Microstructural evolutions and fracture behaviors of a newly developed nickel-base superalloy during creep deformation

Accepted Manuscript Microstructural evolutions and fracture behaviors of a newly developed nickel-base superalloy during creep deformation Luqing Cui, Jinjiang Yu, Jinlai Liu, Xiaofeng Sun PII:

S0925-8388(18)30790-4

DOI:

10.1016/j.jallcom.2018.02.295

Reference:

JALCOM 45171

To appear in:

Journal of Alloys and Compounds

Received Date: 20 November 2017 Revised Date:

23 February 2018

Accepted Date: 25 February 2018

Please cite this article as: L. Cui, J. Yu, J. Liu, X. Sun, Microstructural evolutions and fracture behaviors of a newly developed nickel-base superalloy during creep deformation, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.02.295. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.

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Microstructural evolutions and fracture behaviors of a newly developed nickel-base superalloy during creep deformation Luqing Cui a,b, Jinjiang Yu a,*, Jinlai Liu a, Xiaofeng Sun a Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China

b

School of Materials Science and Engineering, University of Science and Technology of

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a

China, Hefei 230026, China

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ABSTRACT: The creep behaviors of the newly designed nickel-base superalloy M951G

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have been studied over a stress range of 120 to 760 MPa and a temperature range of 700
to 1000
, and the corresponding relationship among testing conditions-microstructural evolutions and fracture characteristics-creep properties has been established. Results show that both the microstructural evolutions and fracture characteristics are dependent on testing

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temperatures and applied stresses. According to X-ray diffraction (XRD) examinations, creep deformations at 700-800 °C and high temperatures low stresses have not any distinctive effect on the acquired patterns. However, deformations at high temperatures relatively higher

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stresses have significantly highlighted the (220)γ&γ′ planes, which may enhance the creep

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strength and lead to a better plasticity for M951G alloy. Furthermore, precipitate dimensional changes and their effects on creep properties also have been widely discussed. At the same testing temperature, creep life of M951G alloy decreases with applied stress increasing, which has been analyzed in detail. Fracture behaviors of M951G alloy are observed by using optical microscope (OM) and scanning electron microscope (SEM), which change from a mixed type of transgranular and intergranular to pure intergranular cracking with the variation of creep

*

Corresponding author. Tel.: +86-24-2397-1713; fax: +86-24-2390-6722; e-mail: [email protected]

ACCEPTED MANUSCRIPT conditions. The values of apparent stress exponents at 700 °C and 1000 °C are calculated to be about 11.5 and 6.4 respectively, the difference of which is due to the changes of the deformation mechanisms, fracture modes and the degeneration of microstructures.

Precipitate dimensional changes; Crystallographic changes

1. Introduction

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Keywords: M951G alloy; Creep behaviors; Microstructural evolution; Fracture modes;

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Nickel-base superalloys are critical materials for turbine blades and vanes in the aircraft

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engines and industries gas engines owing to their outstanding mechanical properties and good corrosion resistance at elevated temperatures [1-6]. The excellent properties of superalloys are mainly due to the precipitation strengthened by γ′ precipitates, solution strengthened by plenty of alloying elements and the microstructure of coherent γ/γ′ interface [1, 7-13].

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With the increase of the turbine engine thrust-weight ratio, the high temperature creep resistances of superalloys need to be improved urgently [14]. Furthermore, during the actual service conditions, the hot components of vanes in engines withstand many extreme

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conditions, among which the major ones are temperature and stress [5, 6, 12, 15, 16].

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Therefore, the high temperature creep property has become one of primary criteria for engineers to evaluate the mechanical properties of the newly developed superalloys. Extensive researches have been conducted to study the creep fracture behaviors of superalloys previously [7, 9, 14, 17, 18]. With the increase of applied stress from 670 MPa to 897 MPa at 700
, the creep fracture mode of a LSHR superalloy changed from intergranular to transgranular [14]. The effect of phosphorus on a newly developed Ni-Fe-Cr base wrought superalloy has been studied by Guan [19], which showed that it took longer time to form 2

ACCEPTED MANUSCRIPT comparable intergranular cracking area with the phosphorus content increasing from 0.002 wt% to 0.027 wt% during creep deformation at 750
/250 MPa. The improving of the grain boundary oxidation resistance owing to the segregation of phosphorus on the grain boundary

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might be responsible for it [19]. Besides, Ref. [19] illustrated that for the newly designed Ni-Fe-Cr base superalloy at 750
, the intergranular fracture mode transformed into transgranular with the increase of strain rate. In addition, the relationship between grain size

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and creep fracture behaviors of superalloy also has been studied. With the refinement of grain

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size, the carbides and eutectics were refined, so the probability of generating microcracks was decreased and creep life had been prolonged [20]. Apart from applied stress, alloying element, strain rate and grain size, the influence of film-hole configuration on the high temperature creep rupture behaviors of the superalloys was also investigated at 980 °C/300 MPa [13].

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Experimental results showed that one- and two-row specimens exhibited longer lives than those without film-holes. However, the creep rupture life decreased with the increase of film-hole rows. Stress concentration and local stress changing from uniaxial to multi-axial

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caused by the existence of film-holes resulted in the variation of creep lives [13]. On the other

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hand, during creep deformation microstructural evolutions such as the average size and morphology of γ′ precipitates [21], the lattice misfit between γ and γ′ phases [22], the grown-up of casting micropores and formation of creep micropores [7] as well as the density and configurations of dislocations in microstructures [6, 16, 23, 24] also change with the creep process going on. Accordingly, these changes will have a great impact on the creep properties of superalloys. As mentioned above, although the creep behaviors of superalloys have been widely 3

ACCEPTED MANUSCRIPT studied, previous studies mainly focus on a certain or a small range of temperatures and stresses on the creep behaviors of superalloys. Few references have been reported related to the effects of a wide range of temperatures and stresses on the creep fracture behaviors and

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microstructural evolutions of an identical nickel-base superalloy. It will certainly to make a contribution for better understanding of the creep performances of superalloys. In present work, the newly developed nickel-base superalloy M951G alloy [7, 25, 26] is chosen to

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systematically study the creep behaviors over a stress range of 120 to 760 MPa at the

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temperatures ranging from 700 to 1000
. The purpose of this work is to investigate the fracture behaviors of M951G alloy with the change of a large range of testing conditions during creep deformation. Furthermore, the influence of microstructural evolutions on creep properties of M951G alloy also has been analyzed in detail.

