The effect of coarsening of γ′ precipitate on creep properties of Ni-based single crystal superalloys during long-term aging

Journal Pre-proof The effect of coarsening of γ′ precipitate on creep properties of Ni-based single crystal superalloys during long-term aging Y.S. Huang, X.G. Wang, C.Y. Cui, J.G. Li, L.H. Ye, G.C. Hou, Y.H. Yang, J.L. Liu, J.D. Liu, Y.Z. Zhou, X.F. Sun PII:

S0921-5093(19)31666-1

DOI:

https://doi.org/10.1016/j.msea.2019.138886

Reference:

MSA 138886

To appear in:

Materials Science & Engineering A

Received Date: 16 September 2019 Revised Date:

25 December 2019

Accepted Date: 26 December 2019

Please cite this article as: Y.S. Huang, X.G. Wang, C.Y. Cui, J.G. Li, L.H. Ye, G.C. Hou, Y.H. Yang, J.L. Liu, J.D. Liu, Y.Z. Zhou, X.F. Sun, The effect of coarsening of γ′ precipitate on creep properties of Nibased single crystal superalloys during long-term aging, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2019.138886. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Y.S. Huang: Writing-Original Draft, Methodology, Investigation X.G. Wang: Conceptualization, Writing - Review & Editing C.Y. Cui: Methodology, Writing Review & Editing J.G. Li: Funding acquisition, Project administration L.H. Ye: Data Curation, Visualization G.C. Hou: Resources Y.H. Yang: Software, Visualization J.L. Liu: Formal analysis, Supervision J.D. Liu: Validation，Supervision Y.Z. Zhou: Project administration X.F. Sun: Funding acquisition, Supervision

The effect of coarsening of γ′ precipitate on creep properties of Ni-based single crystal superalloys during long-term aging Y.S. Huang

a, b

, X.G. Wang

a, ∗

, C.Y. Cui

a, ∗

, J.G. Li a, L.H. Ye a, G.C. Hou a, Y.H.

Yang a, J.L. Liu a, J.D. Liu a, Y.Z. Zhou a, X.F. Sun a, ∗ a

Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua

Road, Shenyang 110016, China b

School of Nano Science and Technology Institute, University of Science and Technology of

China, Suzhou 215123, China

Abstract: The correlation between the coarsening of γ′ precipitate and creep properties during long-term aging at 1000 ºC for 100-1000 h of a fourth-generation nickel-based single crystal superalloy was investigated. The average size of γ′ precipitate increased with the extension of aging time and the morphology of the γ′ precipitate kept cuboidal shape after 1000 h. Experimental observation on the coarsening of γ′ phase demonstrated that the relationship between r − r and t was in accordance with the LSW theory and the coarsening rate k was about 5954 nm3/h. The creep life of aged samples at 1140 ºC/137 MPa increased firstly and then decreased with the exposure time prolonging. The casting pores and creep pores were observed on the fracture surface of samples after creep rupture. Compared with the creep pores, the microcracks formed at the edge of casting pores had greater effect on the creep life. The creep properties of samples during aging were mostly affected by the size of γ′ phase. The promotion of creep life of aged sample was rationalized by the slightly grown γ′ phase, the diffusion of the refractory elements and the accumulation of dislocations in the γ channel. The decrease of creep life was predominantly affected by the seriously coarsened γ′ phase because the significant increase of the size of γ′ phase leaded to a decrease in the resistance of dislocation movement. Keywords: Ni-based single crystal superalloys; Long-term aging; γ′ coarsening; ∗

Corresponding author. E-mail address: [email protected] (X.G. Wang), [email protected] (C.Y. Cui), [email protected] (X.F. Sun) 1

