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The influence of long-term thermal exposure on intermediate temperature brittleness behavior of a Nickel-base superalloy
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
Available online at www.sciencedirect.com
www.elsevier.com/locate/matchar
The influence of long-term thermal exposure on intermediate temperature brittleness behavior of a Nickel-base superalloy Y.H. Yang, J.J. Yu⁎, X.F. Sun, T. Jin, H.R. Guan, Z.Q. Hu Superalloy Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China
AR TIC LE D ATA
ABSTR ACT
Article history:
The morphology of γ' precipitates and intermediate temperature brittleness behavior of
Received 19 September 2011
samples aged at 700 °C and 800 °C for different aging times then tested at 700 °C and
Received in revised form 3 February
800 °C, is investigated. The results show that the γ' precipitates remain cuboidal in shape
2012
and regularly aligned in the matrix after long-term thermal exposure at 700 °C and 800 °C.
Accepted 4 February 2012
The size of secondary γ' precipitates increases greatly after long-term exposure at 800 °C. The growth of γ' precipitates would enhance the possibility of dislocation–dislocation and
Keywords:
dislocation–γ' precipitate interactions, resulting in a high yield stress after long-term ther-
M951 superalloy
mal exposure. For samples tested at 700 °C, the elongation of samples aged at 700 °C and
Long-term thermal exposure
800 °C gradually decreases with increasing aging time. However, for samples tested at
Microstructure
800 °C the elongation of samples aged at 700 °C and 800 °C gradually increases, and the
Intermediate temperature brittleness
elongation is higher for than samples tested at 700 °C. In addition, the γ' precipitates are sheared easily by dislocations after long-term thermal exposure. The intermediate temperature brittleness behavior is not eliminated during long-term thermal exposure at intermediate temperatures, while the intermediate temperature brittleness behavior takes place at the lower temperature. Additionally, γ' precipitates sheared by dislocations is dominant deformation mechanism. © 2012 Elsevier Inc. All rights reserved.
1.
Introduction
Nickel-base superalloys have been developed for decades due to their excellent creep and fatigue strength, and good corrosion resistance at elevated temperatures. Investigations have shown that the mechanical properties are influenced by the size, morphology and volume fraction of γ' precipitates [1–4]. M951 alloy, a cast Ni-base superalloy, is being used for advanced vane applications in aircraft and industrial gas turbines, due to its attractive advantages of low density, low cost, foundry performance and high incipient melting temperature and mechanical properties [5–11]. However, almost all Nickel-base superalloys exhibit a minimum in their tensile ductility. This behavior has been named intermediate temperature brittleness [12]. The intermediate temperature brittleness behavior has been reported in many polycrystalline ⁎ Corresponding author. Tel.: +86 24 23971713; fax: + 86 24 23971758. E-mail address: [email protected] (J.J. Yu). 1044-5803/$ – see front matter © 2012 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2012.02.004
superalloys [12–14]. Previous reports showed that the intermediate temperature brittleness behavior can be eliminated during long-term thermal exposure at elevated temperatures, due to γ' precipitate coarsening [12,15]. The microstructure and related mechanical properties after long-term thermal exposure at elevated temperatures has been extensively investigated [15–17]. The purpose of this work is to investigate the microstructure and intermediate temperature brittleness behavior after long-term thermal exposure at intermediate temperatures.
2.
Experimental Procedures
M951 alloy has the nominal composition (wt. %) of 0.05C, 9.0Cr, 5.0Co, 3.5W, 3.0Mo, 2.2Nb, 5.9Al, 0.02Y, 0.024B, and the
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
31
Fig. 1 – The γ' morphology of M951 alloy after standard heat treatment. (a) γ' precipitates; (b) secondary γ' precipitates.
rest Ni. The master alloy was melted in a VIM25F vacuum induction furnace, and remelted in 10 kg ZG-0.01 vacuum induction furnace, then cast into test bars. The corresponding mold preheating and pouring temperatures were 900 °C and 1450 °C, respectively. The samples were subjected to standard heat treatment, 1100 °C/4 h and air cooled (AC) to room temperature. After standard heat treatment, samples were exposed at 700 °C and 800 °C for 25, 50, 100, 500 and up to 1000 h. Tensile test specimens with a diameter of 5 mm and a gage length of 25 mm were machined longitudinally from the heat-treated bars. Tensile tests were performed at 700 °C and 800 °C. The strain rate was maintained at 0.5 mm/min up to yield and 2.5 mm/min after yield. Each test value, including yield strength and elongation, represented an average of at least two test results. Scanning electron microscope and transmission electron microscopy were used to examine the microstructure and
deformation structure of M951 alloy. The samples for scanning electron microscope were electrolyzed in a solution of 20 ml HNO3 + 40 ml CH3COOH + 340 ml H2O at 7 V. The foils for transmission electron microscopy were cut normal to the stress axis from the ruptured samples. They were prepared by twin-jet thinning electrolytically in a solution of 10% perchloric acid and 90% ethanol at −20 °C and observed using a Jeol 2000FX transmission electron microscope operated at 200 KV.
