Crack formation and microstructure-sensitive propagation in low cycle fatigue of a polycrystalline nickel-based superalloy with different heat treatments

Accepted Manuscript Crack formation and microstructure-sensitive propagation in low cycle fatigue of a polycrystalline nickel-based superalloy with different heat treatments Yahui Liu, Maodong Kang, Yun Wu, Mengmeng Wang, Min Li, Junwei Yu, Haiyan Gao, Jun Wang PII: DOI: Reference:

S0142-1123(17)30404-8 https://doi.org/10.1016/j.ijfatigue.2017.10.012 JIJF 4486

To appear in:

International Journal of Fatigue

Received Date: Revised Date: Accepted Date:

20 July 2017 6 October 2017 20 October 2017

Please cite this article as: Liu, Y., Kang, M., Wu, Y., Wang, M., Li, M., Yu, J., Gao, H., Wang, J., Crack formation and microstructure-sensitive propagation in low cycle fatigue of a polycrystalline nickel-based superalloy with different heat treatments, International Journal of Fatigue (2017), doi: https://doi.org/10.1016/j.ijfatigue. 2017.10.012

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Crack formation and microstructure-sensitive propagation in low cycle fatigue of a polycrystalline nickel-based superalloy with different heat treatments Yahui Liu a, Maodong Kang a, Yun Wu a, Mengmeng Wang a, Min Li a, Junwei Yu a, Haiyan Gao a, Jun Wang a, b, * a. School of Materials Science and Engineering, Shanghai Jiao Tong University, No. 800, Dongchuan Road, Shanghai 200240, China b. Shanghai Key Laboratory of Advanced High-temperature Materials and Precision Forming, Shanghai Jiao Tong University, No. 800, Dongchuan Road, Shanghai 200240, China * Corresponding author. Address: No. 800, Dongchuan Road, Shanghai 200240, China. Tel.: +86 21 54745487; fax: +86 21 54742683; e-mail: [email protected].

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Crack formation and microstructure-sensitive propagation in low cycle fatigue of a polycrystalline nickel-based superalloy with different heat treatments

Abstract Gas turbine engine materials demand high performance of fatigue life under alternative loading. In this paper, crack initiation and propagation in the low cycle fatigue (LCF) of a polycrystalline nickel based superalloys was studied via experimental and numerical methods. LCF experiments were conducted with the specimens subjected to different heat treatment regimes, including of standard heat treatment (SHT), heat isostatic pressing (HIP), and combined heat treatment (HIP+SHT). Simulations based on finite element method (FEM) were implemented with the 3D digital model obtained from synchrotron radiation X-ray computerized tomography (CT) to investigate the effect of microporosity. Results indicated that microporosity was a dominant factor of fatigue crack and the parameters of microporosity including of porosity, effective diameter, and distribution co-contributed to affect fatigue life. The δ stack and carbide were observed to be fatigue crack sites. The HIP specimen had long fatigue cycle life since casting pore and acicular δ were eliminated, while further SHT can reintroduce δ into matrix resulting in high strength and low fatigue life. LCF crack behaviors in HIP+SHT superalloy depended on surface quality, stress concentration, and the quantity of carbide. Keywords Nickel-based superalloy; low cycle fatigue; heat treatment; synchrotron radiation X-ray CT; FEM 2 / 28

1.

Introduction Using in the aggressive alternative loading environment requires excellent fatigue strength

of nickel based superalloy employed as an elevated temperature material in turbine engine. Under the operation conditions of aircraft propulsion, low cycle fatigue (LCF) behaviors caused by high stress are more important than the high cycle fatigue (HCF) for the structure design [1, 2]. The applications of nickel based superalloy in turbine engine consist of two groups: low temperature components (cold section components) and high temperature components (hot section components) [3]. The former case involves support casing, low pressure (LP) compressor blade and so on. The latter case includes combustion chamber components, turbine blade, turbine disk, and so on. Most mass of the aircraft engine is contributed by the low temperature components, including of forgings and castings (monocrystalline, directional crystal and polycrystalline) [4]. It’s necessary to understand the fatigue failure behaviors of the nickel based superalloy at low temperature in the views of material designer and structure designer. Factors influencing on the fatigue crack behavior of nickel based superalloy can be classified in terms of alloy elements [5], microstructure characteristics ( such as micropore and precipitates) [6-11], heat treatment [12], surface finish quality, thermal-mechanical conditions (such as temperature, loading, and residual stress distributions) [13-16], and so on. The strategic elements is an important factor in the fatigue strength of the material, referring the homogeneous distribution of elements. For example, the effect of boron (B) on LCF behavior of Inconel 718 at room temperature was reported by Xiao et al [5]. Born addition into Inconel 718 was observed to improve the deformation and to delay the crack initiation.

