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Effects of stress and temperature on creep behavior of a new third-generation powder metallurgy superalloy FGH100L
Materials Science & Engineering A 776 (2020) 139007
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Effects of stress and temperature on creep behavior of a new third-generation powder metallurgy superalloy FGH100L Tian Tian a, Zhibo Hao a, Changchun Ge a, **, Xinggang Li b, *, Shiqing Peng a, Chonglin Jia c a
Institute of Powder Metallurgy and Advanced Ceramics, School of Materials Science and Engineering, University of Science and Technology Beijing (USTB), Beijing, 100083, China b Academy for Advanced Interdisciplinary Studies and Department of Mechanical and Energy Engineering, Southern University of Science and Technology (SUSTech), Shenzhen, 518055, China c Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials (BIAM), Beijing, 00095, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Third generation PM nickel-based superalloy Creep property Creep mechanism Fracture characteristic
In this paper, a new third generation powder metallurgy nickel-based superalloy, designated as FGH100L, was prepared by Spray Forming, followed by Hot Isostatic Pressing, Isothermal Forging and Heat Treatment. The effects of stress and temperature on the creep properties of the alloy were studied by analyzing the creep mechanism and fracture characteristics. The results show that the creep life of the alloy decreases significantly with the increase of stress and temperature. The creep fracture indicates mixed intergranular and transgranular fracture characteristics for all the tested specimens. However, at low stress or low temperature conditions, the intergranular fracture is more prominent in the fracture source region and almost no shear lip region can be observed in the final rupture region. With the increase of stress and temperature, the transgranular fracture becomes more prominent, and the shear lip region is enlarged. The misorientation within the grain decreases with the increase of stress and temperature, indicating that the higher the stress and the temperature, the lower the local strain the alloy can bear. Under the creep condition of high temperature and high stress, the misori entation is much larger around the original grain boundary than in the interior of the original grain, indicating that the internal strain of the alloy is mainly concentrated near the original grain boundary under this condition.
1. Introduction Nickel-based superalloys are important materials for manufacturing key hot-end components of aero-engines. However, the failure of most load-bearing components in high temperature environment is caused by high temperature creep which results from high temperature and high pressure [1,2]. Creep will lead to excessive plastic deformation or creep stress fracture of components in aero-engine at high temperature, which become more prominent especially with the further increase of frontal turbine temperature. Creep properties and creep behaviors are closely related to the high temperature strength, damage tolerance, structural stability and safety factor of the alloy. Therefore, the study of creep characteristics of superalloys can provide a basis for their life prediction and safe use. Mohammad et al. [3] studied the high-temperature creep behavior of the Inconel-713C nickel-based superalloy used in turbine blades under
850 � C and at different stress levels. Farangis et al. [4] studied the origin of creep dislocation in nickel-based single crystal superalloy under me dium stress and temperature. Wollgramm et al. [5] studied the creep behavior of nickel-based single-crystal superalloy, mainly analyzing the effects of stress and temperature on the minimum creep rate. Tian et al. [6,7] investigated the influence of element Re on the deformation mechanism of γ0 phase in nickel-based single-crystal superalloy during high temperature creep. Zhou et al. [8] studied the evolution rules of sub-structures by interrupting creep tests at different creep stages. Heather et al. [9] studied the effects of super solvus heat-treatments on the microstructures and properties of LSHR, Alloy10 and RR1000 nickel-based superalloys. The aim is to establish the relationship be tween microstructure and mechanical properties such as creep strength and creep crack propagation. Peng et al. [10] studied the creep prop erties of FGH96 nickel-based superalloy under different aging heat treatments at 700 � C and 690 MPa.
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (C. Ge), [email protected] (X. Li). https://doi.org/10.1016/j.msea.2020.139007 Received 22 October 2019; Received in revised form 21 January 2020; Accepted 23 January 2020 Available online 27 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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Xie et al. [11] studied dislocation morphology, creep behavior and dislocation networks of FGH95 Ni-based superalloy during creep pro cesses. Under the creep testing condition of 705 � C and 897 MPa, Jia et al. [12,13] demonstrated the effects of solution treatment and cooling rate on the microstructure, fracture and creep performance of FGH100L nickel-based superalloy. Currently, aero-engine turbine disks are mainly manufactured using Cast & Wrought (C&W) technology and Powder Metallurgy (PM) tech nology [14,15]. Spray Forming (SF) [16], which was developed in the early 1980s to compete with traditional superalloy manufacturing pro cesses, has significantly reduced the processing steps from atomization to final preforming compared with C&W or PM technologies. The SF process eliminates not only remelting and transformation steps in C&W processing, but also the steps of sieving, electrostatic precipitation, tank filling and degassing of the powder in PM processing. Therefore, the SF technology has become an efficient manufacturing alternative. The SF process is followed by hot working, i.e., Hot Isostatic Pressing (HIP) and Isothermal Forging (IF), to make the alloy denser. The heat treatment (HT) system of solid solution and two-stage aging is designed to improve the microstructure of the alloy, and the goal is to improve the overall properties of the alloy. The raw material used in this paper is the new third-generation PM superalloy FGH100L [12,13], which the authors’ team has been researching and developing for many years.
Fig. 1. Processing diagram of creep specimen.
