Influences of a Hot-Working Process on the Microstructural Evolution And

metals

Article Inﬂuences of a Hot-Working Process on the Microstructural Evolution and Creep Performance of a Spray-Formed Nickel-Based Superalloy

Tian Tian 1, Changchun Ge 1,*, Xinggang Li 2,* , Zhibo Hao 1, Shiqing Peng 1 and Chonglin Jia 3

1 Institute of Powder Metallurgy and Advanced Ceramics, School of Materials Science and Engineering, University of Science and Technology Beijing (USTB), Beijing 100083, China; [email protected] (T.T.); [email protected] (Z.H.); [email protected] (S.P.) 2 Academy for Advanced Interdisciplinary Studies and Department of Mechanical and Energy Engineering, Southern University of Science and Technology (SUSTech), Shenzhen 518055, China 3 Science and Technology on Advanced High Temperature Structural Materials Laboratory, Aero Engine Corporation of China Beijing Institute of Aeronautical Materials (BIAM), Beijing 100095, China; [email protected] * Correspondence: [email protected] (C.G.); [email protected] (X.L.); Tel.: +86-010-62334951 (C.G.); +86-755-88018144 (X.L.) Received: 8 February 2020; Accepted: 23 March 2020; Published: 31 March 2020

Abstract: A new third generation nickel-based powder metallurgy (PM) superalloy, designated as FGH100L, was prepared by spray forming. The eﬀects of hot isostatic pressing (HIP) and isothermal forging (IF) processes on the creep performance, microstructure, fracture, and creep deformation mechanism of the alloy were studied. The microstructure and fracture were characterized by optical microscopy (OM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The coupled HIP and IF process improved the creep performance of the alloy under the creep condition of 705 ◦C/897 MPa. As for both the HIPed and IFed alloys, the creep process was dominated by the accumulation of dislocations and stacking faults, cutting through γ0 precipitates. The microstructural evolution was the main factor aﬀecting the creep performance, which mainly manifested as coarsening, splitting, and morphology change of γ0 precipitates. Both the creep fractures of the HIPed and IFed alloys indicated intergranular fracture characteristics. In the former, wedge-shaped cracks usually initiated at the trigeminal intersection of the grain boundaries, while in the latter, cavity cracks generate more easily around the serrated curved grain boundary and carbides.

Keywords: third generation nickel-based PM superalloy; creep; microstructure; creep property; creep mechanism; fracture characteristics

1. Introduction The nickel-based powder metallurgy (PM) superalloys are of the most important materials for manufacturing turbine disks of advanced aeroengines. The failure of most load-bearing components under the condition of high temperature and high pressure is caused by high-temperature creep [1], which can lead to excessive plastic deformation or creep fracture of aeroengine parts. There has been a lot of research work on the creep deformation mechanisms of nickel-based superalloys. The turbine disk of aeroengine can work continuously for hundreds to thousands of hours at 650–750 ◦C during its service. It is predicted that the service temperature requirements of the high-pressure compressor and turbine disk of an aeroengine may rise to over 750–800 ◦C in the future. Therefore, it is necessary to develop high temperature materials that can withstand higher temperature and have better creep and oxidation resistance. Accordingly, it is necessary to continuously improve the preparation process of

Metals 2020, 10, 454; doi:10.3390/met10040454 www.mdpi.com/journal/metals Metals 2020, 10, 454 2 of 21 the alloy. At present, there are three kinds of preparation processes of superalloys for turbine disk. The first one is the traditional casting and forging (C & W) process. For example, the superalloys such as Inconel718, U720Li, and AD730TM, which have service temperature of 650–700 ◦C, are produced by the C & W process [2]. The second one is made by the powder metallurgy (PM) process, that is, vacuum induction melting + argon atomization powder making + electrostatic powder screening and then canning + compacting (hot pressing, hot isostatic pressing, extrusion) + the hot working deformation (isothermal forging, die forging) process. For example, the superalloys, such as René88DT, RR1000, René104 (ME3), Alloy10, LSHR, NR3, and NR6 [3–5], are produced by the PM process. Russia has developed the fifth generation of BBΠ system (BB750Π, BB751Π, BB752Π, BB753Π) PM superalloys [6], and the PM superalloy disks are manufactured by the process of plasma rotation electrode process (PREP) for superalloy powders + direct hot isostatic pressing (As-HIP). FGH4095, FGH4096, FGH4097, FGH4098, and other powder superalloys [4,7–10] developed in China adopt the process of powder metallurgy, and the working temperature ranges from 650 to 750 ◦C. The third one is spray forming + compaction (hot isostatic pressing, extrusion) + the isothermal forging process. Compared with C & W or PM technology, the processing steps of spray forming technology [11] from atomization to final pre-forming are significantly reduced, which not only saves the remelting and transformation steps in the process of C & W processing but also avoids the segregation of alloy elements in the solidification process, as well as the screening, electrostatic precipitation, canning, and degassing steps of powder in PM processing. Therefore, spray forming technology has become a cost-effective and efficient preparation method. In this paper, two processes were employed to prepare a new third generation nickel-based PM superalloy FGH100L, i.e., spray forming + HIP and spray forming + HIP + IF. The FGH100L alloy is developed based on the composition of the LSHR alloy [5], and its designed working temperature is above 700 ◦C. The microstructure and creep mechanism of various single crystal and polycrystalline superalloys have been widely studied. The main creep control parameters of single crystal superalloy are orientation, porosity, and the characteristics of strengthening precipitates. The studies on the creep behavior of polycrystalline superalloys mainly focus on the characteristics of grain and γ0 precipitates, the role of trace elements, and dislocation microstructure. Zhu et al. [12] proposed a physical model of the creep deformation of nickel-based single crystal superalloys, which is sensitive to the chemical composition and microstructure. Zhang et al. [13] studied the creep properties of the alloy M4706 at 900 ◦C. When the applied stress was increased from 300 to 375 MPa, the steady-state creep rate gradually increased, while the creep fracture life decreased. Li et al. [14] studied the effect of orientation on the creep behavior of the third-generation nickel-based single crystal superalloy at 850 ◦C. Zhao et al. [15] studied the deformation and damage characteristics of the 4.5%Re/3.0%Ru nickel-based single crystal superalloy during creep. Song et al. [16] studied the influence of Ru on the microstructure and creep properties of the two alloys. With the increasing amount of Ru, the creep fracture life of the alloy increased significantly. Wu et al. [17] studied the microstructural evolution and creep behavior of an Ni3Al-based as-cast superalloy under a high-temperature heat treatment. Yue et al. [18] studied the high-temperature creep behavior of the third-generation nickel-based single crystal superalloy under the condition of 1100 ◦C/120–174 MPa. Huo et al. [19] studied the microstructure of the alloys with different contents of Co (7.0 wt % and 15.0 wt %), Cr (3.5 wt % and 6.0 wt %), Mo (1.0 wt % and 2.5 wt %), and Ru (2.5 wt % and 4.0 wt %) under the creep condition of 950 ◦C/400 MPa. Tian et al. [20] studied the effect of element Re on the internal deformation mechanism of phase in nickel-based single crystal superalloy during high-temperature creep. Lv et al. [21] studied the superdislocation structure of nickel-based single crystal superalloy at the <100> interface after creeping for 85 h at 1100 ◦C/130 MPa. Mohammad et al. [22] studied the high-temperature creep behavior of Inconel-713C, a nickel-based superalloy used for turbine blades. Wollgramm et al. [23] studied the creep behavior of nickel-based single-crystal superalloy, mainly considering the effects of stress and temperature on the minimum creep rate. Viswanathan et al. [24] studied the substructures of the superalloy René 88DT after creep deformation under two stress levels at 650 ◦C with small strain (0.5%). They found that Metals 2020, 10, 454 3 of 21 microtwinning was the main deformation mode under low stress. At higher stress, a dislocation of 1/2<110> passed through the matrix and the tertiary γ0 precipitate, and a dislocation loop surrounded the larger secondary γ0 precipitate. Xie et al. [7] studied the roles of dislocation morphology, creep behavior, and dislocation network during the creeping of the FGH95 superalloy through a creep test, microstructure observation, and comparative analysis. Peng et al. [8] studied the creep properties of the FGH96 superalloys after different aging treatments under the creep condition of 700 ◦C/690 MPa and found that the creep life decreased with the increase in the volume fraction of tertiary γ0 precipitate. Jia et al. [25,26] studied the effects of solid solution treatment and cooling rate on the microstructure and creep properties of the PM superalloy FGH100L at 705 ◦C/897 MPa. Hot isostatic pressing (HIP) is a kind of manufacturing technology that applies high temperature and high pressure to materials at the same time, making materials produce diffusion connection or densification, and it is gradually used for forming complex parts with integral structure. Isothermal forging (IF) is a kind of advanced hot working technology, which heats the mold to the same temperature as the forging billet, and forms the precision forgings with complex shapes with lower deformation rate and fewer fire times under constant temperature. In this paper, the microstructure and creep properties of the alloys prepared by these two different processes were comparatively studied. Based on the analysis of the fracture characteristics, microstructure, and dislocation configuration of the alloy after creep testing, the creep deformation and fracture mechanisms of the alloy were discussed. It is expected to provide a basis for life prediction and the safe use of the new PM superalloy for turbine disks.

