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Tensile deformation and fracture characteristics of delta-processed Inconel 718 alloy at elevated temperature
Materials Science and Engineering A 528 (2011) 6253–6258
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Tensile deformation and fracture characteristics of delta-processed Inconel 718 alloy at elevated temperature Shi-Hong Zhang ∗ , Hai-Yan Zhang, Ming Cheng Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
a r t i c l e
i n f o
Article history: Received 17 December 2010 Received in revised form 21 April 2011 Accepted 22 April 2011 Available online 30 April 2011 Keywords: Inconel 718 Delta-processed ␦ phase Tensile deformation Fracture
a b s t r a c t Three specimens with different initial ␦ phase contents have been used to study the tensile deformation and fracture characteristics of the delta-processed Inconel 718 alloy at elevated temperatures by tensile tests at 950 ◦ C. The results indicated that the tensile stress–strain curves of the three specimens were the elastic-uniform plastic curves, and there were two deformation processes during the uniform plastic deformation stage of the delta-processed specimens. The strain hardening exponent in the first deformation process was higher than that in the second process, and the value of the strain hardening exponent increased as the content of ␦ phase increased. In addition, the plasticity of the specimen at elevated temperatures decreased as the content of ␦ phase increased. However, the fracture mechanisms for all the specimens were all microvoid coalescence ductile fracture, and the ␦ phase and carbide were the nucleuses for the formation of micropores. The elongation for the specimen with 8.21% preprecipitated ␦ phase was still as high as 80%. Thus, delta-processed Inconel 718 alloy presented excellent high-temperature plasticity. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Inconel 718 is an important material used for aviation and aerospace engines because of its excellent mechanical properties in the temperature range from −253 ◦ C to 650 ◦ C. In order to improve the safety and reliability of engines, it is crucial to obtain the forgings with uniform and fine microstructures. Generally, the forgings with large size and complex shape, such as turbine disks and engine shafts, are manufactured by multi-stage hot working processes. In addition, the microstructure of the alloy is sensitive to the hot deformation parameters. Therefore, the defects of coarse grain or duplex grain often appear in the forgings. As the ␦ phase in Inconel 718 can control grain size through the strong pinning effect, the Delta Process (DP) has been applied to the forging of Inconel 718, which uses an intentional ␦ phase precipitation cycle and subsequent thermomechanical processing to produce uniform fine grain billet and bar stock [1]. A large size billet of Inconel 718 with the average grain size of ASTM 8 and a turbine disk forging with grain size of ASTM 11, have been obtained by the DP process [1–3]. The ␦ phase plays a significant role in the microstructure evolution and mechanical properties for Inconel 718 [4–10]. Yuan [7] and Wang [8] investigated the effect of ␦ phase on the deforma-
∗ Corresponding author. Fax: +86 24 23906831. E-mail address: [email protected] (Shi-Hong Zhang). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.04.074
tion behaviors of delta-processed Inconel 718, and their results showed that the ␦ phase can stimulate the occurrence of dynamic recrystallization. Desvallées et al. [9] found that the ␦ phase had a detrimental effect on the yield strength but no effect on the ultimate tensile strength at the room temperature and 550 ◦ C. Zhang et al. [10] observed that the ␦ phase guided crack to advance in a zigzag way along ␦ phase in grains and at grain boundaries by the in-situ tensile tests at the room temperature. However, in the DP process, to prevent the dissolution of ␦ phase and ensure the occurrence of dynamic recrystallization, the deformation temperature must be controlled between the dissolution temperature of ␦ phase and the dynamic recrystallization temperature. The results in Ref. [11] showed that the dynamic recrystallization temperature of Inconel 718 is higher than 850 ◦ C. Therefore, in the present work, for the DP process of Inconel 718 alloy, high temperature tensile test was carried out to study the effect of ␦ phase on the tensile deformation and fracture characteristics of delta-processed Inconel 718 at elevated temperatures. 2. Experimental procedures The material used in this study is cut from a commercially available Inconel 718 wrought bar with a diameter of 250 mm. The chemical compositions (in wt.%) of the alloy are as follows: C, 0.027; Ni, 53.74; Cr, 17.58; Nb, 5.35; Mo, 3.01; Ti, 0.98; Al, 0.52; Mn, 0.07; Si, 0.009; B, 0.0025; Fe, balance. There is a little difference in microstructure at the center, mid-radius and outer profile.
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Table 1 Heat treatments of the cylinder-specimens used in the tensile tests. Specimen No.