variation of creep rate of the FGH100L alloy, respectively, under different stress loads and different temperatures. As shown in Fig. 2(a), the total creep life of the alloy decreases significantly with the increase of the applied stress load (from 690 MPa to 897 MPa) at 705 � C. When a low stress load of 690 MPa is applied, the steady creep duration of the alloy is the longest, about 300 h. The maximum total creep strain is 24.9%, with a low steady creep rate (0.0016 � 10-2h-1, see Fig. 3(a)) and a long creep life (1286.27 h). As the applied stress increases from 750 MPa to 897 MPa, the total creep strain of the alloy firstly decreases and then increases, and the creep life de creases significantly (see Fig. 2(a)). The creep life under 750 MPa is 920.36 h, and it is reduced to 81.55 h under 897 MPa. With the increase of stress load, the creep rate of the alloy increases, and the duration of the second creep stage, namely the steady creep stage, decreases (see Fig. 3(b), (c) and (d)). The steady-state creep rate of the alloy under a stress load of 897 MPa is the highest (0.1199 � 10-2h-1). Fig. 2(b) shows the creep curves of the FGH100L alloy under different stress loads at 750 � C. As shown in Fig. 2(b), the total creep life of the alloy decreases significantly with the increase of the applied stress load (from 450 MPa to 897 MPa). When a low stress load of 450 MPa is applied, the total creep strain of the alloy is relatively high (37.1%, see Fig. 2(b)), and the alloy has a long creep life (629.18 h). As for the steady creep stage, the alloy has a low creep rate (0.0010 � 10-2h-1, see Fig. 3 (e)) and the steady state creep duration is about 300 h. As the applied stress increases from 750 MPa to 897 MPa, the alloy creep rate rapidly increases and the steady state creep duration reduces (see Fig. 3(f), (g) and (h)). The total creep strain firstly increases and then decreases. The creep life decreases significantly from 62.16 h under 750 MPa to 11.49 h under 897 MPa. The highest creep rate is 0.2705 � 10-2h-1 under 897 MPa. As shown in Fig. 2(a) and (b), the creep properties of the FGH100L alloy are sensitive to the applied stress at a constant temperature. With the decrease of the applied stress, the second and third stages of creep of the alloy are prolonged, and thus the creep life of the alloy increases. On the contrary, the creep life of the alloy decreases with the increase of the applied stress. Fig. 2(c) shows the creep curves of the FGH100L alloy at different temperatures under a stress load of 897 MPa. Under the same stress load, the total creep life of the alloy decreases significantly with the increase of temperature. The creep curve under this condition still has three stages of creep, but the second stage lasts for a short time and then the alloy quickly enters the third stage. Under the condition of 650 � C/897 MPa, the total creep life of the alloy is 1477.42 h, the total strain is 5.2%, and the steady creep rate is 0.0004 � 10-2h-1 (see Fig. 3(i)). As the testing temperature increases from 705 � C, 725 � C–750 � C, the duration of the steady creep stage gradually reduces from 50 h, 12 h–2 h, and the creep rate of the second stage gradually increases from 0.1199 � 10-2h-1,
2. Materials and methods The FGH100L master alloy was prepared by a double smelting pro cess, i.e., vacuum induction smelting plus vacuum consumable remelt ing (VIM þ VAR). The main chemical composition of the alloy 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 alloy ingot (Φ200 mm � H300 mm) was prepared by SF, during which high purity N2 was used as an atomizing gas. Then it was treated by HIP without canning. The ingot was first heated to 1160 � C at a rate of 10 � C/min, holding for 3 h at 150 MPa, and then cooled to room tem perature at a cooling rate of 25 � C/min. Following this, an IF experiment was carried out, in which the forging temperature was 1150 � C, the pressing rate was 0.1 mm/s, and the engineering deformation was about 30.4%. After forging, the ingot was covered with asbestos until it was cooled to room temperature. Samples were taken from the IFed ingot for HT. The HT process was solid solution treatment plus two-stage aging: 1130 � C/1 h/fan quenching þ850 � C/4 h/AC þ 775 � C/8 h/AC. In this paper, the creep tests were carried out according to China national standard GB/T 2039–2012 uniaxial tensile creep test method for metallic materials [17]. The test equipment was the CSS type creep test machine of Changchun Test Machine Factory, and the specimen was non-standard creep specimen with a Ф5 mm round shape. The specific sample size is shown in Fig. 1. The creep temperature range was 650–750 � C, and the stress range was 450–897 MPa. Two samples were taken for creep testing under each condition. The grain structure was observed by a DC300 metallographic optical microscope (OM), and the metallographic etchant was 10 g CuCl2 þ 50 mL HCl þ 50 mL H2O. The fracture characteristics and microstructures were observed by Scanning Electron Microscope (SEM), JSM-6701F and ZEISS EVO®18. The dis tributions of grain orientation and misorientation were obtained by Electron Backscattering Diffraction (EBSD) technique. The electrolytic polishing solution was 20% H2SO4 þ 80% CH3OH and the electrolytic etching solution was 9 g CrO3 þ 90 mL H3PO4 þ 30 mL C2H5OH. 3. Results 3.1. Effects of stress and temperature on creep properties The general creep curve can be roughly divided into three stages: the first stage, steady creep in the second stage and accelerated creep in the third stage. Fig. 2 and Fig. 3 show the creep curves and the temporal 2
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Fig. 2. Creep curves of FGH100L alloy under different stress and temperature conditions: (a) 690–897MPa/705 � C, (b) 450–897MPa/750 � C, (c) 897MPa/ 650–750 � C.
0.1598 � 10-2h-1 to 0.2705 � 10-2h-1 (see Fig. 3(d), (j) and (h)). After entering the third creep stage, the creep rate increases with the increase of temperature, leading to a significant reduction of creep life. It can be seen from Fig. 2(c) that the creep strain of the third stage of the alloy changes greatly with the increase of temperature, which reflects that the influence of temperature on the microstructural changes, crack initia tion and growth, and fracture mode in this stage may be different. Overall, the creep life of the alloy will be reduced when operating at a higher temperature.
Fig. 4(a) and (b) reflect the liner relationships between lgε s and lgσ a at 705 � C and 750 � C, respectively. The slope value of the line (namely the stress index n) is 16.33 and 8.10, respectively, which is closely related to the creep mechanism. In general, the creep deformation of metal materials is very sensitive to the influence of stress factor when the stress index n is larger than 5, which is in accordance with the dislo cation mechanism during creep [21,22]. Therefore, the creep mecha nism of the FGH100L alloy under different stress loads at 705 � C and 750 � C is dislocation creep. Actually, the high n value in the alloy is related to the properties of the grain boundaries and the strengthening effect of γ0 phase which has a large volume fraction in the alloy. Fig. 5 reflects the relationship between lgε_ s and T-1 under a stress load of 897 MPa. According to the slope of the curve, the steady-state creep activation energy was derived as Qc ¼ 339.35 kJ/mol, which is much higher than the self-diffusion activation energy of Ni in austenite (265–285 kJ/mol). This can be attributed to the γ0 precipitates that take up about 52% mass fraction in the FGH100L alloy. The self-diffusion activation energy of polycrystalline nickel-based superalloys is gener ally ~300 kJ/mol [22,23]. The higher Qc value indicates that the creep mechanism of the FGH100L alloy is controlled by the atomic diffusion process. Dislocation climbing, directional vacancy diffusion and grain boundary sliding all involve atomic diffusion processes. The higher the Qc value, the more difficult for creep deformation to occur. Meanwhile, a higher Qc indicates higher creep resistance, higher creep strength and longer creep life of the alloy [18]. �
3.2. Prediction of creep life of FGH100L alloy Creep rate is one of the most important parameters to describe the creep behavior of an alloy. According to the previous analysis, the creep rate of the FGH100L alloy is significantly affected by load stress and service temperature. Generally, Norton’s Power Law is satisfied between the minimum creep rate, creep stress and temperature, as shown in Eq. (1). Therefore, the creep rate of the FGH100L alloy at the steady state stage can be described by Norton’s formula [18–20]: �
ε s ¼ Aσ na expð
Qc = RTÞ
(1)
In Eq. (1), ε_ s is the steady-state creep rate, A is the constant related to the structure of the material, T is the absolute temperature, R is the universal gas constant (8.314 J mol-1K-1), Qc is the apparent creep activation en ergy, n is the stress index, and σ a is the applied stress. At a given temperature, take the logarithm of both sides of Eq. (1) and convert it to Eq. (2): �
lgε s ¼ nlgσ a þ C
3.3. Creep fracture characteristics of FGH100L alloy Fig. 6 shows the creep fracture morphology of the FGH100L alloy under the condition of 705 � C/690 MPa. Fig. 6(a) shows the macroscopic morphology of creep fracture of the alloy. It can be seen that the neckshrinking phenomenon occurs. The fracture surface is flat. There is almost no shear lip in the final rupture region, and the propagation re gion accounts for a large area. The fracture source region demonstrates mixed intergranular and transgranular fracture characteristics, the former being more prominent. “Hole” type cracks can be seen, as indi cated by the arrow in the inset at the upper-right corner of Fig. 6(b). Fig. 6(c) and (d) manifest that there are a large number of steps in the
(2)
In Eq. (2), C is a constant, lgε_ s is linear with lgσ a and the slope of the straight line is n. At a given stress load, take the logarithm of both sides of Eq. (1) and convert it to Eq. (3): �
lgεs ¼
Qc =RT þ C’
(3)
In Eq. (3), C0 is a constant, and lgε_ s has a linear relationship with T-1. 3
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Fig. 3. Creep rate versus creep time for FGH100L alloy under different stress and temperature conditions: (a)–(d) 690–897MPa/705 � C, (e)–(h) 450–897MPa/750 � C, (i)(j) 897MPa/650–725 � C.
