MICROSCOPY RESEARCH AND TECHNIQUE 78:336–342 (2015)
Effect of Titanium Carbonitride (Ti(C,N)) Decomposition on Failure Mechanisms in Inconel 617 alloy RAM KRISHNA,1,2 SARAH V. HAINSWORTH,1* SIMON P. A. GILL,1 AND HELEN V. ATKINSON1 1
Department of Engineering, University of Leicester, University Road, Leicester, LE1 7RH, UK Now at Dalton Cumbrian Facility, The University of Manchester, Westlakes Science and Technology Park, Moor Row, Whitehaven, Cumbria, CA24 3HA, UK
2
KEY WORDS
Carbide decomposition reaction; Creep; Hardness; Intergranular fracture; Inconel 617; TEM
ABSTRACT Titanium Carbonitride (Ti(C,N)) decomposition in Inconel 617 alloy creepexposed at 650 C for 574 hours is reported using analytical electron microscopy techniques. Crenriched M23C6-type carbides enveloped in fine gamma prime particles thought to be precipitated from the decomposition reaction are observed in the alloy. The morphology of the M23C6 carbides is irregular and blocky and the particle size up to 5lm, whereas the morphology of gamma prime particles is mostly spherical and up to 30 nm in size. Intergranular carbides are mostly secondary precipitates of the M23Cc type (M predominantly Cr) and these respond to solution heat treatment and precipitate on the grain boundaries as a result of ageing. The ability of intragranular MX to decompose is sensitive to the N content, high N resists decomposition. Decomposed intragranular MX provides an excess source of C which can react locally with Cr to form heat treatable intragranular fine Cr23C6 precipitates. M6C can segregate in interdendritic locations during melting which may be the reason for high content of Mo in M23C6. These precipitates are generally very small and contribute to an additional hardening effect and are the reason for the onset of voiding and cracking along the grain boundaries that ultimately lead to a reduced creep rupture life. Microsc. Res. Tech. 78:336–342, 2015. V 2015 Wiley Periodicals, Inc. C
INTRODUCTION Future generation steam turbines and their components are expected to confront ultra supercritical conditions i.e., working steam conditions up to 760 C or more and operating pressure up to 35 MPa, in order to improve the thermal efficiency and reduce the environmental emissions. To accommodate the imposed extended working conditions, the currently used ferritic and austenitic alloys will need to be substituted by alloys with better creep strength and environmental resistance. Nickel-based alloys are likely to be considered as a potential alternatives. Inconel 617 alloy (IN617) belongs to the family of nickel-based alloys and is being actively evaluated as a potential candidate alloy for elevated temperature applications owing to its excellent high-temperature strength, adequate metallurgical stability, good resistance to creep, and better resistance to corrosion and oxidation (Penkalla et al., 2001; Viswanathan et al., 2005; Vanstone, 2000). The alloying elements Cr (23 wt%), Co (12 wt%), and Mo (9 wt%) in IN617 contribute towards solid solution strengthening, while small quantities of Ti and Al contribute additional precipitation strengthening (Hosier and Tillack, 1972; Krishna et al., 2010; Mankins et al., 1974; Zhao et al., 2004). Since the inception of this alloy it has been used for high temperature structural application due to its balanced alloying elements. However, degradation in mechanical properties and creep strength during thermal and creep exposure are significant as a result of microstructural evolutions and are C V
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sensitive to operating conditions (Krishna et al., 2010; Zhao et al., 2004). The microstructural evolutions in nickel-based alloys including IN617 comprises gamma prime (g0 ) coarsening (Hosier and Tillack, 1972; Mankins et al., 1974), secondary precipitation of M23C6 carbide network (Krishna et al., 2010; Zhao et al., 2004), nucleation and growth of topologically closed packed (TCP) phase such as m-phases (Kihara et al., 1980; Krishna et al., 2013) and precipitation of discrete phases by TiC or Ti carbonitride degeneration reactions (He et al., 2005; Lvov et al., 2004; Qin et al., 2007; Xuebing et al., 1998). Most of these modifications in microstructure during service are detrimental due to the fact that their occurrence depletes potent solidsolution strengtheners from the austenite g-matrix, which results in softening of the matrix at operating temperature. All of the above modifications are also not very common in the alloy IN617 and literature reports on IN617 are relatively sparse (Davis, 1997; El-Magd et al., 1996; He et al., 2005; Kihara et al., 1980; Krishna et al., 2010, 2013; Lvov et al., 2004; Mankins et al., 1974; Qin et al., 2007; Xuebing et al., 1998; Zhao et al., 2004). However the decomposition of primary TiC and Ti carbonitrides has been reported in some nickel-based alloys. The degeneration reaction *Correspondence to: Sarah V. Hainsworth, Department of Engineering, University of Leicester, University Road, Leicester LE1 7RH, UK. E-mail: [email protected] Received 1 February 2014; accepted in revised form 8 March 2014 REVIEW EDITOR: Dr. Chuanbin Mao DOI 10.1002/jemt.22359 Published online 1 April 2015 in Wiley Online Library (wileyonlinelibrary.com).
EFFECT OF TIC CARBIDE DECOMPOSITION
Elements
Ni
Cr
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TABLE 1. Chemical composition of as-received Inconel 617 alloy (wt%) Mo Co Al Ti Fe
C
Si
W
0.05
0.02
Wt%
Bal.