Heat treatment
␦ phase content (wt.%)
1 2
1040 ◦ C × 1 h water quenching 1040 ◦ C × 1 h water quenching + 900 ◦ C × 8 h water quenching 1040 ◦ C × 1 h water quenching + 900 ◦ C × 32 h water quenching
0 3.79
3
8.21
To minimize the scatter of the initial microstructure, the specimens with a diameter of 18 mm and a height of 60 mm, were cut from an outer diameter circle. To study the effect of ␦ phase content and morphology on the deformation and fracture characteristics, the specimens have been pre-precipitated ␦ phase by different heat treatments to obtain different initial ␦ phase contents, and the corresponding heat treatments and ␦ phase contents are shown in Table 1. The ␦ phase content is measured by the quantitative Xray diffraction technique and the details can be seen in Ref. [12]. Fig. 1 shows the microstructures of the specimens in Table 1. The microstructure of the specimen 1 (in Fig. 1a) is the matrix ␥ and the carbide NbC, and the average grain size is about 70.42 m. There are two morphologies of ␦ phase precipitated in the specimen 2 (in Fig. 1b). The short plate-like ␦ phase distributes in the grain boundaries and twin boundaries, and spherical ␦ phase particles distribute in the interior of grains. In Fig. 1c, the microstructure of the specimen 3, which was solution treated at 1040 ◦ C for 1 h and then aged at 900 ◦ C for 32 h, is the Widmanstätten ␦. Fig. 2a and b shows the shape and size of the tensile specimen and the tensile specimen after machined, respectively. The dissolution temperature of ␦ phase in Inconel 718 rises with the
increasing of the content of Nb [13], and as the temperature is higher than 980 ◦ C, the ␦ phase dissolves largely [14]. In addition, the dynamic recrystallization temperature of Inconel 718 is higher than 850 ◦ C [11]. Therefore, the high temperature tensile test was conducted at a temperature of 950 ◦ C. In the tensile test, the deformation speed employed the collet separation speed, in which a load of 200 MPa min−1 was exerted before the yield and then a speed of 2.8 mm min−1 was applied. The tensile tests at elevated temperatures were carried out on an AG-X 250 kN universal material testing machine. In the tensile tests at elevated temperatures, the specimens were put in the furnace at the room temperature and heated to the test temperature with the heating of furnace. The tensile tests were conducted in the furnace, and the specimens after deformation were cooled to the room temperature in the furnace. The deformed specimens were cleaned in the acetone by the ultrasonator. The morphology of the fracture was observed by a scanning electron microscope (SEM). Then, the deformed specimens were sectioned parallel to the tensile axis for microstructure analysis. The microstructure was observed by an optical microscope (OM), and the morphology of ␦ phase was observed by the SEM.
3. Results and discussion 3.1. Tensile deformation behaviors Fig. 3 shows the true stress–strain curves of the specimens with different initial ␦ contents deformed at 950 ◦ C. It can be seen that the tensile stress–strain curves of the specimens with different initial ␦ contents are all the typical elastic-uniform plastic curves, and there are three deformation stages during the deformation. The
Fig. 1. Microstructures of the specimens in the tensile tests obtained by different heat treatments: (a) 1040 ◦ C × 1 h water quenching (solution); (b) 1040 ◦ C × 1 h water quenching + 900 ◦ C × 8 h water quenching; (c) 1040 ◦ C × 1 h water quenching + 900 ◦ C × 32 h water quenching.
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Fig. 2. Dimensions of the tensile specimen (a) and tensile specimen after machined (b).
250 1: Solution treated withoutδ phase 2: pre-precipitated δ phase of 3.79% 3: pre-precipitated δ phase of 8.21%
True Stress , MPa
200
150
1
2
3
Table 2 Parameters of Hollomon equation for the specimens with different initial ␦ phase contents. Specimen No.
␦ phase content (wt.%)
n
K
1 2
0 3.79
3
8.21
0.375 0.403 0.065 0.494 0.101
560.82 731.38 207.93 1142.17 264.15
100
elastic deformation stage, the relationship between the true stress and true strain fits the Hooke’s law (1): 50
el = Eεel
0 0.0
0.2
0.4
0.6
0.8
1.0
True Strain Fig. 3. True stress–strain curves of the tensile deformed specimens at 950 ◦ C.
