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Microstructure evolution and its influence on deformation mechanisms during high temperature creep of a nickel base superalloy
Materials Science and Engineering A 499 (2009) 445–453
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure evolution and its influence on deformation mechanisms during high temperature creep of a nickel base superalloy Javad Safari a,∗ , Saeed Nategh b,1 a b
Materials Science and Engineering Department, Shahid Chamran University, Ahwaz, Iran Materials Science and Engineering Department, Sharif University of Technology, P.O. Box 11365-9466, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 20 August 2007 Received in revised form 27 August 2008 Accepted 8 September 2008 Keywords: Nickel base superalloys Creep Microstructure Dislocation Deformation mechanism Transmission electron microscopy
a b s t r a c t The interaction of dislocation with strengthening particles, including primary and secondary ␥ , during different stages of creep of Rene-80 was investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). During creep of the alloy at 871 ◦ C under stress of 290 MPa, the dislocation network was formed during the early stages of creep, and the dislocation glide and climb process were the predominant mechanism of deformation. The density of dislocation network became more populated during the later stages of the creep, and at the latest stage of the creep, primary particles shearing were observed alongside with the dislocation glide and climb. Shearing of ␥ particles in creep at 871 ◦ C under stress of 475 MPa was commenced at the earlier creep times and governed the creep deformation mechanism. In two levels of examined stresses, as far as the creep deformation was controlled by glide and climb, creep curves were found to be at the second stage of creep and commence of the tertiary creep, with increasing creep rate, were found to be in coincidence with the particles shearing. Microstructure evolution, with regard to ␥ strengthening particles, led to particles growth and promoted activation of other deformation mechanisms such as dislocation bypassing by orowan loop formation. Dislocation-secondary ␥ particles interaction was detected to be the glide and climb at the early stages of creep, while at the later stages, the dislocation bypassed the secondary precipitation by means of orowan loops formation, as the secondary particle were grown and the mean inter-particle distance increased. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Cast nickel base superalloys containing medium to high volume fractions of the ordered ␥ phase are used as blades and other high temperature components in aero- and power-generating gas turbines [1,2]. High mechanical properties of Ni base superalloys at high temperatures are mostly attributed to the characteristics of the ␥ precipitates. In Ni–Al–Ti system, ␥ is an ordered phase with L12 structure and nominal composition of Ni3 (Al,Ti) that is embedded in the disordered (FCC) ␥ matrix [3–6]. Polycrystalline, and recently directionally solidified, Rene-80 (GE trademark) with bimodal precipitation of ␥ particles, is frequently used as first stage aero-gas turbine blades and is subject to high temperature degradation mechanisms mostly creep and low
∗ Corresponding author. Tel.: +98 611 333 0015; fax: +98 21 2205 7662. E-mail addresses: [email protected] (J. Safari), [email protected] (S. Nategh). 1 Tel.: +98 912 384 2146; fax: +98 21 6600 5717. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.017
cycle fatigue [7–9]. Cast microstructure evolution during different stages of heat treatment and subsequent microstructure degradation due to thermal exposure at high temperatures were studied by the authors and were reported elsewhere [10,11]. Creep deformation mechanisms of superalloys are controlled by the interaction of dislocation with primary and secondary ␥ particles during dislocation creep [12–16]. Lin and Wen [17–19] have performed some investigation on creep and tensile deformation mechanisms of directionally solidified Rene-80. Their results may be summarized as follows. In creep deformation at 760 ◦ C under stress of 618 MPa, ␥ shearing occurs by viscous slip of pairs of superlattice intrinsic stacking faults (SISF)-separated by 1/31 1 2 super partials with net slip vector 1 1 0, while the cores of the 1/31 1 2 superpartials do not appear to be separated [17–19]. Creep deformation at 982 ◦ C under stress of 190 MPa is controlled by glide and climb of dislocation from loose dislocation network around ␥ particles. Almost the same results were obtained by Sajjadi and Nategh on GDT-111 Nickel base superalloy. They referred to this alloy as the new modification of Rene-80 [20].
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J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
Fig. 1. Class-A heat treatment of the Rene-80 per GE-C50TF28 [7].
2. Experiments 2.1. Material and heat treatment The material used in this research was the polycrystalline Rene80 nickel base superalloy [7,8]. Cylindrical bars, 12 mm in diameter and 120 mm in length, were produced by vacuum precision casting. Casting was carried out using previously determined parameters which produced components with acceptable defect levels by the GE standard assessment [7]. Soundness of the samples after casting has been approved by dye penetrant and radiography as the NDT examinations. The bulk composition was determined by XRF analysis. The analysis was repeated at least five times for each casting and the averaged results are presented in Table 1. Samples were then heat treated in a vacuum between 10−3 and 10−4 mbar to the class-A GE specification [7]. This heat treatment consists of four stages; 2 h at 1204 ◦ C (solution heat treatment), 4 h at 1093 ◦ C (primary aging), 4 h at 1054 ◦ C (coating heat treatment), and 16 h at 843 ◦ C (Aging). These stages are shown graphically in Fig. 1.
Fig. 3. Rene-80 creep curves at 871 ◦ C under two stress levels of (A) 290 MPa and (B) 475 MPa.
order to investigate the microstructure evolution and deformation mechanisms during creep, creep tests were interrupted at different stages of tests and samples were prepared for investigation with high resolution scanning electron microscopy (SEM) and transmission electron microscopy (TEM).