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Fig. 4. lgεs �
lgσ a curves of FGH100L alloy at different temperatures and stress loads: (a) 705 � C/690-897 MPa; (b) 750 � C/450-897 MPa.
propagation region and final rupture region, are obviously presented. Overall, the fracture surface is rough and uneven. Compared with Fig. 6 (a), the shear lip of the final rupture region is larger under 897 MPa in Fig. 7(a). The fracture source region of the alloy is characterized by mixed intergranular and transgranular fractures, the latter becoming more prominent. “Wedge” cracks can be seen, as indicated by the arrow in Fig. 7(b). Fig. 7(c) shows that in the propagation region, there are deep dimples, secondary cracks, and small facets where deep slip lines can be seen (see the enlarged area at the lower left corner in Fig. 7(c)). There are shallow dimples in the shear lip region (see Fig. 7(d)). Fig. 8 shows the creep fracture morphology of the FGH100L alloy under the condition of 650 � C/897 MPa. As can be seen in Fig. 8(a), the fracture shows severe neck shrinkage, large propagation region and almost no shear lip in final rupture region. Fig. 8(b) shows that the intergranular fracture is prominent in the fracture source region. There are secondary cracks, many steps, facets (see Fig. 8(c) and its enlarged region), deep dimples, and slip lines (see Fig. 8(d) and its enlarged re gion) in the propagation region. Fig. 9 shows the creep fracture morphology of the FGH100L alloy under the condition of 750 � C/897 MPa. Fig. 9(a) shows that the creep fracture of the alloy is uneven, with fracture source region, propagation region and large shear lip in the final rupture region. Intergranular fracture morphology and long “wedge” cracks are visible, and the transgranular fracture becomes prominent in the fracture source region
Fig. 5. Relationship between lgεs and T-1 of FGH100L alloy at 897 MPa. �
propagation region and slip lines and deep dimples on the facets. Fig. 7 shows the creep fracture morphology of the FGH100L alloy under the condition of 705 � C/897 MPa. As shown in Fig. 7(a)–(d), three characteristic regions of creep fracture, i.e., fracture source region,
Fig. 6. Creep fracture morphology of FGH100L alloy at 705 � C/690 MPa: (a) macroscopic morphology of fracture, (b) fracture source region, (c) steps and slip lines, (d) deep dimples. 5
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Fig. 7. Creep fracture morphology of FGH100L alloy at 705 � C/897 MPa: (a) macroscopic morphology of fracture; (b) fracture source region; (c) deep slip lines; (d) shallow dimples.
Fig. 8. Creep fracture morphology of FGH100L alloy at 650 � C/897 MPa: (a) macroscopic morphology of fracture, (b) fracture source region, (c) steps and facets, (d) deep dimples and slip lines.
(see Fig. 9(b)). In Fig. 9(c), there are a few steps with shallow slip lines and a few secondary cracks. Deep dimples can be seen in the shear lip region (see Fig. 9(d)). Fig. 10 shows the creep fracture morphology of the FGH100L alloy under the condition of 750 � C/450 MPa. It can be seen from Fig. 10(a) that the creep fracture morphology of the alloy is similar to that under 650 � C/897 MPa in Fig. 8(a). Fig. 10(b) shows that the fracture char acteristics of the source region are mixed intergranular and trans granular, the former being prominent. In Fig. 10(c), steps can be seen in the propagation region and the steps are covered with a large number of deep dimples and “hole-shaped” cracks can be seen. According to the analysis above, there are three kinds of conditions when the alloy exhibits good creep properties: (1) low stress and low temperature, (2) high stress and low temperature, and (3) low stress and
high temperature. In these cases, neck-shrinking phenomenon occurs on creep fracture, which shows the characteristics of mixed intergranular and transgranular fracture, the former being prominent. The fracture surface is relatively flat, there is almost no shear lip area, and the expansion region occupies a large area. In case (1) and case (2), there are many steps, planes and sliding lines on the surface. In case (2) and case (3), the “wedge” and “hole” cracks are easily formed. The formation of cracks can be divided into two processes: nucleation and growth. Both wedge-shaped and hole-shaped cracks are related to the slip behavior of grain boundary [24]. 4. Discussion Fig. 11 shows the local misorientation distribution of the FGH100L 6
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Fig. 9. Creep fracture morphology of FGH100L alloy at 750 � C/897 MPa: (a) macroscopic morphology of fracture, (b) fracture source region and cracks, (c) steps, (d) shear lip region.
Fig. 10. Creep fracture morphology of FGH100L alloy at 750 � C/450 MPa: (a) macroscopic morphology of fracture, (b) fracture source region and cracks, (c) steps and deep dimples.
alloy after creep fracture under different stress loads (450 MPa, 897 MPa) and temperatures (650–750 � C). As shown in Fig. 11(a), a large misorientation can be seen within the alloy after testing at 650 � C/897 MPa for 1477.4 h, indicating that there is a large strain inside the alloy. As shown in Fig. 11(b), after testing for 81.5 h under 705 � C/897 MPa, the grain misorientation of the alloy is smaller than that in Fig. 11(a), indicating that the internal strain of the alloy is relatively low. In this case, the strain and dislocation are mostly distributed around grain boundaries. Fig. 11(c) shows that after testing under 725 � C/897 MPa for 26.1 h, many small recrystallized grains are present in the alloy surrounding large original grains. The misorientation of the recrystal lized grains is smaller than that of the original grains, which means that the strain is relatively lower within the recrystallized grains. It can be inferred that the small recrystallized grains relieve strain concentration during creep by grain rotation, while the large original grains relieve strain concentration by grain distortion. As shown in Fig. 11(d), the microstructure of the alloy after testing for 11.5 h at 750 � C/897 MPa is
similar to that in Fig. 11(c), but the difference is that the overall misorientation of the alloy is lower than that in Fig. 11(c). It is shown in Fig. 11(e) that after testing at 750 � C/450 MPa for 629.2 h, the internal misorientation of the alloy is significant, indicating that there are large strains and many dislocations inside the grains. Fig. 11(f) shows that with the increase of temperature from 650 � C to 750 � C at 897 MPa, the grain misorientation in the alloy reduces, indicating that the higher the temperature, the lower the strain that the alloy can bear. This corre sponds with the fact that the higher the temperature, the lower the creep life. This suggests that the grain boundary becomes the weak location of the alloy with the increase of temperature. At 750 � C, the grain misorientation of the alloy decreases with the increase of stress from 450 MPa to 897 MPa, but the internal strain of the alloy is mainly distributed near the grain boundary, which is the weak location at this temperature. When the stress is relatively low, the dislocation multi plication rate in the alloy is slow, which provides enough time for the strain inside the alloy to release. Therefore, dislocations move into the 7
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Fig. 11. Local misorientation distribution of FGH100L alloy after creep fracture under different stress and temperature conditions (EBSD): (a) 650 � C/897 MPa, (b) 705 � C/897 MPa, (c) 725 � C/897 MPa, (d) 750 � C/897 MPa, (e) 750 � C/450 MPa, (f) comparison of local misorientation.
