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Microstructure evolution and mechanical properties of Inconel 740H during aging at 750 °C
Materials Science & Engineering A 589 (2014) 153–164
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Microstructure evolution and mechanical properties of Inconel 740H during aging at 750 1C Chong Yan a,b,n, Liu Zhengdong b, Andy Godfrey a, Liu Wei a, Weng Yuqing b a b
School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China Central Iron&Steel Research Institute, Beijing 100084, China
art ic l e i nf o
a b s t r a c t
Article history: Received 6 May 2013 Received in revised form 16 September 2013 Accepted 18 September 2013 Available online 1 October 2013
The microstructure evolution of Inconel 740H during aging at 750 1C for up to 3000 h was investigated by means of optical microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), small-angle X-ray scattering (SAXS), three-dimensional atom probe (3DAP) analysis and micro-phase analysis. The mechanical properties of samples after aging were also studied. The grain size increased substantially during the whole aging period. Two different types of grain boundary carbides were observed; block-shaped and needle-shaped. Both were identified to be M23C6 by selected area diffraction measurements. The grain boundary carbides did not coarsen significantly during aging, with the weight fraction increasing only from 0.20% to 0.28%. In contrast, a much higher coarsening rate of γ′ precipitates was observed, as evidenced from both TEM observations and SAXS analysis. 3DAP was used to study the elemental partitioning behavior between γ′ precipitates and γ matrix as well as the evolution in width of the γ/γ′ interface. A large increase in the width of γ/γ′ interfaces was seen between 1000 h and 3000 h aging. In addition, for the sample aged at 750 1C for 3000 h, Cr enrichment on the γ matrix side of the γ/γ′ interface was found. Tensile tests at 750 1C of the aged samples showed a gradual decrease in elevated-temperature yield strength after 500 h, when this alloy was over-aged. The critical precipitate size for the transition from precipitate cutting by weakly coupled dislocations to strongly coupled dislocations for Inconel 740H was calculated to be approximately 50 nm, which agrees well with the experiment measurements of elevated-temperature yield strength. The room temperature impact toughness of all samples decreased during aging as the grain size kept growing. & 2013 Elsevier B.V. All rights reserved.
Keywords: Inconel 740H Microstructure evolution γ′ precipitates Mechanical properties Three-dimensional atom probe
1. Introduction As a result of increasing energy demands and accelerated environmental problems, there is an urgent need to improve the thermal efficiency of coal-fired power plants. To achieve this goal, next generation ultra-supercritical power plants, operating at steam temperatures of 700–750 1C and operating pressure up to 37.5 MPa, are under development [1–3]. However, the increased operating parameters place more stringent requirements on the properties of candidate materials and cannot be met by conventional ferritic and austenitic steels. Consequently, there are efforts to replace these materials by Ni-base superalloys, which show a longer creep rupture life and higher corrosion resistance. Ni-based superalloys have already found widespread applications in a number of critical technological areas, especially those involving high temperatures, such as jet-engine turbines and coalfired power plants. These alloys typically exhibit an excellent
n
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balance of properties, including good mechanical properties and ductility (both at room temperature and elevated temperatures), improved fracture toughness and fatigue resistance, as well as enhanced creep and oxidation resistance at high temperatures [4]. One of the most promising candidate Ni-base superalloys for the main steam pipe of 700 1C ultra-supercritical coal-fired power plants is Inconel 740H, which is a modified version of Inconel 740 developed by Special Metals Corp, Huntington, West Virginia, USA. The nominal composition of Inconel 740H is 25Cr, 20Co, 0.5Mo, 1.5Nb, 1.4Ti, 1.4Al, 0.3Mn, 0.03Fe, 0.03C and Ni balance (wt%). Compared with Inconel 740, the ratio of Ti to Al in Inconel 740H is lowered in order to eliminate microstructure instabilities found in Inconel 740 during prolonged thermal aging at 750 1C [5–11]. In addition, the Nb content is also reduced to improve the weldability of the alloy. Since Inconel 740H has only recently been developed, few studies concerning the microstructure evolution and mechanical properties have been carried out so far on this material. Moreover, no reports exist on the detailed microstructure evolution during aging of Inconel 740H at the atomic scale. Microstructural stability is of prime importance for alloys to be used at high temperatures for long periods of time. Inconel 740H
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consists of a γ matrix with a coarse grain size, containing a dispersion of γ′ precipitates inside the grains, as well as carbides located at the grain boundaries. The microstructural evolution during long time aging, including grain size, γ′ precipitates and grain boundary carbides, should be studied at different scales in detail. In addition to traditional optical microscopy and electron microscopy observation, three-dimensional atom probe (3DAP) has been used in this study for the first time. In the last decade, there have been large numbers of studies employing 3DAP to characterize the microstructure of Ni-base superalloys at the atomic scale, including the size, morphology and composition of γ′ precipitates within the γ matrix as a function of varied heat treatments [4,12–15]. In particular, the capability of 3DAP to provide information on elemental partitioning between phases, and on segregation to certain material defects, lend it a
competitive edge in research related to interfaces and grain boundaries in Ni-base superalloys. The change of mechanical properties in Ni-base superalloys with aging time is also an important issue that needs to be investigated. The room temperature impact toughness and microhardness of Inconel 740H have already been studied after aging at 750 1C for 1000 h at different temperatures [6]. Also, the elevated-temperature microhardness of the alloy in the standard heat treatment condition has been studied, and the results compared with room temperature microhardness [6]. However, studies of the evolution of elevatedtemperature strength and of room temperature toughness of Inconel 740H with aging time, as well as their relationship with microstructure have not been reported yet. In this work, tensile tests at 750 1C and impact toughness measurements at room temperature have been carried out for samples aged up to 3000 h at 750 1C. The results are analyzed with respect to the observed changes in the size of the γ′precipitates and the grain size.
2. Experimental procedure
Fig. 1. The geometries of the mechanical test samples (a) elevated-temperature tensile test and (b) room temperature impact toughness test.
The bulk chemical composition of the Inconel 740H sample was analyzed to be 49.27Ni–25.57Cr–19.92Co–1.48Nb–0.50Mo–1.38Ti– 1.47Al–0.34Mn–0.03Fe–0.02C wt%. Samples cut from the forged bars were solution treated at 1150 1C for 30 min and then water quenched. These samples were subsequently aged at 750 1C for 100, 300, 500, 1000 and 3000 h and then air quenched. For convenience, these samples will be subsequently referred to as WQ100, WQ300, WQ500, WQ1000 and WQ3000. After aging, the specimens were ground and polished following standard metallographical methods and then chemically etched in a fresh solution of HCl, HNO3 and H2O volume proportions of 10:1:10 at 50 1C for 15 min. Optical microscope (OM) investigations were conducted using an Olympus GX51 microscope. Scanning electron microscope (SEM) investigations were conducted using a Hitachi 4500 scanning electron microscope. Samples for transmission electron microscope (TEM) investigations were prepared by a standard electro-polishing technique using a double-jet device.
Fig. 2. Optical micrographs of the samples aged at 750 1C for different times.