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However, segregation of refractory elements such as rubidium and molybdenum in the superalloy results in the development of interdendritic brittle phase although it can be eliminated by homogenization process. Microstructure is of paramount importance in the fatigue behavior of the materials, owing to the direct relationship between microstructural characteristics and fatigue crack. Micropores are believed commonly as the micro crack source in the fatigue failure of superalloy casting, because of the severe stress concentration at the internal surface of micropore [11, 17]. The strengtheners in the nickel based superalloy include γ′ and γ′′, both of which precipitates in the γ base. The deformation mechanism of the strengtheners at room temperature depends on planar slip and cross slip, and stack faults can form in the strengtheners and matrix because of the cycle hardening in LCF [18]. Inhomogeneous dislocation deformation between precipitates and matrix causes the micro crack formation inside the brittle precipitates or at the particle interface. Additionally, low interface bonding strength promotes particle interface crack or accelerates the fatigue crack propagation. An et al. [19] reported that the stress concentration around the acicular δ caused by dislocations made the surrounding of δ phase to be crack propagation channels in forged superalloy. This phenomenon may also occur in casting superalloy due to the common acicular δ. Nonmetallic inclusions as crack nucleation sites can also reduce the fatigue life of the superalloy [20-22]. Grain boundaries (GBs) and twin boundaries (TBs) with larger disorientation angle can also increase the local stress concentration and promote fatigue crack nucleation [23]. Configuration was also found by Fu and the fellow that the stress field around micropore can induce twins and determine the fatigue life of TMS-82 superalloy in common [24].

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Heat treatment and HIP can affect the mechanical performance owing the significant microstructure evolution in the processing [1, 25-29]. The effect of SHT on the LCF of Inconel 718 superalloy was studied by Xu et al [1], and the cycle softening of SHT superalloy was reported in the reference. What’s more, mechanical properties including of creep strength, fatigue strength, and toughness can improved significantly by using HIP processing [25]. Moreover, parameters of fatigue experiments including of testing temperature, total strain amplitude, mean strain, have also significant influences on the fatigue life of superalloy [15, 30]. The LCF life decreased with the increasing of testing temperature or total strain range. The current study is aim to investigate the crack initiation and propagation in the LCF of a nickel based superalloy considering the influences of heat treatment processes on the microstructural changings. The crack initiation mechanisms in fatigue experiments were discussed in terms of microporosity, carbides, GBs, and acicular δ. The formation mechanism of micro crack at casting micropores in the material was investigated by using the combination of synchrotronic radiation X-ray CT and FEM. Care was taken in the interact of micropore and interdendritic precipitates in order to understand the their effects on the fatigue fracture behavior of the alloy under consideration at room temperature. 2.

Experimental

2.1. Material The chemical compositions of the polycrystalline nickel based superalloy in this study is shown in Table 1. The superalloy was prepared by investment casting processing where the superalloy ingot re-melted in vacuum induction furnace and formed by ceramic shell. In the experimental procedure (see Fig. 1), the as-cast superalloy was chipped to prepare specimens