2. Materials and Methods The FGH100L master alloy was prepared by a double melting process, i.e., vacuum induction melting plus vacuum consumable remelting (VIM + VAR). The main chemical composition of the FGH100L alloy ingot is (mass fraction, wt%): C 0.04, Cr 12.24, Co 20.90, Mo 2.77, W 4.4, Al 3.48, Ti 3.35, Nb 1.52, Ta 1.47, B 0.023, Zr 0.04, and Ni bal. The FGH100L alloy ingot (diameter: 200 mm, height: 300 mm) was deposited by spray forming. High purity N2 was used as an atomizing gas. The hot isostatic pressing (HIP) was then carried out upon the ingot without canning: the ingot was heated to 1160 ◦C at a rate of 10 ◦C/min, and then held at 150 MPa for 3 h in order to achieve high degree densification. The isothermal forging (IF) experiment was carried out: the forging temperature was 1150 ◦C, the pressing rate was 0.1mm/s, and the engineering deformation was about 30.4%. The samples, taken from the alloy ingots produced by these two different processes, were then subject to solid solution treatment followed by two-stage aging, i.e., 1130 ◦C/1 h/fan quenching + 850 ◦C/4 h/air cooling (AC) + 775 ◦C/8 h/AC. In the rest of the article, the alloy prepared only by the HIP process was designated as the “HIPed alloy” for convenience, while the alloy prepared by the combined HIP and IF process was designated as the “IFed alloy”. The creep tests were carried out on the samples according to China’s national standard GB/T 2039-2012 “Metallic materials—Uniaxial creep testing method in tension” [27]. The test equipment was CSS type creep test machine of Changchun testing machine factory (Changchun, China). The φ5 mm round non-standard creep sample was used, whose specific dimension was shown in Figure1. The creep temperature range of 650–750 ◦C and stress range of 450–897 MPa were selected. Under different temperature and stress conditions, 2–3 samples were taken for creep performance testing. The grain morphology was observed under a metallographic optical microscope XTX-200 (Jiangnan Optical Instrument Factory, Nanjing, China) and the sample etching solution was 10 g CuCl2 + 50 mL HCl + 50 mL H2O. The ImageJ software was used to calculate the size of precipitates (at least 100 data were collected to calculate the average value). The characteristics of γ0 precipitate and grain boundary distribution were observed by JSM-6701F (JEOL Japan Electronics Co., Ltd., Tokyo, Japan) and ZEISS EVO®18 scanning electron microscopy (SEM) (Carl Zeiss AG Co., Ltd., Oberkochen, Germany). The electrolytic polishing solution was 20 vol% H2SO4 + 80 vol% CH3OH, and the electrolytic etching solution was 9 g CrO3 + 90 mL H3PO4 + 30 mL C2H5OH. The dislocations and carbides in the alloy after Metals 2020, 10, x FOR PEER REVIEW 3 of 21 under low stress. At higher stress, a dislocation of 1/2<110> passed through the matrix and the tertiary γ’ precipitate, and a dislocation loop surrounded the larger secondary γ’ precipitate. Xie et al. [7] studied the roles of dislocation morphology, creep behavior, and dislocation network during the creeping of the FGH95 superalloy through a creep test, microstructure observation, and comparative analysis. Peng et al. [8] studied the creep properties of the FGH96 superalloys after different aging treatments under the creep condition of 700 °C/690 MPa and found that the creep life decreased with the increase in the volume fraction of tertiary γ’ precipitate. Jia et al. [25,26] studied the effects of solid solution treatment and cooling rate on the microstructure and creep properties of the PM superalloy FGH100L at 705 °C/897 MPa. Hot isostatic pressing (HIP) is a kind of manufacturing technology that applies high temperature and high pressure to materials at the same time, making materials produce diffusion connection or densification, and it is gradually used for forming complex parts with integral structure. Isothermal forging (IF) is a kind of advanced hot working technology, which heats the mold to the same temperature as the forging billet, and forms the precision forgings with complex shapes with lower deformation rate and fewer fire times under constant temperature. In this paper, the microstructure and creep properties of the alloys prepared by these two different processes were comparatively studied. Based on the analysis of the fracture characteristics, microstructure, and dislocation configuration of the alloy after creep testing, the creep deformation and fracture mechanisms of the alloy were discussed. It is expected to provide a basis for life prediction and the safe use of the new PM superalloy for turbine disks.

Figure 1. Processing diagram of creep specimen.

3. Results

3.1. Creep Properties and Creep Characteristics Generally, the creep curve can be divided into three stages: the primary creep stage, the secondary stage of steady-state creep, and the tertiary stage of accelerated creep. Figure2 shows the creep curves of the HIPed and IFed alloys at 705 ◦C/897 MPa. No distinct primary creep stage can be observed in the creep curves, and the secondary and tertiary stages predominate in these curves. According to Figure2a, under the same creep test conditions, the creep fracture time and strain of the IFed alloy are increased by about 24.6 h and 5.8%, respectively, compared with that of the HIPed alloy. Figure2b shows that the minimum creep speed of the HIPed alloy is 0.1449 10 2 h 1, while the IFed alloy has × − − a smaller steady-state creep rate of 0.1199 10 2 h 1 in Figure2c. Table1 shows the comparison of × − − speciﬁc parameter values. In general, the primary creep behaviors of these two alloys are similar: it is very short and the creep rate decreases with increasing time. During the secondary stage of creep, the creep rate basically remains unchanged with the increase in time, which is the lowest in the whole creep process due to the balance between work hardening and recovery softening. During the steady-state creep stage, the creep rate of the IFed alloy is lower than that of the HIPed alloy, and the creep time of the former is about 20 h longer than that of the latter. In the tertiary creep stage, the creep rate increases with time, and ﬁnally fracture occurs. The increase in creep rate is not only related to the increase in stress, but also to the evolution of the microstructure and the formation and growth of cracks. Metals 2020, 10, 454 5 of 21 Metals 2020, 10, x FOR PEER REVIEW 5 of 21

FigureFigure 2. CreepCreep curves curves of of FGH100L FGH100L alloy alloy at at 705 705 °C◦C/897/897 MPa: ( (aa)) a a comparison comparison of of strainstrain–time–time relationshipsrelationships between between the the two two processes, processes, ( (bb)) the the relationship relationship between between creep creep rate rate and and the the time time of of the the HIPedHIPed alloy, alloy, and ( c) the relationship between creep raterate andand thethe timetime ofof thethe IFedIFed alloy.alloy. Table 1. Comparison of creep properties of FGH100L alloy under diﬀerent processes. Table 1. Comparison of creep properties of FGH100L alloy under different processes. Preparation Process Temperature (◦C) Stress (MPa) Rupture Life (h) Strain (%) PreparationHIP process Temperature 705 (°C) Stress 897 (MPa) Rupture 56.96 life (h) Strain 16.0 (%) HIP 750705 450897 629.1856.96 16.0 37.1 690 1286.27 24.9 HIP + IF HIP + IF 705750 793450 290.46629.18 37.1 7.2 897 81.54 21.9 705 690 1286.27 24.9

The stress-rupture properties of the FGH100L alloys793 under diff290.46erent processes are7.2 compared. Table2 shows that under the condition of 705 ◦C/897 MPa, the average stress-rupture life of the IFed alloy is 1.4 times longer than that of the HIPed alloy, and897 the elongation81.54 after fracture increases21.9 by 2%. The results show that under the same test condition, different processes have obvious effects on the creepThe and stress rupture-rupture properties properties of the of alloys, the FGH100L which may alloys be related under to different the different processes microstructures are compared. of the Tablealloys 2 under shows di thatfferent under processing the condition conditions. of 705 °C/897 MPa, the average stress-rupture life of the IFed alloy is 1.4 times longer than that of the HIPed alloy, and the elongation after fracture increases by Table 2. The stress-rupture properties of FGH100L alloy under different processes. 2%. The results show that under the same test condition, different processes have obvious effects on the creep andPreparation rupture Process properties Test of the Conditions alloys, which Stress-Rupture may be related Life to (h) the different Elongation microstructures (%) of the alloys under different processing conditions. HIP 705 ◦C/897 MPa 47.8 12 HIP + IF 705 ◦C/897 MPa 66.5 14 Table 2. The stress-rupture properties of FGH100L alloy under different processes.