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Slices of 0.3 mm thickness were cut out of the bulk material and mechanically ground to 40–50 μm in thickness. Discs of 3 mm in diameter were then punched from the thinned slice and then electro-polished at 30 1C at a voltage of 16 V, using a solution of 85% ethanol and 15% perchloric acid. TEM investigations were conducted using a JEOL 2010 TEM operated at 200 kV. The weight fractions of the γ′ precipitates and carbides were ascertained through micro-chemical phase analysis methods. The γ′ precipitates were electrolytically extracted in an aqueous solution containing 1% (NH4)2SO4 and 1% C6H8O7.H2O from 0 to 5 1C. The carbides were electrolytically extracted in a methanol solution containing 5% HCl, 5% C3H8O3 and 1% C6H8O7 H2O from 10 to 5 1C. In both cases a current density of 0.05 A/cm2 was used. The particle size distribution and mean radius of the extracted γ′ precipitates were determined using an X-ray diffractro-spectrometer equipped with a Krathy small angle scattering (SAXS) goniometer. Samples for 3DAP tomography studies in the LEAP microscopy were prepared by an electro-polishing method. For this purpose,
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samples of different ageing conditions were first electro-discharge machined into thin wires with a square cross section of 0.7 0.7 mm2. These wires were mechanically ground and subsequently electro-polished in two steps; first with a 75% acetic acidþ 25% perchloric acid solution for the coarse polish, and finally with a 98% butyl cellulose þ2% perchloric acid solution at 12–14 V for the final polish. The 3DAP tomography experiments, were performed using an Imago (now Camera Instruments) local electrode atom probe (LEAP 3000 h) at a residual pressure of 5 10–9 Pa and a specimen temperature of 50 K, and with a pulse repetition frequency of 200 kHz and a pulse-voltage to dc-voltage ratio of 15%. Data analysis was performed using the IVAS 3.6.0 software. Impact testing of Charpy V-notch samples cut from the aged samples, was performed at room temperature. Tensile tests at 750 1C were carried out on aged samples in a hydraulic test system equipped with a temperature-controlled furnace. The temperature fluctuations were kept within 3 1C. The load rate before yielding was 0.5 mm/min and increased to 2.5 mm/min after yielding.
Fig. 3. SEM micrographs and the typical EDS of the grain boundary carbides aged at 750 1C for different times.
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The geometries of the Charpy V-notch samples and tensile test samples are provided in Fig. 1.
3. Results and discussion 3.1. Microstructure evolution 3.1.1. Grain size The microstructures of the WQ100, 300, 500, 1000, 3000 samples as seen in the optical microscope are shown in Fig. 2. The microstructure consists of a coarse grain structure containing annealing twins. It is evident that grain growth took place during aging. The average grain sizes, determined by an intercept method, for each aging time were 108, 112, 120, 137, 151 μm respectively.
3.1.2. Grain boundary carbides The evolution of the grain boundary carbides with aging time is illustrated in Fig. 3. The carbides were identified to be Cr rich M23C6 according to energy dispersive spectrometer (EDS) measurements of the carbides in Fig. 3, which is consistent with previous studies of Inconel 740 by Xie et [5]. There was only marginal coarsening of the carbides at prolonged aging time, with the weight fraction increasing just from 0.20% to 0.28% (Fig. 4(a)). TEM images and selected area diffraction (SAD) patterns of the grain boundary carbides are shown in Fig. 5. In general, the grain boundary carbides exhibited two different shapes; block-shaped most common, and needle-shaped rarely. The needle-shaped carbides appear to be nucleated from the grain boundaries and to grow perpendicularly into the grains, similar to the morphology of η phase reported by Zhao on the microstructure evolution of
Fig. 4. Evolution of phase fraction (wt%) with aging time (a) carbides, (b) γ′ phase.
Fig. 5. TEM micrographs and corresponding SAD patterns of different shape grain boundary M23C6. (a) and (c) block-shaped M23C6 and corresponding SAD pattern, (b) and (d) needle-shaped M23C6 and corresponding SAD pattern.
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Inconel 740 aged at 750 1C [6,7]. However, SAD patterns of the needle-shaped phase confirm that it is Cr rich M23C6 carbides rather than η phase. 3.1.3. γ′ Phase 3.1.3.1. Temporal evolution of γ′ morphology. As the main contributor to high temperature strength in Ni-base superalloys, the morphology, size and chemical composition of γ′ precipitates are of primary importance in an investigation of Inconel 740H. TEM images, together with diffraction pattern showing typical superlattice reflections from γ′ precipitates for the WQ100, 300, 1000, 3000 samples are given in Fig. 6. No γ″ precipitates were detected from either the TEM images or the SAD patterns. The γ′ precipitates exhibited a spheroidal morphology despite of becoming slightly cuboidal after long time aging. To determine precisely the variations in size distribution and mean radius of γ′ precipitates with aging time, SAXS experiments were also carried out. The results are shown in Fig. 7. In general the γ′ precipitates exhibit a one-modal size distribution, with evident coarsening during aging, from a mean precipitate radius increasing of 22.4– 69.2 nm. The γ′ precipitate weight fraction increased from 14.0% to 16.5%, as shown in Fig. 4(b). In the first 100 h, all γ′ precipitates remained below 30 nm in radius, whereas after that γ′ precipitates of size 30–60 nm were detected. In the samples aged longer than 1000 h, even larger γ′ precipitates of size 60–90 nm were detected. It should be noted however that, these larger precipitate sizes were not included in the calculation of the mean precipitate size, as it was assumed these were in fact clusters of several small γ′ precipitates together. At the same time, the number density of precipitates decreased with aging time (Fig. 6). Although the
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morphology of γ′ precipitates can be clearly seen in TEM foils, the elemental distribution between γ′ precipitates and γ matrix, as well as the compositional profile across the γ/γ′ interfaces cannot be precisely obtained due to the small size of the γ′ precipitates. For this purpose, 3DAP studies were carried out to study the microstructure at the atomic scale. 3.1.3.2. Chemical analysis of this alloy. Examples of 3DAP reconstructions from the WQ100 and 3000 samples are shown in Fig. 8 (a) and (b). The size of each reconstruction volume is shown in the figure. The Co rich regions (blue) corresponded to the γ matrix, whereas the red Cr ¼25 at% isosurfaces are used to delineate the outlines of the γ′ precipitates. According to previous work [14], the concentration value used to delineate the Cr isosurfaces was selected to be the average value of the concentrations in the γ matrix and γ′ precipitates (in this case E25%). Owing to the limited size of the reconstruction volume, only part of each γ′ precipitate can be seen in the reconstructed volume, although a few complete small precipitates are seen in the WQ100 sample. Nevertheless, the 3-D morphology of the γ′ precipitates is seen to be near spherical, in agreement with the TEM investigations. We take here the WQ3000 sample as an example to investigate the partitioning behavior of different alloying elements in Inconel 740H. The results are illustrated in Fig. 9. By comparing Figs. 8 and 9, it is seen that different elements have a different partitioning behavior between the γ matrix and the γ′ precipitates. Cr, Co and Mo tend to distribute in the γ matrix, while Al, Ti and Nb preferentially segregate to the γ′ precipitates. This segregation tendency is consistent with other studies of Ni-base superalloys [13,16–18]. It is worth noting that the level of segregation for
Fig. 6. TEM images of γ′ precipitates in samples aged for different times at 750 1C.
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Fig. 7. Size distribution of γ′ measured by SAXS in Inconel 740H at different aging times.
each element was not the same. For example, segregation of Mo in γ matrix was rather low compared with that of Cr and Co. Likewise, the segregation levels for different γ′ segregating elements were also different. A lever rule diagram, first implemented by Blavette et al. [19], was constructed to determine in detail the partitioning behavior as well as partitioning intensity of the solutes (Fig. 10). In the diagram, the nominal concentration for each alloying element minus the solute content in the γ matrix is plotted versus the difference in composition between γ′ precipitates and γ matrix. The slope of the best-fit line passing through the origin for all elements gives the volume fraction of γ′ precipitates. The elements segregating to the γ′ precipitates are seen in the first quadrant of the coordinate system. Points to the right of the diagram correspond to stronger segregation in γ′ precipitates (vice versa for the γ matrix segregating elements). In agreement with the results shown in Fig. 9, the lever rule diagram shows that Al, Ti and Nb tend to segregate in γ′ precipitates, with segregation intensity in descending order. Similarly in the γ matrix the Cr, Co and Mo segregate in descending order. In order to quantify the partitioning behavior of different elements during aging, we define the solute partitioning ratio as γ′ γ γ′ γ K γ ′=γ ¼ C i =C i , where C i and C i are the solute concentrations of element i in the γ′ precipitates and in the γ matrix, respectively.