22.8
9.0
11.9
1.15
0.48
0.35
0.06
Elements
Mn
Cu
Zr
P
V
N
Nb
S
Wt%
0.02
0.005
0.004
0.003
0.001
0.013
0.005
0.005
has either a beneficial or a detrimental effect on creep and mechanical properties depending on the morphology and distribution of the resulting precipitates, their location, and the alloys exposure conditions (El-Magd et al., 1996; He et al., 2005; Qin et al., 2007; Lvov et al., 2004; Xuebing et al., 1998). Primary Ti(C,N) particles are a major source of carbon in most nickel-based alloys below 980 C. However, they decompose slowly during thermal treatment and apparently at a faster rate during creep exposure condition, releasing carbon for several important reactions, including what has been described as the carbide degeneration reaction (Ray et al., 2006; Ross and Sims, 1987). This leads to the formation of Crenriched M23C6 carbides and g0 following the reaction: TiC 1 gmatrix ! M23C6 1 g0 . This reaction has been reported for a range of alloys from the nickel-based family K452, IN738 and Udimet-520 (Qin et al. 2007; Wang et al. 2012). In nickel-based M963 cast superalloy, primary TiC carbide degeneration reaction is observed; but, with the formation of M6C type carbide and g0 particles, TiC 1 g ! M6C 1 g0 (He et al., 2005), in contrast to the degeneration reactions in studies of Qin et al. (2007) and Wang et al. (2012). Further, in nickel-based precipitate hardened superalloys such as IN738 and GTD111 alloys primary TiC carbide degenerates as M23C6 carbides and h phase through the following reaction: TiC 1 g/g0 ! M23C6 1 h. In a very recent study, Wang et al. (2012) reported TiC carbide degeneration in nickel-based K444 superalloy. In this mechanism TiC carbide degenerated into two different carbides, M23C6 and M6C, as described in reaction: TiC 1 g! M23C6 1 M6C 1 h. Further, it was observed that the occurrence of this degeneration reaction is was detrimental to the mechanical properties. Formation of the h phase, M6C and M23C6 precipitates in K444 alloy is attributed to long-term thermal exposure (up to 10,000 h) and higher content of alloy elements (W, Ti, Al, and Hf) in the K444 alloy composition. Thus, the carbide reactions lead to the precipitation of either M23C6, g0 , M6C or h phase at intergranular locations and play a role in determining the long-term mechanical properties, creep performance, and fracture mechanisms of the alloy (He et al., 2005; Kihara et al., 1980; Krishna et al., 2013; Lvov et al., 2004; Qin et al., 2007; Wu et al., 2008). TiC is not in itself detrimental to the mechanical properties of the alloy as discussed in Mankins et al. (1974) study. Nucleation and growth of the finer size M23C6 carbides and g0 precipitates are supposed to be more prominent in strengthening the alloy than larger Ti(C,N) carbides because of the finer size of the degenerated precipitates. However, the formation of M23C6 carbides at and near the grain boundaries introduces brittleness in the alloy and results in failure due to de-cohesion and triple point cracking at the carbide-matrix interface (Lvov et al., 2004; Wang et al., 2012). It is important to note here Microscopy Research and Technique
B 0.0002
that M23C6 and M6C precipitates are chromium (Cr) enriched carbides, where M denotes elements of Cr, Ni, Mo, Co etc. The alloys refered to here are used in the cast condition whereas IN617 is a thermo-mechanically processed alloy which influences its microstructure and properties. In the present study the aim is to develop understanding of the effect of Ti (C,N) carbide degeneration in Inconel 617 alloy in terms of creep failure mechanism and variation in hardness as a mechanical property. In this article we seek to delineate the Ti(C,N) decomposition in creep-exposed IN617 alloy at 650 C using analytical microscopy. Further, formation of g0 and M23C6 through the Ti(C,N) degeneration reaction is illustrated using transmission electron microscopy in the creep-exposed alloy. MATERIAL AND EXPERIMENTAL METHODS The complete nominal chemical composition of the Inconel 617 studied in the present investigation is presented in Table 1. A standard creep specimen was prepared from a solution annealed forged rod of IN617. The ASTM grain size number and hardness of the as-received specimen were 5.0 and 186.6 6 5 HV20, respectively. Forged rods of Inconel 617 solution annealed at 1100 C/3 h/WQ were given a heat treatment at 670 C for 10 h/air cooled (AC), which usually precipitates intergranular primary carbides (Krishna et al., 2010). Creep specimens were then machined from the heattreated forged rods. The creep rupture test was conducted at 650 C under an applied stress in air by Alstom Power (Rugby, UK) and the creep test was continued until fracture occurred. Four different creep ruptured samples were examined for hardness and microstructural modifications using Vickers hardness tests and electron microscopy. X-ray diffractometry (XRD) was performed on asreceived and creep-exposed samples to characterize the phases present in the alloy, using a Philips PW ˚) 1716 diffractometer with CuKa1 radiation (k 5 1.54A in the angular 2u range between 2 to 125 , with a step size of 0.0096 . Vickers hardness tests were conducted using a Vickers hardness testing machine, with a 20 kgf load on ground and polished longitudinal cross-sections of the creep fractured specimen. The failure mechanism of the creep-exposed sample was evaluated using scanning electron microscopy. The standard polishing technique was to grind with successively finer grades of SiC emery papers and diamond suspensions to prepare specimens for metallography. Scanning electron microscopy (SEM) was performed on specimens in the polished but unetched condition. TEM foils were prepared using the standard method as discussed in Krishna et al. (2013).
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Fig. 1. (a) Scanning electron micrograph of as-received IN617 alloy shows Cr-enriched M23C6 carbides and Ti-enriched Ti(C,N) carbonitride. (b) Scanning electron micrograph of creep-exposed IN617 alloy shows M23C6 carbides precipitated along and near to grain boundaries, and slip lines. The globular particles identified in figure 1a as M23C6 could possibly be M6C particles which are not uncommon
in Ni-Cr-Mo alloys and occur in original interdendritic regions often near Ti(CN) as in the figures, particularly 1a) however M6C is Mo enriched in IN617 and EDX showed these to be Cr-rich (see discussion). [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
RESULTS AND DISCUSSION Alloying elements (Cr, Co, Mo, Al, Ti, and Fe) in Inconel 617 alloy are responsible for high temperature properties that occur through the formation of different types of carbides, carbo-nitrides, and intermetallic phases. An additional hardening arises from the localized precipitation of carbides/carbo-nitrides during ageing treatment (Hosier and Tillack, 1972; Krishna et al., 2010; Mankins et al., 1974; Zhao et al., 2004). The morphology (shape, size, and geometry) and distribution of these precipitates play an influential role in controlling the alloy mechanical and creep properties. Embedded intermetallic g0 precipitates in FCC austenite-g matrix and secondary carbide precipitates act as barriers to the movement of dislocations and contribute to hardening (Kihara et al., 1980; Krishna et al., 2010; Zhao et al., 2004). The microstructure of as-received IN617 alloy in the pre-aged condition (10h /670 C) consists of sparsely distributed blocky shape primary carbide particles of Ti(C,N) with an average size of 5 mm, a few Crenriched M23C6 type carbides (M@Cr, Ni, Mo, Co) with an average size of 3 mm and ordered g0 particles with an average size of 12.5 nm (Krishna et al., 2013; Mankins et al., 1974; Zhao et al., 2004). Figure 1a shows the morphology and location of M23C6 and Ti (C,N) carbonitrides. The particles identified as M23C6 were found to be Cr-rich. Akbari-Garakani and Mehdizadeh (2011) and Roy et al (2009) note that two types of carbide can co-exist on the grain boundaries and also within the grains on IN617 after exposure at high temperatures with M23C6 being Cr-rich and M6C having a higher Mo content. Since these particles were found to have Cr contents greater than 50% they are labelled here as M23C6 particles. M6C particles can occur in Mo containing alloys at residual interdendritic sites however reports of M6C in these alloys report them
