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Precipitate evolution during the aging of Super304H steel and its influence on impact toughness
Materials Science & Engineering A 754 (2019) 238–245
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Precipitate evolution during the aging of Super304H steel and its influence on impact toughness
T
Xue Wanga,∗, You Lia, Dongxu Chenb, Jianhua Sunb a Key Laboratory of Hydraulic Machinery Transients (Wuhan University), Ministry of Education, School of Power and Mechanical Engineering, Wuhan University, Wuhan, Hubei, 430072, China b Henan Engineering Co. LTD, Power Construction Corporation of China, Zhengzhou, Henan, 450001, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Super304H steel Aging Precipitate Mechanical properties
Super304H, which is an austenitic heat-resistant steel, was aged at 650 °C for up to 5000 h. Its microstructural evolution during aging was observed by scanning electron microscopy (SEM) equipped with electron backscatter diffraction (EBSD). The precipitate evolution in the aged specimens was analyzed by both X-ray diffraction (XRD) and transmission electron microscopy (TEM) with energy dispersive spectroscopy (EDS) and selected area electron diffraction (SAED). In addition, the precipitation kinetics of M23C6 carbide was established. Hardness and impact tests were performed on the aged specimens. Finally, the relationship between the mechanical properties and microstructural evolution was revealed. After aging, Super304H steel exhibited the same fine grains as in the as-received condition with a Cu-rich phase, secondary Nb(C,N) precipitates inside the grains, and M23C6 carbides mainly at grain boundaries. The Cu-rich phase and the secondary Nb(C,N) precipitated in specimens aged for 500 h; this leads to a marked increase in hardness. Precipitation of the Z phase also occurred in the aged Super304H steel. M23C6 carbides precipitated rapidly during aging, with content approaching 90% of the saturated content in specimens aged for 5000 h. The reduction in impact toughness in Super304H steel was found to be associated with the precipitation of the M23C6 carbide, which is similar to that in the HR3C steel. However, the deterioration was much less significant as compared to HR3C steel because of the discontinuous distribution of carbides at grain boundaries due to the fine grain size.
1. Introduction To improve the thermal efficiency of power plants and reduce CO2 emission, ultra-supercritical boilers (USCBs) have been developed in China and other countries. New grade austenitic heat-resistant steels, such as Super304H and HR3C, are widely used in USCBs (mainly in superheater and reheater tubes), due to their combination of excellent creep and oxidation resistance at high temperatures [1–6]. Super304H is a type of 18–8 austenitic heat-resistant steel, developed from the TP304H steel by reducing the amount of manganese and adding approximately 3.0% copper, 0.45% niobium, and trace amounts of nitrogen [7]. Recently, the microstructural evolution and mechanical properties of the Super304H steel during exposure to high temperatures have attracted considerable research attention. Some studies have concentrated on the high-temperature strengthening mechanism of Super304H steel. Jiang et al. [8] revealed that the Cu-rich phase precipitates dispersedly inside grains during aging at 650 °C, effectively improving the strength of the Super304H steel. By utilizing a three-
∗
dimensional atom probe (3DAP), Chi et al. [9] revealed the precipitation behavior of the Cu-rich phase. The growth kinetics of the Cu-rich phase and its strengthening mechanism have also been investigated [10,11]. Ou et al. [12] found that the fine Nb(C,N) particles that precipitated inside the grains improved the strength of Super304H steel during aging at 650 °C. It can be concluded from these studies that the high-temperature strength of Super304H steel is mainly enhanced by the precipitation strengthening effect of the Cu-rich phase and fine Nb (C,N) particles. However, the precipitation behavior of M23C6 has seldom been explored, particularly with respect to its effects on the mechanical properties of Super304H steel. Wang et al. [13] reported that HR3C steel showed low impact toughness (less than 15 J/cm2) after exposure at 650 °C for 6000 h. Zieliński et al. [14] indicated that the coarse grains in an as-received HR3C steel specimen (grain size of 4) and the precipitation of M23C6 resulted in low ductile properties (impact energy KV < 27 J) after aging at 650 °C for 1000 h. Zhu et al. [15] conducted a study to compare the impact toughness of modified and commercial HR3C steel after
Corresponding author. E-mail address: [email protected] (X. Wang).
https://doi.org/10.1016/j.msea.2019.03.086 Received 15 October 2018; Received in revised form 18 March 2019; Accepted 19 March 2019 Available online 22 March 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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after aging for 500 h and 5000 h was 13.01 μm and 13.61 μm, respectively (Fig. 1b and c). Considering the statistical error of the EBSD instrument, the average size of grains during aging was considered to be stable, which is consistent with reports in previous studies [18,19].
long-term thermal exposure at 650 °C. Their results revealed that coarse and continuous M23C6 carbides at grain boundaries promote a greater reduction in the impact toughness. Therefore, the deterioration of the HR3C impact toughness can mainly be attributed to the precipitation of M23C6 carbides. The carbon content of Super304H steel is higher than that of HR3C steel (in the ASME specification, the carbon content of Super304H steel is 0.07–0.13%, whereas that of HR3C steel is 0.04–0.10%). Precipitates of M23C6 carbides have also been observed at grain boundaries in Super304H steel. However, the effect of M23C6 on the impact toughness of the Super304H steel has been rarely reported. In this study, aging tests at 650 °C for up to 5000 h were performed on a commercial Super304H steel tube to investigate the microstructural evolution during aging. The precipitation of M23C6 carbides and other precipitates and their influence on the impact toughness were thoroughly studied, providing a reference for the safe use of Super304H steel in USCBs.
3.2. Precipitate evolution during aging Fig. 2 shows the XRD results of the as-received and aged specimens. In the as-received specimen, only strong diffraction peaks of γ-Fe and relatively weaker peaks of Nb(C,N) were found. After aging for 500 h, 3000 h, and 5000 h, in addition to the γ-Fe and Nb(C,N) peaks, some M23C6 peaks were observed in the patterns, as shown clearly in the section marked “a1” in Fig. 2a. Some γ-Fe peaks shifted slightly to the right during aging. For instance, as shown in Fig. 2b, the value of 2θ for the (2 2 0) peak increased with aging time, especially in the first 1000 h. In the as-received specimen, most of the alloy elements were dissolved in the austenite matrix, particularly the interstitial solution of carbon and nitrogen atoms. This resulted in a lower lattice constant and lower value of 2θ for the (2 2 0) peak. For instance, the value of 2θ for the (2 2 0) peak of the as-received specimen was 74.4°, which was smaller than the standard value of 74.68° stated on the PDF card. During the aging process, carbon and nitrogen were released from the matrix, which reduces the matrix lattice distortion. This resulted in an increase in the 2θ values (i.e., the diffraction peaks moved gradually to the right). The peak corresponding to the Cu-rich phase was not detected in the XRD test, because of its small amount and size. Furthermore, the Cu-rich phase had a parallel relationship with γ-Fe, leading to an overlap of their diffraction peaks. Fig. 3 shows the SEM images of as-received and aged specimens. Granular precipitates identified as Nb(C,N) by EDS analysis and with different sizes were scattered inside the grains and at grain boundaries in the as-received specimen. The large particles were Nb(C,N) that remained undissolved during high temperature softening treatment. These are referred to as primary Nb(C,N). The small particles probably precipitated during the subsequent cold rolling or final solid-solution treatment. These are referred to as secondary Nb(C,N). After aging at 650 °C for 500 h, the number of precipitates inside the grains tended to increase slightly compared to the as-received specimen. The fresh precipitates distributed uniformly inside grains were secondary Nb(C,N) [20]. From Fig. 3b, it can be seen that precipitates rich in Cr carbides were dense and formed a chain of beads at grain boundaries. After aging for 5000 h, the precipitates of Nb(C,N) inside grains were unaffected, whereas the amount of precipitates rich in Cr carbides at the grain boundaries increased, leading to the formation of chain-like carbides. TEM images of the aged specimens are shown in Fig. 4. In the specimen aged for 500 h, primary Nb(C,N) particles, with diameters greater than 300 nm, as well as secondary Nb(C,N) particles of 40–60 nm were observed inside grains. Carbide particles were distributed uniformly at grain boundaries, which were measured with widths of 50–70 nm (Fig. 4a). These were confirmed to be M23C6 carbides by SAED (Fig. 4b). In addition, numerous very fine particles precipitated dispersedly inside grains. These were identified as Cu-rich
2. Experimental The composition of the investigated Super304H steel was: 0.075C, 0.45Si, 0.76Mn, 18.46Cr, 8.84Ni, 2.84Cu, 0.29Nb, 0.104N, and 0.15Mo (wt%, bal. Fe). A tube with an outer diameter of 42 mm and a wall thickness of 9 mm, was aged at 650 °C for 500 h, 1000 h, 3000 h, and 5000 h. Specimens with a size of 7.5 mm × 10 mm × 55 mm were cut from aged tubes for a Charpy V impact test at room temperature. The hardness was examined using a Brinell hardness tester (320HBS-3000/ 0035) with a load of 750 kgf. Specimens were prepared by mechanical polishing and then etched in the aqua regia. The phase constituents in both the as-received and aged specimens were analyzed by XRD (D/ MAX-2500). The microstructural evolution during aging was observed by SEM equipped with EBSD (MIRA3LMH). Extraction replicas and thin foils were prepared to further identify precipitates using TEM (JEM2010) with SAED and EDS. The equilibrium phase fraction of precipitates in the Super304H steel at 650 °C was calculated using Thermal-Calc. 3. Results 3.1. Microstructural evolution during aging The as-received Super304H steel was found to have very fine grains, with an average diameter of 14.27 μm, as shown in Fig. 1a, which is similar to that reported in the literature [16]. By contrast, HR3C austenitic steel was found to have much larger grains, with a diameter of 70 μm [17]. In general, the fine grains of Super304H steel are derived from advanced softening treatment at a high temperature of 1250–1300 °C. During this process, most of the primary niobium-rich precipitates dissolve and then precipitate again in the subsequent cold rolling process, which effectively hinders the growth of grains by pinning grain boundaries. Considerable twins, which were formed during deformation and the final solution heat-treatment, were observed in the as-received specimen. The average diameter of grains in specimens
Fig. 1. Images of Super304H steel: (a) as-received and aged at 650 °C for (b) 500 h and (c) 5000 h. 239
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Fig. 2. XRD results of aged Super304H steel specimens. (a) XRD pattern and (b) shifting of diffraction angle for the (2 2 0) γ-Fe peak with aging time.
phase and the average values of elements are shown in Table 1. The M23C6 was composed mainly of Cr and Fe, and included additional elements (i.e., Nb, Mn, and Mo). Aging for 500–5000 h resulted in a slight increase in the Cr content and a gradual decrease in the Fe content. In addition to being rich in Nb elements, the Nb(C,N) particles also contained Cr and Fe. During the aging process, the content of Nb decreased. In addition, some nano-sized particles, with a Nb/(Cr + Fe) ratio close to 1, were found in the samples aged for 500 h (the standard deviation was not calculated because of the very low content). These particles were indicated to be Z phase particles, according to thermodynamic calculations. In the samples aged for 5000 h, the size of the Z phase particles remained unchanged, but the quantity increased, as shown in Fig. 5d. Iseda et al. [21] also reported the existence of a Z phase in a Super304H steel tube serviced at 600 °C for 85426 h.
particles. After aging for 5000 h, the primary Nb(C,N) seemed to be stable, whereas the M23C6 particles grew considerably, where the width of the grain boundary increased to 100 nm (Fig. 4c). Fig. 4d shows the diffraction pattern of M23C6 in samples aged for 5000 h. In addition, a coffee-bean shaped Cu-rich phase was observed inside grains, exhibiting an obvious growth from 3 nm to 25 nm compared to samples aged for 500 h. This was consistent with the results of Chi et al. [9]. TEM images of precipitates in specimens of the Super304H steel aged for 500 h and 5000 h are shown in Fig. 5. Although the size and quantity of M23C6 at grain boundaries both increased with increasing aging time, discontinuous distribution remained. Nb(C,N) particles were found inside the grains along with a few nano-sized particles, the compositions of which were similar to those of the Z phase. The compositions of precipitates in the replica samples were determined by EDS. More than five particles were measured for each
Fig. 3. SEM images of Super304H steel: (a) as-received and aged at 650 °C for (b) 500 h, (c) 1000 h, and (d) 5000 h. 240
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Fig. 4. TEM images of Super304H steel aged at 650 °C for: (a), (b) 500 h and (c), (d) 5000 h.
Fig. 5. TEM images of precipitates in the Super304H steel aged at 650 °C for: (a), (b) 500 h and (c), (d) 5000 h. 241
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Table 1 Composition of precipitates determined through EDS in Super304H steel aged at 650 °C for different times (atomic %).
500 h
5000 h
Phase
Cr
Mn
M23C6 Nb(C,N) Z phase M23C6 Nb(C,N) Z phase
77.33 ± 2.56 9.82 ± 4.26 40.48 81.89 ± 3.38 19.32 ± 9.51 40.06 ± 9.16
0.75 0.16 0.28 0.70 0.24 0.24
± 0.48 ± 0.21 ± 0.33 ± 0.12 ± 0.11
Fe
Nb
Mo
20.32 ± 2.27 2.28 ± 1.19 8.66 15.33 ± 2.32 3.76 ± 1.71 6.13 ± 1.21
1.02 ± 2.64 87.55 ± 5.47 50.36 1.36 ± 3.42 76.41 ± 11.43 53.16 ± 8.97
0.58 0.18 0.22 0.71 0.26 0.41
± 0.19 ± 0.17 ± 0.31 ± 0.19 ± 0.15
5000 h, the Nb content of the Nb(C,N) phase decreased, whereas the Cr content increased markedly. This reveals the transformation of the Nb (C,N) phase to the Z phase for a long-term aging process. The Cu-rich phase mainly consists of Cu and few other atoms, with the Cu content approximating 98%. As seen in Table 2, the σ phase in Super304H steel, which is a brittle phase with a complex crystal structure, contains mainly Fe and Cr atoms. The tendency of the σ phase to precipitate in austenitic steels is usually calculated using the following Cr-equivalent equation [22]:
Creq = Cr + 0.31Mn + 1.76Mo + 0.97W + 2.02V + 1.58Si + 2.4Ti (1)
+ 1.76Nb + 1.22Ta − 0.226Ni − 0.177Co
WhenCreq is higher than 17–18%, the precipitation tendency of the σ phase is high [23]. The Creq value of the investigated Super304H steel was 18.18%, showing that the precipitation of the σ phase is highly probable. However, the σ phase was not observed, even in the specimen aged for 5000 h, indicating that the σ phase in Super304H steel requires a longer time to precipitate. Vach et al. [24] reported that the σ phase was found in Super304H steel aged at 650 °C for 87600 h. Moreover, some investigations [25] showed that the σ phase preferentially nucleates in the carbon-depleted region around M23C6 carbides. Iseda et al. [21] observed the σ phase at grain boundaries in a HR3C steel kept at 700 °C and 69 MPa for 88362 h. Thus the precipitation behavior of the σ phase in Super304H steel during longer periods of aging at 650 °C or at higher temperatures should be studied.
Fig. 6. Calculated equilibrium phase of Super304H steel.
3.3. Properties of precipitates Fig. 6 shows the fractions of precipitates at the equilibrium state from 500 °C to 1300 °C in the Super304H steel, which was determined using Thermal-Calc. It can be seen that the σ phase and Cu-rich phase dissolved at approximately 740 °C and 860 °C, respectively, whereas the Z phase and the M23C6 carbides dissolved at 930 °C and 950 °C, respectively. For temperatures higher than 1100 °C, only the Nb(C,N) phase existed, which is consistent with the microstructure of the asreceived Super304H steel, where only Nb(C,N) particles were present. According to Fig. 6, M23C6 carbides as well as Cu-rich, σ, and Z phases existed in the Super304H steel at the 650 °C equilibrium. The amount of the σ phase was the largest (9.41%), followed by the M23C6 carbides (2.09%) and the Z phase (0.51%). It should be noted that at the 650 °C equilibrium state, no Nb(C,N) phase was present in the Super304H steel because it tended to transform to the more stable Nb-rich Z phase during the long-term aging process. The chemical compositions of precipitates at the 650 °C equilibrium state in the Super304H steel are listed in Table 2. Normalized without considering the content of carbon, the contents of Cr and Fe in the M23C6 carbides were 83.98% and 10.91%, respectively. Compared with the experimental results of aged specimens for different times (as listed in Table 1), it can be concluded that the Cr content in M23C6 increased slowly with increasing aging time, whereas the Fe content decreased gradually. Table 2 shows that the Z phase has a chemical formula of Nb (Cr,Fe)N. Table 1 shows that as the aging time increased from 500 h to
3.4. Mechanical properties Fig. 7 shows the hardness and impact toughness of specimens annealed with different aging times. During the early stage of the aging process, the hardness increased rapidly from 157HB for the as-received sample to 191HB after 500 h. For longer aging times, the hardness was nearly stable and reached 200HB after aging for 5000 h. The impact toughness dropped rapidly in the primary exposure period (0–500 h) but changed little in the later exposure period (500–5000 h). Zhu et al. [15] investigated the impact toughness evolution of commercial HR3C steel during aging at 650 °C. Those results are also shown in Fig. 7. The figure shows that the toughness of the HR3C steel dropped much more rapidly than that of the Super304H steel. After exposure for 5000 h, the impact toughness of the HR3C steel fell to only 25 J/cm2, but that of the Super304H steel remained at 160 J/cm2. The fractures of aged Super304H steel were analyzed by SEM and the images are shown in Fig. 8. The as-received sample presented a ductile fracture with large and deep plastic-deformed dimples and some precipitates distributed at the bottom of the dimples. After aging for
Table 2 Chemical compositions of precipitates in Super304H steel at 650 °C equilibrium (atomic %). Phase
Cr
Fe
Nb
N
C
Cu
Mn
Ni
Si
Mo
M23C6 Cu-rich phase Z phase σ phase
66.60 0.01 26.75 41.74
8.65 0.30 6.60 54.45
– – 32.90 –
– – 33.31 –
20.69 – – –
– 97.96 – –
0.13 0.24 – 0.37
0.18 1.39 – 1.73
– – – 1.69
3.74 0.09 0.45 0.02
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the bonding between M23C6 and the grain is weak, cracks may mainly be generated near the M23C6 particles. The increase in the number and decrease in the size of dimples are associated with the increasing amount of M23C6 during aging, revealing that the precipitation of M23C6 is closely related to the decline in impact toughness. Regarding the HR3C steel, Wang et al. [13] reported that the impact fracture was a brittle rock candy-like fracture after aging at 650 °C for 6000 h. Zieliński et al. [14] and Golański et al. [27] indicated the tendency of M23C6 carbides to coagulate during service and to form a continuous network at the grain boundaries, which causes a rapid decrease in impact toughness in HR3C steel. The decrease in impact toughness in the Super304H and HR3C steel during aging was mainly due to the precipitation of M23C6 at grain boundaries. The large difference in embrittlement during aging between these two types of steel may be because of the difference in grain size, which produces marked changes in the distribution and morphology of M23C6 in grain boundaries. This results in different fracture modes.
Fig. 7. Mechanical properties of Super304H and HR3C steels as a function of aging time.
4. Discussion 500 h, the number of dimples increased markedly but were smaller and shallower than those in the as-received sample. In addition, the number of precipitates presented in the dimples also increased. With increasing aging time, the size and depth of the dimples seemed to decrease, but the features of ductile fracture could still be observed in the specimen aged for 5000 h. To understand the correlation between the second phase and the impact fracture, the precipitates in the dimples were analyzed. Large fragmented precipitates with a diameter of 1–2 μm were found at the bottom of the dimples in the as-received specimen and were identified as primary Nb(C,N) (shown in Fig. 9a). These large, brittle precipitates resulted in stress concentration under impact loading and their cracking or detachment from the matrix resulted in lower energy being required for crack initiation and expansion. In the specimen aged for 5000 h, precipitates at the dimples were identified as M23C6 carbides by EDS (shown in Fig. 9b). During aging, secondary Nb(C,N), the Cu-rich phase, and the Z phase precipitated inside grains. These precipitates were very fine and scattered, which could result in precipitation strengthening and may not impair the ductility. In general, the precipitation of M23C6 carbides leads to a depletion of C and Cr atoms in the surrounding matrix [26]. As the strength of the depletion area is relatively low and
4.1. Precipitation kinetics of M23C6 According to the analysis above, the precipitation of M23C6 carbides at grain boundaries is the main reason for the decrease in impact toughness in Super304H steel during aging. To investigate the evolution of impact toughness after long-term exposure, analyzing the precipitation kinetics of M23C6 in Super304H steel is necessary. The precipitation of M23C6 is a process that involves the diffusion and bonding of C and Cr atoms [28,29], which nucleate preferentially at positions with high energy and defects in the matrix, such as grain boundaries, vacancies, and dislocations [30]. As the atomic number of the Nb atom is much higher than that of Cr, the Nb(C,N) particles are white and bright, whereas M23C6 particles are dark gray in the back-scattered electron (BSE) images obtained through SEM. Nano-sized particles such as the Cu-rich phase, Z phase, and so on are difficult to observe in the BSE images. Thus, the M23C6 carbide can be distinguished and its volume fraction can be calculated based on the BSE images. The measured volume fractions of M23C6 in the as-received specimens and those aged for 500 h, 3000 h, and 5000 h are 0, 0.97%, 1.58%, and 1.84%, respectively. The precipitation kinetics of M23C6 is described by the Johnson-
Fig. 8. Fractography of impact samples. (a) as-received and aged at 650 °C for (b) 500 h, (c) 3000 h, and (d) 5000 h. 243
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Fig. 9. Precipitates on the fracture surface of Super304H steel. (a) As-received and (b) aged at 650 °C for 5000 h.
that mainly affected the mechanical properties of the Super304H steel. During the primary stage of aging (0–500 h), the precipitation of large amounts of M23C6 and nano-sized secondary Nb(C,N) particles caused the release of C and N atoms from the matrix (consistent with the right shift of the γ-Fe diffraction peaks), thus weakening the solution strength. However, the Cu-rich phase and secondary Nb(C,N) particles effectively pinned dislocations, thereby improving the precipitation strengthening effectively. Thus, the hardness of the specimen aged for 500 h was significantly higher than that of the as-received specimen. From 500 h to 5000 h, the Cu-rich phase and M23C6 continued to precipitate and coarsen. The nano-sized Cu-rich phase was coherent with the γ-Fe matrix, and its growth intensified the distortion of the lattice, offsetting the reduction of lattice distortion caused by the precipitation of C atoms. Hence, the diffraction peaks of γ-Fe remained nearly unchanged, which was in agreement with the stable hardness values of the Super304H steel during aging from 500 h to 5000 h. The decrease in toughness in the Super304H steel was mainly the result of the precipitation of M23C6. The impact toughness ak and volume fraction of M23C6 (V ) were fitted, as shown in Fig. 11. The function of this fit is shown by the following equation, with an R2 value of 0.99.
Mehl-Avrami equation [31]:
f (t ) = 1 −
exp (−kt n )
(2)
where f (t ) is the transformation fraction of a new phase at a certain temperature after a certain time t , n is the Avrami index, and k is the constant of crystallization rate. Here, f (t ) can also be defined by the ratio between the volume fraction of M23C6 (V (t )) after a certain time t , and the maximum volume fraction of M23C6 (Vm) when fully precipitated. Thus, the following equation can be written based on the J-MA equation:
f (t ) =
V (t ) = 1 − exp (−kt n ) Vm
(3)
where Vm is 2.09%, as calculated by Thermal-Calc at the 650 °C equilibrium state of the Super304H steel. Taking the volume fraction of M23C6 measured at 500 h, 3000 h, and 5000 h as V (t ) , three groups of data were obtained. Based on these data, the value of n and k were calculated to be 0.554 and 0.021, respectively. The J-M-A curve, as shown in Fig. 10 was also obtained. It can be seen that the M23C6 carbides began to precipitate quickly at 50–100 h and the volume fraction of M23C6 at 5000 h was close to 90% of the saturated precipitation content. The M23C6 in grain boundaries at 5000 h widened clearly as compared with that at 500 h (Fig. 4), which coincides with the increased amount of M23C6.
ak = 1/(0.0047 + 0.0012V − 0.0003V 2)
(4)
Fig. 11 shows that the decrease in impact toughness was in good agreement with the increase in the volume fraction of M23C6 in Super304H steel during aging. For a comparison, assuming that the decline of impact toughness of HR3C steel is only affected by M23C6, the values for the HR3C steel was calculated based on data from the literature [13]. These are also shown in Fig. 11. For the same amount of M23C6, the decline in toughness in the Super304H steel was lower than that in the HR3C steel, revealing
4.2. Effect of M23C6 precipitation on impact toughness The grain size in the Super304H steel was stable during aging (Fig. 1). Thus, the increase in hardness and decrease in toughness were mainly attributed to the evolution of precipitates. As the primary Nb (C,N) phase was stable and the Z phase was not significantly present, it was the secondary Nb(C,N) phase, Cu-rich phase, and M23C6 particles
Fig. 11. Effect of the amount of M23C6 phase on Charpy impact value (ak ) in Super304H steel.
Fig. 10. Precipitation kinetics curves of M23C6 phase in Super304H steel. 244
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that the difference in the aging embrittlement between these two types of austenitic steels derived from their different grain sizes. The Super304H steel had relatively fine grains (grain size of 9), whereas the HR3C steel had much coarser grains (grain size of 4) [14]. For the same amount of M23C6, the finer the grains were, the more difficult it was to form a continuous distribution of M23C6 at grain boundaries. As shown in Fig. 3d, although the content of M23C6 in the Super304H steel after 5000 h of aging approximated the amount of saturated precipitation, its distribution in the grain boundaries was discontinuous and appeared as scattered granules, unlike the network distribution of M23C6 in the HR3C steel. Thus, the negative influence of M23C6 precipitation on the toughness was insignificant. Based on the above analysis, the fine grains resulted in a considerably reduced aging embrittlement for the Super304H steel. It should also be noted that no σ phase was observed during aging at 650 °C for 5000 h. However, according to the thermodynamic calculation by Thermal-Calc, the σ phase could potentially precipitate in the Super304H steel. The precipitation of the σ phase at grain boundaries was found to be seriously harmful to the toughness [18]. Therefore, studying the precipitation behavior of the σ phase in the Super304H steel in future works is necessary.
[15]
5. Conclusion
[16]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
1. During aging at 650 °C for different times up to 5000 h, the Super304H steel retained a microstructure with fine grains, consisting of a Cu-rich phase, secondary Nb(C,N), as well as a few Z phase particles that precipitated inside the grains and M23C6 carbides that precipitated at the grain boundaries. 2. The Cu-rich phase and secondary Nb(C,N) particles precipitated in specimens aged for 500 h. These particles retained their small size with prolonged aging time, which resulted in a marked increase in hardness. 3. M23C6 carbides precipitated rapidly during aging, and their amounts approached 90% of the saturated content in the specimen aged for 5000 h. The precipitation of M23C6 in the Super304H steel led to a clear reduction in impact toughness. However, the embrittlement in Super304H steel was much lower than that in HR3C steel. This was because of the discontinuous distribution of M23C6 at the grain boundaries due to the finer grain size in the Super304H steel.
[17]
[18]
[19]
[20]
[21]
[22]
Acknowledgments
[23]
The authors would like to express their gratitude for projects supported by the National Natural Science Foundation of China (51574181; 51374153).
[24]
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Precipitate evolution during the aging of Super304H steel and its influence on impact toughness
T
Xue Wanga,∗, You Lia, Dongxu Chenb, Jianhua Sunb a Key Laboratory of Hydraulic Machinery Transients (Wuhan University), Ministry of Education, School of Power and Mechanical Engineering, Wuhan University, Wuhan, Hubei, 430072, China b Henan Engineering Co. LTD, Power Construction Corporation of China, Zhengzhou, Henan, 450001, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Super304H steel Aging Precipitate Mechanical properties
Super304H, which is an austenitic heat-resistant steel, was aged at 650 °C for up to 5000 h. Its microstructural evolution during aging was observed by scanning electron microscopy (SEM) equipped with electron backscatter diffraction (EBSD). The precipitate evolution in the aged specimens was analyzed by both X-ray diffraction (XRD) and transmission electron microscopy (TEM) with energy dispersive spectroscopy (EDS) and selected area electron diffraction (SAED). In addition, the precipitation kinetics of M23C6 carbide was established. Hardness and impact tests were performed on the aged specimens. Finally, the relationship between the mechanical properties and microstructural evolution was revealed. After aging, Super304H steel exhibited the same fine grains as in the as-received condition with a Cu-rich phase, secondary Nb(C,N) precipitates inside the grains, and M23C6 carbides mainly at grain boundaries. The Cu-rich phase and the secondary Nb(C,N) precipitated in specimens aged for 500 h; this leads to a marked increase in hardness. Precipitation of the Z phase also occurred in the aged Super304H steel. M23C6 carbides precipitated rapidly during aging, with content approaching 90% of the saturated content in specimens aged for 5000 h. The reduction in impact toughness in Super304H steel was found to be associated with the precipitation of the M23C6 carbide, which is similar to that in the HR3C steel. However, the deterioration was much less significant as compared to HR3C steel because of the discontinuous distribution of carbides at grain boundaries due to the fine grain size.
1. Introduction To improve the thermal efficiency of power plants and reduce CO2 emission, ultra-supercritical boilers (USCBs) have been developed in China and other countries. New grade austenitic heat-resistant steels, such as Super304H and HR3C, are widely used in USCBs (mainly in superheater and reheater tubes), due to their combination of excellent creep and oxidation resistance at high temperatures [1–6]. Super304H is a type of 18–8 austenitic heat-resistant steel, developed from the TP304H steel by reducing the amount of manganese and adding approximately 3.0% copper, 0.45% niobium, and trace amounts of nitrogen [7]. Recently, the microstructural evolution and mechanical properties of the Super304H steel during exposure to high temperatures have attracted considerable research attention. Some studies have concentrated on the high-temperature strengthening mechanism of Super304H steel. Jiang et al. [8] revealed that the Cu-rich phase precipitates dispersedly inside grains during aging at 650 °C, effectively improving the strength of the Super304H steel. By utilizing a three-
∗
dimensional atom probe (3DAP), Chi et al. [9] revealed the precipitation behavior of the Cu-rich phase. The growth kinetics of the Cu-rich phase and its strengthening mechanism have also been investigated [10,11]. Ou et al. [12] found that the fine Nb(C,N) particles that precipitated inside the grains improved the strength of Super304H steel during aging at 650 °C. It can be concluded from these studies that the high-temperature strength of Super304H steel is mainly enhanced by the precipitation strengthening effect of the Cu-rich phase and fine Nb (C,N) particles. However, the precipitation behavior of M23C6 has seldom been explored, particularly with respect to its effects on the mechanical properties of Super304H steel. Wang et al. [13] reported that HR3C steel showed low impact toughness (less than 15 J/cm2) after exposure at 650 °C for 6000 h. Zieliński et al. [14] indicated that the coarse grains in an as-received HR3C steel specimen (grain size of 4) and the precipitation of M23C6 resulted in low ductile properties (impact energy KV < 27 J) after aging at 650 °C for 1000 h. Zhu et al. [15] conducted a study to compare the impact toughness of modified and commercial HR3C steel after
Corresponding author. E-mail address: [email protected] (X. Wang).
https://doi.org/10.1016/j.msea.2019.03.086 Received 15 October 2018; Received in revised form 18 March 2019; Accepted 19 March 2019 Available online 22 March 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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after aging for 500 h and 5000 h was 13.01 μm and 13.61 μm, respectively (Fig. 1b and c). Considering the statistical error of the EBSD instrument, the average size of grains during aging was considered to be stable, which is consistent with reports in previous studies [18,19].
long-term thermal exposure at 650 °C. Their results revealed that coarse and continuous M23C6 carbides at grain boundaries promote a greater reduction in the impact toughness. Therefore, the deterioration of the HR3C impact toughness can mainly be attributed to the precipitation of M23C6 carbides. The carbon content of Super304H steel is higher than that of HR3C steel (in the ASME specification, the carbon content of Super304H steel is 0.07–0.13%, whereas that of HR3C steel is 0.04–0.10%). Precipitates of M23C6 carbides have also been observed at grain boundaries in Super304H steel. However, the effect of M23C6 on the impact toughness of the Super304H steel has been rarely reported. In this study, aging tests at 650 °C for up to 5000 h were performed on a commercial Super304H steel tube to investigate the microstructural evolution during aging. The precipitation of M23C6 carbides and other precipitates and their influence on the impact toughness were thoroughly studied, providing a reference for the safe use of Super304H steel in USCBs.
