; Li, J.; Sun, X. Effect of γ’ size on intermediate temperature stress rupture property of the third generation single crystal nickel-base superalloy containing Re. Rare Met. Mater. Eng. 2022, 51, 369–373. [Google Scholar]
  11. Yang, W.; Li, J.; Liu, S.; Zhao, J.; Shi, Z.; Wang, X. Transverse stress rupture properties of a third generation single crystal superalloy at medium and elevated temperatures. J. Mater. Eng. 2020, 48, 139–145. [Google Scholar]
  12. Yu, J.; Li, J.; Fang, X.; Wang, Q.; Liu, S.; Han, M. Influence of secondary γ′ phase evolution on creep properties of single crystal superalloy DD6. J. Mater. Eng. 2023, 51, 60–66. [Google Scholar]
  13. Ma, L.; Wang, D.; Zhang, G.; Shen, J.; Zhang, J. Effect of Hf on microstructure and creep properties of nickel-based bicrystal superalloy with different grain boundaries misorientations. Mater. Sci. Eng. A 2025, 935, 148371. [Google Scholar] [CrossRef]
  14. Xie, H.; Li, J.; Han, M. Effect of stress ratio on high cycle fatigue behavior of a single crystal superalloy. Rare Met. Mater. Eng. 2018, 47, 3381–3386. [Google Scholar]
  15. Shui, L.; Xu, Y.; Hu, Z. Dislocation structure in a single crystal nickel base superalloy during high cycle fatigue at 870 °C. Rare Met. Mater. Eng. 2018, 47, 1054–1058. [Google Scholar] [CrossRef]
  16. Shi, Z.; Zhao, J. High cycle fatigue properties of a single crystal superalloy at different temperatures. Nonferr. Met. Sci. Eng. 2019, 10, 58–63. [Google Scholar]
  17. Luo, Y.; Guo, H.; Zhao, Y.; Zhang, J. Effect of hot isostatic pressing on high-temperature high cycle fatigue properties of a second generation single crystal superalloy DD6. Mater. Mech. Eng. 2016, 40, 51–56. [Google Scholar]
  18. Wei, C.N.; Bor, H.Y.; Chang, L. Effect of hot isostatic pressing on microstructure and mechanical properties of CM-681LC nickel-base superalloy using microcast. Mater. Trans. 2008, 49, 193–201. [Google Scholar] [CrossRef]
  19. Li, J.R.; Zhong, Z.G.; Tang, D.Z.; Liu, S.Z.; Wei, P.; Wei, P.Y.; Wu, Z.T.; Huang, D.; Han, M. A low-cost second generation single crystal superalloy DD6. In Superalloys 2000 (Ninth International Symposium); Pollock, T.M., Kissinger, R.D., Bowman, R.R., Green, K.A., Mclean, M., Olson, S., Schina, J.J., Eds.; TMS (The Minerals, Metals &Materials Societ): Pittsburgh, PA, USA, 2000; pp. 777–783. [Google Scholar]
  20. Li, J.R.; Zhao, J.Q.; Liu, S.Z.; Han, M. Effects of low angle boundaries on the mechanical properties of single crystal superalloy DD6. In Superalloys 2008, Proceedings of the Eleventh International Symposium on Superalloys, Champion, PA, USA, 14–18 September 2008; Reed, R.C., Green, K.A., Eds.; The Minerals, Metals & Materials Society: Warrendale, PA, USA, 2008; pp. 443–450. [Google Scholar]
  21. Xie, H.; Li, J.; Han, M.; Yu, J.; Yang, L.; Yue, X. Effect of over-temperature on microstructure and high cycle fatigue properties of DD6 single crystal superalloy. Rare Met. Mater. Eng. 2022, 47, 2483–2488. [Google Scholar]
  22. Li, J.; Dong, J.; Han, M.; Liu, S. Effects of Sand Blasting on Surface Integrity and High Cycle Fatigue Properties of DD6 single crystal superalloy. Acta Metall. Sin. 2023, 59, 1201–1208. [Google Scholar]
  23. Yu, J.; Yang, Y.; Sun, X.; Guan, H.; Hu, Z. Rotary bending high-cycle fatigue behavior of DD32 single crystal superalloy containing rhenium. J. Mater. Sci. 2012, 47, 4805–4812. [Google Scholar] [CrossRef]
  24. Yi, J.Z.; Torbet, C.J.; Feng, Q.; Pollock, T.M.; Jones, J.W. Ultrasonic fatigue of a single crystal Ni-base superalloy at 1000 °C. J. Mater. Sci. Eng. A 2007, 443, 142–149. [Google Scholar] [CrossRef]
  25. Liu, Y.; Yu, J.J.; Xu, Y.; Sun, X.F.; Guan, H.R.; Hu, Z.Q. High cycle fatigue behavior of a single crystal superalloy at elevated temperatures. Mater. Sci. Eng. A 2007, 454–455, 357–366. [Google Scholar] [CrossRef]
  26. Wasson, A.J.; Fuchs, G.E. The effect of carbide morphologies on elevated temperature tensile and fatigue behavior of a modified single crystal Ni-Base superalloy. In Superalloys 2008, Proceedings of the Eleventh International Symposium on Superalloys, Champion, PA, USA, 14–18 September 2008; Reed, R.C., Green, K.A., Eds.; The Minerals, Metals & Materials Society: Warrendale, PA, USA, 2008; pp. 489–497. [Google Scholar]
  27. Hu, C.; Liu, X.; Tao, C.; Kong, Z. Demage behavior of film holes of DD6 single crystal superalloy by electro-stream machining. Rare Met. Mater. Eng. 2019, 48, 3190–3194. [Google Scholar]