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for LCF and synchrotron radiation X-ray computerized tomography (X-ray CT). The specimens were given three after-treatments: SHT, HIP, and HIP+SHT, respectively. SHT with the regime of 1095°C×2h/air cooling+950°C×1h/air cooling+720°C×8h/air cooling at 56°C/h to 650°C× 8h/air cooling was performed to homogenize the original morphology of microstructure and to increase the number of precipitates. Most of the Laves (a topologically close packed phases, TCP phase) can be transformed to precipitates during the SHT processing. HIP with the regime of 1165°C×4h×140MPa was conducted in order to eliminate the porosity in casting. HIP + SHT was implemented by applying SHT to the HIP superalloy. Finally, pin specimens were prepared for the synchrotron radiation X-ray CT and the 3D reconstruction of the material which can be useful in the subsequent finite element modeling and simulation. 2.2. LCF experiment LCF tests were conducted with the heat treated specimens, and the design of the specimen is illustrated in Fig. 2. LCF specimens were obtained by wire electrical discharge machining (WEDM) and machining. The specimen surface was carefully polished to Ra0.4. The LCF testing was conducted at room temperature on a hydraulic experimental platform MTS810 equipped with computer controlled program, the strain control condition was used. All the LCF experiments were performed in the way of mechanical strain control in axial direction and the strain amplitude ε was set as constant value of 0.4%, where the strain ratio was set as -1. LCF experiments were implemented by triangular wave with loading frequency of f = 0.2Hz. In each cycle, peak and valley values of displacement were recorded, and the stress

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amplitude can be calculated. Additionally, the number of cycles was recorded until the LCF specimen failure occurs. Post LCF tests, morphological observations were performed including of optical microscope (OM) and scanning electron microscope (SEM). Fracture surfaces were sectioned from the fracture LCF specimens for fractography observation, and the fractography was directly observed by using a Raman imaging scanning electron microscopy (RISE Microscopy, TESCAN-MIRA/WITEC) and a field emission scanning electron microscope JSM-7600 (JEOL Ltd., Tokyo, Japan). Thin cylindrical specimens were sectioned perpendicular to the LCF specimen axial direction for metallurgic observation. Metallurgic specimens were mechanically polished and chemically etched by acid solution (100ml alcohol-100ml HCl-5g CuCl2) for metallurgic observation with an Olympus BX51M optical microscope and the SEM JSM-7600. The SHT specimens were twin-jet electropolished by the 10% chloric + 90% ethanol solution in conditions of 30V and -20°C for the observation with a field emission transmission electron microscopy (TEM) JEM2100F (JEOL Ltd., Tokyo, Japan). 2.3. Synchrotron radiation X-ray scanning experiment and FEM Synchrotron radiation X-ray CT experiments were performed with the as-cast pin specimens with diameter of Ø0.5mm. The surfaces of acicular specimens were mechanically polished. The X-ray scanning experiment (Fig. 3) was implemented on the synchrotron radiation beamline station BL13W1 at Shanghai Synchrotron Radiation Facility (SSRF). The electron energy of synchrotron radiation source is 3.5GeV. The superior penetration ability of the X-ray promises the accurate observation of the dense nickel based superalloy. The pixel

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size of CCD camera is 0.65μm×0.65μm×0.65μm, which meets the imaging requirements of microporosity in the superalloy. In the X-ray CT experiment, the pin specimens were located vertically at the rotation center of experimental platform which can rotate horizontally in a range of -180° ~ 180° and translate vertically in a range of -24mm ~ 24mm. A desired scanning position was scanned to build a series of slices with the exposure time 7s per slice. Post processing of the raw data were performed with the programs P3 (SSRF) to obtain X-ray CT slices which were used for 3D reconstruction of material. 3D reconstruction were processed by using ImageJ and Amira software. The final 3D models of alloy were used for micropore statistics and finite element modeling. The 3D model of microporosity was imported into the software ABAQUS to establish the FE model, as shown in Fig. 3. The tensile tests were simulated with the FE model in ABAQUS and fatigue life calculation was conducted in the software fe-safe based on the simulation results of tensile tests. 3.

Results

3.1. Microstructure in superalloy Nickel based superalloy is a multicomponent alloy that contains so many alloy elements and a variety of phases. The nickel γ base is strengthened by γ′ and disk-like γ′′. Some brittle phase can precipitate in the interdendritic zone of casting superalloy, including of Laves, δ, and carbide. What’s more, impurity elements can also be imported into the superalloy to form oxide and nitride even though the vacuum induction heating was employed in practice. Microstructure evolution in the nickel-based superalloy with different heat treatment regime was shown in Fig.4. The metallurgic picture with magnitude of 1000 times were obtained by