3.
Results and discussion
3.1.
Influence of thermal exposure on the γ' precipitates
After the standard heat treatment, the alloy consists of a bimodal γ' precipitate distribution with cuboidal particles of
Fig. 2 – Morphology of γ' precipitates in samples aged for different amounts of time at 700 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
32
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
Fig. 3 – Morphology of secondary γ' precipitates in samples aged for different amounts of time at 700 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
0.73 μm average edge length (Fig. 1a) and spherical particles of 60–70 nm average diameter as shown in Fig. 1(b). Changes of the γ' precipitates at the aging temperature of 700 °C as the function of time are shown in Fig. 2. It is clear that the larger γ' precipitates remain cuboidal in shape and
align regularly in the matrix after 1000 h long-term aging at 700 °C (Fig. 2d). The size of γ' precipitates is not obviously influenced by aging at 700 °C. Fig. 3 exhibits the morphology of secondary γ' precipitates after long-term aging at 700 °C. It can be seen that the
Fig. 4 – Morphology of γ' precipitates in samples aged for different amounts of time at 800 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
33
Fig. 5 – Morphology of secondary γ' precipitates in samples aged for different amounts of time at 800 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
secondary γ' precipitates remain spherical in shape, and the size of these particles is also not obviously influenced by aging at 700 °C. This is due to the low temperature and short aging time which are related to diffusion of the alloying elements. For samples aged at 800 °C for various times, the morphology of the γ' precipitates is shown in Fig. 4. It can be seen that the shape of the γ' precipitates becomes more cuboidal with increasing aging time. The change of the shape of the γ' precipitates is due to the minimizing the elastic energy, which is not only associated with the misfit between γ' precipitates and matrix, but also the size and volume fraction of the γ' precipitates [11]. Fig. 5 exhibits the morphology of the secondary γ' precipitates after long-term aging at 800 °C. It is clear that the
secondary γ' precipitates remain spherical in shape, and the size of these particles is also obviously influenced by aging at 800 °C. With increasing aging time (up to 1000 h), the size of the secondary γ' precipitates increases as shown in Fig. 5(d). The size of secondary γ' precipitates is about 300 nm after exposure at 800 °C for 1000 h. The growth of γ' precipitates is influenced by the diffusion of the alloying elements through the matrix [18]. Because the diffusion rate of alloying elements at 800 °C is faster than at 700 °C, the γ' precipitates grow larger than they do at 700 °C. Consequently, the longer the aging time, the larger the γ' precipitates. Additionally, the growth rates are related to the diffusion distances of alloying elements which are the nearly identical in various directions, so, the shape of γ' precipitates
Fig. 6 – Influence of aging time on tensile property of M951 alloy. (a) tensile properties at 700 °C; (b) tensile properties at 800 °C.
34
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
Fig. 7 – Deformation structures of samples after testing at 700 °C. (a) aging at 700 °C for 25 h; (b) aging at 800 °C for 25 h; (c) aging at 700 °C for 50 h ; (d) aging at 800 °C for 50 h.
remain cuboidal and spherical for secondary γ' precipitates, respectively, and only the size of γ' precipitates slightly increases.
3.2.
Influence Of Thermal Exposure On Tensile Properties
Fig. 6 shows the tensile properties of samples aged at 700 °C and 800 °C for different aging times and then tested at 700 °C and 800 °C. The tensile properties at 700 °C of samples aged at 700 °C and 800 °C with various aging times are shown in Fig. 6(a). It can be seen that the yield stress falls sharply after 50 h aging. Previous investigation showed that the coarsening of γ' precipitates lead to a decrease in yield stress [15–17]. The size and morphology of γ' precipitates after 50 h aging is similar to that after 25 h aging, due to the intermediate aging temperature (700 °C and 800 °C) and short aging time (50 h) in this experiment. Thus, the drop in yield stress might be mainly attributed to a reduction in interfacial energy or misfit stress at γ/γ' interface. The growth of γ' precipitate would also enhance the possibility of dislocation–dislocation and dislocation–γ' precipitates interactions, resulting in high yield stress after long-term thermal exposure [15]. With increasing aging time, the elongation of samples aged at 700 °C gradually decreases when the aging time is less than 500 h. Although the growth of γ' precipitates can enhance yield stress, it also decreases the elongation at the early stage of aging. The trend in elongation of samples aged at 800 °C is inverse to the yield strength. This phenomenon is commonly observed in other superalloys [15,19]. The results of tensile tests at 800 °C of samples aged at 700 °C and 800 °C are shown in Fig. 6(b). The yield stress falls after 50 h aging. Then it increases slightly after 100 h aging, while no obvious difference seen when the aging
time is increased 1000 h, due to the growth of secondary γ' precipitates, which can enhance dislocation–precipitate interactions around secondary γ' precipitates. As for elongation, it shows peak value after 50 h aging, due to the reduction in interfacial energy at γ/γ' interface and coherency strains at the early stage of aging. Then it drops sharply after 100 h aging, and correspondingly increases with the increment of aging time. This is due to the fact that the size of γ' precipitates is larger leading to easier movement of dislocations. Meanwhile, the elongation of samples aged at 800 °C is higher than samples aged at 700 °C. The higher the aging temperature and aging time, the more the interfacial energy and coherency strains are balanced at γ/γ' interface. Therefore, the easier motion of dislocations in the matrix probably leads to higher elongation.
