Reductions of Intergranular Corrosion Resistance and Wear Resistance in a Ni-Cr-Mo-Based Superalloy by Aging-Treatment-Induced Precipitation

: Ni-Cr-Mo-based superalloy is widely used as a key component in many critical environ-ments. To ensure that the manufacturing process does not impact the long-term service performance of these components, the aging precipitation behavior at different temperatures and its effect on intergranular corrosion (IGC) resistance and wear resistance of a Ni-Cr-Mo-based C276 superalloy were investigated. The equilibrium phase diagram was calculated ﬁrst using thermodynamic software to conﬁrm the potential phases. Carbides of M 6 C were found to be formed at grain boundaries after aging at 800–850 ◦ C for short-term treatment. The other two phases ( µ phase and P phase) indicated in the phase diagram were not observed for the samples after aging treatment up to 15 h. Furthermore, double loop electrochemical potentiokinetic reactivation (DL-EPR) tests were conducted to examine the IGC resistance. The degree of sensitization increased with the aging time and severe corrosion was found to occur at grain boundaries. For the ﬁrst time, the inﬂuence of aging treatment on the wear behavior of this superalloy has been speciﬁcally studied. Concerning the hot processing of Ni-Cr-Mo-based C276 superalloy, these results indicate the importance of avoiding high-temperature heat treatment for long periods.


Introduction
Superalloys usually preserve excellent high-temperature performance, resulting in them being widely used in many industries with intense conditions, such as aerospace, chemical, nuclear, etc. C276 superalloy is one of the most promising superalloys belonging to the Ni-Cr-Mo-based family, with good corrosion resistance, high temperature strength, and creep resistance [1][2][3][4]. This superalloy was first developed by dramatically reducing the content of carbon and optimizing the alloying elements [5]. The outstanding advantage of this superalloy is that it can be used directly in a welding state without severe intergranular corrosion (IGC) attack [6,7].
Due to the excellent properties of C276 superalloy, it is an ideal candidate material of some key components in the nuclear industry [8][9][10]. The material was designed to be used in temperatures below 650 • C. However, during the manufacture process, C276 superalloy needs to undergo heat treatments at elevated temperatures for a certain period of time in some cases. Recently, a new technology, named hot vacuum bulging technology, was developed to manufacture a rotor can, which is a key component in nuclear reactor coolant pumps, using C276 superalloy [11,12]. The hot vacuum bulging processing is conducted at 800 • C for 2 h to realize exact forming through creep deformation. Creep age forming (CAF) is an effective forming technology for aluminum alloys, and potentially has applications with superalloys [13]. Therefore, the effect of high-temperature aging on performance during the manufacture process (which is much higher than the service temperature) becomes a main concern for C276 superalloy.
C276 superalloy is used in solution state with single austenite phase. There is no need to apply aging treatment to this superalloy to optimize the precipitation of γ' and/or γ", such as the precipitation-strengthened superalloys [14,15]. Previous investigations into the aging treatment of C276 superalloy focused mainly on testing the stability of performance during long-term service. Liu et al. found that massive block µ phase and M 6 C carbides precipitated in the C276 superalloy after aging at 700 • C for more than 360 h, and the precipitates grew along with the aging time [16]. The tensile test at 700 • C indicated that the aged samples maintained the tensile strength with a small decrease in plasticity until 2160 h. The fracture of the tested sample still presented ductile morphology. Akhter et al. tested the effect of aging treatment on the hardness and impact energy of C276 superalloy [17]. The authors found that Mo-rich µ phase precipitated after aging at 850 • C for 48 h, associated with a dramatic decrease in impact energy. Raghavan et al. have characterized the three distinct second phases, µ phase, M 6 C carbides, and P phase, in the long-term aged C276 superalloy using transmission electron microscopy (TEM) [5]. The chemical compositions are remarkably close for the three kinds of precipitates. The authors defined the precipitates using a transmission electron microscope.
