Influence of thermal shock behavior on microstructure and mechanical properties of IN718 superalloy

Applied Surface Science 484 (2019) 1282–1287

Contents lists available at ScienceDirect

Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc

Full length article

Influence of thermal shock behavior on microstructure and mechanical properties of IN718 superalloy

T

Shu-liang Wanga, Yan-rong Suna,b, Li-jing Dua, Li Liua, Jie Wanga, , Ding-han Xiangc ⁎

a

School of Materials Science and Engineering, Southwest Petroleum University, Xindu Avenue 8#, Sichuan 610500, China CNBM (Chengdu) Optoelectronic Materials Co., Ltd., Sichuan 610500, China c Guangxi Key Laboratory of Information Materials, Guilin University of Electronic Technology, Guilin 541004, China b

ARTICLE INFO

ABSTRACT

Keywords: Thermal shock Microstructure Mechanical properties Oxide films δ phase

The influence of thermal shock behavior between 960 °C and room temperature on microstructure and tensile properties of IN718 alloy after three heat treatments was investigated in this paper. The experiments show that protective and dense oxide films with the composition of Cr2O3, Ti0.95Nb0.95O4 and (Cr, Fe)2O3 are formed on the surface of alloys. The average oxide films increased with the thermal shock cycle. The thermal shock behavior changed the distribution and quantities of δ phases, which have significantly reduced the strength and increased the elongation of IN718 alloy. The decrease of strength was mainly determined by the dissolution of strengthening phase and grain coarsening. High standard treatment alloy shows the best thermal shock resistance since the strength decreased smallest after thermal shock.

1. Introduction IN718 alloy is a precipitation strengthened nickel-based superalloy, which has been proposed to be used in aerospace structures, aero-engine hot section components and rock components due to its excellent mechanical properties and outstanding oxidation resistance in the range of −253–650 °C [1–3]. Compared to the traditional superalloy, it is hardened by an ordered body-centered cubic γ″ (Ni3Nb) phase and an ordered face-centered cubic γ′ (Ni3Al, Ti) phase [4,5]. The phases normally presented in the IN718 alloy are γ″ phase, γ′ phase, δ phase and a few of carbide. The γ″ phase is the major strengthening phase and γ′ phase plays an assistant strengthening role [6]. If IN718 alloy is being kept above 650 °C for a long time, the metastable γ″ phase would transform into an orthogonal thermodynamic equilibrium δ (Ni3Nb) phase [7,8]. Although the crystal structure of δ phase and γ″ phase are different, they have the same chemical composition of Ni3Nb [9]. As the common substrate of thermal barrier coatings (TBCs), IN718 alloy is usually endured with thermal shock under the actual service environment. The change of temperature gradient will cause deformation and cracking in IN 718 alloy. At the same time, oxidation defects also occur in alloy under high temperature. In the process of work, the above problems often caused the scrap of parts [1, 10–14]. However, there are few reports about the influence of thermal shock behavior on IN718 alloy. Some researches have been focused on the effects of active ⁎

elemental addition on the high temperature oxidation resistance of IN718 alloy [15–18]. Therefore, it is important to investigate the thermal shock resistance and oxidation resistance of IN718 alloy. In the present work, the IN718 alloy under three heat treatments was subjected to thermal shock test between 960 °C and room temperature for 40 times. The influence of thermal shock on the variation of δ phase, oxide films and mechanical properties was evaluated. This research will provide an experimental foundation and theoretical basis for the high-temperature application of IN718 alloy. 2. Experimental procedure 2.1. Materials and heat treatment The IN718 alloy used in this research was provided by a research institute. The chemical composition (mass %) of the as-received materials are as follows: Ni 52.23, Cr 19.36, Nb 5.26, Mo 3.07, Ti 0.91, Al 0.73, Mn 0.08, Co 0.021, Si 0.06, C 0.04, P 0.0048, S 0.0035 and Fe balance. Three different heat treatments in Table 1 were applied to alloys in a muffle furnace before thermal shock. According to the TTT curves of IN718 alloy in Fig. 1, the three heat treatment schemes were formulated [19,20]. Cylindrical heat treatment specimens with 7 mm diameter and 180 mm height were machined from alloy bars.