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2. Experimental procedures

The nominal chemical composition of M951G alloy is listed in Table1. The master alloy was remelted up to 1450 °C in a vacuum induction furnace and then poured into the

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preheating mode to form the creep sample bars. And the related heat treatments were

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conducted on the alloy as follows: 1210 °C×4 h, AC→1100 °C×4 h, AC→870 °C×24 h, AC (AC: air cooling).

The specimens for creep tests were machined from the heat-treated bars with a diameter

of 5 mm and a gauge length of 25 mm. The tensile creep tests were carried out over a wide stress range of 120 to 760 MPa under temperatures from 700 °C to 1000 °C, as listed in Table 2. During the creep tests, the temperature variation of each specimen was controlled within ±2 °C by using three thermocouples attached to the gage part. Two or more identical 4

ACCEPTED MANUSCRIPT specimens were tested at each creep condition. The details of creep tests are critical important for the measure of creep properties. Therefore, the detailed information about creep process should be systematically introduced. We selected 1000
/120 MPa as the representative

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testing condition to illustrate the creep deformation process in present study. The black line represents the applied stress and the red line represents the testing temperature, as displayed in Fig. 1. At first, a very small constant stress (10 MPa) is applied on the specimen in order to

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keep the specimen fixed. Meanwhile, the testing temperature begins to rise gradually until the

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required experimental temperature (1000
) is achieved, and then the specimen is kept at 1000
for 1h, as shown in Fig. 1. After that, the required experimental stress (120 MPa) is applied on the specimen until fracture (which means the specimen was deformed at 1000
/120 MPa until fracture). Finally, the ruptured specimen is cooling down to room

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temperature in the air.

The microstructures before and after creep tests were observed by the Zeiss Axio Observer ZIm OM and INSPECT F50 SEM. The INSPECT F50 SEM was equipped with an

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Oxford energy dispersive spectrometer (EDS), which was used to qualitatively identify the

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chemical compositions of precipitated phases in M951G alloy. The specimens for OM and SEM observation were ground, polished and chemically etched in a solution of 20 ml HCl + 5 g CuSO4 + 25 ml H2O. In order to study the crack propagation path, the longitudinal sections of the fracture samples also have been carefully observed. X-ray diffraction technique was also extensively utilized for creep deformed and heat treatment samples in order to assess the crystallographic transformations. The area fractions and mean sizes of carbides and eutectics in the heat treatment microstructure were measured and calculated by Image-Pro Plus 5

ACCEPTED MANUSCRIPT software (IPP). In addition, the volume fraction and average edge length of γ′ precipitates after heat treatment were also measured by IPP. Foils for transmission electron microscopy (TEM) observation were cut normal to the orientation of applied stress from the post-creep

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specimens 5 mm away from the fracture surface. These foils were mechanically thinned to less than 50 µm, and then electro-polished using a twin jet polisher in 10% perchloric acid ethanol solution at -25
, 42 mA. The resulting foils were observed using a JEM 2100

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transmission electron microscopy at 200 kV.

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Table 1 Nominal compositions of M951G alloy (wt%). W

Mo

Nb

Co

6.5

3

2.2

5

Table 2

Cr

Al

Hf

others

Ni

9

6

1.5

0.091

Bal

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The design of creep conditions in present study. Temperature/


800 900

120

150

160

200

250

200

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1000

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700

6

Loading Stress/MPa 650 450 300

360

450

550

650

700

760

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Fig. 1. The detailed information of tensile creep deformation in present study.

3. Results

3.1. Microstructures before creep deformation

Fig. 2 shows the typical microstructures of M951G alloy after heat treatment. The grain,

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grain boundary, eutectic, carbide and γ′ precipitates of M951G alloy are carefully observed by OM and SEM. The distribution of grains in the cross section of the heat-treated bar is

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displayed in Fig. 2a.The statistical grain size of M951G alloy is about 3.273 mm (the standard deviation is 1.435 mm), which means the grain boundary content in present alloy is very low.

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Serrate grain boundary is obtained in the heat-treated microstructures, which still maintain a pronounced dendritic structure, as shown in Fig. 2b. Fig. 2c shows the distribution of eutectics in the microstructure, most of which are located in the interdendritic regions, only a small part along the grain boundaries. The MC carbides, which are identified to be NbC and HfC by EDS (Energy Dispersive Spectroscopy), are large block-like precipitating in the interdendritic regions (Fig. 2d) and chain-like along the grain boundaries (Fig. 2e). The γ′

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in the air and grew during high

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temperature aging (1100
×4 h). The tiny spherical γ′ particles precipitated during air cooling after high temperature aging (1100
) and then grew when aged at 870
for 24 hours. Both the coarse and tiny γ′ particles can improve the strength of superalloy and stabilize

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microstructure by pinning dislocations, grains and grain boundaries [27-29]. In addition, the

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statistical area fractions and the mean sizes of eutectics and carbides along the grain boundary and in the grain interior are also calculated by IPP, which are listed in Table 3. It can be clearly seen from the Table 3 that most of the eutectics and carbides are distributed in the interdendritic regions. Furthermore, the average carbide size in the interdendritic regions is

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much larger than that along the grain boundaries. While the mean size of the eutectics nearly

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the same whether in the grain interior or on the grain boundaries.

Fig. 2. Typical microstructures of M951G alloy after heat treatment. (a) Distribution of grains 8

ACCEPTED MANUSCRIPT on the cross section of the heat-treated bar. (b) Serrate grain boundary. (c) Eutectic distribution in the microstructure. (d) Large and block-like carbides in the interdendritic regions. (e) Tiny and chain-like carbides along the grain boundaries. (f) Coarse cuboidal and

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fine spherical γ′ precipitates in the microstructures. Table 3

interior and along the grain boundaries. Carbides

in

Carbides along the

Eutectics in interdendritic

Eutectics along the

grain boundary

regions

grain boundary

Area fraction (%)

1.11±0.04

Average size (µm)

4.66±2.08

3.2. Creep properties

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M951G alloy interdendritic regions

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The statistical area fractions and the average sizes of eutectics and carbides in the grain

0.35±0.03

4.51±0.47

2.96±0.36

0.62±0.25

25.72±8.14

26.70±10.69

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Fig. 3 shows the creep curves of M951G alloy at different testing conditions. All creep curves consist of the primary, steady-state and accelerated stages. Although the accelerated stage is limited in all creep curves, the lengths of the primary and steady-state stages

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dramatically decrease with the increase of applied stress. In addition, the creep life,