However, the influence of short-term aging treatment on the performance of this kind of superalloy has been given little attention. There are indications that short-term aging at high temperature can also induce precipitation in other Ni-Cr-Mo-based superalloys [18,19]. A similar phenomenon was also observed in C276 superalloy when it was aging treated at 800 • C for a relatively short time [20]. Furthermore, the nuclear reactor coolant pump rotor serves as a protective barrier for the rotor against cyclic erosion of the coolant. The corrosion resistance and wear resistance of C276 superalloy are key utilized material properties for C276 superalloy. For example, there were clear, synergistic effects between corrosion and wear in this alloy under seawater conditions, leading to corrosion-induced wear and wear-induced-corrosion in tribocorrosion processes [21][22][23]. According to Hashim and Duraiselvam [24], the tribological properties of C276 superalloy could be improved by laser surface treatment without sacrificing its corrosion properties. Yilbas and Ali also studied the effect of laser treatment on C276 alloy [25]; they found that laser treatment was effective in increasing the surface microhardness and decreasing the friction coefficient. Nevertheless, it is worth noting that no study investigating the effect of aging treatment on the tribological behavior of C276 superalloy has been reported. Our previous paper [20] revealed that small amounts of precipitation resulted in the deterioration of corrosion resistance, but the relationship between precipitation, corrosion behavior, and wear resistance of C276 superalloy is yet to be elucidated. Therefore, the precipitation behavior at different aging treatment temperatures and its influence on the corrosion and wear resistances are highly deserving of an investigation.
The present study is an extension from our previous report [20]. In this work, the equilibrium phase diagram for C276 superalloy was calculated first. Then, based on the thermal dynamic calculation results, aging treatments at two different temperatures were specifically conducted. The aging-treatment-induced precipitations were carefully characterized by SEM and TEM. Moreover, the corrosion resistance, mechanical property, and wear behavior of aging-treated C276 superalloy was evaluated by double-loop electrochemical potentiokinetic reaction (DL-EPR) test, microhardness measurement and ball-on-plate sliding wear test, respectively.

Materials and Methods
The tested C276 superalloy is a 0.5 mm thick commercial sheet, which is in solution state with equiaxed austenite grain structure. The chemical composition of the material is: 0.004 C, 0.01 Si, 16 [20]. Based on the chemical compositions, thermodynamic software Thermal-Calc (Thermo-Calc Software, Canonsburg, RA, USA) was used to calculate the equilibrium phase diagram. In order to make a comparison with the previous paper [20], the sheets were aging treated in a vacuum furnace at 800 • C and 850 • C for 1-15 h, followed by water quenching.
The samples used for microstructure analysis were cut from these aging-treated sheets. After mechanical polishing, the samples were light etched in aquaregia (HCl: HNO 3 = 3:1) for 1-3 min and then examined with a field emission scanning electron microscope (Zeiss Supra 55, Zeiss, Germany). For TEM observation, disks 3 mm in diameter and 50 µm in thickness were electrochemically polished at −30 • C in a 5% perchloric acid ethanol solution and examined in JEM-2100F (Jeol Ltd., Japan, JEM-2100F).
The DL-EPR test was performed to investigate the corrosion resistance of the aged samples. Square samples measuring 15 × 15 mm were cut and the sample surfaces were mechanically polished. The conventional solution (0.01 mol/L KSCN and 0.5 mol/L H 2 SO 4 ) was found to not be applicable for distinguishing the current peaks in the electrochemical tests. In this work, we introduced a new solution of 6 mol/L HCl to investigate the IGC resistance of C276 superalloy. A saturated calomel electrode (SCE) is served as the reference electrode. The scan started from the potential of −0.3 V SCE to 0.7 V SCE with a rate of 1 mV/s, then reversed.
The microhardness was examined using the MHV−2.0 Vickers (Shenzhen Shunhua Instrument Equipment Co., Ltd., Shenzhen, China) hardness tester with a load of 0.25 N. The influence of aging treatment on the wear resistance of C276 superalloy was studied in an acceleration condition using a ball-on-plate wear testing rig. The detailed information about the wear testing rig can be found in the study by Li et al. [26]. Prior to the wear tests, C276 superalloy plates were ground and polished to a final surface roughness of less than 100 nm [27]. Al 2 O 3 balls with a diameter of 7 mm were selected as the counter-body. Wear tests were performed at room temperature and repeated at least three times for each material. The applied normal load on the ball was 10 N and the sliding speed was 0.92 m/min. The length of the wear stroke was 12 mm and the total reciprocating wear test duration was 30 min. During each test, the friction coefficient was automatically calculated and recorded. The worn surface profile was then examined by an optical 3D surface profiler (Super View W1 type, which is manufactured by Shenzhen Zhongtu Instrument Co., Ltd., Shenzhen, China).