Corresponding author. E-mail address: [email protected] (J. Wang).

https://doi.org/10.1016/j.apsusc.2019.04.047 Received 27 December 2018; Received in revised form 31 March 2019; Accepted 4 April 2019 Available online 13 April 2019 0169-4332/ © 2019 Published by Elsevier B.V.

Applied Surface Science 484 (2019) 1282–1287

S.-l. Wang, et al.

Table 1 Three heat treatments of IN718 alloy before thermal shock. HST ST DA

960 °C × 2 h/air cooling (AC) + 720 °C × 8 h/furnace cooling (FC) at 55 °C/h to 620 °C + 620 °C × 8 h/air cooling (AC) 960 °C × 1 h/(AC) + 720 °C × 8 h/(FC) at 55 °C/h to 620 °C + 620 °C × 8 h/(AC) 720 °C × 8 h/(FC) at 55 °C/h to 620 °C + 620 °C × 8 h/(AC)

Note: AC-air cooling; FC-furnace cooling.

treatment: put the samples into the furnace maintained at 960 °C for 20 min, followed by air cooling for 15 min. The test was repeated for 40 times. 2.3. Mechanical testing The tensile experiments were performed on an SHT-4605 type universal testing machine. The standard proportional tensile specimens with 50 mm gauge length for 10 mm diameter, the detail sketch is shown in Fig. 2. The tensile tests were conducted according to GB/T 228-2002 [21], the loading velocity was 2 mm/min. 2.4. Microstructure analysis The samples for optical microscopy and scanning electron microscopy examination were prepared by using standard metallographic preparation techniques and etched in the mixture of 5 g CuCl2 + 100 ml HCl +100 ml C2H5OH. ZEISS EVO MA15-scanning electron microscopy equipped with energy dispersive x-ray spectroscopy was used to characterize the morphology and compositions of the surface oxide films, the distribution of δ phase. The near-surface microstructures were evaluated with the SEM. In addition, DX-2700 type x-ray diffraction was applied to analysis the phase composition of oxide films. Then the software of Smile View was used to measure the thickness of oxide films. Quantitative analysis of δ phase was also performed by the Image-Pro Plus software. The micro-hardness was measured by using an HXD-2000TM/LCD Micro Hardness Tester with a load of 200 gf and a loading time of 20 s.

Fig. 1. The TTT curves of IN718 alloy.

3. Results and discussion 3.1. Microstructure and percentages of δ phase

Fig. 2. The detail shape and size of the tensile specimen.

The original optical microstructure (OM) and SEM microstructures of the as-received IN718 alloy are shown in Fig. 3, in which the grain was equiaxed crystal and the average grain size is about 11–12 μm. The microstructure consists of the γ matrix, δ phase and a little carbide. The δ phase distributed at austenite boundary unevenly. Fig. 4 illustrates the microstructures of IN718 alloy before and after thermal shock cycle for 40 times. The morphology and distribution of δ phase in three alloys were changed after thermal shock and the grain

2.2. Thermal shock test Before thermal shock, the pre-treatment of these samples included removing surface oxide films with silicon carbide papers of 400-grit, degreasing with acetone to remove machining greases and oils, rinsing with distilled water, sequentially. The alloys for thermal shock test with dimensions of ϕ10 mm × 5 mm. The experiments involved a two-step

(a)

(b)

twins Carbid

carbide

γ phase δ phase 10 µm

20 µm

Fig. 3. The OM and SEM microstructures of the as-received IN718 alloy (a) OM micrograph; (b) SEM micrograph. 1283

Applied Surface Science 484 (2019) 1282–1287

S.-l. Wang, et al.

Fig. 4. SEM images show the microstructure of IN718 alloy. Before thermal shock: (a) ST (c) HST (e) DA; after thermal shock: (b) ST (d) HST (f) DA.