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elongation and minimum creep rate for all the specimens at various testing conditions are listed in Table 4. It can be clearly seen that at the same testing temperature with the increase of applied stress, the minimum creep rate is increased, while the creep life is decreased. Apart from that, M951G alloy exhibits relatively better ductility at higher temperatures. For metals and alloys creep at high temperatures, the relationship between the steady-state creep rate (ߝሶ௦௦ ) and applied stress (ߪ௔ ) at a given temperature (T) can be described as follows [30]: 9

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ߝሶ௦௦ =Aߪ௔௡ exp(-ୖ୘)

(1)

where A is a constant related to the given material. R is the gas constant. Q is the apparent creep activation energy, the value of which is close to that for self-diffusion. n is the apparent

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stress exponent, which is between 4 and 6 for simple alloys. For M951G alloy, as shown in Fig. 3e, different values of apparent stress exponents have been obtained at 700
and 1000
. Furthermore, the values of n at 700
(11.594), 800
(11.110) and 900
under relatively

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higher applied stresses (10.699) are nearly the same. Similarly, close values of the apparent

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stress exponents are obtained when creep deformations at 900
under low applied stresses (5.427) and at 1000
(6.381). The variation of apparent stress exponents with the changes of creep conditions will be analyzed in detail in Section 4.5.

In order to compare the difference in minimum creep rate with the testing conditions, the

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relationship between the steady-state creep rate and time to rupture of M951G alloy has been presented in Fig. 3f. For nickel-base superalloy creep deformation at high temperatures, the correlation of the minimum creep rate and creep life (‫ݐ‬௙ ) is commonly described by

(2)

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௠ ߝሶ௦௦ ‫ݐ‬௙ =c

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Monkman-Grant law [31]:

where m and c are constants. m is around 0.8-1.7, and c is around 0.07-1.7. This relationship is independent of temperature, applied stress and chemical compositions of alloy, which can be used to estimate the long creep life from the minimum creep rate obtained by short time creep test. In present study, the Monkman-Grant relationship employed for M951G alloy is shown in Fig. 3f. When specimens are deformed at the same minimum creep rate, with the increase of testing temperature, the creep life of M951G alloy distinctly decreases. It 10

ACCEPTED MANUSCRIPT illustrates that thermal activation at high temperatures has significant impact on creep properties of superalloy. On the other hand, the data points at each temperature distribute in a

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well straight line, and data of m and c at different temperatures are listed in Table 5.

Fig. 3. (a)-(d) Creep curves of M951G alloy at 700, 800, 900 and 1000
respectively. (e) The

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variation of apparent stress exponents for M951G alloy with the changes of testing temperatures. (f) The minimum creep rate of M951G alloy as a function of creep life under

Table 4

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different testing conditions.

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The creep life, elongation and minimum creep rate for all the specimens at various testing conditions.

Temperature(
)

Applied stress(MPa)

Creep life(h)

Elongation(%)

minimum creep rate(h-1)

700

650

693±36

2.24±1.26

(1.39±0.21)*10-5

700

700

274±21

2.93±1.42

(6.25±0.08)*10-5

700

760

155±13

2.01±1.63

(8.55±0.85)*10-5

800

450

757±24

5.84±1.20

(3.76±0.17)*10-5

800

550

121±14

5.67±2.31

(2.74±0.18)*10-4

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20±5

6.46±1.53

(2.27±0.11)*10-3

900

200

787±33

3.29±1.56

(2.81±0.29)*10-5

900

250

333±17

4.78±1.05

(6.76±0.06)*10-5

900

300

125±13

5.50±2.07

(2.08±0.10)*10-4

900

360

37±4

5.61±1.34

(8.75±0.07)*10-4

900

450

10±3

6.25±2.54

1000

120

294±32

4.13±1.72

1000

150

134±13

5.81±1.52

1000

160

74±11

6.86±1.15

1000

200

20±1

6.77±0.52

(7.34±0.16)*10-3

(5.80±0.32)*10-5 (2.13±0.15)*10-4

(5.23±0.38)*10-4

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Table 5

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800

(1.42±0.17)*10-3

Creep parameters in the Monkman-Grant relationship for M951G alloy under different temperatures. 700


m

0.7984

c

0.2068

3.3. Fractographs

800


900


1000


0.8862

0.7925

0.8523

0.2924

0.2358

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Parameter

0.2617

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3.3.1. Microstructures after creep deformation at 700

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The typical fractographs of M951G alloy for tensile creep tests at 700
under different applied stresses are shown in Fig. 4. As can be seen in Fig. 4a-4c, the rupture mode is cleavage like on the whole. Some slip planes can be seen on the fracture surfaces, which are considered to be formed by the slipping of dislocations or stacking faults on ሼ111ሽ planes in nickel-base superalloy [32]. In addition, microcracks also have been observed in the vicinity of the slip planes, as displayed in Fig. 4d-4f, illustrating that there is a certain relationship between the formation of microcracks and the slipping of dislocations on ሼ111ሽ planes.

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ACCEPTED MANUSCRIPT Apart from that, a great many broken carbides are observed on the fracture surfaces (as shown in Fig. 4g-4i), which suggest that the formation of microcracks at 700
may also relate to the broken carbides. The longitudinal sections of the fracture samples are also observed by SEM

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and OM in order to study the crack propagation path, as displayed in Fig. 4j-4o. It can be found that the propagation of the cracks is a mixed type of transgranular and intergranular for

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mode increases with the enhancement of applied stress.

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all the specimens tested at 700
. It should be noted that the proportion of the transgranular

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Fig. 4. The typical fractographs of M951G alloy for tensile creep tests at 700
.

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(a)(d)(g)(j)(m), (b)(e)(h)(k)(n), (c)(f)(i)(l)(o) Specimens tested at 700
under 650, 700 and 760 MPa, respectively.

3.3.2. Microstructures after creep deformation at 800 Compared with the fracture surfaces of M951G alloy at 700
, similar deformation

microstructures are observed when tensile creep at 800
, as shown in Fig. 5. Many broken carbides (Fig. 5c and 5h) as well as the mixed fracture mode of transgranular and intergranular (Fig. 5d, 5e, 5i and 5j) are also observed. It is worth noting that the proportion of 13

ACCEPTED MANUSCRIPT the transgranular mode also increases with the enhancement of applied stress, as displayed in Fig. 5e and 5j. However, there is one obvious difference should be noted that a number of small dimples are formed on the fracture surfaces (Fig. 5b and 5g), which reveals that the

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fracture mode changes from brittle cleavage fracture to ductile quasi-cleavage type fracture

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with the increase of testing temperature.