Preparation process Test conditions Stress-rupture life (h) Elongation (%)

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3.2.Metals Creep 2020, Microstructure10, x FOR PEER REVIEW Characteristics 6 of 21 Table3 shows the density and relative density of the FGH100L alloys under these two processes, HIP 705 °C/897 MPa 47.8 12 which are measured by the drainage method [28]. According to Table3, the relative densities of the HIPed and IFed alloys are 98.44% and 99.16%, respectively. Through the HIP treatment, the porosities in the as-sprayedHIP + preforms IF can be705 closed °C/897 and MPa the densities can66.5 be increased. The HIP14 temperature is the key factor affecting the effect of porosities. Too high a temperature can easily cause the complete dissolution of the large γ0 precipitates at the grain boundary, reduce the resistance of grain boundary migration,3.2. Creep M increaseicrostructure the recrystallizationCharacteristics area, and lead to the gradual phagocytosis and growth of the grains.Table 3 Too shows low the a temperature density and relative is not conducive density of to the the FGH100L full diffusion alloys of under solute these atoms two around processes, the porosities,which are measured so that the by porosities the drainage cannot method be completely [28]. According closed and to fullTable densification 3, the relative cannot densities be achieved. of the γ InHIPed this paper,and IFed the alloys HIP temperature are 98.44% and is 1160 99.16%,◦C, whichrespectively. is close Through to 1170 ◦ theC, at HIP which treatment, the 0 precipitates the porosities in thein the FGH100L as-sprayed alloy preforms is completely can be dissolved. closed and At the the densities same time, can be the increased. densification The eHIPffect temperature of the alloy is obviousthe key factor under affecting a pressure the of effect 150 MPa. of porosities. The IF temperature Too high a temperature is 1150 ◦C, and can the easily deformation cause theamount complete is 30.4%.dissolution Relatively of the largelarge plasticγ’ precipitates deformation at the occurs grain inboundary, the alloy, reduce the storage the resistance energy of of alloy grain deformation boundary ismigration, sufficient, increase and the the dynamic recrystallization recrystallization area, and occurs lead toto refinethe gradual the grain. phagocytosis In addition, and the growth IF can of also the breakgrains. the Too prior low particle a temperature boundary is (PPB)not conducive and carbide, to the and full close diffusion the internal of solute microscopic atoms around porosities the ofporosities, the alloy. so Therefore, that the the porosities number cannot of microscopic be comp porositiesletely closed in the and IFed full alloy densification was further cannot reduced, be andachieved. the densification In this paper, effect the was HIP better temperature than that is in 1160 the HIPed°C, which alloy. is close to 1170 °C, at which the γ’ precipitates in the FGH100L alloy is completely dissolved. At the same time, the densification effect Table 3. Density and relative density of FGH100L alloys under different processes. of the alloy is obvious under a pressure of 150 MPa. The IF temperature is 1150 °C, and the deformation amountPreparation is 30.4%. Process RelativelyAverage large Density plastic/(g deformationcm 3) Relative occurs Density in the/ %alloy, the storage · − energy of alloy deformationSF is sufficient, and the dynamic – recrystallization 97.33 occurs to refine the grain. In addition, the IF can alsoHIP break the prior particle 8.23 boundary (PPB) and carbide, 98.44 and close the internal microscopic porositiesHIP of +theIF alloy. Therefore, the 8.29 number of microscopic 99.16 porosities in the IFed alloy was further reduced, and the densification effect was better than that in the HIPed alloy.

Figure3 shows the initial morphology of grains and γ0 precipitates in both of HIPed and IFed Table 3. Density and relative density of FGH100L alloys under different processes. alloys before creep. It can be seen from Figure3a that there are a few recrystallized grains in the HIPed alloy. The grainsPreparation are generally process polygonal Average and the density/(g·cm grain boundaries–3) tendRelative to be density/% flat. The grain size in the HIPed alloy ranges from 12 to 120.49 µm, and the average grain size is 40.73 µm. Figure3b shows that a large number of fineSF recrystallized grains can– be found in the IFed alloy,97.33 and the average size is 4.33 µm. The larger grain size ranges from 4.96 to 53.99 µm, with an average grain size of 24.25 µm. HIP 8.23 98.44 Dynamic recrystallization occurs during the IF process, which can refine grains and form curved grain boundaries, and the morphologyHIP+IF of grains changes8.29 from polygon to nearly99.16 spherical.

Figure 3. Cont.

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Figure 3. Initial microstructures of FGH100L alloy before creep: (a) optical microscopy (OM) images Figure 3. Initial microstructures of FGH100L alloy before creep: (a) optical microscopy (OM) images of HIPedof HIPed alloy, alloy, (b) ( OMb) OM images images of IFedof IFed alloy, alloy, (c,e ()c SEM) and images (e) SEM of images HIPed of alloy, HIPed and alloy, (d,f) SEMand ( imagesd) and ( off) IFedSEM alloy.images of IFed alloy.

Figure 33c,e shows show the the initialγ0 morphology morphology in of the grains HIPed and alloy. γ’ precipitates The average in both size of of HIPed the primary and IFedγ0 precipitatesalloys before is creep. 1.98 µ mIt incan Figure be seen3c. Thesefrom primaryFigure 3aγ 0thatprecipitates there are are a few in irregular recrystallized strip or grains sheet shapesin the andHIPed are alloy. mainly The distributed grains are on generally the grain polygonal boundaries. and Figure the grain3e shows boundaries that a large tend number to be flat. of secondaryThe grain γsize0 precipitates in the HIPed are alloy dispersed ranges in from the 12 grains, to 120.49 with μm, an averageand the average size of 0.84 grainµm. size During is 40.73 the μm. splitting Figure process,3b shows the that secondary a large γnumber0 precipitates of fine mainly recrystallized split into grains triangular can orbe cubic found shapes, in the due IFed to thealloy, interaction and the betweenaverage size the adjacentis 4.33 μm.γ0 Theprecipitates. larger grain The size tertiary rangesγ0 fromprecipitates 4.96 to are53.99 formed μm, with during an average the aging grain process size andof 24.25 act μm. as supplementary Dynamic recrystallization strengthening. occurs The during big tertiary the IF process,γ0 precipitates which can (B-tertiary refine grainsγ0) are and mainly form distributedcurved grain in boundaries, the grains andand the appear morphology in spherical of grains or cubic changes shapes from with polygon a size to of nearly less than spherical. 0.14 µm. The smallFigure tertiary 3c,e showγ0 precipitates the γ’ morphology (S-tertiary inγ 0 the) are HIPed distributed alloy. inThe the average phase boundary size of the between primary the γ’ primaryprecipitates and is secondary 1.98 μm inγ 0Figureprecipitates 3 c. These and primary the matrix, γ’ precipitates with a size ofare about in irregular 0.06 µm strip and or a finesheet spherical shapes shapeand are (see mainly in the distributed amplified on area the in grain Figure boundaries.3e). The morphology Figure 3e shows of γ0 thatprecipitate a large number is determined of secondary by the interfaceγ’ precipitates and elastic are dispersed strain energy. in the The grains, change with in the an supersaturation average size of of 0.84alloying μm. elements During inthe the splitting matrix causedprocess, by the the secondary decrease in γ’ temperature precipitates is main the mainly split factor into controlling triangular the or nucleation cubic shapes, of γ0 dueprecipitate, to the whileinteraction the di betweenffusion of the alloying adjacent elements γ’ precipitates. in the matrix The tertiary is the main γ’ precipitates factor controlling are formed the growth during of theγ0 precipitateaging process [10 ].and act as supplementary strengthening. The big tertiary γ’ precipitates (B-tertiary γ’) are mainlyFigure 3distributedd,f show the in theγ0 morphology grains and appear in the IFedin spherical alloy. In or Figure cubic3 d,shapes there with are a a small size of number less than of incomplete0.14 μm. The re-dissolved small tertiary primary γ’ precipitatesγ0 precipitates (S-tertiary distributed γ’) are in the distributed grain boundary, in the phase with an boundary average sizebetween of 1.75 theµ primarym in irregular and secondary strip or sheet γ’ precipitates shapes. In and Figure the3 matrix,f, a large with number a size of of cubic about secondary 0.06 μm andγ0 precipitatesa fine spherical are distributedshape (see in the grainsamplified with area an averagein Figure size 3e). of The 0.52 morphologyµm. There are of alsoγ’ precipitate two sizes ofis determined by the interface and elastic strain energy. The change in the supersaturation of alloying tertiary γ0 precipitates in the IFed alloy. The big tertiary γ0 precipitates are mainly distributed in the grainselements and in appear the matrix in spherical caused and by cubic the decrease shapes with in temperature an average size is theof 0.08 mainµm. factor The small controlling tertiary theγ0 precipitatesnucleation of are γ’ ofprecipitate, fine spherical while shape the diffusion with a size of ofalloying about elements 0.01 µm and in the are matrix distributed is the inmain the phasefactor boundarycontrolling between the growth the primaryof γ’ precipitate and secondary [10]. γ0 precipitates and the matrix. Figure 3d,f show the γ’ morphology in the IFed alloy. In Figure 3d, there are a small number of incomplete re-dissolved primary γ’ precipitates distributed in the grain boundary, with an average size of 1.75 μm in irregular strip or sheet shapes. In Figure 3f, a large number of cubic secondary γ’