Creep properties. 1. Introduction Nickel-base single crystal superalloys are manufactured as gas turbine engine components because of their excellent high temperature strength, oxidation and corrosion resistance and microstructural stability [1, 2]. During long-term service under elevated temperature, the microstructures of superalloys generally degrade with the increase of exposure time and temperature, such as the coarsening of γ′ phase, the formation of TCP phase, precipitation of secondary carbides, enlarging of the defect pores and so on, which are responsible for the degradation of important mechanical properties of the alloy [3-7]. The excellent elevated temperature creep strength of Ni-based single crystal superalloys is derived from the γ′ phase with high volume fraction that is precipitated coherently from γ matrix [8, 9]. The morphology, average size, volume fraction and distribution of the γ′ precipitate [10-12] and microstructure, elastic modulus and lattice misfit of the γ/γ′ interface [13-15] play a vital key in determining mechanical properties of superalloys. The size of γ′ particle increases with exposure temperature and time during thermal exposure, which evolves in two ways: increasing its dimensions by the Ostwald ripening if the γ′ precipitate is still cubical; forming directional clusters of γ′ rafts in another case [16]. The γ′ coarsening is driven by the reduction of the γ/γ′ interface area, lattice mismatch strain of the γ/γ′ interface and modulus misfit [17]. As a result of the coarsening of γ′, the precipitate coherence is lost, and high temperature strength is significantly degraded. It has been demonstrated that the coarsening of γ′ precipitate is the main factor for the decrease of creep properties in CMSX-10 during long-term thermal exposure [17]. The single crystal alloy materials have been developed to the fourth generation, and the temperature bearing capacity has been increased to 1140-1150 ºC. However, its long-term serviceability temperature is about 1000-1100 ºC owing to the uneven heating of the blade. Considering that superalloys service prolonged periods under high temperature and complex stress, the high creep properties are essential. The favorable high temperature structural stability and 2

comprehensive mechanical properties have always been the objective pursued by researchers. The high temperature service behavior of turbine blade can be partly simulated by long-term aging, which helps to investigate the microstructural evolution the alloy [18]. This approach provides a reference for the design of the alloy and can evaluate the residual life of the alloy. The effect of long-term aging on the γ′ coarsening and creep properties of the alloy is investigated and the correlation between the γ′ coarsening and creep properties is discussed in this paper. 2. Experimental procedure The cast Ni-based superalloys containing Re and Ru was smelted via vacuum induction melting. The chemical compositions of which were shown in Table 1. The method for preparing single crystal rods was a crystal selection method with a withdrawal rate of 6 mm/min, which was performed in a directional solidification vacuum furnace ZGD-2. The single crystal rods with a deviation of [001] within 10° in the longitudinal direction were selected for the heat treatment regime: 1325 ºC/16 h + 1333 ºC/16 h, AC→1150 ºC/4 h, AC → 870 ºC/24 h, AC The creep test sticks 76 mm in length and small pieces 8-10 mm in thickness were cut from the heat-treated rods. These sticks and pieces were exposed isothermally at 1000 ºC for 100 h, 200 h, 500 h and 1000 h to study the effect of long-term aging on microstructure and creep properties of alloys. The creep test samples prepared from both the heat-treated and aged bars were carried out at 1140 ºC/137 MPa. Table 1. Chemical compositions (wt.%) of the experimental superalloys.

Element

Cr

Co

Al

Mo + W + Ta

Re

Ru

Ni

Nominal

4

12

5.9

15

5.4

3

Bal.

Measured

3.97

12.02

5.81

15.03

5.44

3.01

Bal.

The microstructure of the alloy after full heat treatment as well as long-term aging were observed by the small pieces in an INSPECT F50 scanning electron microscope (SEM). The fracture surface and longitudinal section microstructure of ruptured samples after creep rupture were also observed by SEM. To observe the morphology of γ′ precipitates clearly, the samples were grinded, polished and finally etched in a 3

solution of 20 g CuSO4 + 100 ml HCl + 5 ml H2SO4 + 80 ml H2O. Image-Pro Plus software was used to measure the average size of γ′ phase, and over 4 micrographs and three hundred γ′ particles were selected. After creep rupture, the specimens were cut into 600 µm thick disc normal to the longitudinal direction to observe the dislocation configurations by JEM 2100 transmission electron microscope (TEM). To avoid the necked region, the cutting position was about 5 mm away from the fracture surface. These foils were grinded down to abount 40 µm and then were electrochemically thinned in a solution of 10% HClO4 and 90% methanol at -27 ºC and 20 V by twin jet polisher. 3. Results 3.1. Initial microstructure The microstructure of the alloy after standard heat treatment was shown in Fig. 1(a). There was a two-phase microstructure in the samples: the cuboidal γ′ precipitate and γ matrix. The cuboidal γ′ phases with high volume fraction were uniformly distributed in the γ matrix. According to the measurement by software, the average size of γ′ was about 0.252 µm and the volume fraction of γ′ phase was about 66.53%. As shown in Fig. 1(b), it was found that the size of the majority of γ′ precipitate was about 0.2-0.3 µm in sample after full heat treatment.