Fig. 5. The typical fractographs of M951G alloy for tensile creep tests at 800
. (a)-(e) 800
/450 MPa. (f)-(j) 800
/650 MPa.

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3.3.3. Microstructures after creep deformation at 900 The typical fractographs and longitudinal sections of the fracture samples for M951G

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alloy after creep deformation at 900
under different applied stresses are shown in Fig. 6. Similar to the fracture characteristics obtained at 800
(Fig. 5), the deformation

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microstructures of M951G alloy after creep test at 900
under relatively higher stress (450 MPa) consist in a great number of broken carbides (Fig. 6h), a certain number of dimples (Fig. 6i) and a few micropores (Fig. 6g). However, with the decrease of applied stress, significantly different deformation microstructures are observed after creep deformation at 900
/200 MPa, as presented in Fig. 6a-6d. A considerable number of micropores (Fig. 6b) and dimples (Fig. 6d) existed on the fracture surface, and no broken carbides were observed. On the other hand,

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may be related with the interlinking of the coarse micropores.

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Fig. 6e and 6j show the longitudinal sections of the fracture samples, which are used to observe the fracture modes of M951G alloy after different creep tests. It can be clearly seen that the fracture mode after creep deformation at low applied stress (200MPa) is pure

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intergranular cracking (Fig. 6e), while it changes to a mixed fracture mode of transgranular

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and intergranular as the applied stress increases to a high level (450 MPa), as shown in Fig.

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6j.

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Fig. 6. The typical fractographs of M951G alloy for tensile creep tests at 900
. (a)-(e) 900
/200 MPa. (f)-(j) 900
/450 MPa.

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3.3.4. Microstructures after creep deformation at 1000 Fig. 7 shows the fracture surfaces and the longitudinal sections of the fracture samples at

1000
. Different to the flat fracture surfaces after creep deformation at 700
, the fracture surfaces of M951G alloy is relatively rougher at 1000
. Various degrees of necking deformations occur after high temperature creep tests (Fig. 7a and 7d). In addition, a great number of large dimples are formed on the fracture surfaces, and some of the micropores are

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ACCEPTED MANUSCRIPT linked together to form the microcracks (Fig. 7b and 7e). Therefore, it can be speculated that the tensile creep fracture mode is microvoid accumulation type. Apart from that, as shown in Fig. 7c and 7f, a number of coarse micropores can be seen in the vicinity of grain boundaries,

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and the fracture mode is pure intergranular cracking under all applied stresses at 1000
, which is the same as that of M951G alloy after creep deformation at 900
under low applied

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stress (Fig. 6e).

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Fig. 7. The typical fractographs of M951G alloy for tensile creep tests at 1000
. (a)-(c) 1000
/120 MPa. (d)-(f) 1000
/200 MPa.

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4. Discussion

4.1. The crystallographic changes Fig. 8 illustrates the XRD profiles of M951G alloy in different states. It can be seen that

both the heat-treatment and deformed samples consist of γ, γ′ and MC (M=Hf, Nb) carbide. Correspondingly, the summary of the intensity ratios (I/Imax) for all crystallographic planes in M951G alloy after creep deformations under different conditions is listed in Table 6. Before

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ACCEPTED MANUSCRIPT creep test, (111)γ&γ′ planes are the main slip planes in the heat-treatment specimen (Fig. 8a). After fracture at low and intermediate temperatures (700-800 °C), as shown in Fig. 8b-8e, creep deformation neither has obvious influence on (111)γ&γ′ nor on other ones. Similar

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phenomenon also has been observed at high temperatures low applied stresses (900 °C/200 MPa and 1000 °C/120 MPa), as demonstrated in Fig. 8f-8h. However, deformations under high temperatures relatively higher applied stresses (900 °C/450 MPa and 1000 °C/200 MPa)

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have significantly highlighted the (220)γ&γ′ planes, as illustrated in Fig. 8g-8i. The

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augmentation of the (220)γ&γ′ planes of M951G alloy at 900 °C/450 MPa and 1000 °C/200 MPa illustrates that greater extent of grain rotation occurs during creep deformations under high temperatures relatively higher applied stresses. In addition, since (220) plane is not the close-packed plane in face-centered cubic metal, the resistance for dislocations to slip on (220)

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plane is much higher than that on (111) plane, which means the creep strength of M951G alloy can be enhanced by the higher degree activation of (220) planes. On the other hand, the stress concentration along the grain boundaries can be easier released by the higher extent

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activation of secondary slip systems at higher applied stress, and it may lead to a better

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plasticity for M951G alloy, which is consistent with the higher creep elongations measured at 900 °C/450 MPa and 1000 °C/200 MPa (as shown in Table 4).

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Fig. 8. X-ray diffraction profiles of M951G alloy after creep deformations under different conditions. (a) Heat-treatment specimen. (b) 700 °C/650 MPa. (c) 700 °C/760 MPa. (d) 800 °C/450 MPa. (e) 800 °C/650 MPa. (f) 900 °C/200 MPa. (g) 900 °C/450 MPa. (h)

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Table 6

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1000 °C/120 MPa. (i) 1000 °C/200 MPa.

The intensity ratios (I/Imax) of crystallographic planes in M951G alloy in different states. M951G

(111)MC

(200)MC

(111)γ&γ′

(200)γ&γ′

(220)MC

(311)MC

(220)γ&γ′

Heat-treatment

17

17

100

25

10

6

29

700
/650 MPa

23

25

100

31

27

18

25

700
/760 MPa

10

8

100

26

5

5

12

800
/450 MPa

9

9

100

23

5

5

12

800
/650 MPa

9

7

100

6

3

6

15

900
/200 MPa

10

9

100 18

10

4

5

10

ACCEPTED MANUSCRIPT 900
/450 MPa

6

4

100

20

4

2

91

1000
/120 MPa

4

5

100

24

5

3

15

1000
/200 MPa

19

19

100

18

26

19

86

4.2. The effect of solute atoms diffusion on precipitate dimensional changes and

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creep behavior

As is known to all, the diffusion of solute atoms can be strongly accelerated by the

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increase of testing temperature, which has a significant impact on the microstructural evolutions and creep performances of alloys. Fig. 9 shows the typical γ′ morphologies of