2.2. Creep tests Creep tests at 871 ◦ C and two different stress levels, 290 and 475 MPa were carried out according to the GE C50 TF28 specification and ASTM E139. Samples temperature was controlled within ±2 ◦ C by attached thermocouples and 1 h soaking time was applied on samples before applying the creep stress. Creep extension was measured by the extensometers placed on samples projections. In
Fig. 2. Microstructure of Rene-80 after GE class-A heat treatment.
Fig. 4. Creep rate versus normalized creep time for creep tests at 871 ◦ C under stress of (A) 290 MPa and (B) 475 MPa.
J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
447
Table 1 Comparison of Rene-80 standard composition and experimental analysis result for one casting. wt%
Ni
Cr
Co
Ti
Al
Mo
W
C
B
Zr
P
S
Min [7] Max [7] Exp.a
Bal. Bal. Bal.
13.70 14.30 14.24
9.00 10.00 9.45
4.80 5.20 5.12
2.80 3.20 3.05
3.70 4.30 3.96
3.70 4.30 4.12
0.15 0.19 0.17
0.01 0.02 0.014
0.02 0.1 0.07
– 0.015 0.010
– 0.0075 0.0050
a
Average of at least five analyses.
Fig. 5. Microstructure evolution of Rene-80 during creep at 871 ◦ C under stress of 290 MPa after (A) 0.1%, (B) 0.5%, (C) 1.7%, (D) 4% creep strain and (E) after creep rupture of samples.
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Table 2 Creep data of Rene-80 at 871 ◦ C at different stresses. Stress, MPa
Rupture time, h
Min. creep rate, s−1
Reduction of area, %
Elongation, %
Percentage of secondary creep life, %
290 475
629 8.43
1.90E−08 1.25E−06
14.12 14.23
5.51 5.74
72 22
2.3. Microstructure characterization Samples were sectioned for preparation of metallography and microstructure evolution. Microstructure was studied by field emission high resolution SEM and TEM. Sample preparation for electron microscopy was carried out by the lapping, electro polishing with a solution of 10% HCl + methanol at −40 ◦ C and 25 V for 20 s and electro etching with 170 mL phosphoric acid + 10 mL sulfuric acid and 16 g chromium trioxide at room temperature and 6 V for 4 s. The H3 PO4 –H2 SO4 –CrO3 etchant attacks the gamma solid solution and forms an anodic film on the gamma areas. This anodic film permits the differentiation between ␥ and ␥ in image analysis. The fact that this etch attacks gamma, makes it the preferred etch for bringing out the very fine ␥ [10]. As the quantitative metallographic investigation is basically a statistical measurement, a large number of images at different magnifications for all conditions of aging were analyzed. However, only one of each is shown here as being representative. Measurement of the ␥ particle characteristics was performed on the SEM photos using the UTHSCSA Image Tool software, a semi-quantitative image analyzer program
(see web site address: http://ddsdx.uthscsa.edu/dig/itdesc.html for the software information and download).
3. Results and discussion Fig. 2 shows the microstructure of Rene-80 superalloy after completion of class-A heat treatment according to GE C50 TF28. Microstructure of Rene-80 consists of dual precipitation of ␥ strengthening particles in the austenitic ␥ matrix. Primary ␥ particles with cubic morphology, which have been formed at the primary aging stage of heat treatment, have 450 ± 50 nm average edge size and occupy 40–45 vol.% of the alloy. Secondary ␥ precipitations with average diameter of 55–75 nm have spherical morphology and occupy 8–12 vol.%. The influence of the possible impurities on the precipitation of the secondary particles was not investigated in this study. However, as the nucleation of the secondary phase in this alloy is homogeneous not so much influences could be foreseen on the precipitations distribution or average size [11].
Fig. 6. Microstructure evolution of Rene-80 during creep at 871 ◦ C under stress of 475 MPa after (A) 0.1%, (B) 0.5%, (C) 1.7% creep strain and (D) after creep rupture.