3.3.
Deformation Structures
3.3.1.
Tensile at 700 °C
Fig. 7 shows the deformation structures of samples aged at 700 °C and 800 °C for various times and tested at 700 °C. Most of dislocations are in the γ matrix and dislocations tangle at the γ/γ' interface (Fig. 7a), impeding further motion. The dislocation density is higher in samples aged at 800 °C for 25 h, as shown in Fig. 7(b). In addition, there are also a few dislocations cutting into the γ' precipitates and stacking faults are observed in the γ' precipitates (Fig. 7b). These stacking faults are observed in some alloys and have been investigated by researchers. Their results show that the formation of stacking faults is attributed to the dislocation of type a/3 < 112 > shearing the γ' precipitates [20]. The tangled dislocations impede dislocation movement in the matrix and inhibit further plastic deformation. Then, another deformation mechanisms, γ'
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
35
Fig. 8 – Deformation structures of samples aged at 700 °C for 1000 h after testing at 700 °C. (a) dislocations within the γ' precipitates; (b) an array of dislocations.
precipitates sheared by stacking faults, is operative at high stress levels. Consequently, the shearing of γ' precipitates by stacking faults contributes to the high yield stress and low elongation of samples aged at 800 °C for 25 h. Fig. 7(c) shows the deformation of structures of samples aged at 700 °C for 50 h. It can be seen that the dislocations are mainly in the γ matrix and there are tangles at the γ/γ' interface. In addition, a few dislocations cutting into the γ' precipitates are also observed in Fig. 7(c), indicating that it is difficult for dislocations to move in the matrix and further plastic deformation, thus the samples show low elongation. Slip bands can be observed in samples aged at 800 °C for 50 h as shown in Fig. 7(d). The slip bands are mostly constrained in the matrix and γ' precipitates, which indicates that the γ' shearing by dislocation is occurring in the slip bands. The observation of slip bands is evidence of homogeneous deformation, and indicates that the samples undergo homogeneous deformation during deformation at 700 °C, thus, the samples exhibit high elongation. Comparing to samples aged for a shorter time, the dislocations in samples aged for 1000 h are mainly distributed within the γ' precipitates (Fig. 8a) due to the high antiphase boundary energy [21]. Additionally, the array of dislocation is also observed (Fig. 8b), which is another evidence of γ' shearing. It is suggested that the γ' precipitates are sheared very easily by dislocations after long-term thermal exposure due to the lower strength of γ' precipitates.
3.3.2.
strength of γ' precipitates. During the long-term exposure to heat, the growth of γ' precipitates is influenced by the diffusion of alloying elements, the growth of γ' precipitates would lead to the removal of strengthening elements from γ' precipitates by diffusion. Additionally, the interface energy
Tensile at 800 °C
Fig. 9 shows the deformation structures of samples aged at 700 °C and 800 °C for various times and tested at 800 °C. Although the tensile strengths are very similar, the deformation structures are quite different at this condition. High dislocation density leads to tangles at the interface of γ/γ' precipitates, and stacking faults are observed in γ' precipitates (Fig. 9a). By increasing the aging temperature to 800 °C, dislocation arrays are observed in the γ' precipitates (Fig. 9b), indicating that γ' precipitates are sheared easily by dislocations after aging for 50 h at 800 °C. With an increase in aging time, slip bands are formed in the γ' precipitates, indicating that γ' precipitates are sheared by dislocations in the slip band (Fig. 9c). It is also found that dislocation loops are left in the γ/γ' interface as shown in Fig. 9(c). These microstructures demonstrate that the increase of aging temperature and aging times can weaken the
Fig. 9 – Deformation structures of samples after testing at 800 °C. (a) aging at 700 °C for 50 h; (b) aging at 800 °C for 50 h; (c) aging at 700 °C for 1000 h.
36
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
and strains at γ/γ' interface are much lower which allows dislocations to cut the γ' precipitates easily. The longer the aging times, the more balanced localized plastic flow, and easier motion for dislocation in matrix [15]. Therefore, the samples exhibit high elongation after long-term exposure. Since the γ' precipitates are weakened after long-term thermal exposure and bimodal γ' precipitate distribution exists in M951 alloy, the high yield strength after long-term thermal exposure could be due to the fine secondary γ' precipitates within the matrix channels and γ' precipitate shearing [22].