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M951G alloy after creep deformations at different conditions. At low and intermediate temperatures (700 °C-800 °C), the γ′ precipitates maintain a well cubic morphology (Fig. 9b-9e). However, when the temperature rises to a relatively higher level (900 °C-1000 °C), varying degrees of N-type rafting [33] are formed as presented in Fig. 9f-9i. It is well known

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that the directional diffusion of solute atoms is responsible for rafting [34, 35]. Fig. 10b shows the schematic illustration of directional coarsening by preferential dissolution of the

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precipitates adjacent to horizontal channels. The precipitate alloying elements diffuse to vertical channels, and simultaneously matrix alloying elements diffuse to horizontal channels

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[34, 35]. For M951G alloy as shown in Fig. 10a, elements such as W, Mo, Co, Cr prefer to segregate in γ matrix, while Hf and Nb elements are partitioned to γ′ phase. Therefore, during directional coarsening, Nb and Hf atoms in horizontal channels tend to diffuse to vertical channels. On the contrary, Co, Cr and Mo atoms in vertical channels tend to diffuse to horizontal channels, as presented in Fig. 10b. The diffusion of alloying elements during creep deformation also has effect on precipitate dimensional changes. Table 7 shows the γ′ phase volume fraction (F(γ′)), 19

ACCEPTED MANUSCRIPT precipitate dimensions and average width of matrix channels after different creep tests. Similar to the report in Ref. [36], it can be clearly seen that at the same applied stress with the increase of testing temperature or at the same testing temperature with the decrease of applied

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stress, the γ′ phase volume fraction is decreased, while the channel width, average γ′ length and aspect ratio of γ′ phase are increased. This trend is even more pronounced, especially the testing temperature above 900 °C. It is also interesting to note that the mean width of γ′ phase

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under all creep tests in present study is thinner than that of the heat treatment sample. The

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coarsening of the γ′ phase perpendicular to the external stress direction and the decrease of the γ′ phase volume fraction during creep tests may be responsible for it. Compared to the initial microstructure (Fig. 9a), the γ′ precipitates of M951G alloy after creep deformations at 700 °C and 800 °C basically retain cubic shape and some corners

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appear to be slightly degenerated (Fig. 9b-9e), thus the effects of precipitate dimensional changes on the creep properties of M951G alloy at low and intermediate temperatures is very limited. However, when creep deformation at high temperature region (900 °C-1000 °C)

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several differences can be found (Fig. 9f-9i) as follows: (1) all or part of the coarse γ′

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precipitates change from cubes to rafts; (2) the finer γ′ precipitates are dissolved, and the width of γ channels is significantly increased; (3) the γ/γ' misfit stresses are relaxed by depositing dislocations along γ/γ' interfaces. When creep deformations at 1000 °C and at 900 °C under relatively lower applied stress (200 MPa) in present study, the dominant deformation mechanism is a combined process of slip and climb of dislocations in matrix channels (Fig. 11f-11h). The wider γ channels and the dissolution of tiny γ′ precipitates allow dislocations easier to slip and cross-slip in matrix channels. However, with the formation of γ′ 20

ACCEPTED MANUSCRIPT rafts the resistance for dislocations to climb over the γ′ precipitates is increased. Therefore, the effects of precipitate dimensional changes on creep properties of M951G alloy at 1000 °C and at 900 °C under relatively lower applied stress are complex and unclear. On the other

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hand, when creep deformation at 900 °C/450 MPa, the dominant deformation mechanism changes to shearing of γ′ precipitates by stacking faults and antiphase boundaries (APBs), as shown in Fig. 11e. The release of the γ/γ' misfit stress and repeatedly shearing of γ′

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[37], which is harmful for alloy creep properties.

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precipitates will result in a decrease of the resistance for dislocations to cut into γ′ precipitates

Fig. 9. Microstructures of M951G alloy after creep tests at different conditions. (a)

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Heat-treatment. (b) 700 °C/760 MPa. (c) 700 °C/650 MPa. (d) 800 °C/650 MPa. (e) 800 °C/450 MPa. (f) 900 °C/450 MPa. (g) 900 °C/200 MPa. (h) 1000 °C/200 MPa. (i)

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1000 °C/120 MPa.

Fig. 10. (a) The element segregation between γ and γ′ phases in M951G alloy determined by

21

ACCEPTED MANUSCRIPT EPMA. The definitions of K=Cγ′/Cγ and K-1=Cγ/Cγ′ are the ratio of alloying elements in the γ′ phase to γ phase and the γ phase to γ′ phase. (b) Schematic illustration of directional coarsening.

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Table 7 The γ′ phase volume fraction (F(γ′)), precipitate dimensions and average width of matrix channels after each creep test. M951G

F(γ′)

Channel width/nm

γ′ length/µm

γ′ width/µm

γ′ form

HT

0.622±0.034

84.2±9.4

0.46±0.10

0.46±0.08

1

Cuboids

700
/ 760 MPa

0.614±0.041

84.9±8.6

0.46±0.10

0.45±0.12

1.02

Cuboids

700
/650 MPa

0.610±0.033

85.2±9.2

800
/ 650 MPa

0.604±0.045

88.3±7.2

800
/ 450 MPa

0.595±0.043

93.0±3.0

900
/ 450 MPa

0.575±0.035

161.8±42.7

900
/ 200 MPa

0.533±0.043

290.8±59.1

1000
/200 MPa

0.529±0.029

1000
/120 MPa

0.550±0.037

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Aspect ratio

0.46±0.10

0.99

Cuboids

0.44±0.09

0.42±0.08

1.05

Cuboids

0.45±0.09

0.43±0.11

1.05

Cuboids

0.72±0.55

0.40±0.10

1.80

Cuboids/plates

4.32±1.74

0.37±0.10

11.68

plates

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0.45±0.08

315.7±72.2

6.30±1.81

0.45±0.09

14.0

plates

312.5±80.1

7.81±1.57

0.44±0.11

17.75

plates

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HT–Heat treatment sample.