Thermodynamic Calculation Results
The equilibrium phase diagram based on the chemical compositions was calculated using thermodynamic software Thermal-Calc. The molar fractions (NP) of phase constituents varying with the temperature is shown in Figure 1. It infers that the µ phase, P phase, and M 6 C are supposed to precipitate from the matrix at high temperatures; µ phase can precipitate at temperatures lower than 772 • C, P phase can precipitate in the temperature range of 598-1098 • C, M 6 C can form at temperatures lower than 1111 • C, and the content increases with the decreasing temperature. The maximum content for the precipitation of M 6 C is 0.14 mol%. The equilibrium chemical compositions of the precipitations are listed in Table 1. The precipitations are mainly composed of Cr, Mo, W, and Ni.  It should be noted that the equilibrium phase diagram is calculated based on the thermodynamic analysis. It needs time for the diffusion of alloy atoms to form the precipitates. Three kinds of precipitates have been reported to be formed in Ni-Cr-Mo-based superalloy [5,28]. However, the aging treatments were conducted at 650-900 °C for 1000 h in the investigation of Raghavan et al. [5]. Akhter et al. heated the samples up to 240 h to study the effects of aging treatment on the hardness and impact property [17]. They also mentioned that μ phase cannot form during the aging treatment at 650 °C within 500 h, while M6C may be produced after 10 h of aging.

Precipitate Observations
In our previous paper [20], the precipitation behavior of C276 superalloy during aging at 800 °C was studied. In order to understand the influence of aging temperature, the precipitation behavior of C276 superalloy during aging at 850 °C was studied and compared with the results of aging at 800 °C. Figure 2 shows the backscattered electron (BSE) images of the aging-treated samples with different aging temperatures. One can determine that some second phases precipitated along the austenite grain boundaries. By comparison, the number of precipitates increases with the increase in aging time and aging temperature. Furthermore, the morphologies of the precipitates seem very similar for all the samples. The amounts of precipitates were observed to be very low for the samples aged within 4 h. Energy dispersive spectrometer (EDS) analysis was used to detect the compositions of the precipitates, as shown in Figure 3. It can be seen that the precipitates are rich in Mo and W as compared to the base material. The compositions are similar to the thermodynamic calculation results as listed in Table 1. However, it is hard to distinguish the precipitates based on the chemical compositions.
TEM was used to further confirm the precipitates, as shown in Figure 4. The precipitates were formed at grain boundaries with a twinning boundary. The diffraction pattern indicates that the precipitation is a face-centered cubic (FCC) structure with a lattice constant of a = 1.11 nm. M6C is reported to have a cubic crystal structure with a = 1.08-1.12 nm [5]. μ phase (A6B7 type) has a hexagonal close packed structure with a = 0.476-0.479  It should be noted that the equilibrium phase diagram is calculated based on the thermodynamic analysis. It needs time for the diffusion of alloy atoms to form the precipitates. Three kinds of precipitates have been reported to be formed in Ni-Cr-Mo-based superalloy [5,28]. However, the aging treatments were conducted at 650-900 • C for 1000 h in the investigation of Raghavan et al. [5]. Akhter et al. heated the samples up to 240 h to study the effects of aging treatment on the hardness and impact property [17]. They also mentioned that µ phase cannot form during the aging treatment at 650 • C within 500 h, while M 6 C may be produced after 10 h of aging.