grain by sliding, thus enlarging the grain misorientation in the alloy and improving creep life. Fig. 12 shows the local orientation distribution of the FGH100L alloy after creep fracture under different stress loads (450 MPa, 897 MPa) and temperatures (650–750 � C). It can be seen that the grain orientation is random, and the color gradient change indicates that the misorientation exists in the alloy grains. There are many fine recrystallized grains around the large original grains and the orientation of the recrystallized grains is randomly distributed. In Fig. 12 (a) the color in some grains changes gradually from boundary to interior, indicating a large misorientation. In Fig. 12(b), (c) and (d), the color within grains is relatively uniform, indicating a small misorientation, while the color
near the boundary changes suddenly, exhibiting a large misorientation. It can be seen from Figs. 11 and 12 that, under the same stress load, the internal stress inside the alloy changes from grain interior to grain boundary after creep fracture as the temperature increases. Nickel-based superalloys have face-centered cubic (FCC) structure, in which slip is an important mode of deformation. The grain boundary slip (GBS) of fine grains that are obtained after sub-solid solution treatment is very important in the deformation of polycrystalline su peralloys during creep at high temperatures [25]. Combining the creep performance in Figs. 2 and 3 with the grain misorientation results in Figs. 11 and 12, the strain rate accelerates sharply in the acceleration phase under high stress load. When the strain rate is too fast for the grain 8
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Fig. 12. Local orientation distribution of FGH100L alloy after creep fracture under different stress and temperature conditions (EBSD): (a) 650 � C/897 MPa, (b) 705 � C/897 MPa, (c) 725 � C/897 MPa, (d) 750 � C/897 MPa, (e) 750 � C/450 MPa.
boundary to adjust, intragranular slip plays a significant role in creep deformation. In this case, the strain tends to concentrate in the slip zone. Since microcracks are easy to form at the location where the strain is concentrated, the secondary cracks nucleate along the slip zone under high stress. During creep, grain boundary deformation is initially acti vated, and when creep strain rate increases, grain matrix deformation coordinates with GBS [26,27].
(3) Under the stress load of 897 MPa, the grain misorientation of the alloy decreases with the increase of temperature from 650 � C to 750 � C. At 750 � C, the grain misorientation of the alloy decreases with the increase of stress load from 450 MPa to 897 MPa. This indicates that under the high temperature or high stress condi tions, the local strain the alloy can bear reduces greatly. With the increase of the temperature or the stress load, the grain misori entation around the original grain boundary becomes more sig nificant than within the original grain, indicating larger strain concentration near the original grain boundary. At the low tem perature or stress conditions, the difference in misorientation distribution between grain boundary and interior is not very large.
5. Conclusions (1) The creep properties of the FGH100L alloy, including creep life, creep strain and creep rate, are very sensitive to the applied creep stress and temperature. Based on Norton’s Power Law, the value of the stress index n of the FGH100L alloy at 705 � C and 750 � C is 16.33 and 8.10, respectively, indicating the dislocation creep mechanism. Under the stress load of 897 MPa, the activation energy Qc of the FGH100L alloy during steady-state creep is 339.35 kJ/mol, indicating the atomic-diffusion controlled creep process. (2) Mixed intergranular and transgranular fracture characteristics were observed under all the testing conditions. At low stress or low temperature conditions, the intergranular fracture is more prominent. With the increase of stress and temperature, the transgranular fracture becomes more prominent. This may be ascribed to the high strain rate under the high stress or high temperature conditions, in which the strain rate is too fast for the grain boundary to adjust, so that the intragranular slip becomes more important in creep deformation.
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Tian Tian: Data curation, Methodology, Investigation, Writing original draft. Zhibo Hao: Data curation, Investigation. Changchun Ge: Project administration, Supervision, Writing - review & editing. Xing gang Li: Data curation, Methodology, Supervision, Writing - review & editing. Shiqing Peng: Data curation. Chonglin Jia: Data curation.
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Acknowledgements
[10] Z.C. Peng, G.F. Tian, J. Jiang, M.Z. Li, Y. Chen, J.W. Zou, F.P.E. Dunne, Mechanistic behaviour and modelling of creep in powder metallurgy FGH96 nickel superalloy, Mater. Sci. Eng. 676 (2016) 441–449. [11] J. Xie, S.G. Tian, L.J. Shang, X.M. Zhou, Creep behaviors and role of dislocation network in a powder metallurgy Ni-based superalloy during medium-temperature, Mater. Sci. Eng. 606 (2014) 304–312. [12] C.L. Jia, F.L. Zhang, K. Wei, S.M. Lv, Effective solution treatment can result in improved creep performance of superalloys, J. Alloys Compd. 770 (2019) 166–174. [13] C.L. Jia, Y. Li, F.L. Zhang, S.M. Lv, The creep behavior of a disk superalloy under different stress conditions, Met. Powder Rep. 73 (2018) 94–100. [14] G. Raisson, Evolution of PM nickel base superalloy processes and products, Powder Metall. 1 (2008) 10–13. [15] C. Crozet, A. Devaux, R. Forestier, S. Charmond, M. Hueller, D. Helm, W. Buchmann, Effect of ingot size on microstructure and properties of the new advanced AD730TM superalloy, in: Proceedings of the 13th International Symposium of Superalloys, TMS, New York, 2016, pp. 437–440. [16] H. Henein, V. Uhlenwinkel, U. Fritsching, Metal Sprays and Spray Deposition, Springer, Cham, 2017, pp. 497–520. [17] China national standardization administration, GBT 2039-2012 Uniaxial Tensile Creep Test Method for Metallic Materials, China Standard Press, Beijing, 2012, pp. 1–38. [18] J.T. Guo, C. Yuan, J.S. Hou, Creep and creep-fatigue-environment interaction and mechanisms of superalloys, Chin. J. Nonferrous Metals 3 (2011) 25–42. [19] D. McLean, Mechanical Properties of Metals, John Wiley & Sons Inc, 1962, pp. 308–311. [20] W.G. Kim, J.Y. Park, I. Ekaputra, S.J. Kim, M.H. Kim, Y.W. Kim, Creep deformation and rupture behavior of Alloy 617, Eng. Fail. Anal. 58 (2015) 441–451. [21] O.D. Sherby, P.M. Burke, Mechanical behavior of crystalline solids at elevated temperature, Prog. Mater. Sci. 13 (1968) 323–390. [22] S.H. Mo, J.H. Yu, J.J. Wang, Mechanical Properties of Engineering Materials, Peking, University Press, 2013, pp. 200–202. [23] S.J. Davies, S.P. Jeffs, M.P. Coleman, R.J. Lancaster, Effects of heat treatment on microstructure and creep properties of a laser powder bed fused nickel superalloy, Mater. Des. 159 (2018) 39–46. [24] G.L. Chen, Superalloy Theory, Metallurgical Industry Press, Beijing, 1988, pp. 239–243. [25] K.X. Dong, C. Yuan, S. Gao, C.L. Jia, J.T. Guo, C.C. Ge, Creep properties of a powder metallurgy disk superalloy at 700� C, J. Mater. Res. 32 (2017) 624–633. [26] A. Soula, Y. Renollet, D. Boivin, J.L. Pouchou, D. Locq, P. Caron, Y. Br�echet, Analysis of high-temperature creep deformation in a polycrystalline nickel-base superalloy, Mater. Sci. Eng. 510–511 (2009) 301–306. [27] T.G. Langdon, Grain boundary sliding as a deformation mechanism during creep, Phil. Mag. 22 (2006) 689–700.
