2019 Volume 60 Issue 5 Pages 824-829
The age-hardening behavior of the Fe–Ni-based alloy HR6W was investigated in the temperature range between 973 K and 1073 K. A two-step increase of hardness was detected for the alloy at every aging temperature; the first increase of hardness results from the precipitation of M23C6 carbides, and the second increase corresponds to precipitation of the C14–Fe2W Laves phase. The time–temperature–precipitation diagram for the alloy was established on the basis of the results of hardness measurements and microstructure observations, where the precipitation of the C14–Fe2W Laves phase was slower than that of the M23C6 carbides by three orders of magnitude and the nose temperature of the Laves phase was greater than 1073 K. The M23C6 carbides precipitated with a plate-like morphology along grain boundaries at the early stage of aging, followed by the precipitation of the C14–Fe2W Laves phase with a granular morphology with increasing aging time. The M23C6 carbides and C14–Fe2W Laves phase are aligned under the stress condition because of their precipitation on the dislocations introduced during creep deformation.
This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 83 (2019) 30–35.
Fig. 6 TTP diagram of HR6W, showing the precipitation start time of the M23C6 and Laves phases in the matrix phase.
Society is demanding improvements in the efficiency of coal-fired power generation facilities to reduce the environmental burden by reducing the facilities’ carbon dioxide emissions.1) Enhancement of the energy conversion efficiency in a power plant is achieved by increasing steam temperature and pressure in the boiler.2) In an advanced ultra-supercritical (A-USC) thermal power plant, in which the steam temperature is set above 973 K, the ferritic and/or austenitic heat-resistant steels that have been conventionally used for boiler piping are not well suited from the viewpoints of high-temperature creep and oxidation characteristics.3) As candidate materials for A-USC boiler piping, Ni-based superalloys4,5) and Fe–Ni-based alloys6–8) have attracted attention in recent years. In such a social background, the low-cost Fe–Ni-based HR6W (Ni–23Cr–7W–25Fe) alloy with excellent workability has been developed.9,10) The high-temperature mechanical and oxidation properties of HR6W have been investigated to determine the suitability of this alloy for high-temperature structural components.11,12)
Thermodynamic calculations have indicated that the M23C6 carbide and C14–Fe2W Laves phase are the main equilibrium phases of HR6W in the temperature range between 973 K and 1073 K, which is the envisaged operating temperature of the alloy.13–15) However, the precipitation kinetics of both phases in this temperature range has not yet been fully clarified. This study aims to clarify the precipitation behavior of the Fe–Ni-based alloy HR6W during aging treatment in the temperature range between 973 K and 1073 K, with emphasis on the following three points. First, the precipitation kinetics of the M23C6 carbide and C14–Fe2W Laves phase is clarified by examining the age-hardening behavior and microstructure evolution during the aging treatment and the time–temperature–precipitation (TTP) diagram is established for HR6W in this temperature range. Second, the microstructure evolution of grain boundaries during high-temperature exposure is clarified for HR6W through observation of the precipitation microstructure near the grain boundaries. Third, the effect of applied stress on the precipitation microstructure is clarified for HR6W through observation of the microstructure of creep-tested specimen.
The alloy used in this study was the Fe–Ni-based alloy HR6W; its composition is shown in the Table 1. This alloy contains Cr, Fe and W as the main components in amounts of 23.23, 23.00, and 7.81 mass%, respectively, together with small amounts of Ti and Nb below 0.2 mass%. The C concentration of the alloy is 0.08 mass%, and its Ni content is evaluated at approximately 44 mass%. The alloy ingots of 50 kg were produced via vacuum induction melting and then hot-forged and hot-rolled at temperatures between 1173 K and 1423 K to form 15-mm-thick billets. The produced billets were subjected to a solution treatment at 1493 K for 0.5 h, followed by water quenching; the aging treatment was then carried out at 973–1073 K for 1–10000 h. After this solution treatment, the grain size of the γ matrix was approximately 120 µm. The hardness measurements were conducted using a micro-Vickers hardness tester; the load was set at 9.8 N, and the holding time was constant at 10 s. Care was taken to position the indenter of the hardness tester within γ grains of the alloy specimen for the hardness measurements. Furthermore, a full-sized creep-test piece was cut from the solution-treated alloy. The creep test was carried out at 973 K under an applied stress of 120 MPa, followed by interruption at 2872 h, and the interrupted test piece was used as a received material.