Values of K γ ′=γ above 1 indicates γ′ segregating elements, and for γ segregating elements, the value is below 1. Moreover an increasing value of K γ ′=γ corresponds to an increasing tendency for segregation to the γ′ phase. The values of K γ ′=γ for different elements during aging are summarized in Fig. 11. For the γ segregating elements, an interesting phenomenon was observed, namely that the values of K γ ′=γ for Cr, Co and Mo all reached a peak value at 300 h, which means that, the γ′ precipitates contained highest contents of Cr, Co, Mo after aging at 750 1C for 300 h. The reason for this phenomenon might be as follows. Elements such as Cr, Co and Mo that segregate to γ matrix diffuse from γ′ precipitates into the γ matrix during growth of γ′ precipitates. As observed in Fig. 7, a sudden growth in size of γ′ precipitates was seen between 100 h and 300 h. This growth may leave insufficient time for diffusion of Cr, Co and Mo into the γ matrix. After 300 h aging, the growth rate of the γ′ precipitates slowed down, leaving sufficient time for diffusion. It is suggested therefore that, the combined effects of size growth and repartitioning by diffusion lead to an enrichment of Cr, Co and Mo inside the γ′ precipitates at 300 h. The trends of K γ ′=γ for Ti and Al with aging time were opposite, indicating that with prolonged aging time, Ti in the γ′ phase was gradually replaced by Al. The trend of K γ ′=γ with aging time, however, had nothing in common with other elements.
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Fig. 8. 3-D reconstruction showing Co atoms in blue and isosurfaces for Cr ¼ 25 at% in red for more clear visualization of the γ′ precipitates for (a) WQ100 and (b) WQ3000. (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
3.1.3.3. Temporal evolution of the γ/γ′ interface. The compositional profiles across the γ/γ′ interfaces for the WQ100, 300, 1000, 3000 samples were determined using a powerful analysis tool, the proxigram, for the various γ′ precipitates shown in the reconstruction volumes in Fig. 8. Widely employed in other studies using 3DAP, proxigrams have been shown to be useful for the study of any internal surfaces, even those involving a complex topology [14,15]. In this study proxigrams for the γ/γ′interfaces were constructed based on isosurfaces with a threshold value of 25 at% Cr. The results are shown in Fig. 12, where only the main alloying elements (Cr, Co, Mo, Ti, Al, Nb) are shown for simplicity. The proxigrams show that the compositional profile across the γ/γ′ interface is not abrupt but
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transitionary. The width of the transitionary region is defined as the interface width. According to Yoon [19], the choice of Cr threshold value used to fix the isosurfaces only influences the location of the interface, but not the width of the interface. The width of the interface in each proxigram was determined from the Al composition profile, between the steady-state γ′ and γ. The width of the γ/γ′ interfaces was determined using a 10–90% of the plateau method (long-range γ and γ′ compositions) [14,15]. The results are shown in Fig. 12. The interface width in the WQ100 sample was about 2.0 nm, similar to the results of Rene88DT [14]. The interface width was relatively stable over the first 1000 h with a large increase between 1000 h and 3000 h, which is believed to be associated with the growth of the precipitates and the misfit strain around the interface. It is worth noting that in the WQ3000 sample, an enrichment of Cr was found on the γ matrix side of the γ/γ′ interface, (Fig. 13). This phenomenon was not found in the other aged samples. The width of the enriched region was about 2.5 nm. Yoon observed a similar phenomenon in the commercial Ni-base superalloy RENE N6, but the reason for this enrichment was not discussed [19]. In addition, the enrichment of Co at the interface was also detected in RENE N6. In contrast this was not found in the WQ3000 sample investigated in the present study. Two possible explanations have been proposed for the enrichment of Cr at the γ/γ′ interfaces. The first is based on a diffusion mechanism. As shown in Fig. 10, Cr segregates strongly to the γ matrix. Therefore, with the increase of γ′ precipitate size, excess Cr inside the γ′ precipitates must be excluded and must diffuse across the γ/γ′ interface into the γ matrix. A low diffusion coefficient of Cr across the interface might therefore lead to the enrichment of Cr at the interface. However, this explanation cannot account for the absence of Cr enrichment at earlier aging times. The other explanation is associated with the suggestion that a new phase Cr-rich may form at the γ/γ′ interface. It has been reported that a Ni2Cr-type long-range ordered phase can be expected in Ni–Al alloys with high Cr and Mo content. In addition to the binary Ni–Cr system such a phase has also been found in ternary Ni–Cr–Mo [20,21] and Ni–Cr–Al alloys [22]. The Ni2Cr superstructure (Pt2Mo type), with orthorhombic unit cell, forms in the Ni–Cr system below 590 1C. It has been shown that addition of Mo to binary Ni2Cr significantly increases the critical temperature for the formation of this phase, e.g. in Ni–25Mo–8Cr, the critical temperature of Ni2(Cr,Mo) phase is 750 1C [23,24]. Additionally Svoboda and his co-workers found the existence of Ni2Cr at the γ/γ′ interface in a Ni–Cr–Al–Mo system aged at 600 1C for 3000 h [25]. Sundararaman in Alloy 625 also found similar results [26]. Intermetallic phase Ni2(Cr,Mo) precipitates with Pt2Mo-type structure have also been observed, in addition to that of the γ″ phase, in Alloy 625 after prolonged ( 70,000 h) service at temperatures close to but less than 600 1C. Identification of the nature of the Cr enrichment at the γ/γ′ interface will therefore be the subject of further research. 3.2. Mechanical properties 3.2.1. Tensile strength at 750 1C The results of tensile tests at 750 1C of Inconel 740H aged at 750 1C for 100, 300, 500, 1000 and 3000 h are shown in Fig. 14. The yield strength of Inconel 740H at 750 1C increased in the first 500 h and then decreased gradually with further increase in aging time. The elongation increased during the whole aging. The changes in yield strength can be rationalized on the basis of the growth of γ′ precipitates based on two previously reported models [27,28]. In both cases, it is believed that the critical resolved shear stress (CRSS) is determined by the force necessary to move two coupled edge dislocations in the 〈110〉 direction on the {111} plane through the γ′ precipitates. The first model [27] describes the CRSS for cutting through small γ′ precipitates in a disordered matrix by
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Fig. 9. Elemental partition between γ and γ′ phase of sample aged at 750 1C for 3000 h.
two weakly coupled dislocations while the second [28] describes the same motion for cutting of larger γ′ precipitates by strongly coupled dislocations. In the first model, the CRSS for two weakly coupled dislocations is given by 1 3 1 Γ 2 bdf 2 1 Γ f ð1Þ Δτ0 ¼ A 2 T 2 b b where Γ is the anti-phase boundary energy of the γ′ precipitates in the {111} plane, b is the Burgers vector of the edge dislocation in the γ matrix, d is the γ′ precipitates diameter, f is the volume fraction of the γ′ precipitates, T is the line tension of the dislocation and A is a numerical factor depending on the morphology of the particles. For spherical particles A¼0.72 [29]. The line tension is given by 2
T¼
Gb 2
ð2Þ
where G is the elastic shear modulus and is taken as 60 GPa in this study.
For cutting of larger precipitates by strongly coupled dislocation pairs, the CRSS is given by 1
Tf 2 ω Δτ0 ¼ 0:86 bd
!
1:28dΓ ωT
12
ð3Þ
where w is a constant accounting for the elastic repulsion between the strongly paired dislocations, and which is of the order of unity. For quantitative calculation, f is taken as the average phase fraction of γ′ precipitates (15%), w is taken as 1 for simplicity, and Γ is taken to be 0.42 J/m2. Fig. 15 is a plot of γ′ precipitate diameter d versus CRSS Δτ0 for the two dislocation mechanisms. For precipitate cutting by weakly coupled dislocations, Δτ0 monotonically increases with increasing precipitate diameter. In contrast Δτ0 gradually decreases with increasing precipitate diameter for the mechanism of precipitates cutting by strongly coupled dislocations. The resultant Δτ0 for a given d is the lower value of the two, since dislocation activity follows the mechanism that provides the least resistance to glide.
