Effect of Withdrawal Rates on Microstructure and Creep Strength of a Single Crystal Superalloy Processed by LMC

Transcription

1 J. Mater. Sci. Technol., 21, 26(4), Effect of Withdrawal Rates on Microstructure and Creep Strength of a Single Crystal Superalloy Processed by LMC Chengbao Liu 1), Jian Shen 1), Jian Zhang 2) and Langhong Lou 1) 1) Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 1116, China 2) National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 1116, China [Manuscript received December 18, 28, in revised form July 17, 29] A nickel base single crystal (SC) superalloy was directionally solidified using liquid metal cooling (LMC) process at various withdrawal rates. The microstructure was refined as increasing the withdrawal rate from 3 to 12 mm/min. However, higher withdrawal rate of 15 mm/min induced the formation of stray grains. Size and volume fraction of the eutectics were found to decrease with the increasing in withdrawal rate. After solution heat treatment at 125 C, un-dissolved eutectic was observed in specimens. High temperature creep rupture life was observed to be very sensitive to the fraction of these remaining eutectics. Creep rupture tests at 1 C/235 MPa showed that refined microstructure and low fraction of the remaining eutectic lead to significant improvement of the rupture life. KEY WORDS: Eutectic; Single crystal superalloy; Liquid-metal-cooling; Withdrawal rate 1. Introduction Directionally solidified (DS) and single crystal (SC) blades and vanes are widely used in aero and power generation gas turbines [1,2]. These components are mainly produced by high rate solidification (HRS) process, where the heated mold was gradually lowered from the heating furnace thus providing a uniaxial thermal gradient. The heat extraction occurred by the thermal conduction through the casting to the water chill plate and radiation through the mold shell to vacuum chamber. Recently, due to the difficulties in producing large DS and SC components used in industrial gas turbines (IGTs), considerable work has been concentrated on the liquid metal cooling (LMC) process, where high thermal gradient and therefore improved productivity and lower costs can be achieved. In this process, the heat was removed by immersing the casting and mold into a container of liquid metal coolant as they were withdrawn from the furnace. Since 198 s, LMC us- Corresponding author. Prof., Ph.D.; Tel.: ; address: jianzhang@imr.ac.cn (J. Zhang). ing aluminum as the coolant has been employed by the former Soviet Union to produce DS and SC blades for aero engines [3]. Improved mechanical properties and microstructural stability have been reported [3,4]. Beginning from 199s, the LMC process re-attracted the attentions in Europe and the US due to the increasing demand of large DS and SC components for IGT applications [5]. Because of the lower melting temperature, liquid tin has significantly higher heat flux than aluminum as a cooling medium, and can therefore provide higher thermal gradient and greater cooling rates [6]. Moreover, liquid aluminum was observed to attack the superalloy casting at 7 C, while liquid tin showed no harmful effect on the surface of the casting [7]. For these reasons, liquid tin was chosen as the cooling medium in the LMC furnaces developed in recent years [8]. Most of the previous work concerning the LMC process has been focused on the processing optimization of large DS castings. Comparative study and thermal simulation of large DS castings processed by HRS and LMC has revealed that LMC process provides higher thermal gradient and therefore refined