curves in the beginning stage of the deformation are straight line, which indicates the elastic deformation, and as the stress reaches the yield strength, the tensile curves are parabola, which indicates the uniform plastic deformation. Then, as the stress reaches the ultimate tensile strength, necking occurs, which indicates the nonuniform plastic deformation. In addition, there is a large and sharp stress dropping in the curve 1, which has also been found by Wang et al. [15]. According to Ref. [15], the reason for a large and sharp stress dropping in the curve 1 is the work softening. There is obvious work-hardening process for the specimens with different initial ␦ contents tensile deformed at 950 ◦ C. In the
where el and εel are the elastic stress and elastic strain, respectively, and E is the Young’s modulus. And in the uniform plastic deformation stage, the relationship between the true stress and true strain satisfies the Hollomon equation (2): p = Kεnp
3 5.2
where p and εp are the plastic stress and plastic strain, respectively, K is the strength factor, n is the strain hardening exponent. The strain hardening exponent n responses the strain hardening behavior of the material during the uniform plastic deformation. The value of n determines the maximum uniform plastic strain before the occurrence of necking, which is an important performance parameter during the tensile deformation. Because the strains in the elastic deformation stage of the specimens with different initial ␦ contents are very small, the elastic deformation can be neglected and the strain ε can be instead of the plastic εp in Eq. (2). Taking logarithm of the two sides of Eq. (2), it can be described as Eq. (3): ln = ln K + n ln ε
2
lnσ
5.1
5.0
1
4.9
1: Solution treated withoutδ phase 2: pre-precipitated δ phase of 3.79% 3: pre-precipitated δ phase of 8.21%
4.8 -4.0
(2)
(3)
where and ε are the true stress and true strain in the uniform plastic deformation stage, respectively. Fig. 4 shows the relationships between the ln and ln ε of the specimens with different initial ␦ contents in the uniform plastic
5.3
-4.2
(1)
-3.8
-3.6
lnε
-3.4
-3.2
-3.0
Fig. 4. ln versus ln ε during the uniform deformation.
-2.8
Fig. 5. Deformed specimens after the tensile deformation at 950 ◦ C.
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Fig. 6. Fracture morphologies of specimens with different initial ␦ phase contents at 950 ◦ C: (a) specimen 1, (b) specimen 2 and (c) specimen 3.
Fig. 7. Fracture morphologies and microstructures in longitudinal profile of specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C: (a) specimen 1, (b) specimen 2 and (c) specimen 3.
S.-H. Zhang et al. / Materials Science and Engineering A 528 (2011) 6253–6258
deformation stage. It can be clearly seen that the ln and ln ε of the solution treated specimen 1 meets a linear relationship. However, the relationships between the ln and ln ε of the delta-processed specimens 2 and 3 both satisfy two lines with different slopes, which has also been found in the tensile deformation of dual phase steel [16,17]. The feature with two lines relation indicates that there are two deformation processes in the uniform plastic deformation stage of the delta-processed specimens 2 and 3. The values of K and n in the Hollomon equation (2) for the specimens with different initial ␦ contents are shown in Table 2. The value of n for the solution treated specimen 1 is 0.375. For the deltaprocessed specimens 2 and 3, n has two values, and the value of n in the first deformation process is higher than that in the second deformation process. Furthermore, the values of n in the two deformation processes increase as the ␦ phase content increases. The value of n depends on the microstructures of the specimens closely. In Fig. 1, the microstructure of the solution treated specimen 1 is the matrix ␥ and a little of carbide NbC. During the tensile deformation of the specimen 1, the matrix ␥ is mainly deformed structure. For the delta-processed specimens 2 and 3, the microstructures are the mixture structures consisting of the matrix ␥ and the ␦ phase. Therefore, the deformation structures are different during the different deformation processes of the deltaprocessed specimens 2 and 3. As the ␦ phase is the hard brittle phase relative to the matrix ␥, the matrix ␥ is deformed preferentially during the tensile deformation. Because the crystal structure of matrix ␥ is the body-centered cubic, the main motion of dislocations in the matrix ␥ is the plane slipping when the strain is small, and the dislocation hardening is the main deformation mechanism. However, the moving dislocations are pinned in the vicinity of ␦ phase when meeting the ␦ phase, which can be confirmed by Ref. [5]. The amount of the piled dislocations increases with the ␦ phase content increasing, so the value of n increases. As the strain increases, the piled dislocations in the phase boundary between the matrix ␥ and ␦ phase would pass through the barrier and make the ␦ phase deform, which would lead to the relaxation of the stress concentration in the phase boundary and reduce the strain incompatibility between the matrix ␥ and ␦ phase. Therefore, the value of n in the second deformation process decreases. The elongation is an important index for the plastic deformation ability of the material, and the high elongation responses the good plasticity. Fig. 5 shows the deformed specimens after the tensile deformation at 950 ◦ C. It can be clearly seen that the elongation of the solution treated specimen 1 is the highest. The ␦ phase content of the specimens 1, 2 and 3 is 0, 3.79 and 8.21%, respectively. And the corresponding elongation is 173.6, 120 and 80%, respectively. This indicates that the plasticity of the specimens decreases with the ␦ phase content increasing. Therefore, the existence of the ␦ phase would reduce the plasticity of Inconel 718 alloy at elevated temperatures. 