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Fig. 10. Lever rule diagram of Inconel 740H(WQ1000). The nominal concentration (Cn) for each element (Cr, Al, Ti, etc.) minus the solute content in the γ matrix (Cγ) is plotted as a function of the difference in composition between γ and γ′ (Cγ Cγ′). The slope gives the molar fraction of γ′ present in Inconel 740H.
Fig. 11. Partitioning ratios, K γ′=γ , for main elements in Inconel 740H plotted as a function of aging time.
For any given alloy, there exists a critical value of d that gives the highest flow stress. For Inconel 740H, the optimum value of d was calculated to be around 50 nm (shown in Fig. 15), which coincides well with the experimental results of yield strength shown in Fig. 14. In the first 500 h aging, when the diameter of γ′ precipitates is smaller than 50 nm, the precipitates cutting by weakcoupled pairs of dislocations is the dominant mechanism and the yield strength increases with the increasing precipitate diameter. However, after 500 h as the diameter of γ′ precipitates exceeds the critical value of d, the mechanism changes to cutting by strongcoupled pairs of dislocations. As a result, the yield strength of Inconel 740H decreases monotonically with longer aging time.
time is evident. For the sample aged at 750 1C for 3000 h, the value decreased to less than 20 J, which would present a huge threat to the safety of use in operation. As reported by Zhao [6], the fractographs of Inconel 740 samples heat-treated at 750 1C indicated a clearly brittle fracture with a localized mixed-mode behavior and a clear intergranular fracture behavior. Therefore, the impact toughness of the aged samples will be influence by the change in grain size and by the size of the grain boundary carbides. As mentioned in Sections 3.1.1 and 3.1.2, during aging the size of grain boundary carbides remained unchanged, whereas the grain size increased. Therefore, it can be concluded that the drastic loss of impact toughness at room temperature for aged samples is mainly due to the growth in grain size.
3.2.2. Impact toughness at room temperature The results of the impact toughness measurements at room temperature of Inconel 740H aged at 750 1C for 100, 300, 500, 1000 and 3000 h are shown in Fig. 16. A loss of impact toughness with aging
4. Conclusions The microstructure evolution and mechanical properties of Inconel 740H during aging at 750 1C up to 3000 h have
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Fig. 12. Proxigrams showing the compositions profiles of Cr, Co, Mo, Ti, Al and Nb atoms across the γ/γ′ interfaces for (a) WQ100, (b) WQ300, (c) WQ1000 and (d) WQ3000 conditions.
900 34 850 32
750 28 700 26 650 24 600
Strength/MPa
Elongation/%
800 30
22 550 20
A
500
18 450 100
Aging time/h1000
Fig. 14. The change of mechanical properities of Inconel 740H with aging time.
Fig. 13. Proxigram across the γ/γ′ interface of WQ3000, note that γ is on the right and γ′ is on the left.
been investigated with a combination of OM, SEM, TEM, 3DAP, SAXS, micro-chemical analysis and elevated-temperature tensile tests, as well as room temperature impact toughness
measurements. The main conclusions of this study are as follows: (1) The average initial grain size of Inconel 740H is above 100 μm and evident grain growth takes place during aging. The grain boundaries are fully decorated with Cr-rich M23C6 carbides, which does not exhibit significant coarsening during aging up
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Fig. 15. Theoretical critical resolved shear stress versus γ′ precipitates size relationships in Inconel 740H deformed at 750 1C.
around the γ/γ′ interface. Enrichment of Cr at the γ matrix side of γ/γ′ interface is found for the sample aged for 3000 h. (4) There exists a critical value for the diameter of γ′ precipitates that gives optimum yield strength of Inconel 740H. When the diameter of γ′ precipitates is below the critical value, the yield strength increases monotonically with increasing precipitate size, while when the size is above the critical value, the yield strength decreases with the increase of precipitate size. The critical value for Inconel 740H was calculated to be around 50 nm, which agrees well with the experimental results. A loss of room temperature impact toughness of Inconel 740H with aging time was evident. The main reason for this was attributed to the increase in grain size during aging.
Acknowledgments
Fig. 16. The change of impact toughness at room temperature of Inconel 740H with aging time.
to 3000 h, with the carbide weight fraction increasing from 0.20% to 0.28%. Two different forms of grain boundary carbides, namely block-shaped and needle-shaped, are observed. Both are identified to be M23C6 by selected area diffraction measurements. (2) TEM and 3DAP image data, together with the results of SAXS experiments, indicated that size distribution of the spheroidal γ′ precipitates inside each grain is one-modal and that the mean precipitate radius increases rapidly during aging at 750 1C. The elemental partitioning behaviors have been studied with the help of 3DAP measurements and the results analyzed using lever rule diagrams and the solute partitioning ratio K γ ′=γ . It was found that, Cr, Co, Mo tend to distribute in the γ matrix, while Ti, Al, Nb preferentially segregate to the γ′ precipitates. The partitioning intensities for different elements are not the same and vary with aging time. (3) The compositional profiles across the γ/γ′ interfaces were investigated using proxigrams and the compositional widths of the interface determined for different aging times. An increase in the γ/γ′ interface width was found between 1000 h and 3000 h aging, which was believed to be associated with the growth of the γ′ precipitates and the misfit strain
The authors are grateful to the financial support from the HighTech Research and Development Project (No. 2102AA03A501) supported by the National Ministry of Science and Technology of China.
Appendix A. Supplementary materials Supplementary data associated with this article can be found in the online version at http://dx.doi.org/10.1016/j.msea.2013.09.076. References [1] R. Viswanthan, J.F. Henry, J. Tanzosh, G. Stanko, J. Mater. Eng. Perform. 14 (2005) 281–292. [2] R. Viswanthan, K. Coleman, U. Rao, Int. J. Pressure Ves. Pip. 83 (2006) 778–783. [3] R. Viswanthan, R. Purget, P. Rawl, Adv. Mater. Process. (2008) 41–45. [4] T.Y. Hwang, R. Banerjee, J. Tiley, R. Srinivasan, G.B. Viswanthan, H.L. Fraser, Metall. Mater. Trans A40 (2009) 24–35. [5] S. Zhao, X. Xie, G.D. Smith, S.J. Patel, Mater. Lett. 58 (2004) 1784–1787. [6] S. Zhao, X. Xie, G.D. Smith, S.J. Patel, Mater. Sci. Eng. 355A (2002) 96–105. [7] S. Zhao, X. Xie, G.D. Smith, S.J. Patel, Mater. Des. 27 (2006) 1120–1127. [8] N.D. Evans, P.J. Maziasz, R.W. Swindeman, G.D. Smith, Scr. Mater. 51 (2004) 503–507. [9] J.P. Shingledecker, G.M. Pharr, Metall. Mater. Trans. A43 (2012) 1902–1911. [10] J. Cowen, P.E. Christopher, P.D. Danielson, Jablonski, J. Mater. Eng. Perform. 20 (2011) 1078–1083. [11] J.-H. Oh, B.-G. Yoo, I.-C. Choi, M.L. Santella, J.-I. Jang., J. Mater. Res. 26 (2011) 1253–1259. [12] J. Tiley, G.B. Viswanthan, R. Srinivasan, R. Banerjee, D.M. Dimiduk, H.L. Fraser, Acta Mater. 57 (2009) 2538–2549.