  28. Liu, M.; Zou, T.; ZhangYu, T.; Ni, S.; Pei, Y.; Wang, Q.; Zhang, H.; Liu, Y.; Wang, Q. Crack initiation and lifetime assessment of MAR-M247 nickel-based superalloy in the very high cycle fatigue regime at 550 °C. Int. J. Fatigue 2025, 21, 109176. [Google Scholar] [CrossRef]
  29. Rémy, L.; Geuffrard, M.; Alam, A.; Köster, A.; Fleury, E. Effects of microstructure in high temperature fatigue: Lifetime to crack initiation of a single crystal superalloy in high temperature low cycle fatigue. Int. J. Fatigue 2013, 57, 37–49. [Google Scholar] [CrossRef]
  30. Zhao, Z.; Zhang, F.; Dong, C.; Yang, X.; Chen, B. Initiation and early-stage growth of internal fatigue cracking under very high-cycle fatigue regime at high temperature. Metall. Mater. Trans. A 2020, 51, 1575–1592. [Google Scholar] [CrossRef]
  31. Yang, M.; Zhou, C.; Zhao, Z.; Shen, Y.; Pei, H.; Zhao, M. Effect of oxidation on the competition between internal and external fatigue crack initiation of Ni-based single crystal superalloy. Int. J. Fatigue 2025, 197, 108930. [Google Scholar] [CrossRef]
  32. Utada, S.; Ormastroni, L.M.B.; Rame, J.; Villechaise, P.; Cormier, J. VHCF life of AM1 Ni-based single crystal superalloy after pre-deformation. Int. J. Fatigue 2021, 148, 106224. [Google Scholar] [CrossRef]
  33. Ormastroni, L.M.B.; Suave, L.M.; Cervellon, A.; Villechaise, P.; Cormier, J. LCF, HCF and VHCF life sensitivity to solution heat treatment of a third-generation Ni-based single crystal superalloy. Int. J. Fatigue 2020, 130, 105247. [Google Scholar] [CrossRef]
  34. Cervellon, A.; Hémery, S.; Kürnsteiner, B.; Gault, P.; Kontis, P.; Cormier, J. Crack initiation mechanisms during very high cycle fatigue of Ni-based single crystal superalloys at high temperature. Acta Mater. 2020, 188, 131–144. [Google Scholar] [CrossRef]
  35. Cervellon, A.; Cormier, J.; Mauget, F.; Hervier, Z.; Nadot, Y. Very high cycle fatigue of Ni-based single crystal superalloy at high temperature. Metall. Mater. Trans. A 2018, 49, 3938–3950. [Google Scholar] [CrossRef]
  36. Li, J.; Xie, H.; Han, M.; Liu, S. High cycle fatigue behavior of second generation single crystal superalloy. Acta Metall. Sin. 2019, 55, 1195–1203. [Google Scholar]
  37. Murakami, Y. Metal Fatigue: Effects of Small Defects and Nonmetallic Inclusions, 2nd ed.; Academic Press: Cambridge, MA, USA, 2019. [Google Scholar]
  38. Murakami, M.; Endo, M. Effects of defects, inclusions and inhomogeneities on fatigue strength. Int. J. Fatigue 1994, 16, 163–182. [Google Scholar] [CrossRef]
  39. Brundidge, C.L.; Pollock, T.M. Processing to fatigue properties: Benefits of high gradient casting for single crystal airfoils. In Superalloys 2012: 12th International Symposium on Superalloys; Huron, E.S., Reed, R.C., Hardy, M.C., Mills, M.J., Montero, R.E., Portella, P.D., Telesman, J., Eds.; TMS (The Minerals, Metals & Materials Societ): Pittsburgh, PA, USA, 2012; pp. 379–385. [Google Scholar]
  40. Ma, A.; Shi, H.; Gu, J.; Chen, G.; Luesebrink, O.; Hardersd, H. In-situ observations of the effects of orientation and carbide on low cycle fatigue crack propagation in a single crystal superalloy. Procedia Eng. 2010, 188, 2287–2295. [Google Scholar] [CrossRef]
  41. Huang, Y.; Shen, J.; Wang, D.; Xie, G.; Lu, Y.; Lou, L.; Zhang, J. Formation of sliver defect in Ni-based single crystal superalloy. Metall. Mater. Trans. A 2020, 51, 99–103. [Google Scholar] [CrossRef]
Metals 15 01385 g001
Figure 1. Schematic of smooth specimen for high cycle fatigue test (mm).
Figure 1. Schematic of smooth specimen for high cycle fatigue test (mm).
Metals 15 01385 g001
Metals 15 01385 g002
Figure 2. High cycle fatigue performance of DD6 alloy under different solidification conditions.
Figure 2. High cycle fatigue performance of DD6 alloy under different solidification conditions.
Metals 15 01385 g002
Metals 15 01385 g003
Figure 3. High cycle fatigue fraction of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 450 MPa, Nf = 5.49 × 105; (b) HRS, 1500 °C, σa = 550 MPa, Nf = 7.35 × 104; (c) HRS, 1560 °C, σa = 500 MPa, Nf = 1.99 × 106; (d) HRS, 1560 °C, σa = 600 MPa, Nf = 3.56 × 105; (e) HRS, 1590 °C, σa = 450 MPa, Nf = 3.46 × 106; (f) HRS, 1590 °C, σa = 600 MPa, Nf = 3.28 × 105; (g) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106; (h) LMC, 1590 °C, σa = 600 MPa, Nf = 4.04 × 105.