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using OM. The precipitates at interdendritic and grain boundaries (GBs) can be clearly observed. In the as cast case (Fig. 4a), interdendritic precipitates are irregular shape Laves and acicular δ. The Laves phases form through eutectic reaction and consume a large number of the reinforce elements, such as Mo, Nb, and Ta in the nickel based superalloy. While acicular δ (orthorhombic, NI3Nb) has the same chemical compositions with the strengthener γ′′ (bct, Ni3Nb). In the case of SHT, Laves phase decomposition caused the formation of acicular δ, coarse γ′′ (black spots) in the interdendritic region, as shown in Fig. 4(b). A large number of acicular δ precipitated at interdendritic and intergranular zone. Under the elevated temperature and extreme high pressure in HIP, the acicular δ reduced significantly while the carbides were relatively stable (Fig. 4c). Therefore, the carbide has better thermodynamic stability than other precipitates. The further SHT after HIP process can form acicular δ again at intergranular zone. However, the precipitated δ in Fig. 4(d) was much smaller than the primary δ in the Fig. 4(a) and the secondary δ in the Fig. 4(b). Carbide was always observed in all cases. The SEM pictures (Fig. 5a) of SHT superalloy show the morphologies of the interdendritic precipitates. It can be seen that the acicular δ has several sub layer. Since the formation of acicular δ needs a large number of niobium, the depleted area of niobium forms around the acicular δ. Coarse γ′ and γ′′ precipitate in the vicinity of the acicular δ during the SHT processing, as shown in Fig. 5(b). The interdendritic disk-like γ′′ precipitates with diameter of ~0.6μm and thickness of ~0.1μm are particles which form in the certain crystal planes, as

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shown in Fig. 5(c). Smaller precipitates γ′′ were also observed in the matrix, as shown in Fig. 5(d). TEM observation reveals that the morphologies of precipitates in the superalloy after SHT processing, as shown in Fig. 6. Fig. 6(a) shows a large number of the strengthener γ′′ in matrix and the diffraction patterns reveals the coherent relationship between the nickel base and strengtheners. Fig. 6(a) gives the diffraction pattern of blocky carbide in the vicinity of δ particles, and the large spots was the diffraction pattern of surrounding mass, the γ-nickel. Micrograph of the δ particle provides the confirmation to the fact that the δ particle contains multilayer substructures, as shown in Fig. 6(c). HR-TEM of the SHT specimens in Fig. 6(d) reveals the depletion of the strengthener γ′′ in the interface zone between δ and γ-base, and the interface zone has a width of 5nm.

3.2. LCF life and fractography The effect of heat treatment on the number of cycles to fracture in the LCF test and the stress amplitude are illuminated in Fig. 7. In the case of SHT specimens, the LCF test performs as high strength and low fatigue cycle life. This phenomenon is caused by the precipitation of strengthening particles such as the γ′ and disk-like γ′′ in the interdendritic regions, as shown in Fig. 5. Simultaneously, the decreasing of Laves phases can also contribute to improve the fatigue strength. In the case of HIP specimens, the cycles to fracture increases significantly while the stress amplitude reduces to the lower level than the case of SHT. Long time exposition in elevated temperature and extreme high pressure results in the microstructural homogenization of the HIP specimens. Coarse interdendritic precipitates and strengtheners decompose under the temperature close to melting point. Only the steady carbide can remain in the matrix of HIP case. Nevertheless elimination of microporosity in

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the HIP case increase the densification of the superalloy. Due to the regeneration of acicular δ in the HIP+SHT case, the strength of the HIP+SHT superalloy increases while the fatigue cycle life of the HIP+SHT superalloy decreases compared with the HIP case. The continuous decreasing of the stress amplitude in the case of SHT specimen indicates that the high strength SHT exhibited cycle softening after the first 8 cycles, while the HIP with lower strength exhibited cycle hardening. Under the cycle strain loading, the HIP+SHT specimen exhibited both cycle softening and cycle hardening, resulting in the stress amplitude fluctuating up and down. To understand the crack mechanism in the LCF experiment of the superalloy, the fractography were observed by using SEM. Fig. 8 gives the LCF fractography of the SHT specimen. Microporosity was observed in the crack initiation zone and crack propagation zone on the fracture surface, as shown in Fig. 8 (a). Secondary cracks initiating from microporosity surface was observed as shown in Fig. 8(b) and (e). That provides an evidence to the fact that micropores in the subsurface layer of test bar are the crack source. In detail, Fig. 8 (e) shows the distribution of acicular δ particles on the surface of microporosity relates to micro crack initiation. The characteristics of the crack propagation zone include the fracture morphology of dendrite (Fig. 8c) and the plastic fatigue striations on the fracture surface of secondary dendritic arms (Fig. 8f). This kind of fatigue striations recorded crack movement under the alternative loading in the LCF experiments. When the axial strain reaches to the tensile peak, the fatigue crack opens and moves a step toward. When the axial strain shaft to the compression peak, the fatigue crack close to form a striation on the two fracture surface. The cyclic loading remains cyclic striations on the failure surface. In addition,

