Tensile deformation behavior of superalloy 718 at elevated temperatures

Tensile deformation behavior of superalloy 718 at elevated temperatures

Journal of Alloys and Compounds 471 (2009) 331–335 Contents lists available at ScienceDirect Journal of Alloys and Compounds journal homepage: www.e...

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Journal of Alloys and Compounds 471 (2009) 331–335

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom

Tensile deformation behavior of superalloy 718 at elevated temperatures Y. Wang, W.Z. Shao, L. Zhen ∗ , C. Yang, X.M. Zhang School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China

a r t i c l e

i n f o

Article history: Received 2 November 2007 Received in revised form 21 March 2008 Accepted 23 March 2008 Available online 8 May 2008 Keywords: High temperature alloys Structure Strain

a b s t r a c t Uniaxial tensile deformation behavior of superalloy 718 at various temperatures and strain rates was studied. The flow stress decreases with increasing temperature and decreasing strain rate. The flow oscillations in stress at lower strain rates are attributed to the occurrence of cyclic dynamic recrystallization (DRX). The activation energy for superalloy 718 under hot tensile deformation is determined to be 450.2 kJ/mol. The undissolved delta phases play an important role in the elongation behavior. The existence of a certain amount of undissolved particle-like delta phases can promote the DRX process as well as suppress the initiation and propagation of grain boundary cracks, leading to the increase of the elongations. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Nickel-base superalloy 718 has been one of the most frequently used alloys in aircraft engines because of its good mechanical properties at elevated temperatures up to 650 ◦ C. With the development of modern aeroengines, advances in design of the critical components require that the properties of a material be used to their full extent. Better control of the thermomechanical processing is of paramount importance to obtain superior performances for superalloy 718. In the past several decades, the hot working characteristics of superalloy 718 have been extensively studied using hot torsion or hot compression experiments. Different types of models have been proposed to predict the flow stress as well as the grain size and recrystallization fraction [1–9]. Meanwhile, there has been considerable interest in examining the possibility of superplasticity in superalloys, and the superplasticity has now been reported in the temperature range of 900–1020 ◦ C for superalloy 718 with finer initial grain size [10–12]. However, little understanding is available on the fundamental tensile deformation and fracture behavior of superalloy 718 at elevated temperatures. Uniaxial tension test is the most common mechanical test since it can not only provide the strength properties of the material but also the ductility data. The flow stress and fracture behavior obtained from uniaxial tension is valuable for estimating the effects of various factors on the formability under more complex loading conditions. In the present study, uniaxial tension tests were performed on superalloy 718 with coarse grain size at various strain rates and temperatures. The tensile deformation behavior was stud-

ied to contribute to the database of tensile properties for superalloy 718 as well as to provide better understanding of the effects of temperature and strain rate on its fracture behavior. Efforts were also made to analyze the correlation between deformation microstructure and tensile properties. 2. Experimental procedures The chemical compositions (wt.%) of superalloy 718 used in this investigation are as follows: Cr, 18.09; Fe, 17.69; Nb + Ta, 5.43; Mo, 3.07; Ti, 0.97; Al, 0.46; Co, 0.18; Si, 0.078; Mn, 0.065; Cu, 0.065; C, 0.040; S < 0.001; P < 0.007; and Ni, balance. Button-head tensile specimens 25 mm in gauge length and 5 mm in diameter were machined from an as-received wrought billet with the tensile axes parallel to the pressing direction. The specimens were solution treated for 30 min at 1100 ◦ C and then mechanically polished before tensile tests. The microstructure of the annealed superalloy 718 consists of equiaxed grains with an average size of 176 ␮m and lamella-like straight annealing twins, as seen from Fig. 1. The uniaxial tensile tests were carried out in a screw-driven universal testing machine at temperatures ranging from 950 ◦ C to 1050 ◦ C with strain rates of 5 × 10−4 s−1 to 10−2 s−1 until failure. Prior to the tensile test, all specimens were put inside the furnace at the temperature of 900 ◦ C and then heated to the selected testing temperature holding for 15 min in order to stabilize the structure. Typical samples were sectioned near the fracture of the tensile specimens, mechanically polished and etched electrolytically with a solution consisting of 13% HF, 7% HNO3 and 80% HCl for optical metallographic examination. The foils for Transmission Electron Microscope (TEM) examination were prepared firstly by hand grinding from a thickness of 0.3–0.05 mm and then thinned using a twin-jet technique in the electrolyte of 10% solution of HClO4 in ethanol. TEM observation was performed in a Philips TENCAI-20 microscope operated at 200 kV.