sizes were increased. It can be concluded from Fig. 3(a),(c),(e) that some short rod-like and particle shaped δ phases precipitated along the boundaries of ST, DA and HST alloys after three heat treatments. While some needle-shaped δ phases located inside the grain of HST. The δ phase in DA alloy was from the as-received IN718 alloy. After thermal shock, Fig. 3(a),(c),(e) shows some needle shaped δ phases, which can be observed at the ST and DA alloy boundaries. In HST alloy, δ phases inside the grain disappeared, the short rod-like shaped δ phases along the grain boundaries transformed into particle shaped uniformly. There are many defects in the grain and twin boundaries, which can reduce the energy required to nucleate the precipitated phase [22,23]. Thus, the δ phase precipitated in boundary firstly. The amount of δ phase increased with increasing solid solution time and precipitated inside grain gradually. The percentage of δ phases before and after thermal shock was quantified with Image-Pro Plus in Table 2. After thermal shock, δ phase precipitated in HST alloy was up to 10.31%, while in ST and DA alloy was only 4.52% and 4.43% respectively. The solvus temperatures of γ″ and γ′ phase were 915 °C and 900 °C respectively and the precipitation temperature of δ phase was range from 780 °C to 980 °C. Hence, there was only austenite matrix and δ phase presented when the temperature

over 915 °C and below 980 °C [24]. This could explain why the amount of δ phases increased with increasing solid solution time, and δ phase can also be observed inside the grains after HST heat treatment. The highest aging temperature of DA was 720 °C, which is not in the range of the precipitation temperature of δ phase, so nearly no δ phase precipitated after DA treatment. Owing to thermal shock cycle is holding at 960 °C, this temperature is suitable for δ phase precipitated from the matrix and the amounts of δ phases were increased after thermal shock. Hence, the percentage of δ phases in ST, HST and DA were increased to 6.44%, 13.98% and 5.47% respectively.

Table 2 Percentages of δ phases in IN718 alloy before and after thermal shock.

1284

Samples

ST

HST

DA

Area ratio of δ phases (%) before thermal shock Area ratio of δ phases (%) after thermal shock

4.52 6.44

10.31 13.98

4.43 5.47

Applied Surface Science 484 (2019) 1282–1287

S.-l. Wang, et al.

Fig. 5. Oxide surface morphology of IN718 alloy after thermal shock for 40 times. (a) ST; (b) HST; (c) DA. 8000

O O

8000

2000

Cr

Wt. 21.0 51.7 1.27 8.23 1.57 3.24 7.95

At. 39.1 29.6 Cr 0.68 9.07 0.8 0.09 0.1

Ni

Fe

Ti

HST

ST

6000

Elemen Wt. At.% O 19.27 42.6 Cr 67.62 46.0 Fe 1.88 1.19 Ni 1.84 1.11 Ti 4.18 3.09

4000

Counts

Elemen O Cr Fe Al Al Ni Si Ti Si

Counts

Counts

4000

O

6000

DA

6000

Elemen Wt. At. O 23.71 43.1 Cr 61.11 34.2 Fe 1.68 0.88 Ni 0.94 0.47 Ti 2.01 1.22 Al 0.64 0.05

4000

2000

2000 C C

Nb

Cr

0 0

2

4

6

C Ni

8

0

2

Energy/keV

Fe

Ti

0

Nb

Cr

Ti Al

4

Al

Fe

Cr

Ti

Ti

Cr

Fe Fe

0

6

8

0

2

Energy/keV

4

6

8

Energy/keV

Fig. 6. EDS spectrum of oxide films on IN718 alloys after thermal shock. (a) DA; (b) ST; (c) HST.

to Cr2O3, Cr1.3Fe0.7O3 and Al0.5CNi3Ti0.5. Fig. 8 presents the cross-section morphology after thermal shock at 960 °C for 40 times. A large number of holes were formed at the interface between the matrix metal and oxide layer. During the formation of surface oxide film, because of the different diffusion rates of various atoms and ions, the cationic vacancies diffusion toward the vicinity of interface. In order to lower the energy, these vacancies will gather to grow up and results in the formation of holes. However, holes have significant influence on the adherence of oxide films. The oxide films may be easily fall off from the surface of alloy, when the porosity size reaches a certain critical value. Fig. 9 gives the average thickness of the oxide films, it demonstrated that the mean thickness of oxide films was obviously increased with thermal shock cycle. No continuous oxide films were formed on the surface of IN718 alloy due to fewer thermal shock cycles. After the same cycle of thermal shock, the oxide films after ST was thicker than the two others, and which follows as: DA > HST. The oxide films thickness of ST alloy was increased from 0.92 μm to 2.41 μm from 10th to 40th cycle, the growth rate was faster than HST and DA alloy. According to the previous literature [25], the Cr usually accounted for about 20% of the total IN718 alloy, which can provide a good oxidation resistance by formation of a dense and rugged Cr2O3 oxide films on the surface. Generally, Cr2O3 formed through Cr element diffused along the grain and phase boundaries. IN718 alloy contains plentiful Fe, which reduced the oxidation resistance of alloy and reacted with O, Cr then formed Cr1.3Fe0.7O3. A lot of published literature showed that the oxidation mechanism of IN718 alloy was determined by element diffusion [26–29]. The oxidation process was primarily controlled by the inward diffusion of O2−. In high temperature, the surface oxidation of IN718 alloy was caused by contact with air and the diffusion of oxygen atom along the grain boundary or phase boundary.