J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
Fig. 3 shows the creep curves at 871 ◦ C under 290 and 475 MPa stresses respectively and Table 2 summarizes the creep data obtained from these creep tests. As it is evident from Fig. 3 and Fig. 4 and data from Table 2, increase in creep stress resulted in reduction of the creep life and increase in minimum creep rate. No regular changes were observed in reduction of area and elongation values. According to the creep results, sample at 781 ◦ C under 290 MPa had 629 h creep life in which 72% was the percentage of secondary creep life with steady creep rate. While at the same temperature under stress of 475 MPa samples showed 22% of its creep life at the secondary stage. Fig. 4 obtained from creep curves shows the change of creep rate, dε/dt versus normalized creep time, t/tf and compares the results of creep at 871 ◦ C under stresses of 290 and 475 MPa, respectively. As it can be seen from Fig. 4, with increase in creep-normalized time, creep rate in primary creep region showed decrease and then secondary creep was dominated. Extent of the secondary creep region was longer in creep at 290 MPa and shorter at 475 MPa. The reason for this behavior would be sought in the dominant mechanisms controlling creep deformation. Secondary creep rate or steady state creep rate is one of the most important areas for materials researchers and scientists. Some researchers believe in steady state creep rate and suggested that during the secondary creep, samples experience secondary creep with constant creep rate and life of this stage dominate almost most of the creep life [14,16,26,27]. While McLean and Dyson [28–30] proposed that superalloys have very short secondary creep region and actually there is just one transient region of secondary creep which can be considered as the minimum creep rate. They argued that as the microstructure of superalloys is not thermodynamically stable at high temperatures, it would be subject to degradation dur-
449
ing creep and result in the acceleration of the secondary creep rate leading to commence of tertiary creep. In this research, microstructure evolutions during different stages of the creep tests after completing of 0.1, 0.5, 1.7, and 4% creep strains as well as after creep rupture were investigated by SEM and TEM. Fig. 5 shows the microstructure evolution obtained by SEM at different stages of creep test at 871 ◦ C and 290 MPa. Dimensional measurement of primary and secondary ␥ particles from microstructure after heat treatment and after different stages of creep tests revealed that both primary and secondary particles were undergone coarsening at creep temperature, 871 ◦ C. The same microstructure evolution was detected in thermal exposure of Rene-80 by the authors [11]. Coarsening of ␥ as the secondary phase particle, by ostwald ripening mechanism, has been investigated by many researchers according to LSW theory and subsequent modifications [22–25]. In Rene-80 coalescence of primary ␥ particles during growth had significant influence on particle coarsening and accelerated the microstructure evolution substantially [11]. Fig. 5 shows that spherical secondary ␥ particles have grown and coarsened according to ostwald ripening mechanism from the very beginning of the creep temperature at 871 ◦ C and stress of 290 MPa. No coalescence was observed during coarsening at this condition for secondary precipitates (Fig. 5). The main feature of the evolution was that the density or population of secondary ␥ particles decreased substantially after 1.7% creep strain. As the creep life at 871 ◦ C and 475 MPa (Fig. 3) is relatively short in comparison to creep life at 871 ◦ C and 290 MPa, 8.3 and 629 h, respectively, microstructure evolution was not as much as the former creep test (compare Fig. 5 and Fig. 6). It can be seen from Fig. 6 that in this condition again primary ␥ particles growth and
Fig. 7. TEM microstructure during creep at 871 ◦ C under a stress of 290 MPa after 0.5% creep strain, (A) dislocation glide and climb and (B) dislocation pair movements in superalloy containing ordered particles (schematically) [29].
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Fig. 8. TEM microstructure during creep at 871 ◦ C under a stress of 290 MPa after 4% creep strain including particle shearing and secondary particle bypassing by orowan loop formation.
Fig. 9. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 0.1% creep strain.
J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
coalescence were observed during creep. Although secondary ␥ particles showed coarsening during creep, but creep life was not long enough to cause high inter-particle spacing (similar to what did happen after creep at 871 ◦ C and 290 MPa). In order to determine the creep mechanisms at different creep conditions, TEM investigation was performed on the interrupted creep samples at 871 ◦ C under stress of 290 MPa and the results after 0.5% and 4% creep strains as well as after rupture were reported (Fig. 7). Fig. 7A shows TEM microstructure of the samples after 0.5% creep strain which is equal to 40 h creep life. Dislocations during creep of Rene-80 at 871 ◦ C and 290 MPa were spread in matrix within the corridors between primary ␥ . Dislocations in the matrix gained pair configuration during movement, which has been well described by Nembach and coworkers [31,32]. This dislocation configuration was depicted schematically in Fig. 7B. As it can be seen from Fig. 7A glide of pair dislocations in corridors between ␥ particles in this condition of creep deformation is controlling the strain of the creep and time consuming thermally activated climb of dislocations over primary and secondary strengthening particles determine the creep rate. No sign of particle shearing by dislocation was observed after 0.5% creep strain at 871 ◦ C under stress of 290 MPa. Secondary ␥ had enough growth (Figs. 5 and 7) and dislocation bypassing by formation of orowan looping on the secondary particles would be another mechanism of deformation besides glide and climb. Fig. 8 shows the TEM microstructure after 4% creep strain during creep at 871 ◦ C and 290 MPa. Shearing of primary particles and dislocation bypassing of secondary ␥ particles by orowan loop formation were observed at this stage of creep. According to Figs. 3 and 4 it can be seen that 4% creep strain is the strain in which creep curve enter to the tertiary stage and creep
451
rate is no more constant. McLean [14] described that as long as the creep deformation is controlled by dislocation glide and climb at high temperatures, growth of the particles would not have influence on the creep rate and secondary creep stage with steady rate will be observed. He also mentioned that if growth of particles allows inter-particle spacing to be widening enough to lead dislocation bypassing by orowan looping, creep will enter to the tertiary stage and creep rate will steadily increase. Table 2 shows that crept sample at 871 ◦ C and 290 MPa will experience nearly 22% of its creep life as the tertiary creep, which means tertiary creep began after 78% of creep life equivalent to almost 4% creep strain. Then conclusion can be drawn that in dislocationparticle interaction point of view (Fig. 8), tertiary creep with ever increasing of creep rate will be commenced while particle shearing or particle by-passing by orowan looping were commenced. Fig. 9 shows the TEM microstructure of the creep samples during tests at 871 ◦ C and 475 MPa after 0.1% creep strain. As it can be seen from Fig. 9, after 0.1% creep strain dislocation gliding and propagating is evidence in matrix between ␥ particles (Fig. 9A) and shearing of particles has just started (Fig. 9B). Dislocation started to build up and form dislocation jungle (Fig. 10A) with increase in the creep strain up to 1.7% and as the result particle shearing was observed in Fig. 10B. Requirement of the particle shearing by the matrix dislocation is the built up of the dislocation behind the precipitates and exerting enough shear stress for dissociation of the dislocation and then subsequent particle shearing will happen [21]. Particle sharing in the similar condition was investigated by Lin and Wen [19–21] and they found that particle shearing is achieved by the super lattice intrinsic stacking fault (SISF) dislocation in this alloy.