as those that have a higher Mo content than Cr content. The precipitates are identified with Energy Dispersive X-Ray (EDX) analysis using a Princeton Gamma Tech system attached to SEM. Typical EDX analyses for primary M23C6 and TiC carbides in asreceived sample are shown in Figure 2 (carbon and nitrogen were excluded from the quantitative analysis). Primary TiC carbides may also contain some nitrogen and can form both carbides and nitrides which have isostructural cubic structures with lattice parameters a 5 0.43–0.47 nm and a 5 0.424 nm, respectively (Davis, 1997). Thus it is more accurate to refer to carbo-nitrides, Ti(C,N). Primary carbo-nitrides Ti(C,N) of blocky morphology are observed at or near to grain boundaries in the asreceived sample (Fig. 1a). In the creep-exposed sample M23C6 (Cr-enriched) carbides are observed along grain boundaries (Fig. 1b), (and can also be found along intragranular slip lines). M23C6 carbides, carbo-nitrides Ti(C,N) and g matrix are also identified by analysis of the XRD traces (although the precise identification of the M23C6 is difficult because the metallic M elements substitute for each other). Peak identification has been performed R software (for further details please see using PeakFitV Krishna et al., 2010). An XRD spectrum from the creep-exposed sample is shown in Figure 3. The peak positions of TiN/TiC are labelled and whilst there is a large peak at 95 that is consistent with TiN/TiC the supporting peaks at ~41 and ~60 are less distinguishable, this could be owing to the relatively small TiN/TiC content in the alloy making it difficult to resolve the peaks above the backgound. In order to resolve this issue, TEM analysis was used to confirm the presence of these particles. The Ti(C,N) degeneration in creep-exposed IN617 alloy was corroborated using transmission electron microscopy (TEM) Figure 4. The decomposition of a large primary Ti(C,N) particle near to a grain boundary is Microscopy Research and Technique
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Fig. 2. (a) Energy dispersive X-ray analysis of Ti(C,N) precipitate found inside the grain in the as-received sample. Inset shows an elemental analysis of Ti(C,N) precipitate, marked with an arrow. (b) Energy dispersive X-ray analysis of a Cr-enriched M23C6 precipitate
observed intragranularly in the as-received sample. Inset shows an elemental analysis of the precipitate, marked with an arrow. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
shown in Figure 4(a). Adjacent to the Ti(C,N) are numerous M23C6 carbide particles that apparently precipitated during the creep exposure. The application of creep stresses, (in this case at intermediate level), is thought
to play an important role in accelerating Ti (C,N) carbide degeneration. However, in alloys IN738 and GTD111, long-term thermal exposure was sufficient to result in carbide degeneration. Both alloys showed carbide degeneration in high temperature range of 1080 to 1120 C with 26,000 hours of thermal exposure (Lvov et al., 2004). Formation of h phase instead of g0 has also been found in the carbide degeneration reaction in few nickel-based alloys (Lvov et al., 2004; Qin et al., 2007; Wang et al., 2012). This tends to occur with extended exposure hours and high alloying contents of Ti and Ta in IN738 and GTD111 alloys. It is inferred that the higher (Ti1Ta)/Al ratio suppresses the formation of g0 phase and promotes the h phase formation in the condition of extended thermal exposure. Here the Ti(C,N) carbide degeneration reaction is observed in creep-exposed IN617 alloy after short exposure of 574 h at 650 C and results in g0 and M23C6 precipitation. results in g0 and M23C6 precipitation. The size of g0 and M23C6 carbides varied in the range of 10 to 30 nm Krishna et al (2013) and 0.1 to 10 lm, respectively. Figure 4 (b) shows a typical bright field (BF) TEM micrograph of M23C6 carbide with g0 particles apparently precipitated around it. The presence of M23C6 carbide and g0 precipitates was confirmed by
Fig. 3. X-Ray diffraction spectrum (Intensity versus 2h) of creepexposed IN617 alloy. XRD peaks identified with different phases present in the alloy.
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Fig. 4. (a) Bright Field (BF) TEM micrograph shows degeneration of Ti(C,N) during creep exposure in IN617 alloy. Inset shows SAD pattern revealing fundamental reflections corresponding to TiC/TiN and M23C6 carbides. (b) BF TEM micrograph shows M23C6 carbide and fine g0 particles. The fine g0 particles are decorating the surface of the
M23C6 carbide and provide evidence for the distinctive carbide degeneration reaction TiC + g ! M23C6 + g0 . Inset shows the selected area electron diffraction pattern revealing fundamental reflections corresponding to the M23C6 carbide and g0 particles.
the selected area electron diffraction (SAD) pattern inset acquired from the precipitates. g0 particles and M23C6 carbide are predicted to occur in IN617 alloy from thermodynamic considerations. A time-temperature-transformation (TTT) diagram computed from JMatPro (Krishna et al. 2010; Krishna et al. 2013) is shown in Fig. 5 (a) (Krishna et al., 2013). TTT diagram is substantiated in Figure 5a. Figure 5b shows the time at which phases will appear in the alloy. It can be observed that g0 particles start to appear first in the alloy at around 10–100 h. Then, the M23C6 carbides will appear at around 50–2,000 h. Even without MX degradation the chemical nature of the alloy will lead to the formation of intragranular g0 and intergranular M23C6 precipitates upon ageing at
an appropriate temperature e.g. 670 C. The g0 is much more difficult to detect at low total volume fraction as it is more homogeneously distributed than the grain boundary M23C6. The other nonequilibrium phases can also be observed in the diagram such as M6C, M7C3, Eta, and P, but these are not observed in the creep sample used for this study. Also these will not be discussed in this article because they are beyond the scope of this article. In the degeneration reaction, the surrounding matrix enriches in Ti and C and the Cr content of the matrix is reduced as the Cr reacts to form M23C6:
Fig. 5. (a) Predicted TTT diagram for IN617 alloy. The thermodynamic equilibrium phases in the present exposed condition are the Ti(C,N), M23C6, and g0 and the rest of the phases are nonequilibrium (This graph was previously published in Krishna et al., 2013). (b) Thermodynamic prediction of the volume fraction of various phases
with time at a holding temperature of 650 C. Note that the creep time for the sample considered here is 574 hours. Therefore mu-phase is not expected to be observed. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
TiC 1g ! M23 C6 1g0
(1)
Microscopy Research and Technique
EFFECT OF TIC CARBIDE DECOMPOSITION