Fig. 7. X-ray diffraction (XRD) pattern of oxide films on IN718 alloy under three heat treatments after thermal shock.

3.2. Oxide films Oxide films were formed on the surface of IN718 alloy after thermal shock cycle for 40 times, the oxide morphologies are presented in Fig. 5. The oxide films were dark grey and grew along with the outer scratches densely, which were composed of irregularly shaped particles with different sizes. However, the bonding between the particles was not tight enough and there were obvious voids. All the oxide films did not come adrift from the alloy's surface, and there were no visible cracks are observed, which indicates that the adhesion strength between oxide film and matrix is pretty strong. The compositions of the films were measured by EDS, and the results are shown in Fig. 6. It is seen that the oxide films is rich in O, the content of Cr is about 60%，and there are small amount of Ni, Fe, Ti and Al. combined with the X-ray patterns of the surface oxide films in Fig. 7, it could be further confirmed that the major components belong

3.3. Mechanical properties The tensile properties of IN718 alloy before and after thermal shock 1285

Applied Surface Science 484 (2019) 1282–1287

S.-l. Wang, et al.

Fig. 8. SEM micrographs showing the surface structure of IN718 alloy after thermal shock. (a) DA; (b) ST; (c) HST.

Oxide scale thickness (mm)

2.5

Table 3 Room temperature tensile properties of IN718 superalloy.

ST HST DA

2.0

Sample ST DA HST ST (thermal shock) DA (thermal shock) HST (thermal shock)

1.5

δb/Mpa

δ0.2/Mpa

Elongation, %

1510 1450 1430 1350 1280 1160

995 950 90 485 455 465

32.2 29.7 33.6 52.2 55.5 49.2

Note: δ0.2-yield strength; δb-high tensile strength.

1.0

800 700

0.5

Before thermal shock After thermal shock

600

20 30 Thermal shock (times)

40

Hardness/HV

10

Fig. 9. Average thickness of surface oxide films on IN718 alloy after thermal shock.

tests are pictured in Table 3. The δ0.2 and δb of IN718 alloy after thermal shock are dropped dramatically. The δb and δ0.2 of ST alloy are appreciably higher than the two other alloys. The data demonstrated that the δb of ST, HST and DA for IN718 alloy after thermal shock decrease by 34.31%, 32.87% and 34.48%, respectively. The values of δ0.2 decrease by 64.07%, 59.91% and 64.45%. However, the elongations after thermal shock are significantly increased. The increase in elongation after thermal shock can be explained by no γ″ and γ′ phase serving as dislocation barrier. Three samples after thermal shock were found to be with a very low hardness as shown in Fig. 10. It is seen that the hardness of IN718 alloys

500 400 300 200 100 0 DA

ST

HST

Fig. 10. Hardness of IN718 alloys before and after thermal shock.

after thermal shock declines abruptly compared with before thermal shock. The hardness of DA, ST and HST alloys decreases from 630.6, 584, 578 to 257, 245 and 223 respectively after thermal shock. The 1286

Applied Surface Science 484 (2019) 1282–1287

S.-l. Wang, et al.

hardness of IN718 alloys are mainly determined by the number of precipitated strengthening phases. The more the number of strengthening phases, the higher the hardness value of alloys. Then change of grain size has little effect on the hardness value. The result is consistent with Table 2. The mechanical properties of IN718 alloy are major determined by the volume fraction and particle size of γ″ phase and γ′ phase [30–32]. The reasons for the strength degradation after thermal shock are mainly as follows. Firstly, during thermal shock, the dissolution of γ″ and γ′ phase into the matrix resulted in the strength would not reach the maximum strength condition. Secondly, the grain coarsening behavior and existence of thermal stress also reduced the strength of alloys. Thirdly, the thermal shock not only caused the increase of surface roughness but also lead to stressing concentration due to the microscopic holes on the surface, which have a significant influence on the strength. However, the oxide films formed during thermal shock was extremely thin, and the impact on strength can be ignored.