Fig. 10. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 1.7% creep strain.
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Fig. 11. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 8.7% creep strain (creep rupture).
Further creep strain was dominated by the particle shearing by the SISF separated dislocation and it can be seen from Fig. 11, TEM after fracture of the specimen equal to 8.7% creep strain, that the interaction of the SISF dislocation inside of the ␥ particles is highly populated. 4. Conclusion Creep tests on Rene-80 nickel base superalloy were conducted at 871 ◦ C under stresses of 290 and 475 MPa. Microstructure evolution and creep deformation mechanisms were investigated during different stages of the creep tests. Conclusion could be summarized as the followings: (1) GE Class-A heat treatment on the Rene-80 produces a dual precipitation of ␥ particles including primary cubical particles with average edge size of 425 ± 25 nm and secondary spherical ␥ particles with average 70 nm in diameter. (2) Creep test at 871 ◦ C under stresses of 290 and 475 MPa caused substantial evolution in microstructure. (3) Secondary ␥ particles showed particle growth at early creep times while with further times at 871 ◦ C, these particles dissolved according to the LSW theory and scarified themselves for growth of the larger primary cubical precipitate. (4) Primary ␥ particles showed substantial coalescence during creep tests which was the dominant in growth mechanism of the primary ␥ . (5) Creep deformation at 871 ◦ C under stress of 290 MPa was dominantly controlled by the glide and climb of the matrix dislocations. As soon as particle cutting or by-passing by the
dislocation were commenced, creep curve has been entered to the tertiary creep stage. (6) Creep samples at 871 ◦ C under stress of 475 MPa were experienced particle shearing very earlier when comparing normalized creep life. This behavior was approved by the TEM investigation and it was found that it was in coincident with commence of the tertiary creep stage. Acknowledgements The authors are grateful to Professor M. Mclean, Dr. M. Ardakani from Imperial College University London for the provision of SEM and TEM laboratory facilities and valuable guide ness and British Council for financial support. References [1] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, Inc., 1987. [2] M.J. Donachie, S.J. Donachie, Superalloys a Technical Guide, 2nd ed., ASM International, 2002. [3] P. Beardmore, R.G. Davies, T.L. Johnston, Trans. Metal. Soc. AIME 254 (1969) 1537–1545. [4] D.P. Pope, S.S. Ezz, Int. Met. Rev. 29 (1984) 136–167. [5] B.F. Dyson, M. McLean, Proceedings on “Microstructural Stability of Creep resistant alloys for High Temperature Plant application”, IOM, London, 1998. [6] A.K. Koul, W. Wallace, Metall. Trans. A. 13A (1982) 673–675. [7] General Electric Specification no. GE C50 TF28, Rene-80 Investment Vacuum Cast Turbine Blades and Vanes, Oct. 1976. [8] W.D. Klopp, Nonferrous Alloys, Code 4214, NASA Publication, 1984. [9] Earl W. Ross, Met. Prog. (1971) 93–94. [10] J. Safari, S. Nategh, M. McLean, Mater. Sci. Technol. 22 (2006) 888–898. [11] J. Safari, S. Nategh, J. Mater. Proc. Technol. 176 (2006) 240–250. [12] D.-C. Madeleine, The Microstructure of Superalloys, OPA (Overseas Publishers Association), Amsterdam B.V., 1997.
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[17]
[18]
[19] [20] [21]
A. Kelly, R.B. Nicholson, Prog. Mater. Sci. (1963) 151–383. D. McLean, Deformation at high temperatures, Metall. Rev. 7 (1962) 481–527. R. Lagneborg, The Donald McLean Symposium, 1996, pp. 109–119. L.M. Brown an, R.K. Ham, Strengthening Methods in Crystals, Chapter II, Dislocation-Particle Interactions, Applied Science Publishers Ltd., London, 1971, pp. 9–130. Dongliang (T.L. Lin), Mao Wen, Dislocation structure due to high temperature deformation in ␥ phase of a nickel-base superalloy, Mater. Sci. Eng. A, 113 (1989) 207–214. Dongliang Lin (T.L. Lin), Mao Wen, Dislocation structure due to high temperature deformation in ␥ phase of a nickel-base superalloy, Acta Metall. 37 (1989) 3099–3105. T.L. Lin, Mao Wen, Mater. Sci. Eng. A 128 (1990) 23–31. S. Nategh, S.A. Sajjadi, Mater. Sci. Eng. A 339 (2003) 103–108. S.A. Sajjadi, S. Nategh, Mater. Sci. Eng. A 307 (2001) 158–164.