3.4.
This work is financially supported by the National Basic Research Program (973 Program) of China under grant no. 2010CB631200 (2010CB631206), and the National Natural Science Foundation of China (NSFC) under grant nos. 50931004 and U 1037601. The authors are grateful for those supports.
REFERENCES
Change In Ductility Minimum
Intermediate temperature brittleness behavior takes place in this alloy at 800 °C [9]. Intermediate temperature brittleness behavior takes place at 700 °C after long-term exposure (Fig. 6) in this experiment. That is to say, the intermediate temperature brittleness behavior cannot be eliminated by long-term thermal exposure. However long-term thermal exposure can lower the temperature at which the intermediate temperature brittleness behavior takes place. According to Fiore [23], there are a number of causes of ductility minimums in alloys. In the superalloys the instability of γ' precipitates, deformation mechanisms and embrittlement of grain boundaries are the most important reasons. The first factor is not important for the intermediate temperature brittleness behavior of M951 alloy. The γ' precipitates are stable below 900 °C [11], and the MC carbide has good stability [24]. Furthermore, the dominant deformation is γ' cutting at intermediate temperatures, it indicates that the change in intermediate temperature brittleness behavior must be mainly attributed to the embrittlement of grain boundaries in the present experiment. The trace elements tend to migrate to the grain boundaries and weaken it due to the change in grain boundary electronic structure [25]. Then the grain boundaries decohesion occurs easily by an applied tensile stress leading to the embrittlement of grain boundaries. When the tensile test is performed at 800 °C, the stress concentration obviously cannot reach the critical value for grain boundary decohesion due to gliding of dislocations in slip bands [12]. However, the large stress concentrations at the grain boundary can easily reach the critical value for grain boundaries decohesion at 700 °C.
4.
Acknowledgements
Conclusions
In summary, the size of secondary γ' precipitates increases greatly after long-term exposure at 800 °C. The growth of γ' precipitates is diffusion. The yield stress drops sharply at the early stage of aging, due to the release of interfacial energy and coherency strains. It increases after long-term exposure due to the effect of dislocation–dislocation and dislocation–γ' precipitate interactions. Generally, the elongation increases with an increase in aging time, because of the balanced interfacial energy and coherency strains and the weakened γ' precipitates which are sheared easily by dislocations. The intermediate temperature brittleness behavior is not eliminated during long-term thermal exposure at intermediate temperatures, but takes place at lower temperatures.
[1] Reed PAS, Sinclair I, Wu XD. Fatigue crack path prediction in UDIMET 720 nickel-based alloy single crystals. Metall Mater Trans A 2000;31:109–23. [2] Koble M, Neuking K, Eggeler G. Dislocation reactions and microstructural instability during 1025 °C shear creep testing of superalloy single crystals. Mater Sci Eng A 1997;234–236: 877–9. [3] Sczerzenie F, Maurer GE. Developments in disc materials. Mater Sci Technol 1987;3:733–42. [4] Zheng J, Powell BE. Effect of stress ratio and test methods on fatigue crack growth rate for nickel based superalloy Udimet 720. Int J Fatigue 1999;21:507–13. [5] Lian ZW, Yu JJ, Sun XF, Guan HR, Hu ZQ. Influence of heat treatments on the microstructure and tensile properties of cast Nickel-based superalloy M951. Rare Met Mater Eng 2008;37:798–802. [6] Yu JJ. Thermal fatigue behaviour of superalloy M951. J Iron Steel Res 2003;15:179–82. [7] Zhou PJ, Yu JJ, Sun XF, Guan HR, Hu ZQ. Role of yttrium in the microstructure and mechanical properties of a boron-modified nickel-based superalloy. Scr Mater 2007;57: 643–6. [8] Zhou PJ, Yu JJ, Sun XF, Guan HR, Hu ZQ. The role of boron on a conventional nickel-based superalloy. Mater Sci Eng A 2008;491:159–63. [9] Lian ZW, Yu JJ, Sun XF, Guan HR, Hu ZQ. Temperature dependence of tensile behavior of Ni-based superalloy M951. Mater Sci Eng A 2008;489:227–33. [10] Zhou PJ, Yu JJ, Sun XF, Guan HR, Hu ZQ. Influence of boron on mechanical properties of M 951 alloy. Trans Nonferrous Met Soc China 2005;15:86–9. [11] Yu JJ, Lian ZW, Sun XF, Guan HR, Hu ZQ. Properties and microstructures of M951 alloy after long-term exposure. Mater Sci Eng A 2010;527:1896–902. [12] He LZ, Zheng Q, Sun XF, Hou GC, Guan HR, Hu ZQ. Low ductility at intermediate temperature of Ni–base superalloy M963. Mater Sci Eng A 2004;380:340–8. [13] Jensen RR, Tien JK. Temperature and strain rate dependence of stress–strain behavior in a nickel-base superalloy. Metall Trans A 1985;16A:1049–68. [14] Copley SM, Kear BH, Rowe GM. The temperature and orientation dependence of yielding in Mar-M200 single crystals. Mater Sci Eng A 1972;10:87–92. [15] Mostaghim RS, Asgari S. The influence of thermal exposure on the γ′ precipitates characteristics and tensile behavior of superalloy IN-738LC. J Mater Process Technol 2004;147: 343–50. [16] Acharya MV, Fuchs GE. The effect of long-term thermal exposures on the microstructure and properties of CMSX-10 single crystal Ni-base superalloys. Mater Sci Eng A 2004;381: 143–53. [17] Hou JS, Guo JT, Yang GX, Zhou LZ, Qin XZ, Ye HQ. The microstructural instability of a hot corrosion resistant superalloy during long-term exposure. Mater Sci Eng A 2008;498:349–58.