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[13] J.Y. Hwang, S. Nag, A.R.P. Singh, R. Srinivasan, J. Tiley, G.B. Viswanthan, H.L. Fraser, R. Banerjee, Metall. Mater. Trans. A40 (2009) 3059–3068. [14] J.Y. Hwang, S. Nag, A.R.P. Singh, R. Srinivasan, J. Tiley, G.B. Viswanthan, H.L. Fraser, R. Banerjee, Scr. Mater. 61 (2009) 92–95. [15] R. Srinivasan, R. Banerjee, J.Y. Hwang, G.B. Viswanthan, J. Tiley, D.M. Dimiduk, H.L. Fraser, Phys. Rev. Lett. 102 (2009) 086101. [16] E. Cadel, D. Lemarchand, S. Chambreland, D. Blavette, Acta Mater. 50 (2002) 957–966. [17] Y. Zhang, N. Wanderka, G. Shumacher, R. Schneider, W. Niumann, Acta Mater. 48 (2000) 2787–2793. [18] P.J. Warren, A. Cerezo, G.D.W. Smith, Mater. Sci. Eng. 250 A (1998) 88–92. [19] D. Blavette, A. Buchon, S. Chambreland, in: H.E. Exner, V. Schumacher (Eds.), Proceedings of the Euromat Conference, Aachen, 1989, pp. 419–424.
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K.E. Yoon, Ph.D. Thesis, Northwestern University, Evanston, Illinois, 2004. J. Karmazin, J. Krejci, J. Zeman, Mater. Sci. Eng. A183 (1994) 103–109. J. Zeman, Ph.D. Thesis, Institute of Physical Metallurgy, CAS, Brno, 1989. S.K. Srivastava, B.E. Lewis, in: H. Chen, V.K. Vasudevan (Eds.), The Minerals, Metal and Material Society, TMS, Warrendale, PA, 1992, pp. 141–149. M. Kumar, V.K. Vasudevan, in: H. Chen, V.K. Vasudevan (Eds.), The Minerals, Metal and Material Society, TMS, Warrendale, PA, 1992, pp. 131–139. J. Bursik, M. Svoboda, Scr. Mater. 39 (1998) 1107–1112. M. Sundararaman, L. Kumar, G.E. Prasad, P. Mukhopadhyay, S. Banerjee, Metall. Mater. Trans. A30 (1999) 41–52. L.M. Brown, R.K. Ham, Applied Science, Elsevier Applied Science, London, 1971. W. Huther, B. Reppich, Z. Metall. 69 (1978) 628–634. B. Reppich, Acta. Metall. 30 (1982) 87–94.
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Microstructure evolution and mechanical properties of Inconel 740H during aging at 750 1C Chong Yan a,b,n, Liu Zhengdong b, Andy Godfrey a, Liu Wei a, Weng Yuqing b a b
School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China Central Iron&Steel Research Institute, Beijing 100084, China
art ic l e i nf o
a b s t r a c t
Article history: Received 6 May 2013 Received in revised form 16 September 2013 Accepted 18 September 2013 Available online 1 October 2013
The microstructure evolution of Inconel 740H during aging at 750 1C for up to 3000 h was investigated by means of optical microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), small-angle X-ray scattering (SAXS), three-dimensional atom probe (3DAP) analysis and micro-phase analysis. The mechanical properties of samples after aging were also studied. The grain size increased substantially during the whole aging period. Two different types of grain boundary carbides were observed; block-shaped and needle-shaped. Both were identified to be M23C6 by selected area diffraction measurements. The grain boundary carbides did not coarsen significantly during aging, with the weight fraction increasing only from 0.20% to 0.28%. In contrast, a much higher coarsening rate of γ′ precipitates was observed, as evidenced from both TEM observations and SAXS analysis. 3DAP was used to study the elemental partitioning behavior between γ′ precipitates and γ matrix as well as the evolution in width of the γ/γ′ interface. A large increase in the width of γ/γ′ interfaces was seen between 1000 h and 3000 h aging. In addition, for the sample aged at 750 1C for 3000 h, Cr enrichment on the γ matrix side of the γ/γ′ interface was found. Tensile tests at 750 1C of the aged samples showed a gradual decrease in elevated-temperature yield strength after 500 h, when this alloy was over-aged. The critical precipitate size for the transition from precipitate cutting by weakly coupled dislocations to strongly coupled dislocations for Inconel 740H was calculated to be approximately 50 nm, which agrees well with the experiment measurements of elevated-temperature yield strength. The room temperature impact toughness of all samples decreased during aging as the grain size kept growing. & 2013 Elsevier B.V. All rights reserved.
Keywords: Inconel 740H Microstructure evolution γ′ precipitates Mechanical properties Three-dimensional atom probe
1. Introduction As a result of increasing energy demands and accelerated environmental problems, there is an urgent need to improve the thermal efficiency of coal-fired power plants. To achieve this goal, next generation ultra-supercritical power plants, operating at steam temperatures of 700–750 1C and operating pressure up to 37.5 MPa, are under development [1–3]. However, the increased operating parameters place more stringent requirements on the properties of candidate materials and cannot be met by conventional ferritic and austenitic steels. Consequently, there are efforts to replace these materials by Ni-base superalloys, which show a longer creep rupture life and higher corrosion resistance. Ni-based superalloys have already found widespread applications in a number of critical technological areas, especially those involving high temperatures, such as jet-engine turbines and coalfired power plants. These alloys typically exhibit an excellent
n
Corresponding author. Tel.: þ 86 10 6277 3460; fax: þ 86 10 6277 2853. E-mail address: [email protected] (C. Yan).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.09.076
balance of properties, including good mechanical properties and ductility (both at room temperature and elevated temperatures), improved fracture toughness and fatigue resistance, as well as enhanced creep and oxidation resistance at high temperatures [4]. One of the most promising candidate Ni-base superalloys for the main steam pipe of 700 1C ultra-supercritical coal-fired power plants is Inconel 740H, which is a modified version of Inconel 740 developed by Special Metals Corp, Huntington, West Virginia, USA. The nominal composition of Inconel 740H is 25Cr, 20Co, 0.5Mo, 1.5Nb, 1.4Ti, 1.4Al, 0.3Mn, 0.03Fe, 0.03C and Ni balance (wt%). Compared with Inconel 740, the ratio of Ti to Al in Inconel 740H is lowered in order to eliminate microstructure instabilities found in Inconel 740 during prolonged thermal aging at 750 1C [5–11]. In addition, the Nb content is also reduced to improve the weldability of the alloy. Since Inconel 740H has only recently been developed, few studies concerning the microstructure evolution and mechanical properties have been carried out so far on this material. Moreover, no reports exist on the detailed microstructure evolution during aging of Inconel 740H at the atomic scale. Microstructural stability is of prime importance for alloys to be used at high temperatures for long periods of time. Inconel 740H
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consists of a γ matrix with a coarse grain size, containing a dispersion of γ′ precipitates inside the grains, as well as carbides located at the grain boundaries. The microstructural evolution during long time aging, including grain size, γ′ precipitates and grain boundary carbides, should be studied at different scales in detail. In addition to traditional optical microscopy and electron microscopy observation, three-dimensional atom probe (3DAP) has been used in this study for the first time. In the last decade, there have been large numbers of studies employing 3DAP to characterize the microstructure of Ni-base superalloys at the atomic scale, including the size, morphology and composition of γ′ precipitates within the γ matrix as a function of varied heat treatments [4,12–15]. In particular, the capability of 3DAP to provide information on elemental partitioning between phases, and on segregation to certain material defects, lend it a
competitive edge in research related to interfaces and grain boundaries in Ni-base superalloys. The change of mechanical properties in Ni-base superalloys with aging time is also an important issue that needs to be investigated. The room temperature impact toughness and microhardness of Inconel 740H have already been studied after aging at 750 1C for 1000 h at different temperatures [6]. Also, the elevated-temperature microhardness of the alloy in the standard heat treatment condition has been studied, and the results compared with room temperature microhardness [6]. However, studies of the evolution of elevatedtemperature strength and of room temperature toughness of Inconel 740H with aging time, as well as their relationship with microstructure have not been reported yet. In this work, tensile tests at 750 1C and impact toughness measurements at room temperature have been carried out for samples aged up to 3000 h at 750 1C. The results are analyzed with respect to the observed changes in the size of the γ′precipitates and the grain size.