[22] [23] [24] [25] [26] [27] [28]
[29] [30] [31] [32]
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A.J. Ardell, Acta Metall. 16 (1968) 511–516. A.C. Lund, P.W. Voorhees, Acta Mater. 50 (2002) 2085–2098. C.K.L. Davies, P. Nash, R.N. Stevens, Acta Metall. 28 (1980) 179–189. R.A. MacKay, M.N. Nathal, Acta Metal. Mater. 38 (1990) 993–1005. G.R. Leverant, B.H. Kear, J.M. Oblak, Metall. Trans. 4 (1973) 355–362. G.R. Leverant, B.H. Kear, Metall. Trans. 1 (1970) 491–498. B.F. Dyson, M. McLean, in: A. Strang, J. Cawley, G.W. Greenwood (Eds.), Proceedings of the Conference on Microstructural Stability of Creep Resistant Alloys for High Temperature Plant Applications, Institute of Materials, London, 1998, pp. 371–393. B.F. Dyson, M. McLean, J. Eng. Mater. Technol. 122 (2000) 273–278. B.F. Dyson, M. McLean, Acta Metall. 31 (1983) 17–27. E. Nembach, Particle Strengthening of Metals and Alloys, John Wiley & Sons, Inc., 1997. V. Mohles, D. Ronnpagel, E. Nembach, Comput. Mater. Sci. 16 (1999) 144–150.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure evolution and its influence on deformation mechanisms during high temperature creep of a nickel base superalloy Javad Safari a,∗ , Saeed Nategh b,1 a b
Materials Science and Engineering Department, Shahid Chamran University, Ahwaz, Iran Materials Science and Engineering Department, Sharif University of Technology, P.O. Box 11365-9466, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 20 August 2007 Received in revised form 27 August 2008 Accepted 8 September 2008 Keywords: Nickel base superalloys Creep Microstructure Dislocation Deformation mechanism Transmission electron microscopy
a b s t r a c t The interaction of dislocation with strengthening particles, including primary and secondary ␥ , during different stages of creep of Rene-80 was investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). During creep of the alloy at 871 ◦ C under stress of 290 MPa, the dislocation network was formed during the early stages of creep, and the dislocation glide and climb process were the predominant mechanism of deformation. The density of dislocation network became more populated during the later stages of the creep, and at the latest stage of the creep, primary particles shearing were observed alongside with the dislocation glide and climb. Shearing of ␥ particles in creep at 871 ◦ C under stress of 475 MPa was commenced at the earlier creep times and governed the creep deformation mechanism. In two levels of examined stresses, as far as the creep deformation was controlled by glide and climb, creep curves were found to be at the second stage of creep and commence of the tertiary creep, with increasing creep rate, were found to be in coincidence with the particles shearing. Microstructure evolution, with regard to ␥ strengthening particles, led to particles growth and promoted activation of other deformation mechanisms such as dislocation bypassing by orowan loop formation. Dislocation-secondary ␥ particles interaction was detected to be the glide and climb at the early stages of creep, while at the later stages, the dislocation bypassed the secondary precipitation by means of orowan loops formation, as the secondary particle were grown and the mean inter-particle distance increased. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Cast nickel base superalloys containing medium to high volume fractions of the ordered ␥ phase are used as blades and other high temperature components in aero- and power-generating gas turbines [1,2]. High mechanical properties of Ni base superalloys at high temperatures are mostly attributed to the characteristics of the ␥ precipitates. In Ni–Al–Ti system, ␥ is an ordered phase with L12 structure and nominal composition of Ni3 (Al,Ti) that is embedded in the disordered (FCC) ␥ matrix [3–6]. Polycrystalline, and recently directionally solidified, Rene-80 (GE trademark) with bimodal precipitation of ␥ particles, is frequently used as first stage aero-gas turbine blades and is subject to high temperature degradation mechanisms mostly creep and low
∗ Corresponding author. Tel.: +98 611 333 0015; fax: +98 21 2205 7662. E-mail addresses: [email protected] (J. Safari), [email protected] (S. Nategh). 1 Tel.: +98 912 384 2146; fax: +98 21 6600 5717. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.017
cycle fatigue [7–9]. Cast microstructure evolution during different stages of heat treatment and subsequent microstructure degradation due to thermal exposure at high temperatures were studied by the authors and were reported elsewhere [10,11]. Creep deformation mechanisms of superalloys are controlled by the interaction of dislocation with primary and secondary ␥ particles during dislocation creep [12–16]. Lin and Wen [17–19] have performed some investigation on creep and tensile deformation mechanisms of directionally solidified Rene-80. Their results may be summarized as follows. In creep deformation at 760 ◦ C under stress of 618 MPa, ␥ shearing occurs by viscous slip of pairs of superlattice intrinsic stacking faults (SISF)-separated by 1/31 1 2 super partials with net slip vector 1 1 0, while the cores of the 1/31 1 2 superpartials do not appear to be separated [17–19]. Creep deformation at 982 ◦ C under stress of 190 MPa is controlled by glide and climb of dislocation from loose dislocation network around ␥ particles. Almost the same results were obtained by Sajjadi and Nategh on GDT-111 Nickel base superalloy. They referred to this alloy as the new modification of Rene-80 [20].
446
J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
Fig. 1. Class-A heat treatment of the Rene-80 per GE-C50TF28 [7].