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[18] Ges A, Fornaro O, Palalio H. Long term coarsening of γ′ precipitates in a Ni-base superalloy. J Mater Sci 1997;32: 3687–91. [19] Luo ZP, Wu ZT, Miller DJ. The dislocation microstructure of a nickel-base single-crystal superalloy after tensile fracture. Mater Sci Eng A 2003;354:358–68. [20] Caron P, Khan T, Veyssiere P. On precipitate shearing by superlattice stacking faults in superalloys. Philos Mag A 1988;57:859–75. [21] Milligan WW, Antolovich SD. Yielding and deformation behavior of the single crystal superalloy PWA 1480. Metall Trans A 1987;18:85–95.
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[22] Balikci E, Mirshams RA, Raman A. Tensile strengthening in the nickel-base superalloy IN738LC. J Mater Eng Perform 2000;9:324–9. [23] Fiore NF. Rev. Mid-range ductility minimum in Ni-base superalloys. High Temp Mater 1975;2:373–408. [24] Xia PC, Yu JJ, Sun XF, Guan HR, Hu ZQ. The influence of thermal exposure on the microstructure and stress rupture property of DZ951 nickel-base alloy. J Alloys Compd 2007;443: 125–31. [25] Eberhart ME, Vvedensky DD. Localized grain-boundary electronic states and intergranular fracture. Phys Rev Lett 1987;58:61–4.
Available online at www.sciencedirect.com
www.elsevier.com/locate/matchar
The influence of long-term thermal exposure on intermediate temperature brittleness behavior of a Nickel-base superalloy Y.H. Yang, J.J. Yu⁎, X.F. Sun, T. Jin, H.R. Guan, Z.Q. Hu Superalloy Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China
AR TIC LE D ATA
ABSTR ACT
Article history:
The morphology of γ' precipitates and intermediate temperature brittleness behavior of
Received 19 September 2011
samples aged at 700 °C and 800 °C for different aging times then tested at 700 °C and
Received in revised form 3 February
800 °C, is investigated. The results show that the γ' precipitates remain cuboidal in shape
2012
and regularly aligned in the matrix after long-term thermal exposure at 700 °C and 800 °C.
Accepted 4 February 2012
The size of secondary γ' precipitates increases greatly after long-term exposure at 800 °C. The growth of γ' precipitates would enhance the possibility of dislocation–dislocation and
Keywords:
dislocation–γ' precipitate interactions, resulting in a high yield stress after long-term ther-
M951 superalloy
mal exposure. For samples tested at 700 °C, the elongation of samples aged at 700 °C and
Long-term thermal exposure
800 °C gradually decreases with increasing aging time. However, for samples tested at
Microstructure
800 °C the elongation of samples aged at 700 °C and 800 °C gradually increases, and the
Intermediate temperature brittleness
elongation is higher for than samples tested at 700 °C. In addition, the γ' precipitates are sheared easily by dislocations after long-term thermal exposure. The intermediate temperature brittleness behavior is not eliminated during long-term thermal exposure at intermediate temperatures, while the intermediate temperature brittleness behavior takes place at the lower temperature. Additionally, γ' precipitates sheared by dislocations is dominant deformation mechanism. © 2012 Elsevier Inc. All rights reserved.
1.
Introduction
Nickel-base superalloys have been developed for decades due to their excellent creep and fatigue strength, and good corrosion resistance at elevated temperatures. Investigations have shown that the mechanical properties are influenced by the size, morphology and volume fraction of γ' precipitates [1–4]. M951 alloy, a cast Ni-base superalloy, is being used for advanced vane applications in aircraft and industrial gas turbines, due to its attractive advantages of low density, low cost, foundry performance and high incipient melting temperature and mechanical properties [5–11]. However, almost all Nickel-base superalloys exhibit a minimum in their tensile ductility. This behavior has been named intermediate temperature brittleness [12]. The intermediate temperature brittleness behavior has been reported in many polycrystalline ⁎ Corresponding author. Tel.: +86 24 23971713; fax: + 86 24 23971758. E-mail address: [email protected] (J.J. Yu). 1044-5803/$ – see front matter © 2012 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2012.02.004
superalloys [12–14]. Previous reports showed that the intermediate temperature brittleness behavior can be eliminated during long-term thermal exposure at elevated temperatures, due to γ' precipitate coarsening [12,15]. The microstructure and related mechanical properties after long-term thermal exposure at elevated temperatures has been extensively investigated [15–17]. The purpose of this work is to investigate the microstructure and intermediate temperature brittleness behavior after long-term thermal exposure at intermediate temperatures.
