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Microstructure evolution of a Ni–Cr–W superalloy during long-term aging at high temperatures
Materials Characterization 106 (2015) 283–291
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Microstructure evolution of a Ni–Cr–W superalloy during long-term aging at high temperatures Ho Jung Lee a, Hyunmyung Kim a, Donghoon Kim b, Changheui Jang a,⁎ a b
Department of Nuclear and Quantum Engineering, Korea Advanced Institute of Science and Technology, 305-701 Daejeon, Republic of Korea Agency for Defense Department, 305-600 Daejeon, Republic of Korea
a r t i c l e
i n f o
Article history: Received 29 April 2015 Received in revised form 19 June 2015 Accepted 19 June 2015 Available online 20 June 2015 Keywords: Ni–Cr–W superalloy Long-term aging Microstructure evolution Carbide dissociation Tensile elongation
a b s t r a c t The effects of long-term exposure to high temperature (800 and 900 °C for up to 20,000 h) on the microstructure evolution of a Ni–Cr–W superalloy (Alloy 230) were investigated. After long-term aging at 800 °C, extensive precipitation of the secondary Cr-rich M23C6 carbides was observed. In addition, the internal dissociation of W-rich primary η-M6C carbides was observed, such as, W-rich M6C → α-W, Cr-rich M23C6, and W-depleted Ni-base matrix. After long-term aging at 900 °C, precipitation of Cr- and Ni-rich phase and α-W phase was observed in some areas at the expense of the secondary Cr-rich M23C6 carbides. Meanwhile, the dissociation of the primary M6C carbides was less significant at 900 °C. Long-term aging at both temperatures resulted in a decrease in tensile elongation while an increase in strength was minimal after aging at 900 °C. The degradation of tensile property depended on the evolution of Cr-rich M23C6 carbides, Cr- and Ni-rich precipitates, and α-W phases but not on the primary M6C carbides. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Among generation-IV (Gen-IV) nuclear reactors under development, a very high temperature gas-cooled reactor (VHTR) will be operated at the temperature range of 800–900 °C for up to 60 years [1]. For the application of an intermediate heat exchanger (IHX) in a VHTR, a solid solution hardened Ni–Cr–W superalloy, Alloy 230, is considered one of the candidate materials because of the superior mechanical properties and oxidation resistance at high temperature [2–4]. It has been reported that the W-rich primary M6C carbides are formed in Alloy 230 during solidification process [2,5]. They contribute to additional increase in high temperature strength and creep resistance by pinning grain boundaries and thereby inhibiting grain boundary sliding [6,7]. As mechanical properties are affected by various microstructural features, evaluation of the microstructural changes associated with long-term exposure at high temperature is important. Previously, the effects of microstructure on mechanical properties were investigated for Ni-base superalloys with similar Cr content to Alloy 230 [8–13]. That is, various types of carbides (M23C6 and M6C) and precipitates (γ′, CrMo(C,N), and TiN for Alloy 617 and γ″, δ, and Ni2(Cr,Mo) for Alloy 625) were formed during high temperature exposure depending on exposure temperatures and chemical compositions. The presence of those precipitates resulted in hardening and reduction of ductility and creep resistance by disturbing dislocation movement.
⁎ Corresponding author. E-mail address: [email protected] (C. Jang).
http://dx.doi.org/10.1016/j.matchar.2015.06.018 1044-5803/© 2015 Elsevier Inc. All rights reserved.
However, in case of alloy 230, there are limited studies on the effects of the long-term aging on the mechanical properties of Alloy 230 [5,11, 14–17]. Previously, formation of the Cr-rich M23C6 carbides at the boundary of the primary M6C carbides was also reported when exposed to the temperature ranges of 650–1000 °C [16,18]. Moreover, Veverkova [16] and Wu [19] reported the breakdown of the primary carbides after long-term exposure at the temperature ranges of 650–870 °C. The issue in other Ni-base superalloys has also been reported recently to assure long-term stability for high temperature applications [20,21]. However, the mechanism and nature of the phases formed in Alloy 230 were not properly characterized in the previous studies. Therefore, in this study, the long-term aging behaviors of Alloy 230 at the anticipated operating temperatures of a VHTR were assessed focusing on the microstructure evolution such as the formation and dissociation of carbides. Then the consequent changes in the tensile properties associated with such microstructure evolution were investigated. 2. Materials and experimental The chemical composition of a commercial grade Alloy 230 used in the study was analyzed by inductively coupled plasma mass spectroscopy (ICP-MS) method and the result is listed in Table 1. The test material was provided as solution annealed at 1177 °C for 0.5 h followed by water quenched. For the aging process, blocks of Alloy 230 were exposed at 800 and 900 °C for 10,000 and 20,000 h in a box furnace. After aging, the blocks were removed from the furnaces and aircooled to room temperature. The specimens for the microstructure observation and tensile test were taken at 2 mm from the surface of the
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Table 1 Chemical composition of Alloy 230 used in the study (wt.%). Ni
Cr
W
Mo
Fe
Mn
Si
Al
Co
C
La
Bal.
22.3
13.9
1.37
1.34
0.5
0.28
0.25
0.15
0.1
0.012
block to avoid environmentally induced degradations such as internal oxidation and carburization/decarburization. The microstructure features including the carbides and other phases were observed using a scanning electron microscope (SEM, Magellan 400) with back scattered electron (BSE) mode at 10 kV. In SEM sample preparation, specimens were mechanically polished down to 1 μm using SiC papers and diamond polishing compound and ultra-sonically cleaned in ethanol. The polished samples were etched in solutions of hydrogen chloride and 2 vol.% bromine in methanol in sequence to reveal grain boundaries. The volume fractions of primary and secondary carbides were measured using an image analyzer (TOMORO analySIS TS) as follow: 1) Low magnification SEM and BSE images acquired by a dual mode of SEM. 2) Total volume fractions of primary and secondary carbides were measured with SEM images because the contrast of those phases was different from the matrix. 3) Volume fraction of the primary carbides was measured using BSE images because primary carbides with high atomic number are easily distinguished from secondary carbides and matrix. 4) Volume fraction of the secondary carbides was estimated by excluding volume fraction of primary carbides from total carbide volume fraction measured in step 2. However, the image analyzer could not accurately distinguish phases with very thin shape or of similar contrast, which will be discussed in Section 3.2. Also, the grain size of Alloy 230 before and after the aging was measured with circle intercept method following the procedure of ASTM E112 using the image analyzer. In addition, an X-ray diffractometer (XRD, RIGAKU D/MAX-2500) and a transmission electron microscopy (TEM, Titan cubed G2 60-300) equipped with energy dispersive spectroscopy (EDS) and selective area diffraction (SAD) at 300 kV were used to identify the phases before and after the aging. The TEM samples were prepared by a focused ion beam (FIB, Helios Nanolab 450 F1) with area of 10 × 10 μm and around 50 nm in thickness. For the tensile tests, mini-sized tensile specimens with 16 mm in length and 0.5 mm in thickness were machined. The tests were conducted at room temperature under displacement control (strain rate of 3.33 × 10− 4/s) following the procedures of ASTM E8/ E8M-13a. Due to the mini-size of the specimen, the extensometer was not attached. Instead, the strain was estimated from the displacement divided by the gage length of mini-size specimen. In the XRD analysis, in addition to the bulk samples, phase extracted samples were used in order to clearly identify the peaks from the carbides without the interference of those from matrix. In the phase extraction process, samples were mechanical polished to eliminate impurities on the surface and then ultrasonically cleaned in ethanol. The samples were soaked into the solution of 10 vol.% bromine in ethanol until all metallic phases were dissolved. Then the solution containing
the extracted particles was filtered through the porous filter paper to collect the particles selectively. Samples were analyzed with a high applied power of 40 kV and 300 mA. The incident angle of the X-ray beam was 4° in the 2 theta mode within the range of 20 to 80°.
3. Results and discussion 3.1. Microstructure evolutions during aging at 800 °C Fig. 1 shows the microstructure of an as-annealed Alloy 230. Grain boundaries are covered with fine discrete carbides which are known as the secondary Cr-rich M23C6 formed during the cooling process from solution annealing [7,11,22]. In addition, heavy and globular shaped primary carbides (bright color in BSE image in Fig. 1(b)) are randomly distributed on the grain boundaries and within the grains with various sizes (1 to 20 μm). In previous studies, these carbides were reported as the primary W-rich M6C type carbide in the form of η-Ni3W3C through TEM/SAD [2,11,18], XRD [16,22] and synchrotron radiation [23] analyses. The volume fractions of the carbides are measured by an image analyzer and listed in Table 2. As summarized in Table 2, volume fraction of W-rich M6C type carbide is about 4%, while the amount of Cr-rich M23C6 grain boundary carbides too thin to be quantitatively measured. After 10,000 and 20,000 h aging at 800 °C, significant microstructure changes are observed as compared to the as-annealed condition as shown in Fig. 2. For the specimens aged at 800 °C, the extensive precipitation of the secondary Cr-rich M23C6 carbides (gray color in BSE image in Fig. 2(b) and (d)) is observed on the grain boundaries as well as within the grains. Also, the amount of the secondary carbides increases with aging time. Such behaviors are in good agreement with the previous reports on the formation of the secondary Cr-rich M23C6 carbides in Ni-base superalloys exposed at various aging temperature ranges [10, 15,24]. On the other hand, the volume fraction of the primary M6C carbides is not significantly altered even after 20,000 h aging at 800 °C. Instead, areas of different contrasts are observed inside of the primary M6C carbides boundary for the aged specimens as shown in SEM/BSE micrographs (Fig. 2(b) and (d)), indicating the possibility of ‘internal dissociation’ during high temperature aging. It is similar to the so-called ‘breakdown’ phenomenon previously reported by Veverkova [16] and Wu [19]. They observed the breakdown of the primary M6C carbides in Alloy 230 after over 10,000 h of isothermal aging in the temperature ranges of 700–870 °C. Wu suggested the transformed phase as W2C carbide based on the EDS analysis and SEM observation [19]. On the other hand, the crystal structure and chemical composition of the primary M6C carbides were reported the same as those found in the as-annealed condition after a short-term exposure up to 3000 h in the temperature range of 650–1000 °C [2,11,16]. Such observations suggest that the transformation phenomenon is a slow process, possibly associated with the diffusion of slow-moving element. The phase transformation within the
Fig. 1. SEM micrographs of the as-annealed Alloy 230; (a) low and (b) high magnifications (SEM/BSE images).
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Table 2 Volume fractions of various phases present in Figs. 1–3. As-annealed Alloy 230
Prior to aging 10,000 h aging 20,000 h aging a b c d
Aging at 900 °Cd
Aging at 800 °C
W-rich M6C
Cr-rich M23C6
~4%
a
W-rich M6Cb
Cr-rich M23C6
W-rich M6C & α-Wc
Cr-rich M23C6
4–5% 4–5%
7–10% 13–14%
3–4% 3–5%
5–6% 5–8%
Measurement could not be made due to very low volume fraction. Volume fraction of the phases within the prior M6C carbides. Volume fraction of un-dissociated and dissociated M6C carbides plus α-W on the grain boundary. Volume fraction of Cr- and Ni-rich precipitates was not measured as it was not distinguishable from the matrix.
primary M6C carbides, or internal dissociation, will be further analyzed and discussed in the later sections.
3.2. Microstructure evolutions during aging at 900 °C The microstructures of Alloy 230 after 10,000 and 20,000 h aging at 900 °C are shown in Fig. 3. Unlike those after aging at 800 °C, the precipitation of the secondary phases is rather irregular after aging at 900 °C. That is, along with the areas with extensive precipitation of the secondary M23C6 carbides (for example, both side of Fig. 3(b)), there are areas with very little secondary M23C6 carbides (carbide depleted region, middle of Fig. 3(b)). In the areas with very little secondary carbides, a number of large and elongated phases are developed along the grain boundaries, which are identified as Cr- and Ni-rich (52 Cr and 38 Ni in at.%) and W-rich (83 at.% W) precipitates by the point TEM/EDS analysis (details will be provided in Sections 3.4 and 3.5). The Cr- and Ni-rich phases are indicated by arrows in the figures as the contrast is weak in the SEM photos, probably due to the fact that the number of electron per atom is similar in both Cr- and Ni-rich phases and matrix. Meanwhile, W-rich precipitates are bright and easily distinguishable from Cr- and Ni-rich precipitates. Precipitation of Cr- and Ni-rich phases in Alloy 230 was previously reported and described as the ‘pool-like precipitates’ [16] or ‘intermetallic phase of Cr0.8Ni0.2’ [11] after long-term exposure at above 950 °C including 500 h exposure at 1000 °C. In addition, the dissociation of some primary M6C carbides is observed mostly in the carbide depleted region while most of them remain similar morphology to the as-annealed condition (Fig. 3(b) and (c)).
It is interesting that, compared to the specimens aged at 800 °C, the volume fraction of the secondary M23C6 carbides are much smaller in the specimens aged at 900 °C, which is expected from the existence of carbide free zone in Fig. 3. Instead, both Cr- and Ni-rich precipitates and W-rich precipitates are present in large quantity, though they were not quantitatively analyzed because the contrast of Cr- and Nirich precipitates and W-rich precipitates is similar to that of the matrix and primary M6C carbides respectively. Meanwhile, as aging time increases from 10,000 h to 20,000 h, the amount of the secondary M23C6 carbides increases but only by a small degree (Fig. 3(a), (b), and Table 2). Also the dissociation of the primary M6C carbides further progresses with aging time, while the total amount of dissociated and nondissociated M6C carbides remains almost the same (Table 2). Table 3 shows the measured grain size of as-annealed and aged Alloy 230 at 800 and 900 °C up to 20,000 h. The grain sizes are almost the same irrespective of the aging conditions and the ASTM grain size number is around 5.5. Whittenberger [17] also reported that the grain size of Alloy 230 was not changed after aging at 820 °C up to 22,500 h. It is thought that the precipitation of thick and continuous Cr-rich M23C6 grain boundary carbides act as efficient obstacles against grain growth. 3.3. XRD analysis for phase identification XRD analyses were performed to identify the phases in Alloy 230 before and after the long-term aging and the results are shown in Fig. 4. In the figure, the XRD results of both bulk and phase-extracted samples are presented. In the as-annealed Alloy 230 (Fig. 4(a)), high intensity peaks of matrix and M6C carbides are detected indicating that the M6C
Fig. 2. SEM micrographs of Alloy 230 after (a), (b) 10,000 h and (c), (d) 20,000 h aging at 800 °C (SEM/BSE images).
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Fig. 3. SEM micrographs of Alloy 230 after (a), (b), (c) 10,000 h and (d), (e), (f) 20,000 h aging at 900 °C (SEM/BSE images).
carbides are dominant precipitates. Along with M6C carbides, the existence of small amount of M23C6 carbides is confirmed as the shallow peaks of them are detected for the phase extracted sample (Fig. 4(d)). After 20,000 h aging at 800 °C, the peaks of M6C carbides are not detected while those of M23C6 carbides are present for the bulk sample as shown in Fig. 4(b). The XRD result on the phase extracted sample also suggests that M6C carbides are barely present while M23C6 carbides are abundant (Fig. 4(e)). Such observation is consistent with the microstructure (Fig. 2) showing the extensive formation of the secondary M23C6 carbides on the grain boundaries and within the grains during the aging. What is more distinctive in Fig. 4(b) is the presence of very strong peaks of α-W phase which is not detected in the as-annealed Alloy 230. It should be noted that though the peaks of α-W phase are much stronger than those of M23C6 carbides, the amount of α-W phase is not as much as that of M23C6 carbides. It is clear in Fig. 3 that the amount of M23C6 carbides is much greater than that of α-W phase. The strong peaks of α-W phase could have come from the high diffraction intensity of high atomic number element like W because the atomic scattering factors increase with number of electrons. The α-W phase found in the Alloy 230 aged at 800 °C could have resulted from the breakdown or ‘dissociation’ of W-rich M6C carbides as
mentioned in Section 3.1. Further analysis was performed for verification and the results are discussed in details in Section 3.4. Meanwhile, for the sample aged at 900 °C for 20,000 h, all the peaks of matrix, M6C and M23C6 carbide, and α-W phase are present in the XRD result of the bulk sample (Fig. 4(c)). Compared to the XRD result of the 800 °C aged sample, the peak intensities of α-W phase are relatively low whereas strong peaks of M6C are present along with those of M23C6 carbide. This suggests that the existence of the peaks of α-W phase and M6C could be associated with the dissociated as well as non-dissociated primary M6C carbides respectively. 3.4. Dissociation of M6C carbides during aging at 800 °C The precipitates found in Alloy 230 in the as-annealed condition and after long-term aged at 800 °C are further characterized using TEM, EDS and SAD analyses. As there are extensive studies on the formation of the Cr-rich M23C6 carbides in Ni-base alloys during high temperature aging and the results are well documented [2,11,14–17], we focused on the phases associated with the dissociation of W-rich M6C carbides. Fig. 5(a) and (b) represents cross-sectional TEM micrographs showing a W-rich M6C carbide (primary carbide) in the as-annealed condition
Table 3 ASTM grain size number of as-annealed and aged Alloy 230.
Exposure time ASTM grain size number
As-annealed Alloy 230
Aging at 800 °C
5.5
10,000 h 5.5
Aging at 900 °C 20,000 h 5.7
10,000 h 5.6
20,000 h 5.5
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and isolated α-(W, Mo) phase near MC or Ni-rich η-phase in the γ′ forming Ni-base superalloy (3.5 wt.% W) after 10,000 h aging at 800 and 900 °C. Therefore, the formation of fairly large α-W phases via the dissociation of W-rich primary M6C carbides seems quite unique. The SAD patterns of the W-depleted regions were also analyzed. According to the SAD patterns and EDS analyses, the Cr-rich region adjacent to the outer boundary of W-rich phases (point 5 in Fig. 5(b)) is identified as M23C6 carbide (M being Cr, W and Mo) with FCC crystal structure. The Cr and Ni-rich region (point 6 in Fig. 5(b)) squeezed between W-rich phases is also identified as M23C6 carbide with FCC crystal structure, but the lattice parameter is slightly larger due to the partitioning of significant amount of Ni. Finally, the SAD pattern of the Ni-rich regions (point 7 in Fig. 5(b)) within a prior M6C carbide boundary matches that of the Ni-base matrix (point 8 in Fig. 5(b)) with lattice parameters of 3.60–3.62 Å. Also, the chemical composition of the Ni-rich region within the prior M6C carbide is very close to that of the Ni-base matrix. Considering all the analysis results mentioned above, we can conclude that the primary W-rich M6C carbides in the as-annealed Alloy 230 are internally dissociated during the long-term aging at 800 °C as represented below: W‐rich M6 C→α‐W; Cr‐rich M23 C6 ; and W‐depleted Ni‐base matrix: ð1Þ
Fig. 4. XRD peaks for bulk samples of Alloy 230 (a) in the as-annealed condition, (b) after aging at 800 °C for 20,000 h, (c) after aging at 900 °C for 20,000 h, and for phase extracted samples of Alloy 230 (d) in the as-annealed condition, and (e) after aging at 800 °C for 20,000 h.
and a dissociated carbide (area corresponding to the prior M6C carbide) after 20,000 h aging at 800 °C, respectively. In the as-annealed condition, the uniformity of the composition within the primary M6C carbides is verified with TEM/EDS mapping images in Fig. 5(c). To the contrary, for the 800 °C aged specimen, three compositionally distinct regions (or phases) are evident within the perimeter of the prior M6C carbides as shown in Fig. 5(d). First, the W-rich phases are located in the middle and on the boundary of prior M6C carbide as continuous or isolated regions. The W-rich phases (points 2, 3, and 4 in Fig. 5(b), whose point EDS results are listed in Table 4) are more than 82 at.% rich in W with small amount of Cr, Ni, and Mo. Meanwhile, other regions are depleted of W. The W-depleted regions can be further divided into two phases: 1) isolated phases which are rich in Cr with small amount of Ni (points 5 and 6 in Fig. 5(b) and phases in blue in Fig. 5(d)) formed adjacent to W-rich phases and 2) W-depleted matrix regions which are rich in Ni and Cr (points 7 and 8 in Fig. 5(b) and bluish-green regions in Fig. 5(d)) formed inside and outside of the prior M6C carbides. The chemical compositions of points shown in Fig. 5(a) and (b) were measured using point EDS analysis and the results are summarized in Table 4. The crystal structures of the phases found in Fig. 5(a) and (b) were analyzed by SAD patterns and the results are shown in Fig. 5(e) and Table 4. In the as-annealed condition, the primary M6C carbide (point 1 in Fig. 5(b)) is FCC crystal structure containing around 30 W, 25 Cr, 35 Ni, and 7 Mo (in at.%). Considering the chemical composition, it is further identified as the M3W3C type carbide (M being Ni or Mo). The SAD patterns of the W-rich phases in the perimeter of the prior M6C carbides (point 2, 3 and 4 in Fig. 5(b)) indicate that they are metallic α-W phase with BCC crystal structure. The results of TEM/EDS and SAD analyses are in good agreement of the XRD results shown in Fig. 4(b), which support the extensive presence of metallic α-W phase in the aged specimen. Such extensive formation of metallic α-W phase during high temperature exposure has been rarely reported. Previously, Qin et al. [25] observed that the formation of very fine (b1 μm)
It has been known that the primary carbides such as M6C and MC could be transformed to M23C6 and/or M6C carbides during high temperature exposure [18,26,27]. Such phase transformation phenomena could be related to the instability of the primary carbides. Li et al. calculated the ground state properties of η-M3W3C (M being Fe, Co, or Ni) carbides and indicated that those carbides have a thermodynamically stable structure [28]. Meanwhile, a high content of Cr in M6C would weaken the interatomic bonds and ultimately decrease the stability of carbides [18,29]. Therefore, the formation of the large metallic α-W phase within the prior M6C carbides could be explained in view of the stability of the M6C carbides and different diffusivities of constituent elements. That is, during the long-term exposure at 800 °C, carbide forming elements such as C, Ni, and Cr in the meta-stable primary carbides would diffuse into surrounding Ni-base matrix [29]. Since C diffuses more rapidly than other elements, it moves easily but could be trapped on the carbide/matrix interface. Then C reacts with Cr within the matrix or from the dissociated carbides, resulting in the formation of more stable carbides such as Cr-rich M23C6 on the carbide/matrix interface. As the inside of the primary carbides are depleted of C, Cr, and Ni, α-W would begin to form and coarsen, which further release Cr and Ni to nearby regions. Thus, regions of W-depleted Ni-base matrix are formed within the boundary of the prior M6C carbides. In some areas trapped between the α-W phases, such as point 6 in Fig. 5(b), Cr-rich M23C6 carbide is formed but with more partitioning of Ni which is expelled from the surrounding metallic α-W phase. Eventually, through these processes, the meta-stable W-rich M6C (in Fig. 5(a)) carbides are internally dissociated into regions of metallic α-W, Cr-rich M23C6, and W-depleted Ni-base matrix as shown in Fig. 5. 3.5. Additional phase precipitation during aging at 900 °C As mentioned in Section 3.2, in addition to the primary M6C carbides, dissociated prior M6C carbides, and secondary Cr-rich M23C6, additional phases such as Cr- and Ni-rich precipitates and W-rich precipitates are found in Alloy 230 aged at 900 °C (Fig. 3). For Alloy 230 aged at 900 °C, the characteristics of the phases formed by the dissociation of the primary M6C carbides are similar to those found in the specimens aged at 800 °C and are not repeated in this section. However, the reason for the partial dissociation of primary M6C carbides at 900 °C remains presently unclear. Further investigation is needed including thermodynamic analysis. Meanwhile, additional phases such as Cr- and Ni-rich
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Fig. 5. TEM micrographs of (a) the primary M6C in the as-annealed condition and (b) internally dissociated prior M6C carbide after 20,000 h aging at 800 °C, TEM/EDS mapping micrographs of (c) the primary M6C in the as-annealed condition and (d) internally dissociated prior M6C carbide after aging at 800 °C for 20,000 h, and (e) SAD patterns of points in (a) and (b).
precipitates and W-rich precipitates found in the specimens aged at 900 °C are characterized. The precipitation of the Cr- and Ni-rich precipitates has been previously reported [11,16], but its characteristics are still unclear. In this study, the attempt to identify the Cr- and Ni-rich precipitates by XRD analysis was not successful as its volume fraction was not enough to show peaks in the XRD result of bulk sample (Fig. 4(c)). As an additional effort to identify the phase, the selected area diffraction (SAD) pattern of the Cr- and Ni-rich precipitates and EDS analyses were performed and the results are shown in Fig. 6. Previously, D. Kim [11] reported these precipitates as an intermetallic phase of Cr0.8Ni0.2 with BCC structure based on the selected area diffraction (SAD) result
for the Alloy 230 aged at 1000 °C for 500 h. However, the SAD pattern shown in Fig. 6 is not consistent with the previous result, even though the chemical compositions are very similar in both analyses. Furthermore, we could not find phases that would match to the SAD pattern shown in Fig. 6. Similarly, J. Veverkova [16] also attempted SAD analysis to identify the Cr- and Ni-rich precipitates, but failed to find any matching phases. On the other hand, K. Kuo [30] and K. Kikuchi [31] observed the existence of tetragonal structured σ-Cr8Ni5W at the grain boundary in Ni–Cr–W system by XRD analysis. Nonetheless, it is clear that the Cr- and Ni-rich precipitates are transformed phases from Crrich M23C6 grain boundary carbides because they occupy all the grain
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Table 4 The results of point EDS analysis (in at.%) in Fig. 5(a) and (b). Point in Fig. 5 C W Cr Ni Mo Possible phase Crystal structure and space group Lattice parameter (Å)
+1 5.4 28.5 26.2 33.3 6.5 η-M6C FCC, Fd3m 11.38
+2 – 84.1 3.8 3.7 8.3 α-W
+3 – 82.6 7.0 3.9 6.5
+4 – 84.8 4.7 2.0 8.5
BCC, Im3m
+5 – 4.8 85.5 4.4 5.2 M23C6 FCC,
+6 7.7 6.7 55.1 24.7 4.5 M23C6 FCC,
+7 +8 – – 4.4 4.2 24.6 23.5 70.5 71.3 0.5 1.0 Matrix FCC,
3.19
Fm3m 10.87
Fm3m 10.97
Fm3m 3.60–3.62
boundaries after short exposure period [2,14,16]. Further researches are required to identify the precipitate. W-rich precipitates on the grain boundaries formed after long-term aging at 900 °C was characterized using TEM, EDS and SAD analyses. Fig. 7 represents a cross-sectional TEM micrograph of the W-rich precipitate on the grain boundary within the carbide depleted region in Fig. 3(e). The point EDS result shows that it mainly contains around 83 W and 10 Mo (in at.%). This chemical composition is quite similar to that of the α-W phase formed in the dissociated prior M6C carbides shown in Table 4. Also, the SAD pattern inserted in Fig. 7 matches to that of the metallic α-W phase with BCC crystal structure (lattice parameter of 3.16 Å). The formation of the α-W phase on the grain boundaries might be associated with the carbide depletion near the Cr- and Ni-rich precipitates. That is, during the shorter exposure time at 900 °C, Cr-rich M23C6 carbides are formed on the grain boundaries rather than W-rich phases due to the slower diffusion of W compared to Cr [2,15]. However, as further aging progresses, the abovementioned Cr- and Ni-rich precipitates and Crdepleted region are formed in some areas at the expense of Cr-rich M23C6 carbides, though the mechanisms are not well understood [16]. Then, W within the matrix would diffuse to the grain boundaries in the Cr-depleted region and precipitate as the α-W phase during long-term exposure. As the diffusion of W is required, precipitation of the α-W phase would be much faster at higher temperature, which is supported by the precipitation of discontinuous α-W on the grain boundaries for Ni-base alloys with 18 wt.% W exposed for 1 h at 1200 °C [32].
Fig. 6. TEM micrograph, SAD pattern, and chemical compositions of Cr- and Ni-rich precipitates formed after aging at 900 °C for 10,000 h.
Fig. 7. TEM micrograph, SAD pattern, and chemical compositions of W-rich precipitate formed on the grain boundary after aging at 900 °C for 20,000 h.
3.6. Changes in tensile properties after aging at 800 and 900 °C The effects of microstructure evolution during high temperature aging on the tensile properties were assessed at room temperature. Fig. 8 shows the stress–strain curves of tensile tests for as-annealed and aged specimens. In addition, the results of tensile tests are summarized in Table 5. As shown in the figure, the long-term aging (10,000 and 20,000 h) at 800 °C results in the increase in ultimate tensile strength (UTS) and decrease in elongation for Alloy 230. Meanwhile, the longterm aging at 900 °C results in larger decrease in elongation, but decrease in UTS is very small. Fig. 9 shows the cross-sectional microstructures of tensile fractured specimens. In the as-annealed condition, fracture mode is the mixture of the trans- and inter-granular (Fig. 9(a)). However, as Cr-rich M23C6 grain boundary carbides are extensively formed and thickened, inter-granular fracture mode is dominant in the aged specimens (Fig. 9(b) and (c)). Meanwhile, the Cr-rich M23C6 carbides formed within the grain would increase strength instead of lowering the elongation, which is consistent with the tensile test results such that specimen aged at 800 °C (with more Cr-rich M23C6 carbides within the grain) shows higher UTS value than that aged at 900 °C (with less Cr-rich M23C6 carbides within the grain).
Fig. 8. The room temperature tensile stress–strain curves of Alloy 230 after aging at 800 and 900 °C for up to 20,000 h.
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Table 5 The results of room temperature tensile test of as-annealed and aged Alloy 230.
Exposure time Ultimate tensile strength (MPa) Elongation (%)
As-annealed Alloy 230
Aging at 800 °C
843 ± 12 60 ± 1
10,000 h 938 ± 2 45 ± 1
It is also interesting to mention that some primary M6C carbides near the fractured area are cracked in the as-received condition (Fig. 9(a)), although the dissociated and non-dissociated primary M6C carbides are mostly remained un-cracked for the aged specimens (Fig. 9(b) and (c)). Previously, K. Mo [23] reported the fracture of the primary M6C carbide in the as-annealed Alloy 230 was resulted from significant dislocation movement at the transition from elastic to plastic regime when tensile loading was applied. He explained that a great number of dislocations were piled at the large and non-deformable primary M6C carbide. It could not endure large internal stress and begin to rupture upon early yielding. However, in case of the aged specimens in this study, a presence of extensive precipitation of the secondary M23C6 carbides on the grain boundaries and within the grains would relieve internal stress around the primary M6C carbides by blocking dislocation movement. This explains the absence of any cracked primary M6C carbides in the specimen aged at 800 °C for 20,000 h (Fig. 9(b)). On the other hand, for the specimen aged at 900 °C for 20,000 h, some large Cr- and Ni-rich precipitates are cracked as shown in Fig. 9(d). Local cracking of Cr- and Ni-rich precipitates was reported during tensile or creep tests and explained with the brittle nature of the phase [11,16]. Therefore, the cracking of Cr- and Ni-rich precipitates as well as grain boundary carbides contributed to the premature failure of the specimens aged at 900 °C, which resulted in the lower UTS values than the specimens aged at 800 °C as shown in Fig. 8. For both aging temperatures, a degradation of tensile property is observed after 10,000 and 20,000 h exposure, but with different degrees. After aging at 800 °C, the typical tensile property change, such as increase in strength and decrease in elongation, was observed. Such change could be well explained with the microstructure evolution, especially, the precipitation of the secondary Cr-rich M23C6 carbides. It seems that the dissociation of the primary M6C carbide has minimal effects on tensile properties. After aging at 900 °C, formation of Cr-
Aging at 900 °C 20,000 h 935 ± 6 44 ± 1
10,000 h 824 ± 7 38 ± 2
20,000 h 823 ± 19 37 ± 4
and Ni-rich precipitates affects the tensile properties. That is, formation of brittle Cr- and Ni-rich precipitates is associated with the carbide free region and grain boundary α-W phases, which results in lowering elongation as well as strength as observed in tensile tests. Therefore, for the aged specimens, the degradation of tensile property would depend on the evolution of Cr-rich M23C6 carbides, Cr- and Ni-rich precipitates, and α-W phases but not on the primary M6C carbides.
4. Conclusions The effects of long-term exposure to high temperature (800 and 900 °C for up to 20,000 h) on the microstructure evolution of a Ni–Cr– W superalloy (Alloy 230) were investigated. Based on the analysis and tests, the following conclusions are drawn. 1. After long-term aging at 800 °C, extensive precipitation of the secondary Cr-rich M23C6 carbides was observed on the grain boundaries as well as within the grains. In addition, W-rich primary η-M6C carbides were internally dissociated into three distinctive phases, such that W-rich M6C → α-W, Cr-rich M23C6, and W-depleted Nibase matrix. 2. After long-term aging at 900 °C, precipitation of Cr- and Ni-rich intermetallic and α-W phase was observed in some areas at the expense of the secondary Cr-rich M23C6 carbides. Meanwhile, the dissociation of the primary M6C carbides was less significant. 3. Long-term aging at both temperature resulted in decrease in tensile elongation of Alloy 230 while increase in strength was minimal after aging at 900 °C. The degradation of tensile property depends on the evolution of Cr-rich M23C6 carbides, Cr- and Nirich precipitates, and α-W phases but not on the primary M6 C carbides.
Fig. 9. Cross-sectional microstructure of room temperature tensile tested Alloy 230; (a) as-annealed and after 20,000 h aging at (b) 800 and (c), and (d) 900 °C.
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Acknowledgments This study was supported by the MSIP/NRF (Nuclear R&D Program 2013M2A8A6035683) of Republic of Korea. References [1] J.K. Wright, Next Generation Nuclear Plant Intermediate Heat Exchanger Materials Research and Development Plan. INL/EXT-08-14107, Idaho National Laboratory, 2008. [2] H.M. Tawancy, D.L. Klarstrom, M.F. Rothman, Development of a new nickel-base superalloy, J. Met. 36 (1984) 58–62. [3] W. Ren, R. Swindeman, A review on current status of alloys 617 and 230 for Gen IV nuclear reactor internals and heat exchangers, J. Press. Vessel. Technol. 131 (2009) (044002-1-15). [4] C. Jang, D. Kim, D. Kim, I. Sah, W.S. Ryu, Y.S. Yoo, Oxidation behaviors of wrought nickel-base superalloys in various high temperature environments, Trans. Nonferrous Met. Soc. China 21 (2011) 1524–1531. [5] L. Zheng, C.Q. Gu, Y.R. Zheng, Investigation of the solidification behavior of a new Rucontaining cast Ni-base superalloy with high W content, Scr. Mater. 50 (2004) 435–439. [6] H.M. Tawancy, High temperature creep behaviour of an Ni–Cr–W–B alloy, J. Mater. Sci. 27 (1992) 6481–6489. [7] K.S. Vecchio, M.D. Fitzpatrick, D. Klarstrom, Influence of subsolvus thermomechanical properties of HAYNES 230 alloy, Metall. Mater. Trans. A 26 (1995) 673–689. [8] J.K. Benz, L.J. Carroll, J.K. Wright, R.N. Wright, T.M. Lillo, Threshold stress creep behavior of Alloy 617 at intermediate temperatures, Metall. Mater. Trans. A 45 (2014) 3010–3022. [9] W.L. Mankins, J.C. Hosier, T.H. Bassford, Microstructure and phase stability of INCONEL Alloy 617, Metall. Trans. 5 (1974) 2579–2589. [10] S. Kihara, J.B. Newkirk, A. Ohtomo, Y. Saiga, Morphological changes of carbides during creep and their effects on the creep properties of Inconel 617 at 1000 °C, Metall. Trans. A 11 (1980) 1019–1031. [11] D. Kim, I. Sah, C. Jang, Effects of high temperature aging in an impure helium environment on low temperature embrittlement of Alloy 617 and Haynes 230, J. Nucl. Mater. 405 (2010) 9–16. [12] L.M. Suave, J. Cormier, P. Villechaise, A. Soula, Z. Hervier, D. Bertheau, J. Laigo, Microstructure evolutions during thermal aging of Alloy 625: impact of temperature and forming process, Metall. Mater. Trans. A 45 (2014) 2963–2977. [13] L.M. Suave, D. Bertheau, J. Cormier, P. Villechaise, A. Soula, Z. Hervier, J. Laigo, Impact of microstructural evolutions during thermal aging of Alloy 625 on its monotonic mechanical properties, MATEC Web of Conference, 2014 (14:21001-1-6). [14] D. Kim, I. Sah, C. Jang, Effects of aging in high temperature helium environments on room temperature tensile properties of nick-base superalloys, Mater. Sci. Eng. A 528 (2011) 1713–1720.
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Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Microstructure evolution of a Ni–Cr–W superalloy during long-term aging at high temperatures Ho Jung Lee a, Hyunmyung Kim a, Donghoon Kim b, Changheui Jang a,⁎ a b
Department of Nuclear and Quantum Engineering, Korea Advanced Institute of Science and Technology, 305-701 Daejeon, Republic of Korea Agency for Defense Department, 305-600 Daejeon, Republic of Korea
a r t i c l e
i n f o
Article history: Received 29 April 2015 Received in revised form 19 June 2015 Accepted 19 June 2015 Available online 20 June 2015 Keywords: Ni–Cr–W superalloy Long-term aging Microstructure evolution Carbide dissociation Tensile elongation
a b s t r a c t The effects of long-term exposure to high temperature (800 and 900 °C for up to 20,000 h) on the microstructure evolution of a Ni–Cr–W superalloy (Alloy 230) were investigated. After long-term aging at 800 °C, extensive precipitation of the secondary Cr-rich M23C6 carbides was observed. In addition, the internal dissociation of W-rich primary η-M6C carbides was observed, such as, W-rich M6C → α-W, Cr-rich M23C6, and W-depleted Ni-base matrix. After long-term aging at 900 °C, precipitation of Cr- and Ni-rich phase and α-W phase was observed in some areas at the expense of the secondary Cr-rich M23C6 carbides. Meanwhile, the dissociation of the primary M6C carbides was less significant at 900 °C. Long-term aging at both temperatures resulted in a decrease in tensile elongation while an increase in strength was minimal after aging at 900 °C. The degradation of tensile property depended on the evolution of Cr-rich M23C6 carbides, Cr- and Ni-rich precipitates, and α-W phases but not on the primary M6C carbides. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Among generation-IV (Gen-IV) nuclear reactors under development, a very high temperature gas-cooled reactor (VHTR) will be operated at the temperature range of 800–900 °C for up to 60 years [1]. For the application of an intermediate heat exchanger (IHX) in a VHTR, a solid solution hardened Ni–Cr–W superalloy, Alloy 230, is considered one of the candidate materials because of the superior mechanical properties and oxidation resistance at high temperature [2–4]. It has been reported that the W-rich primary M6C carbides are formed in Alloy 230 during solidification process [2,5]. They contribute to additional increase in high temperature strength and creep resistance by pinning grain boundaries and thereby inhibiting grain boundary sliding [6,7]. As mechanical properties are affected by various microstructural features, evaluation of the microstructural changes associated with long-term exposure at high temperature is important. Previously, the effects of microstructure on mechanical properties were investigated for Ni-base superalloys with similar Cr content to Alloy 230 [8–13]. That is, various types of carbides (M23C6 and M6C) and precipitates (γ′, CrMo(C,N), and TiN for Alloy 617 and γ″, δ, and Ni2(Cr,Mo) for Alloy 625) were formed during high temperature exposure depending on exposure temperatures and chemical compositions. The presence of those precipitates resulted in hardening and reduction of ductility and creep resistance by disturbing dislocation movement.
⁎ Corresponding author. E-mail address: [email protected] (C. Jang).
http://dx.doi.org/10.1016/j.matchar.2015.06.018 1044-5803/© 2015 Elsevier Inc. All rights reserved.
However, in case of alloy 230, there are limited studies on the effects of the long-term aging on the mechanical properties of Alloy 230 [5,11, 14–17]. Previously, formation of the Cr-rich M23C6 carbides at the boundary of the primary M6C carbides was also reported when exposed to the temperature ranges of 650–1000 °C [16,18]. Moreover, Veverkova [16] and Wu [19] reported the breakdown of the primary carbides after long-term exposure at the temperature ranges of 650–870 °C. The issue in other Ni-base superalloys has also been reported recently to assure long-term stability for high temperature applications [20,21]. However, the mechanism and nature of the phases formed in Alloy 230 were not properly characterized in the previous studies. Therefore, in this study, the long-term aging behaviors of Alloy 230 at the anticipated operating temperatures of a VHTR were assessed focusing on the microstructure evolution such as the formation and dissociation of carbides. Then the consequent changes in the tensile properties associated with such microstructure evolution were investigated. 2. Materials and experimental The chemical composition of a commercial grade Alloy 230 used in the study was analyzed by inductively coupled plasma mass spectroscopy (ICP-MS) method and the result is listed in Table 1. The test material was provided as solution annealed at 1177 °C for 0.5 h followed by water quenched. For the aging process, blocks of Alloy 230 were exposed at 800 and 900 °C for 10,000 and 20,000 h in a box furnace. After aging, the blocks were removed from the furnaces and aircooled to room temperature. The specimens for the microstructure observation and tensile test were taken at 2 mm from the surface of the
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Table 1 Chemical composition of Alloy 230 used in the study (wt.%). Ni
Cr
W
Mo
Fe
Mn
Si
Al
Co
C
La
Bal.
22.3
13.9
1.37
1.34
0.5
0.28
0.25
0.15
0.1
0.012
block to avoid environmentally induced degradations such as internal oxidation and carburization/decarburization. The microstructure features including the carbides and other phases were observed using a scanning electron microscope (SEM, Magellan 400) with back scattered electron (BSE) mode at 10 kV. In SEM sample preparation, specimens were mechanically polished down to 1 μm using SiC papers and diamond polishing compound and ultra-sonically cleaned in ethanol. The polished samples were etched in solutions of hydrogen chloride and 2 vol.% bromine in methanol in sequence to reveal grain boundaries. The volume fractions of primary and secondary carbides were measured using an image analyzer (TOMORO analySIS TS) as follow: 1) Low magnification SEM and BSE images acquired by a dual mode of SEM. 2) Total volume fractions of primary and secondary carbides were measured with SEM images because the contrast of those phases was different from the matrix. 3) Volume fraction of the primary carbides was measured using BSE images because primary carbides with high atomic number are easily distinguished from secondary carbides and matrix. 4) Volume fraction of the secondary carbides was estimated by excluding volume fraction of primary carbides from total carbide volume fraction measured in step 2. However, the image analyzer could not accurately distinguish phases with very thin shape or of similar contrast, which will be discussed in Section 3.2. Also, the grain size of Alloy 230 before and after the aging was measured with circle intercept method following the procedure of ASTM E112 using the image analyzer. In addition, an X-ray diffractometer (XRD, RIGAKU D/MAX-2500) and a transmission electron microscopy (TEM, Titan cubed G2 60-300) equipped with energy dispersive spectroscopy (EDS) and selective area diffraction (SAD) at 300 kV were used to identify the phases before and after the aging. The TEM samples were prepared by a focused ion beam (FIB, Helios Nanolab 450 F1) with area of 10 × 10 μm and around 50 nm in thickness. For the tensile tests, mini-sized tensile specimens with 16 mm in length and 0.5 mm in thickness were machined. The tests were conducted at room temperature under displacement control (strain rate of 3.33 × 10− 4/s) following the procedures of ASTM E8/ E8M-13a. Due to the mini-size of the specimen, the extensometer was not attached. Instead, the strain was estimated from the displacement divided by the gage length of mini-size specimen. In the XRD analysis, in addition to the bulk samples, phase extracted samples were used in order to clearly identify the peaks from the carbides without the interference of those from matrix. In the phase extraction process, samples were mechanical polished to eliminate impurities on the surface and then ultrasonically cleaned in ethanol. The samples were soaked into the solution of 10 vol.% bromine in ethanol until all metallic phases were dissolved. Then the solution containing
the extracted particles was filtered through the porous filter paper to collect the particles selectively. Samples were analyzed with a high applied power of 40 kV and 300 mA. The incident angle of the X-ray beam was 4° in the 2 theta mode within the range of 20 to 80°.
3. Results and discussion 3.1. Microstructure evolutions during aging at 800 °C Fig. 1 shows the microstructure of an as-annealed Alloy 230. Grain boundaries are covered with fine discrete carbides which are known as the secondary Cr-rich M23C6 formed during the cooling process from solution annealing [7,11,22]. In addition, heavy and globular shaped primary carbides (bright color in BSE image in Fig. 1(b)) are randomly distributed on the grain boundaries and within the grains with various sizes (1 to 20 μm). In previous studies, these carbides were reported as the primary W-rich M6C type carbide in the form of η-Ni3W3C through TEM/SAD [2,11,18], XRD [16,22] and synchrotron radiation [23] analyses. The volume fractions of the carbides are measured by an image analyzer and listed in Table 2. As summarized in Table 2, volume fraction of W-rich M6C type carbide is about 4%, while the amount of Cr-rich M23C6 grain boundary carbides too thin to be quantitatively measured. After 10,000 and 20,000 h aging at 800 °C, significant microstructure changes are observed as compared to the as-annealed condition as shown in Fig. 2. For the specimens aged at 800 °C, the extensive precipitation of the secondary Cr-rich M23C6 carbides (gray color in BSE image in Fig. 2(b) and (d)) is observed on the grain boundaries as well as within the grains. Also, the amount of the secondary carbides increases with aging time. Such behaviors are in good agreement with the previous reports on the formation of the secondary Cr-rich M23C6 carbides in Ni-base superalloys exposed at various aging temperature ranges [10, 15,24]. On the other hand, the volume fraction of the primary M6C carbides is not significantly altered even after 20,000 h aging at 800 °C. Instead, areas of different contrasts are observed inside of the primary M6C carbides boundary for the aged specimens as shown in SEM/BSE micrographs (Fig. 2(b) and (d)), indicating the possibility of ‘internal dissociation’ during high temperature aging. It is similar to the so-called ‘breakdown’ phenomenon previously reported by Veverkova [16] and Wu [19]. They observed the breakdown of the primary M6C carbides in Alloy 230 after over 10,000 h of isothermal aging in the temperature ranges of 700–870 °C. Wu suggested the transformed phase as W2C carbide based on the EDS analysis and SEM observation [19]. On the other hand, the crystal structure and chemical composition of the primary M6C carbides were reported the same as those found in the as-annealed condition after a short-term exposure up to 3000 h in the temperature range of 650–1000 °C [2,11,16]. Such observations suggest that the transformation phenomenon is a slow process, possibly associated with the diffusion of slow-moving element. The phase transformation within the
Fig. 1. SEM micrographs of the as-annealed Alloy 230; (a) low and (b) high magnifications (SEM/BSE images).
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Table 2 Volume fractions of various phases present in Figs. 1–3. As-annealed Alloy 230
Prior to aging 10,000 h aging 20,000 h aging a b c d
Aging at 900 °Cd
Aging at 800 °C
W-rich M6C
Cr-rich M23C6
~4%
a
W-rich M6Cb
Cr-rich M23C6
W-rich M6C & α-Wc
Cr-rich M23C6
4–5% 4–5%
7–10% 13–14%
3–4% 3–5%
5–6% 5–8%
Measurement could not be made due to very low volume fraction. Volume fraction of the phases within the prior M6C carbides. Volume fraction of un-dissociated and dissociated M6C carbides plus α-W on the grain boundary. Volume fraction of Cr- and Ni-rich precipitates was not measured as it was not distinguishable from the matrix.
primary M6C carbides, or internal dissociation, will be further analyzed and discussed in the later sections.
3.2. Microstructure evolutions during aging at 900 °C The microstructures of Alloy 230 after 10,000 and 20,000 h aging at 900 °C are shown in Fig. 3. Unlike those after aging at 800 °C, the precipitation of the secondary phases is rather irregular after aging at 900 °C. That is, along with the areas with extensive precipitation of the secondary M23C6 carbides (for example, both side of Fig. 3(b)), there are areas with very little secondary M23C6 carbides (carbide depleted region, middle of Fig. 3(b)). In the areas with very little secondary carbides, a number of large and elongated phases are developed along the grain boundaries, which are identified as Cr- and Ni-rich (52 Cr and 38 Ni in at.%) and W-rich (83 at.% W) precipitates by the point TEM/EDS analysis (details will be provided in Sections 3.4 and 3.5). The Cr- and Ni-rich phases are indicated by arrows in the figures as the contrast is weak in the SEM photos, probably due to the fact that the number of electron per atom is similar in both Cr- and Ni-rich phases and matrix. Meanwhile, W-rich precipitates are bright and easily distinguishable from Cr- and Ni-rich precipitates. Precipitation of Cr- and Ni-rich phases in Alloy 230 was previously reported and described as the ‘pool-like precipitates’ [16] or ‘intermetallic phase of Cr0.8Ni0.2’ [11] after long-term exposure at above 950 °C including 500 h exposure at 1000 °C. In addition, the dissociation of some primary M6C carbides is observed mostly in the carbide depleted region while most of them remain similar morphology to the as-annealed condition (Fig. 3(b) and (c)).
It is interesting that, compared to the specimens aged at 800 °C, the volume fraction of the secondary M23C6 carbides are much smaller in the specimens aged at 900 °C, which is expected from the existence of carbide free zone in Fig. 3. Instead, both Cr- and Ni-rich precipitates and W-rich precipitates are present in large quantity, though they were not quantitatively analyzed because the contrast of Cr- and Nirich precipitates and W-rich precipitates is similar to that of the matrix and primary M6C carbides respectively. Meanwhile, as aging time increases from 10,000 h to 20,000 h, the amount of the secondary M23C6 carbides increases but only by a small degree (Fig. 3(a), (b), and Table 2). Also the dissociation of the primary M6C carbides further progresses with aging time, while the total amount of dissociated and nondissociated M6C carbides remains almost the same (Table 2). Table 3 shows the measured grain size of as-annealed and aged Alloy 230 at 800 and 900 °C up to 20,000 h. The grain sizes are almost the same irrespective of the aging conditions and the ASTM grain size number is around 5.5. Whittenberger [17] also reported that the grain size of Alloy 230 was not changed after aging at 820 °C up to 22,500 h. It is thought that the precipitation of thick and continuous Cr-rich M23C6 grain boundary carbides act as efficient obstacles against grain growth. 3.3. XRD analysis for phase identification XRD analyses were performed to identify the phases in Alloy 230 before and after the long-term aging and the results are shown in Fig. 4. In the figure, the XRD results of both bulk and phase-extracted samples are presented. In the as-annealed Alloy 230 (Fig. 4(a)), high intensity peaks of matrix and M6C carbides are detected indicating that the M6C
Fig. 2. SEM micrographs of Alloy 230 after (a), (b) 10,000 h and (c), (d) 20,000 h aging at 800 °C (SEM/BSE images).
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Fig. 3. SEM micrographs of Alloy 230 after (a), (b), (c) 10,000 h and (d), (e), (f) 20,000 h aging at 900 °C (SEM/BSE images).
carbides are dominant precipitates. Along with M6C carbides, the existence of small amount of M23C6 carbides is confirmed as the shallow peaks of them are detected for the phase extracted sample (Fig. 4(d)). After 20,000 h aging at 800 °C, the peaks of M6C carbides are not detected while those of M23C6 carbides are present for the bulk sample as shown in Fig. 4(b). The XRD result on the phase extracted sample also suggests that M6C carbides are barely present while M23C6 carbides are abundant (Fig. 4(e)). Such observation is consistent with the microstructure (Fig. 2) showing the extensive formation of the secondary M23C6 carbides on the grain boundaries and within the grains during the aging. What is more distinctive in Fig. 4(b) is the presence of very strong peaks of α-W phase which is not detected in the as-annealed Alloy 230. It should be noted that though the peaks of α-W phase are much stronger than those of M23C6 carbides, the amount of α-W phase is not as much as that of M23C6 carbides. It is clear in Fig. 3 that the amount of M23C6 carbides is much greater than that of α-W phase. The strong peaks of α-W phase could have come from the high diffraction intensity of high atomic number element like W because the atomic scattering factors increase with number of electrons. The α-W phase found in the Alloy 230 aged at 800 °C could have resulted from the breakdown or ‘dissociation’ of W-rich M6C carbides as
mentioned in Section 3.1. Further analysis was performed for verification and the results are discussed in details in Section 3.4. Meanwhile, for the sample aged at 900 °C for 20,000 h, all the peaks of matrix, M6C and M23C6 carbide, and α-W phase are present in the XRD result of the bulk sample (Fig. 4(c)). Compared to the XRD result of the 800 °C aged sample, the peak intensities of α-W phase are relatively low whereas strong peaks of M6C are present along with those of M23C6 carbide. This suggests that the existence of the peaks of α-W phase and M6C could be associated with the dissociated as well as non-dissociated primary M6C carbides respectively. 3.4. Dissociation of M6C carbides during aging at 800 °C The precipitates found in Alloy 230 in the as-annealed condition and after long-term aged at 800 °C are further characterized using TEM, EDS and SAD analyses. As there are extensive studies on the formation of the Cr-rich M23C6 carbides in Ni-base alloys during high temperature aging and the results are well documented [2,11,14–17], we focused on the phases associated with the dissociation of W-rich M6C carbides. Fig. 5(a) and (b) represents cross-sectional TEM micrographs showing a W-rich M6C carbide (primary carbide) in the as-annealed condition
Table 3 ASTM grain size number of as-annealed and aged Alloy 230.
Exposure time ASTM grain size number
As-annealed Alloy 230
Aging at 800 °C
5.5
10,000 h 5.5
Aging at 900 °C 20,000 h 5.7
10,000 h 5.6
20,000 h 5.5
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and isolated α-(W, Mo) phase near MC or Ni-rich η-phase in the γ′ forming Ni-base superalloy (3.5 wt.% W) after 10,000 h aging at 800 and 900 °C. Therefore, the formation of fairly large α-W phases via the dissociation of W-rich primary M6C carbides seems quite unique. The SAD patterns of the W-depleted regions were also analyzed. According to the SAD patterns and EDS analyses, the Cr-rich region adjacent to the outer boundary of W-rich phases (point 5 in Fig. 5(b)) is identified as M23C6 carbide (M being Cr, W and Mo) with FCC crystal structure. The Cr and Ni-rich region (point 6 in Fig. 5(b)) squeezed between W-rich phases is also identified as M23C6 carbide with FCC crystal structure, but the lattice parameter is slightly larger due to the partitioning of significant amount of Ni. Finally, the SAD pattern of the Ni-rich regions (point 7 in Fig. 5(b)) within a prior M6C carbide boundary matches that of the Ni-base matrix (point 8 in Fig. 5(b)) with lattice parameters of 3.60–3.62 Å. Also, the chemical composition of the Ni-rich region within the prior M6C carbide is very close to that of the Ni-base matrix. Considering all the analysis results mentioned above, we can conclude that the primary W-rich M6C carbides in the as-annealed Alloy 230 are internally dissociated during the long-term aging at 800 °C as represented below: W‐rich M6 C→α‐W; Cr‐rich M23 C6 ; and W‐depleted Ni‐base matrix: ð1Þ
Fig. 4. XRD peaks for bulk samples of Alloy 230 (a) in the as-annealed condition, (b) after aging at 800 °C for 20,000 h, (c) after aging at 900 °C for 20,000 h, and for phase extracted samples of Alloy 230 (d) in the as-annealed condition, and (e) after aging at 800 °C for 20,000 h.
and a dissociated carbide (area corresponding to the prior M6C carbide) after 20,000 h aging at 800 °C, respectively. In the as-annealed condition, the uniformity of the composition within the primary M6C carbides is verified with TEM/EDS mapping images in Fig. 5(c). To the contrary, for the 800 °C aged specimen, three compositionally distinct regions (or phases) are evident within the perimeter of the prior M6C carbides as shown in Fig. 5(d). First, the W-rich phases are located in the middle and on the boundary of prior M6C carbide as continuous or isolated regions. The W-rich phases (points 2, 3, and 4 in Fig. 5(b), whose point EDS results are listed in Table 4) are more than 82 at.% rich in W with small amount of Cr, Ni, and Mo. Meanwhile, other regions are depleted of W. The W-depleted regions can be further divided into two phases: 1) isolated phases which are rich in Cr with small amount of Ni (points 5 and 6 in Fig. 5(b) and phases in blue in Fig. 5(d)) formed adjacent to W-rich phases and 2) W-depleted matrix regions which are rich in Ni and Cr (points 7 and 8 in Fig. 5(b) and bluish-green regions in Fig. 5(d)) formed inside and outside of the prior M6C carbides. The chemical compositions of points shown in Fig. 5(a) and (b) were measured using point EDS analysis and the results are summarized in Table 4. The crystal structures of the phases found in Fig. 5(a) and (b) were analyzed by SAD patterns and the results are shown in Fig. 5(e) and Table 4. In the as-annealed condition, the primary M6C carbide (point 1 in Fig. 5(b)) is FCC crystal structure containing around 30 W, 25 Cr, 35 Ni, and 7 Mo (in at.%). Considering the chemical composition, it is further identified as the M3W3C type carbide (M being Ni or Mo). The SAD patterns of the W-rich phases in the perimeter of the prior M6C carbides (point 2, 3 and 4 in Fig. 5(b)) indicate that they are metallic α-W phase with BCC crystal structure. The results of TEM/EDS and SAD analyses are in good agreement of the XRD results shown in Fig. 4(b), which support the extensive presence of metallic α-W phase in the aged specimen. Such extensive formation of metallic α-W phase during high temperature exposure has been rarely reported. Previously, Qin et al. [25] observed that the formation of very fine (b1 μm)
It has been known that the primary carbides such as M6C and MC could be transformed to M23C6 and/or M6C carbides during high temperature exposure [18,26,27]. Such phase transformation phenomena could be related to the instability of the primary carbides. Li et al. calculated the ground state properties of η-M3W3C (M being Fe, Co, or Ni) carbides and indicated that those carbides have a thermodynamically stable structure [28]. Meanwhile, a high content of Cr in M6C would weaken the interatomic bonds and ultimately decrease the stability of carbides [18,29]. Therefore, the formation of the large metallic α-W phase within the prior M6C carbides could be explained in view of the stability of the M6C carbides and different diffusivities of constituent elements. That is, during the long-term exposure at 800 °C, carbide forming elements such as C, Ni, and Cr in the meta-stable primary carbides would diffuse into surrounding Ni-base matrix [29]. Since C diffuses more rapidly than other elements, it moves easily but could be trapped on the carbide/matrix interface. Then C reacts with Cr within the matrix or from the dissociated carbides, resulting in the formation of more stable carbides such as Cr-rich M23C6 on the carbide/matrix interface. As the inside of the primary carbides are depleted of C, Cr, and Ni, α-W would begin to form and coarsen, which further release Cr and Ni to nearby regions. Thus, regions of W-depleted Ni-base matrix are formed within the boundary of the prior M6C carbides. In some areas trapped between the α-W phases, such as point 6 in Fig. 5(b), Cr-rich M23C6 carbide is formed but with more partitioning of Ni which is expelled from the surrounding metallic α-W phase. Eventually, through these processes, the meta-stable W-rich M6C (in Fig. 5(a)) carbides are internally dissociated into regions of metallic α-W, Cr-rich M23C6, and W-depleted Ni-base matrix as shown in Fig. 5. 3.5. Additional phase precipitation during aging at 900 °C As mentioned in Section 3.2, in addition to the primary M6C carbides, dissociated prior M6C carbides, and secondary Cr-rich M23C6, additional phases such as Cr- and Ni-rich precipitates and W-rich precipitates are found in Alloy 230 aged at 900 °C (Fig. 3). For Alloy 230 aged at 900 °C, the characteristics of the phases formed by the dissociation of the primary M6C carbides are similar to those found in the specimens aged at 800 °C and are not repeated in this section. However, the reason for the partial dissociation of primary M6C carbides at 900 °C remains presently unclear. Further investigation is needed including thermodynamic analysis. Meanwhile, additional phases such as Cr- and Ni-rich
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Fig. 5. TEM micrographs of (a) the primary M6C in the as-annealed condition and (b) internally dissociated prior M6C carbide after 20,000 h aging at 800 °C, TEM/EDS mapping micrographs of (c) the primary M6C in the as-annealed condition and (d) internally dissociated prior M6C carbide after aging at 800 °C for 20,000 h, and (e) SAD patterns of points in (a) and (b).
precipitates and W-rich precipitates found in the specimens aged at 900 °C are characterized. The precipitation of the Cr- and Ni-rich precipitates has been previously reported [11,16], but its characteristics are still unclear. In this study, the attempt to identify the Cr- and Ni-rich precipitates by XRD analysis was not successful as its volume fraction was not enough to show peaks in the XRD result of bulk sample (Fig. 4(c)). As an additional effort to identify the phase, the selected area diffraction (SAD) pattern of the Cr- and Ni-rich precipitates and EDS analyses were performed and the results are shown in Fig. 6. Previously, D. Kim [11] reported these precipitates as an intermetallic phase of Cr0.8Ni0.2 with BCC structure based on the selected area diffraction (SAD) result
for the Alloy 230 aged at 1000 °C for 500 h. However, the SAD pattern shown in Fig. 6 is not consistent with the previous result, even though the chemical compositions are very similar in both analyses. Furthermore, we could not find phases that would match to the SAD pattern shown in Fig. 6. Similarly, J. Veverkova [16] also attempted SAD analysis to identify the Cr- and Ni-rich precipitates, but failed to find any matching phases. On the other hand, K. Kuo [30] and K. Kikuchi [31] observed the existence of tetragonal structured σ-Cr8Ni5W at the grain boundary in Ni–Cr–W system by XRD analysis. Nonetheless, it is clear that the Cr- and Ni-rich precipitates are transformed phases from Crrich M23C6 grain boundary carbides because they occupy all the grain
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Table 4 The results of point EDS analysis (in at.%) in Fig. 5(a) and (b). Point in Fig. 5 C W Cr Ni Mo Possible phase Crystal structure and space group Lattice parameter (Å)
+1 5.4 28.5 26.2 33.3 6.5 η-M6C FCC, Fd3m 11.38
+2 – 84.1 3.8 3.7 8.3 α-W
+3 – 82.6 7.0 3.9 6.5
+4 – 84.8 4.7 2.0 8.5
BCC, Im3m
+5 – 4.8 85.5 4.4 5.2 M23C6 FCC,
+6 7.7 6.7 55.1 24.7 4.5 M23C6 FCC,
+7 +8 – – 4.4 4.2 24.6 23.5 70.5 71.3 0.5 1.0 Matrix FCC,
3.19
Fm3m 10.87
Fm3m 10.97
Fm3m 3.60–3.62
boundaries after short exposure period [2,14,16]. Further researches are required to identify the precipitate. W-rich precipitates on the grain boundaries formed after long-term aging at 900 °C was characterized using TEM, EDS and SAD analyses. Fig. 7 represents a cross-sectional TEM micrograph of the W-rich precipitate on the grain boundary within the carbide depleted region in Fig. 3(e). The point EDS result shows that it mainly contains around 83 W and 10 Mo (in at.%). This chemical composition is quite similar to that of the α-W phase formed in the dissociated prior M6C carbides shown in Table 4. Also, the SAD pattern inserted in Fig. 7 matches to that of the metallic α-W phase with BCC crystal structure (lattice parameter of 3.16 Å). The formation of the α-W phase on the grain boundaries might be associated with the carbide depletion near the Cr- and Ni-rich precipitates. That is, during the shorter exposure time at 900 °C, Cr-rich M23C6 carbides are formed on the grain boundaries rather than W-rich phases due to the slower diffusion of W compared to Cr [2,15]. However, as further aging progresses, the abovementioned Cr- and Ni-rich precipitates and Crdepleted region are formed in some areas at the expense of Cr-rich M23C6 carbides, though the mechanisms are not well understood [16]. Then, W within the matrix would diffuse to the grain boundaries in the Cr-depleted region and precipitate as the α-W phase during long-term exposure. As the diffusion of W is required, precipitation of the α-W phase would be much faster at higher temperature, which is supported by the precipitation of discontinuous α-W on the grain boundaries for Ni-base alloys with 18 wt.% W exposed for 1 h at 1200 °C [32].
Fig. 6. TEM micrograph, SAD pattern, and chemical compositions of Cr- and Ni-rich precipitates formed after aging at 900 °C for 10,000 h.
Fig. 7. TEM micrograph, SAD pattern, and chemical compositions of W-rich precipitate formed on the grain boundary after aging at 900 °C for 20,000 h.
3.6. Changes in tensile properties after aging at 800 and 900 °C The effects of microstructure evolution during high temperature aging on the tensile properties were assessed at room temperature. Fig. 8 shows the stress–strain curves of tensile tests for as-annealed and aged specimens. In addition, the results of tensile tests are summarized in Table 5. As shown in the figure, the long-term aging (10,000 and 20,000 h) at 800 °C results in the increase in ultimate tensile strength (UTS) and decrease in elongation for Alloy 230. Meanwhile, the longterm aging at 900 °C results in larger decrease in elongation, but decrease in UTS is very small. Fig. 9 shows the cross-sectional microstructures of tensile fractured specimens. In the as-annealed condition, fracture mode is the mixture of the trans- and inter-granular (Fig. 9(a)). However, as Cr-rich M23C6 grain boundary carbides are extensively formed and thickened, inter-granular fracture mode is dominant in the aged specimens (Fig. 9(b) and (c)). Meanwhile, the Cr-rich M23C6 carbides formed within the grain would increase strength instead of lowering the elongation, which is consistent with the tensile test results such that specimen aged at 800 °C (with more Cr-rich M23C6 carbides within the grain) shows higher UTS value than that aged at 900 °C (with less Cr-rich M23C6 carbides within the grain).
Fig. 8. The room temperature tensile stress–strain curves of Alloy 230 after aging at 800 and 900 °C for up to 20,000 h.
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Table 5 The results of room temperature tensile test of as-annealed and aged Alloy 230.
Exposure time Ultimate tensile strength (MPa) Elongation (%)
As-annealed Alloy 230
Aging at 800 °C
843 ± 12 60 ± 1
10,000 h 938 ± 2 45 ± 1
It is also interesting to mention that some primary M6C carbides near the fractured area are cracked in the as-received condition (Fig. 9(a)), although the dissociated and non-dissociated primary M6C carbides are mostly remained un-cracked for the aged specimens (Fig. 9(b) and (c)). Previously, K. Mo [23] reported the fracture of the primary M6C carbide in the as-annealed Alloy 230 was resulted from significant dislocation movement at the transition from elastic to plastic regime when tensile loading was applied. He explained that a great number of dislocations were piled at the large and non-deformable primary M6C carbide. It could not endure large internal stress and begin to rupture upon early yielding. However, in case of the aged specimens in this study, a presence of extensive precipitation of the secondary M23C6 carbides on the grain boundaries and within the grains would relieve internal stress around the primary M6C carbides by blocking dislocation movement. This explains the absence of any cracked primary M6C carbides in the specimen aged at 800 °C for 20,000 h (Fig. 9(b)). On the other hand, for the specimen aged at 900 °C for 20,000 h, some large Cr- and Ni-rich precipitates are cracked as shown in Fig. 9(d). Local cracking of Cr- and Ni-rich precipitates was reported during tensile or creep tests and explained with the brittle nature of the phase [11,16]. Therefore, the cracking of Cr- and Ni-rich precipitates as well as grain boundary carbides contributed to the premature failure of the specimens aged at 900 °C, which resulted in the lower UTS values than the specimens aged at 800 °C as shown in Fig. 8. For both aging temperatures, a degradation of tensile property is observed after 10,000 and 20,000 h exposure, but with different degrees. After aging at 800 °C, the typical tensile property change, such as increase in strength and decrease in elongation, was observed. Such change could be well explained with the microstructure evolution, especially, the precipitation of the secondary Cr-rich M23C6 carbides. It seems that the dissociation of the primary M6C carbide has minimal effects on tensile properties. After aging at 900 °C, formation of Cr-
Aging at 900 °C 20,000 h 935 ± 6 44 ± 1
10,000 h 824 ± 7 38 ± 2
20,000 h 823 ± 19 37 ± 4
and Ni-rich precipitates affects the tensile properties. That is, formation of brittle Cr- and Ni-rich precipitates is associated with the carbide free region and grain boundary α-W phases, which results in lowering elongation as well as strength as observed in tensile tests. Therefore, for the aged specimens, the degradation of tensile property would depend on the evolution of Cr-rich M23C6 carbides, Cr- and Ni-rich precipitates, and α-W phases but not on the primary M6C carbides.
4. Conclusions The effects of long-term exposure to high temperature (800 and 900 °C for up to 20,000 h) on the microstructure evolution of a Ni–Cr– W superalloy (Alloy 230) were investigated. Based on the analysis and tests, the following conclusions are drawn. 1. After long-term aging at 800 °C, extensive precipitation of the secondary Cr-rich M23C6 carbides was observed on the grain boundaries as well as within the grains. In addition, W-rich primary η-M6C carbides were internally dissociated into three distinctive phases, such that W-rich M6C → α-W, Cr-rich M23C6, and W-depleted Nibase matrix. 2. After long-term aging at 900 °C, precipitation of Cr- and Ni-rich intermetallic and α-W phase was observed in some areas at the expense of the secondary Cr-rich M23C6 carbides. Meanwhile, the dissociation of the primary M6C carbides was less significant. 3. Long-term aging at both temperature resulted in decrease in tensile elongation of Alloy 230 while increase in strength was minimal after aging at 900 °C. The degradation of tensile property depends on the evolution of Cr-rich M23C6 carbides, Cr- and Nirich precipitates, and α-W phases but not on the primary M6 C carbides.
Fig. 9. Cross-sectional microstructure of room temperature tensile tested Alloy 230; (a) as-annealed and after 20,000 h aging at (b) 800 and (c), and (d) 900 °C.
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