Precipitate Observations
In our previous paper [20], the precipitation behavior of C276 superalloy during aging at 800 • C was studied. In order to understand the influence of aging temperature, the precipitation behavior of C276 superalloy during aging at 850 • C was studied and compared with the results of aging at 800 • C. Figure 2 shows the backscattered electron (BSE) images of the aging-treated samples with different aging temperatures. One can determine that some second phases precipitated along the austenite grain boundaries. By comparison, the number of precipitates increases with the increase in aging time and aging temperature. Furthermore, the morphologies of the precipitates seem very similar for all the samples. The amounts of precipitates were observed to be very low for the samples aged within 4 h. Energy dispersive spectrometer (EDS) analysis was used to detect the compositions of the precipitates, as shown in Figure 3. It can be seen that the precipitates are rich in Mo and W as compared to the base material. The compositions are similar to the thermodynamic calculation results as listed in Table 1. However, it is hard to distinguish the precipitates based on the chemical compositions.
cipitate is supposed to be M6C, which is consistent with the previous report [20]. No other types of precipitates were observed in TEM analysis. The chemical composition analysis result corresponds to the thermodynamic calculation result in Section 3.1. Therefore, both the thermodynamic calculation and experimental investigation confirm the precipitation of M6C carbides in the aging-treated C276 superalloy. Increasing the aging temperature to 850 °C does not introduce other precipitates during aging.   1, μ phase is supposed to precipitate at a temperature below 772 °C. Therefore, the precipitate is supposed to be M6C, which is consistent with the previous report [20]. No other types of precipitates were observed in TEM analysis. The chemical composition analysis result corresponds to the thermodynamic calculation result in Section 3.1. Therefore, both the thermodynamic calculation and experimental investigation confirm the precipitation of M6C carbides in the aging-treated C276 superalloy. Increasing the aging temperature to 850 °C does not introduce other precipitates during aging.   TEM was used to further confirm the precipitates, as shown in Figure 4. The precipitates were formed at grain boundaries with a twinning boundary. The diffraction pattern indicates that the precipitation is a face-centered cubic (FCC) structure with a lattice constant of a = 1.11 nm. M 6 C is reported to have a cubic crystal structure with a = 1.08-1.12 nm [5]. µ phase (A 6 B 7 type) has a hexagonal close packed structure with a = 0.476-0.479 nm and c = 2.57-2.59 nm, while P has a tetragonal crystal structure [5,16]. Moreover, Raghavan et al. reported that µ phase and M 6 C were the main precipitates for C276 superalloy when aging at 650 • C-900 • C up to 1000 h [5]. Based on the phase diagram in Figure 1, µ phase is supposed to precipitate at a temperature below 772 • C. Therefore, the precipitate is supposed to be M 6 C, which is consistent with the previous report [20]. No other types of precipitates were observed in TEM analysis. The chemical composition analysis result corresponds to the thermodynamic calculation result in Section 3.1. Therefore, both the thermodynamic calculation and experimental investigation confirm the precipitation of M 6 C carbides in the aging-treated C276 superalloy. Increasing the aging temperature to 850 • C does not introduce other precipitates during aging. Metals 2021, 11, x FOR PEER REVIEW 6 of 12

IGC Behavior
DL-EPR test was employed to evaluate the degrees of sensitization of the samples aging treated at 850 °C. The representative DL-EPR curves are shown in Figure 5 [20]. The values in Table 2 indicate that this solution can reflect the degree of sensitization.