The microstructure of specimens was observed by field-emission scanning electron microscopy (FE-SEM), which was performed on a cross section parallel to the forging and rolling direction of the billets. The embedded samples were subjected to standard mechanical polishing, followed by electrolytic etching using a supersaturated chromic acid phosphate solution in a hot-water bath; the current was 40 mA, and the etching time was 30 s. In the FE-SEM observations, secondary electron imaging (SEI) was conducted at an accelerating voltage of 15.0–20.0 kV. For microstructure observations by transmission electron microscopy (TEM), thin films cut from cubic test pieces were shaped into disk-like samples with a diameter of 3 mm and a thickness of 150 µm by mechanical polishing. These samples were electrolytically polished using a standard twin-jet polisher and a solution of methanol and perchloric acid (9:1); the polishing temperature was 243 K, and the polishing current was 20 mA. The perforated foils were examined using a transmission electron microscope equipped with a double-tilt goniometer stage; the microscope was operated at 200 kV.
Hardness was measured with a micro-Vickers hardness tester to clarify the age-hardening behavior of HR6W. The age-hardening curves of HR6W at 973 K and 1073 K are shown in Fig. 1. Note that each plot is the average of five hardness data, and the distribution of the five data is indicated by an error bar. The hardness of the as-solution-treated alloy is Hv 158. At 973 K, the hardness begins to increase between 3 h and 10 h. After reaching the maximum value of Hv 205 at 300 h, the hardness slightly decreases with increasing aging time to Hv 200 at 1000 h. When the aging time is further increased, the hardness increases again to Hv 213 at 10000 h.
Age-hardening behavior of HR6W at 973 K and 1073 K.
At 1073 K, the hardness begins to increase before 1 h; after reaching the maximum value of Hv 187 at 4 h, the hardness decreases with increasing aging time, becoming Hv 175 at 30 h. As the aging time increases, the hardness increases continuously again, reaching Hv 203 at 3000 h. As previously described, in the age-hardening curves for HR6W, a two-step increase of hardness is observed at both 973 K and 1073 K. With increasing aging temperature, the hardness at both the first and second hardness peaks decreases and the aging time at which the hardness begins to increase becomes shorter.
3.2 Precipitation microstructure within grainsThe increase of hardness in the age-hardening curves is generally accepted to correspond to the onset of the precipitation of secondary phases.16) The SEM SEI of the aged alloy at 1073 K for 3 h, corresponding to the first hardness peak in the age-hardening curve at 1073 K, is shown in Fig. 2. Fine granular particles smaller than 100 nm precipitate uniformly within the γ grains. The SEM SEI of the aged alloy at 1073 K for 3000 h, which corresponds to the second hardness peak at 1073 K, is shown in Fig. 3. The size of the granular particles within γ grains becomes approximately 100 nm; in addition, rod-like precipitates with a length of approximately 1–3 µm are detected. As previously described, during the aging treatment of HR6W at 1073 K, fine granular particles first precipitate at short aging times. The size of the granular particles increases with increasing aging time, and the rod-like precipitates appear at longer aging times.
SEM SEI of HR6W aged at 1073 K for 3 h.
SEM SEI of HR6W aged at 1073 K for 3000 h.
The fine granular precipitates and rod-like precipitates observed in the aged alloys were identified from the selected-area diffraction patterns (SADPs) obtained by TEM. The TEM dark-field image (DFI) of the alloy aged at 1023 K for 10000 h is shown in Fig. 4, together with its SADP with the incident beam direction of B = [011], in which the fine granular precipitates are fully coarsened. The granular precipitates exhibit a rectangular shape with a length of approximately 100 nm. In addition, the following characteristics are obtained from the superlattice spots of SADP derived from the granular precipitates. First, the interface between the granular precipitates and the γ matrix is located on the {111} plane of the γ matrix. Second, the SADP derived from the granular precipitates is similar to that derived from the γ matrix, which indicates that the granular precipitates have a cubic crystal structure and that the {111} plane of the granular precipitates is parallel to {111}γ. Third, the superlattice reflections derived from the granular precipitates are positioned at one-third of the fundamental reflections of the γ matrix, which indicates that the lattice constant of the granular precipitates is three times larger than that of the γ matrix. These three aforementioned characteristics are consistent with those of Ni-based superalloys with the precipitation of M23C6 carbides.17,18) From this result, the fine granular precipitates are deductively identified as M23C6 carbides and the following orientation relationship is satisfied between the M23C6 carbides and γ matrix in HR6W: ($1\bar{1}1$)M23C6//($1\bar{1}1$)γ and ($\bar{1}\bar{1}1$)M23C6//($\bar{1}\bar{1}1$)γ.19)
TEM DFI of M23C6 carbides observed in HR6W aged at 1023 K for 10000 h, taken with B = [011].