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the grain boundaries can significantly affect the width of plastic fatigue striations, as shown in Fig. 8(f). Ultimately, intergranular micropore was also observed in the propagation zone. Another kind of fatigue striations, cleavage striations, were observed in the final rupture zone on the fracture surface, see Fig. 8 (d). The cleavage striations in different directions are conjoined through fatigue steps. In the case of HIP specimens, see Fig. 9(a), the LCF fractography includes the crack initiation zone, crack propagation zone and final rupture zone. Fig. 9 (b) shows intergranular fracture morphology in the crack initiation zone. Since the microporosity and brittle precipitates were eliminated in HIP process, the GBs were apt to be crack site during fatigue experiment. Crack propagating along the low strength GBs leads to the superalloy failure. Fig. 9 (c) gives the plastic fatigue striations in crack propagation zone in different direction. Microstructural features in alloy, such as GBs, twin boundaries (TBs), carbide, and inclusion, can significantly change the direction and width of fatigue striations. Fig. 9 (d) shows the fatigue striations in the final rupture zone. The edges of fatigue striations in the final rupture zone deepened alternatively, which may by caused by the instability of the material. Tearing ridges are formed in the direction of crack propagation. No microporosity was observed on the fatigue surface of the HIP specimens. A fact can be obtained from the fractography of this case is that plastic fracture dominates the LCF fracture of the HIP superalloy. Fig.10 (a) gives the macro fractography of the HIP+SHT specimen. The characteristics of the three fatigue zones can be descripted as: (1) the radial river pattern in crack initiation zone centered on surface defects (Fig. 10b) while fine fatigue striations formed in this zone; (2) numerous fatigue striations formed in the crack initiation zone, and a lots of TBs were also

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observed to cause the symmetrical fatigue striations with a uniform width on the fracture surface; (3) large scale tearing burrs and secondary cracks were observed in the smooth final rapture zone. Obviously, the carbide and carbonitride caused the formation of smooth fracture surface on the fractography and broke the inhomogeneity of fatigue striations. According to the microstructure observation in Fig. 4 (d), carbide and small scale δ particles are the main precipitates in the HIP+SHT alloy beside the strengthener γ′ and γ′′. However, no evidence was found to support the small scale δ to be crack sites or propagation channel. 3.3. 3D reconstructive model and simulation results In order to understand the influence of the casting microporosity on LCF in the study, X-ray CT has been performed by using synchrotronic radiation source at SSRF. Digital volume of the superalloy was obtained via 3D reconstruction and used for subsequent numerical simulation by FEM software. Variation of the porosity was measured with the as-cast superalloy specimens, revealing that the solidification conditions has important influence on the materials densification, as shown in Fig. 11 (a). It’s recognized that smoothening of the 3D reconstruction model was conducted by mathematical algorithm. Statistic results of the 3D digital models in terms of porosity, max diameter of pore, and number of pore are illuminated in Fig. 11 (b), indicating that the densification of the superalloy increases from no.1 to no.4. Numerical simulation was conducted to investigate the fatigue properties of the nickel based superalloy by using the 3D reconstructive models obtained from X-ray CT. Finite element model was established by the software ABAQUS CAE and the material properties were set as follows: density ρ=8220 kg/m3, Young's modulus E=186GPa, and Poisson's ratio