Fig. 6. (a) Hardness profile (hardness versus distance from the head) of the creep-exposed sample is shown. (b) Micrograph of creepexposed IN617 fracture surface showing intergranular cracking along equiaxed grain boundaries and wedge type cracks. The applied tensile stress is uniaxial. (c) Fractograph of creep fracture surface illus-
TiC 1 gðNi ; Cr ; Mo ; Co ; Al ; Ti Þ ! Cr 16 ðNi ; Mo ; Co Þ7 C6 1 Ni 3 ðTi ; Al Þ
(2)
Given that IN617 is primarily strengthened by solid solution hardening of the g matrix, withdrawal of Cr from the matrix to form the M23C6 could have a detrimental effect quite apart from the effect of the precipitates formed in the reaction. The volume fractions involved however are very small. The envelopes of g0 particles surrounding the Crenriched M23C6 carbides probably suggests that the nucleation and growth is interconnected. The formation of g0 depletes the Ni, Al, Ti hence increasing the effective concentration of chromium and carbon contents. The nucleation and growth of Cr-enriched M23C6 carbides in turn depletes the carbide - Cr and C and increases the effective concentration of Ni, Al, and Ti for the precipitation of g0 particles. These Microscopy Research and Technique
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trating intergranular fracture initiation at triple point grain boundaries, Crack propagation has occurred along the grain boundaries. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
processes of nucleation and growth of M23C6 carbides and g0 particles are considerably accelerated under the creep exposure condition and can be observed in IN617 only after 574 h/650 C. However, degeneration reactions can be sluggish in thermally-exposed conditions as observed in IN738 and GTD111 alloys. Moreover, primary TiC carbide degeneration has not been observed in IN617 alloy after short or long thermal and creep exposures (Mankins et al., 1974; Wu et al., 2008). Wu et al. (2008) reported M6C carbide degeneration reaction in IN617 alloy aged at 871 C under the condition of prolonged ageing for 51,850 h, facilitated through the reaction: M6C 1g ! M23C6 1 g0 . The hardness of the failed sample was measured and the creep failure mechanisms were investigated. Results for hardness along the creep sample are shown in Figure 6 (a) and a longitudinal cross section of the fracture surface and a fractograph of the creep-failed sample are shown in Figure 6 (b) and (c). In the as-received condition, a low hardness value was
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observed, but there is a significant rise in hardness in the creep-exposed condition. The gauge hardness has increased to 334 Hv in comparison with 187 Hv for the as-received sample. (Fig. 6a). The rise in hardness observed in the sample is attributed to potentially precipitation hardening mechanism and work hardening (the extensive dislocation activity is indicated by the slip lines in Fig. 6 (b)). The failure mechanism of the sample was intergranular fracture by decohesion of the grain boundaries (Figs 6b and 6c). The brittle carbides formed through the degeneration reaction at grain boundaries have provided potential paths for crack progression. CONCLUSIONS The Ti (C,N) carbide decomposition reaction in creep-exposed Inconel 617 alloy has been illustrated by TEM analysis. There is evidence of the degeneration reaction, TiC + g ! M23C6 + g0 for the sample which has been creep exposed for 574 hours at 650 C with intermediate creep stress. The hardness of the failed sample significantly increased for creep-exposed as compared to as-received sample. The rise in hardness observed in the creep-exposed sample is associated with precipitation of M23C6 carbides and fine g0 particles and also with work hardening. Moreover, formation of M23C6 carbides at and near the grain boundaries contributed to creep failure due to grain boundary de-cohesion and triple point cracking. ACKNOWLEDGMENTS We thank ALSTOM Power, UK for supplying creep– exposed samples and University of Leicester’s Advanced Microscopy Centre (AMC) for assistance in experimental work. Gordon McColvin of Alstom Power UK is thanked for his comments on the manuscript. REFERENCES Akbari-Garakani M, Mehdizadeh M. 2011. Effect of long-term service exposure on microstructure and mechanical properties of Alloy 617. Materials and Design 32:2695–2700. Davis JR. 1997. ASM International, Material Park, OH 44073-0002, USA. El-Magd E, Nicolini G, and Farag M. 1996. Effect of carbide precipitation on the creep behaviour of Alloy 800HT in the Temperature Range 700 to 900 . Metall Mater Trans A 27:747–756.
He LZ, Zheng Q, Sun XF, Hou GC, Guan HR, Hu ZQ. 2005. M23C6 precipitation behavior in a Ni-based superalloy M963. J Mater Sci 40:2959–2964. Hosier JC, Tillack DJ. 1972. Inconel alloy 617—A new high temperature alloy. Met Eng Quart 12:51–55. Kihara S, Newkirk JB, Ohtomo A, Saiga Y. 1980. Morphlogical changes of carbides during creep and their effects on the creep properties of INCONEL 617 at 1000 C. Metall Trans A 11: 1019–1031. Krishna R, Hainsworth SV, Atkinson HV, Strang A. 2010. Microstructural analysis of creep exposed IN617 alloy. Mater Sci Technol 26: 797–802. Krishna R, Hainsworth SV, Gill SPA, Strang A, Atkinson HV. 2013. Topologically close-packed l phase precipitation in creep-exposed inconel 617 alloy. Metall Mater Trans A 44:1419–1429. Lvov G, Levit VI, Kaufman MJ. 2004. Mechanism of primary MC carbide decomposition in Ni-base superalloys. Metall Mater Trans A 35:1669–1679. Mankins WL, Hosier JC, Bassford TH. 1974. Microstructure and phase stability of Inconel 617. Metall Trans 5:2579–2590. Penkalla HJ, Wosik J, Fischer W, Schubert F. 2001. Structural investigation of candidate materials for turbine disc applications beyond 700 C, Proc. 5th International Symposium on Superalloys 718, 625, 706, and Various Derivatives, Loria EA, editor. Warrendale, PA: TMS. pp. 279–290. Qin XZ, Guo JT, Yuan C, Chen CL, Ye HQ. 2007. Effect of long-term thermal exposure on the microstructure and properties of a cast Ni-base superalloy. Metall Mater Trans A 38:3014–3022. Roy AK, Hasan MH, Pal J. 2009 Creep deformation of Alloys 617 and 276 at 750–950 C. Mater Sci Eng A 520:184–188. Ray AK, Singh SR, Swaminathan J, Roy PK, Tiwari YN, Bose SC, Ghosh RN. Structure property correlation study of a service exposed first stage turbine blade in a thermal power plant, 2006. Mater Sci Eng A-Struct 419:225–232. Ross EW, Sims CT. 1987. Nickel-base alloys. New York: Wiley. Vanstone RW. 2000. Advanced (700 C) Pulverised Fuel Power Plan. Proc. 5th International Charles Parsons Turbine Conference: Parsons 2000: Advanced Materials for 21st Century Turbine and Power Plants. Strang A, Banks WM, Conroy RD, McColvin GM, Neal JC, Simpson S, editors. London, UK: IOM Communications. pp. 91–97. Viswanathan R, Henry JF, Tanzosh J, Stanko G, Shingledecker J, Vitalis B, Purgert R. 2005. U.S. program on materials technology for ultra-supercritical coal power plants. J Mater Eng Perform 14: 281–292. Wang J, Zhou L, Qin X, Sheng L, Hou J, Guo J. 2012. Primary MC decomposition and its effects on the rupture behaviors in hotcorrosion resistant Ni-based superalloy K444. Mater Sci Eng AStruct 553:14–21. Wu Q, Song H, Swindeman R, Shingledecker J, Vasudevan V. 2008. Microstructure of long-term aged IN617 Ni-base superalloy. Metal Mater Trans A 39:2569–2585. Xuebing H, Yan K, Huihua Z, Yun Z, Zhuangqi H. 1998. Influence of heat treatment on the microstructure of a unidirectional Ni-base superalloy. Mater Lett 36:210–213. Zhao S, Xie X, Smith GD, Patel SJ. 2004. Gamma prime coarsening and age hardening behaviors in a new nickel base superalloy. Mater Lett 58:1784–1787.