Fig. 11. Deformation mechanisms of M951G alloy under different creep conditions. (a) 700 °C/760 MPa. (b) 700 °C/650 MPa. (c) 800 °C/650 MPa. (d) 800 °C/450 MPa. (e) 900 °C/450 MPa. (f) 900 °C/200 MPa. (g) 1000 °C/200 MPa. (h) 1000 °C/120 MPa. 22

ACCEPTED MANUSCRIPT 4.3 The effect of applied stress on the creep life of M951G alloy As shown in Fig. 3, the creep curves of M951G alloy under all testing conditions have a shape with primary, secondary and accelerated creep stages. At the short primary stage, with

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the increase of creep time, the creep strain rate gradually decreases until the minimum value is reached. At the long steady-state stage, the specimen is deformed with a nearly constant creep rate; then creep strain rate significantly increases leading to rupture. The variation of

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creep strain rate with the advance of creep process allows us to estimate the durations of

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primary, steady-state and accelerated stages. Defining steady-state creep as deformation with a creep rate, which deviates from the minimum creep rate not more than 100 %. For example, as displayed in Fig. 13a, when tensile creep deformation at 800
/450 MPa, the minimum strain rate is 2.67*10-5/h, hence the critical creep rate of steady-state stage is 5.34*10-5/h (2

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times of the minimum strain rate). According to the above definition, the primary stage is defined as the time ranging from beginning of the creep test to the deformation with the minimum creep rate (0→A in Fig. 12); the accelerated stage is defined as the time ranging

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from the creep rate above 2 times of minimum creep rate to rupture (B→C in Fig. 12). The

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period between the primary and accelerated stages is the steady-state stage (A→B in Fig. 12). Fig. 13 shows the relationship between creep strain rate and creep time of M951G alloy

under different testing conditions. In addition, the lengths of the primary, steady-state and accelerated stages for M951G alloy under different creep conditions are also measured, which are presented in Fig. 14. At the same testing temperature with the increase of applied stress, the lengths of primary, steady-state and accelerated stages are dramatically decreased. The decrease of creep life can be rationalized as follows. 23

ACCEPTED MANUSCRIPT At 700
and 800
, as shown in Fig. 9, the γ′ precipitates maintain a well cubic shape and the width of γ channels remains relatively narrow and nearly not change. In addition, as displayed in Fig. 11, the dominant deformation mechanism at 700
and 800
is shearing of

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the γ′ precipitates by isolated stacking faults and APBs coupled dislocation pairs. Furthermore, by comparing Fig. 11a and 11b or Fig. 11c and 11d, it can be clearly observed that at the same testing temperature with the increase of applied stress much more tangled dislocations are

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piled up in the γ channels and the density of shearing dislocations (including isolated stacking

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faults and APBs) in γ′ precipitates also significantly increases. The repeated shearing of γ′ precipitates by dislocations through the ordered lattice during creep deformation leads to a mechanical disordering, which weakens the order contribution to strengthening of alloy, sharply accelerates the creep rate and reduces the creep life of specimen [37-39]. On the other

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hand, higher applied stress corresponding to higher strain rate, which will result in inhomogeneous distribution of dislocations, local stress concentration and heterogeneous deformation. Therefore, it is certain that the creep life decreasing is also related to strain rate.

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Apart from the mechanical disordering of γ′ precipitates by dislocation shearing and creep

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strain rate, the stress concentration along the grain boundaries also should be fully considered. As is known to all, the orientations and slip systems of grains on both sides of the grain boundary in polycrystalline alloy are different. During the creep deformation, the slipping of dislocations on {111} planes in different grains will eventually accumulate at the grain boundary. The higher the applied stress, the greater the stress concentration along the grain boundaries, and the easier it is to lead to creep fracture of the alloy. According to the above discussion, it is concluded that the mechanical disordering of γ′ precipitates, creep strain rate 24

ACCEPTED MANUSCRIPT and stress concentration along the grain boundaries are the main reasons for the decrease of creep life with the increase of applied stress (at the same testing temperature) for M951G alloy at 700
and 800
. It also should be noted that, as shown in Fig. 11e, shearing of γ′

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precipitates by dislocations also is the dominant way for M951G alloy plastic deformation at 900
relatively higher applied stress. Therefore, the reasons mentioned above are also responsible for the shorter creep life of M951G alloy at 900
high applied stress.

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At 900
low applied stress and 1000
, as shown in Fig. 11f-11h, a considerable

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number of dense and regular dislocation networks extensively distribute at the γ/γ′ interfaces, while only a few APBs coupled dislocation pairs can be occasionally observed in the γ′ precipitates. Therefore, the primary deformation mechanism changes from shearing of γ′ precipitates by dislocations to a combined process of slipping and climbing of dislocations in

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γ matrix channels with the testing temperature increasing. Similar to the creep deformations at 700
and 800
, the raising of creep strain rate and stress concentration along the grain boundaries with the increase of applied stress are also responsible for the decrease of creep

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life at relatively higher temperatures. Apart from that, other factors such as the grown-up

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micropores and grain boundary sliding also should be taken into consideration. As presented in Fig. 6e, Fig. 7c and 7f, when creep deformations at 1000
and 900
low applied stress, a certain number of grown-up micropores can be observed in the vicinity of the grain boundaries. The formation of grown-up micropores can reduce the effective contact area between the grains and increase the stress concentration along the grain boundaries, which finally accelerates the fracture of specimens. On the other hand, the formation of grown-up micropores may also promote the occurrence of grain boundary sliding when creep 25

ACCEPTED MANUSCRIPT deformation at high temperatures, and this trend will be enhanced as the applied stress increases. It is harmful for alloy creep properties. In summary, for M951G alloy creep deformations at 900
low applied stress and 1000
, the stress concentration along the grain

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sliding are the major reasons for the decrease of creep life.

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boundaries, creep strain rate, the formation of grown-up micropores as well as grain boundary

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Fig. 12. Deviation from the minimum strain rate versus creep time. Defining steady-state creep as deformation with a creep rate, which deviates from the minimum creep rate not more

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than 100 %. The creep curve can be divided into three parts: (i) Primary stage: 0→A. (ii) Steady-state stage: A→B. (iii) Accelerated stage (B→C).

26

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SC

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ACCEPTED MANUSCRIPT

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Fig. 13. The relationship between creep strain rate and creep time of M951G alloy under different testing conditions. (a)-(c) Creep at 700
under 650 MPa, 700 MPa and 760 MPa,

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respectively. (d)-(f) Creep at 800
under 450 MPa, 550 MPa and 650 MPa, respectively. (g)-(i) Creep at 900
under 250 MPa, 360 MPa and 450 MPa, respectively. (j)-(l) Creep at 1000
under 150 MPa, 160 MPa and 200 MPa, respectively.

27

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ACCEPTED MANUSCRIPT

Fig. 14. The lengths of the primary, steady-state and accelerated stages for M951G alloy

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under different creep conditions. (a) 700
. (b) 800
. (c) 900
. (d) 1000
.