2. Experiments 2.1. Material and heat treatment The material used in this research was the polycrystalline Rene80 nickel base superalloy [7,8]. Cylindrical bars, 12 mm in diameter and 120 mm in length, were produced by vacuum precision casting. Casting was carried out using previously determined parameters which produced components with acceptable defect levels by the GE standard assessment [7]. Soundness of the samples after casting has been approved by dye penetrant and radiography as the NDT examinations. The bulk composition was determined by XRF analysis. The analysis was repeated at least five times for each casting and the averaged results are presented in Table 1. Samples were then heat treated in a vacuum between 10−3 and 10−4 mbar to the class-A GE specification [7]. This heat treatment consists of four stages; 2 h at 1204 ◦ C (solution heat treatment), 4 h at 1093 ◦ C (primary aging), 4 h at 1054 ◦ C (coating heat treatment), and 16 h at 843 ◦ C (Aging). These stages are shown graphically in Fig. 1.
Fig. 3. Rene-80 creep curves at 871 ◦ C under two stress levels of (A) 290 MPa and (B) 475 MPa.
order to investigate the microstructure evolution and deformation mechanisms during creep, creep tests were interrupted at different stages of tests and samples were prepared for investigation with high resolution scanning electron microscopy (SEM) and transmission electron microscopy (TEM).
2.2. Creep tests Creep tests at 871 ◦ C and two different stress levels, 290 and 475 MPa were carried out according to the GE C50 TF28 specification and ASTM E139. Samples temperature was controlled within ±2 ◦ C by attached thermocouples and 1 h soaking time was applied on samples before applying the creep stress. Creep extension was measured by the extensometers placed on samples projections. In
Fig. 2. Microstructure of Rene-80 after GE class-A heat treatment.
Fig. 4. Creep rate versus normalized creep time for creep tests at 871 ◦ C under stress of (A) 290 MPa and (B) 475 MPa.
J. Safari, S. Nategh / Materials Science and Engineering A 499 (2009) 445–453
447
Table 1 Comparison of Rene-80 standard composition and experimental analysis result for one casting. wt%
Ni
Cr
Co
Ti
Al
Mo
W
C
B
Zr
P
S
Min [7] Max [7] Exp.a
Bal. Bal. Bal.
13.70 14.30 14.24
9.00 10.00 9.45
4.80 5.20 5.12
2.80 3.20 3.05
3.70 4.30 3.96
3.70 4.30 4.12
0.15 0.19 0.17
0.01 0.02 0.014
0.02 0.1 0.07
– 0.015 0.010
– 0.0075 0.0050
a
Average of at least five analyses.
Fig. 5. Microstructure evolution of Rene-80 during creep at 871 ◦ C under stress of 290 MPa after (A) 0.1%, (B) 0.5%, (C) 1.7%, (D) 4% creep strain and (E) after creep rupture of samples.
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Table 2 Creep data of Rene-80 at 871 ◦ C at different stresses. Stress, MPa
Rupture time, h
Min. creep rate, s−1
Reduction of area, %
Elongation, %
Percentage of secondary creep life, %
290 475
629 8.43
1.90E−08 1.25E−06
14.12 14.23
5.51 5.74
72 22
2.3. Microstructure characterization Samples were sectioned for preparation of metallography and microstructure evolution. Microstructure was studied by field emission high resolution SEM and TEM. Sample preparation for electron microscopy was carried out by the lapping, electro polishing with a solution of 10% HCl + methanol at −40 ◦ C and 25 V for 20 s and electro etching with 170 mL phosphoric acid + 10 mL sulfuric acid and 16 g chromium trioxide at room temperature and 6 V for 4 s. The H3 PO4 –H2 SO4 –CrO3 etchant attacks the gamma solid solution and forms an anodic film on the gamma areas. This anodic film permits the differentiation between ␥ and ␥ in image analysis. The fact that this etch attacks gamma, makes it the preferred etch for bringing out the very fine ␥ [10]. As the quantitative metallographic investigation is basically a statistical measurement, a large number of images at different magnifications for all conditions of aging were analyzed. However, only one of each is shown here as being representative. Measurement of the ␥ particle characteristics was performed on the SEM photos using the UTHSCSA Image Tool software, a semi-quantitative image analyzer program
(see web site address: http://ddsdx.uthscsa.edu/dig/itdesc.html for the software information and download).
3. Results and discussion Fig. 2 shows the microstructure of Rene-80 superalloy after completion of class-A heat treatment according to GE C50 TF28. Microstructure of Rene-80 consists of dual precipitation of ␥ strengthening particles in the austenitic ␥ matrix. Primary ␥ particles with cubic morphology, which have been formed at the primary aging stage of heat treatment, have 450 ± 50 nm average edge size and occupy 40–45 vol.% of the alloy. Secondary ␥ precipitations with average diameter of 55–75 nm have spherical morphology and occupy 8–12 vol.%. The influence of the possible impurities on the precipitation of the secondary particles was not investigated in this study. However, as the nucleation of the secondary phase in this alloy is homogeneous not so much influences could be foreseen on the precipitations distribution or average size [11].
Fig. 6. Microstructure evolution of Rene-80 during creep at 871 ◦ C under stress of 475 MPa after (A) 0.1%, (B) 0.5%, (C) 1.7% creep strain and (D) after creep rupture.