3.2. Fracture characteristics Fig. 6 shows the fracture morphologies of the specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C. The fracture morphologies of the specimens 1, 2 and 3 are all the typical equiaxed dimples. In addition, plastic deformation occurs in the specimens 1, 2 and 3 before the fracture. Therefore, the fracture mechanisms for the specimens with different initial ␦ phase contents are all the microvoid coalescence ductile fracture. The formation of micropores is an important process for the microvoid coalescence fracture, and the time of the micropores formation determines the plasticity of the material largely. Fig. 7 shows the fracture morphologies and microstructures in the longitudinal profile of the specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C. It can be seen
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that the micropore fracture trends to occur in the phase boundary between the matrix ␥ and ␦ phase or the carbide during the deformation. Therefore, the ␦ phase and carbide are the nucleuses for the formation of micropores, which is in line with that the micropores always form around the inclusion, second-phase and the boundaries between the matrix and the hard brittle phase. Furthermore, by the in-situ tensile test at room temperature, Zhang et al. [10] found that the ␦ phase could pin the glide bands generated during the deformation and played the guide in the crack propagation. Compared to the matrix ␥, the ␦ phase and carbide are hard brittle phases. During the deformation, the dislocations slip in the gliding plane of the matrix ␥ preferentially, and the existence of ␦ phase or carbide can impede the slipping of dislocations. Therefore, there is stress concentration in the vicinity of ␦ phase or carbide. As the stress reaches the bond strength of the phase boundaries between the matrix ␥ and ␦ phase or carbide, or the break limit of ␦ phase or carbide, the microspores fracture occurs. In addition, the dislocations are pinned in the vicinity of plate-like ␦ phase easily during the deformation, and the contact area between the matrix ␥ and plate-like ␦ phase is large. These would cause the higher tensile stress generated during the deformation, so the time of the micropores formation is earlier and the plasticity of the specimen is worse. Therefore, as the ␦ phase content increases, the high-temperature plasticity of the alloy reduces. In addition, there is the occurrence of dynamic recrystallization in the fractures of the specimens with different initial ␦ phase contents, which has also been observed in the tensile test of the solution treated Inconel 718 alloy at elevated temperature [15]. There are a few of spherical ␦ phase particles precipitating in the solution treated specimen 1. The initial plate-like ␦ phase has been disappeared in the delta-processed specimens 2 and 3, and spherical ␦ phase particles appear in the grain boundary of the recrystallized grains. The evolution of the ␦ phase during the deformation has been discussed in the Ref. [7]. 4. Conclusions (1) The tensile stress–strain curves of the specimens with different initial ␦ phase content were all the typical elastic-uniform plastic curves. (2) In the uniform plastic deformation stage, the strain hardening exponent n for the solution treated specimen 1 were 0.375, and there were two different values of n for the delta-processed specimens 2 and 3. In addition, the value of n in the first deformation process was higher than that in the second deformation process, and the value of n increased with the ␦ phase content increasing. (3) The elongation of the specimens with the initial ␦ phase contents of 0, 3.79 and 8.21% were 173.6, 124 and 80%, respectively, and as the ␦ phase content increases, the plasticity of Inconel 718 alloy at elevated temperatures reduces. (4) The tensile fracture mechanisms for the specimens with different initial ␦ phase contents were all microvoid coalescence ductile fracture, and the ␦ phase and carbide were the nucleuses for the formation of micropores. Acknowledgement This work has been supported by the National Natural Science Foundation of China with the Grant Number: 50834008. References [1] C. Ruiz, A. Obabueki, K. Gillespie, in: E.A. Loria (Ed.), Superalloys 1992, TMS, 1992, p. 33.
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[2] A.W. Dix, J.M. Hyzak, R.P. Singh, in: E.A. Loria (Ed.), Superalloys 1992, TMS, 1992, p. 23. [3] P.R. Bhowal, J.J. Schirra, in: E.A. Loria (Ed.), Superalloys 718, 625, 706 and Various Derivatives, TMS, 2001, p. 193. [4] D.Y. Cai, W.H. Zhang, P.L. Nie, et al., Materials Characterization 58 (2007) 220. [5] H. Yoshida, T. Hatta, T. Hironaka, et al., in: J.M.A. César, A.D. Santos (Eds.), Proceedings of the 9th International Conference on Numerical Methods in Industrial Processes, 2007, p. 987. [6] H.J. Lv, C.G. Yao, K.F. Zhang, X.C. Jia, Materials for Mechanical Engineering 27 (15) (2003) (in Chinese). [7] H. Yuan, W.C. Liu, Materials Science and Engineering A 408 (2005) 281. [8] Y. Wang, L. Zhen, W.Z. Shao, L. Yang, X.M. Zhang, Journal of Alloys and Compounds 44 (2009) 341. [9] Y. Desvallées, M. Bouzidi, F. Bois, et al., in: E.A. Loria (Ed.), Superalloys 718, 625, 706 and Various Derivatives, TMS, 1994, p. 281.