Fig. 1. The microstructure of sample after full heat treatment: (a) SEM micrograph of microstructure (b) the average size and distribution of the γ′ phase.

3. 2. Microstructure after long-term aging The microstructures of samples after thermal exposure at 1000 ºC for 100 h, 200 h, 4

500 h and 1000 h were shown in Fig.2. It was obvious that the size of γ′ phase progressively increased with the aging time and the morphology of the γ′ phase kept cuboidal shape after 1000 h, which was demonstrated that the experimental alloys had preferable stability of microstructure. In addition, the width of γ channel broadened with the increase of aging time. The average γ′ size was about 0.281 µm and volume fraction of γ′ phase was about 62.95 % for 100 h, 0.310 µm and 61.65 % for 200 h, 0.351 µm and 63.81 % for 500 h, 0.405 µm and 62.01 % for 1000 h, as shown in Fig. 3. Compared with the heat-treated microstructure, the volume fraction of γ′ phase after aging decreased slightly but remained at about 62 %. With the extension of exposure time, it was found that the proportion of small γ′ phase reduced gradually and the amount of large γ′ phase constantly increased, which indicated that the growth of large γ′ phases was proceeded by annexing the small ones. As was seen in Fig. 2(b-d), the disappearing γ channel between the two adjacent γ′ phases was observed, which was related with the coarsening process of γ′ phase.

Fig. 2. SEM micrographs of the microstructure of samples after long-term aging at 1000 ºC for (a) 100 h, (b) 200 h, (c) 500 h, (d) 1000 h.

5

Fig. 3. The size and distribution of γ′ phase after long-term aging at 1000 ºC for: (a) 100 h, (b) 200 h, (c) 500 h, (d) 1000 h.

3.3. Creep behavior 3.3.1. The creep properties The creep curves of the heat-treated sample and samples aged at 1000 ºC for 100 h, 500 h and 1000 h at 1140 ºC/137 MPa were shown in Fig. 4. The creep curves exhibited three stages and the life of the second stage accounted for most of the whole creep process. The elongation of aged samples increased significantly with respect to that of heat-treated sample. An interesting observation was that the creep life of sample aged at 1000 ºC for 100 h was slightly longer than that of heat-treated sample and then the creep life was slowly decreased with the exposure time prolonging. The variation of creep life was associated with the γ′ coarsening, diffusion of the elements and the accumulation of dislocations in the γ channel. Further discussion was required to explain these phenomena. 6

Fig. 4. The creep life of heat-treated sample and samples aged at 1000 ºC for different time at 1140 ºC and 137 MPa.

3.3.2. Microstructure of samples after creep rupture The microstructure of samples after creep rupture at 1140 ºC and 137 MPa were shown in Fig. 5. It was observed that several vanishing γ matrixes embedded in the γ′ phase, as marked by triangle. The dislocation networks were observed on the γ/γ′ interface, as marked by arrows in Fig. 5(d-f). The dislocation networks observed by SEM were slightly different from those observed by TEM. Under TEM condition, the dislocation network was the boundary between the distorted and non- distorted regions, and the interfacial dislocation networks observed by SEM was some grooves. Various studies demonstrated that the dislocation networks were pivotal part in the creep process [19, 20], therefore, the dislocation networks observed by SEM would provide a complement for TEM analysis. During the early stage of creep process, the most obvious feature was a high creep rate. Microscopically, the massive a/2<101>{111} dislocations from different slip systems initiated, and the 60° mixed dislocations were left on the γ/γ′ interface because of the cross slip of the screw dislocations [21]. Since there were few 7