Figure 3. High cycle fatigue fraction of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 450 MPa, Nf = 5.49 × 105; (b) HRS, 1500 °C, σa = 550 MPa, Nf = 7.35 × 104; (c) HRS, 1560 °C, σa = 500 MPa, Nf = 1.99 × 106; (d) HRS, 1560 °C, σa = 600 MPa, Nf = 3.56 × 105; (e) HRS, 1590 °C, σa = 450 MPa, Nf = 3.46 × 106; (f) HRS, 1590 °C, σa = 600 MPa, Nf = 3.28 × 105; (g) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106; (h) LMC, 1590 °C, σa = 600 MPa, Nf = 4.04 × 105.
Metals 15 01385 g003
Metals 15 01385 g004
Figure 4. River pattern in the high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Figure 4. River pattern in the high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Metals 15 01385 g004
Metals 15 01385 g005
Figure 5. Stepped step structure in the high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Figure 5. Stepped step structure in the high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Metals 15 01385 g005
Metals 15 01385 g006
Figure 6. Morphologies of the HCF crack initiation zones for the DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Figure 6. Morphologies of the HCF crack initiation zones for the DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Metals 15 01385 g006
Metals 15 01385 g007
Figure 7. Microporosity in crack propagation zone of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 500 MPa, Nf = 7.44 × 105; (b) HRS, 1560 °C, σa = 500 MPa, Nf = 9.51 × 105; (c) HRS, 1590 °C, σa = 500 MPa, Nf = 1.38 × 106; (d) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106.
Figure 7. Microporosity in crack propagation zone of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 500 MPa, Nf = 7.44 × 105; (b) HRS, 1560 °C, σa = 500 MPa, Nf = 9.51 × 105; (c) HRS, 1590 °C, σa = 500 MPa, Nf = 1.38 × 106; (d) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106.
Metals 15 01385 g007
Metals 15 01385 g008
Figure 8. Microporosity in heat-treated DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Figure 8. Microporosity in heat-treated DD6 alloy under different solidification conditions: (a) HRS, 1500 °C; (b) HRS, 1560 °C; (c) HRS, 1590 °C; (d) LMC, 1590 °C.
Metals 15 01385 g008
Metals 15 01385 g009
Figure 9. Size characteristics of microporosity in heat-treated DD6 alloy under different solidification conditions.
Figure 9. Size characteristics of microporosity in heat-treated DD6 alloy under different solidification conditions.
Metals 15 01385 g009
Metals 15 01385 g010
Figure 10. Microstructure of section near high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 500 MPa, Nf = 7.44 × 105; (b) HRS, 1560 °C, σa = 500 MPa, Nf = 9.51 × 105; (c) HRS, 1590 °Cs, σa = 500 MPa, Nf = 1.38 × 106; (d) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106.
Figure 10. Microstructure of section near high cycle fatigue fracture of DD6 alloy under different solidification conditions: (a) HRS, 1500 °C, σa = 500 MPa, Nf = 7.44 × 105; (b) HRS, 1560 °C, σa = 500 MPa, Nf = 9.51 × 105; (c) HRS, 1590 °Cs, σa = 500 MPa, Nf = 1.38 × 106; (d) LMC, 1590 °C, σa = 500 MPa, Nf = 1.86 × 106.
Metals 15 01385 g010
Metals 15 01385 g011
Figure 11. Dislocation morphology of DD6 alloy after high cycle fatigue fracture under different solidification conditions: (a) HRS, 1500 °C, σa = 550 MPa, Nf = 7.35 × 104; (b) HRS, 1560 °C, σa = 550 MPa, Nf = 3.83 × 105; (c) HRS, 1590 °Cs, σa = 550 MPa, Nf = 5.21 × 105; (d) LMC, 1590 °C, σa = 550 MPa, Nf = 7.14 × 106.
Figure 11. Dislocation morphology of DD6 alloy after high cycle fatigue fracture under different solidification conditions: (a) HRS, 1500 °C, σa = 550 MPa, Nf = 7.35 × 104; (b) HRS, 1560 °C, σa = 550 MPa, Nf = 3.83 × 105; (c) HRS, 1590 °Cs, σa = 550 MPa, Nf = 5.21 × 105; (d) LMC, 1590 °C, σa = 550 MPa, Nf = 7.14 × 106.
Metals 15 01385 g011
Table 1. Nominal composition of the high-aluminum Ni-based superalloy (wt. %).
Table 1. Nominal composition of the high-aluminum Ni-based superalloy (wt. %).
CCrCoMoWTaReNbAlHfNi
0.0064.39.02.08.07.52.00.55.60.1bal.
Table 2. Directional solidification processes under diverse solidification conditions.
Table 2. Directional solidification processes under diverse solidification conditions.
Solidification MethodPouring Temperature (°C)Withdrawing Rate (mm/min)
HRS15003.5
1560
1590
LMC15903.5
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