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Fig. 3 shows the creep curves at 871 ◦ C under 290 and 475 MPa stresses respectively and Table 2 summarizes the creep data obtained from these creep tests. As it is evident from Fig. 3 and Fig. 4 and data from Table 2, increase in creep stress resulted in reduction of the creep life and increase in minimum creep rate. No regular changes were observed in reduction of area and elongation values. According to the creep results, sample at 781 ◦ C under 290 MPa had 629 h creep life in which 72% was the percentage of secondary creep life with steady creep rate. While at the same temperature under stress of 475 MPa samples showed 22% of its creep life at the secondary stage. Fig. 4 obtained from creep curves shows the change of creep rate, dε/dt versus normalized creep time, t/tf and compares the results of creep at 871 ◦ C under stresses of 290 and 475 MPa, respectively. As it can be seen from Fig. 4, with increase in creep-normalized time, creep rate in primary creep region showed decrease and then secondary creep was dominated. Extent of the secondary creep region was longer in creep at 290 MPa and shorter at 475 MPa. The reason for this behavior would be sought in the dominant mechanisms controlling creep deformation. Secondary creep rate or steady state creep rate is one of the most important areas for materials researchers and scientists. Some researchers believe in steady state creep rate and suggested that during the secondary creep, samples experience secondary creep with constant creep rate and life of this stage dominate almost most of the creep life [14,16,26,27]. While McLean and Dyson [28–30] proposed that superalloys have very short secondary creep region and actually there is just one transient region of secondary creep which can be considered as the minimum creep rate. They argued that as the microstructure of superalloys is not thermodynamically stable at high temperatures, it would be subject to degradation dur-
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ing creep and result in the acceleration of the secondary creep rate leading to commence of tertiary creep. In this research, microstructure evolutions during different stages of the creep tests after completing of 0.1, 0.5, 1.7, and 4% creep strains as well as after creep rupture were investigated by SEM and TEM. Fig. 5 shows the microstructure evolution obtained by SEM at different stages of creep test at 871 ◦ C and 290 MPa. Dimensional measurement of primary and secondary ␥ particles from microstructure after heat treatment and after different stages of creep tests revealed that both primary and secondary particles were undergone coarsening at creep temperature, 871 ◦ C. The same microstructure evolution was detected in thermal exposure of Rene-80 by the authors [11]. Coarsening of ␥ as the secondary phase particle, by ostwald ripening mechanism, has been investigated by many researchers according to LSW theory and subsequent modifications [22–25]. In Rene-80 coalescence of primary ␥ particles during growth had significant influence on particle coarsening and accelerated the microstructure evolution substantially [11]. Fig. 5 shows that spherical secondary ␥ particles have grown and coarsened according to ostwald ripening mechanism from the very beginning of the creep temperature at 871 ◦ C and stress of 290 MPa. No coalescence was observed during coarsening at this condition for secondary precipitates (Fig. 5). The main feature of the evolution was that the density or population of secondary ␥ particles decreased substantially after 1.7% creep strain. As the creep life at 871 ◦ C and 475 MPa (Fig. 3) is relatively short in comparison to creep life at 871 ◦ C and 290 MPa, 8.3 and 629 h, respectively, microstructure evolution was not as much as the former creep test (compare Fig. 5 and Fig. 6). It can be seen from Fig. 6 that in this condition again primary ␥ particles growth and
Fig. 7. TEM microstructure during creep at 871 ◦ C under a stress of 290 MPa after 0.5% creep strain, (A) dislocation glide and climb and (B) dislocation pair movements in superalloy containing ordered particles (schematically) [29].
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Fig. 8. TEM microstructure during creep at 871 ◦ C under a stress of 290 MPa after 4% creep strain including particle shearing and secondary particle bypassing by orowan loop formation.
Fig. 9. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 0.1% creep strain.