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λ=0.33. The bottom end of the cylindrical superalloy model was applied with the boundary condition of displacement d = 0 and the top end of the model was applied a small tensile load of 1.0×10-7m in +Z direction (the axial direction of the specimens). Only the elastic properties were considered in the study because the tensile load was too small to cause plastic deformation of the superalloy. A simulation step of implicit dynamic procedure was subjected to calculate under the boundary conditions. The predicted Mises stress distribution of the FE model no.1 in Fig. 11 was obtained as shown in Fig. 12(a). A horizontal path 1 and a vertical path 2 through the micropore with the maximum stress value were created, as shown in Fig. 12 (a). Fig. 12 (c) illuminated the Mises stress along path 1 and path 2, indicating that the width and height of the pore profile are 58.03μm and 16.08μm, respectively. Stress concentration appears in horizontal direction (path 1) at the location of the casting pore while stress relaxation occurs in the vertical direction (path 2). Subsequently, the calculated stress field was submitted to calculate fatigue life by the platform fe-safe. In preprocessing, both the stress field and a triangular wave were used to create an elastic loading block, then whole model was analyzed. The surface roughness was set as 0.6
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the max diameter of the pore is parallel to the loading direction, the fatigue life in the vicinity of pore keeps at average level. The predicted LCF life distributions on the surface of micropores in Fig. 13 provide useful information to understand the influence of the morphology and distribution of microporosity on fatigue life. 4.

Discussions

4.1. Role of casting microporosity Micropores in the crack initiation zone and crack propagation zone and fatigue cracks initiating from the micropore were observed on the fracture surface on SHT superalloy, see Fig. 14 (a), indicating that the micropores in superalloy casting were the crack sites during the LCF experiments. Parameters influencing the distribution of stress and fatigue life include porosity (ratio of micropore volume to whole alloy volume), pore volume, the maximum equivalent diameter, structural complexity (the number of the channels in micropore), spacing between two micropores, and the length from microporosity to material surface. As an instance, the predicted fatigue cycle life of the nickel based superalloy studied here was affected by those parameters, as shown in Fig. 12. Increasing of the structural complexity, micropore volume, the maximum equivalent diameter, spacing between two micropores, and the length from microporosity to material surface, results in the decreasing of the predicted fatigue cycle life and the largest loading stress of the nickel based superalloy. In addition, the microporosity prefers to form at the interdendritic region where alloy elements segregate to form precipitates including of Laves, carbides, and δ particles. Combination formation of the microporosity and those precipitates were commonly observed in the as-cast and SHT superalloy. In Fig. 14 (b), the microcracks initiating from

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microporosity pass through the interdendritic region where the fracture of precipitates (namely carbides and δ stacks) were observed, as shown in Fig. 14 (c) and (d). 4.2. Role of carbide and δ phase Carbide was observed in all the superalloys. The shape of carbide in the nickel based superalloy includes two types: blocky and dendrite. Firstly, the blocky carbide containing a certain level of nitrogen has an octahedral crystal structure and generally nuclei grows based on an alumina embryo in center. While the dendrite carbide grows based on the blocky carbide and well develops into dendritic morphology due to the local carbon enrichment. Electron diffraction spectrum (EDS) of the broken dendrite carbide in the SHT superalloy reveals that the compositions of MC carbide include a high level of niobium, as shown in Fig. 15. Under the alternative loading conditions here, a major crack along the dendrite carbide connects several parallel minor cracks throughout the carbide. It is interesting to note that the large scale carbide prefers to break itself rather than spread the crack to the surrounding material. The TEM picture and diffraction pattern of the blocky carbide in SHT superalloy reveals the certain coherent relationship between blocky carbide and nickel base, as shown in Fig. 6 (b). Some dimples distributing around the carbide indicates that the plastic deformation of surrounding mass may prevent the interior crack in carbide from propagating towards matrix. As one of the important interdendritic or intergranular precipitates, acicular δ particles with were observed in the as cast superalloy, SHT superalloy, and HIP+SHT superalloy. However, the differences in quantity and morphology of the acicular δ particles in the superalloys, as shown in Fig. 4 (a), (b) and (d), result in the different crack mechanism of the δ particles. The