Microscopy Research and Technique
Effect of Titanium Carbonitride (Ti(C,N)) Decomposition on Failure Mechanisms in Inconel 617 alloy RAM KRISHNA,1,2 SARAH V. HAINSWORTH,1* SIMON P. A. GILL,1 AND HELEN V. ATKINSON1 1
Department of Engineering, University of Leicester, University Road, Leicester, LE1 7RH, UK Now at Dalton Cumbrian Facility, The University of Manchester, Westlakes Science and Technology Park, Moor Row, Whitehaven, Cumbria, CA24 3HA, UK
2
KEY WORDS
Carbide decomposition reaction; Creep; Hardness; Intergranular fracture; Inconel 617; TEM
ABSTRACT Titanium Carbonitride (Ti(C,N)) decomposition in Inconel 617 alloy creepexposed at 650 C for 574 hours is reported using analytical electron microscopy techniques. Crenriched M23C6-type carbides enveloped in fine gamma prime particles thought to be precipitated from the decomposition reaction are observed in the alloy. The morphology of the M23C6 carbides is irregular and blocky and the particle size up to 5lm, whereas the morphology of gamma prime particles is mostly spherical and up to 30 nm in size. Intergranular carbides are mostly secondary precipitates of the M23Cc type (M predominantly Cr) and these respond to solution heat treatment and precipitate on the grain boundaries as a result of ageing. The ability of intragranular MX to decompose is sensitive to the N content, high N resists decomposition. Decomposed intragranular MX provides an excess source of C which can react locally with Cr to form heat treatable intragranular fine Cr23C6 precipitates. M6C can segregate in interdendritic locations during melting which may be the reason for high content of Mo in M23C6. These precipitates are generally very small and contribute to an additional hardening effect and are the reason for the onset of voiding and cracking along the grain boundaries that ultimately lead to a reduced creep rupture life. Microsc. Res. Tech. 78:336–342, 2015. V 2015 Wiley Periodicals, Inc. C
INTRODUCTION Future generation steam turbines and their components are expected to confront ultra supercritical conditions i.e., working steam conditions up to 760 C or more and operating pressure up to 35 MPa, in order to improve the thermal efficiency and reduce the environmental emissions. To accommodate the imposed extended working conditions, the currently used ferritic and austenitic alloys will need to be substituted by alloys with better creep strength and environmental resistance. Nickel-based alloys are likely to be considered as a potential alternatives. Inconel 617 alloy (IN617) belongs to the family of nickel-based alloys and is being actively evaluated as a potential candidate alloy for elevated temperature applications owing to its excellent high-temperature strength, adequate metallurgical stability, good resistance to creep, and better resistance to corrosion and oxidation (Penkalla et al., 2001; Viswanathan et al., 2005; Vanstone, 2000). The alloying elements Cr (23 wt%), Co (12 wt%), and Mo (9 wt%) in IN617 contribute towards solid solution strengthening, while small quantities of Ti and Al contribute additional precipitation strengthening (Hosier and Tillack, 1972; Krishna et al., 2010; Mankins et al., 1974; Zhao et al., 2004). Since the inception of this alloy it has been used for high temperature structural application due to its balanced alloying elements. However, degradation in mechanical properties and creep strength during thermal and creep exposure are significant as a result of microstructural evolutions and are C V
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sensitive to operating conditions (Krishna et al., 2010; Zhao et al., 2004). The microstructural evolutions in nickel-based alloys including IN617 comprises gamma prime (g0 ) coarsening (Hosier and Tillack, 1972; Mankins et al., 1974), secondary precipitation of M23C6 carbide network (Krishna et al., 2010; Zhao et al., 2004), nucleation and growth of topologically closed packed (TCP) phase such as m-phases (Kihara et al., 1980; Krishna et al., 2013) and precipitation of discrete phases by TiC or Ti carbonitride degeneration reactions (He et al., 2005; Lvov et al., 2004; Qin et al., 2007; Xuebing et al., 1998). Most of these modifications in microstructure during service are detrimental due to the fact that their occurrence depletes potent solidsolution strengtheners from the austenite g-matrix, which results in softening of the matrix at operating temperature. All of the above modifications are also not very common in the alloy IN617 and literature reports on IN617 are relatively sparse (Davis, 1997; El-Magd et al., 1996; He et al., 2005; Kihara et al., 1980; Krishna et al., 2010, 2013; Lvov et al., 2004; Mankins et al., 1974; Qin et al., 2007; Xuebing et al., 1998; Zhao et al., 2004). However the decomposition of primary TiC and Ti carbonitrides has been reported in some nickel-based alloys. The degeneration reaction *Correspondence to: Sarah V. Hainsworth, Department of Engineering, University of Leicester, University Road, Leicester LE1 7RH, UK. E-mail: [email protected] Received 1 February 2014; accepted in revised form 8 March 2014 REVIEW EDITOR: Dr. Chuanbin Mao DOI 10.1002/jemt.22359 Published online 1 April 2015 in Wiley Online Library (wileyonlinelibrary.com).
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Elements
Ni
Cr
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TABLE 1. Chemical composition of as-received Inconel 617 alloy (wt%) Mo Co Al Ti Fe
C
Si
W
0.05
0.02
Wt%
Bal.