[10] Y. Zhang, X.B. Huang, Y. Wang, et al., in: E.A. Loria (Ed.), Superalloys 718, 625, 706 and Various Derivatives, TMS, 1997, p. 229. [11] D. Zhao, P.K. Chaudhury, in: E.A. Loria (Ed.), Superalloys 718, 625, 706 and Various Derivatives, TMS, 1994, p. 303. [12] H.Y. Zhang, S.H. Zhang, M. Cheng, Materials Characterization 61 (2010) 49. [13] R.E. Schafrik, D.D. Ward, J.R. Groh, in: E.A. Loria (Ed.), Superalloys 718, 625, 706 and Various Derivatives, TMS, 2001, p. 1. [14] J.P. Hu, Ph.D. Dissertation, Steel Research Institute, 1999 (in Chinese). [15] Y. Wang, W.Z. Shao, L. Zhen, C. Yang, X.M. Zhang, Journal of Alloys And Compounds 471 (2009) 331. [16] J.K. Liu, M. Jin, Z.H. Jiang, L.J. Zhang, Chinese Journal of Mechanical Engineering 26 (71) (1990) (in Chinese). [17] M.H. Cai, H. Ding, J.S. Zhang, L. Li, Z.Y. Tang, Chinese Journal of Materials Research 23 (83) (2009) (in Chinese).
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Tensile deformation and fracture characteristics of delta-processed Inconel 718 alloy at elevated temperature Shi-Hong Zhang ∗ , Hai-Yan Zhang, Ming Cheng Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
a r t i c l e
i n f o
Article history: Received 17 December 2010 Received in revised form 21 April 2011 Accepted 22 April 2011 Available online 30 April 2011 Keywords: Inconel 718 Delta-processed ␦ phase Tensile deformation Fracture
a b s t r a c t Three specimens with different initial ␦ phase contents have been used to study the tensile deformation and fracture characteristics of the delta-processed Inconel 718 alloy at elevated temperatures by tensile tests at 950 ◦ C. The results indicated that the tensile stress–strain curves of the three specimens were the elastic-uniform plastic curves, and there were two deformation processes during the uniform plastic deformation stage of the delta-processed specimens. The strain hardening exponent in the first deformation process was higher than that in the second process, and the value of the strain hardening exponent increased as the content of ␦ phase increased. In addition, the plasticity of the specimen at elevated temperatures decreased as the content of ␦ phase increased. However, the fracture mechanisms for all the specimens were all microvoid coalescence ductile fracture, and the ␦ phase and carbide were the nucleuses for the formation of micropores. The elongation for the specimen with 8.21% preprecipitated ␦ phase was still as high as 80%. Thus, delta-processed Inconel 718 alloy presented excellent high-temperature plasticity. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Inconel 718 is an important material used for aviation and aerospace engines because of its excellent mechanical properties in the temperature range from −253 ◦ C to 650 ◦ C. In order to improve the safety and reliability of engines, it is crucial to obtain the forgings with uniform and fine microstructures. Generally, the forgings with large size and complex shape, such as turbine disks and engine shafts, are manufactured by multi-stage hot working processes. In addition, the microstructure of the alloy is sensitive to the hot deformation parameters. Therefore, the defects of coarse grain or duplex grain often appear in the forgings. As the ␦ phase in Inconel 718 can control grain size through the strong pinning effect, the Delta Process (DP) has been applied to the forging of Inconel 718, which uses an intentional ␦ phase precipitation cycle and subsequent thermomechanical processing to produce uniform fine grain billet and bar stock [1]. A large size billet of Inconel 718 with the average grain size of ASTM 8 and a turbine disk forging with grain size of ASTM 11, have been obtained by the DP process [1–3]. The ␦ phase plays a significant role in the microstructure evolution and mechanical properties for Inconel 718 [4–10]. Yuan [7] and Wang [8] investigated the effect of ␦ phase on the deforma-
∗ Corresponding author. Fax: +86 24 23906831. E-mail address: [email protected] (Shi-Hong Zhang). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.04.074
tion behaviors of delta-processed Inconel 718, and their results showed that the ␦ phase can stimulate the occurrence of dynamic recrystallization. Desvallées et al. [9] found that the ␦ phase had a detrimental effect on the yield strength but no effect on the ultimate tensile strength at the room temperature and 550 ◦ C. Zhang et al. [10] observed that the ␦ phase guided crack to advance in a zigzag way along ␦ phase in grains and at grain boundaries by the in-situ tensile tests at the room temperature. However, in the DP process, to prevent the dissolution of ␦ phase and ensure the occurrence of dynamic recrystallization, the deformation temperature must be controlled between the dissolution temperature of ␦ phase and the dynamic recrystallization temperature. The results in Ref. [11] showed that the dynamic recrystallization temperature of Inconel 718 is higher than 850 ◦ C. Therefore, in the present work, for the DP process of Inconel 718 alloy, high temperature tensile test was carried out to study the effect of ␦ phase on the tensile deformation and fracture characteristics of delta-processed Inconel 718 at elevated temperatures. 2. Experimental procedures The material used in this study is cut from a commercially available Inconel 718 wrought bar with a diameter of 250 mm. The chemical compositions (in wt.%) of the alloy are as follows: C, 0.027; Ni, 53.74; Cr, 17.58; Nb, 5.35; Mo, 3.01; Ti, 0.98; Al, 0.52; Mn, 0.07; Si, 0.009; B, 0.0025; Fe, balance. There is a little difference in microstructure at the center, mid-radius and outer profile.