2.
Experimental Procedures
M951 alloy has the nominal composition (wt. %) of 0.05C, 9.0Cr, 5.0Co, 3.5W, 3.0Mo, 2.2Nb, 5.9Al, 0.02Y, 0.024B, and the
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
31
Fig. 1 – The γ' morphology of M951 alloy after standard heat treatment. (a) γ' precipitates; (b) secondary γ' precipitates.
rest Ni. The master alloy was melted in a VIM25F vacuum induction furnace, and remelted in 10 kg ZG-0.01 vacuum induction furnace, then cast into test bars. The corresponding mold preheating and pouring temperatures were 900 °C and 1450 °C, respectively. The samples were subjected to standard heat treatment, 1100 °C/4 h and air cooled (AC) to room temperature. After standard heat treatment, samples were exposed at 700 °C and 800 °C for 25, 50, 100, 500 and up to 1000 h. Tensile test specimens with a diameter of 5 mm and a gage length of 25 mm were machined longitudinally from the heat-treated bars. Tensile tests were performed at 700 °C and 800 °C. The strain rate was maintained at 0.5 mm/min up to yield and 2.5 mm/min after yield. Each test value, including yield strength and elongation, represented an average of at least two test results. Scanning electron microscope and transmission electron microscopy were used to examine the microstructure and
deformation structure of M951 alloy. The samples for scanning electron microscope were electrolyzed in a solution of 20 ml HNO3 + 40 ml CH3COOH + 340 ml H2O at 7 V. The foils for transmission electron microscopy were cut normal to the stress axis from the ruptured samples. They were prepared by twin-jet thinning electrolytically in a solution of 10% perchloric acid and 90% ethanol at −20 °C and observed using a Jeol 2000FX transmission electron microscope operated at 200 KV.
3.
Results and discussion
3.1.
Influence of thermal exposure on the γ' precipitates
After the standard heat treatment, the alloy consists of a bimodal γ' precipitate distribution with cuboidal particles of
Fig. 2 – Morphology of γ' precipitates in samples aged for different amounts of time at 700 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
32
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
Fig. 3 – Morphology of secondary γ' precipitates in samples aged for different amounts of time at 700 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
0.73 μm average edge length (Fig. 1a) and spherical particles of 60–70 nm average diameter as shown in Fig. 1(b). Changes of the γ' precipitates at the aging temperature of 700 °C as the function of time are shown in Fig. 2. It is clear that the larger γ' precipitates remain cuboidal in shape and
align regularly in the matrix after 1000 h long-term aging at 700 °C (Fig. 2d). The size of γ' precipitates is not obviously influenced by aging at 700 °C. Fig. 3 exhibits the morphology of secondary γ' precipitates after long-term aging at 700 °C. It can be seen that the
Fig. 4 – Morphology of γ' precipitates in samples aged for different amounts of time at 800 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
33
Fig. 5 – Morphology of secondary γ' precipitates in samples aged for different amounts of time at 800 °C. (a) 25 h; (b) 50 h; (c) 100 h; (d) 1000 h.
secondary γ' precipitates remain spherical in shape, and the size of these particles is also not obviously influenced by aging at 700 °C. This is due to the low temperature and short aging time which are related to diffusion of the alloying elements. For samples aged at 800 °C for various times, the morphology of the γ' precipitates is shown in Fig. 4. It can be seen that the shape of the γ' precipitates becomes more cuboidal with increasing aging time. The change of the shape of the γ' precipitates is due to the minimizing the elastic energy, which is not only associated with the misfit between γ' precipitates and matrix, but also the size and volume fraction of the γ' precipitates [11]. Fig. 5 exhibits the morphology of the secondary γ' precipitates after long-term aging at 800 °C. It is clear that the
secondary γ' precipitates remain spherical in shape, and the size of these particles is also obviously influenced by aging at 800 °C. With increasing aging time (up to 1000 h), the size of the secondary γ' precipitates increases as shown in Fig. 5(d). The size of secondary γ' precipitates is about 300 nm after exposure at 800 °C for 1000 h. The growth of γ' precipitates is influenced by the diffusion of the alloying elements through the matrix [18]. Because the diffusion rate of alloying elements at 800 °C is faster than at 700 °C, the γ' precipitates grow larger than they do at 700 °C. Consequently, the longer the aging time, the larger the γ' precipitates. Additionally, the growth rates are related to the diffusion distances of alloying elements which are the nearly identical in various directions, so, the shape of γ' precipitates
Fig. 6 – Influence of aging time on tensile property of M951 alloy. (a) tensile properties at 700 °C; (b) tensile properties at 800 °C.