2. Experimental procedure
Fig. 1. The geometries of the mechanical test samples (a) elevated-temperature tensile test and (b) room temperature impact toughness test.
The bulk chemical composition of the Inconel 740H sample was analyzed to be 49.27Ni–25.57Cr–19.92Co–1.48Nb–0.50Mo–1.38Ti– 1.47Al–0.34Mn–0.03Fe–0.02C wt%. Samples cut from the forged bars were solution treated at 1150 1C for 30 min and then water quenched. These samples were subsequently aged at 750 1C for 100, 300, 500, 1000 and 3000 h and then air quenched. For convenience, these samples will be subsequently referred to as WQ100, WQ300, WQ500, WQ1000 and WQ3000. After aging, the specimens were ground and polished following standard metallographical methods and then chemically etched in a fresh solution of HCl, HNO3 and H2O volume proportions of 10:1:10 at 50 1C for 15 min. Optical microscope (OM) investigations were conducted using an Olympus GX51 microscope. Scanning electron microscope (SEM) investigations were conducted using a Hitachi 4500 scanning electron microscope. Samples for transmission electron microscope (TEM) investigations were prepared by a standard electro-polishing technique using a double-jet device.
Fig. 2. Optical micrographs of the samples aged at 750 1C for different times.
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Slices of 0.3 mm thickness were cut out of the bulk material and mechanically ground to 40–50 μm in thickness. Discs of 3 mm in diameter were then punched from the thinned slice and then electro-polished at 30 1C at a voltage of 16 V, using a solution of 85% ethanol and 15% perchloric acid. TEM investigations were conducted using a JEOL 2010 TEM operated at 200 kV. The weight fractions of the γ′ precipitates and carbides were ascertained through micro-chemical phase analysis methods. The γ′ precipitates were electrolytically extracted in an aqueous solution containing 1% (NH4)2SO4 and 1% C6H8O7.H2O from 0 to 5 1C. The carbides were electrolytically extracted in a methanol solution containing 5% HCl, 5% C3H8O3 and 1% C6H8O7 H2O from 10 to 5 1C. In both cases a current density of 0.05 A/cm2 was used. The particle size distribution and mean radius of the extracted γ′ precipitates were determined using an X-ray diffractro-spectrometer equipped with a Krathy small angle scattering (SAXS) goniometer. Samples for 3DAP tomography studies in the LEAP microscopy were prepared by an electro-polishing method. For this purpose,
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samples of different ageing conditions were first electro-discharge machined into thin wires with a square cross section of 0.7 0.7 mm2. These wires were mechanically ground and subsequently electro-polished in two steps; first with a 75% acetic acidþ 25% perchloric acid solution for the coarse polish, and finally with a 98% butyl cellulose þ2% perchloric acid solution at 12–14 V for the final polish. The 3DAP tomography experiments, were performed using an Imago (now Camera Instruments) local electrode atom probe (LEAP 3000 h) at a residual pressure of 5 10–9 Pa and a specimen temperature of 50 K, and with a pulse repetition frequency of 200 kHz and a pulse-voltage to dc-voltage ratio of 15%. Data analysis was performed using the IVAS 3.6.0 software. Impact testing of Charpy V-notch samples cut from the aged samples, was performed at room temperature. Tensile tests at 750 1C were carried out on aged samples in a hydraulic test system equipped with a temperature-controlled furnace. The temperature fluctuations were kept within 3 1C. The load rate before yielding was 0.5 mm/min and increased to 2.5 mm/min after yielding.
Fig. 3. SEM micrographs and the typical EDS of the grain boundary carbides aged at 750 1C for different times.
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The geometries of the Charpy V-notch samples and tensile test samples are provided in Fig. 1.
3. Results and discussion 3.1. Microstructure evolution 3.1.1. Grain size The microstructures of the WQ100, 300, 500, 1000, 3000 samples as seen in the optical microscope are shown in Fig. 2. The microstructure consists of a coarse grain structure containing annealing twins. It is evident that grain growth took place during aging. The average grain sizes, determined by an intercept method, for each aging time were 108, 112, 120, 137, 151 μm respectively.
3.1.2. Grain boundary carbides The evolution of the grain boundary carbides with aging time is illustrated in Fig. 3. The carbides were identified to be Cr rich M23C6 according to energy dispersive spectrometer (EDS) measurements of the carbides in Fig. 3, which is consistent with previous studies of Inconel 740 by Xie et [5]. There was only marginal coarsening of the carbides at prolonged aging time, with the weight fraction increasing just from 0.20% to 0.28% (Fig. 4(a)). TEM images and selected area diffraction (SAD) patterns of the grain boundary carbides are shown in Fig. 5. In general, the grain boundary carbides exhibited two different shapes; block-shaped most common, and needle-shaped rarely. The needle-shaped carbides appear to be nucleated from the grain boundaries and to grow perpendicularly into the grains, similar to the morphology of η phase reported by Zhao on the microstructure evolution of
Fig. 4. Evolution of phase fraction (wt%) with aging time (a) carbides, (b) γ′ phase.
Fig. 5. TEM micrographs and corresponding SAD patterns of different shape grain boundary M23C6. (a) and (c) block-shaped M23C6 and corresponding SAD pattern, (b) and (d) needle-shaped M23C6 and corresponding SAD pattern.
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Inconel 740 aged at 750 1C [6,7]. However, SAD patterns of the needle-shaped phase confirm that it is Cr rich M23C6 carbides rather than η phase. 3.1.3. γ′ Phase 3.1.3.1. Temporal evolution of γ′ morphology. As the main contributor to high temperature strength in Ni-base superalloys, the morphology, size and chemical composition of γ′ precipitates are of primary importance in an investigation of Inconel 740H. TEM images, together with diffraction pattern showing typical superlattice reflections from γ′ precipitates for the WQ100, 300, 1000, 3000 samples are given in Fig. 6. No γ″ precipitates were detected from either the TEM images or the SAD patterns. The γ′ precipitates exhibited a spheroidal morphology despite of becoming slightly cuboidal after long time aging. To determine precisely the variations in size distribution and mean radius of γ′ precipitates with aging time, SAXS experiments were also carried out. The results are shown in Fig. 7. In general the γ′ precipitates exhibit a one-modal size distribution, with evident coarsening during aging, from a mean precipitate radius increasing of 22.4– 69.2 nm. The γ′ precipitate weight fraction increased from 14.0% to 16.5%, as shown in Fig. 4(b). In the first 100 h, all γ′ precipitates remained below 30 nm in radius, whereas after that γ′ precipitates of size 30–60 nm were detected. In the samples aged longer than 1000 h, even larger γ′ precipitates of size 60–90 nm were detected. It should be noted however that, these larger precipitate sizes were not included in the calculation of the mean precipitate size, as it was assumed these were in fact clusters of several small γ′ precipitates together. At the same time, the number density of precipitates decreased with aging time (Fig. 6). Although the
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morphology of γ′ precipitates can be clearly seen in TEM foils, the elemental distribution between γ′ precipitates and γ matrix, as well as the compositional profile across the γ/γ′ interfaces cannot be precisely obtained due to the small size of the γ′ precipitates. For this purpose, 3DAP studies were carried out to study the microstructure at the atomic scale. 3.1.3.2. Chemical analysis of this alloy. Examples of 3DAP reconstructions from the WQ100 and 3000 samples are shown in Fig. 8 (a) and (b). The size of each reconstruction volume is shown in the figure. The Co rich regions (blue) corresponded to the γ matrix, whereas the red Cr ¼25 at% isosurfaces are used to delineate the outlines of the γ′ precipitates. According to previous work [14], the concentration value used to delineate the Cr isosurfaces was selected to be the average value of the concentrations in the γ matrix and γ′ precipitates (in this case E25%). Owing to the limited size of the reconstruction volume, only part of each γ′ precipitate can be seen in the reconstructed volume, although a few complete small precipitates are seen in the WQ100 sample. Nevertheless, the 3-D morphology of the γ′ precipitates is seen to be near spherical, in agreement with the TEM investigations. We take here the WQ3000 sample as an example to investigate the partitioning behavior of different alloying elements in Inconel 740H. The results are illustrated in Fig. 9. By comparing Figs. 8 and 9, it is seen that different elements have a different partitioning behavior between the γ matrix and the γ′ precipitates. Cr, Co and Mo tend to distribute in the γ matrix, while Al, Ti and Nb preferentially segregate to the γ′ precipitates. This segregation tendency is consistent with other studies of Ni-base superalloys [13,16–18]. It is worth noting that the level of segregation for
Fig. 6. TEM images of γ′ precipitates in samples aged for different times at 750 1C.