22.8
9.0
11.9
1.15
0.48
0.35
0.06
Elements
Mn
Cu
Zr
P
V
N
Nb
S
Wt%
0.02
0.005
0.004
0.003
0.001
0.013
0.005
0.005
has either a beneficial or a detrimental effect on creep and mechanical properties depending on the morphology and distribution of the resulting precipitates, their location, and the alloys exposure conditions (El-Magd et al., 1996; He et al., 2005; Qin et al., 2007; Lvov et al., 2004; Xuebing et al., 1998). Primary Ti(C,N) particles are a major source of carbon in most nickel-based alloys below 980 C. However, they decompose slowly during thermal treatment and apparently at a faster rate during creep exposure condition, releasing carbon for several important reactions, including what has been described as the carbide degeneration reaction (Ray et al., 2006; Ross and Sims, 1987). This leads to the formation of Crenriched M23C6 carbides and g0 following the reaction: TiC 1 gmatrix ! M23C6 1 g0 . This reaction has been reported for a range of alloys from the nickel-based family K452, IN738 and Udimet-520 (Qin et al. 2007; Wang et al. 2012). In nickel-based M963 cast superalloy, primary TiC carbide degeneration reaction is observed; but, with the formation of M6C type carbide and g0 particles, TiC 1 g ! M6C 1 g0 (He et al., 2005), in contrast to the degeneration reactions in studies of Qin et al. (2007) and Wang et al. (2012). Further, in nickel-based precipitate hardened superalloys such as IN738 and GTD111 alloys primary TiC carbide degenerates as M23C6 carbides and h phase through the following reaction: TiC 1 g/g0 ! M23C6 1 h. In a very recent study, Wang et al. (2012) reported TiC carbide degeneration in nickel-based K444 superalloy. In this mechanism TiC carbide degenerated into two different carbides, M23C6 and M6C, as described in reaction: TiC 1 g! M23C6 1 M6C 1 h. Further, it was observed that the occurrence of this degeneration reaction is was detrimental to the mechanical properties. Formation of the h phase, M6C and M23C6 precipitates in K444 alloy is attributed to long-term thermal exposure (up to 10,000 h) and higher content of alloy elements (W, Ti, Al, and Hf) in the K444 alloy composition. Thus, the carbide reactions lead to the precipitation of either M23C6, g0 , M6C or h phase at intergranular locations and play a role in determining the long-term mechanical properties, creep performance, and fracture mechanisms of the alloy (He et al., 2005; Kihara et al., 1980; Krishna et al., 2013; Lvov et al., 2004; Qin et al., 2007; Wu et al., 2008). TiC is not in itself detrimental to the mechanical properties of the alloy as discussed in Mankins et al. (1974) study. Nucleation and growth of the finer size M23C6 carbides and g0 precipitates are supposed to be more prominent in strengthening the alloy than larger Ti(C,N) carbides because of the finer size of the degenerated precipitates. However, the formation of M23C6 carbides at and near the grain boundaries introduces brittleness in the alloy and results in failure due to de-cohesion and triple point cracking at the carbide-matrix interface (Lvov et al., 2004; Wang et al., 2012). It is important to note here Microscopy Research and Technique
B 0.0002
that M23C6 and M6C precipitates are chromium (Cr) enriched carbides, where M denotes elements of Cr, Ni, Mo, Co etc. The alloys refered to here are used in the cast condition whereas IN617 is a thermo-mechanically processed alloy which influences its microstructure and properties. In the present study the aim is to develop understanding of the effect of Ti (C,N) carbide degeneration in Inconel 617 alloy in terms of creep failure mechanism and variation in hardness as a mechanical property. In this article we seek to delineate the Ti(C,N) decomposition in creep-exposed IN617 alloy at 650 C using analytical microscopy. Further, formation of g0 and M23C6 through the Ti(C,N) degeneration reaction is illustrated using transmission electron microscopy in the creep-exposed alloy. MATERIAL AND EXPERIMENTAL METHODS The complete nominal chemical composition of the Inconel 617 studied in the present investigation is presented in Table 1. A standard creep specimen was prepared from a solution annealed forged rod of IN617. The ASTM grain size number and hardness of the as-received specimen were 5.0 and 186.6 6 5 HV20, respectively. Forged rods of Inconel 617 solution annealed at 1100 C/3 h/WQ were given a heat treatment at 670 C for 10 h/air cooled (AC), which usually precipitates intergranular primary carbides (Krishna et al., 2010). Creep specimens were then machined from the heattreated forged rods. The creep rupture test was conducted at 650 C under an applied stress in air by Alstom Power (Rugby, UK) and the creep test was continued until fracture occurred. Four different creep ruptured samples were examined for hardness and microstructural modifications using Vickers hardness tests and electron microscopy. X-ray diffractometry (XRD) was performed on asreceived and creep-exposed samples to characterize the phases present in the alloy, using a Philips PW ˚) 1716 diffractometer with CuKa1 radiation (k 5 1.54A in the angular 2u range between 2 to 125 , with a step size of 0.0096 . Vickers hardness tests were conducted using a Vickers hardness testing machine, with a 20 kgf load on ground and polished longitudinal cross-sections of the creep fractured specimen. The failure mechanism of the creep-exposed sample was evaluated using scanning electron microscopy. The standard polishing technique was to grind with successively finer grades of SiC emery papers and diamond suspensions to prepare specimens for metallography. Scanning electron microscopy (SEM) was performed on specimens in the polished but unetched condition. TEM foils were prepared using the standard method as discussed in Krishna et al. (2013).
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Fig. 1. (a) Scanning electron micrograph of as-received IN617 alloy shows Cr-enriched M23C6 carbides and Ti-enriched Ti(C,N) carbonitride. (b) Scanning electron micrograph of creep-exposed IN617 alloy shows M23C6 carbides precipitated along and near to grain boundaries, and slip lines. The globular particles identified in figure 1a as M23C6 could possibly be M6C particles which are not uncommon
in Ni-Cr-Mo alloys and occur in original interdendritic regions often near Ti(CN) as in the figures, particularly 1a) however M6C is Mo enriched in IN617 and EDX showed these to be Cr-rich (see discussion). [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
RESULTS AND DISCUSSION Alloying elements (Cr, Co, Mo, Al, Ti, and Fe) in Inconel 617 alloy are responsible for high temperature properties that occur through the formation of different types of carbides, carbo-nitrides, and intermetallic phases. An additional hardening arises from the localized precipitation of carbides/carbo-nitrides during ageing treatment (Hosier and Tillack, 1972; Krishna et al., 2010; Mankins et al., 1974; Zhao et al., 2004). The morphology (shape, size, and geometry) and distribution of these precipitates play an influential role in controlling the alloy mechanical and creep properties. Embedded intermetallic g0 precipitates in FCC austenite-g matrix and secondary carbide precipitates act as barriers to the movement of dislocations and contribute to hardening (Kihara et al., 1980; Krishna et al., 2010; Zhao et al., 2004). The microstructure of as-received IN617 alloy in the pre-aged condition (10h /670 C) consists of sparsely distributed blocky shape primary carbide particles of Ti(C,N) with an average size of 5 mm, a few Crenriched M23C6 type carbides (M@Cr, Ni, Mo, Co) with an average size of 3 mm and ordered g0 particles with an average size of 12.5 nm (Krishna et al., 2013; Mankins et al., 1974; Zhao et al., 2004). Figure 1a shows the morphology and location of M23C6 and Ti (C,N) carbonitrides. The particles identified as M23C6 were found to be Cr-rich. Akbari-Garakani and Mehdizadeh (2011) and Roy et al (2009) note that two types of carbide can co-exist on the grain boundaries and also within the grains on IN617 after exposure at high temperatures with M23C6 being Cr-rich and M6C having a higher Mo content. Since these particles were found to have Cr contents greater than 50% they are labelled here as M23C6 particles. M6C particles can occur in Mo containing alloys at residual interdendritic sites however reports of M6C in these alloys report them