3.2. Precipitate evolution during aging Fig. 2 shows the XRD results of the as-received and aged specimens. In the as-received specimen, only strong diffraction peaks of γ-Fe and relatively weaker peaks of Nb(C,N) were found. After aging for 500 h, 3000 h, and 5000 h, in addition to the γ-Fe and Nb(C,N) peaks, some M23C6 peaks were observed in the patterns, as shown clearly in the section marked “a1” in Fig. 2a. Some γ-Fe peaks shifted slightly to the right during aging. For instance, as shown in Fig. 2b, the value of 2θ for the (2 2 0) peak increased with aging time, especially in the first 1000 h. In the as-received specimen, most of the alloy elements were dissolved in the austenite matrix, particularly the interstitial solution of carbon and nitrogen atoms. This resulted in a lower lattice constant and lower value of 2θ for the (2 2 0) peak. For instance, the value of 2θ for the (2 2 0) peak of the as-received specimen was 74.4°, which was smaller than the standard value of 74.68° stated on the PDF card. During the aging process, carbon and nitrogen were released from the matrix, which reduces the matrix lattice distortion. This resulted in an increase in the 2θ values (i.e., the diffraction peaks moved gradually to the right). The peak corresponding to the Cu-rich phase was not detected in the XRD test, because of its small amount and size. Furthermore, the Cu-rich phase had a parallel relationship with γ-Fe, leading to an overlap of their diffraction peaks. Fig. 3 shows the SEM images of as-received and aged specimens. Granular precipitates identified as Nb(C,N) by EDS analysis and with different sizes were scattered inside the grains and at grain boundaries in the as-received specimen. The large particles were Nb(C,N) that remained undissolved during high temperature softening treatment. These are referred to as primary Nb(C,N). The small particles probably precipitated during the subsequent cold rolling or final solid-solution treatment. These are referred to as secondary Nb(C,N). After aging at 650 °C for 500 h, the number of precipitates inside the grains tended to increase slightly compared to the as-received specimen. The fresh precipitates distributed uniformly inside grains were secondary Nb(C,N) [20]. From Fig. 3b, it can be seen that precipitates rich in Cr carbides were dense and formed a chain of beads at grain boundaries. After aging for 5000 h, the precipitates of Nb(C,N) inside grains were unaffected, whereas the amount of precipitates rich in Cr carbides at the grain boundaries increased, leading to the formation of chain-like carbides. TEM images of the aged specimens are shown in Fig. 4. In the specimen aged for 500 h, primary Nb(C,N) particles, with diameters greater than 300 nm, as well as secondary Nb(C,N) particles of 40–60 nm were observed inside grains. Carbide particles were distributed uniformly at grain boundaries, which were measured with widths of 50–70 nm (Fig. 4a). These were confirmed to be M23C6 carbides by SAED (Fig. 4b). In addition, numerous very fine particles precipitated dispersedly inside grains. These were identified as Cu-rich
2. Experimental The composition of the investigated Super304H steel was: 0.075C, 0.45Si, 0.76Mn, 18.46Cr, 8.84Ni, 2.84Cu, 0.29Nb, 0.104N, and 0.15Mo (wt%, bal. Fe). A tube with an outer diameter of 42 mm and a wall thickness of 9 mm, was aged at 650 °C for 500 h, 1000 h, 3000 h, and 5000 h. Specimens with a size of 7.5 mm × 10 mm × 55 mm were cut from aged tubes for a Charpy V impact test at room temperature. The hardness was examined using a Brinell hardness tester (320HBS-3000/ 0035) with a load of 750 kgf. Specimens were prepared by mechanical polishing and then etched in the aqua regia. The phase constituents in both the as-received and aged specimens were analyzed by XRD (D/ MAX-2500). The microstructural evolution during aging was observed by SEM equipped with EBSD (MIRA3LMH). Extraction replicas and thin foils were prepared to further identify precipitates using TEM (JEM2010) with SAED and EDS. The equilibrium phase fraction of precipitates in the Super304H steel at 650 °C was calculated using Thermal-Calc. 3. Results 3.1. Microstructural evolution during aging The as-received Super304H steel was found to have very fine grains, with an average diameter of 14.27 μm, as shown in Fig. 1a, which is similar to that reported in the literature [16]. By contrast, HR3C austenitic steel was found to have much larger grains, with a diameter of 70 μm [17]. In general, the fine grains of Super304H steel are derived from advanced softening treatment at a high temperature of 1250–1300 °C. During this process, most of the primary niobium-rich precipitates dissolve and then precipitate again in the subsequent cold rolling process, which effectively hinders the growth of grains by pinning grain boundaries. Considerable twins, which were formed during deformation and the final solution heat-treatment, were observed in the as-received specimen. The average diameter of grains in specimens
Fig. 1. Images of Super304H steel: (a) as-received and aged at 650 °C for (b) 500 h and (c) 5000 h. 239
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Fig. 2. XRD results of aged Super304H steel specimens. (a) XRD pattern and (b) shifting of diffraction angle for the (2 2 0) γ-Fe peak with aging time.
phase and the average values of elements are shown in Table 1. The M23C6 was composed mainly of Cr and Fe, and included additional elements (i.e., Nb, Mn, and Mo). Aging for 500–5000 h resulted in a slight increase in the Cr content and a gradual decrease in the Fe content. In addition to being rich in Nb elements, the Nb(C,N) particles also contained Cr and Fe. During the aging process, the content of Nb decreased. In addition, some nano-sized particles, with a Nb/(Cr + Fe) ratio close to 1, were found in the samples aged for 500 h (the standard deviation was not calculated because of the very low content). These particles were indicated to be Z phase particles, according to thermodynamic calculations. In the samples aged for 5000 h, the size of the Z phase particles remained unchanged, but the quantity increased, as shown in Fig. 5d. Iseda et al. [21] also reported the existence of a Z phase in a Super304H steel tube serviced at 600 °C for 85426 h.
particles. After aging for 5000 h, the primary Nb(C,N) seemed to be stable, whereas the M23C6 particles grew considerably, where the width of the grain boundary increased to 100 nm (Fig. 4c). Fig. 4d shows the diffraction pattern of M23C6 in samples aged for 5000 h. In addition, a coffee-bean shaped Cu-rich phase was observed inside grains, exhibiting an obvious growth from 3 nm to 25 nm compared to samples aged for 500 h. This was consistent with the results of Chi et al. [9]. TEM images of precipitates in specimens of the Super304H steel aged for 500 h and 5000 h are shown in Fig. 5. Although the size and quantity of M23C6 at grain boundaries both increased with increasing aging time, discontinuous distribution remained. Nb(C,N) particles were found inside the grains along with a few nano-sized particles, the compositions of which were similar to those of the Z phase. The compositions of precipitates in the replica samples were determined by EDS. More than five particles were measured for each
Fig. 3. SEM images of Super304H steel: (a) as-received and aged at 650 °C for (b) 500 h, (c) 1000 h, and (d) 5000 h. 240
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Fig. 4. TEM images of Super304H steel aged at 650 °C for: (a), (b) 500 h and (c), (d) 5000 h.
Fig. 5. TEM images of precipitates in the Super304H steel aged at 650 °C for: (a), (b) 500 h and (c), (d) 5000 h. 241
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Table 1 Composition of precipitates determined through EDS in Super304H steel aged at 650 °C for different times (atomic %).
500 h
5000 h
Phase
Cr
Mn
M23C6 Nb(C,N) Z phase M23C6 Nb(C,N) Z phase
77.33 ± 2.56 9.82 ± 4.26 40.48 81.89 ± 3.38 19.32 ± 9.51 40.06 ± 9.16
0.75 0.16 0.28 0.70 0.24 0.24
± 0.48 ± 0.21 ± 0.33 ± 0.12 ± 0.11
Fe
Nb
Mo
20.32 ± 2.27 2.28 ± 1.19 8.66 15.33 ± 2.32 3.76 ± 1.71 6.13 ± 1.21
1.02 ± 2.64 87.55 ± 5.47 50.36 1.36 ± 3.42 76.41 ± 11.43 53.16 ± 8.97
0.58 0.18 0.22 0.71 0.26 0.41
± 0.19 ± 0.17 ± 0.31 ± 0.19 ± 0.15
5000 h, the Nb content of the Nb(C,N) phase decreased, whereas the Cr content increased markedly. This reveals the transformation of the Nb (C,N) phase to the Z phase for a long-term aging process. The Cu-rich phase mainly consists of Cu and few other atoms, with the Cu content approximating 98%. As seen in Table 2, the σ phase in Super304H steel, which is a brittle phase with a complex crystal structure, contains mainly Fe and Cr atoms. The tendency of the σ phase to precipitate in austenitic steels is usually calculated using the following Cr-equivalent equation [22]:
Creq = Cr + 0.31Mn + 1.76Mo + 0.97W + 2.02V + 1.58Si + 2.4Ti (1)
+ 1.76Nb + 1.22Ta − 0.226Ni − 0.177Co
WhenCreq is higher than 17–18%, the precipitation tendency of the σ phase is high [23]. The Creq value of the investigated Super304H steel was 18.18%, showing that the precipitation of the σ phase is highly probable. However, the σ phase was not observed, even in the specimen aged for 5000 h, indicating that the σ phase in Super304H steel requires a longer time to precipitate. Vach et al. [24] reported that the σ phase was found in Super304H steel aged at 650 °C for 87600 h. Moreover, some investigations [25] showed that the σ phase preferentially nucleates in the carbon-depleted region around M23C6 carbides. Iseda et al. [21] observed the σ phase at grain boundaries in a HR3C steel kept at 700 °C and 69 MPa for 88362 h. Thus the precipitation behavior of the σ phase in Super304H steel during longer periods of aging at 650 °C or at higher temperatures should be studied.
Fig. 6. Calculated equilibrium phase of Super304H steel.
3.3. Properties of precipitates Fig. 6 shows the fractions of precipitates at the equilibrium state from 500 °C to 1300 °C in the Super304H steel, which was determined using Thermal-Calc. It can be seen that the σ phase and Cu-rich phase dissolved at approximately 740 °C and 860 °C, respectively, whereas the Z phase and the M23C6 carbides dissolved at 930 °C and 950 °C, respectively. For temperatures higher than 1100 °C, only the Nb(C,N) phase existed, which is consistent with the microstructure of the asreceived Super304H steel, where only Nb(C,N) particles were present. According to Fig. 6, M23C6 carbides as well as Cu-rich, σ, and Z phases existed in the Super304H steel at the 650 °C equilibrium. The amount of the σ phase was the largest (9.41%), followed by the M23C6 carbides (2.09%) and the Z phase (0.51%). It should be noted that at the 650 °C equilibrium state, no Nb(C,N) phase was present in the Super304H steel because it tended to transform to the more stable Nb-rich Z phase during the long-term aging process. The chemical compositions of precipitates at the 650 °C equilibrium state in the Super304H steel are listed in Table 2. Normalized without considering the content of carbon, the contents of Cr and Fe in the M23C6 carbides were 83.98% and 10.91%, respectively. Compared with the experimental results of aged specimens for different times (as listed in Table 1), it can be concluded that the Cr content in M23C6 increased slowly with increasing aging time, whereas the Fe content decreased gradually. Table 2 shows that the Z phase has a chemical formula of Nb (Cr,Fe)N. Table 1 shows that as the aging time increased from 500 h to
3.4. Mechanical properties Fig. 7 shows the hardness and impact toughness of specimens annealed with different aging times. During the early stage of the aging process, the hardness increased rapidly from 157HB for the as-received sample to 191HB after 500 h. For longer aging times, the hardness was nearly stable and reached 200HB after aging for 5000 h. The impact toughness dropped rapidly in the primary exposure period (0–500 h) but changed little in the later exposure period (500–5000 h). Zhu et al. [15] investigated the impact toughness evolution of commercial HR3C steel during aging at 650 °C. Those results are also shown in Fig. 7. The figure shows that the toughness of the HR3C steel dropped much more rapidly than that of the Super304H steel. After exposure for 5000 h, the impact toughness of the HR3C steel fell to only 25 J/cm2, but that of the Super304H steel remained at 160 J/cm2. The fractures of aged Super304H steel were analyzed by SEM and the images are shown in Fig. 8. The as-received sample presented a ductile fracture with large and deep plastic-deformed dimples and some precipitates distributed at the bottom of the dimples. After aging for
Table 2 Chemical compositions of precipitates in Super304H steel at 650 °C equilibrium (atomic %). Phase
Cr
Fe
Nb
N
C
Cu
Mn
Ni
Si
Mo
M23C6 Cu-rich phase Z phase σ phase
66.60 0.01 26.75 41.74
8.65 0.30 6.60 54.45
– – 32.90 –
– – 33.31 –
20.69 – – –
– 97.96 – –
0.13 0.24 – 0.37
0.18 1.39 – 1.73
– – – 1.69
3.74 0.09 0.45 0.02
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the bonding between M23C6 and the grain is weak, cracks may mainly be generated near the M23C6 particles. The increase in the number and decrease in the size of dimples are associated with the increasing amount of M23C6 during aging, revealing that the precipitation of M23C6 is closely related to the decline in impact toughness. Regarding the HR3C steel, Wang et al. [13] reported that the impact fracture was a brittle rock candy-like fracture after aging at 650 °C for 6000 h. Zieliński et al. [14] and Golański et al. [27] indicated the tendency of M23C6 carbides to coagulate during service and to form a continuous network at the grain boundaries, which causes a rapid decrease in impact toughness in HR3C steel. The decrease in impact toughness in the Super304H and HR3C steel during aging was mainly due to the precipitation of M23C6 at grain boundaries. The large difference in embrittlement during aging between these two types of steel may be because of the difference in grain size, which produces marked changes in the distribution and morphology of M23C6 in grain boundaries. This results in different fracture modes.
Fig. 7. Mechanical properties of Super304H and HR3C steels as a function of aging time.
4. Discussion 500 h, the number of dimples increased markedly but were smaller and shallower than those in the as-received sample. In addition, the number of precipitates presented in the dimples also increased. With increasing aging time, the size and depth of the dimples seemed to decrease, but the features of ductile fracture could still be observed in the specimen aged for 5000 h. To understand the correlation between the second phase and the impact fracture, the precipitates in the dimples were analyzed. Large fragmented precipitates with a diameter of 1–2 μm were found at the bottom of the dimples in the as-received specimen and were identified as primary Nb(C,N) (shown in Fig. 9a). These large, brittle precipitates resulted in stress concentration under impact loading and their cracking or detachment from the matrix resulted in lower energy being required for crack initiation and expansion. In the specimen aged for 5000 h, precipitates at the dimples were identified as M23C6 carbides by EDS (shown in Fig. 9b). During aging, secondary Nb(C,N), the Cu-rich phase, and the Z phase precipitated inside grains. These precipitates were very fine and scattered, which could result in precipitation strengthening and may not impair the ductility. In general, the precipitation of M23C6 carbides leads to a depletion of C and Cr atoms in the surrounding matrix [26]. As the strength of the depletion area is relatively low and
4.1. Precipitation kinetics of M23C6 According to the analysis above, the precipitation of M23C6 carbides at grain boundaries is the main reason for the decrease in impact toughness in Super304H steel during aging. To investigate the evolution of impact toughness after long-term exposure, analyzing the precipitation kinetics of M23C6 in Super304H steel is necessary. The precipitation of M23C6 is a process that involves the diffusion and bonding of C and Cr atoms [28,29], which nucleate preferentially at positions with high energy and defects in the matrix, such as grain boundaries, vacancies, and dislocations [30]. As the atomic number of the Nb atom is much higher than that of Cr, the Nb(C,N) particles are white and bright, whereas M23C6 particles are dark gray in the back-scattered electron (BSE) images obtained through SEM. Nano-sized particles such as the Cu-rich phase, Z phase, and so on are difficult to observe in the BSE images. Thus, the M23C6 carbide can be distinguished and its volume fraction can be calculated based on the BSE images. The measured volume fractions of M23C6 in the as-received specimens and those aged for 500 h, 3000 h, and 5000 h are 0, 0.97%, 1.58%, and 1.84%, respectively. The precipitation kinetics of M23C6 is described by the Johnson-
Fig. 8. Fractography of impact samples. (a) as-received and aged at 650 °C for (b) 500 h, (c) 3000 h, and (d) 5000 h. 243
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Fig. 9. Precipitates on the fracture surface of Super304H steel. (a) As-received and (b) aged at 650 °C for 5000 h.
that mainly affected the mechanical properties of the Super304H steel. During the primary stage of aging (0–500 h), the precipitation of large amounts of M23C6 and nano-sized secondary Nb(C,N) particles caused the release of C and N atoms from the matrix (consistent with the right shift of the γ-Fe diffraction peaks), thus weakening the solution strength. However, the Cu-rich phase and secondary Nb(C,N) particles effectively pinned dislocations, thereby improving the precipitation strengthening effectively. Thus, the hardness of the specimen aged for 500 h was significantly higher than that of the as-received specimen. From 500 h to 5000 h, the Cu-rich phase and M23C6 continued to precipitate and coarsen. The nano-sized Cu-rich phase was coherent with the γ-Fe matrix, and its growth intensified the distortion of the lattice, offsetting the reduction of lattice distortion caused by the precipitation of C atoms. Hence, the diffraction peaks of γ-Fe remained nearly unchanged, which was in agreement with the stable hardness values of the Super304H steel during aging from 500 h to 5000 h. The decrease in toughness in the Super304H steel was mainly the result of the precipitation of M23C6. The impact toughness ak and volume fraction of M23C6 (V ) were fitted, as shown in Fig. 11. The function of this fit is shown by the following equation, with an R2 value of 0.99.
Mehl-Avrami equation [31]:
f (t ) = 1 −
exp (−kt n )
(2)
where f (t ) is the transformation fraction of a new phase at a certain temperature after a certain time t , n is the Avrami index, and k is the constant of crystallization rate. Here, f (t ) can also be defined by the ratio between the volume fraction of M23C6 (V (t )) after a certain time t , and the maximum volume fraction of M23C6 (Vm) when fully precipitated. Thus, the following equation can be written based on the J-MA equation:
f (t ) =
V (t ) = 1 − exp (−kt n ) Vm
(3)
where Vm is 2.09%, as calculated by Thermal-Calc at the 650 °C equilibrium state of the Super304H steel. Taking the volume fraction of M23C6 measured at 500 h, 3000 h, and 5000 h as V (t ) , three groups of data were obtained. Based on these data, the value of n and k were calculated to be 0.554 and 0.021, respectively. The J-M-A curve, as shown in Fig. 10 was also obtained. It can be seen that the M23C6 carbides began to precipitate quickly at 50–100 h and the volume fraction of M23C6 at 5000 h was close to 90% of the saturated precipitation content. The M23C6 in grain boundaries at 5000 h widened clearly as compared with that at 500 h (Fig. 4), which coincides with the increased amount of M23C6.
ak = 1/(0.0047 + 0.0012V − 0.0003V 2)
(4)
Fig. 11 shows that the decrease in impact toughness was in good agreement with the increase in the volume fraction of M23C6 in Super304H steel during aging. For a comparison, assuming that the decline of impact toughness of HR3C steel is only affected by M23C6, the values for the HR3C steel was calculated based on data from the literature [13]. These are also shown in Fig. 11. For the same amount of M23C6, the decline in toughness in the Super304H steel was lower than that in the HR3C steel, revealing
4.2. Effect of M23C6 precipitation on impact toughness The grain size in the Super304H steel was stable during aging (Fig. 1). Thus, the increase in hardness and decrease in toughness were mainly attributed to the evolution of precipitates. As the primary Nb (C,N) phase was stable and the Z phase was not significantly present, it was the secondary Nb(C,N) phase, Cu-rich phase, and M23C6 particles
Fig. 11. Effect of the amount of M23C6 phase on Charpy impact value (ak ) in Super304H steel.
Fig. 10. Precipitation kinetics curves of M23C6 phase in Super304H steel. 244
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that the difference in the aging embrittlement between these two types of austenitic steels derived from their different grain sizes. The Super304H steel had relatively fine grains (grain size of 9), whereas the HR3C steel had much coarser grains (grain size of 4) [14]. For the same amount of M23C6, the finer the grains were, the more difficult it was to form a continuous distribution of M23C6 at grain boundaries. As shown in Fig. 3d, although the content of M23C6 in the Super304H steel after 5000 h of aging approximated the amount of saturated precipitation, its distribution in the grain boundaries was discontinuous and appeared as scattered granules, unlike the network distribution of M23C6 in the HR3C steel. Thus, the negative influence of M23C6 precipitation on the toughness was insignificant. Based on the above analysis, the fine grains resulted in a considerably reduced aging embrittlement for the Super304H steel. It should also be noted that no σ phase was observed during aging at 650 °C for 5000 h. However, according to the thermodynamic calculation by Thermal-Calc, the σ phase could potentially precipitate in the Super304H steel. The precipitation of the σ phase at grain boundaries was found to be seriously harmful to the toughness [18]. Therefore, studying the precipitation behavior of the σ phase in the Super304H steel in future works is necessary.
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5. Conclusion
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[8]
[9]
[10]
[11]
[12]
[13]
[14]
1. During aging at 650 °C for different times up to 5000 h, the Super304H steel retained a microstructure with fine grains, consisting of a Cu-rich phase, secondary Nb(C,N), as well as a few Z phase particles that precipitated inside the grains and M23C6 carbides that precipitated at the grain boundaries. 2. The Cu-rich phase and secondary Nb(C,N) particles precipitated in specimens aged for 500 h. These particles retained their small size with prolonged aging time, which resulted in a marked increase in hardness. 3. M23C6 carbides precipitated rapidly during aging, and their amounts approached 90% of the saturated content in the specimen aged for 5000 h. The precipitation of M23C6 in the Super304H steel led to a clear reduction in impact toughness. However, the embrittlement in Super304H steel was much lower than that in HR3C steel. This was because of the discontinuous distribution of M23C6 at the grain boundaries due to the finer grain size in the Super304H steel.
[17]
[18]
[19]
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Acknowledgments
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The authors would like to express their gratitude for projects supported by the National Natural Science Foundation of China (51574181; 51374153).
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