dislocations in the γ channel at the beginning of the creep deformation, various dislocations could be rapidly initiated and proliferated in the matrix and exhibited a high creep rate. With the progress of the creep deformation proceeded, these dislocations began to accumulate on the γ/γ′ interface, which resulted in the difficulty of dislocation movement. Eventually, the dislocation reactions of the interfacial dislocations from different slip systems occurred under the interaction of temperature, applied stress, mismatch stress and stress field between dislocations, and then gradually evolved into interfacial dislocation networks [19]. During steady state of creep, several dislocations began to cut into the γ′ rafts and then formed the superdislocations. There were two main types of superdislocation: a<101> super-dislocation (marked as region A in Fig. 6) and a<010> superdislocation (marked as region B in Fig. 6). It was generally believed that the a<101> superdislocation was composed of two a/2<101> dislocations with the same Burgers vector [22]. Moreover, the a<101> superdislocation was a superdislocation pair and the spacing of them was determined by the antiphase boundary (APB) energy. The a<010> superdislocation was formed by meeting and reaction of the two a/2<011> matrix dislocations with different Burgers vectors at the γ/γ′ interface [23, 24]. Morphologically, the a<010> superdislocations were long and straight. The tertiary creep stage was characterized by a dramatically increasing creep strain and then ruptured within a short time as a result. The interfacial dislocation networks were no longer steady, therefore, the dislocations were clearly bowed and moved in this stage. In addition, the casting defect pores greatly affected the final fracture of the alloy. The microcracks that caused the fatal damage on the alloy mostly initiated at these pores and propagated rapidly in the third stage of creep. The dislocation configurations of the creep ruptured specimens with different exposure time were showed in Fig. 6. The characteristic of the dislocation configurations was that the dislocation networks were formed at the γ/γ′ interface in all exposed samples, regardless of the exposure time. The dislocation networks were formed by different Burgers vectors at the γ/γ′ interface, as shown in Fig. 5(d-f). It was observed that the interfacial dislocation networks of all samples ruptured after creep deformation were regular and dense. The superdislocations cutting into the γ′ phase was 8

a main restrict factor that leaded to creep rupture [25, 26]. These dislocation networks acted as an obstacle to effectively hinder the dislocations from cutting into γ′ phase and forming the superdislocations [27].

Fig. 5. SEM micrographs of samples aged at 1000 ºC for (a, d) 100 h, (b, e) 500 h and (c ,f) 1000 h ruptured after creep deformation under 1140 ºC/137 MPa: (a-c) the γ/γ′ microstructure and (d-f) the interfacial dislocation networks.

Fig. 6. Dislocation configurations of samples aged with different time at 1000 ºC ruptured after creep at 1140 ºC/137 MPa: (a) standard heat treatment, (b) 100 h, (c) 500 h, (d) 1000 h.

3.3.3. The pores in samples after creep rupture 9

Fig.7(a-c) showed the fracture morphologies of aged samples, the micrographs illustrated that the fracture mode of the alloy was micropore aggregation fracture. The process of the fracture mode was that the microcracks formed firstly at the edge of the micropores and then propagated under stress, which leaded to the fracture of the alloy. A mass of pores and microcracks were observed on the edge of fracture surface in all samples after long-term aging, as showed in Fig. 7(d-f). The fracture surface of single crystal superalloys after high temperature creep existed two types of pores: casting pores and creep pores [28]. During the solidification of the alloy, the casting pores were generally formed in the interdendritic area. The creep pore was a special type of pore formed in the process of creep deformation, and it nucleated on the γ/γ′ interface and constantly grew throughout the creep process. Obviously, the casting pores were more detrimental to fracture of the alloy than the creep pores. The microcracks caused by creep pores were much finer, which could not cause fatal damage to the alloy. The creep strain increased rapidly at the beginning of the tertiary creep, and then the applied stress increased on account of the reduction of effective area. The microcracks initiated at the corners of these pores in the necked regions of samples, then continuously propagated and interconnected under stress, resulting in the final fracture of alloys.