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Figure4 shows the morphology of grains and γ0 precipitates in both of HIPed and IFed alloys after creep at 705 ◦C/897 MPa. The samples were taken around the fracture surface along the longitudinal direction (parallel to the stress direction). The ImageJ software was used to calculate the size of the precipitates in two different process state alloys. Figure4a shows that the grain size ranges from 15.22 to 133.13 µm in the HIPed alloy after creep, and the average grain size is 54.44 µm. The recrystallized grains in the HIPed alloy after creep are larger and more evenly distributed than those before creep. Figure4b shows that a large number of fine recrystallized grains can still be seen in the IFed alloy after creep. The smaller recrystallized grains range from 2.64 to 15.78 µm after creep, and the average grain size is 8.31 µm, which is slightly larger than that of recrystallized grains before creep. The larger grains range from 7.24 to 107.63 µm, and the average grain size is 37.24 µm. Table4 shows the statistics of average grain size and γ0 precipitate size before and after creep of FGH100L alloy in different processes. According to Table4, the grain size and γ0 precipitate size of both the HIPed and IFed alloys tend to increase after creep. It can be seen from Figure4c that after creep, part of primary γ0 precipitates coarsened in the HIPed alloy, with an average size of 2.32 µm, distributing on the grain boundaries. In Figure4d,e, some secondary γ0 precipitates are coarsened to form a raft structure. These secondary γ0 precipitates, with an average size of 1.18 µm, are mainly distributed in the crystal with cubic shapes. Compared with that before creep, there is only one size of cubic tertiary γ0 precipitates in the alloy after creep. These tertiary γ0 precipitates, with an average size of 0.23 µm, are mainly distributed in the phase boundary between the primary and secondary γ0 precipitates and the matrix. The phenomenon of the coarsening and growing of tertiary γ0 precipitates is obvious. A large number of tertiary γ0 precipitates tend to gather in a separate wide region, as indicated by the arrows in Figure4c,d. Compared with the microstructure in Figure3c,e, it is found that, in the microstructure after creep, the γ0 precipitates grow along a preferred direction under the action of stress and temperature. This coarsening behavior directly affects the creep fatigue life of the alloy, and the creep fracture usually occurs along the coarsening direction. Figure4f–h shows the morphology of γ0 precipitates in the IFed alloy after creep. Compared with the HIPed alloy, the IFed alloy has smaller sized γ0 precipitates after creep, as shown in Figure4f and Table4. Compared with Figure3d,f, the γ0 precipitate morphology in the IFed alloy changed after creep. The size of the primary γ0 precipitate does not change significantly, and the average size is 1.72 µm. These primary γ0 precipitates are distributed on the grain boundaries with an irregular long strip shape. The average size of the secondary γ0 precipitate is 0.62 µm. As shown in Figure4g, these secondary γ0 precipitates are mainly in a cubic shape and distributed in the crystal, but no raft structure is observed. The size of the tertiary γ0 precipitate is significantly larger than that before creep, and there are two different sizes of tertiary γ0 precipitates, as shown in Figure4h. The division and coarsening of the big tertiary γ0 precipitates, with an average size of 0.24 µm, occur simultaneously. Part of these big tertiary γ0 precipitates coarsen to form raft structures in a cubic shape. These big tertiary γ0 precipitates are mainly distributed within the grain. The small tertiary γ0 precipitates, in a spherical shape and with an average size of 0.06 µm, tend to coarsen. These small tertiary γ0 precipitates are distributed in the phase boundary between the primary and secondary γ0 precipitates and the matrix. Metals 2020, 10, x FOR PEER REVIEW 8 of 21 precipitates are distributed in the grains with an average size of 0.52 μm. There are also two sizes of tertiary γ’ precipitates in the IFed alloy. The big tertiary γ’ precipitates are mainly distributed in the grains and appear in spherical and cubic shapes with an average size of 0.08 μm. The small tertiary γ’Metals precipitates2020, 10, 454 are of fine spherical shape with a size of about 0.01 μm and are distributed in the phase9 of 21 boundary between the primary and secondary γ’ precipitates and the matrix.

Figure 4. Microstructures of FGH100L alloy after creep at 705 ◦C/897 MPa: (a) OM images of HIPed alloy, (b) OM images of IFed alloy, (c–e) SEM images of HIPed alloy, and (f–h) SEM images of IFed alloy.

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Table 4. Statistics of average grain size and γ0 precipitate size before and after the creep of FGH100L alloy in diﬀerent processes.

Before Creep After Creep Process Grain Primary Secondary Grain Primary Secondary Tertiary γ0 Tertiary γ0 Size γ0 γ0 Size γ0 γ0 B-tertiary γ 0.14 µm; HIP 40.73 µm 1.98 µm 0.84 µm 0 54.44 µm 2.32 µm 1.18 µm 0.23 µm S-tertiary γ0 0.06 µm B-tertiary B-tertiary γ 0.08 µm; γ 0.24 µm; IF 24.25 µm 1.75 µm 0.52 µm 0 37.24 µm 1.72 µm 0.62 µm 0 S-tertiary S-tertiary γ0 0.01 µm γ0 0.06 µm

According to the literature [29–31], after the creep of the general nickel-based superalloy, the γ0 precipitates in the alloy will directionally coarsen to form raft structures, which corresponds to the ﬁrst stage of macroscopic alloy creep. The directional coarsening of γ0 precipitate is a spontaneous process of energy reduction, so it is beneﬁcial to the increase of creep strength. During the second stage of steady-state creep, the raft structure is basically unchanged, but the dislocation network at the interface will undergo a process from formation to gradual dissipation. In the third stage of creep, the raft structure will coarsen violently. At this time, some holes will form and gradually expand at pores and phase interfaces, which will eventually lead to material fracture.