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Fig. 7. Size distribution of γ′ measured by SAXS in Inconel 740H at different aging times.
each element was not the same. For example, segregation of Mo in γ matrix was rather low compared with that of Cr and Co. Likewise, the segregation levels for different γ′ segregating elements were also different. A lever rule diagram, first implemented by Blavette et al. [19], was constructed to determine in detail the partitioning behavior as well as partitioning intensity of the solutes (Fig. 10). In the diagram, the nominal concentration for each alloying element minus the solute content in the γ matrix is plotted versus the difference in composition between γ′ precipitates and γ matrix. The slope of the best-fit line passing through the origin for all elements gives the volume fraction of γ′ precipitates. The elements segregating to the γ′ precipitates are seen in the first quadrant of the coordinate system. Points to the right of the diagram correspond to stronger segregation in γ′ precipitates (vice versa for the γ matrix segregating elements). In agreement with the results shown in Fig. 9, the lever rule diagram shows that Al, Ti and Nb tend to segregate in γ′ precipitates, with segregation intensity in descending order. Similarly in the γ matrix the Cr, Co and Mo segregate in descending order. In order to quantify the partitioning behavior of different elements during aging, we define the solute partitioning ratio as γ′ γ γ′ γ K γ ′=γ ¼ C i =C i , where C i and C i are the solute concentrations of element i in the γ′ precipitates and in the γ matrix, respectively.
Values of K γ ′=γ above 1 indicates γ′ segregating elements, and for γ segregating elements, the value is below 1. Moreover an increasing value of K γ ′=γ corresponds to an increasing tendency for segregation to the γ′ phase. The values of K γ ′=γ for different elements during aging are summarized in Fig. 11. For the γ segregating elements, an interesting phenomenon was observed, namely that the values of K γ ′=γ for Cr, Co and Mo all reached a peak value at 300 h, which means that, the γ′ precipitates contained highest contents of Cr, Co, Mo after aging at 750 1C for 300 h. The reason for this phenomenon might be as follows. Elements such as Cr, Co and Mo that segregate to γ matrix diffuse from γ′ precipitates into the γ matrix during growth of γ′ precipitates. As observed in Fig. 7, a sudden growth in size of γ′ precipitates was seen between 100 h and 300 h. This growth may leave insufficient time for diffusion of Cr, Co and Mo into the γ matrix. After 300 h aging, the growth rate of the γ′ precipitates slowed down, leaving sufficient time for diffusion. It is suggested therefore that, the combined effects of size growth and repartitioning by diffusion lead to an enrichment of Cr, Co and Mo inside the γ′ precipitates at 300 h. The trends of K γ ′=γ for Ti and Al with aging time were opposite, indicating that with prolonged aging time, Ti in the γ′ phase was gradually replaced by Al. The trend of K γ ′=γ with aging time, however, had nothing in common with other elements.
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Fig. 8. 3-D reconstruction showing Co atoms in blue and isosurfaces for Cr ¼ 25 at% in red for more clear visualization of the γ′ precipitates for (a) WQ100 and (b) WQ3000. (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
3.1.3.3. Temporal evolution of the γ/γ′ interface. The compositional profiles across the γ/γ′ interfaces for the WQ100, 300, 1000, 3000 samples were determined using a powerful analysis tool, the proxigram, for the various γ′ precipitates shown in the reconstruction volumes in Fig. 8. Widely employed in other studies using 3DAP, proxigrams have been shown to be useful for the study of any internal surfaces, even those involving a complex topology [14,15]. In this study proxigrams for the γ/γ′interfaces were constructed based on isosurfaces with a threshold value of 25 at% Cr. The results are shown in Fig. 12, where only the main alloying elements (Cr, Co, Mo, Ti, Al, Nb) are shown for simplicity. The proxigrams show that the compositional profile across the γ/γ′ interface is not abrupt but
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transitionary. The width of the transitionary region is defined as the interface width. According to Yoon [19], the choice of Cr threshold value used to fix the isosurfaces only influences the location of the interface, but not the width of the interface. The width of the interface in each proxigram was determined from the Al composition profile, between the steady-state γ′ and γ. The width of the γ/γ′ interfaces was determined using a 10–90% of the plateau method (long-range γ and γ′ compositions) [14,15]. The results are shown in Fig. 12. The interface width in the WQ100 sample was about 2.0 nm, similar to the results of Rene88DT [14]. The interface width was relatively stable over the first 1000 h with a large increase between 1000 h and 3000 h, which is believed to be associated with the growth of the precipitates and the misfit strain around the interface. It is worth noting that in the WQ3000 sample, an enrichment of Cr was found on the γ matrix side of the γ/γ′ interface, (Fig. 13). This phenomenon was not found in the other aged samples. The width of the enriched region was about 2.5 nm. Yoon observed a similar phenomenon in the commercial Ni-base superalloy RENE N6, but the reason for this enrichment was not discussed [19]. In addition, the enrichment of Co at the interface was also detected in RENE N6. In contrast this was not found in the WQ3000 sample investigated in the present study. Two possible explanations have been proposed for the enrichment of Cr at the γ/γ′ interfaces. The first is based on a diffusion mechanism. As shown in Fig. 10, Cr segregates strongly to the γ matrix. Therefore, with the increase of γ′ precipitate size, excess Cr inside the γ′ precipitates must be excluded and must diffuse across the γ/γ′ interface into the γ matrix. A low diffusion coefficient of Cr across the interface might therefore lead to the enrichment of Cr at the interface. However, this explanation cannot account for the absence of Cr enrichment at earlier aging times. The other explanation is associated with the suggestion that a new phase Cr-rich may form at the γ/γ′ interface. It has been reported that a Ni2Cr-type long-range ordered phase can be expected in Ni–Al alloys with high Cr and Mo content. In addition to the binary Ni–Cr system such a phase has also been found in ternary Ni–Cr–Mo [20,21] and Ni–Cr–Al alloys [22]. The Ni2Cr superstructure (Pt2Mo type), with orthorhombic unit cell, forms in the Ni–Cr system below 590 1C. It has been shown that addition of Mo to binary Ni2Cr significantly increases the critical temperature for the formation of this phase, e.g. in Ni–25Mo–8Cr, the critical temperature of Ni2(Cr,Mo) phase is 750 1C [23,24]. Additionally Svoboda and his co-workers found the existence of Ni2Cr at the γ/γ′ interface in a Ni–Cr–Al–Mo system aged at 600 1C for 3000 h [25]. Sundararaman in Alloy 625 also found similar results [26]. Intermetallic phase Ni2(Cr,Mo) precipitates with Pt2Mo-type structure have also been observed, in addition to that of the γ″ phase, in Alloy 625 after prolonged ( 70,000 h) service at temperatures close to but less than 600 1C. Identification of the nature of the Cr enrichment at the γ/γ′ interface will therefore be the subject of further research. 3.2. Mechanical properties 3.2.1. Tensile strength at 750 1C The results of tensile tests at 750 1C of Inconel 740H aged at 750 1C for 100, 300, 500, 1000 and 3000 h are shown in Fig. 14. The yield strength of Inconel 740H at 750 1C increased in the first 500 h and then decreased gradually with further increase in aging time. The elongation increased during the whole aging. The changes in yield strength can be rationalized on the basis of the growth of γ′ precipitates based on two previously reported models [27,28]. In both cases, it is believed that the critical resolved shear stress (CRSS) is determined by the force necessary to move two coupled edge dislocations in the 〈110〉 direction on the {111} plane through the γ′ precipitates. The first model [27] describes the CRSS for cutting through small γ′ precipitates in a disordered matrix by
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Fig. 9. Elemental partition between γ and γ′ phase of sample aged at 750 1C for 3000 h.