3. Results and discussion 3.1. Flow stress behavior

∗ Corresponding author. Tel.: +86 451 86412133; fax: +86 451 86413922. E-mail address: [email protected] (L. Zhen). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.03.082

Typical true stress–true strain curves for superalloy 718 obtained from the uniaxial tension tests at various temperatures

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Fig. 3. Effect of strain rate on the true stress–true strain curves of superalloy 718 obtained at a temperature of 980 ◦ C. Fig. 1. Optical microstructure for the initial annealed superalloy 718.

are shown in Fig. 2. The increase in the deformation temperature results in a decrease in the flow stress. At the temperatures higher than 950 ◦ C, the flow curves show rapid work hardening at an initial stage of deformation followed by work softening at a small strain and then change gradually with a small stress fluctuation until final fracture. Such features of the flow curves without clear peak stresses differ from the relevant reports for superalloy 718 compressed in the same deformation conditions [1–5], which is closely related to the difference in stress orientation. Fig. 3 shows the flow curves of the alloy deformed at 980 ◦ C at strain rates ranging from 5 × 10−4 s−1 to 10−2 s−1 . It is obvious that the flow stress decreases with the decrease of strain rate. However, the configuration of the tensile flow curves is quite different at different strain rates. The curve corresponding to lower strain rate shows the formation of regular oscillations in stress during the work softening stage. The flow curve oscillates almost at a constant level at the strain rate of 5 × 10−4 s−1 , as seen from Fig. 3. Different mechanisms can lead to the appearance of flow oscillations, such as dynamic recrystallization (DRX), dynamic strain aging (DSA) and flow localization [13]. Generally, DRX is believed to be the main mechanism during high temperature deformation of superalloy 718, which seems to be responsible for the flow oscillations in this study. But DSA can also occur during

Fig. 2. Tensile true stress–true strain curves of superalloy 718 obtained at various temperatures with a constant strain rate of 0.01 s−1 .

Fig. 4. Dependence of the strain rate on the peak stress at various temperatures.

monotonic tensile tests of superalloy 718, as noted by Hale et al. [14], at a low strain rate at the temperature ranging from 200 ◦ C to 650 ◦ C. Thomas et al. [9] also found the phenomenon of serrated flow in stress versus strain curves due to DSA appeared at the temperature ranging from 900 ◦ C to 1050 ◦ C during hot compression of superalloy 718. Therefore, it is difficult to determine the mechanisms of the flow oscillations only according to the shape of the stress–strain curves, and detailed discussion is still required.

Fig. 5. Variation of elongation to fracture as a function of temperature.

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Fig. 6. Typical microstructures of superalloy 718 deformed at 5 × 10−4 s−1 . (a) 950 ◦ C, (b) and (c) 980 ◦ C, (d) 1000 ◦ C, and (e) 1050 ◦ C.

3.2. Stress exponent and activation energy The stress exponent (n) can be calculated according to the following Eq. (1) [15]: n=

d ln ε˙ d ln 

(1)

In order to derive the stress exponent, the dependence of strain rate on the peak stress was plotted at each tensile temperature on a log–log scale (Fig. 4). It may be noted that the regressive straight line can be plotted at every temperature, but the slope is not quite similar at every set of tests. The stress exponent obtained at four temperatures, 950 ◦ C, 980 ◦ C, 1000 ◦ C and 1050 ◦ C are 5.3, 4.7, 5.7

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and 4.7, respectively. The higher stress exponent appears at the temperature of 1000 ◦ C and the exponent value at the temperature of 980 ◦ C is the same as that at 1050 ◦ C. It was reported [9,16] that an exponent value close to 5 can be obtained when the deformation is controlled by the glide and climb of dislocations. In the present case, the deviation of the stress exponent at 1000 ◦ C is an indication that some strain hardening is induced at the deformation temperature and the additional mechanism is acting on the deformation behavior. The relationship between the peak stress and the deformation parameters for superalloy 718 can be expressed by a hyperbolicsine Arrhenius-type function [17]: ε˙ = A[sinh(˛P )]n exp