3.3. Creep Fracture Characteristics Figures5 and6 show the creep fracture morphology of the FGH100L alloy. The creep fracture life of the HIPed and IFed alloys is 56.96 h and 81.54 h, respectively. It can be seen clearly from Figures5a and6a that there are three characteristic regions of creep fracture in both of the alloys, namely, the fracture source region, the propagation region, and the ﬁnal fracture region (i.e., the shear lip region). Overall, the fracture surfaces are rough and uneven. The shear lip region of the creep fracture in the IFed alloy (Figure6a) is larger than that in the HIPed alloy (Figure5a), indicating that the IFed alloy has better plasticity. Figure5b demonstrates intergranular fracture characteristics in the fracture source region of the HIPed alloy. Since the grain boundary is the origin of the crack, most of the creep cavities are generated in the grain boundary perpendicular to the tensile stress. Moreover, under the higher stress load, the cracks are initiated at the trigeminal intersection of the grain boundaries, which are shown as wedge-shaped cracks (see the ampliﬁed area in Figure5b). With the increase in stress, the wedge-shaped cracks widen. Figure5c shows small platforms, deep dimples, and secondary cracks in the propagation region. Figure5d shows shallow dimples in the shear lip region. Metals 2020, 10, 454 11 of 21

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Figure 5. Fracture morphology of HIPed alloy after creep at 705 °C/897 MPa: (a) low magnification Figure 5. Fracture morphology of HIPed alloy after creep at 705 ◦C/897 MPa: (a) low magnification Figure 5. Fracture morphology of HIPed alloy after creep at 705 °C/897 MPa: (a) low magnification morphology, ((bb)) fracturefracture sourcesource region,region, ( (cc)) propagation propagation region, region, ( d(d)) shear shear lip lip region. region. morphology, (b) fracture source region, (c) propagation region, (d) shear lip region. Figure6 b 6b shows shows the thefracture fracture source source region region of the IFed of the alloy, IFed indicating alloy, the indicating feature of the intergranular feature of fractureintergranularFigure and 6b the fracture shows hole-type and the cracks. the fracture hole In - Figuretype source cracks.6 c, region there In are ofFigure deepthe 6c,IFed dimples, there alloy, are secondary indicating deep dimples, cracks, the featureand secondary small of intergranularstepscracks, with and deep small fracture interlaced steps and with slip the traces de holeep - ininterlacedtype the cracks. propagation slip In traces Figure region in 6c, the (see there propagation the are area deep indicated region dimples, by (see the secondary arrowthe area in cracks,theindicated figure). and by In smallthe Figure arrow steps6d, in the with the deep figure). de dimplesep interlaced In Figure (as shown 6d, slip the in traces thedeep magnified in dimples the propagation area(as shown in the figure)in region the magnified can (see be the seen areaarea in indicatedthein the shear figure) lipby the region. can arrow be seen in thein the figure). shear In lip Figure region. 6d, the deep dimples (as shown in the magnified area in the figure) can be seen in the shear lip region.

Figure 6. Cont.

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Figure 6. Fracture morphology of IFed alloy after creep at 705 °C/897 MPa: (a) low magnification Figure 6. FractureFracture morphology morphology of of IFed IFed alloy alloy after creep at 705 ◦°CC/897/897 MPa: ( a)) low magnificationmagnification morphology, (b) the fracture source region, (c) the propagation region, the arrow shows small steps morphology, ( b) the fracture source region, ( c) the propagation region region,, the arrow shows small steps with deep interlaced slip traces, and (d) the shear lip region. with deep interlaced slip traces, and ( d) the shear lip region. FiguresFigure 77 andand Figure8 show 8 the sho fracturew the fracture morphology morphology of the of stress-rupture the stress-rupture samples samples of the of HIPed the HIPed and Figure 7 and Figure 8 show the fracture morphology of the stress-rupture samples of the HIPed IFedand IFed alloys, alloys, and and the the endurance endurance life life of theseof these two two alloys alloys is 47.8is 47.8h h and and 66.5 66.5h, h, respectively.respectively. According and IFed alloys, and the endurance life of these two alloys is 47.8h and 66.5h, respectively. According to FiguresFigure 77aa andand8 Figurea, the fracture 8a, the surfacesfracture ofsurfaces the two of samples the two are samples rough are and rough uneven and like uneven a sponge like as a to Figure 7a and Figure 8a, the fracture surfaces of the two samples are rough and uneven like a asponge whole, as and a whole, the fracture and the includes fracture theincludes fracture the source fracture region, source propagation region, propagation region and region final and fracture final sponge as a whole, and the fracture includes the fracture source region, propagation region and final regionfracture (i.e., region shear (i.e., lip shear region). lip region). fracture region (i.e., shear lip region). In Figure 77b,b, intergranularintergranular fracturefracture morphologymorphology andand longlong wedge-shapedwedge-shaped crackscracks cancan bebe seenseen inin In Figure 7b, intergranular fracture morphology and long wedge-shaped cracks can be seen in the f fractureracture source source region region of of the the HIPed HIPed alloy alloy (in (in the the amplified amplified region region at the at the upper upper right right corner corner in the in the fracture source region of the HIPed alloy (in the amplified region at the upper right corner in the thefigure). figure). In Figure In Figure 7c,d,7c,d, small small planes, planes, tearing tearing edges, edges, steps, steps, and and deep deep dimples dimples can can be be seen in in the the figure). In Figure 7c,d, small planes, tearing edges, steps, and deep dimples can be seen in the propagation region. Figure 7 7ee showsshows thatthat therethere areare shallowshallow dimplesdimples in in the the shear shear lip lip region. region. propagation region. Figure 7e shows that there are shallow dimples in the shear lip region.

Figure 7. Cont.

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Figure 7. Fracture morphology of HIPed alloy after the stress-rupture at 705 °C/897 MPa: (a) low Figure 7. FractureFracture morphology of of HIPed alloy after the stress stress-rupture-rupture at 705 °CC//897897 MPa: ( a) low magnification morphology, (b) the fracture source region, (c) steps and ◦ small planes in the magnificationmagniﬁcation morphology, morphology, (b ) (b the) the fracture fracture source source region, region, (c) steps (c and) steps small and planes small in the planes propagation in the propagation region, (d) tearing edges and steps in the propagation region, and (e) the shear lip region. propagationregion, (d) tearing region, edges (d) tearing and steps edges in and the propagationsteps in the propagation region, and (region,e) the shear and (e lip) the region. shear lip region.

InIn FigureFigure8 8b,b,8b, intergranular intergranularintergranular fracture fracturefracture features featuresfeatures are areare demonstrated demonstrateddemonstrated in theinin thethe fracture fracturefracture source sourcesource region regionregion of the ofof theIFedthe IFedIFed alloy alloyalloy and andand cavity cavitycavity cracks crackscracks can becancan found bebe foundfound (as shown (as(as shownshown in the inin ampliﬁed thethe amplifiedamplified area at areaarea the aa lowertt thethe lowerlower left corner leftleft cornercorner in the inﬁgure).in thethe figure).figure). This is ThisThis because isis becausebecause the IFed thethe IFed alloyIFed alloyalloy has serrated, hashas serrated,serrated, curved curvedcurved grain graingrain boundaries, boundaries,boundaries, so the soso cavitythethe cavitycavity cracks crackscracks are areeasier easier to germinate to germinate and and expand expand at at the the grain grain boundary boundary and and carbide carbide interface. interface. Figure Figure8 8c,dc,d show show the the featuresfeatures of of small small planes planes,planes,, tearing tearing edges, edges, and and deep deep dimples dimples in in the the propagation propagation region. region. Figure Figure 8 8e8ee shows showsshows deep dimplesdimples inin the the shear shear lip lip region. region. The The above above analyses analyses show show that that the rupture the rupture fractures fractures of both of HIPed both HIPed and IFed alloys are intergranular fractures under the stress-rupture condition of 705 °C/897 HIPedand IFed and alloys IFed are alloys intergranular are intergranular fractures fractures under the under stress-rupture the stress- conditionrupture conditi of 705on◦C /of897 705 MPa. °C/897 MPa.

Figure 8. Cont.