Oxid. Met. 55 (2001) 505–521. [6] M. Sundararaman, P. Mukhopadhyay, S. Banerjee, Precipitation of the δ-Ni3Nb phase in two nickel base superalloys, Metall. Trans. A. 19 (1988) 453–465. [7] D.Y. Cai, W.H. Zhang, P.L. Nie, W.C. Liu, M. Yao, Dissolution kinetics of δ phase and its influence on the notch sensitivity of Inconel 718, Mater. Charact. 58 (2007) 220–225. [8] K. Kulawik, P.A. Buffat, A. Kruk, A.M. Wusatowska-Sarnek, A.C. Filemonowicz, Imaging and characterization of γ′ and γ″ nanoparticles in Inconel 718 EDX element mapping and FIB-SEM tomography, Mater. Sci. Charact. 100 (2015) 74–80. [9] M.Q. Li, W.B. Ju, Y.Y. Lin, X.J. Wang, Y. Niu, Deformation behaviour of GH4169 nickel based superalloy in isothermal compression, Mater. Sci. Tech-Lond. 24 (2013) 1195–1198. [10] J.H. He, γ″ precipitation in Inconel 718, J. Mater. Sci. Technol. 10 (1994) 293–303. [11] X.H. Zhang, Y. Zhang, B.H. Tian, J.C. An, Z. Zhao, A.A. Volinsky, Y. Liu, K.X. Song, Arc erosion behavior of the Al2O3-Cu/(W, Cr) electrical contacts, Compos. Part B 160 (2019) 110–118. [12] R. Cozar, A. Pineau, Morphology of γ″ and γ′ precipitation and thermal stability of Inconel 718 type alloys, Metall. Trans. 4 (1973) 47–59. [13] Y.P. Mei, Y.C. Liu, C.X. Liu, C. Li, L.M. Yu, Q.Y. Guo, H.J. Li, Effects of cold rolling on the precipitation kinetics and the morphology evolution of intermediate in Inconel 718 alloy, J. Alloys Compd. 649 (2015) 949–960. [14] B.J. Wang, Y. Zhang, B.H. Tian, V. Yakubov, J.C. An, K.X. Song, L.H. Li, M. Fu, Effects of Ce and Y addition on microstructure evolution and precipitation of Cu-Mg alloy hot deformation, J. Alloys Compd. 781 (2019) 118–130. [15] L. Geng, Y.S. Na, N.K. Park, Oxidation behavior of alloy 718 at a high temperature, Mater. Des. 28 (2007) 978–981. [16] V. Garat, J. Deleume, J.M. Cloue, E. Andrieu, High temperature intergranular oxidation of alloy 718, Superalloy Various Deriv. (2005) 550–569. [17] K.A. Al-hatab, M.A. Al-bukhaiti, U. Krupp, M. Kantehm, Cyclic oxidation behavior of IN718 superalloy in air at high temperature, Oxid. Met. 75 (2011) 209–238. [18] D.M. England, A.V. Virkar, Oxidation kinetics of some nickel-based superalloy foils and electronic resistance of the oxide scale formed in air part I, J. Electrochem. Soc. 146 (1999) 3196–3202. [19] Y.C. Liu, Q.Y. Guo, C. Li, Y.P. Mei, X.S. Zhou, Y. Huang, H.J. Li, Recent progress on evolution of precipitates in Inconel 718 superalloy, Acta Metall. Sin. 152 (2016) 1259–1266. [20] R.B. Li, M. Yao, W.C. Liu, X.C. He, Effect of direct aging treatment on structure of Inconel 718 alloy, Heat Treat. Met. 27 (2002) 21–22. [21] Y. Huang, T.G. Langdon, The evolution of delta-phase in a superplastic Inconel 718 alloy, J. Mater. Sci. 42 (2007) 421–427. [22] T. Cui, J.Q. Wang, X.X. Wang, W.H. Wang, Z.X. Wang, H.C. Yang, High temperature oxidation of superalloy GH4169, Corros. Sci. Prot. Technol. 16 (2004) 53–57. [23] F. Abe, H. Araki, M. Okada, The role of aluminum and titanium on the oxidation process of a nickel-base superalloy in steam at 800 °C, Oxid. Met. 27 (1987) 21–36. [24] M.H. Yoo, Stability of super-dislocations and shear faults in L12 ordered alloys, Acta Metall. 35 (1987) 1559–1569. [25] N. Hussain, K.A. Shahid, I.H. Khan, S. Rahman, Oxidation of high-temperature alloys (superalloys) at elevated temperatures in air.II, Oxid. Met. 41 (1994) 251–269. [26] L. Kumar, R. Venkataramani, M. Sundararaman, P. Mukhopadhyay, S.P. Garg, Oxidation of Inconel 718 in air at high temperatures, Oxid. Met. 45 (1996) 221–224. [27] C. Juillet, A. Oudriss, J. Balmain, X. Feaugas, F. Pedraza, Characterization and oxidation resistance of additive manufactured and forged IN718 Ni-based superalloy, Corros. Sci. 142 (2018) 266–276. [28] J.P. Collier, H.W. Song, K.T. Johk, J.C. Phillips, The effect of varying Al, Ti, and Nb content on the phase stability of Inconel 718, Metall. Trans. A. 19 (1988) 1657–1666. [29] Y.T. Chen, A.C. Yeh, M.Y. Li, S.M. Kuo, Effects of processing routes on room temperature tensile strength and elongation for Inconel 718, Mater. Des. 119 (2017) 235–243. [30] J.H. Song, H. Huh, The effect of strain rate on the material characteristics of nickelbased superalloy Inconel 718, Key Eng. Mater. 340-341 (2007) 283–288. [31] J.H. Du, X.D. Lü, Q. Deng, Effect of heat treatment on microstructure and mechanical properties of GH4169 superalloy, Rare Metal Mater. Eng. 43 (2014) 1830–1834. [32] L.D. Lu, J.H. Du, Q. Deng, High temperature structure stability of GH4169 superalloy, Mater. Sci. Eng. A 559 (2013) 623–628.