The authors thank the group of Prof. Udo Fritsching and Dr. Volker Uhlenwinkel of Bremen University in Germany, General Manager Zhang Yu-Chun of Fushun Special Steel Shares Co., Ltd., Ye Jun-Qing of Guizhou Anda Aviation Forging Co., Ltd. and associate professor Zheng Wei-Wei of State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing for their valuable contri butions. The work was supported by the National Natural Science Foundation of China (No. 51171016) and Shenzhen Science and Tech nology Innovation Commission under (No. JSGG20180508152608855). References [1] T.P. Gabb, J. Telesman, P.T. Kantzos, Kenneth OConnor, Characterization of the Temperature Capabilities of Advanced Disk Alloy ME3, NASA Glenn Research Center, 2002, pp. 1–51. NASA/TM-2002-211796. [2] J. Jia, Y. Tao, Y.W. Zhang, Y. Zhang, J.T. Liu, Recent development of third generation P/M superalloy Ren� e104, Powder Metall. Ind. 3 (2007) 36–43. [3] A. Mohammad, A. Mahboobeh, Evaluation of high-temperature creep behavior in Inconel-713C nickel based superalloy considering effects of stress levels, Mater. Sci. Eng. 689 (2017) 298–305. [4] F. Ram, Z.M. Li, S. Zaefferer, S.M.H. Haghighat, Z.L. Zhu, D. Raabe, R.C. Reed, On the origin of creep dislocations in a Ni-base, single-crystal superalloy: an ECCI, EBSD, and dislocation dynamics-based study, Acta Mater. 109 (2016) 151–161. [5] P. Wollgramm, H. Buck, K. Neuking, A.B. Parsa, S. Schuwalow, J. Rogal, R. Drautz, G. Eggeler, On the role of Re in the stress and temperature dependence of creep of Ni-base single crystal superalloys, Mater. Sci. Eng. 628 (2015) 382–395. [6] S.G. Tian, J. Wu, D. Shu, Y. Su, H.C. Yu, B.J. Qian, Influence of element Re on deformation mechanism within γ’ phase of single crystal nickel-based superalloys during creep at elevated temperatures, Mater. Sci. Eng. 616 (2014) 260–267. [7] S.G. Tian, S. Zhang, F.S. Liang, A.N. Li, J.J. Li, Microstructure evolution and analysis of a single crystal nickel-based superalloy during compressive creep, Mater. Sci. Eng. 528 (2011) 4988–4993. [8] H.J. Zhou, H. Chang, Q. Feng, Transient minimum creep of a γ’ strengthened Cobase single-crystal superalloy at 900� C, Scripta Mater. 135 (2017) 84–87. [9] H.J. Sharpe, A. Saxena, Effect of microstructure on high-temperature mechanical behavior of nickel-base superalloys for turbine disc applications, Adv. Mater. Res. 278 (2011) 259–264.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Effects of stress and temperature on creep behavior of a new third-generation powder metallurgy superalloy FGH100L Tian Tian a, Zhibo Hao a, Changchun Ge a, **, Xinggang Li b, *, Shiqing Peng a, Chonglin Jia c a
Institute of Powder Metallurgy and Advanced Ceramics, School of Materials Science and Engineering, University of Science and Technology Beijing (USTB), Beijing, 100083, China b Academy for Advanced Interdisciplinary Studies and Department of Mechanical and Energy Engineering, Southern University of Science and Technology (SUSTech), Shenzhen, 518055, China c Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials (BIAM), Beijing, 00095, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Third generation PM nickel-based superalloy Creep property Creep mechanism Fracture characteristic
In this paper, a new third generation powder metallurgy nickel-based superalloy, designated as FGH100L, was prepared by Spray Forming, followed by Hot Isostatic Pressing, Isothermal Forging and Heat Treatment. The effects of stress and temperature on the creep properties of the alloy were studied by analyzing the creep mechanism and fracture characteristics. The results show that the creep life of the alloy decreases significantly with the increase of stress and temperature. The creep fracture indicates mixed intergranular and transgranular fracture characteristics for all the tested specimens. However, at low stress or low temperature conditions, the intergranular fracture is more prominent in the fracture source region and almost no shear lip region can be observed in the final rupture region. With the increase of stress and temperature, the transgranular fracture becomes more prominent, and the shear lip region is enlarged. The misorientation within the grain decreases with the increase of stress and temperature, indicating that the higher the stress and the temperature, the lower the local strain the alloy can bear. Under the creep condition of high temperature and high stress, the misori entation is much larger around the original grain boundary than in the interior of the original grain, indicating that the internal strain of the alloy is mainly concentrated near the original grain boundary under this condition.
1. Introduction Nickel-based superalloys are important materials for manufacturing key hot-end components of aero-engines. However, the failure of most load-bearing components in high temperature environment is caused by high temperature creep which results from high temperature and high pressure [1,2]. Creep will lead to excessive plastic deformation or creep stress fracture of components in aero-engine at high temperature, which become more prominent especially with the further increase of frontal turbine temperature. Creep properties and creep behaviors are closely related to the high temperature strength, damage tolerance, structural stability and safety factor of the alloy. Therefore, the study of creep characteristics of superalloys can provide a basis for their life prediction and safe use. Mohammad et al. [3] studied the high-temperature creep behavior of the Inconel-713C nickel-based superalloy used in turbine blades under
850 � C and at different stress levels. Farangis et al. [4] studied the origin of creep dislocation in nickel-based single crystal superalloy under me dium stress and temperature. Wollgramm et al. [5] studied the creep behavior of nickel-based single-crystal superalloy, mainly analyzing the effects of stress and temperature on the minimum creep rate. Tian et al. [6,7] investigated the influence of element Re on the deformation mechanism of γ0 phase in nickel-based single-crystal superalloy during high temperature creep. Zhou et al. [8] studied the evolution rules of sub-structures by interrupting creep tests at different creep stages. Heather et al. [9] studied the effects of super solvus heat-treatments on the microstructures and properties of LSHR, Alloy10 and RR1000 nickel-based superalloys. The aim is to establish the relationship be tween microstructure and mechanical properties such as creep strength and creep crack propagation. Peng et al. [10] studied the creep prop erties of FGH96 nickel-based superalloy under different aging heat treatments at 700 � C and 690 MPa.