The TEM BFI of a rod-like precipitate is shown in Fig. 5, together with its SADP. To minimize the effect of the γ matrix on the SADP, we conducted the observation by focusing on the rod-like precipitate positioned at the edge of the perforation of the TEM film sample (Fig. 5(a)). The incident beam direction is B = [001] (Fig. 5(b)), and the indices of some spots in the SADP are presented in Fig. 5(c). In the TEM BFI, stacking faults are observed at a high density within the rod-like precipitate. Correspondingly, the superlattice reflections in the SADP derived from the rod-like precipitates are streaked perpendicular to the stacking fault planes. This characteristic feature is consistent with that observed in the high-chromium ferritic steels with the precipitation of the C14–Fe2W Laves phase.20,21) On the basis of this result, the rod-like precipitate is deduced as the C14 Laves phase with a hexagonal crystal structure, and the following orientation relationship is satisfied between the C14 Laves phase and the γ matrix in HR6W: (110)γ//(0001)Laves and [100]γ//[$1\bar{1}00$]Laves.
TEM BFI (a) and corresponding SADP (b) of the Laves phase observed in HR6W aged at 1023 K for 10000 h, taken with B = [001] and g = 200. A schematic of the SADP, together with the indices, is shown in (c).
The time required to initiate precipitation of the M23C6 carbide and C14–Fe2W Laves phase was identified from the age-hardening curves at 973 K and 1073 K in Fig. 1; the results are summarized as the TTP diagram in Fig. 6, together with the results obtained at 1023 K. The results of the microstructure observation of the precipitates under each aging condition are included in the figure as square symbols. The precipitation of the M23C6 carbide starts before 1 h at 1073 K and 1023 K. At 973 K, the precipitation of the M23C6 carbide is delayed to between 3 h and 10 h. By contrast, the precipitation of the C14–Fe2W Laves phase occurs between 1000 h and 3000 h at 973 K. Precipitation occurs earlier with increasing aging temperature, starting between 30 h and 100 h at 1073 K. Because the aging conditions at 1073 K for 100 h and 1023 K for 300 h are positioned immediately after the precipitation onset of the C14–Fe2W Laves phase, we inferred that the Laves phase was difficult to clearly detect in the aged alloys.
TTP diagram of HR6W, showing the precipitation start time of the M23C6 and Laves phases in the matrix phase.
On the basis of the aforementioned results, the precipitation onset of the C14–Fe2W Laves phase is approximately three orders of magnitude slower than that of the M23C6 carbide in the temperature range between 973 K and 1073 K and that the nose temperature of the Laves phase is greater than 1073 K. Note that a substantial increase of hardness is detected between 1000 h and 3000 h in the age-hardening curve at 1073 K, as shown in Fig. 1, indicative of the precipitation of the third phase for HR6W other than the M23C6 carbide and C14–Fe2W Laves phase. However, the third precipitation phase was not observed in the alloy specimen aged at 1073 K for 3000 h (Fig. 3). Clarifying the nose temperature of the C curves for HR6W, which represents the onset of the precipitation of the M23C6 carbide and the C14–Fe2W Laves phase, is beyond the scope of the present work.
3.3 Precipitation microstructure around grain boundariesIn the previous sections, the Vickers hardness and precipitation microstructure within the γ grains of HR6W subjected to the aging treatment were investigated and the TTP diagram of the alloy was established. In this section, the microstructure evolution around grain boundaries of HR6W during the aging treatment is investigated. The SEM SEI near the grain boundaries of the aged alloy at 1073 K for 3 h, which corresponds to the first peak of hardness in the age-hardening curve, is shown in Fig. 7. The precipitates with a plate-like morphology and a length of approximately 2 µm are placed along the grain boundaries, and fine granular M23C6 carbides with a size of approximately 100 nm precipitate at a high density near the grain boundaries. For the aged alloy, the precipitation density of M23C6 carbides around grain boundaries is higher than that within γ grains (Fig. 2). The M23C6 carbides are considered to precipitate earlier near the grain boundaries compared with their precipitation within γ grains. The TEM DFI of the M23C6 carbides precipitated at the grain boundaries for the aged alloy at 973 K for 10 h is shown in Fig. 8 with the incident beam direction of B = [001]. We observed that the M23C6 carbides on the grain boundaries precipitate with a plate-like morphology so as to cover the grain boundaries.
SEM SEI around grain boundaries of HR6W aged at 1073 K for 3 h.
TEM DFI of M23C6 precipitates at grain boundaries of HR6W aged at 973 K for 10 h, taken with B = [001].
We investigated whether the M23C6 carbides that precipitated on the grain boundary possess an orientation relationship with any of the γ grains. The TEM BFI of the alloy aged at 973 K for 30 h is shown in Fig. 9, where the SADPs were collected in the immediate vicinity of the grain boundary for both γ grains. Figure 9(a) and (b) are the same field of view, and the γ grains on the upper side (a) and the lower side (b) across the grain boundary were photographed with each crystal zone axis. Note that the incident beam direction on the upper grain (Fig. 9(a)) is B = [001] and that the reciprocal lattice vector is g = 200, whereas B = [$\bar{1}10$] and g = [111] are set for the lower grain (Fig. 9(b)). In both γ grains, numerous dislocations and M23C6 carbides are observed near the grain boundary. From the SADPs, the M23C6 carbides are identified as precipitating coherently with respect to the γ matrix for each grain. In some cases, the M23C6 carbides extend from the grain boundary into the γ grain interior, as indicated by red arrowheads, whereas the M23C6 carbides with a plate-like morphology precipitate along the grain boundary in the other cases, as indicated by blue arrowheads. From the aforementioned results, the M23C6 carbides on the grain boundary are inferred to precipitate coherently with either γ grain and to grow toward the γ grain interior while retaining the coherency.
TEM BFI of M23C6 precipitates around grain boundaries of HR6W aged at 973 K for 30 h, taken with B = [001], g = 200 (a) and B = [$\bar{1}10$], g = 111 (b). Red arrowheads indicate the M23C6 precipitates within grains; blue arrowheads are the M23C6 precipitates at grain boundaries.
The SEM SEI near the grain boundaries for HR6W aged at 1073 K for 3000 h, which is the highest aging temperature and the longest aging time in this study, is shown in Fig. 10. Part of the precipitates along the grain boundaries retains a plate-like morphology, as indicated by yellow arrowheads, whereas the other part of the precipitates exhibits a granular morphology (red arrowheads). That is, for HR6W, the M23C6 carbides precipitate with a plate-like morphology along the grain boundaries at the early stage of aging and the C14–Fe2W Laves phase precipitates adjacently to the M23C6 carbides with increasing aging time. At that time, the M23C6 carbides retain their plate-like morphology, whereas the C14–Fe2W Laves phase exhibits a granular morphology; thus, plate-like and granular precipitates are located alternatingly on the grain boundaries.
SEM SEI around grain boundaries of HR6W aged at 1073 K for 3000 h. Grain-boundary precipitates consist of a plate-like region (yellow arrowheads) and a granular-shaped region (red arrowheads).
To clarify the effect of stress application on the precipitation microstructure of HR6W, the microstructure of the alloy crept at 973 K under a stress of 120 MPa was observed. The SEM SEI for the HR6W creep-interrupted at 2872 h is shown in Fig. 11. Note that, according to the TTP diagram of HR6W shown in Fig. 6, both the M23C6 carbide and the C14–Fe2W Laves phase are presumed to precipitate during aging at 973 K for 2872 h. Figure 11 shows that the precipitation occurs in parallel lines, which indicates that the dislocations introduced during the creep deformation act as effective precipitation sites. The SEM SEI obtained by magnifying the precipitates in a linear fashion is shown in Fig. 12, which reveals that two kinds of precipitates; fine granular precipitates with a length of approximately 100 nm and coarse rod-like precipitates with a length of approximately 600 nm, are adjacent to each other. From the size and morphology of each precipitate, we inferred that the fine granular precipitate is the M23C6 carbide and the coarse rod-like precipitate is the C14–Fe2W Laves phase.
SEM SEI of HR6W creep-interrupted at 973 K/120 MPa for 2872 h.
Magnified view of intragranular precipitates observed in HR6W creep-interrupted at 973 K/120 MPa for 2872 h.
In the creep-interrupted specimen, the following process is considered to be responsible for the M23C6 carbides and C14–Fe2W Laves phase precipitating linearly adjacent to each other. First, dislocations are introduced into γ grain interiors during the creep deformation. Because the introduced dislocations are generated from identical dislocation sources, they are located on the same slip plane and are parallel to each other. Second, the M23C6 carbides precipitate on the introduced dislocations within γ grains and the dislocation is pinned by the M23C6 carbides. Third, the total amount of W included in the M23C6 carbides increases as the M23C6 carbides grow and coarsen with time because the W content of the M23C6 carbides is relatively high for HR6W.13) Fourth, the C14–Fe2W Laves phase precipitates on the pinned dislocations and grows by incorporating W from the M23C6 carbides. Via the four aforementioned steps, the fine M23C6 carbides and coarse C14–Fe2W Laves phase precipitate linearly adjacent to each other in the creep-deformed alloy.
Hardness measurements and microstructure observations were performed for the HR6W Fe–Ni-based alloy at temperatures between 973 K and 1073 K, and the effect of applied stress on the microstructure evolution of the alloy was also investigated. The results obtained in this study are summarized as follows.
The alloy samples used in this study were manufactured and provided by Nippon Steel & Sumitomo Metal Co. The authors would like to thank Dr. Keiji Kubushiro of IHI Co. for the extensive cooperation in promoting this research and Prof. Susumu Onaka and Prof. Yoshisato Kimura of Tokyo Institute of Technology and Prof. Yoji Miyajima of Kanazawa University for kind assistance with the microstructure observation using electron microscopy.