two weakly coupled dislocations while the second [28] describes the same motion for cutting of larger γ′ precipitates by strongly coupled dislocations. In the first model, the CRSS for two weakly coupled dislocations is given by 1 3 1 Γ 2 bdf 2 1 Γ f ð1Þ Δτ0 ¼ A 2 T 2 b b where Γ is the anti-phase boundary energy of the γ′ precipitates in the {111} plane, b is the Burgers vector of the edge dislocation in the γ matrix, d is the γ′ precipitates diameter, f is the volume fraction of the γ′ precipitates, T is the line tension of the dislocation and A is a numerical factor depending on the morphology of the particles. For spherical particles A¼0.72 [29]. The line tension is given by 2
T¼
Gb 2
ð2Þ
where G is the elastic shear modulus and is taken as 60 GPa in this study.
For cutting of larger precipitates by strongly coupled dislocation pairs, the CRSS is given by 1
Tf 2 ω Δτ0 ¼ 0:86 bd
!
1:28dΓ ωT
12
ð3Þ
where w is a constant accounting for the elastic repulsion between the strongly paired dislocations, and which is of the order of unity. For quantitative calculation, f is taken as the average phase fraction of γ′ precipitates (15%), w is taken as 1 for simplicity, and Γ is taken to be 0.42 J/m2. Fig. 15 is a plot of γ′ precipitate diameter d versus CRSS Δτ0 for the two dislocation mechanisms. For precipitate cutting by weakly coupled dislocations, Δτ0 monotonically increases with increasing precipitate diameter. In contrast Δτ0 gradually decreases with increasing precipitate diameter for the mechanism of precipitates cutting by strongly coupled dislocations. The resultant Δτ0 for a given d is the lower value of the two, since dislocation activity follows the mechanism that provides the least resistance to glide.
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Fig. 10. Lever rule diagram of Inconel 740H(WQ1000). The nominal concentration (Cn) for each element (Cr, Al, Ti, etc.) minus the solute content in the γ matrix (Cγ) is plotted as a function of the difference in composition between γ and γ′ (Cγ Cγ′). The slope gives the molar fraction of γ′ present in Inconel 740H.
Fig. 11. Partitioning ratios, K γ′=γ , for main elements in Inconel 740H plotted as a function of aging time.
For any given alloy, there exists a critical value of d that gives the highest flow stress. For Inconel 740H, the optimum value of d was calculated to be around 50 nm (shown in Fig. 15), which coincides well with the experimental results of yield strength shown in Fig. 14. In the first 500 h aging, when the diameter of γ′ precipitates is smaller than 50 nm, the precipitates cutting by weakcoupled pairs of dislocations is the dominant mechanism and the yield strength increases with the increasing precipitate diameter. However, after 500 h as the diameter of γ′ precipitates exceeds the critical value of d, the mechanism changes to cutting by strongcoupled pairs of dislocations. As a result, the yield strength of Inconel 740H decreases monotonically with longer aging time.
time is evident. For the sample aged at 750 1C for 3000 h, the value decreased to less than 20 J, which would present a huge threat to the safety of use in operation. As reported by Zhao [6], the fractographs of Inconel 740 samples heat-treated at 750 1C indicated a clearly brittle fracture with a localized mixed-mode behavior and a clear intergranular fracture behavior. Therefore, the impact toughness of the aged samples will be influence by the change in grain size and by the size of the grain boundary carbides. As mentioned in Sections 3.1.1 and 3.1.2, during aging the size of grain boundary carbides remained unchanged, whereas the grain size increased. Therefore, it can be concluded that the drastic loss of impact toughness at room temperature for aged samples is mainly due to the growth in grain size.
3.2.2. Impact toughness at room temperature The results of the impact toughness measurements at room temperature of Inconel 740H aged at 750 1C for 100, 300, 500, 1000 and 3000 h are shown in Fig. 16. A loss of impact toughness with aging
4. Conclusions The microstructure evolution and mechanical properties of Inconel 740H during aging at 750 1C up to 3000 h have
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Fig. 12. Proxigrams showing the compositions profiles of Cr, Co, Mo, Ti, Al and Nb atoms across the γ/γ′ interfaces for (a) WQ100, (b) WQ300, (c) WQ1000 and (d) WQ3000 conditions.
900 34 850 32
750 28 700 26 650 24 600
Strength/MPa
Elongation/%
800 30
22 550 20
A
500
18 450 100
Aging time/h1000
Fig. 14. The change of mechanical properities of Inconel 740H with aging time.
Fig. 13. Proxigram across the γ/γ′ interface of WQ3000, note that γ is on the right and γ′ is on the left.
been investigated with a combination of OM, SEM, TEM, 3DAP, SAXS, micro-chemical analysis and elevated-temperature tensile tests, as well as room temperature impact toughness
measurements. The main conclusions of this study are as follows: (1) The average initial grain size of Inconel 740H is above 100 μm and evident grain growth takes place during aging. The grain boundaries are fully decorated with Cr-rich M23C6 carbides, which does not exhibit significant coarsening during aging up
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Fig. 15. Theoretical critical resolved shear stress versus γ′ precipitates size relationships in Inconel 740H deformed at 750 1C.
around the γ/γ′ interface. Enrichment of Cr at the γ matrix side of γ/γ′ interface is found for the sample aged for 3000 h. (4) There exists a critical value for the diameter of γ′ precipitates that gives optimum yield strength of Inconel 740H. When the diameter of γ′ precipitates is below the critical value, the yield strength increases monotonically with increasing precipitate size, while when the size is above the critical value, the yield strength decreases with the increase of precipitate size. The critical value for Inconel 740H was calculated to be around 50 nm, which agrees well with the experimental results. A loss of room temperature impact toughness of Inconel 740H with aging time was evident. The main reason for this was attributed to the increase in grain size during aging.
Acknowledgments
Fig. 16. The change of impact toughness at room temperature of Inconel 740H with aging time.
to 3000 h, with the carbide weight fraction increasing from 0.20% to 0.28%. Two different forms of grain boundary carbides, namely block-shaped and needle-shaped, are observed. Both are identified to be M23C6 by selected area diffraction measurements. (2) TEM and 3DAP image data, together with the results of SAXS experiments, indicated that size distribution of the spheroidal γ′ precipitates inside each grain is one-modal and that the mean precipitate radius increases rapidly during aging at 750 1C. The elemental partitioning behaviors have been studied with the help of 3DAP measurements and the results analyzed using lever rule diagrams and the solute partitioning ratio K γ ′=γ . It was found that, Cr, Co, Mo tend to distribute in the γ matrix, while Ti, Al, Nb preferentially segregate to the γ′ precipitates. The partitioning intensities for different elements are not the same and vary with aging time. (3) The compositional profiles across the γ/γ′ interfaces were investigated using proxigrams and the compositional widths of the interface determined for different aging times. An increase in the γ/γ′ interface width was found between 1000 h and 3000 h aging, which was believed to be associated with the growth of the γ′ precipitates and the misfit strain
The authors are grateful to the financial support from the HighTech Research and Development Project (No. 2102AA03A501) supported by the National Ministry of Science and Technology of China.
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