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S.-H. Zhang et al. / Materials Science and Engineering A 528 (2011) 6253–6258
Table 1 Heat treatments of the cylinder-specimens used in the tensile tests. Specimen No.
Heat treatment
␦ phase content (wt.%)
1 2
1040 ◦ C × 1 h water quenching 1040 ◦ C × 1 h water quenching + 900 ◦ C × 8 h water quenching 1040 ◦ C × 1 h water quenching + 900 ◦ C × 32 h water quenching
0 3.79
3
8.21
To minimize the scatter of the initial microstructure, the specimens with a diameter of 18 mm and a height of 60 mm, were cut from an outer diameter circle. To study the effect of ␦ phase content and morphology on the deformation and fracture characteristics, the specimens have been pre-precipitated ␦ phase by different heat treatments to obtain different initial ␦ phase contents, and the corresponding heat treatments and ␦ phase contents are shown in Table 1. The ␦ phase content is measured by the quantitative Xray diffraction technique and the details can be seen in Ref. [12]. Fig. 1 shows the microstructures of the specimens in Table 1. The microstructure of the specimen 1 (in Fig. 1a) is the matrix ␥ and the carbide NbC, and the average grain size is about 70.42 m. There are two morphologies of ␦ phase precipitated in the specimen 2 (in Fig. 1b). The short plate-like ␦ phase distributes in the grain boundaries and twin boundaries, and spherical ␦ phase particles distribute in the interior of grains. In Fig. 1c, the microstructure of the specimen 3, which was solution treated at 1040 ◦ C for 1 h and then aged at 900 ◦ C for 32 h, is the Widmanstätten ␦. Fig. 2a and b shows the shape and size of the tensile specimen and the tensile specimen after machined, respectively. The dissolution temperature of ␦ phase in Inconel 718 rises with the
increasing of the content of Nb [13], and as the temperature is higher than 980 ◦ C, the ␦ phase dissolves largely [14]. In addition, the dynamic recrystallization temperature of Inconel 718 is higher than 850 ◦ C [11]. Therefore, the high temperature tensile test was conducted at a temperature of 950 ◦ C. In the tensile test, the deformation speed employed the collet separation speed, in which a load of 200 MPa min−1 was exerted before the yield and then a speed of 2.8 mm min−1 was applied. The tensile tests at elevated temperatures were carried out on an AG-X 250 kN universal material testing machine. In the tensile tests at elevated temperatures, the specimens were put in the furnace at the room temperature and heated to the test temperature with the heating of furnace. The tensile tests were conducted in the furnace, and the specimens after deformation were cooled to the room temperature in the furnace. The deformed specimens were cleaned in the acetone by the ultrasonator. The morphology of the fracture was observed by a scanning electron microscope (SEM). Then, the deformed specimens were sectioned parallel to the tensile axis for microstructure analysis. The microstructure was observed by an optical microscope (OM), and the morphology of ␦ phase was observed by the SEM.
3. Results and discussion 3.1. Tensile deformation behaviors Fig. 3 shows the true stress–strain curves of the specimens with different initial ␦ contents deformed at 950 ◦ C. It can be seen that the tensile stress–strain curves of the specimens with different initial ␦ contents are all the typical elastic-uniform plastic curves, and there are three deformation stages during the deformation. The
Fig. 1. Microstructures of the specimens in the tensile tests obtained by different heat treatments: (a) 1040 ◦ C × 1 h water quenching (solution); (b) 1040 ◦ C × 1 h water quenching + 900 ◦ C × 8 h water quenching; (c) 1040 ◦ C × 1 h water quenching + 900 ◦ C × 32 h water quenching.
S.-H. Zhang et al. / Materials Science and Engineering A 528 (2011) 6253–6258
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Fig. 2. Dimensions of the tensile specimen (a) and tensile specimen after machined (b).
250 1: Solution treated withoutδ phase 2: pre-precipitated δ phase of 3.79% 3: pre-precipitated δ phase of 8.21%
True Stress , MPa
200
150
1
2
3
Table 2 Parameters of Hollomon equation for the specimens with different initial ␦ phase contents. Specimen No.