2 C.B. Liu et al.: J. Mater. Sci. Technol., 21, 26(4), Fig. 1 Stray grains in bars solidified at 15 mm/min: (a) vertical surface, (b) cross section at 12 mm from chill microstructure. The alleviated segregation and cast defects, as well as better production yield for large component with shorter casting time was expected in Sn assisted LMC process [8 12]. Improved mechanical performance of the LMC alloys, especially fatigue properties was also reported [13 15]. More recently, our preliminary work has demonstrated that LMC could also be an alternative way to develop stronger alloys since the content of alloying elements could be higher as a result of the reduced micro-segregation [8,16]. A hot corrosion resistant SC alloy DD1 was developed based on this concept [17]. The aim of the present paper is to study the effect of withdrawal rates on the as-cast microstructure, heat treatment response and creep strength of DD1 alloy processed by LMC. 2. Experimental The nickel base DD1 SC superalloy used in the present experiments contains 13Cr, 4Co, 7.8(Al+Ti), 11 13(Ta+W+Mo), with minor C and B, and Ni in balance. The single crystal alloy with <1> orientation was directionally solidified by LMC technique using liquid Sn as cooling medium at different withdrawal rates: 3, 6, 9, 12 and 15 mm/min. Each casting produced four SC bars with 16 mm in diameter and 2 mm in length. The bars were macro etched by 7%HCl+3% peroxide solution. Samples for microstructural examination were cut from the midst of the bars. They were polished and electrolytically etched with 1 ml H 3 PO 4 +9 ml H 2 O reagent. From randomly chosen 2 pieces of optical microscopy photos, the primary dendrite arm spacing (PDAS), secondary dendrite arm spacing (SDAS), volume fractions and size of eutectic were analyzed by the Image- Pro Plus software. Solution heat treatment was carried out at 125 C/3 h/air cooling (AC), followed by two steps aging treatment at 11 C/5 h/ac+87 C/24 h/ac. The heat treated specimens were machined into creep rupture samples with a gauge length of 25 mm and a diameter of 5 mm. Creep rupture tests were performed at 1 C/235 MPa, with the loading direction parallel to the <1> direction of SC. Metallographic examinations of the heat treated and creep ruptured samples were performed by optical microscopy and scanning electron microscopy (SEM). 3. Results and Discussion Macro etching revealed that the highest withdrawal rate applied in the present work led to stray grain formation. After the successful grain selection at the withdrawal rate of 15 mm/min, two to three grains nucleated at the up-middle part of the SC bars (Fig. 1), and grew through the rest part of the castings. All the other castings performed at withdrawal rates of 3 12 mm/min resulted in SC structures. The following characterization was therefore only focused on these SC bars. Metallographic microstructure (Fig. 2) of the single crystal bars cast at different withdrawal rates showed that the alloys had dendrite structure. No stray grains or low angle grain boundary was observed. PDAS and SDAS measured in these samples were summarized in Fig. 3. The eutectic morphology and its size and fraction were compared in Fig. 4 and 5. It is clear from Figs. 2 5 that increasing the withdrawal rate obviously leads to the refinement of microstructure. Solution heat treatment at 125 C/3 h cannot completely remove the as-cast γ/γ eutectics. However, Fig. 6(a) and (b) clearly demonstrate that the residual eutectics in the sample solidified at higher withdrawal rate were much smaller in size. The fraction of residual eutectics was also reduced to a very low level in the sample solidified at the withdrawal rate of 12 mm/min. Uniformly distributed γ particles with similar shape and size precipitated after aging in both samples. Figure 6(c) and (d) compare the morphology of γ and residual eutectics at inter-

3 38 C.B. Liu et al.: J. Mater. Sci. Technol., 21, 26(4), Fig. 2 Dendrite structure of the cast specimen solidified at different withdrawal rates: (a) 3 mm/min, (b) 6 mm/min, (c) 9 mm/min, and (d) 12 mm/min 2 PDAS 6 mm/min SDAS 9 mm/min 12mm/min Secondary dendrite arm spacing / Primary dendrite arm spacing / m 3 mm/min m 6 Fig. 3 Comparison of PDAS and SDAS in samples solidified at different withdrawal rates dendritic area in these samples. It is worth to note that the γ channels within the residual eutectics were much wider than that of the matrix after heat treatment. The creep rupture life and elongation of samples solidified at different withdrawal rates were listed in Table 1. The creep rupture life increased significantly as increasing the withdrawal rates. It is interesting to note that most of the secondary cracks in this alloy were associated with the residual eutectics (Fig. 7(a) and (b)). SEM images revealed the raft structure of the γ precipitates after creep rupture tests. Secondary cracks were observed either at the matrix/eutectic interface (Fig. 7(c)), or within the residual γ/γ eutectics (Fig. 7(d)). Fig. 4 Eutectic island of the cast specimen solidified at different withdrawal rates: (a) 3 mm/min, (b) 6 mm/min, (c) 9 mm/min, and (d) 12 mm/min

4 C.B. Liu et al.: J. Mater. Sci. Technol., 21, 26(4), Eutectic size / m mm/min 6 mm/min 9 mm/min Size Fraction 12 mm/min Eutectic volume fraction / % Fig. 5 Size and fraction of the eutectic in the alloy solidified at different withdrawal rates Table 1 The 1 C/235 MPa creep rupture life and elongation of alloys solidified at different withdrawal rates Withdrawal rate Rupture life/h Elongation/% 3 mm/min 45.1± ±4.5 6 mm/min 55.2± ±4.4 9 mm/min 62.1± ± mm/min 79.8± ±2.7 The as-cast microstructure was refined as increasing the withdrawal rate. Comparing to the typical PDAS and SDAS achieved in conventional HRS process, which was 3 and 7 µm, respectively [8,9], LMC technique provides a much higher thermal gradient and causes finer dendritic structure. Since the eutectic formed during the latter stage of solidification, the nucleation site and growth room of eutectic were all restricted by the as-solidified dendrites. The increasing withdrawal rates induced finer dendrite structure, and therefore smaller eutectic pools. The volume fraction of eutectic was related to the amount of residual liquid after dendrite arm formation. At higher withdrawal rate (higher cooling rate), the fraction of residual liquid at interdendritic area was reduced, therefore the volume fraction of eutectic also decreased with increasing withdrawal rate. It was reported that the un-dissolved eutectic in SC superalloys was prone to the formation of cleavage fracture during tensile [21] and high cycle fatigue [22]. Our present results showed that the interface of matrix/residual eutectic was also vulnerable site for crack initiation during creep rupture tests (Fig. 7(c)). Moreover, raft structure was observed in the residual eutectics after creep, but in a larger scale, i.e. the γ channel was much wider than that of the general raft structures found in the matrix. The wider matrix channel is deformed easily than the narrower one during high temperature creep [23]. Therefore, residual eutectics were generally detrimental to the creep properties. The residual eutectics may additionally weaken the mechanical properties since it contains high amount of γ forming elements, Ta, Al and Ti. Fig. 6 Microstructure of the interdendritic area showing the un-dissolved eutectic and uniformly distributed γ precipitates after heat treatment. Sample obtained at the withdrawal rates of (a) and (c) 6 mm/min, (b) and (d) 12 mm/min