10

Fig. 7. SEM images of samples aged at 1000 ºC for (a, d) 100 h, (b, e) 500 h and (c, f) 1000 h ruptured after creep at 1140 ºC/137 MPa: (a-c) fracture morphologies; (d-f) fracture surfaces; (g) distribution of the pores, enlarged images of (h) casting pores and (h) creep pores.

4. Discussion 4.1. The coarsening of γ′ phase The coarsening of γ′ phase is driven by the decrease of interface area and lattice mismatch strain of the γ/γ′ interface, and a decrease in modulus misfit [17]. It is believed that if its shape is still regular, the γ′ phase will increase its dimensions via the Ostwald ripening, which is a process controlled by the diffusion [29, 30]. As mentioned above, the growth of large γ′ phase is proceeded by merging with the adjacent small one. The Gibbs-Thomson theorem is the thermodynamic basis for studying the Ostwald ripening. The theorem builds a relationship between the radius of curvature of the interface and equilibrium concentration of solute atoms near the interface [31], which can be described as Eq. (1). 11

=

1+

Where

and

∙

(1) are the equilibrium concentration of particles with radius

and

radius of infinity in the matrix,

is interfacial energy of the γ/γ′ interface,

is the

molar volume of γ′ precipitate,

is the ideal gas constant (8.314 J⋅K-1mol-1),

is the

absolute temperature. The Eq. (1) shows that the concentration of solute around the small particles is greater than the concentration of solute around large particles, so solute atoms around small particles will diffuse to the vicinity of large particles. The diffusion will destroy balance of the solute concentration around the particles, resulting in the dissolution of small particles and the enlargement of large particles. Lifshitz, Slyozov [32] and Wagner [33] proposed a mathematical model describing Ostwald ripening, whose theory is called the LSW theory. The coarsening process of γ′ phase on time can be described by the following Eq. (2). − =

$

Where

=

(2)

!" #

(3)

is the average radius of γ′ particles after full heat treatment,

radius of γ′ phases after aging for time ,

is the average

is the LSW theory coarsening rate constant,

which can be expressed as Eq. (3) and D is the diffusion coefficient of the specified phase. Since the γ′ particles keep cuboidal shape after aging at 1000 ºC for 1000 h, the average radius of γ′ particles r =a/2, wherea is the average cube edge width. Fig. 8 plots the relationship between the average radius of the γ′ precipitate and aging time. The relationship between

−

and

is an approximately linear line, which

demonstrates that the γ′ coarsening follows the LSW theory in this work. The coarsening rate

is 5954 nm3/h in accordance with the slope of line.

According to Eq. (2) and Eq. (3), it may conclude that at any moment in the Ostwald ripening process, there is a critical particle radius r>

%

grow, particles with r <

%

%,

particles with radius

dissolve, invariant particles with r =

%.

It is

depicted in Fig. 3 that the proportion of the large size of γ′ precipitate continuously increases and the proportion of the small γ′ phase decreases with the extension of aging 12

time. Schematic illustration for the coarsening of γ′ phases during long-term aging is plotted in Fig. 9. As shown in Fig. 2(b-d), it is found that the γ channel between a large γ′ phase and a small γ′ phase is gradually disappearing during long-term aging. One reason for vanishing channel is that the γ channel is covered by moving dislocations. Another is that the γ channel slowly dissolves by virtue of the diffusion of elements during aging. Therefore, a larger particle is formed by the merge of the two adjacent γ′ particles, increasing of the average size of γ′ phase and the width of γ channel.

Fig. 8. Effect of long-term aging on average size of γ′ precipitate of the alloy: the correlation between ( − ) and aging time.