as those that have a higher Mo content than Cr content. The precipitates are identified with Energy Dispersive X-Ray (EDX) analysis using a Princeton Gamma Tech system attached to SEM. Typical EDX analyses for primary M23C6 and TiC carbides in asreceived sample are shown in Figure 2 (carbon and nitrogen were excluded from the quantitative analysis). Primary TiC carbides may also contain some nitrogen and can form both carbides and nitrides which have isostructural cubic structures with lattice parameters a 5 0.43–0.47 nm and a 5 0.424 nm, respectively (Davis, 1997). Thus it is more accurate to refer to carbo-nitrides, Ti(C,N). Primary carbo-nitrides Ti(C,N) of blocky morphology are observed at or near to grain boundaries in the asreceived sample (Fig. 1a). In the creep-exposed sample M23C6 (Cr-enriched) carbides are observed along grain boundaries (Fig. 1b), (and can also be found along intragranular slip lines). M23C6 carbides, carbo-nitrides Ti(C,N) and g matrix are also identified by analysis of the XRD traces (although the precise identification of the M23C6 is difficult because the metallic M elements substitute for each other). Peak identification has been performed R software (for further details please see using PeakFitV Krishna et al., 2010). An XRD spectrum from the creep-exposed sample is shown in Figure 3. The peak positions of TiN/TiC are labelled and whilst there is a large peak at 95 that is consistent with TiN/TiC the supporting peaks at ~41 and ~60 are less distinguishable, this could be owing to the relatively small TiN/TiC content in the alloy making it difficult to resolve the peaks above the backgound. In order to resolve this issue, TEM analysis was used to confirm the presence of these particles. The Ti(C,N) degeneration in creep-exposed IN617 alloy was corroborated using transmission electron microscopy (TEM) Figure 4. The decomposition of a large primary Ti(C,N) particle near to a grain boundary is Microscopy Research and Technique
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Fig. 2. (a) Energy dispersive X-ray analysis of Ti(C,N) precipitate found inside the grain in the as-received sample. Inset shows an elemental analysis of Ti(C,N) precipitate, marked with an arrow. (b) Energy dispersive X-ray analysis of a Cr-enriched M23C6 precipitate
observed intragranularly in the as-received sample. Inset shows an elemental analysis of the precipitate, marked with an arrow. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
shown in Figure 4(a). Adjacent to the Ti(C,N) are numerous M23C6 carbide particles that apparently precipitated during the creep exposure. The application of creep stresses, (in this case at intermediate level), is thought
to play an important role in accelerating Ti (C,N) carbide degeneration. However, in alloys IN738 and GTD111, long-term thermal exposure was sufficient to result in carbide degeneration. Both alloys showed carbide degeneration in high temperature range of 1080 to 1120 C with 26,000 hours of thermal exposure (Lvov et al., 2004). Formation of h phase instead of g0 has also been found in the carbide degeneration reaction in few nickel-based alloys (Lvov et al., 2004; Qin et al., 2007; Wang et al., 2012). This tends to occur with extended exposure hours and high alloying contents of Ti and Ta in IN738 and GTD111 alloys. It is inferred that the higher (Ti1Ta)/Al ratio suppresses the formation of g0 phase and promotes the h phase formation in the condition of extended thermal exposure. Here the Ti(C,N) carbide degeneration reaction is observed in creep-exposed IN617 alloy after short exposure of 574 h at 650 C and results in g0 and M23C6 precipitation. results in g0 and M23C6 precipitation. The size of g0 and M23C6 carbides varied in the range of 10 to 30 nm Krishna et al (2013) and 0.1 to 10 lm, respectively. Figure 4 (b) shows a typical bright field (BF) TEM micrograph of M23C6 carbide with g0 particles apparently precipitated around it. The presence of M23C6 carbide and g0 precipitates was confirmed by
Fig. 3. X-Ray diffraction spectrum (Intensity versus 2h) of creepexposed IN617 alloy. XRD peaks identified with different phases present in the alloy.
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Fig. 4. (a) Bright Field (BF) TEM micrograph shows degeneration of Ti(C,N) during creep exposure in IN617 alloy. Inset shows SAD pattern revealing fundamental reflections corresponding to TiC/TiN and M23C6 carbides. (b) BF TEM micrograph shows M23C6 carbide and fine g0 particles. The fine g0 particles are decorating the surface of the
M23C6 carbide and provide evidence for the distinctive carbide degeneration reaction TiC + g ! M23C6 + g0 . Inset shows the selected area electron diffraction pattern revealing fundamental reflections corresponding to the M23C6 carbide and g0 particles.
the selected area electron diffraction (SAD) pattern inset acquired from the precipitates. g0 particles and M23C6 carbide are predicted to occur in IN617 alloy from thermodynamic considerations. A time-temperature-transformation (TTT) diagram computed from JMatPro (Krishna et al. 2010; Krishna et al. 2013) is shown in Fig. 5 (a) (Krishna et al., 2013). TTT diagram is substantiated in Figure 5a. Figure 5b shows the time at which phases will appear in the alloy. It can be observed that g0 particles start to appear first in the alloy at around 10–100 h. Then, the M23C6 carbides will appear at around 50–2,000 h. Even without MX degradation the chemical nature of the alloy will lead to the formation of intragranular g0 and intergranular M23C6 precipitates upon ageing at
an appropriate temperature e.g. 670 C. The g0 is much more difficult to detect at low total volume fraction as it is more homogeneously distributed than the grain boundary M23C6. The other nonequilibrium phases can also be observed in the diagram such as M6C, M7C3, Eta, and P, but these are not observed in the creep sample used for this study. Also these will not be discussed in this article because they are beyond the scope of this article. In the degeneration reaction, the surrounding matrix enriches in Ti and C and the Cr content of the matrix is reduced as the Cr reacts to form M23C6:
Fig. 5. (a) Predicted TTT diagram for IN617 alloy. The thermodynamic equilibrium phases in the present exposed condition are the Ti(C,N), M23C6, and g0 and the rest of the phases are nonequilibrium (This graph was previously published in Krishna et al., 2013). (b) Thermodynamic prediction of the volume fraction of various phases
with time at a holding temperature of 650 C. Note that the creep time for the sample considered here is 574 hours. Therefore mu-phase is not expected to be observed. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
TiC 1g ! M23 C6 1g0
(1)
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EFFECT OF TIC CARBIDE DECOMPOSITION
Fig. 6. (a) Hardness profile (hardness versus distance from the head) of the creep-exposed sample is shown. (b) Micrograph of creepexposed IN617 fracture surface showing intergranular cracking along equiaxed grain boundaries and wedge type cracks. The applied tensile stress is uniaxial. (c) Fractograph of creep fracture surface illus-
TiC 1 gðNi ; Cr ; Mo ; Co ; Al ; Ti Þ ! Cr 16 ðNi ; Mo ; Co Þ7 C6 1 Ni 3 ðTi ; Al Þ
(2)
Given that IN617 is primarily strengthened by solid solution hardening of the g matrix, withdrawal of Cr from the matrix to form the M23C6 could have a detrimental effect quite apart from the effect of the precipitates formed in the reaction. The volume fractions involved however are very small. The envelopes of g0 particles surrounding the Crenriched M23C6 carbides probably suggests that the nucleation and growth is interconnected. The formation of g0 depletes the Ni, Al, Ti hence increasing the effective concentration of chromium and carbon contents. The nucleation and growth of Cr-enriched M23C6 carbides in turn depletes the carbide - Cr and C and increases the effective concentration of Ni, Al, and Ti for the precipitation of g0 particles. These Microscopy Research and Technique
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trating intergranular fracture initiation at triple point grain boundaries, Crack propagation has occurred along the grain boundaries. [Color figure can be viewed in the online issue, which is available at wileyonlinelibrary.com.]