␦ phase content (wt.%)
n
K
1 2
0 3.79
3
8.21
0.375 0.403 0.065 0.494 0.101
560.82 731.38 207.93 1142.17 264.15
100
elastic deformation stage, the relationship between the true stress and true strain fits the Hooke’s law (1): 50
el = Eεel
0 0.0
0.2
0.4
0.6
0.8
1.0
True Strain Fig. 3. True stress–strain curves of the tensile deformed specimens at 950 ◦ C.
curves in the beginning stage of the deformation are straight line, which indicates the elastic deformation, and as the stress reaches the yield strength, the tensile curves are parabola, which indicates the uniform plastic deformation. Then, as the stress reaches the ultimate tensile strength, necking occurs, which indicates the nonuniform plastic deformation. In addition, there is a large and sharp stress dropping in the curve 1, which has also been found by Wang et al. [15]. According to Ref. [15], the reason for a large and sharp stress dropping in the curve 1 is the work softening. There is obvious work-hardening process for the specimens with different initial ␦ contents tensile deformed at 950 ◦ C. In the
where el and εel are the elastic stress and elastic strain, respectively, and E is the Young’s modulus. And in the uniform plastic deformation stage, the relationship between the true stress and true strain satisfies the Hollomon equation (2): p = Kεnp
3 5.2
where p and εp are the plastic stress and plastic strain, respectively, K is the strength factor, n is the strain hardening exponent. The strain hardening exponent n responses the strain hardening behavior of the material during the uniform plastic deformation. The value of n determines the maximum uniform plastic strain before the occurrence of necking, which is an important performance parameter during the tensile deformation. Because the strains in the elastic deformation stage of the specimens with different initial ␦ contents are very small, the elastic deformation can be neglected and the strain ε can be instead of the plastic εp in Eq. (2). Taking logarithm of the two sides of Eq. (2), it can be described as Eq. (3): ln = ln K + n ln ε
2
lnσ
5.1
5.0
1
4.9
1: Solution treated withoutδ phase 2: pre-precipitated δ phase of 3.79% 3: pre-precipitated δ phase of 8.21%
4.8 -4.0
(2)
(3)
where and ε are the true stress and true strain in the uniform plastic deformation stage, respectively. Fig. 4 shows the relationships between the ln and ln ε of the specimens with different initial ␦ contents in the uniform plastic
5.3
-4.2
(1)
-3.8
-3.6
lnε
-3.4
-3.2
-3.0
Fig. 4. ln versus ln ε during the uniform deformation.
-2.8
Fig. 5. Deformed specimens after the tensile deformation at 950 ◦ C.
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Fig. 6. Fracture morphologies of specimens with different initial ␦ phase contents at 950 ◦ C: (a) specimen 1, (b) specimen 2 and (c) specimen 3.
Fig. 7. Fracture morphologies and microstructures in longitudinal profile of specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C: (a) specimen 1, (b) specimen 2 and (c) specimen 3.
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deformation stage. It can be clearly seen that the ln and ln ε of the solution treated specimen 1 meets a linear relationship. However, the relationships between the ln and ln ε of the delta-processed specimens 2 and 3 both satisfy two lines with different slopes, which has also been found in the tensile deformation of dual phase steel [16,17]. The feature with two lines relation indicates that there are two deformation processes in the uniform plastic deformation stage of the delta-processed specimens 2 and 3. The values of K and n in the Hollomon equation (2) for the specimens with different initial ␦ contents are shown in Table 2. The value of n for the solution treated specimen 1 is 0.375. For the deltaprocessed specimens 2 and 3, n has two values, and the value of n in the first deformation process is higher than that in the second deformation process. Furthermore, the values of n in the two deformation processes increase as the ␦ phase content increases. The value of n depends on the microstructures of the specimens closely. In Fig. 1, the microstructure of the solution treated specimen 1 is the matrix ␥ and a little of carbide NbC. During the tensile deformation of the specimen 1, the matrix ␥ is mainly deformed structure. For the delta-processed specimens 2 and 3, the microstructures are the mixture structures consisting of the matrix ␥ and the ␦ phase. Therefore, the deformation structures are different during the different deformation processes of the deltaprocessed specimens 2 and 3. As the ␦ phase is the hard brittle phase relative to the matrix ␥, the matrix ␥ is deformed preferentially during the tensile deformation. Because the crystal structure of matrix ␥ is the body-centered cubic, the main motion of dislocations in the matrix ␥ is the plane slipping when the strain is small, and the dislocation hardening is the main deformation mechanism. However, the moving dislocations are pinned in the vicinity of ␦ phase when meeting the ␦ phase, which can be confirmed by Ref. [5]. The amount of the piled dislocations increases with the ␦ phase content increasing, so the value of n increases. As the strain increases, the piled dislocations in the phase boundary between the matrix ␥ and ␦ phase would pass through the barrier