Fig. 9. Schematic illustration for coarsening process of γ′ phase during long-term aging. 13

4.2. The correlation between the γ′ coarsening and creep properties The high temperature strength of superalloys is directly derived from the ordered γ′ phases that is the main strengthening phase of superalloys [8]. The morphology, size, volume fraction and distribution of γ′ precipitate play a crucial key role in improving creep life of the alloy because the γ′ phase can act as an effective obstacle for preventing the deformation of the alloy and movement of dislocations [10-12, 34]. It is shown in Fig. 2 that the γ′ phase keeps cuboidal shape in sample after aging at 1000 ºC for 1000 h. It has been demonstrated that the creep life of Ni-based single crystal superalloys increases at first and then decreases with the increase of the volume fraction of γ′ precipitate [9]. When the volume fraction of γ′ phase is between 60 % and 70 %, the alloy attains the best creep properties. In this work, the volume fraction of γ′ phase of aged samples was about 62-66 %. Moreover, the previous study [27] has suggested that the alloy with dense and regular interfacial dislocation networks has superior creep properties. The interfacial dislocation networks of all aged samples after creep deformation were regular and dense in this work (Fig. 6). Therefore, the size of γ′ precipitate is the main factor affecting the creep properties of alloys after aging. Neumeier et al. [35] have concluded that single crystal superalloys containing Re and Ru with a γ′ phase size of about 0.32 µm attain optimum creep properties at 1100 ºC/137 MPa. Generally, when the size of γ′ phase is much small, the most of γ′ phase are spherical shape and not well aligned. The investigation by Nathal [36] has demonstrated that the γ′ precipitates tend to change into cuboidal shape with the coarsening of γ′ phase, which helps to minimizing the coherency stresses. The more orderly rafting microstructure with fewer terminations of γ′ rafts per unit area can be formed by these cubic precipitates, which can enhance the creep resistance. Therefore, the alloy with an optimum γ′ size has better creep properties than the alloy with a γ′ size below the optimum size. However, when the γ′ size exceeds the optimum γ′ size, the creep life of the alloy decreases. For one thing, the width of the γ channel gradually increases with the coarsening of γ′ phase during long-term thermal exposure. When the applied stress is not enough to activate the Orowan bypass mechanism or the 14

dislocation cutting mechanism, dislocations can only overcome the obstacle of the γ′ phase by thermally activated climbing under high temperature and low stress. Brown and Ham [37] proposed that the slip dislocations can bypass the γ′ phase through the local climb or general climb, and the threshold stresses of two ways [38] (*+% and *,% ) are closely related to the threshold stress for Orowan bowing [39] (*-. ): *+% = *-. ⁄√2 *,% = 2 *-. =

.4

*-. ⁄2

5.

In 67

(4) . 4

(5)

7

(6)

:

?

λ = d =2 >@ − 1A

(7)

where 2 is the volume fraction of γ′ phase, G is the shear modulus, b is the dislocation Burgers vector,

is the dislocation core radius, B is the interparticle distance, d is

the average edge length of γ′ particles. According to Eq. (6) and (7), the volume fraction 2 can be regarded as a constant, the γ′ size d enlarges with the extension of exposure time, the γ channel width B continuously increases, thus the Orowan stress decreases. Therefore, the threshold stress of dislocation climb is also decreased. The dislocations are more likely to bypass the obstacle of the γ′ phase in a short time, reducing creep strength of alloys. Besides, the number of the γ/γ′ interface decreases with the enlargement of γ′ phase size. The proportion of interfacial dislocation networks is also reduced, which will decrease strengthening effect of interfacial dislocation networks on the alloy. As shown in Fig. 4, it is interesting to find that the alloy after aging for 100 h has a superior creep life than the alloy after standard heat treatment. There are several reasons for this phenomenon. For one thing, with the coarsening of γ′ phase, the γ′ size gradually approaches the optimum γ′ size for creep properties of experimental alloys. Consequently, the alloy after aging for 100 h has a longer creep life compared with the alloy with initial microstructure. For another, the experimental superalloys are mainly enhanced by a large number of refractory elements, but the diffusion rate of Ru and Re elements in the alloy is much low, resulting in the inhomogeneous diffusion in alloys 15