processes of nucleation and growth of M23C6 carbides and g0 particles are considerably accelerated under the creep exposure condition and can be observed in IN617 only after 574 h/650 C. However, degeneration reactions can be sluggish in thermally-exposed conditions as observed in IN738 and GTD111 alloys. Moreover, primary TiC carbide degeneration has not been observed in IN617 alloy after short or long thermal and creep exposures (Mankins et al., 1974; Wu et al., 2008). Wu et al. (2008) reported M6C carbide degeneration reaction in IN617 alloy aged at 871 C under the condition of prolonged ageing for 51,850 h, facilitated through the reaction: M6C 1g ! M23C6 1 g0 . The hardness of the failed sample was measured and the creep failure mechanisms were investigated. Results for hardness along the creep sample are shown in Figure 6 (a) and a longitudinal cross section of the fracture surface and a fractograph of the creep-failed sample are shown in Figure 6 (b) and (c). In the as-received condition, a low hardness value was
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observed, but there is a significant rise in hardness in the creep-exposed condition. The gauge hardness has increased to 334 Hv in comparison with 187 Hv for the as-received sample. (Fig. 6a). The rise in hardness observed in the sample is attributed to potentially precipitation hardening mechanism and work hardening (the extensive dislocation activity is indicated by the slip lines in Fig. 6 (b)). The failure mechanism of the sample was intergranular fracture by decohesion of the grain boundaries (Figs 6b and 6c). The brittle carbides formed through the degeneration reaction at grain boundaries have provided potential paths for crack progression. CONCLUSIONS The Ti (C,N) carbide decomposition reaction in creep-exposed Inconel 617 alloy has been illustrated by TEM analysis. There is evidence of the degeneration reaction, TiC + g ! M23C6 + g0 for the sample which has been creep exposed for 574 hours at 650 C with intermediate creep stress. The hardness of the failed sample significantly increased for creep-exposed as compared to as-received sample. The rise in hardness observed in the creep-exposed sample is associated with precipitation of M23C6 carbides and fine g0 particles and also with work hardening. Moreover, formation of M23C6 carbides at and near the grain boundaries contributed to creep failure due to grain boundary de-cohesion and triple point cracking. ACKNOWLEDGMENTS We thank ALSTOM Power, UK for supplying creep– exposed samples and University of Leicester’s Advanced Microscopy Centre (AMC) for assistance in experimental work. Gordon McColvin of Alstom Power UK is thanked for his comments on the manuscript. REFERENCES Akbari-Garakani M, Mehdizadeh M. 2011. Effect of long-term service exposure on microstructure and mechanical properties of Alloy 617. Materials and Design 32:2695–2700. Davis JR. 1997. ASM International, Material Park, OH 44073-0002, USA. El-Magd E, Nicolini G, and Farag M. 1996. Effect of carbide precipitation on the creep behaviour of Alloy 800HT in the Temperature Range 700 to 900 . Metall Mater Trans A 27:747–756.
He LZ, Zheng Q, Sun XF, Hou GC, Guan HR, Hu ZQ. 2005. M23C6 precipitation behavior in a Ni-based superalloy M963. J Mater Sci 40:2959–2964. Hosier JC, Tillack DJ. 1972. Inconel alloy 617—A new high temperature alloy. Met Eng Quart 12:51–55. Kihara S, Newkirk JB, Ohtomo A, Saiga Y. 1980. Morphlogical changes of carbides during creep and their effects on the creep properties of INCONEL 617 at 1000 C. Metall Trans A 11: 1019–1031. Krishna R, Hainsworth SV, Atkinson HV, Strang A. 2010. Microstructural analysis of creep exposed IN617 alloy. Mater Sci Technol 26: 797–802. Krishna R, Hainsworth SV, Gill SPA, Strang A, Atkinson HV. 2013. Topologically close-packed l phase precipitation in creep-exposed inconel 617 alloy. Metall Mater Trans A 44:1419–1429. Lvov G, Levit VI, Kaufman MJ. 2004. Mechanism of primary MC carbide decomposition in Ni-base superalloys. Metall Mater Trans A 35:1669–1679. Mankins WL, Hosier JC, Bassford TH. 1974. Microstructure and phase stability of Inconel 617. Metall Trans 5:2579–2590. Penkalla HJ, Wosik J, Fischer W, Schubert F. 2001. Structural investigation of candidate materials for turbine disc applications beyond 700 C, Proc. 5th International Symposium on Superalloys 718, 625, 706, and Various Derivatives, Loria EA, editor. Warrendale, PA: TMS. pp. 279–290. Qin XZ, Guo JT, Yuan C, Chen CL, Ye HQ. 2007. Effect of long-term thermal exposure on the microstructure and properties of a cast Ni-base superalloy. Metall Mater Trans A 38:3014–3022. Roy AK, Hasan MH, Pal J. 2009 Creep deformation of Alloys 617 and 276 at 750–950 C. Mater Sci Eng A 520:184–188. Ray AK, Singh SR, Swaminathan J, Roy PK, Tiwari YN, Bose SC, Ghosh RN. Structure property correlation study of a service exposed first stage turbine blade in a thermal power plant, 2006. Mater Sci Eng A-Struct 419:225–232. Ross EW, Sims CT. 1987. Nickel-base alloys. New York: Wiley. Vanstone RW. 2000. Advanced (700 C) Pulverised Fuel Power Plan. Proc. 5th International Charles Parsons Turbine Conference: Parsons 2000: Advanced Materials for 21st Century Turbine and Power Plants. Strang A, Banks WM, Conroy RD, McColvin GM, Neal JC, Simpson S, editors. London, UK: IOM Communications. pp. 91–97. Viswanathan R, Henry JF, Tanzosh J, Stanko G, Shingledecker J, Vitalis B, Purgert R. 2005. U.S. program on materials technology for ultra-supercritical coal power plants. J Mater Eng Perform 14: 281–292. Wang J, Zhou L, Qin X, Sheng L, Hou J, Guo J. 2012. Primary MC decomposition and its effects on the rupture behaviors in hotcorrosion resistant Ni-based superalloy K444. Mater Sci Eng AStruct 553:14–21. Wu Q, Song H, Swindeman R, Shingledecker J, Vasudevan V. 2008. Microstructure of long-term aged IN617 Ni-base superalloy. Metal Mater Trans A 39:2569–2585. Xuebing H, Yan K, Huihua Z, Yun Z, Zhuangqi H. 1998. Influence of heat treatment on the microstructure of a unidirectional Ni-base superalloy. Mater Lett 36:210–213. Zhao S, Xie X, Smith GD, Patel SJ. 2004. Gamma prime coarsening and age hardening behaviors in a new nickel base superalloy. Mater Lett 58:1784–1787.
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