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acicular δ particles is minority phase in the as cast superalloy and has not significant implication on the mechanical properties. Fatigue cracks propagated through the stack of the large scale acicular δ particles in the SHT superalloy, which was obviously different from the case of HIP+SHT superalloy. The small scale acicular δ particles in the HIP+SHT superalloy was not observed to be crack source nor crack propagation channel. The smooth fracture surface of grain boundary was also observed in the polycrystalline superalloy indicating the large crack rate at the GB. The fatigue striations at the GB interface reveals that the GB fracture occurs after cyclic deformation of the grains. When the casting defect (e.g. cast pore) was eliminated in HIP superalloy, grain boundaries play the role of crack source in LCF of the material at room temperature, as shown in Fig. 9 (a). Since the movement of GB during HIP processing is commonly hindered at carbide and inclusion, those particles at GB promote the crack of GB during cycle loading. While small scale acicular δ forming at GBs during the SHT process after HIP can improve the strength of GBs, promoting the fracture mechanism transferring from GB fracture to surface fracture caused by surface quality, as shown in Fig. 10 (a). 4.3. LCF crack mechanism of the superalloys The precipitates and microporosity in the nickel based superalloy can be significantly changed by using SHT and HIP processes, resulting in the different fatigue properties and crack mechanisms of the superalloys. In summary, the SHT process can enhance the fracture strength of the superalloy since the brittle interdendritic Laves transforms into δ particles to strengthen the GBs. However, the microporosity in the superalloy cannot be eliminated by

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SHT process. In the LCF of the SHT superalloy, subsurface microporosity in the specimen was observed to cause primary crack. The HIP can eliminate microporosity, TCP, and δ particles in the superalloy. While the decreasing of the δ particles results in the reduction of the metal strength. As an additional case, the HIP+SHT reintroduces smaller scale δ particles at the GBs and interdendritic zone to improve the strength. Ductile fracture dominates the fatigue crack behaviors of the HIP superalloy and the HIP+SHT superalloy. Experimental confirmations suggest the GB crack mechanism in HIP superalloy and the surface crack mechanism in HIP+SHT superalloy. The carbides are difficult to change by the three processes, so carbides are the common fatigue crack sources in the all superalloys, especially the dendritic carbide. 5.

Conclusions Fatigue crack mechanisms of a heat treated polycrystalline nickel-based superalloy was

investigated with a combination of LCF experiments and numerical simulation. Conclusions can be summarized as follows: (1) The effects of heat treatment processes on the fatigue properties of the nickel based superalloy can be described as following: the SHT process enhance the strength by transforming the Laves into δ particles, while HIP process can improve the fatigue life by eliminating casting microporosity. The combination of HIP and SHT can balance the strength and fatigue life. (2) Micropore cracking is the main crack mechanism in the LCF of the SHT superalloy. Results finite element simulation reveal that the stress concentration on the surface of micropore causes the fracture and low LCF life at interdendritic region of micropore.

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While the elimination of micropore results in the GB crack mechanism of the HIP superalloy. Reintroduction of small scale δ particles in the HIP+SHT superalloy improve the GB strength, leading to the surface crack mechanism. (3) Both blocky carbide and dendritic carbide prefer to crack under cyclic loading, while plastic deformation of surrounding mass hinders the crack propagation from carbide to nickel base. The fracture of dendritic carbide exhibits a major central crack connecting several parallel minor cracks throughout the particle. (4) The quantity and morphology of δ particles can significantly affect LCF crack mechanism of the polycrystalline superalloy. The large scale δ particles in the SHT superalloy provides LCF crack propagation channel while the small scale δ particles can hinder cracking along GBs. Acknowledgment This study was supported by the Project funded by China Postdoctoral Science Foundation (No. 2016M591668) and the Aeronautical Science Foundation of China (No. 2016ZE57010). The authors would like to thank Dr. Han Guo and Dr. Rongchang Chen from Shanghai Synchrotron Radiation Facility (SSRF) for their help in conducting X-ray CT experiments. References [1] Xu J, Huang Z, Jiang L. Effect of heat treatment on low cycle fatigue of IN718 superalloy at the elevated temperatures. Materials Science and Engineering: A 2017; 690: 137-45. [2] Kunz L, Lukáš P, Konečná R. High-cycle fatigue of Ni-base superalloy Inconel 713LC. Int J Fatigue 2010; 32(6): 908-13.