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Figure 8. Fracture morphology of IFed alloy after the stress-rupture condition at 705 °C/897 MPa: (a) Figure 8. Fracture morphology of IFed alloy after the stress-rupture condition at 705 ◦C/897 MPa: (lowa)Figure low magnification magnification 8. Fracture morphology, morphology morphology, of(b IFed) ( bthe) thealloy fracture fracture after sourcethe source stress region, region,-rupture (c (c) condition)small small planes planes at 705 in in °C the /897 propagation MPa: (a) region,low magnification (d)) tearingtearing edges morphology, andand smallsmall (b planesplanes) the fracture inin thethe propagationpropagationsource region, region,region, (c) small ((ee)) thetheplanes shearshear in liplipthe region.region. propagation region, (d) tearing edges and small planes in the propagation region, (e) the shear lip region. 4. Discussion 4. Discussion 4.1. Dislocation C Configurationonfiguration and CreepCreep DeformationDeformation MechanismMechanism 4.1. Dislocation Configuration and Creep Deformation Mechanism The TEM analysesanalyses ofof thethe specimensspecimens of of the the HIPed HIPed alloy alloy after after creep creep at at 705 705◦C °C/897/897 MPa MPa are are shown shown in Figurein Figure9The. Figure9. TEM Figure9 analysesa shows9a shows of that the that thespecimens the creep creep deformation of deformation the HIPed mechanism alloy mechanism after creep of of the theat HIPed705 HIPed °C/897 alloy alloy MPa are are are dislocation dislocation shown andin stackingFigure 9. faultFigure cutting 9a shows the thatγγ’0 precipitateprecipitate. the creep deformation. Figure 99bb showsshows mechanism thethe continuouscontinuous of the HIPed andand alloy widewide are stackingstacking dislocation faultsfaults cuttingand stacking the γγ matrixmatrix fault cutting and and γ’γ the0precipitate.precipitate. γ’ precipitate The The. γ’Figure γprecipitates0 precipitates 9b shows are the are obstacles continuous obstacles to dislocationand to dislocation wide stacking movement. movement. faults As Ascreepcutting creep proceeds, proceeds,the γ matrix the the number and number γ’ precipitate. of ofdislocations dislocations The γ’ increases. precipitates increases. Dislocations Dislocations are obstacles can can to bedislocation be seen seen in in themovement. the local local reg regions Asions cuttingcreep γγ’proceeds,0 precipitates the number andand formingforming of dislocations stackingstacking faults increases.faults extending extending Dislocations in in di differentff erentcan be directions seendirections in the (Figure (Figurelocal 9regc) 9c) andions and at theat cuttingthe intersection intersection γ’ precipitates of of partial partial and stacking stackingforming fault faultstacking and and anti-phasefaults anti-phase extending boundary boundary in different (APB), (APB), directions asas indicatedindicated (Figure byby the 9c) arrowsand at the intersection of partial stacking fault and anti-phase boundary (APB), as indicated by the arrows in Figure 99d.d. In In the the later later stage stage of of creep, creep, th thee accumulation accumulation of ofa large a large number number of ofdislocations dislocations near near the in Figure 9d. In the later stage of creep, the accumulation of a large number of dislocations near the thegrain grain boundary boundary creates creates stress stress concentration concentration so so that that the the grain grain boundary boundary can can slide slide more more easily. easily. grain boundary creates stress concentration so that the grain boundary can slide more easily. Meanwhile, the atoms on the grain boundary didiffuseffuse rapidly at a highhigh temperature.temperature. In addition, the Meanwhile, the atoms on the grain boundary diffuse rapidly at a high temperature. In addition, the HIPed alloy has flatflat and smooth grain boundaries, which are easier to slide under stress. When the HIPed alloy has flat and smooth grain boundaries, which are easier to slide under stress. When the sliding grain boundaries are blocked at the trigeminaltrigeminal intersection of the graingrain boundaries,boundaries, cavities sliding grain boundaries are blocked at the trigeminal intersection of the grain boundaries, cavities occur due to stress concentration, and the cavities combine with each other to form wedge-shaped occuroccur due due to to stress stress concentrat concentration,ion, andand thethe cavitiescavities combinecombine with with each each other other to to form form wedge wedge-shaped-shaped cracks. The increase in the applied stress or test temperature will accelerate the dislocation movement cracks.cracks. The The increase increase in in the the applied applied stress stress oror test temperature will will accelerate accelerate the the dislocation dislocation movement movement and grain boundary sliding, which will accelerate the initiation and propagation of grain boundary andand grain grain boundary boundary sliding, sliding, which which willwill accelerateaccelerate the initiationinitiation and and propagation propagation of of g raingrain boundary boundary cracks in the later stage of creep, and finally accelerate the occurrence of creep fracture. crackscracks in inthe the later later stage stage of of creep, creep, and and finallyfinally accelerate thethe occurrence occurrence of of creep creep fracture. fracture.

Figure 9. Cont.

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Figure 9. Transmission electron microscopy (TEM) images of dislocation morphology in HIPed alloy FigureFigure 9. Transmission9. Transmission electron electron microscopy microscopy (TEM)(TEM) imagesimages of dislocation morphology morphology in in HIPed HIPed alloy alloy afterafter creep creep under under the the condition condition of of 705 705 °C °C/897/897 MPa: MPa: ((aa)) dislocationsdislocations andand stackingstacking faults, faults, ( b()b )and and (c )( c) after creep under the condition of 705 ◦C/897 MPa: (a) dislocations and stacking faults, (b,c) stacking stackingstacking faults, faults, and and (d )( dthe) the anti anti-phase-phase boundary boundary (APB). (APB). faults, and (d) the anti-phase boundary (APB).

TheTheThe creep creep deformation deformation ofof of thethe the IFed IFed IFed alloy alloy alloy is is dominatedis dominated dominated by by dislocation dislocation and andand stacking stacking stacking fault fault fault cutting cutting cutting the the theγ’ precipitate,γ’ precipitate, which which hinders hinders the the c reepcreep deformation. deformation. FigureFigure 10a showsshows that that stacking stacking faults faults cut cut the the γ0 precipitate, which hinders the creep deformation. Figure 10a shows that stacking faults cut the γ0 γ’ precipitateγ’ precipitate and and a largea large number number of of dislocations dislocations accumulate accumulate around the the γ’ γ’ precipitate precipitate to to form form precipitate and a large number of dislocations accumulate around the γ0 precipitate to form stacking stacking faults. Figure 10b shows the continuous wide stacking faults cutting the γ matrix and γ’ stackingfaults. Figure faults. 10 Figureb shows 10b the shows continuous the continuous wide stacking wide faults stacking cutting faults the cuttingγ matrix the and γ matrixγ precipitate. and γ’ precipitate. Figure 10c shows a large number of dislocations packed around the γ’ precipitate.0 At the precipitate.Figure 10c shows Figure a 10c large shows number a large of dislocations number of dislocations packed around packed the γ aroundprecipitate. the γ’ Atprecipitate. the late stage At the of late stage of creep, the γ’ precipitate raft demonstrably thickens,0 and the creep deformation latecreep, stage the γ ofprecipitate creep, the raft γ’ demonstrably precipitate raft thickens, demonstrably and the creep thickens, deformation and the mechanism creep deformation is mainly mechanism0 is mainly the dislocation cutting mechanism. In addition, with the further increase in the mechanism is mainly the dislocation cutting mechanism. In addition, with the further increase in the thenumber dislocation of dislocations cutting mechanism. at the end In of addition, creep, dislocations with the further pile up increase near the in serrated the number grain of boundaries dislocations number of dislocations at the end of creep, dislocations pile up near the serrated grain boundaries at theand end carbides of creep, (Figure dislocations 10d), which pile can up block near dislocation the serrated movement. grain boundaries Moreover, and at the carbides same time, (Figure stress 10d), and carbides (Figure 10d), which can block dislocation movement. Moreover, at the same time, stress whichconcentration can block dislocationis generated. movement. When the Moreover, yield strength at the sameis reached, time, stress the cavity concentration cracks initiate is generated. and concentration is generated. When the yield strength is reached, the cavity cracks initiate and Whenpropagate the yield on strengththe grain isboundary reached, and the at cavity the interface cracks initiate between and carbides propagate and matrix. on the grainIt should boundary be noted and propagateat thethat interface grain on the boundaries between grain boundary carbides and carbides and and at matrix. blockthe interface the It should propagation between be noted carbides of that cracks. grainand Figure matrix. boundaries 10e It should is andthe be energycarbides noted thatblockdispersive grain the propagation boundaries spectrometer of and cracks. (EDS) carbides Figureof the blockcarbide 10e is the thein Figureenergy propagation 10d. dispersive According of cracks. spectrometer to the Figure analysis (EDS)10e results is of the the of carbideenergy EDS dispersivein Figureand select 10 spectrometerd.edAccording area electron (EDS) to diffraction the of analysis the carbide (SAED) results in, theFigure of EDScarbide 10d. and mainly selectedAccording contains area to electronthe C, Cr,analysis Co, diﬀ Ni,raction results Ta, W, (SAED),of and EDS and selected area electron diffraction (SAED), the carbide mainly contains C, Cr, Co, Ni, Ta, W, and theMo carbide elements, mainly and contains is an M23 C,C6- Cr,type Co, carbide. Ni, Ta, W, and Mo elements, and is an M23C6-type carbide. Mo elements, and is an M23C6-type carbide.