34
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Fig. 7 – Deformation structures of samples after testing at 700 °C. (a) aging at 700 °C for 25 h; (b) aging at 800 °C for 25 h; (c) aging at 700 °C for 50 h ; (d) aging at 800 °C for 50 h.
remain cuboidal and spherical for secondary γ' precipitates, respectively, and only the size of γ' precipitates slightly increases.
3.2.
Influence Of Thermal Exposure On Tensile Properties
Fig. 6 shows the tensile properties of samples aged at 700 °C and 800 °C for different aging times and then tested at 700 °C and 800 °C. The tensile properties at 700 °C of samples aged at 700 °C and 800 °C with various aging times are shown in Fig. 6(a). It can be seen that the yield stress falls sharply after 50 h aging. Previous investigation showed that the coarsening of γ' precipitates lead to a decrease in yield stress [15–17]. The size and morphology of γ' precipitates after 50 h aging is similar to that after 25 h aging, due to the intermediate aging temperature (700 °C and 800 °C) and short aging time (50 h) in this experiment. Thus, the drop in yield stress might be mainly attributed to a reduction in interfacial energy or misfit stress at γ/γ' interface. The growth of γ' precipitate would also enhance the possibility of dislocation–dislocation and dislocation–γ' precipitates interactions, resulting in high yield stress after long-term thermal exposure [15]. With increasing aging time, the elongation of samples aged at 700 °C gradually decreases when the aging time is less than 500 h. Although the growth of γ' precipitates can enhance yield stress, it also decreases the elongation at the early stage of aging. The trend in elongation of samples aged at 800 °C is inverse to the yield strength. This phenomenon is commonly observed in other superalloys [15,19]. The results of tensile tests at 800 °C of samples aged at 700 °C and 800 °C are shown in Fig. 6(b). The yield stress falls after 50 h aging. Then it increases slightly after 100 h aging, while no obvious difference seen when the aging
time is increased 1000 h, due to the growth of secondary γ' precipitates, which can enhance dislocation–precipitate interactions around secondary γ' precipitates. As for elongation, it shows peak value after 50 h aging, due to the reduction in interfacial energy at γ/γ' interface and coherency strains at the early stage of aging. Then it drops sharply after 100 h aging, and correspondingly increases with the increment of aging time. This is due to the fact that the size of γ' precipitates is larger leading to easier movement of dislocations. Meanwhile, the elongation of samples aged at 800 °C is higher than samples aged at 700 °C. The higher the aging temperature and aging time, the more the interfacial energy and coherency strains are balanced at γ/γ' interface. Therefore, the easier motion of dislocations in the matrix probably leads to higher elongation.
3.3.
Deformation Structures
3.3.1.
Tensile at 700 °C
Fig. 7 shows the deformation structures of samples aged at 700 °C and 800 °C for various times and tested at 700 °C. Most of dislocations are in the γ matrix and dislocations tangle at the γ/γ' interface (Fig. 7a), impeding further motion. The dislocation density is higher in samples aged at 800 °C for 25 h, as shown in Fig. 7(b). In addition, there are also a few dislocations cutting into the γ' precipitates and stacking faults are observed in the γ' precipitates (Fig. 7b). These stacking faults are observed in some alloys and have been investigated by researchers. Their results show that the formation of stacking faults is attributed to the dislocation of type a/3 < 112 > shearing the γ' precipitates [20]. The tangled dislocations impede dislocation movement in the matrix and inhibit further plastic deformation. Then, another deformation mechanisms, γ'
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
35
Fig. 8 – Deformation structures of samples aged at 700 °C for 1000 h after testing at 700 °C. (a) dislocations within the γ' precipitates; (b) an array of dislocations.
precipitates sheared by stacking faults, is operative at high stress levels. Consequently, the shearing of γ' precipitates by stacking faults contributes to the high yield stress and low elongation of samples aged at 800 °C for 25 h. Fig. 7(c) shows the deformation of structures of samples aged at 700 °C for 50 h. It can be seen that the dislocations are mainly in the γ matrix and there are tangles at the γ/γ' interface. In addition, a few dislocations cutting into the γ' precipitates are also observed in Fig. 7(c), indicating that it is difficult for dislocations to move in the matrix and further plastic deformation, thus the samples show low elongation. Slip bands can be observed in samples aged at 800 °C for 50 h as shown in Fig. 7(d). The slip bands are mostly constrained in the matrix and γ' precipitates, which indicates that the γ' shearing by dislocation is occurring in the slip bands. The observation of slip bands is evidence of homogeneous deformation, and indicates that the samples undergo homogeneous deformation during deformation at 700 °C, thus, the samples exhibit high elongation. Comparing to samples aged for a shorter time, the dislocations in samples aged for 1000 h are mainly distributed within the γ' precipitates (Fig. 8a) due to the high antiphase boundary energy [21]. Additionally, the array of dislocation is also observed (Fig. 8b), which is another evidence of γ' shearing. It is suggested that the γ' precipitates are sheared very easily by dislocations after long-term thermal exposure due to the lower strength of γ' precipitates.