4. Conclusions In the present work, thermal shock behavior changes the microstructure and reduces the strength of the IN718 alloy. The different solvus and precipitate temperatures of γ″, γ′ and δ phase cause the δ phase increased after thermal shock, while γ″ and γ′ phase dissolve into the matrix. Protective and dense oxide films with the composition of Cr2O3, Ti0.95Nb0.95O4 and (Cr, Fe)2O3 are formed on the surface of alloys. The oxidation mechanism of IN718 was determined by the inward diffusion of O2−. The strength of three alloys after thermal shock decline sharply while the elongation increase, which are mainly determined by the dissolution of strengthening phase and grain coarsening. HST alloy shows the best thermal shock resistance since the strength decreased smallest after thermal shock. Acknowledgments This work was supported by the National Natural Science Foundation of China [51374180 and 51602269], Guangxi Key Laboratory of Information Materials (Guilin University of Electronic Technology), and Scientific Research Foundation and Opening Foundation (X151518KCL23) of Southwest Petroleum University. References [1] M. Jouiad, E. Marin, R.S. Devarapalli, J. Cormier, F. Ravaux, C. Le Gall, J.M. Franchetd, Microstructure and mechanical properties evolution of alloy 718 during isothermal and thermal cycling over-aging, Mater. Des. 102 (2016) 284–296. [2] T.D. Lee, W.H. Hou, Development of fine-grained structure and the mechanical properties of nickel-based superalloy 718, Mat. Sci. Eng. Atruct. 555 (2012) 13–20. [3] H.B. Dong, G.C. Wang, Effect of deformation process on superplasticity of Inconel 718 alloy, Rare Metal Mater. Eng. 44 (2015) 0298–0302. [4] M. Anderson, A.L. Thielin, F. Bridier, P. Bocher, J. Savoie, δ phase precipitation in Inconel 718 and associated mechanical properties, Mater. Sci. Eng. A 679 (2017) 48–55. [5] G.A. Greene, C.C. Finfrock, Oxidation of Inconel 718 in air at high temperatures,

1287