Figure 10. Cont.

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Figure 10. TEMTEM images images of ofdislocation dislocation morphology morphology in IFed in IFedalloy alloyafter creep after creepat 705 °C at/897 705 MPa:◦C/897 (a), MPa: (b), (anda–c) (c dislocations) dislocations and and stacking stacking faults, faults, ( d(d)) carbides carbides and and selectedselected areaarea electronelectron didiffractionﬀraction (SAED),(SAED), and (e)) thethe energyenergy dispersivedispersive spectrometerspectrometer (EDS)(EDS) ofof carbidecarbide inin (Figured). 10 (d).

Under the creep condition of 705 ◦°CC//897897 MPa, the creep deformation mechanism of both HIPed γ and IFed IFed alloys alloys is dislocation is dislocation and theand stacking the stacking fault cutting fault the cutting0 precipitate, the γ’ precipitate, which makes which dislocations makes γ γ easierdislocations to decompose. easier to decompose. The continuous The continuous wide stacking wide faultsstacking running faults throughrunning through the matrix the γ andmatrix0 precipitateand γ’ precipitate are present are present in the alloy.in the Underalloy. Under the action the action of temperature of temperature and stress, and stress, the state the of state the of grain the boundariesgrain boundaries and carbides and carbides in the secondaryin the secondary and tertiary and tertia stagesry of stages creep of have creep a great have influence a great influence on the final on fracturethe final offracture the alloy. of the This alloy. is because This is because the micro-cracks, the micro- whichcracks, cause which fatal cause damage fatal damage to the alloy, to the mostly alloy, germinatemostly germinate at the trigeminal at the trigeminal intersection intersection of the grain of the boundaries grain boundaries and at the and interface at the interface where carbides where combinecarbides withcombine the matrix, with the and matrix, expand and rapidly expand in rapidly the tertiary in the stage tertiary of creep. stage In of addition, creep. In corrosion, addition, oxidation,corrosion, oxidation, and other actionsand other in theactions process in the of process high-temperature of high-temperature creep are alsocreep important are also important factors to enhancefactors to the enhance alloydamage the alloy to d theamage final to fracture. the final fracture. 4.2. Creep Damage and Endurance Life Prediction 4.2. Creep Damage and Endurance Life Prediction Figure 11 shows the creep curves of the IFed alloy under different stress and temperature Figure 11 shows the creep curves of the IFed alloy under different stress and temperature conditions. Figure 11a shows that under the same stress load of 897 MPa, the endurance life of the conditions. Figure 11a shows that under the same stress load of 897 MPa, the endurance life of the alloy decreases significantly with the temperature increasing from 650 to 750 C. It can be seen from the alloy decreases significantly with the temperature increasing from 650 to 750◦ °C. It can be seen from creep curve that the primary creep stage cannot be observed obviously, and the secondary creep stage the creep curve that the primary creep stage cannot be observed obviously, and the secondary creep lasts for a short time and then quickly enters the tertiary creep stage. Under the creep condition of stage lasts for a short time and then quickly enters the tertiary creep stage. Under the creep condition 650 C/897 MPa, the endurance life of FGH100L alloy is 1477.42 h, and the total strain is 5.22%. As the of 650◦ °C/897 MPa, the endurance life of FGH100L alloy is 1477.42 h, and the total strain is 5.22%. As test temperature increases, the duration of the secondary stage of creep gradually decreases, which is the test temperature increases, the duration of the secondary stage of creep gradually decreases, which is about 50, 12, and 2h at 650, 705, and 750 °C, respectively. At the tertiary stage of creep, the higher the testing temperature, the greater the creep speed, and the shorter the endurance life. The

Metals 2020, 10, 454 17 of 21 about 50, 12, and 2 h at 650, 705, and 750 C, respectively. At the tertiary stage of creep, the higher Metals 2020, 10, x FOR PEER REVIEW ◦ 17 of 21 the testing temperature, the greater the creep speed, and the shorter the endurance life. The creep straincreep of strain the tertiary of the stage tertiary varies stage significantly, varies sig whichnificantly, reflects which that temperature reflects that plays temperature an important plays role an onimportant the structure role on evolution, the structure crack evolution, initiation crack and initiation propagation, and propagation, and fracture and mode. fracture Figure mode. 11b showsFigure the11b creep shows curves the creep of the curves alloy underof the dialloyfferent under temperature different temperature and stress conditions. and stress At conditions. 750 ◦C/450 At MPa, 750 the°C/450 total MPa, creep the strain total of creep the alloy strain is 37.07%,of the alloy and is the 37.07%, alloy has and a lowthe alloy steady-state has a low creep steady rate- andstate a creep long endurancerate and a lifelong of endurance about 629.18 life h. of At about 705 ◦ C629.18/690 MPa, h. At the 705 alloy’s °C/690 steady-state MPa, the creepalloy’s lasts steady for- thestate longest creep timelasts offor about the longest 300h with time a lowof about creep 300h rate. with The maximuma low cree creepp rate. strain The maximum is 24.99%, andcreep the strain endurance is 24.99%, life isand 1286.27 the endurance h. life is 1286.27 h.

Figure 11. Figure 11. CreepCreep curves curves of IFed of IFed alloy alloy under under different diff erent stress stress and temperature and temperature conditions: conditions: (a) 897 (a) 897 MPa/650–750 C and (b) 450–897 MPa/705–750 C. MPa/650–750 °C and◦ (b) 450–897 MPa/705–750 °C. ◦ The above analysis shows that the creep rate of the FGH100L alloy is significantly affected by The above analysis shows that the creep rate of the FGH100L alloy is significantly affected by load load stress and temperature. Generally, the relationship among the minimum creep rate, creep stress, stress and temperature. Generally, the relationship among the minimum creep rate, creep stress, and and temperature complies with Norton’s formula [32–34]. temperature complies with Norton’s formula [32–34]. n ε· s = Aσa exp( Qc/RT) (1)  − n  s = σ exp(−Q / RT) (1) In Equation (1), ε· s is the minimum creep rate,a A is the constantc related to the material microstructure, T is the absolute temperature, R is the gas constant (8.314 J/mol K), Qc is the creep activation energy,  s · n is theIn stress Equation exponent, (1), andis σ thea is the minimum applied creep stress. rate, A is the constant related to the material microstructure,At a given T temperature, is the absolute take temperature, the logarithm R is of the both gas sidesconstant of Equation(8.314 J/mol·K), (1) and Q convertc is the creep it to Equationactivation (2): energy, n is the stress exponent, and σa is the applied stress. At a given temperature, take the logarithmlgε· s = nlg ofσa both+ c sides of Equation (1) and convert it (2) to Equation (2): In Equation (2), c is a constant. The relationship between lgε· s and lgσa is linear in the double logarithmic coordinate, and the slope of the straight line is the stress exponent n. (2) When the applied stress is fixed, Equationlg ε s = (1)n canlg σ bea + convertedc to Equation (3):