 −Q  RT

(2)

where A (s−1 ), ˛ (MPa−1 ) and n are constants, Q, the activation energy of deformation (J/mol), R, the universal gas constant, T, the ˙ the strain rate (s−1 ) and  p deformation temperature in Kelvin, ε, the peak stress (MPa). According to the data obtained from the hot tensile experiments (Fig. 4), the value of Q can be obtained by fitting the experimental data and regression analysis. The activation energy for superalloy 718 under hot tension is determined to be 450.2 kJ/mol, which is comparable to the previously reported value about superalloy 718 under hot compression [4,8]. This value is well above that for the self-diffusion of nickel (278.2 kJ/mol [11]), which is believed to attribute to the occurrence of DRX [18]. 3.3. Tensile elongation Elongation values obtained at various temperatures and strain rates are presented in Fig. 5. It may be noted that the high temperature ductility of superalloy 718 is strongly dependent on the deformation parameters at which the tests were conducted. The elongation increases with the decrease of strain rate and the maximum of elongation appears at the strain rate of 5 × 10−4 s−1 . As we know, DSA is a time dependent strengthening or hardening phenomenon, which manifests itself in the form of serrated stress–strain curves, yield stress plateaus, reduced tensile elongation, and so on. Therefore, the possibility of the flow oscillations caused by DSA can be excluded due to the higher elongations at the strain rate of 5 × 10−4 s−1 . The elongation behavior at lower strain rate can be preliminary attributed to the flow oscillations deriving from the cyclic DRX. Corresponding to the lower stress exponent, the elongation at each strain rate exhibits a maximum at the temperature of 980 ◦ C, which is believed to be concerned with the effect of the undissolved delta phases precipitated during the heating and holding periods on the tensile deformation behavior. 3.4. Deformation microstructure Typical microstructures of the specimens tested under different tensile conditions are shown in Fig. 6. It can be seen that a certain amount of short needle-shaped precipitates are distributed along the grain boundaries or inside the grains at the temperature of 950 ◦ C (Fig. 6a), which are determined to be delta phases. The delta phases were precipitated during heating and holding periods due to their wide range of precipitation temperature (780 ◦ C–980 ◦ C) [19]. The size and fraction of these precipitates observed from Fig. 6a also suggests that the precipitation process is favored by deformation mechanism. There is no apparent change for the original grains from the optical microscope observation, indicating that dynamic recovery is the main softening mechanism of the material deformed at the temperature of 950 ◦ C. At the temperature of 980 ◦ C, small recrystallized grains have already developed along some initial grain boundaries (Fig. 6b).

Fig. 7. TEM micrograph of sample deformed at 980 ◦ C and 5 × 10−4 s−1 .

The oscillations observed in superalloy 718 of the present study, deformed at 5 × 10−4 s−1 should indeed be attributed to the dynamic recrystallization process. There are a few undissolved particle-like delta phases distributed along the orginal grain boundaries, as seen from Fig. 6c. A bright field TEM micrograph, showing the delta particles and dislocations for sample deformed at 980 ◦ C and 5 × 10−4 s−1 is given in Fig. 7. It may be noted that high density dislocations are tangled and stored at the delta particles showed by a white arrow as an example. It is understandable that high density dislocations stored at the vicinity of the delta particles were the source of driving force of DRX, leading to the occurrence of particle stimulated nucleation (PSN) of recrystallization [20]. Therefore, besides the original grain boundaries, the undissolved delta phases at the temperature of 980 ◦ C could as well serve as the potential nucleation centres where the growth of recrystallized grains starts promoting the DRX process. Meanwhile, it was reported [21,22] that the existence of certain size and density of globular delta phases at grain boundaries can block intergranular crack propagation. The ductile region near delta phases due to the existence of denuded zone of ␥ can relax the stress concentration at grain boundaries, then to retard grain boundary crack initiation and propagation. Thus, the existence of the certain amount of undissolved particle-like delta phases should be responsible for the increase of elongation as well as the decrease of stress exponent at the temperature of 980 ◦ C. Fig. 6d shows the optical micrograph of a sample deformed at 1000 ◦ C and 5 × 10−4 s−1 . The amount of new recrystallized grains seems to be less than that in the sample deformed at 980 ◦ C. It may be noted that the undissolved delta phases still exist due to the temperature not exceeding its complete solvus (about 1038 ◦ C), but its size and amount are smaller than that at 980 ◦ C. Let us believe that at the higher exponent at the temperature of 1000 ◦ C some flow hardening existed, resulting in the lower elongations to fracture. Two reasons can be considered to explain the hardening in this alloy: First, smaller size and amount of undissolved delta particles probably promote some hardening, and second the flow localization or plastic instability occurs at the temperature of 1000 ◦ C. Further investigations are still necessary for providing the evidences of the explanation for the tensile behavior at the temperature of 1000 ◦ C. The results from this study suggested that the delta