IGC Behavior
DL-EPR test was employed to evaluate the degrees of sensitization of the samples aging treated at 850 • C. The representative DL-EPR curves are shown in Figure 5. The current density peaks in the anodic scan loop (I a ) and reverse scan loop (I r ) were determined by the curves. Then, the ratio of I r /I a can be obtained to evaluate the sensitization of IGC. Table 2 lists the determined values for the parameters I a , I r , and I r /I a . The I a values varied little with the aging time. However, the I r increased along with the aging. Since a new solution was used, the degree of sensitization was used to compare the relative IGC resistance of the samples. It was found that the degree of sensitization of the tested C276 superalloy increases with the aging time. This corresponds to the precipitate observation results, in which the number of precipitates increases with the progression of the aging treatment. Here, a new solution was used to evaluate the IGC behavior of C276 superalloy, which resulted in a large difference between these results and those reported previously [20]. The values in Table 2 indicate that this solution can reflect the degree of sensitization.   Figure 6 shows the microstructure of samples after DL-EPR tests. The as-received sample shows homogeneous etched morphology after testing. However, for the aged sample, severe corrosion at grain boundaries can be easily defined, as shown in Figure 6b. This indicates that C276 superalloy is sensitive to IGC attack after aging treatment, which agrees well with the DL-EPR curve analysis results.   Figure 6 shows the microstructure of samples after DL-EPR tests. The as-received sample shows homogeneous etched morphology after testing. However, for the aged sample, severe corrosion at grain boundaries can be easily defined, as shown in Figure 6b. This indicates that C276 superalloy is sensitive to IGC attack after aging treatment, which agrees well with the DL-EPR curve analysis results. Figure 5. DL-EPR curves of C276 superalloy after aging treated at 850 °C for different times. Figure 6 shows the microstructure of samples after DL-EPR tests. The as-received sample shows homogeneous etched morphology after testing. However, for the aged sample, severe corrosion at grain boundaries can be easily defined, as shown in Figure 6b. This indicates that C276 superalloy is sensitive to IGC attack after aging treatment, which agrees well with the DL-EPR curve analysis results. The precipitates of M6C carbides after aging treatment are rich in Mo and W as detected by EDS analysis. This results in the formation of Mo-and W-depleted regions adjacent to the grain boundaries [29]. These regions are supposed to be easily attacked by corrosion during etching [30][31][32]. This implies that the precipitation of carbides during aging at high temperatures for higher than 4 h is detrimental to the IGC resistance of C276 superalloy. Moreover, the grain boundaries show different corrosion resistances in Figure  6b. Stratulat and Luo et al. have noticed this phenomenon in stainless steel and aluminum alloy [33,34]. They showed that the grain boundaries with high stored energy, which is related to the grain misorientation angle and dislocation density, were easily IGC attacked.

Microhardness and Wear Behavior
The microhardness of C276 superalloy after aging treatment at 850 °C was measured with the results displayed in Figure 7. The microhardness is shown to increase from about 330 to 360 HV after aging treatment for 4 h, and then declines very slightly with further increasing the aging time. The aging treatment at 800 °C led to very similar microhardness The precipitates of M 6 C carbides after aging treatment are rich in Mo and W as detected by EDS analysis. This results in the formation of Mo-and W-depleted regions adjacent to the grain boundaries [29]. These regions are supposed to be easily attacked by corrosion during etching [30][31][32]. This implies that the precipitation of carbides during aging at high temperatures for higher than 4 h is detrimental to the IGC resistance of C276 superalloy. Moreover, the grain boundaries show different corrosion resistances in Figure 6b. Stratulat and Luo et al. have noticed this phenomenon in stainless steel and aluminum alloy [33,34]. They showed that the grain boundaries with high stored energy, which is related to the grain misorientation angle and dislocation density, were easily IGC attacked.

Microhardness and Wear Behavior
The microhardness of C276 superalloy after aging treatment at 850 • C was measured with the results displayed in Figure 7. The microhardness is shown to increase from about 330 to 360 HV after aging treatment for 4 h, and then declines very slightly with further increasing the aging time. The aging treatment at 800 • C led to very similar microhardness in the alloy with the aging treatment at 850 • C. Coupling the analysis with the precipitation observations in Figure 3, the precipitation of M 6 C carbides is supposed to increase the microhardness of the aged C276 superalloy. Such a phenomenon of age hardening has also been found in a C276 alloy coatings by Mulligan and co-authors [35], who interpreted the increase in hardness to the formation of precipitates rich in Mo and Cr. It is worth noting that precipitation mainly occurs at grain boundaries and there is a slight grain growth during aging treatment, as evident from Figure 2; thus, a very small reduction in the microhardness is observed when the alloy is aging treated after 15 h. cient of about 0.95 is achieved in the sample that has the greatest microhardness of about 360 HV corresponding to the aging treatment at 850 °C for 4 h. These increased friction coefficients in Figure 8 should be attributed to the formation of hard precipitates as displayed in Figure 2. It is obvious that the presence of precipitates not only results in the increase in friction coefficient, but also leads to the fluctuations of friction coefficient during steady state wear process in Figure 8a.  In previous studies [21][22][23][24][25][26][27][36][37][38], it has been reported that the wear behavior of a metallic material can be affected by many factors, such as its microstructure, phase constitution, mechanical properties, and the sliding wear conditions (including normal load, velocity, and atmosphere). Based on the precipitation results in Figure 2 and microhardness results in Figure 7, the C276 superalloy samples after aging treatment at 850 • C for 4 and 15 h were specifically selected for further wear property evaluations. Figure 8a shows the typical friction coefficient evolution histories during sliding wear as a function of time. The results reveal that the friction coefficient increases significantly during the early stage of sliding wear and then changes very slightly, indicating a relatively steady-state wear process. Figure 8b shows the average friction coefficient as a function of microhardness. Of note, the average friction coefficient in this paper is calculated by averaging the results obtained during the steady state wear stage from all three repeated wear tests for each condition. It can be seen from the figure that the average friction coefficient of C276 superalloy without aging treatment is about 0.6. The value is consistent with the typical friction coefficients for dry sliding wear of metallic materials [37]. It is interesting to observe that the friction coefficient increases first after aging treatment for 4 h and then drops with further increasing the aging time. The friction coefficient evolution tendency is consistent with that of microhardness. As shown in Figure 8b, the average friction coefficient of C276 superalloy increases when its microhardness is increased. The maximum friction coefficient of about 0.95 is achieved in the sample that has the greatest microhardness of about 360 HV corresponding to the aging treatment at 850 • C for 4 h. These increased friction coefficients in Figure 8 should be attributed to the formation of hard precipitates as displayed in Figure 2. It is obvious that the presence of precipitates not only results in the increase in friction coefficient, but also leads to the fluctuations of friction coefficient during steady state wear process in Figure 8a. In order to reveal the influence of aging treatment on the wear resistance and wear mechanisms of C276 superalloy, the 3D wear track surface morphologies after wear test are characterized with the results shown in Figure 9. Obvious furrow marks associated In order to reveal the influence of aging treatment on the wear resistance and wear mechanisms of C276 superalloy, the 3D wear track surface morphologies after wear test are characterized with the results shown in Figure 9. Obvious furrow marks associated with wear debris can be found in all the samples as labelled in the figures, indicating the severe abrasion wear. The samples with and without aging treatment show a great difference in the morphologies of furrow. Narrow and shallow wear marks develop in the C276 superalloy without aging treatment. On the contrary, remarkably wider and deeper grooves are found in the aged treated samples. To further analyze the surface morphologies after wear, the surface profiles were detected and drawn together in Figure 10. The wear track width and depth are about 373 µm and 5.5 µm, respectively, for the as-received alloy while increasing to about 1490-1530 µm and 21-22 µm, respectively, for the aged samples. It is evident that aging treatment of C276 superalloy leads to much larger wear tracks, indicating the obvious reductions of its wear resistance.
Although the microhardness increases after aging treatment, the wear resistance was reduced. Based on the observations on the wear tracks in Figure 9, it is estimated that the M 6 C carbides may fall off from the material and increase the wear volume in the form of a third body during abrasive friction. This effect is different from the fine precipitates in the matrix, in which the fine scale precipitates can improve the wear resistance in the form of precipitation strengthening [38][39][40]. The increase in microhardness of C276 superalloy after aging treatment is caused by the formation of M 6 C carbides. However, these precipitates mainly formed at grain boundaries, as confirmed in Figures 2-4. The precipitation at grain boundaries could deteriorate the mechanical properties of the matrix. Therefore, to preserve the good wear resistance of the C276 superalloy, the selection of aging treatment temperature and time should be very carefully selected. It is worth noting that this study is the first report considering the influence of aging treatment on the wear behavior of C276 superalloy, and a more comprehensive investigation on the mechanisms and sliding-wear-induced microstructure changes will be reported in our future research.    Based on the above analysis, it is found that the precipitation, corrosion resistance, microhardness, friction coefficient, and wear resistance of C276 superalloy can be greatly affected by aging treatment. Particularly, both the corrosion resistance and wear resistance deteriorated obviously for the C276 superalloy after aging longer than 4 h. Since this superalloy is mainly used in nuclear industry which faces cyclic load and erosion of