4.4. The creep fracture modes

The creep tests of M951G alloy were carried out at different creep conditions, obviously

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different fracture mechanisms were observed, as presented in Fig. 4-Fig. 7. At 700-800 °C

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and 900 °C relatively higher applied stress, a mixed fracture mode of transgranular and intergranular is observed (Fig. 4-Fig. 6). However, the fracture mode changed to pure intergranular cracking at 1000 °C and 900 °C low applied stress (Fig. 6-Fig. 7). 4.4.1. At 700-800 °C and 900 °C higher applied stress Similar fracture characteristics and mechanisms were observed at 700-800 °C and 900 °C relatively higher applied stress (Fig. 4-Fig. 6). It can be clearly observed that although all fracture modes after creep tests at 700
are mixed type of transgranular and intergranular, 28

ACCEPTED MANUSCRIPT the proportion of the transgranular mode increases with the increase of applied stress. It can be rationalized as follows: (i) As shown in Fig. 4a to 4c, the fracture surfaces of M951G alloy at 700
under different applied stresses are cleavage-like on the whole and some slip planes

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also can be observed on the local regions of fracture surfaces (Fig. 4d-4f), which are considered to be formed by the slipping of dislocations on the {111} planes in nickel-base superalloy [32]. It is well known that there is an orientation deviation between different grains

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in the traditional casting equiaxed superalloy. Therefore, various slip systems will be activated

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in different grains during creep deformation. The expansion of dislocations and stacking faults on various {111} planes in different grains will eventually encounter grain boundaries and cause stress concentration. With the increase of creep strain, the local stress concentration along the grain boundaries also gradually increases. Therefore, the stress concentration along

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the grain boundaries resulted from the slipping of dislocations on the various {111} planes in different grains finally leads to the formation of intergranular microcracks. (ii) For traditional casting nickel-base superalloy, the main origins of microcracks are carbides and eutectics [20],

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which are consistent with the observations in present study (Fig. 4g, 4h, 4i). In addition, the

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statistical area fractions and the average sizes of eutectics and carbides of M951G alloy are listed in Table 3. It can be seen that the majority of the carbides and eutectics are distributed in grain interior and the average size of the carbides in the grain interior is much larger than that on the grain boundary. Thus, the microcracks have more probability to initiate in the grain interior, which finally leads to the occurrence of transgranular fracture. The schematic illustration of the mixed fracture mode of transgranular and intergranular is illustrated in Fig. 15a. (iii) Apart from microstructure and temperature, other conditions such as applied stress 29

ACCEPTED MANUSCRIPT and creep strain rate also should not be ignored. High stress is corresponding with high strain rate. As reported in Ref. [14], at low strain rate, grain boundary sliding is the main contributing parameter to creep deformation. However, when the strain rate is high enough

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that the grain boundary adjustment can not catch up with deformation in the grain interior, large deformation will be localized in grain interior and lead to the formation of transgranular cracking. Similar phenomenon also observed in a newly designed Ni-Fe-Cr-based wrought

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superalloy [14], with the increase of creep strain rate the fracture mode varies from

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intergranular to intragranular. Therefore, in present study with the increase of applied stress at 700
, the increase of the creep strain rate may be responsible for the enhancement of the transgranular microcrack proportion. Furthermore, the remarkable stress concentration along the grain boundaries, broken carbides and eutectics and the variation of creep strain rate also

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account for the mixed fracture mode of transgranular and intergranular at 800
and 900
relatively higher applied stress (Fig. 5d, 5e, 5i, 5j and 6j). However, it should be noted that at 800
and 900
relatively higher applied stress a number of small dimples are presented on

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the fracture surfaces, as shown in Fig. 5b, 5g and 6i, illustrating that thermal activation may

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also has certain contribution to the fracture mechanism. 4.4.2. At 1000 °C and 900 °C low applied stress At 1000 °C and 900 °C low applied stress, as shown in Fig. 6e, Fig. 7c and 7f, a number

of coarse micropores can be seen along the grain boundaries, which are the origins of the microcracks. Moreover, with the creep process going on at high temperatures, more and more vacancies can be generated, and the accumulation of the vacancies may exceed the equilibrium value at corresponding temperatures. In order to obtain the minimum value of the 30

ACCEPTED MANUSCRIPT total free energy, the superfluous vacancies will diffuse to and accumulate along the grain boundary [40]. The directional diffusion of vacancies can lead to the coalescence of casting micropores along the grain boundary, which gradually leads to interlinking between the

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grown-up micropores (Fig. 6c, Fig. 7b and 7e) and finally results in the occurrence of intergranular fracture. The mode of the fracture behavior at 1000 °C and 900 °C low applied

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stress is illustrated in Fig. 15b.

High temperatures.

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Fig. 15. The models of the fracture behaviors. (a) Low and intermediate temperatures. (b)

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4.5. The apparent stress exponent

As displayed in Fig. 3e, the values of the apparent stress exponent (n) at different creep

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conditions are significantly different, which indicates that different creep behaviors may take place under different testing conditions. This phenomenon also has been observed in some other nickel-base superalloys during creep deformation at 700 °C-1000 °C, such as GTD-111 [41], M963 [42] and DZ951 [32] and so on. In addition, by comparison of these superalloys, the nominal compositions (as listed in Table 8) and processing routes are significantly different. Therefore, it is reasonable to infer that the variation of n values with the changes of

31

ACCEPTED MANUSCRIPT testing temperatures is a common phenomenon for nickel-base superalloys during creep tests, regardless of nominal compositions and processing routes. As reported in Ref. [32], the transition of dominant deformation mechanism from shearing of γ′ precipitates by

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dislocations to climbing of dislocations in γ matrix channels is supposed to be responsible for the variation of n values from 7.1 to 16.9. Similarly, in the M963 alloy the changes of deformation mechanisms with the increase of testing temperature also have been observed,

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and simultaneously n values vary from 9.5 (800 °C) to 3.9 (975 °C). In present study, when

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creep deformation at 700 °C-800 °C and at 900 °C under relatively higher applied stress, as shown in Fig. 11a-11e, the dominant deformation mechanism is dislocations shearing of the γ′ precipitates via APBs and stacking faults. However, when creep deformation at 1000 °C and at 900 °C under low applied stress, a considerable number of dense and regular dislocation

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networks extensively distribute at the γ/γ′ interfaces, while only a few APBs coupled dislocation pairs can be occasionally observed in γ′ precipitates (Fig. 11f-11h). Furthermore, it is generally believed that the dislocation networks are formed by slipping and climbing of