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phases played an important role in the hot deformation behavior of superalloy 718, which was quite dependent on its size and amount. Although a large amount of DRX structure is obtained at the temperature of 1050 ◦ C (Fig. 6e), the coarsening of dynamic recrystallized grains after DRX lead to the elongations equal to or even lower than that at the temperature of 980 ◦ C. 4. Conclusions The hot tensile deformation behavior of superalloy 718 was investigated at temperatures from 950 ◦ C to 1100 ◦ C and strain rates from 5 × 10−4 s−1 to 10−2 s−1 . The temperature and strain rate dependence of flow stress as well as elongation were obtained in the present study. The cyclic DRX leads to the occurrence of flow stress oscillations at lower strain rates. The stress exponent changes irregularly with the increase of temperature. The elongation behavior at various temperatures is closely related to the role of the undissolved delta phases in the hot deformation, which is quite dependent on the size and amount of the particles. The activation energy (Q) of deformation for superalloy 718 under hot tension is determined to be 450.2 kJ/mol, which is comparable to the value of superalloy 718 under hot compression. References [1] M.J. Weis, M.C. Mataya, S.W. Thompson, Superalloy 718—Metallurgy and Applications, The Minerals, Metals and Materials Society, Warrendale, PA, 1989, pp. 135–154.

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[2] R. Srinivasan, V. Ramnarayan, U. Deshpande, Metall. Trans. 24A (1993) 2061–2069. [3] L.X. Zhou, T.N. Baker, Mater. Sci. Eng. A 177 (1994) 1–9. [4] C.I. Garcia, G.D. Wang, D.E. Camus, E.A. Loria, A.J. DeArdo, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals and Materials Society, Warrendale, PA, 1994, pp. 293–302. [5] J.M. Zhang, Z.Y. Gao, J.Y. Zhuang, Met. Mater. Trans. 30A (1999) 2701–2713. [6] S.C. Medeiros, Y.V.R.K. Prasad, W.G. Frazier, R. Srinivasan, Mater. Sci. Eng. A293 (2000) 198–207. [7] S.C. Medeiros, Y.V.R.K. Prasad, W.G. Frazier, N. Srinivasan, Scripta Mater. 42 (2000) 17–23. [8] H. Yuan, W.C. Liu, Mater. Sci. Eng. A408 (2005) 281–289. [9] A. Thomas, M. El-Wahabi, J.M. Cabrera, J.M. Prado, J. Mats. Proc. Tech. 177 (2006) 469–472. [10] G.D. Smith, S.R. Gregory, Y. Ma, Y. Li, T.G. Langdon, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals and Materials Society, Warrendale, PA, 1997, pp. 303–314. [11] B.P. Kashyap, M.C. Chaturvedi, Scripta Mater. 43 (2000) 429–433. [12] X. Han, L.J. Wu, H. Xia, R.G. Liu, S.G. Wang, Z.L. Chen, J. Mats. Proc. Tech. 137 (2003) 17–20. ˜ O.A. Ruano, Scripta Mater. 38 (1998) 1717–1723. [13] M. Eddahbi, F. Carreno, [14] C.L. Hale, W.S. Rollings, M.L. Weaver, Mater. Sci. Eng. A300 (2001) 153–164. [15] D.H. Kim, J.H. Kim, J.W. Sa, Y.S. Lee, C.K. Park, S.I. Moon, Mater. Sci. Eng. A 483–484 (2008) 262–265. [16] H.J. Frost, M.F. Ashby, Deformation—Mechanism Maps. The Plasticity and Creep of Metals and Ceramics, Pergamon Press, Oxford, 1982, pp. 54–55. [17] C.M. Sellars, W.J. Tegart, Membr. Sci. Rev. Met. 63 (1966) 731–735. [18] L. Briottet, J.J. Jonas, F. Montheillet, Acta Mater. 44 (1996) 1665–1672. [19] M. Sundararaman, P. Mukhopadhyay, S. Banerjee, Metall. Trans. 19A (1988) 453–465. [20] F.J. Humphreys, P.N. Kalu, Acta Metall. 35 (1987) 2815–2829. [21] S.Q. Li, J.Y. Zhuang, J.Y. Yang, Q. Deng, J.H. Du, Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals and Materials Society, Warrendale, PA, 1994, pp. 545–555. [22] W. Chen, M.C. Chaturvedi, Acta Mater. 45 (1997) 2735–2746.