after full heat treatment. Thus, the refractory elements continue to diffuse in the alloy during the early stage of aging at 1000 ºC, which can reduce or eliminate segregation. The solution strengthening and precipitation strengthening effects are improved by the diffusion. In addition, it is well known that the alloy can obtain high strength through plastic deformation, which is called work hardening. During the plastic deformation, the movement of dislocations is impeded by the increase of dislocation density and the entanglement of dislocations, which is an important cause of work hardening. Accordingly, after aging at 1000 ºC for a while, the misfit dislocation may accumulate in the γ channel, which can act as dislocation source and strengthen the γ/γ′ interface to hinder the motion of dislocations. And the interfacial dislocation network in 100 h sample after creep rupture are relatively more regular than the other aged samples. For the above reasons, the creep life of sample after aging for 100 h is longer than that of the heat-treated sample. With the exposure time prolonging, the size of γ′ phase and the width of γ channel increase continuously, which leads to a decrease in the resistance of slip dislocations to bypass the obstacles of γ′ phase. Besides, with the coarsening of γ′ phase, the number of the γ/γ′ interface in samples reduces, further decreasing the strengthening effect of interfacial dislocation networks on the γ/γ′ interface. Therefore, the creep life of the alloy is progressively reduced. 5. Conclusion (1) After long-term aging at 1000 ºC, the average size of γ′ phase increases with the aging time and the morphology of γ′ phase keeps cuboidal shape after 1000 h. It is found that the γ′ coarsening of the experimental alloy is consistent with the LSW theory and the coarsening rate constant

is about 5954 nm3/h.

(2) Two types of the pores: casting pores and creep pores are observed on the fracture surface after creep deformation. Compared with the creep pores, the microcracks caused by the casting pores are more detrimental to the creep life of alloys. (3) After aging at 1000 ºC for 100-1000 h, the creep life at 1140 ºC/137 MPa increases firstly and then decreases with the extension of aging time. The creep properties of alloys during long-term aging are mainly affected by the size of γ′ 16

phase. (4) The increase of creep life of aged sample is attributed to the slightly grown γ′ phase whose size is close to the optimum γ′ size corresponding to the creep properties, the diffusion of refractory elements and the accumulation of dislocations in the γ channel. The seriously coarsened γ′ phase leads to the decrease of the resistance of dislocation movement, deteriorating the creep properties. Acknowledgements The financial support provided by National Science and Technology Major Project (2017-VI-0002-0072), National Key R&D Program of China under Grant No. 2017YFA0700704, National Natural Science Foundation of China (NSFC) under Grant Nos. 51601192, 51671188, State Key Lab of Advanced Metals and Materials Open Fund under Grant No.2018-Z07. References [1] M.J. Donachie, M. American Society for, Superalloys, American Society for Metals, Metals Park, Ohio, 1984. [2] J. Smialek, G.M. Meier, High Temperature Oxidation in Superalloy, High-temperature oxidation Superalloys II; 2 (1987). [3] R.A. Stevens, P.E.J. Flewitt, The effects of γ′ precipitate coarsening during isothermal aging and creep of the nickel-base superalloy IN-738, Materials Science and Engineering 37(3) (1979) 237-247. [4] W. Sha, Quantification of age hardening in maraging steels and an Ni-base superalloy, Scripta Materialia 42(6) (2000) 549-553. [5] G. Lvov, V.I. Levit, M.J. Kaufman, Mechanism of primary MC carbide decomposition in Ni-base superalloys, Metallurgical and Materials Transactions A 35(6) (2004) 1669-1679. [6] Q.Z. Chen, N. Jones, D.M. Knowles, The microstructures of base/modified RR2072 SX superalloys and their effects on creep properties at elevated temperatures, Acta Materialia 50(5) (2002) 1095-1112. [7] A.K. Koul, R. Castillo, Assessment of sercive induced microstructural damage and ITS rejuvenationI in turbine-blades, Metallurgical Transactions a-Physical Metallurgy and Materials Science 19(8) (1988) 2049-2066. [8] H. Long, H. Wei, Y. Liu, S. Mao, J. Zhang, S. Xiang, Y. Chen, W. Gui, Q. Li, Z. Zhang, X. Han, Effect of lattice misfit on the evolution of the dislocation structure in Ni-based single crystal superalloys during thermal exposure, Acta Materialia 120 (2016) 95-107. [9] T. Murakumo, T. Kobayashi, Y. Koizumi, H. Harada, Creep behaviour of Ni-base single-crystal superalloys with various γ′ volume fraction, Acta Materialia 52(12) (2004) 3737-3744. 17

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Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: