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Effects of long term thermal exposure on microstructure and mechanical property evolution of a new wrought Ni–Fe based superalloy
Materials and Design 105 (2016) 66–74
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Effects of long term thermal exposure on microstructure and mechanical property evolution of a new wrought Ni–Fe based superalloy Hongfei Yin a,⁎, Yimin Gao a, Yuefeng Gu b a b
State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi'an Jiaotong University, Xi'an, Shaanxi Province 710049, PR China Xi'an Thermal Power Research Institute Co. Ltd., Xi'an, Shaanxi Province 710032, PR China
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Article history: Received 2 March 2016 Received in revised form 28 April 2016 Accepted 17 May 2016 Available online 18 May 2016 Keywords: High-temperature superalloy Mechanical property Thermal exposure Microstructure α-Cr phase
a b s t r a c t A new wrought Ni-Fe based superalloy was designed for advanced ultra-supercritical boiler tubes beyond 700 °C. In this study, the recrystallization treatment of the alloy following cold rolling and the microstructural evolution following 500 h, 1000 h and 3000 h at 750 °C were studied. The optimization of heat treatment was performed by thermodynamic calculation and the optimum combination was determined, i.e. annealing treatment at 1150 °C for 45 min, followed by a two-step aging treatment, at 810 °C for 1 h and at 770 °C for 16 h. Following this process, the M(C,N) precipitates and Cr-rich M23C6 carbides were formed at grain boundaries. Following 3000 h thermal exposure at 750 °C, the coarsening rate of the spherical γ precipitates was consistent with the LSW model and similar to that of the IN 740H alloy, while the variation in hardness values was relatively low, as the exposure time was increased. The α-Cr phase was observed after 1000 h at 750 °C and during the long term thermal exposure the coarsening of α-Cr phase occurred. The formation of α-Cr phase enhanced the intragranular yield strength and hardness values at room temperature and the stability of the mechanical properties was partially improved. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction The 700 °C advanced ultra-supercritical (A-USC) technology is an effective way to improve thermal efficiency and reduce CO2 emissions and energy consumption of coal fired power plants [1]. As the steam parameters of the 700 °C A-USC power plant have been established at 700 °C and 35 MPa, the temperature of the boiler tubes can reach 750 °C, or even higher [2,3]. Thus, the austenitic heat resistant stainless steels cannot meet these requirements and Ni-based superalloys should be employed in an A-USC boiler. The IN 740H [4–6] and CCA617 [7–9] alloys are already considered as potential candidates for use in the 700 °C A-USC power plant industry. In Japan, several types of Ni–Fe based alloys have been developed for the 700 °C A-USC power plant industry in order to reduce the cost, such as the HR6W alloy [10,11], but it seems that the rupture life of these alloys cannot meet the requirements needed. In China, a new Ni–Fe based alloy (designated as ST) , which contains 25–30 wt% Fe, without Co additions, has been developed for the 700 ° C A-USC power plant industry in order to reduce the cost, compared to IN 740H alloy, and the properties of this alloy are expected to be close to IN 740H alloy [12,13]. Microstructure stability of these alloys is essential, in order to maintain their mechanical properties for a long ⁎ Corresponding author. E-mail address: [email protected] (H. Yin).
http://dx.doi.org/10.1016/j.matdes.2016.05.059 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
service time, as the components used in an A-USC power plant are employed at high temperatures [14,15]. Therefore, the optimization of the heat treatment process, following cold rolling, has been performed, and the microstructure, microindention, hardness and strength evolution of a new Ni–Fe superalloy were studied, following long term thermal exposure at 750 °C. 2. Experimental procedure The new Ni–Fe based alloy (30 kg), namely ST-alloy, was prepared in a vacuum induction melting furnace. The chemical composition of the experimental alloy is presented in Table 1; the IN 740H alloy was also added for comparison. The cast ingot was homogenized, hot forged and rolled to a thickness of 20 mm at 1150 °C. The hot then rolled plate was cold rolled to a thickness of 10 mm, at room temperature. Samples were cut from the cold rolled plate, and heat treatments for recrystallization experiments at 850 °C, 900 °C, 950 °C, 1000 °C and 1050 ° C were carried out to explore the recrystallization temperature, which is an important factor of hot rolling process. Also, the recrystallized samples were solution treated at 1150 °C for 45 min, followed by air cooling (AC), and then the two-step aging treatment was performed, at 810 °C for 1 h and at 770 °C for 16 h, with subsequent AC. The two step aging treatment was selected according to the thermodynamic simulation of the equilibrium phases using JMatPro software (version 8.0). Finally, following the optimized heat treatment process the samples were
Bal. Bal. 24–32 1.0 0.01–0.05 – – 20 0.003–0.005 0.001
21–27 25
N Co B
Cr
Fe
Ni
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subjected to long term thermal exposure experiments at 750 °C for 500 h, 1000 h and 3000 h. All the specimens were ground and polished following the standard metallographic methods and then chemically etched with the Marbel solution (10 g of CuSO4 + 50 mL of hydrochloric acid + 50 mL of H2O). The various microstructures were assessed via optical microscopy (Leica DVM2500), scanning electron microscopy (Hitachi S4800), X-ray Diffractometer (SHIMADZU 7000) and transmission electron microscopy (JEM-3010). The samples for transmission electron microscopy (TEM) examination were electro-polished using a double-jet apparatus. Sheets of 0.3 mm in thickness were sectioned from the bulk material and mechanically ground down to 40–50 μm in thickness. Discs of 3 mm in diameter were then punched from the thinned sheets and then electro-polished at −30 °C at 16 V, using a solution of 90% ethanol and 10% perchloric acid [6]. Following the microstructural characterization, micro-hardness tests were carried out for each experimental condition, using a microindention hardness tester (MHVD-1000IS), at a load of 200 g. Tensile tests were performed in a universal tensile testing machine (MTS E45) both at room temperature and 750 °C, at the strain rate of 2.5 × 10−4 s−1. There were 3 tensile specimens at each test temperature, and the results were the average values. The fracture surfaces were examined via scanning electron microscopy (SEM) immediately after tensile testing.
0.02–0.05 0.03 0.8–1.6 1.4 0.8–1.6 1.4 0.2–0.6 0.5 0.1–0.5 0.6 0.2–0.7 0.26 ST IN740H
0.09–0.16 0.20
0.3–0.7 1.0
Al Nb Si Mn Alloy
Table 1 Chemical compositions of the ST-alloy and IN 740H alloy (wt.%).
W
Mo
Ti
C
3. Results and discussion 3.1. Thermodynamic simulations Thermodynamic calculations were carried out to predict the precipitation temperatures of the various phases formed in the new alloy. Fig. 1 shows the equilibrium phase diagram based on the chemical composition of the alloy. It seems that the main phases are similar to that of the precipitation-strengthened Ni-base superalloys [16]. The main strengthening phase is γ , which precipitates at 935 °C. The M(C,N) phase started to precipitate since the initiation of the solidification process and the M23C6 carbide was formed at 855 °C. There are only a few of MB2 phases forming as the temperature decreases. The α-Cr phase, which is usually considered detrimental in mechanical properties, is also predicted to appear at 750 °C [17]. Up to date, the effects of α-Cr phase on Ni–Fe based alloys are still not fully understood. As a conclusion, based on the above simulation results, the solution and aging heat treatment for the ST alloy were determined. 3.2. Hardness The measured hardness values of the ST alloy at various recrystallization annealing temperatures and following long term thermal exposure at 750 °C is presented in Fig. 2. Following cold rolling, the hardness value of the alloy was 410 ± 8 HV. The hardness values were significantly decreased following annealing heat treatment, as seen in Fig. 2(a). The lowest hardness value, i.e. 196 ± 8 HV, was recorded at 1100 °C. Following solution and aging treatment, the hardness was approximately 310 ± 6 HV. Long term thermal exposure up to 3000 h at 750 °C resulted in a slight increase of hardness values (see Fig. 2(b)). 3.3. Microstructural evolution Fig. 3 shows the microstructures observed in the ST alloy following recrystallization annealing experiments at 950 °C, 1000 °C and 1050 ° C, obtained by optical microscopy. The average grain size of the ST alloy after cold rolling was 18 μm, and there was no obvious change of the average grain size after recrystallization annealing treatment at 950 °C. However, the microstructure produced during cold rolling was altered and the average grain size was approximately 70 μm, following recrystallization annealing treatment at 1000 °C. It is also evident that
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Fig. 1. Thermodynamic phase diagram of the ST alloy: (a) overall diagram; (b) partial diagram.
Fig. 2. (a) Hardness values of the ST alloy after recrystallization annealing treatments at various temperatures; (b) Hardness values of the ST alloy after long term thermal exposure at 750 °C.
Fig. 3. Optical micrographs of the ST alloy following: (a) cold rolled condition; recrystallization annealing treatments at (b) 950 °C; (c) at 1000 °C; (d) at 1050 °C.
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grain growth took place following recrystallization annealing treatment at 1050 °C, and the average grain size was about 100 μm. As presented in Fig. 4, following the solution treatment and the two steps of aging treatment (1145 °C/45 min/AC + 810 °C/1 h/ AC + 770 °C/16 h/AC), the average grain size was approximately 130 μm, with Cr-rich M23C6 type carbides distributed at grain boundaries and γ′ precipitates homogeneously dispersed within the γ matrix. Fig. 4(d) is a bright-field TEM image of a M23C6 type carbide at grain boundaries, having the face-centered cubic structure. Fig. 4(e) shows a selected area diffraction pattern direction, which can be (SAD) of the M23C6 carbide along the [012] indexed educing lattice parameter values of a = b = c = 1.08 nm. The energy dispersive spectrum (EDS) from the M23C6 is presented in Fig. 4(f). The γ precipitates had an average size of approximately 37 nm. Following the thermal exposure at 750 °C, the grain size and morphology of carbides had no significant variation and the main differences in the observed microstructures were the coarsening of γ precipitates and the precipitation of α-Cr phase. As presented in Fig. 5, the determined average diameters of γ' particles following 500, 1000 and 3000 h were 63 nm, 77 nm and 101 nm, respectively. The needle-like α-Cr phase, with a body-centered cubic structure and 3–15 μm long, was formed in the specimen following thermal exposure for 1000 h. When the heat treatment time was further increased, the α-Cr phase particles became coarser without forming the previously
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observed network structure. The average size of the α-Cr phase particles was increased to 5–30 μm long, following thermal exposure for 3000 h, as presented in Fig. 6(a) and (b). Fig. 6(c) is a bright-field TEM image of a α-Cr phase particle within a grain and the selected area dif fraction (SAD) along the [111] direction is presented in Fig. 6(d), which can be indexed educing lattice parameter values of a = b = c = 0.298 nm. The energy dispersive spectrum (EDS) and X-ray Diffraction (XRD) from the α-Cr phase is presented in Fig. 6(e) and (f). 3.4. Tensile strength at room temperature and 750 °C The results of tensile testing at room temperature and 750 °C of the ST alloy following thermal exposure at 750 °C for 0 (as-aged), 1000 and 3000 h are presented in Fig. 7. The average yield strength of the ST alloy at room temperature was decreased after the first 1000 h. With further increase of the thermal exposure time the yield strength was gradually increased, however, the average ultimate tensile strength at room temperature did not significantly decrease during the thermal exposure at 750 °C, remaining above 1080 MPa. However, both the yield strength and the ultimate tensile strength at 750 °C were gradually decreased as the thermal exposure time was increased. Additional SEM micrograph results are presented in Figs. 8 and 9. It was observed that the finer network of precipitations contained dimples and macro-voids,
Fig. 4. SEM and TEM micrographs with EDS analyses of the ST alloy following the optimized heat treatment: (a) low magnification SEM image; (b) TiN and M23C6 carbides distribution; (c) (f) EDS spectrum of M23C6 carbides. γ precipitates distribution; (d) M23C6 carbides precipitated at the grain boundaries; (e) SAD of M23C6 carbides along [012];
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Fig. 5. The SEM micrographs of the as-aged ST alloy after thermal exposure at 750 °C: (a) 500 h; (b) 1000 h; (c) 3000 h.
following fracture of the ST alloy at room temperature, and the fracture mode was ductile fracture with dimples. As the thermal exposure time was increased, inter-granular fracture occurred (shown in Fig. 8(b) and (c)). The fracture mechanism was a combination of ductile fracture and inter-granular fracture, following tensile testing at room temperature. However, inter-granular fracture could be considered as the primary fracture mechanism at 750 °C.
3.5. Discussion The high temperature recovery of the ST alloy following cold rolling started approximately between 900–950 °C, as the hardness of the matrix was significantly decreased to 280 HV, with insignificant microstructural changes. The new equiaxed grains had already been formed at 1000 °C. Further growth of the new equiaxed grain had already occurred with further reduction of hardness at 1050 °C. Therefore, the initial recrystallization temperature of the ST alloy is between 950 °C and 1000 °C. The microstructure of the ST alloy samples, following the heat-treatment, is presented in Fig. 3. Irregular-shaped TiN and (Nb,Ti)C phases formed within the grains and at the grain boundaries following solution heat treatment, as revealed by the energy dispersive spectroscopy (EDS) analysis, and the Cr-rich M23C6 carbides were discretely distributed at the grain boundaries due to the aging treatment at 810 °C for 1 h. The γ phase was completely precipitated within the γ matrix due to the aging treatment at 770 °C for 16 h. These results were consistent with the simulation of equilibrium phase performed by JMatPro 8.0. Following thermal exposure at 750 °C for 3000 h, both the grain size and the morphology of M23C6 carbides had no significant change; however, the coarsening of γ precipitates from 37 nm to 101 nm occurred in the absence of applied stress. The morphology of γ precipitates was kept spherical during the thermal exposure. On the contrary, for the IN 740 alloy, the morphology of γ precipitates was changed from spherical to cubic shape, during the long term thermal exposure. It is well known that γ' particles with high surface energy and low coherent strains are normally associated with spherical shape [18,19]. Therefore,
the size of γ' particles of the ST alloy were increased in order to minimize the surface area, so that surface energy could be reduced [20]. The trend of size distribution of γ' particles between the ST alloy and IN 740H is presented in Fig. 10(a) and the initial size of γ' phase during thermal exposure for 0 h was defined as do. As presented in Fig. 10(b), the coarsening rate of γ' particles match well with that of the classic diffusion controlled growth model (LSW), which was similar to the IN 740H alloy. The majority of previous studies support that the coarsening of γ' precipitates can be described according to the LSW model and the power law expression is as follows [21]:
3
d
3 −do
¼
! 64σDC o V 2m ∙t 9RT
ð1Þ
where d is the average diameter of γ' particles, σ is the surface energy of γ'/γ, D is the diffusion coefficient, Co is the solubility of γ'-forming elements in γ matrix, Vm is the molar volume of the particles, R is the gas constant and T is the absolute temperature during thermal exposure. Although it is found that α-Cr formation is associated with δ (Ni3Nb) phase formation due to Cr enrichment, which is caused by Ni, Al or Nb depletion from the matrix [22,23], α-Cr phase was formed in the ST alloy without the δ (Ni3Nb) phase following thermal exposure at 750 ° C. This comes in agreement with the thermodynamic calculations of equilibrium phases (Fig. 2). The microhardness values in the grains mainly depend on both the intragranular precipitates and the γ matrix. Previous studies have also found that the microhardness values of most precipitation strengthening superalloys increase rapidly in the beginning and then continuously decrease, with the increasing of thermal exposure time, as the γ' particles coarse above a critical size [24]. However, the α-Cr phase, as a brittle phase, was observed within the grains of the ST alloy during thermal exposure. Previous studies have demonstrated the effect of the distributed α-Cr phase (Fig. 6) on intragranular strengthening [25] and that its hardness value may reach up to 3 times higher than that of the γ matrix [26]. In the present study, the reduction of hardness values, which resulted from coarsening of γ'
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Fig. 6. SEM and TEM micrographs with EDS and XRD data of the ST alloy following thermal exposure at 750 °C for 1000 h: (a) and 3000 h (b); (c) α-Cr phase precipitated in the grain; (d) SAD of α-Cr phase along [111]; (e) EDS spectrum of α-Cr phase; (f) XRD spectrum of α-Cr phase in the ST alloy.
particles, might be compensated by the precipitation of α-Cr phase, so the hardness remained stable during thermal exposure up to 3000 h. According to the morphology of the fracture surfaces of the tensile testing samples, intergranular fracture occurred following tensile testing at 750 °C. Also, the grain boundary failure occurred as the thermal exposure time was increased, due to the transition from the transgranular to the intergranular fracture mode, following tensile testing at room temperature. The average yield strength of the ST alloy at room temperature was decreased after the first 1000 h, due to the γ' particles coarsening, following thermal exposure at 750 °C. The α-Cr phase particles, which are incoherent with the γ matrix, became coarser as the thermal exposure time was increased from 1000 h up to 3000 h and the α-Cr precipitates might have inhibited the dislocation motions within the grains. When dislocations glide in the primary slip plane (i.e., (111) plane) and approach α-Cr precipitates, α-Cr precipitates impeded the gliding of dislocations. It is definitely difficult for the gliding dislocation to cut through the α-Cr precipitate. Firstly, dislocations may partly bypass α-Cr precipitates via Orowan mechanism, and then dislocation loops accumulate at the interface between the matrix and α-Cr precipitates. Because of the dislocation loops pile-up around α-Cr precipitates, higher critical shear stress is necessary for dislocations to bypass α-Cr precipitates [27]. Secondly, much higher driving force is required for dislocations to cross slip to other slip systems, especially to non-primary slip systems, in order to avoid the obstacle of the α-Cr precipitate.
Therefore, the formation of α-Cr phase in the ST alloy made plastic deformation more difficult [13]. The average yield strength of the ST alloy at room temperature was increased due to the substantial precipitation strengthening effect of the α-Cr phase within the grains, though the yield strength was decreased with the coarsening of the γ' particles. The variation of the average yield strength at room temperature was also consistent with the variation of hardness along the thermal exposure. The dispersed α-Cr phase has partially contributed to the stability of the mechanical properties. If the α-Cr precipitates had formed a network of precipitations within the grains or films at the grain boundaries, the alloy would be significantly more brittle, causing deterioration of the mechanical properties. Thus, the effects of the α-Cr phase on the comprehensive mechanical properties of the ST alloy related to the size and distribution of α-Cr precipitates. However, there was insignificant contribution of the α-Cr precipitates to the strength at high temperature, as the fracture mechanism had changed at 750 °C.
4. Conclusions The variation of microstructures and mechanical properties of a new wrought Ni–Fe based superalloy, designed for advanced ultra-supercritical boiler tubes, following optimized heat treatment and long term thermal exposure at 750 °C was studied.
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Fig. 7. Comparison of tensile properties of the ST alloy at room temperature and 750 °C thermal exposure for 3000 h: (a) yield strength; (b) ultimate tensile strength; (c) elongation; (d) area reduction.
(1) The initial recrystallization temperature of this alloy is between 950 °C and 1000 °C. Following the heat treatment optimization the hardness values reached 310 HV. The average grain size of the ST alloy was 130 μm with Cr-rich M23C6 carbides distributed at grain boundaries, while the average size of γ' precipitates was 37 nm.
(2) Following thermal exposure at 750 °C, both the grain size and morphology of M23C6 carbides had no significant change; the coarsening of γ' precipitates from 37 nm to 101 nm and the precipitation of α-Cr occurred. The coarsening rate of γ' particles matches the classic LSW model and the γ' particles kept their spherical morphology. The α-Cr phase was observed following thermal
Fig. 8. SEM fractographs of the tensile tested samples at room temperature: (a) thermal exposure for 0 h; (b) thermal exposure for 1000 h; (c) thermal exposure for 3000 h.
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Fig. 9. SEM fractographs of the tensile tested samples at 750 °C: (a) thermal exposure for 0 h; (b) thermal exposure for 1000 h; (c) thermal exposure for 3000 h.
Fig. 10. (a) Plot of γ' precipitate size during thermal exposure at 750 °C; (b) LSW model.
exposure for 1000 h and the size of needle-like α-Cr phase was significantly increased as thermal exposure time increased. (3) The hardness value of this alloy maintained at approximately 310 HV during thermal exposure up to 3000 h due to the precipitation of α-Cr phase in the grains, though coarsening of γ' particles may result in the reduction of hardness values. The new alloy demonstrated high strength. The grain boundaries of this new alloy were the weaker locations at 750 °C and the grain boundary weakening occurred during thermal exposure. The dispersed α-Cr phase enhanced intragranular strength, implying that the reduction of yield strength at room temperature, which resulted from the coarsening of γ' particles, might have been compensated by the precipitation of α-Cr phase. The dispersed α-Cr particles within the grain could partially improve the stability of mechanical properties during thermal exposure up to 3000 h.
These properties combined with low-cost make this new alloy appropriate for potential use in A-USC power plants.
Acknowledgments This work has been financially supported by China Huaneng Group (Grant No. ZA-12-HKR04), Huaneng Power International Inc. (Grant No. ZA-14-HKR01) and the National Nature Science Foundation of China under Grant No. 51401164.
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Effects of long term thermal exposure on microstructure and mechanical property evolution of a new wrought Ni–Fe based superalloy Hongfei Yin a,⁎, Yimin Gao a, Yuefeng Gu b a b
State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi'an Jiaotong University, Xi'an, Shaanxi Province 710049, PR China Xi'an Thermal Power Research Institute Co. Ltd., Xi'an, Shaanxi Province 710032, PR China
a r t i c l e
i n f o
Article history: Received 2 March 2016 Received in revised form 28 April 2016 Accepted 17 May 2016 Available online 18 May 2016 Keywords: High-temperature superalloy Mechanical property Thermal exposure Microstructure α-Cr phase
a b s t r a c t A new wrought Ni-Fe based superalloy was designed for advanced ultra-supercritical boiler tubes beyond 700 °C. In this study, the recrystallization treatment of the alloy following cold rolling and the microstructural evolution following 500 h, 1000 h and 3000 h at 750 °C were studied. The optimization of heat treatment was performed by thermodynamic calculation and the optimum combination was determined, i.e. annealing treatment at 1150 °C for 45 min, followed by a two-step aging treatment, at 810 °C for 1 h and at 770 °C for 16 h. Following this process, the M(C,N) precipitates and Cr-rich M23C6 carbides were formed at grain boundaries. Following 3000 h thermal exposure at 750 °C, the coarsening rate of the spherical γ precipitates was consistent with the LSW model and similar to that of the IN 740H alloy, while the variation in hardness values was relatively low, as the exposure time was increased. The α-Cr phase was observed after 1000 h at 750 °C and during the long term thermal exposure the coarsening of α-Cr phase occurred. The formation of α-Cr phase enhanced the intragranular yield strength and hardness values at room temperature and the stability of the mechanical properties was partially improved. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction The 700 °C advanced ultra-supercritical (A-USC) technology is an effective way to improve thermal efficiency and reduce CO2 emissions and energy consumption of coal fired power plants [1]. As the steam parameters of the 700 °C A-USC power plant have been established at 700 °C and 35 MPa, the temperature of the boiler tubes can reach 750 °C, or even higher [2,3]. Thus, the austenitic heat resistant stainless steels cannot meet these requirements and Ni-based superalloys should be employed in an A-USC boiler. The IN 740H [4–6] and CCA617 [7–9] alloys are already considered as potential candidates for use in the 700 °C A-USC power plant industry. In Japan, several types of Ni–Fe based alloys have been developed for the 700 °C A-USC power plant industry in order to reduce the cost, such as the HR6W alloy [10,11], but it seems that the rupture life of these alloys cannot meet the requirements needed. In China, a new Ni–Fe based alloy (designated as ST) , which contains 25–30 wt% Fe, without Co additions, has been developed for the 700 ° C A-USC power plant industry in order to reduce the cost, compared to IN 740H alloy, and the properties of this alloy are expected to be close to IN 740H alloy [12,13]. Microstructure stability of these alloys is essential, in order to maintain their mechanical properties for a long ⁎ Corresponding author. E-mail address: [email protected] (H. Yin).
http://dx.doi.org/10.1016/j.matdes.2016.05.059 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
service time, as the components used in an A-USC power plant are employed at high temperatures [14,15]. Therefore, the optimization of the heat treatment process, following cold rolling, has been performed, and the microstructure, microindention, hardness and strength evolution of a new Ni–Fe superalloy were studied, following long term thermal exposure at 750 °C. 2. Experimental procedure The new Ni–Fe based alloy (30 kg), namely ST-alloy, was prepared in a vacuum induction melting furnace. The chemical composition of the experimental alloy is presented in Table 1; the IN 740H alloy was also added for comparison. The cast ingot was homogenized, hot forged and rolled to a thickness of 20 mm at 1150 °C. The hot then rolled plate was cold rolled to a thickness of 10 mm, at room temperature. Samples were cut from the cold rolled plate, and heat treatments for recrystallization experiments at 850 °C, 900 °C, 950 °C, 1000 °C and 1050 ° C were carried out to explore the recrystallization temperature, which is an important factor of hot rolling process. Also, the recrystallized samples were solution treated at 1150 °C for 45 min, followed by air cooling (AC), and then the two-step aging treatment was performed, at 810 °C for 1 h and at 770 °C for 16 h, with subsequent AC. The two step aging treatment was selected according to the thermodynamic simulation of the equilibrium phases using JMatPro software (version 8.0). Finally, following the optimized heat treatment process the samples were
Bal. Bal. 24–32 1.0 0.01–0.05 – – 20 0.003–0.005 0.001
21–27 25
N Co B
Cr
Fe
Ni
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subjected to long term thermal exposure experiments at 750 °C for 500 h, 1000 h and 3000 h. All the specimens were ground and polished following the standard metallographic methods and then chemically etched with the Marbel solution (10 g of CuSO4 + 50 mL of hydrochloric acid + 50 mL of H2O). The various microstructures were assessed via optical microscopy (Leica DVM2500), scanning electron microscopy (Hitachi S4800), X-ray Diffractometer (SHIMADZU 7000) and transmission electron microscopy (JEM-3010). The samples for transmission electron microscopy (TEM) examination were electro-polished using a double-jet apparatus. Sheets of 0.3 mm in thickness were sectioned from the bulk material and mechanically ground down to 40–50 μm in thickness. Discs of 3 mm in diameter were then punched from the thinned sheets and then electro-polished at −30 °C at 16 V, using a solution of 90% ethanol and 10% perchloric acid [6]. Following the microstructural characterization, micro-hardness tests were carried out for each experimental condition, using a microindention hardness tester (MHVD-1000IS), at a load of 200 g. Tensile tests were performed in a universal tensile testing machine (MTS E45) both at room temperature and 750 °C, at the strain rate of 2.5 × 10−4 s−1. There were 3 tensile specimens at each test temperature, and the results were the average values. The fracture surfaces were examined via scanning electron microscopy (SEM) immediately after tensile testing.
0.02–0.05 0.03 0.8–1.6 1.4 0.8–1.6 1.4 0.2–0.6 0.5 0.1–0.5 0.6 0.2–0.7 0.26 ST IN740H
0.09–0.16 0.20
0.3–0.7 1.0
Al Nb Si Mn Alloy
Table 1 Chemical compositions of the ST-alloy and IN 740H alloy (wt.%).
W
Mo
Ti
C
3. Results and discussion 3.1. Thermodynamic simulations Thermodynamic calculations were carried out to predict the precipitation temperatures of the various phases formed in the new alloy. Fig. 1 shows the equilibrium phase diagram based on the chemical composition of the alloy. It seems that the main phases are similar to that of the precipitation-strengthened Ni-base superalloys [16]. The main strengthening phase is γ , which precipitates at 935 °C. The M(C,N) phase started to precipitate since the initiation of the solidification process and the M23C6 carbide was formed at 855 °C. There are only a few of MB2 phases forming as the temperature decreases. The α-Cr phase, which is usually considered detrimental in mechanical properties, is also predicted to appear at 750 °C [17]. Up to date, the effects of α-Cr phase on Ni–Fe based alloys are still not fully understood. As a conclusion, based on the above simulation results, the solution and aging heat treatment for the ST alloy were determined. 3.2. Hardness The measured hardness values of the ST alloy at various recrystallization annealing temperatures and following long term thermal exposure at 750 °C is presented in Fig. 2. Following cold rolling, the hardness value of the alloy was 410 ± 8 HV. The hardness values were significantly decreased following annealing heat treatment, as seen in Fig. 2(a). The lowest hardness value, i.e. 196 ± 8 HV, was recorded at 1100 °C. Following solution and aging treatment, the hardness was approximately 310 ± 6 HV. Long term thermal exposure up to 3000 h at 750 °C resulted in a slight increase of hardness values (see Fig. 2(b)). 3.3. Microstructural evolution Fig. 3 shows the microstructures observed in the ST alloy following recrystallization annealing experiments at 950 °C, 1000 °C and 1050 ° C, obtained by optical microscopy. The average grain size of the ST alloy after cold rolling was 18 μm, and there was no obvious change of the average grain size after recrystallization annealing treatment at 950 °C. However, the microstructure produced during cold rolling was altered and the average grain size was approximately 70 μm, following recrystallization annealing treatment at 1000 °C. It is also evident that
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Fig. 1. Thermodynamic phase diagram of the ST alloy: (a) overall diagram; (b) partial diagram.
Fig. 2. (a) Hardness values of the ST alloy after recrystallization annealing treatments at various temperatures; (b) Hardness values of the ST alloy after long term thermal exposure at 750 °C.
Fig. 3. Optical micrographs of the ST alloy following: (a) cold rolled condition; recrystallization annealing treatments at (b) 950 °C; (c) at 1000 °C; (d) at 1050 °C.
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grain growth took place following recrystallization annealing treatment at 1050 °C, and the average grain size was about 100 μm. As presented in Fig. 4, following the solution treatment and the two steps of aging treatment (1145 °C/45 min/AC + 810 °C/1 h/ AC + 770 °C/16 h/AC), the average grain size was approximately 130 μm, with Cr-rich M23C6 type carbides distributed at grain boundaries and γ′ precipitates homogeneously dispersed within the γ matrix. Fig. 4(d) is a bright-field TEM image of a M23C6 type carbide at grain boundaries, having the face-centered cubic structure. Fig. 4(e) shows a selected area diffraction pattern direction, which can be (SAD) of the M23C6 carbide along the [012] indexed educing lattice parameter values of a = b = c = 1.08 nm. The energy dispersive spectrum (EDS) from the M23C6 is presented in Fig. 4(f). The γ precipitates had an average size of approximately 37 nm. Following the thermal exposure at 750 °C, the grain size and morphology of carbides had no significant variation and the main differences in the observed microstructures were the coarsening of γ precipitates and the precipitation of α-Cr phase. As presented in Fig. 5, the determined average diameters of γ' particles following 500, 1000 and 3000 h were 63 nm, 77 nm and 101 nm, respectively. The needle-like α-Cr phase, with a body-centered cubic structure and 3–15 μm long, was formed in the specimen following thermal exposure for 1000 h. When the heat treatment time was further increased, the α-Cr phase particles became coarser without forming the previously
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observed network structure. The average size of the α-Cr phase particles was increased to 5–30 μm long, following thermal exposure for 3000 h, as presented in Fig. 6(a) and (b). Fig. 6(c) is a bright-field TEM image of a α-Cr phase particle within a grain and the selected area dif fraction (SAD) along the [111] direction is presented in Fig. 6(d), which can be indexed educing lattice parameter values of a = b = c = 0.298 nm. The energy dispersive spectrum (EDS) and X-ray Diffraction (XRD) from the α-Cr phase is presented in Fig. 6(e) and (f). 3.4. Tensile strength at room temperature and 750 °C The results of tensile testing at room temperature and 750 °C of the ST alloy following thermal exposure at 750 °C for 0 (as-aged), 1000 and 3000 h are presented in Fig. 7. The average yield strength of the ST alloy at room temperature was decreased after the first 1000 h. With further increase of the thermal exposure time the yield strength was gradually increased, however, the average ultimate tensile strength at room temperature did not significantly decrease during the thermal exposure at 750 °C, remaining above 1080 MPa. However, both the yield strength and the ultimate tensile strength at 750 °C were gradually decreased as the thermal exposure time was increased. Additional SEM micrograph results are presented in Figs. 8 and 9. It was observed that the finer network of precipitations contained dimples and macro-voids,
Fig. 4. SEM and TEM micrographs with EDS analyses of the ST alloy following the optimized heat treatment: (a) low magnification SEM image; (b) TiN and M23C6 carbides distribution; (c) (f) EDS spectrum of M23C6 carbides. γ precipitates distribution; (d) M23C6 carbides precipitated at the grain boundaries; (e) SAD of M23C6 carbides along [012];
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Fig. 5. The SEM micrographs of the as-aged ST alloy after thermal exposure at 750 °C: (a) 500 h; (b) 1000 h; (c) 3000 h.
following fracture of the ST alloy at room temperature, and the fracture mode was ductile fracture with dimples. As the thermal exposure time was increased, inter-granular fracture occurred (shown in Fig. 8(b) and (c)). The fracture mechanism was a combination of ductile fracture and inter-granular fracture, following tensile testing at room temperature. However, inter-granular fracture could be considered as the primary fracture mechanism at 750 °C.
3.5. Discussion The high temperature recovery of the ST alloy following cold rolling started approximately between 900–950 °C, as the hardness of the matrix was significantly decreased to 280 HV, with insignificant microstructural changes. The new equiaxed grains had already been formed at 1000 °C. Further growth of the new equiaxed grain had already occurred with further reduction of hardness at 1050 °C. Therefore, the initial recrystallization temperature of the ST alloy is between 950 °C and 1000 °C. The microstructure of the ST alloy samples, following the heat-treatment, is presented in Fig. 3. Irregular-shaped TiN and (Nb,Ti)C phases formed within the grains and at the grain boundaries following solution heat treatment, as revealed by the energy dispersive spectroscopy (EDS) analysis, and the Cr-rich M23C6 carbides were discretely distributed at the grain boundaries due to the aging treatment at 810 °C for 1 h. The γ phase was completely precipitated within the γ matrix due to the aging treatment at 770 °C for 16 h. These results were consistent with the simulation of equilibrium phase performed by JMatPro 8.0. Following thermal exposure at 750 °C for 3000 h, both the grain size and the morphology of M23C6 carbides had no significant change; however, the coarsening of γ precipitates from 37 nm to 101 nm occurred in the absence of applied stress. The morphology of γ precipitates was kept spherical during the thermal exposure. On the contrary, for the IN 740 alloy, the morphology of γ precipitates was changed from spherical to cubic shape, during the long term thermal exposure. It is well known that γ' particles with high surface energy and low coherent strains are normally associated with spherical shape [18,19]. Therefore,
the size of γ' particles of the ST alloy were increased in order to minimize the surface area, so that surface energy could be reduced [20]. The trend of size distribution of γ' particles between the ST alloy and IN 740H is presented in Fig. 10(a) and the initial size of γ' phase during thermal exposure for 0 h was defined as do. As presented in Fig. 10(b), the coarsening rate of γ' particles match well with that of the classic diffusion controlled growth model (LSW), which was similar to the IN 740H alloy. The majority of previous studies support that the coarsening of γ' precipitates can be described according to the LSW model and the power law expression is as follows [21]:
3
d
3 −do
¼
! 64σDC o V 2m ∙t 9RT
ð1Þ
where d is the average diameter of γ' particles, σ is the surface energy of γ'/γ, D is the diffusion coefficient, Co is the solubility of γ'-forming elements in γ matrix, Vm is the molar volume of the particles, R is the gas constant and T is the absolute temperature during thermal exposure. Although it is found that α-Cr formation is associated with δ (Ni3Nb) phase formation due to Cr enrichment, which is caused by Ni, Al or Nb depletion from the matrix [22,23], α-Cr phase was formed in the ST alloy without the δ (Ni3Nb) phase following thermal exposure at 750 ° C. This comes in agreement with the thermodynamic calculations of equilibrium phases (Fig. 2). The microhardness values in the grains mainly depend on both the intragranular precipitates and the γ matrix. Previous studies have also found that the microhardness values of most precipitation strengthening superalloys increase rapidly in the beginning and then continuously decrease, with the increasing of thermal exposure time, as the γ' particles coarse above a critical size [24]. However, the α-Cr phase, as a brittle phase, was observed within the grains of the ST alloy during thermal exposure. Previous studies have demonstrated the effect of the distributed α-Cr phase (Fig. 6) on intragranular strengthening [25] and that its hardness value may reach up to 3 times higher than that of the γ matrix [26]. In the present study, the reduction of hardness values, which resulted from coarsening of γ'
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Fig. 6. SEM and TEM micrographs with EDS and XRD data of the ST alloy following thermal exposure at 750 °C for 1000 h: (a) and 3000 h (b); (c) α-Cr phase precipitated in the grain; (d) SAD of α-Cr phase along [111]; (e) EDS spectrum of α-Cr phase; (f) XRD spectrum of α-Cr phase in the ST alloy.