In Equation (2), c is a constant. Thelg εrelationship·s = Q /RT between+ c0 lg s and lgσa is linear in the double(3) − c logarithmic coordinate, and the slope of the straight line is the stress exponent n. 1 In Equation (3), c0 is a constant. So, lgε· s has a linear relationship with T− . The slope of the linear When the applied stress is fixed,1 Equation (1) can be converted to Equation (3): relationship between lgε· s and T− represents the creep activation energy. Figure 12 shows the linear relationship between lgε· s and lgσa of the IFed alloy under the creep  condition of 705 ◦C/690–897 MPa. The minimum creep rate of the alloy increases with the increase(3 in) lg  s = −Qc / RT + c' the applied stress. By calculating the slope of the straight line in the ﬁgure, it can be concluded that the γ n valueIn Equation is 16.33, which (3), c’ isis mainlya constant. related So, tolg thes has0 precipitationa linear relationship strengthening with T−1 because. The slope a large of the number linear of dispersed γ0 precipitates hinder the movement of dislocations. The properties of grain boundaries relationship between lg s and T−1 represents the creep activation energy. and other factors have an important inﬂuence on the creep deformation of the alloy as well. Figure 12 shows the linear relationship between lg s and lgσa of the IFed alloy under the creep condition of 705 °C/690–897 MPa. The minimum creep rate of the alloy increases with the increase in the applied stress. By calculating the slope of the straight line in the figure, it can be concluded that the n value is 16.33, which is mainly related to the γ’ precipitation strengthening because a large

Metals 2020, 10, x FOR PEER REVIEW 18 of 21 Metals 2020, 10, x FOR PEER REVIEW 18 of 21 number of dispersed γ’ precipitates hinder the movement of dislocations. The properties of grain numberboundaries of dispersed and other γ’ factors precipitates have an hinder important the movement influence onof dislocations.the creep deformation The properties of the of alloy grain as boundariesMetalswell. 2020 , 10 and, 454 other factors have an important influence on the creep deformation of the alloy18 of as 21 well.

Figure 12. Relationship of the lg s -lgσa of IFed alloy at 705 °C/690–897 MPa. Figure 12. Relationship of the lgε· s-lgσa of IFed alloy at 705 ◦C/690–897 MPa. Figure 12. Relationship of the lg s -lgσa of IFed alloy at 705 °C/690–897 MPa. 1 TheThe linearlinear relationshiprelationship betweenbetween lglgε· s sand and T T−−1 is obtained, asas shownshown inin FigureFigure 1313.. The steady-statesteady-state creepcreepThe activation activation linear relationship energy energy Q c Qofc between theof theIFed IFed alloylg salloy underand T under−1 the is creepobtained, the condition creep as shown condition of 897 in MPa Figure of / 897650–750 13. MPa/650 The◦C steady is– calculated750- state°C is creeptocalculated be 339.35 activation to kJ be/mol. 339.35 energy The kJ/mol. self-di Qc of ff theTheusion IFedself activation- diffusion alloy under energy activation the of creep pure energy conditionNi in of austenite pure of Ni 897 in is 265–285austenite MPa/650 kJ– is750/mol, 265 °C– and285 is calculatedthatkJ/mol, of polycrystalline and to thatbe 339.35 of polycrystalline nickel-basedkJ/mol. The self superalloynickel-diffusion-based is about activationsuperalloy 300 kJ energy /ismol about [35 of]. pure300 The kJ/mol Ni FGH100L in austenite [35]. alloy The is hasFGH100L 265 a– high285 kJQalloy/mol,c value has and asa therehighthat of Q is cpolycrystalline avalue high-volume as there nickel fractionis a high-based of-volumeγ 0superalloyprecipitates fraction is in aboutof the γ’ alloy,precipitates300 kJ/mol which [35]. stronglyin the The alloy, aFGH100Lffects which the alloyprocessstrongly has of affectsa atom high the di Qffc usionprocessvalue by as of hinderingthere atom isdiffusion a dislocationhigh- volumeby hindering climbing, fraction dislocation vacancy of γ’ precipitates directional climbing, diinvacancy fftheusion, alloy, directional and which grain stronglyboundarydiffusion, affects and sliding. grain the Theprocess boundary higher of atom thesliding. Q diffusionc value, The higher the by greaterhindering the Qc the value, dislocation creep the resistance greater climbing, the of creep the vacancy alloy resistance and directional the of less the diffusion,pronealloy and it is and tothe creep grainless deformation.prone boundary it is sliding.to Therefore, creep The deformation. higher the alloy the will QTherefore,c possessvalue, the higherthe greater alloy creep thewill strength creep possess resistance and higher longer ofcreep the life alloyendurancestrength and a thend [32 longer].less prone life endurance it is to creep [32]. deformation. Therefore, the alloy will possess higher creep strength and longer life endurance [32].

1 Figure 13. Relationship of the lgε· s-T− of IFed alloy at 897 MPa/650–750 ◦C. Figure 13. Relationship of the lg s-T−1 of IFed alloy at 897 MPa/650–750 °C.

5. Conclusions Figure 13. Relationship of the lg s-T−1 of IFed alloy at 897 MPa/650–750 °C. 5. Conclusions 1. Under the creep condition of 705 C/897 MPa, the creep rupture time and strain of the IFed alloy 5. Conclusions ◦ 1. Underare about the 24.6creep h condition and 5.8% of higher, 705 °C respectively,/897 MPa, the compared creep rupture with thattimeof and the strain HIPed of alloy,the IFed and alloy the 1. Underareendurance about the creep24.6h life of conditionand the 5.8% former ofhigher, 705 is 1.4 °C respectively, times/897 MPa, longer the compared than creep that rupture of with the latter.thattime ofand the strain HIPed of thealloy, IFed and alloy the

2. areenduranceThere about is only 24.6h life one of and the size 5.8% former of thehigher, is cubic 1.4 respectively, times shaped longer tertiary comparedthanγ 0thatprecipitates ofwith the thatlatter. in of the the HIPed HIPed alloy alloy, after and creep. the endurance life of the former is 1.4 times longer than that of the latter. The phenomenon of coarsening and growing of tertiary γ0 precipitates is obvious. Some secondary γ0 precipitates coarsen to form raft structures. Metals 2020, 10, 454 19 of 21

3. There are two sizes of tertiary γ0 precipitates in the IFed alloy after creep. The division and coarsening of the big tertiary γ0 precipitates occurred simultaneously. Part of the big tertiary γ0 precipitates coarsened to form raft structures, with a cubic shape, that are mainly distributed in the grains. The small tertiary γ0 precipitates coarsened and were distributed in the phase boundary between the primary and secondary γ0 precipitates and the matrix. After creep, the secondary γ0 precipitates are distributed in the grains, but no raft structure of secondary γ0 precipitates is observed. 4. The fracture source areas of both HIP and IFed alloys demonstrate intergranular fracture characteristics. The diﬀerence is that the cracks in the HIPed alloy usually initiate at the trigeminal intersection of the grain boundaries and produce wedge-shaped cracks. However, the IFed alloy has serrated, curved grain boundaries, which allows cavity cracks to more easily germinate at the grain boundary and carbide interfaces.

5. Under the condition of 705 ◦C and 897 MPa, the creep deformation mechanism of both HIPed and IFed alloys is dislocation and the stacking fault cutting the γ0 precipitate. Continuous wide stacking faults running through the γ matrix and γ0 precipitate are present in both alloys. 6. The creep properties of the FGH100L alloy is sensitive to stress and temperature. Under the condition of 705 ◦C/690–897 MPa, the stress exponent n of the IFed alloy is 16.33. Under the condition of 897 MPa/650–750 ◦C, the creep activation energy Qc of the IFed alloy is 339.35 kJ/mol.

Author Contributions: Conceptualization, C.G. and X.L.; Data curation, T.T., X.L., Z.H., S.P. and C.J.; Investigation, T.T., Z.H. and S.P.; Project administration, C.G.; Supervision, C.G. and X.L.; Writing—original draft, T.T. and Z.H.; Writing—review & editing, C.G., X.L., S.P. and C.J. All authors have read and agreed to the published version of the manuscript. Funding: This research was funded by the National Natural Science Foundation of China, grant number 51171016; Shenzhen Science and Technology Innovation Commission, grant number JSGG20180508152608855. The APC was funded by the National Natural Science Foundation of China. Acknowledgments: The authors thanks the group of Udo Fritsching and Volker Uhlenwinkel of Bremen University in Germany, Zhang Yu-Chun of Fushun Special Steel Shares Co., Ltd., Ye Jun-Qing of Guizhou Anda Aviation Forging Co., Ltd., Zheng Wei-Wei of State Key Laboratory for Advanced Metals and Materials, and Chen Jin of School of Metallurgical and Ecological Engineering in University of Science and Technology Beijing, for their valuable contributions. Conﬂicts of Interest: All authors declare no conﬂict of interest. All funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

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