Xie, H.; Luo, Y.; Zhao, Y.; Yang, Z. Effect of Solidification Conditions on High-Cycle Fatigue Behavior in DD6 Single-Crystal Superalloy. Metals 2025, 15, 1385. https://doi.org/10.3390/met15121385

AMA Style

Xie H, Luo Y, Zhao Y, Yang Z. Effect of Solidification Conditions on High-Cycle Fatigue Behavior in DD6 Single-Crystal Superalloy. Metals. 2025; 15(12):1385. https://doi.org/10.3390/met15121385

Chicago/Turabian Style

Xie, Hongji, Yushi Luo, Yunsong Zhao, and Zhenyu Yang. 2025. "Effect of Solidification Conditions on High-Cycle Fatigue Behavior in DD6 Single-Crystal Superalloy" Metals 15, no. 12: 1385. https://doi.org/10.3390/met15121385

APA Style

Xie, H., Luo, Y., Zhao, Y., & Yang, Z. (2025). Effect of Solidification Conditions on High-Cycle Fatigue Behavior in DD6 Single-Crystal Superalloy. Metals, 15(12), 1385. https://doi.org/10.3390/met15121385

Note that from the first issue of 2016, this journal uses article numbers instead of page numbers. See further details here.

Article Metrics

Back to TopTop