5 31 C.B. Liu et al.: J. Mater. Sci. Technol., 21, 26(4), Fig. 7 Morphology of secondary cracks formed in the vicinity of remaining eutectics after creep rupture test at 1 C/235 MPa (loading direction indicated at the lower left corner). Sample obtained at the withdrawal rates of (a) and (d) 6 mm/min, (b) and (c) 12 mm/min 4. Summary Using the LMC process, a wider withdrawal rate range could be adapted to produce single crystal alloy with refined microstructure. The PDAS, SDAS, eutectic size and volume fraction decreased with increasing the withdrawal rate. The refined as-cast microstructure leads to better heat treatment response and therefore improved high temperature creep rupture properties. It is expected that with further optimization of the LMC parameters and heat treatment procedure, improved mechanical properties of DD1 SC alloy can be achieved. REFERENCES [1 ] P. Caron and T. Khan: Aerosp. Sci. Technol., 1999, (3), 513. [2 ] G.E. Fuchs: Mater. Sci. Eng., 21, 3A, 52. [3 ] Y.A. Bondarenko and E.N. Kablov: Metal Sci. Heat Treatment, 22, 44, 288. [4 ] E.N. Kablov, V.V. Gerasimov, V.A. Dubrovskii and E.M. Visik: Mater. Eng., 1996, 5, 16. (in Chinese) [5 ] M. Konter and M. Thumann: J. Mater. Process. Technol., 21, 117, 386. [6 ] A.F. Giamei and J.G. Tschinkel: Metall. Mater. Trans., 1976, 7A, [7 ] A. Lohmueller, W. Esser, J. Grossmann, M. Hoerdler, J. Preuhs and R.F. Singer: in Superalloys 2, eds. T.M. Pollock, R.D. Kissinger, R.R. Bowman, K.A. Green, M. Mclean, S. Olson and J.J. Schirra, TMS, Warrendale, PA, 2, 181. [8 ] J. Zhang and L.H. Lou: J. Mater. Sci. Technol., 27, 23(3), 289. [9 ] A.J. Eillott, S. Tin, W.T. King, S.C. Huang, M.F.X. Gigliotti and T.M. Pollock: Metall. Mater. Trans., 24, 35A, [1] R.C. Reed: in The Superalloys Fundamentals and Applications, Cambridge University Press, 26, 136. [11] A.J. Eillott and T.M. Pollock: Metall. Mater. Trans., 27, 38A, 871. [12] A. Kermanpur, N. Varahram, P. Davami and M. Pappaz: Metall. Mater. Trans., 2, 31B, [13] M. Lamm, A. Volek, O. Luesebrink and R.F. Singer: in Mater. Adv. Power Eng., 26, eds. J.Lecomte-Beckers, M.Carton, F.Schubert and P.J.Ennis, Froschungszentrum Juelich GmbH, Juelich, 26, 334. [14] S. Balsone, G.J. Feng, L. Peterson and J. Schaeffer: in Solidification Processes and Microstructures-A Symposium in Honor of Wilfried Kurz, eds. M. Rappaz, C. Beckermann and R. Trivedi, TMS, 24, 77. [15] W.S. Tang, J. Shen, D.W. Wang, J. Zhang and L.H. Lou: in Proc. 138th Annual Meeting & Exhibition, TMS (The Minerals, Metals Materials Society), San Francisco, California, USA, 29, in press. [16] T. Zhao, D. Wang, J. Zhang, G. Chen and L.H. Lou: J. Mater. Sci. Technol., 29, 25(3), 361. [17] C.B. Liu, J. Zhang and L.H. Lou: China Patent, (in Chinese) [18] Y.Z. Lu, D.W. Wang, J. Zhang and L.H. Lou: Foundry, 29, 3, 245. (in Chinese) [19] N. D Souza, M.G. Ardakani, M. McLean and B.A. Shollock: Metall. Mater. Trans., 2, 31A, [2] J.C. Brice: J. Cryst. Growth, 1971, 1, 25. [21] W.S. Walston, I.M. Bernstein and A.W. Thompson: Metall. Trans., 1991, 22A(6), [22] P. Luká, L. Kunz and M. Svoboda: Mater. Sci. Eng., 24, , 55. [23] M.F. Kniepmeier, U. Hemmersmeier, T. Kuttner and T. Link: Scripta Metall. Mater., 1994, 3, 1275.