3.3.2.
strength of γ' precipitates. During the long-term exposure to heat, the growth of γ' precipitates is influenced by the diffusion of alloying elements, the growth of γ' precipitates would lead to the removal of strengthening elements from γ' precipitates by diffusion. Additionally, the interface energy
Tensile at 800 °C
Fig. 9 shows the deformation structures of samples aged at 700 °C and 800 °C for various times and tested at 800 °C. Although the tensile strengths are very similar, the deformation structures are quite different at this condition. High dislocation density leads to tangles at the interface of γ/γ' precipitates, and stacking faults are observed in γ' precipitates (Fig. 9a). By increasing the aging temperature to 800 °C, dislocation arrays are observed in the γ' precipitates (Fig. 9b), indicating that γ' precipitates are sheared easily by dislocations after aging for 50 h at 800 °C. With an increase in aging time, slip bands are formed in the γ' precipitates, indicating that γ' precipitates are sheared by dislocations in the slip band (Fig. 9c). It is also found that dislocation loops are left in the γ/γ' interface as shown in Fig. 9(c). These microstructures demonstrate that the increase of aging temperature and aging times can weaken the
Fig. 9 – Deformation structures of samples after testing at 800 °C. (a) aging at 700 °C for 50 h; (b) aging at 800 °C for 50 h; (c) aging at 700 °C for 1000 h.
36
MA TE RI A L S CH A R A CT ER IZ A TI O N 6 6 (2 0 1 2) 3 0– 3 7
and strains at γ/γ' interface are much lower which allows dislocations to cut the γ' precipitates easily. The longer the aging times, the more balanced localized plastic flow, and easier motion for dislocation in matrix [15]. Therefore, the samples exhibit high elongation after long-term exposure. Since the γ' precipitates are weakened after long-term thermal exposure and bimodal γ' precipitate distribution exists in M951 alloy, the high yield strength after long-term thermal exposure could be due to the fine secondary γ' precipitates within the matrix channels and γ' precipitate shearing [22].
3.4.
This work is financially supported by the National Basic Research Program (973 Program) of China under grant no. 2010CB631200 (2010CB631206), and the National Natural Science Foundation of China (NSFC) under grant nos. 50931004 and U 1037601. The authors are grateful for those supports.
REFERENCES
Change In Ductility Minimum
Intermediate temperature brittleness behavior takes place in this alloy at 800 °C [9]. Intermediate temperature brittleness behavior takes place at 700 °C after long-term exposure (Fig. 6) in this experiment. That is to say, the intermediate temperature brittleness behavior cannot be eliminated by long-term thermal exposure. However long-term thermal exposure can lower the temperature at which the intermediate temperature brittleness behavior takes place. According to Fiore [23], there are a number of causes of ductility minimums in alloys. In the superalloys the instability of γ' precipitates, deformation mechanisms and embrittlement of grain boundaries are the most important reasons. The first factor is not important for the intermediate temperature brittleness behavior of M951 alloy. The γ' precipitates are stable below 900 °C [11], and the MC carbide has good stability [24]. Furthermore, the dominant deformation is γ' cutting at intermediate temperatures, it indicates that the change in intermediate temperature brittleness behavior must be mainly attributed to the embrittlement of grain boundaries in the present experiment. The trace elements tend to migrate to the grain boundaries and weaken it due to the change in grain boundary electronic structure [25]. Then the grain boundaries decohesion occurs easily by an applied tensile stress leading to the embrittlement of grain boundaries. When the tensile test is performed at 800 °C, the stress concentration obviously cannot reach the critical value for grain boundary decohesion due to gliding of dislocations in slip bands [12]. However, the large stress concentrations at the grain boundary can easily reach the critical value for grain boundaries decohesion at 700 °C.
4.
Acknowledgements
Conclusions
In summary, the size of secondary γ' precipitates increases greatly after long-term exposure at 800 °C. The growth of γ' precipitates is diffusion. The yield stress drops sharply at the early stage of aging, due to the release of interfacial energy and coherency strains. It increases after long-term exposure due to the effect of dislocation–dislocation and dislocation–γ' precipitate interactions. Generally, the elongation increases with an increase in aging time, because of the balanced interfacial energy and coherency strains and the weakened γ' precipitates which are sheared easily by dislocations. The intermediate temperature brittleness behavior is not eliminated during long-term thermal exposure at intermediate temperatures, but takes place at lower temperatures.
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