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (C. Ge), [email protected] (X. Li). https://doi.org/10.1016/j.msea.2020.139007 Received 22 October 2019; Received in revised form 21 January 2020; Accepted 23 January 2020 Available online 27 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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Xie et al. [11] studied dislocation morphology, creep behavior and dislocation networks of FGH95 Ni-based superalloy during creep pro cesses. Under the creep testing condition of 705 � C and 897 MPa, Jia et al. [12,13] demonstrated the effects of solution treatment and cooling rate on the microstructure, fracture and creep performance of FGH100L nickel-based superalloy. Currently, aero-engine turbine disks are mainly manufactured using Cast & Wrought (C&W) technology and Powder Metallurgy (PM) tech nology [14,15]. Spray Forming (SF) [16], which was developed in the early 1980s to compete with traditional superalloy manufacturing pro cesses, has significantly reduced the processing steps from atomization to final preforming compared with C&W or PM technologies. The SF process eliminates not only remelting and transformation steps in C&W processing, but also the steps of sieving, electrostatic precipitation, tank filling and degassing of the powder in PM processing. Therefore, the SF technology has become an efficient manufacturing alternative. The SF process is followed by hot working, i.e., Hot Isostatic Pressing (HIP) and Isothermal Forging (IF), to make the alloy denser. The heat treatment (HT) system of solid solution and two-stage aging is designed to improve the microstructure of the alloy, and the goal is to improve the overall properties of the alloy. The raw material used in this paper is the new third-generation PM superalloy FGH100L [12,13], which the authors’ team has been researching and developing for many years.
Fig. 1. Processing diagram of creep specimen.
variation of creep rate of the FGH100L alloy, respectively, under different stress loads and different temperatures. As shown in Fig. 2(a), the total creep life of the alloy decreases significantly with the increase of the applied stress load (from 690 MPa to 897 MPa) at 705 � C. When a low stress load of 690 MPa is applied, the steady creep duration of the alloy is the longest, about 300 h. The maximum total creep strain is 24.9%, with a low steady creep rate (0.0016 � 10-2h-1, see Fig. 3(a)) and a long creep life (1286.27 h). As the applied stress increases from 750 MPa to 897 MPa, the total creep strain of the alloy firstly decreases and then increases, and the creep life de creases significantly (see Fig. 2(a)). The creep life under 750 MPa is 920.36 h, and it is reduced to 81.55 h under 897 MPa. With the increase of stress load, the creep rate of the alloy increases, and the duration of the second creep stage, namely the steady creep stage, decreases (see Fig. 3(b), (c) and (d)). The steady-state creep rate of the alloy under a stress load of 897 MPa is the highest (0.1199 � 10-2h-1). Fig. 2(b) shows the creep curves of the FGH100L alloy under different stress loads at 750 � C. As shown in Fig. 2(b), the total creep life of the alloy decreases significantly with the increase of the applied stress load (from 450 MPa to 897 MPa). When a low stress load of 450 MPa is applied, the total creep strain of the alloy is relatively high (37.1%, see Fig. 2(b)), and the alloy has a long creep life (629.18 h). As for the steady creep stage, the alloy has a low creep rate (0.0010 � 10-2h-1, see Fig. 3 (e)) and the steady state creep duration is about 300 h. As the applied stress increases from 750 MPa to 897 MPa, the alloy creep rate rapidly increases and the steady state creep duration reduces (see Fig. 3(f), (g) and (h)). The total creep strain firstly increases and then decreases. The creep life decreases significantly from 62.16 h under 750 MPa to 11.49 h under 897 MPa. The highest creep rate is 0.2705 � 10-2h-1 under 897 MPa. As shown in Fig. 2(a) and (b), the creep properties of the FGH100L alloy are sensitive to the applied stress at a constant temperature. With the decrease of the applied stress, the second and third stages of creep of the alloy are prolonged, and thus the creep life of the alloy increases. On the contrary, the creep life of the alloy decreases with the increase of the applied stress. Fig. 2(c) shows the creep curves of the FGH100L alloy at different temperatures under a stress load of 897 MPa. Under the same stress load, the total creep life of the alloy decreases significantly with the increase of temperature. The creep curve under this condition still has three stages of creep, but the second stage lasts for a short time and then the alloy quickly enters the third stage. Under the condition of 650 � C/897 MPa, the total creep life of the alloy is 1477.42 h, the total strain is 5.2%, and the steady creep rate is 0.0004 � 10-2h-1 (see Fig. 3(i)). As the testing temperature increases from 705 � C, 725 � C–750 � C, the duration of the steady creep stage gradually reduces from 50 h, 12 h–2 h, and the creep rate of the second stage gradually increases from 0.1199 � 10-2h-1,
2. Materials and methods The FGH100L master alloy was prepared by a double smelting pro cess, i.e., vacuum induction smelting plus vacuum consumable remelt ing (VIM þ VAR). The main chemical composition of the alloy 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 alloy ingot (Φ200 mm � H300 mm) was prepared by SF, during which high purity N2 was used as an atomizing gas. Then it was treated by HIP without canning. The ingot was first heated to 1160 � C at a rate of 10 � C/min, holding for 3 h at 150 MPa, and then cooled to room tem perature at a cooling rate of 25 � C/min. Following this, an IF experiment was carried out, in which the forging temperature was 1150 � C, the pressing rate was 0.1 mm/s, and the engineering deformation was about 30.4%. After forging, the ingot was covered with asbestos until it was cooled to room temperature. Samples were taken from the IFed ingot for HT. The HT process was solid solution treatment plus two-stage aging: 1130 � C/1 h/fan quenching þ850 � C/4 h/AC þ 775 � C/8 h/AC. In this paper, the creep tests were carried out according to China national standard GB/T 2039–2012 uniaxial tensile creep test method for metallic materials [17]. The test equipment was the CSS type creep test machine of Changchun Test Machine Factory, and the specimen was non-standard creep specimen with a Ф5 mm round shape. The specific sample size is shown in Fig. 1. The creep temperature range was 650–750 � C, and the stress range was 450–897 MPa. Two samples were taken for creep testing under each condition. The grain structure was observed by a DC300 metallographic optical microscope (OM), and the metallographic etchant was 10 g CuCl2 þ 50 mL HCl þ 50 mL H2O. The fracture characteristics and microstructures were observed by Scanning Electron Microscope (SEM), JSM-6701F and ZEISS EVO®18. The dis tributions of grain orientation and misorientation were obtained by Electron Backscattering Diffraction (EBSD) technique. The electrolytic polishing solution was 20% H2SO4 þ 80% CH3OH and the electrolytic etching solution was 9 g CrO3 þ 90 mL H3PO4 þ 30 mL C2H5OH. 3. Results 3.1. Effects of stress and temperature on creep properties The general creep curve can be roughly divided into three stages: the first stage, steady creep in the second stage and accelerated creep in the third stage. Fig. 2 and Fig. 3 show the creep curves and the temporal 2
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Fig. 2. Creep curves of FGH100L alloy under different stress and temperature conditions: (a) 690–897MPa/705 � C, (b) 450–897MPa/750 � C, (c) 897MPa/ 650–750 � C.
0.1598 � 10-2h-1 to 0.2705 � 10-2h-1 (see Fig. 3(d), (j) and (h)). After entering the third creep stage, the creep rate increases with the increase of temperature, leading to a significant reduction of creep life. It can be seen from Fig. 2(c) that the creep strain of the third stage of the alloy changes greatly with the increase of temperature, which reflects that the influence of temperature on the microstructural changes, crack initia tion and growth, and fracture mode in this stage may be different. Overall, the creep life of the alloy will be reduced when operating at a higher temperature.