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coalescence were observed during creep. Although secondary ␥ particles showed coarsening during creep, but creep life was not long enough to cause high inter-particle spacing (similar to what did happen after creep at 871 ◦ C and 290 MPa). In order to determine the creep mechanisms at different creep conditions, TEM investigation was performed on the interrupted creep samples at 871 ◦ C under stress of 290 MPa and the results after 0.5% and 4% creep strains as well as after rupture were reported (Fig. 7). Fig. 7A shows TEM microstructure of the samples after 0.5% creep strain which is equal to 40 h creep life. Dislocations during creep of Rene-80 at 871 ◦ C and 290 MPa were spread in matrix within the corridors between primary ␥ . Dislocations in the matrix gained pair configuration during movement, which has been well described by Nembach and coworkers [31,32]. This dislocation configuration was depicted schematically in Fig. 7B. As it can be seen from Fig. 7A glide of pair dislocations in corridors between ␥ particles in this condition of creep deformation is controlling the strain of the creep and time consuming thermally activated climb of dislocations over primary and secondary strengthening particles determine the creep rate. No sign of particle shearing by dislocation was observed after 0.5% creep strain at 871 ◦ C under stress of 290 MPa. Secondary ␥ had enough growth (Figs. 5 and 7) and dislocation bypassing by formation of orowan looping on the secondary particles would be another mechanism of deformation besides glide and climb. Fig. 8 shows the TEM microstructure after 4% creep strain during creep at 871 ◦ C and 290 MPa. Shearing of primary particles and dislocation bypassing of secondary ␥ particles by orowan loop formation were observed at this stage of creep. According to Figs. 3 and 4 it can be seen that 4% creep strain is the strain in which creep curve enter to the tertiary stage and creep
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rate is no more constant. McLean [14] described that as long as the creep deformation is controlled by dislocation glide and climb at high temperatures, growth of the particles would not have influence on the creep rate and secondary creep stage with steady rate will be observed. He also mentioned that if growth of particles allows inter-particle spacing to be widening enough to lead dislocation bypassing by orowan looping, creep will enter to the tertiary stage and creep rate will steadily increase. Table 2 shows that crept sample at 871 ◦ C and 290 MPa will experience nearly 22% of its creep life as the tertiary creep, which means tertiary creep began after 78% of creep life equivalent to almost 4% creep strain. Then conclusion can be drawn that in dislocationparticle interaction point of view (Fig. 8), tertiary creep with ever increasing of creep rate will be commenced while particle shearing or particle by-passing by orowan looping were commenced. Fig. 9 shows the TEM microstructure of the creep samples during tests at 871 ◦ C and 475 MPa after 0.1% creep strain. As it can be seen from Fig. 9, after 0.1% creep strain dislocation gliding and propagating is evidence in matrix between ␥ particles (Fig. 9A) and shearing of particles has just started (Fig. 9B). Dislocation started to build up and form dislocation jungle (Fig. 10A) with increase in the creep strain up to 1.7% and as the result particle shearing was observed in Fig. 10B. Requirement of the particle shearing by the matrix dislocation is the built up of the dislocation behind the precipitates and exerting enough shear stress for dissociation of the dislocation and then subsequent particle shearing will happen [21]. Particle sharing in the similar condition was investigated by Lin and Wen [19–21] and they found that particle shearing is achieved by the super lattice intrinsic stacking fault (SISF) dislocation in this alloy.
Fig. 10. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 1.7% creep strain.
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Fig. 11. TEM microstructure of creep sample at 871 ◦ C and 475 MPa after 8.7% creep strain (creep rupture).
Further creep strain was dominated by the particle shearing by the SISF separated dislocation and it can be seen from Fig. 11, TEM after fracture of the specimen equal to 8.7% creep strain, that the interaction of the SISF dislocation inside of the ␥ particles is highly populated. 4. Conclusion Creep tests on Rene-80 nickel base superalloy were conducted at 871 ◦ C under stresses of 290 and 475 MPa. Microstructure evolution and creep deformation mechanisms were investigated during different stages of the creep tests. Conclusion could be summarized as the followings: (1) GE Class-A heat treatment on the Rene-80 produces a dual precipitation of ␥ particles including primary cubical particles with average edge size of 425 ± 25 nm and secondary spherical ␥ particles with average 70 nm in diameter. (2) Creep test at 871 ◦ C under stresses of 290 and 475 MPa caused substantial evolution in microstructure. (3) Secondary ␥ particles showed particle growth at early creep times while with further times at 871 ◦ C, these particles dissolved according to the LSW theory and scarified themselves for growth of the larger primary cubical precipitate. (4) Primary ␥ particles showed substantial coalescence during creep tests which was the dominant in growth mechanism of the primary ␥ . (5) Creep deformation at 871 ◦ C under stress of 290 MPa was dominantly controlled by the glide and climb of the matrix dislocations. As soon as particle cutting or by-passing by the
dislocation were commenced, creep curve has been entered to the tertiary creep stage. (6) Creep samples at 871 ◦ C under stress of 475 MPa were experienced particle shearing very earlier when comparing normalized creep life. This behavior was approved by the TEM investigation and it was found that it was in coincident with commence of the tertiary creep stage. Acknowledgements The authors are grateful to Professor M. Mclean, Dr. M. Ardakani from Imperial College University London for the provision of SEM and TEM laboratory facilities and valuable guide ness and British Council for financial support. References [1] C.T. Sims, N.S. Stoloff, W.C. Hagel, Superalloys II, John Wiley & Sons, Inc., 1987. [2] M.J. Donachie, S.J. Donachie, Superalloys a Technical Guide, 2nd ed., ASM International, 2002. [3] P. Beardmore, R.G. Davies, T.L. Johnston, Trans. Metal. Soc. AIME 254 (1969) 1537–1545. [4] D.P. Pope, S.S. Ezz, Int. Met. Rev. 29 (1984) 136–167. [5] B.F. Dyson, M. McLean, Proceedings on “Microstructural Stability of Creep resistant alloys for High Temperature Plant application”, IOM, London, 1998. [6] A.K. Koul, W. Wallace, Metall. Trans. A. 13A (1982) 673–675. [7] General Electric Specification no. GE C50 TF28, Rene-80 Investment Vacuum Cast Turbine Blades and Vanes, Oct. 1976. [8] W.D. Klopp, Nonferrous Alloys, Code 4214, NASA Publication, 1984. [9] Earl W. Ross, Met. Prog. (1971) 93–94. [10] J. Safari, S. Nategh, M. McLean, Mater. Sci. Technol. 22 (2006) 888–898. [11] J. Safari, S. Nategh, J. Mater. Proc. Technol. 176 (2006) 240–250. [12] D.-C. Madeleine, The Microstructure of Superalloys, OPA (Overseas Publishers Association), Amsterdam B.V., 1997.
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