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dislocations in γ matrix. Thus, the primary deformation mechanism changes from shearing of

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dislocations in γ′ precipitates to a combined process of slipping and climbing of dislocations in matrix channels with the changes of creep conditions. Aside from the deformation mechanisms, other aspects such as fracture modes also should be fully considered to explain the difference of n values. As shown in Fig. 4 to Fig. 7, there is a transition of fracture modes at different temperature regions. Creep deformation at 700-800 °C and at 900 °C under higher applied stress, the fracture mode is a mixed type of transgranular and intergranular (Fig. 4-Fig. 6); while at 1000 °C it turns to pure intergranular cracking (Fig. 7). Compared with the low 32

ACCEPTED MANUSCRIPT and intermediate temperature creep deformation, when creep at high temperature the stress is relatively low while the life is much shorter (as listed in Table 4), which illustrates that the rapid diffusion of atoms and vacancies at high temperature has significant influence on the

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creep properties of M951G alloy. Ultimately, it should be noted that the microstructure changes of M951G alloy during creep deformation, as mentioned in Section 4.2, may also have a partial influence over the variation of n values. In conclusion, the variation of n values

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from 11.594 to 6.381 for M951G alloy during 700-1000 °C creep deformation is due to the

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changes of the deformation mechanisms, fracture modes and the degeneration of microstructures. Ultimately, the schematic map of the microstructural evolutions and fracture behaviors of M951G alloy over a wide range of creep conditions has been summarized and shown in Fig. 16.

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Table 8

Chemical compositions of GTD-111 [41], M963 [42] and DZ951 [32] superalloys. Y

Cr

Co

Ti

W

Al

Ta

Mo

Fe

C

B

Nb

Zr

Ce

Ni

GTD-111

―

13.5

9.5

4.75

3.8

3.3

2.7

1.5

0.23

0.09

0.01

―

―

―

Bal

M963

0.01

8.89

10.0

2.55

10.1

6.0

―

1.6

―

0.15

0.03

1.10

0.031

0.02

Bal

DZ951

―

9.0

5.0

―

3.0

6.0

―

3.0

―

0.05

―

2.2

―

―

Bal

AC C

EP

Alloys

33

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ACCEPTED MANUSCRIPT

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Fig. 16. The summary map of the microstructural evolutions and fracture behaviors of M951G alloy.

5. Conclusions

The creep tests of M951G alloy were carried out over a wide stress range of 120 to 760

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MPa in the temperatures ranging from 700 °C to 1000 °C. Τhe fracture characteristics and microstructural evolutions of M951G alloy under different creep conditions have been

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analyzed in detail. The main conclusions obtained in present study can be listed as follows: (1) The creep curves of M951G alloy obtained at all creep conditions consist of the primary,

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the steady-state and the tertiary stages. At the same testing temperature, with the increase of applied stress the minimum creep rate is increased, while the lengths of primary and the steady-state stages as well as the creep life are significantly decreased. Apart from that, M951G alloy exhibits relatively better ductility at higher temperatures. (2) XRD examinations show that both the heat-treatment and deformed samples consist of γ, γ′ and MC (M=Hf, Nb) carbide. Creep deformations at 700-800 °C and high temperatures

34

ACCEPTED MANUSCRIPT (900-1000 °C) low applied stresses have not any distinctive effect on the acquired patterns. However, deformations under high temperatures relatively higher applied stresses (900 °C/450 MPa and 1000 °C/200 MPa) have significantly highlighted the (220)γ&γ′ planes,

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which may enhance the creep strength and lead to a better plasticity for M951G alloy. (3) At 700-800 °C, the γ′ precipitates basically retain cubic shape and slight degeneration occurs at the corners, which has very limited effect on the creep properties of M951G alloy.

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Significant microstructural evolutions occur after creep deformations at relatively higher

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temperatures (900-1000 °C). At 1000 °C and 900 °C low applied stress region, the wider γ channel and the dissolution of tiny γ′ precipitates allow dislocations easier to slip and cross-slip in matrix channels, while the formation of γ′ rafts increase the resistance for dislocations to climb over the γ′ precipitates. On the other hand, the release of the γ/γ' misfit

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stress and repeatedly shearing of γ′ precipitates by dislocations at 900 °C higher applied stress result in a decrease of the resistance for dislocations to cut into γ′ precipitates, which is harmful for alloy creep properties.

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(4) At 700-800 °C and 900 °C high applied stress region, the mechanical disordering of γ′

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precipitates, creep strain rate and stress concentration along the grain boundaries are the main reasons for the decrease of creep life with the increase of applied stress at the same testing temperature. At 1000 °C and 900 °C low applied stress region, apart from the stress concentration along the grain boundaries and creep strain rate, the formation of grown-up micropores as well as grain boundary sliding are also responsible for the reduction of creep life as the applied stress increases. (5) At 700-800 °C and 900 °C high applied stress region, the fracture mode is a mixed type 35

ACCEPTED MANUSCRIPT of transgranular and intergranular, which is due to the remarkable stress concentration along the grain boundaries, broken carbides and eutectics and the variation of creep strain rate. At 1000 °C and 900 °C low applied stress region, the fracture mode turns to pure intergranular

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cracking. The coarsening and interlinking of the micropores along the grain boundary lead to the occurrence of intergranular fracture.

(6) The values of apparent stress exponents are calculated to be about 11.594 and 6.381 at

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700 °C and 1000 °C, respectively. The changes of the deformation mechanisms, fracture

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modes and the degeneration of microstructures are supposed to be responsible for it.

Acknowledgments

This work was financially supported by the National Natural Science Foundation of China (NSFC) under Nos. 50931004, 51571196 and 51771191. The authors are grateful for

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those supports.

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[2] S. Yamasaki, M. Mitsuhara, H. Nakashima, Deformation Microstructure and Fracture Behavior in Creep-Exposed Alloy 617, Mater. Trans. 58(3) (2017) 442-449. [3] Z.H. Yao, S.C. Wu, J.X. Dong, Q.Y. Yu, M.C. Zhang, G.W. Han, Constitutive behavior and

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ACCEPTED MANUSCRIPT Highlights Relationship among testing conditions-microstructural evolution-creep properties is established.




Crystallographic changes and their influences on creep properties are studied.




Precipitate dimensional changes and their influences on creep properties are investigated.




Deformation mechanisms, fracture modes and γʹ morphologies are important for creep properties.

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