particles, might be compensated by the precipitation of α-Cr phase, so the hardness remained stable during thermal exposure up to 3000 h. According to the morphology of the fracture surfaces of the tensile testing samples, intergranular fracture occurred following tensile testing at 750 °C. Also, the grain boundary failure occurred as the thermal exposure time was increased, due to the transition from the transgranular to the intergranular fracture mode, following tensile testing at room temperature. The average yield strength of the ST alloy at room temperature was decreased after the first 1000 h, due to the γ' particles coarsening, following thermal exposure at 750 °C. The α-Cr phase particles, which are incoherent with the γ matrix, became coarser as the thermal exposure time was increased from 1000 h up to 3000 h and the α-Cr precipitates might have inhibited the dislocation motions within the grains. When dislocations glide in the primary slip plane (i.e., (111) plane) and approach α-Cr precipitates, α-Cr precipitates impeded the gliding of dislocations. It is definitely difficult for the gliding dislocation to cut through the α-Cr precipitate. Firstly, dislocations may partly bypass α-Cr precipitates via Orowan mechanism, and then dislocation loops accumulate at the interface between the matrix and α-Cr precipitates. Because of the dislocation loops pile-up around α-Cr precipitates, higher critical shear stress is necessary for dislocations to bypass α-Cr precipitates [27]. Secondly, much higher driving force is required for dislocations to cross slip to other slip systems, especially to non-primary slip systems, in order to avoid the obstacle of the α-Cr precipitate.
Therefore, the formation of α-Cr phase in the ST alloy made plastic deformation more difficult [13]. The average yield strength of the ST alloy at room temperature was increased due to the substantial precipitation strengthening effect of the α-Cr phase within the grains, though the yield strength was decreased with the coarsening of the γ' particles. The variation of the average yield strength at room temperature was also consistent with the variation of hardness along the thermal exposure. The dispersed α-Cr phase has partially contributed to the stability of the mechanical properties. If the α-Cr precipitates had formed a network of precipitations within the grains or films at the grain boundaries, the alloy would be significantly more brittle, causing deterioration of the mechanical properties. Thus, the effects of the α-Cr phase on the comprehensive mechanical properties of the ST alloy related to the size and distribution of α-Cr precipitates. However, there was insignificant contribution of the α-Cr precipitates to the strength at high temperature, as the fracture mechanism had changed at 750 °C.
4. Conclusions The variation of microstructures and mechanical properties of a new wrought Ni–Fe based superalloy, designed for advanced ultra-supercritical boiler tubes, following optimized heat treatment and long term thermal exposure at 750 °C was studied.
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Fig. 7. Comparison of tensile properties of the ST alloy at room temperature and 750 °C thermal exposure for 3000 h: (a) yield strength; (b) ultimate tensile strength; (c) elongation; (d) area reduction.
(1) The initial recrystallization temperature of this alloy is between 950 °C and 1000 °C. Following the heat treatment optimization the hardness values reached 310 HV. The average grain size of the ST alloy was 130 μm with Cr-rich M23C6 carbides distributed at grain boundaries, while the average size of γ' precipitates was 37 nm.
(2) Following thermal exposure at 750 °C, both the grain size and morphology of M23C6 carbides had no significant change; the coarsening of γ' precipitates from 37 nm to 101 nm and the precipitation of α-Cr occurred. The coarsening rate of γ' particles matches the classic LSW model and the γ' particles kept their spherical morphology. The α-Cr phase was observed following thermal
Fig. 8. SEM fractographs of the tensile tested samples at room temperature: (a) thermal exposure for 0 h; (b) thermal exposure for 1000 h; (c) thermal exposure for 3000 h.
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Fig. 9. SEM fractographs of the tensile tested samples at 750 °C: (a) thermal exposure for 0 h; (b) thermal exposure for 1000 h; (c) thermal exposure for 3000 h.
Fig. 10. (a) Plot of γ' precipitate size during thermal exposure at 750 °C; (b) LSW model.
exposure for 1000 h and the size of needle-like α-Cr phase was significantly increased as thermal exposure time increased. (3) The hardness value of this alloy maintained at approximately 310 HV during thermal exposure up to 3000 h due to the precipitation of α-Cr phase in the grains, though coarsening of γ' particles may result in the reduction of hardness values. The new alloy demonstrated high strength. The grain boundaries of this new alloy were the weaker locations at 750 °C and the grain boundary weakening occurred during thermal exposure. The dispersed α-Cr phase enhanced intragranular strength, implying that the reduction of yield strength at room temperature, which resulted from the coarsening of γ' particles, might have been compensated by the precipitation of α-Cr phase. The dispersed α-Cr particles within the grain could partially improve the stability of mechanical properties during thermal exposure up to 3000 h.
These properties combined with low-cost make this new alloy appropriate for potential use in A-USC power plants.
Acknowledgments This work has been financially supported by China Huaneng Group (Grant No. ZA-12-HKR04), Huaneng Power International Inc. (Grant No. ZA-14-HKR01) and the National Nature Science Foundation of China under Grant No. 51401164.
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