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Captions of Table and Figure

Table 1 Chemical compositions of the nickel-based superalloy in the study.

Fig. 1. Block diagram of the experimental procedure and simulation processing. Fig. 2. Dimension of the LCF test specimen in the study. Fig. 3. Synchrotron radiation X-ray scanning experiment and 3D reconstructive model for FEM simulation. Fig. 4. Microstructure evolution in the nickel-based superalloy with different heat treatment regimens: (a) the as-cast case, (b) SHT, (c) HIP, (d) HIP+SHT. Fig. 5. SEM micrographs of the SHT superalloy: (a) the interdendritic precipitates, (b) the γ′ and disk-like γ′′ beside acicular δ, (c) the disk-like γ′′ distributes in different directions, and (d) γ′′ in nickel matrix. Fig. 6. TEM micrographs of the SHT superalloy: (a) the γ-base and disk-like γ′′ and diffraction pattern //<111>, (b) blocky carbide and acicular δ, (c) multilayer substructure in acicular δ, and (d) the interface zone between δ and γ-base. Fig. 7. The stress amplitude as a function of the cycle number in the LCF experiments. Fig. 8. LCF fractography of the SHT superalloy: (a) the macro fracture surface of the testing bar, (b) crack initiating from the microporosity in the subsurface layer, (c) morphology of crack propagation zone, (d) cleavage striations in the final rupture zone, (e) the acicular δ

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particles on the surface of microporosity, and (f) plastic fatigue striations in crack propagation zone. Fig. 9. LCF fractography of the HIP superalloy: (a) the crack initiating from intergranular fracture, (b) the crack initiation zone, (c) fatigue striations in crack propagation zone in different directions, (d) fatigue striations in the final rupture zone. Fig. 10. LCF fractography of the HIP+SHT superalloy: (a) the macro morphology of the fatigue test bar with typical river pattern, (b) the crack initiation zone on the surface of test bar, (c) the ductile fatigue striations in crack propagation zone, and (d) secondary crack in the final rapture zone. Fig. 11. (a) 3D reconstructive model of the as cast superalloy, and (b) statistic results of the porosity, pore max diameter, and pore number. Fig. 12. (a) The predicted Mises stress distribution in the tensile simulation and (b) the predicted fatigue life of the superalloy no.1, (c) Mises stress distribution along path 1 and path 2 at a cast pore, and (d) fatigue life distribution along path 3 and path 4. Fig. 13. The predicted fatigue cycle life distribution on the surface of microporosity with different morphology (pore volume decreases from pore 1 to pore 15), and the load direction is +Z axis. Fig. 14. Effects of the microporosity on the fatigue crack behavior in the SHT superalloy: (a) a cast micropore, (b) cracks initiating from the micropore, (c) cracks propagate through a carbide particle, and (d) cracks in the vicinity of δ particles.

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Fig. 15. Composition analysis of carbide by EDS: (a) fracture surface of carbide, (b) composition distributions along the line 1, (c) total spectrum, and (d) weight percentage of chemical elements.

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Highlights •

Fractography of LCF experiments at room temperature was analyzed in detail.

•

Stress concentration around casting micropore was analyzed by X-ray CT and FEM.

•

Combination of micropore and precipitates dominates the LCF crack initiation behaviors.

•

Effects of heat treatment process on the LCF of a superalloy was summarized.

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Table 2 Chemical compositions of the nickel-based superalloy in the study. C

Cr

Ni

Mo

Nb+Ta

Al

Ti

Fe

wt. %

0.05

19.0

53.0

3.0

5.0

0.50

0.90

rest

at. %

0.243 21.313 52.668 1.824

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2.375

1.081 1.097 rest

LCF

Investment casting process

Heat treatment and HIP process

Superalloy casting chipping

Fatigue specimens

Acicular bars preparation

Low cycle fatigue experiment

Synchrotron radiation X-ray CT

FEM

OM/SEM observation

3D model reconstruction

ABAQUS CAE simulation

Fatigue crack mechanisms

XCT

fe-safe fatigue life calculation Simulation results analysis

Fig. 1. Block diagram of the experimental procedure and simulation processing.