and make the ␦ phase deform, which would lead to the relaxation of the stress concentration in the phase boundary and reduce the strain incompatibility between the matrix ␥ and ␦ phase. Therefore, the value of n in the second deformation process decreases. The elongation is an important index for the plastic deformation ability of the material, and the high elongation responses the good plasticity. Fig. 5 shows the deformed specimens after the tensile deformation at 950 ◦ C. It can be clearly seen that the elongation of the solution treated specimen 1 is the highest. The ␦ phase content of the specimens 1, 2 and 3 is 0, 3.79 and 8.21%, respectively. And the corresponding elongation is 173.6, 120 and 80%, respectively. This indicates that the plasticity of the specimens decreases with the ␦ phase content increasing. Therefore, the existence of the ␦ phase would reduce the plasticity of Inconel 718 alloy at elevated temperatures. 3.2. Fracture characteristics Fig. 6 shows the fracture morphologies of the specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C. The fracture morphologies of the specimens 1, 2 and 3 are all the typical equiaxed dimples. In addition, plastic deformation occurs in the specimens 1, 2 and 3 before the fracture. Therefore, the fracture mechanisms for the specimens with different initial ␦ phase contents are all the microvoid coalescence ductile fracture. The formation of micropores is an important process for the microvoid coalescence fracture, and the time of the micropores formation determines the plasticity of the material largely. Fig. 7 shows the fracture morphologies and microstructures in the longitudinal profile of the specimens with different initial ␦ phase contents after the tensile deformation at 950 ◦ C. It can be seen
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that the micropore fracture trends to occur in the phase boundary between the matrix ␥ and ␦ phase or the carbide during the deformation. Therefore, the ␦ phase and carbide are the nucleuses for the formation of micropores, which is in line with that the micropores always form around the inclusion, second-phase and the boundaries between the matrix and the hard brittle phase. Furthermore, by the in-situ tensile test at room temperature, Zhang et al. [10] found that the ␦ phase could pin the glide bands generated during the deformation and played the guide in the crack propagation. Compared to the matrix ␥, the ␦ phase and carbide are hard brittle phases. During the deformation, the dislocations slip in the gliding plane of the matrix ␥ preferentially, and the existence of ␦ phase or carbide can impede the slipping of dislocations. Therefore, there is stress concentration in the vicinity of ␦ phase or carbide. As the stress reaches the bond strength of the phase boundaries between the matrix ␥ and ␦ phase or carbide, or the break limit of ␦ phase or carbide, the microspores fracture occurs. In addition, the dislocations are pinned in the vicinity of plate-like ␦ phase easily during the deformation, and the contact area between the matrix ␥ and plate-like ␦ phase is large. These would cause the higher tensile stress generated during the deformation, so the time of the micropores formation is earlier and the plasticity of the specimen is worse. Therefore, as the ␦ phase content increases, the high-temperature plasticity of the alloy reduces. In addition, there is the occurrence of dynamic recrystallization in the fractures of the specimens with different initial ␦ phase contents, which has also been observed in the tensile test of the solution treated Inconel 718 alloy at elevated temperature [15]. There are a few of spherical ␦ phase particles precipitating in the solution treated specimen 1. The initial plate-like ␦ phase has been disappeared in the delta-processed specimens 2 and 3, and spherical ␦ phase particles appear in the grain boundary of the recrystallized grains. The evolution of the ␦ phase during the deformation has been discussed in the Ref. [7]. 4. Conclusions (1) The tensile stress–strain curves of the specimens with different initial ␦ phase content were all the typical elastic-uniform plastic curves. (2) In the uniform plastic deformation stage, the strain hardening exponent n for the solution treated specimen 1 were 0.375, and there were two different values of n for the delta-processed specimens 2 and 3. In addition, the value of n in the first deformation process was higher than that in the second deformation process, and the value of n increased with the ␦ phase content increasing. (3) The elongation of the specimens with the initial ␦ phase contents of 0, 3.79 and 8.21% were 173.6, 124 and 80%, respectively, and as the ␦ phase content increases, the plasticity of Inconel 718 alloy at elevated temperatures reduces. (4) The tensile fracture mechanisms for the specimens with different initial ␦ phase contents were all microvoid coalescence ductile fracture, and the ␦ phase and carbide were the nucleuses for the formation of micropores. Acknowledgement This work has been supported by the National Natural Science Foundation of China with the Grant Number: 50834008. References [1] C. Ruiz, A. Obabueki, K. Gillespie, in: E.A. Loria (Ed.), Superalloys 1992, TMS, 1992, p. 33.
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