Fig. 4(a) and (b) reflect the liner relationships between lgε s and lgσ a at 705 � C and 750 � C, respectively. The slope value of the line (namely the stress index n) is 16.33 and 8.10, respectively, which is closely related to the creep mechanism. In general, the creep deformation of metal materials is very sensitive to the influence of stress factor when the stress index n is larger than 5, which is in accordance with the dislo cation mechanism during creep [21,22]. Therefore, the creep mecha nism of the FGH100L alloy under different stress loads at 705 � C and 750 � C is dislocation creep. Actually, the high n value in the alloy is related to the properties of the grain boundaries and the strengthening effect of γ0 phase which has a large volume fraction in the alloy. Fig. 5 reflects the relationship between lgε_ s and T-1 under a stress load of 897 MPa. According to the slope of the curve, the steady-state creep activation energy was derived as Qc ¼ 339.35 kJ/mol, which is much higher than the self-diffusion activation energy of Ni in austenite (265–285 kJ/mol). This can be attributed to the γ0 precipitates that take up about 52% mass fraction in the FGH100L alloy. The self-diffusion activation energy of polycrystalline nickel-based superalloys is gener ally ~300 kJ/mol [22,23]. The higher Qc value indicates that the creep mechanism of the FGH100L alloy is controlled by the atomic diffusion process. Dislocation climbing, directional vacancy diffusion and grain boundary sliding all involve atomic diffusion processes. The higher the Qc value, the more difficult for creep deformation to occur. Meanwhile, a higher Qc indicates higher creep resistance, higher creep strength and longer creep life of the alloy [18]. �
3.2. Prediction of creep life of FGH100L alloy Creep rate is one of the most important parameters to describe the creep behavior of an alloy. According to the previous analysis, the creep rate of the FGH100L alloy is significantly affected by load stress and service temperature. Generally, Norton’s Power Law is satisfied between the minimum creep rate, creep stress and temperature, as shown in Eq. (1). Therefore, the creep rate of the FGH100L alloy at the steady state stage can be described by Norton’s formula [18–20]: �
ε s ¼ Aσ na expð
Qc = RTÞ
(1)
In Eq. (1), ε_ s is the steady-state creep rate, A is the constant related to the structure of the material, T is the absolute temperature, R is the universal gas constant (8.314 J mol-1K-1), Qc is the apparent creep activation en ergy, n is the stress index, and σ a is the applied stress. At a given temperature, take the logarithm of both sides of Eq. (1) and convert it to Eq. (2): �
lgε s ¼ nlgσ a þ C
3.3. Creep fracture characteristics of FGH100L alloy Fig. 6 shows the creep fracture morphology of the FGH100L alloy under the condition of 705 � C/690 MPa. Fig. 6(a) shows the macroscopic morphology of creep fracture of the alloy. It can be seen that the neckshrinking phenomenon occurs. The fracture surface is flat. There is almost no shear lip in the final rupture region, and the propagation re gion accounts for a large area. The fracture source region demonstrates mixed intergranular and transgranular fracture characteristics, the former being more prominent. “Hole” type cracks can be seen, as indi cated by the arrow in the inset at the upper-right corner of Fig. 6(b). Fig. 6(c) and (d) manifest that there are a large number of steps in the
(2)
In Eq. (2), C is a constant, lgε_ s is linear with lgσ a and the slope of the straight line is n. At a given stress load, take the logarithm of both sides of Eq. (1) and convert it to Eq. (3): �
lgεs ¼
Qc =RT þ C’
(3)
In Eq. (3), C0 is a constant, and lgε_ s has a linear relationship with T-1. 3
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Fig. 3. Creep rate versus creep time for FGH100L alloy under different stress and temperature conditions: (a)–(d) 690–897MPa/705 � C, (e)–(h) 450–897MPa/750 � C, (i)(j) 897MPa/650–725 � C.
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Fig. 4. lgεs �
lgσ a curves of FGH100L alloy at different temperatures and stress loads: (a) 705 � C/690-897 MPa; (b) 750 � C/450-897 MPa.
propagation region and final rupture region, are obviously presented. Overall, the fracture surface is rough and uneven. Compared with Fig. 6 (a), the shear lip of the final rupture region is larger under 897 MPa in Fig. 7(a). The fracture source region of the alloy is characterized by mixed intergranular and transgranular fractures, the latter becoming more prominent. “Wedge” cracks can be seen, as indicated by the arrow in Fig. 7(b). Fig. 7(c) shows that in the propagation region, there are deep dimples, secondary cracks, and small facets where deep slip lines can be seen (see the enlarged area at the lower left corner in Fig. 7(c)). There are shallow dimples in the shear lip region (see Fig. 7(d)). Fig. 8 shows the creep fracture morphology of the FGH100L alloy under the condition of 650 � C/897 MPa. As can be seen in Fig. 8(a), the fracture shows severe neck shrinkage, large propagation region and almost no shear lip in final rupture region. Fig. 8(b) shows that the intergranular fracture is prominent in the fracture source region. There are secondary cracks, many steps, facets (see Fig. 8(c) and its enlarged region), deep dimples, and slip lines (see Fig. 8(d) and its enlarged re gion) in the propagation region. Fig. 9 shows the creep fracture morphology of the FGH100L alloy under the condition of 750 � C/897 MPa. Fig. 9(a) shows that the creep fracture of the alloy is uneven, with fracture source region, propagation region and large shear lip in the final rupture region. Intergranular fracture morphology and long “wedge” cracks are visible, and the transgranular fracture becomes prominent in the fracture source region
Fig. 5. Relationship between lgεs and T-1 of FGH100L alloy at 897 MPa. �
propagation region and slip lines and deep dimples on the facets. Fig. 7 shows the creep fracture morphology of the FGH100L alloy under the condition of 705 � C/897 MPa. As shown in Fig. 7(a)–(d), three characteristic regions of creep fracture, i.e., fracture source region,
Fig. 6. Creep fracture morphology of FGH100L alloy at 705 � C/690 MPa: (a) macroscopic morphology of fracture, (b) fracture source region, (c) steps and slip lines, (d) deep dimples. 5
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Fig. 7. Creep fracture morphology of FGH100L alloy at 705 � C/897 MPa: (a) macroscopic morphology of fracture; (b) fracture source region; (c) deep slip lines; (d) shallow dimples.
Fig. 8. Creep fracture morphology of FGH100L alloy at 650 � C/897 MPa: (a) macroscopic morphology of fracture, (b) fracture source region, (c) steps and facets, (d) deep dimples and slip lines.
(see Fig. 9(b)). In Fig. 9(c), there are a few steps with shallow slip lines and a few secondary cracks. Deep dimples can be seen in the shear lip region (see Fig. 9(d)). Fig. 10 shows the creep fracture morphology of the FGH100L alloy under the condition of 750 � C/450 MPa. It can be seen from Fig. 10(a) that the creep fracture morphology of the alloy is similar to that under 650 � C/897 MPa in Fig. 8(a). Fig. 10(b) shows that the fracture char acteristics of the source region are mixed intergranular and trans granular, the former being prominent. In Fig. 10(c), steps can be seen in the propagation region and the steps are covered with a large number of deep dimples and “hole-shaped” cracks can be seen. According to the analysis above, there are three kinds of conditions when the alloy exhibits good creep properties: (1) low stress and low temperature, (2) high stress and low temperature, and (3) low stress and
high temperature. In these cases, neck-shrinking phenomenon occurs on creep fracture, which shows the characteristics of mixed intergranular and transgranular fracture, the former being prominent. The fracture surface is relatively flat, there is almost no shear lip area, and the expansion region occupies a large area. In case (1) and case (2), there are many steps, planes and sliding lines on the surface. In case (2) and case (3), the “wedge” and “hole” cracks are easily formed. The formation of cracks can be divided into two processes: nucleation and growth. Both wedge-shaped and hole-shaped cracks are related to the slip behavior of grain boundary [24]. 4. Discussion Fig. 11 shows the local misorientation distribution of the FGH100L 6
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Fig. 9. Creep fracture morphology of FGH100L alloy at 750 � C/897 MPa: (a) macroscopic morphology of fracture, (b) fracture source region and cracks, (c) steps, (d) shear lip region.
Fig. 10. Creep fracture morphology of FGH100L alloy at 750 � C/450 MPa: (a) macroscopic morphology of fracture, (b) fracture source region and cracks, (c) steps and deep dimples.
alloy after creep fracture under different stress loads (450 MPa, 897 MPa) and temperatures (650–750 � C). As shown in Fig. 11(a), a large misorientation can be seen within the alloy after testing at 650 � C/897 MPa for 1477.4 h, indicating that there is a large strain inside the alloy. As shown in Fig. 11(b), after testing for 81.5 h under 705 � C/897 MPa, the grain misorientation of the alloy is smaller than that in Fig. 11(a), indicating that the internal strain of the alloy is relatively low. In this case, the strain and dislocation are mostly distributed around grain boundaries. Fig. 11(c) shows that after testing under 725 � C/897 MPa for 26.1 h, many small recrystallized grains are present in the alloy surrounding large original grains. The misorientation of the recrystal lized grains is smaller than that of the original grains, which means that the strain is relatively lower within the recrystallized grains. It can be inferred that the small recrystallized grains relieve strain concentration during creep by grain rotation, while the large original grains relieve strain concentration by grain distortion. As shown in Fig. 11(d), the microstructure of the alloy after testing for 11.5 h at 750 � C/897 MPa is
similar to that in Fig. 11(c), but the difference is that the overall misorientation of the alloy is lower than that in Fig. 11(c). It is shown in Fig. 11(e) that after testing at 750 � C/450 MPa for 629.2 h, the internal misorientation of the alloy is significant, indicating that there are large strains and many dislocations inside the grains. Fig. 11(f) shows that with the increase of temperature from 650 � C to 750 � C at 897 MPa, the grain misorientation in the alloy reduces, indicating that the higher the temperature, the lower the strain that the alloy can bear. This corre sponds with the fact that the higher the temperature, the lower the creep life. This suggests that the grain boundary becomes the weak location of the alloy with the increase of temperature. At 750 � C, the grain misorientation of the alloy decreases with the increase of stress from 450 MPa to 897 MPa, but the internal strain of the alloy is mainly distributed near the grain boundary, which is the weak location at this temperature. When the stress is relatively low, the dislocation multi plication rate in the alloy is slow, which provides enough time for the strain inside the alloy to release. Therefore, dislocations move into the 7
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Fig. 11. Local misorientation distribution of FGH100L alloy after creep fracture under different stress and temperature conditions (EBSD): (a) 650 � C/897 MPa, (b) 705 � C/897 MPa, (c) 725 � C/897 MPa, (d) 750 � C/897 MPa, (e) 750 � C/450 MPa, (f) comparison of local misorientation.
grain by sliding, thus enlarging the grain misorientation in the alloy and improving creep life. Fig. 12 shows the local orientation distribution of the FGH100L alloy after creep fracture under different stress loads (450 MPa, 897 MPa) and temperatures (650–750 � C). It can be seen that the grain orientation is random, and the color gradient change indicates that the misorientation exists in the alloy grains. There are many fine recrystallized grains around the large original grains and the orientation of the recrystallized grains is randomly distributed. In Fig. 12 (a) the color in some grains changes gradually from boundary to interior, indicating a large misorientation. In Fig. 12(b), (c) and (d), the color within grains is relatively uniform, indicating a small misorientation, while the color
near the boundary changes suddenly, exhibiting a large misorientation. It can be seen from Figs. 11 and 12 that, under the same stress load, the internal stress inside the alloy changes from grain interior to grain boundary after creep fracture as the temperature increases. Nickel-based superalloys have face-centered cubic (FCC) structure, in which slip is an important mode of deformation. The grain boundary slip (GBS) of fine grains that are obtained after sub-solid solution treatment is very important in the deformation of polycrystalline su peralloys during creep at high temperatures [25]. Combining the creep performance in Figs. 2 and 3 with the grain misorientation results in Figs. 11 and 12, the strain rate accelerates sharply in the acceleration phase under high stress load. When the strain rate is too fast for the grain 8
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Fig. 12. Local orientation distribution of FGH100L alloy after creep fracture under different stress and temperature conditions (EBSD): (a) 650 � C/897 MPa, (b) 705 � C/897 MPa, (c) 725 � C/897 MPa, (d) 750 � C/897 MPa, (e) 750 � C/450 MPa.
boundary to adjust, intragranular slip plays a significant role in creep deformation. In this case, the strain tends to concentrate in the slip zone. Since microcracks are easy to form at the location where the strain is concentrated, the secondary cracks nucleate along the slip zone under high stress. During creep, grain boundary deformation is initially acti vated, and when creep strain rate increases, grain matrix deformation coordinates with GBS [26,27].
(3) Under the stress load of 897 MPa, the grain misorientation of the alloy decreases with the increase of temperature from 650 � C to 750 � C. At 750 � C, the grain misorientation of the alloy decreases with the increase of stress load from 450 MPa to 897 MPa. This indicates that under the high temperature or high stress condi tions, the local strain the alloy can bear reduces greatly. With the increase of the temperature or the stress load, the grain misori entation around the original grain boundary becomes more sig nificant than within the original grain, indicating larger strain concentration near the original grain boundary. At the low tem perature or stress conditions, the difference in misorientation distribution between grain boundary and interior is not very large.
5. Conclusions (1) The creep properties of the FGH100L alloy, including creep life, creep strain and creep rate, are very sensitive to the applied creep stress and temperature. Based on Norton’s Power Law, the value of the stress index n of the FGH100L alloy at 705 � C and 750 � C is 16.33 and 8.10, respectively, indicating the dislocation creep mechanism. Under the stress load of 897 MPa, the activation energy Qc of the FGH100L alloy during steady-state creep is 339.35 kJ/mol, indicating the atomic-diffusion controlled creep process. (2) Mixed intergranular and transgranular fracture characteristics were observed under all the testing conditions. At low stress or low temperature conditions, the intergranular fracture is more prominent. With the increase of stress and temperature, the transgranular fracture becomes more prominent. This may be ascribed to the high strain rate under the high stress or high temperature conditions, in which the strain rate is too fast for the grain boundary to adjust, so that the intragranular slip becomes more important in creep deformation.
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Tian Tian: Data curation, Methodology, Investigation, Writing original draft. Zhibo Hao: Data curation, Investigation. Changchun Ge: Project administration, Supervision, Writing - review & editing. Xing gang Li: Data curation, Methodology, Supervision, Writing - review & editing. Shiqing Peng: Data curation. Chonglin Jia: Data curation.
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