Interpretation for the fast sigma phase precipitation in the high intensity shot peened nanocrystallined Super304H stainless steel

Accepted Manuscript Full Length Article Interpretation for the fast sigma phase precipitation in the high intensity shot peened nanocrystallined Super304H stainless steel Qingwen Zhou, Ruikun Wang, Zhijun Zheng, Yan Gao PII: DOI: Reference:

S0169-4332(18)32232-3 https://doi.org/10.1016/j.apsusc.2018.08.097 APSUSC 40134

To appear in:

Applied Surface Science

Received Date: Revised Date: Accepted Date:

8 March 2018 20 July 2018 9 August 2018

Please cite this article as: Q. Zhou, R. Wang, Z. Zheng, Y. Gao, Interpretation for the fast sigma phase precipitation in the high intensity shot peened nanocrystallined Super304H stainless steel, Applied Surface Science (2018), doi: https://doi.org/10.1016/j.apsusc.2018.08.097

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Interpretation for the fast sigma phase precipitation in the high intensity shot peened nanocrystallined Super304H stainless steel

Qingwen Zhoua, Ruikun Wangb, Zhijun Zhengc, Yan Gaoa,* a

School of Materials Science and Engineering, South China University of Technology, Guangzhou 510641, P R China

b

School of Mechanical and Electric Engineering，Guangzhou University, Guangzhou 510006,P R China c

School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou 510641, P R China

Further information related to this article please send to Prof. Y. Gao School of Materials Science and Engineering South China University of Technology Guangzhou 510641, PR China E-mail: [email protected] Tel: 86 20-87114219， Fax: 86 20-87111950

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Abstract: Surface shot peening (SP) with an intensity of 0.5 MPa/12 min plus 650 oC aging was performed on the Super304H austenitic stainless steel with an attempt to decrease its intergranular corrosion susceptibility more effectively. It was found that the local stress/strain concentration with higher energy at some typical sites caused by over saturated cold deformation induced the early nucleation of sigma phase and the enhanced chromium diffusion in nanocrystallized microstructure promoted its fast growth in the material. The critical shot peening time to trigger fast precipitation of sigma phase is between 8~12 min under 0.5 MPa for the Super304H steel. The occurrence of massive sigma phase during aging brought a positive move of the reactivation potential of chromium depleted zones around sigma phase compared to M 23C6. Keywords: Super304H; Nanocrystallization; Resensitization; Sigma phase.

1. Introduction Super304H austenitic stainless steel is widely used in thermal power plants, especially ultra supercritical units, as reheater and superheater tubes due to its excellent high temperature performance [1][2][3][4]. However, austenitic stainless steel is always susceptible to intergranular corrosion due to the formation of chromium depleted zones caused by M23C6 precipitation at grain boundaries at sensitization temperatures (450-800℃) [5][6][7]. For the Super304H steel with higher carbon content to ensure its high temperature strength, more M23C6 is likely to precipitate under the service process which will bring higher intergranular corrosion sensitivity and unexpected risk to the stable operation of the units [7][8][9]. There are two possible routes to solve the problem of high intergranular corrosion susceptibility of stainless steels, composition optimization [10][11][12] and technology design and

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control [6][13][14][15][16][17]. In composition optimization, carbide forming elements may be added to play a certain role in carbon sequestration and inhibit the precipitation of M23C6 [7]. As for process technology control, appropriate high temperature softening [17] and solution treatment may reduce the precipitation of M 23C6, but it is impossible to fundamentally avoid the formation of chromium depleted zones through these measures. In our previous study [14], it has been found that a 650℃/24h aging treatment of the nanostructured Super304H steel by 0.5 MPa/8 min shot peening can ensure a low intergranular corrosion susceptibility at subsequent service stage. In the actual application, however, the desensitization process is expected to be done in a shorter time to ensure its feasibility of low cost and acceptable oxidation. When the shot peening parameter was increased into 0.5 MPa/20 min, however, a significant amount of sigma phase appeared during aging and brought a deterioration of intergranular corrosion susceptibility (IGCS) [18]. The possible reason for the fast precipitation of sigma phase may be related to the higher degree of cold deformation. The condition for the formation of sigma phase has been discussed a lot and many factors have been found to influence it [19]. For example the grain size and grain shape affects the density of nucleation sites [20][21] and a higher crystallographic misorientation between the austenite phase and the ferrite phase favors sigma precipitation [22]. Increasing the degree of cold deformation will increase the degree of both residual stress and refined grain size [20] and subsequently the precipitation behavior of the material such as sigma phase will be affected obviously. Of all kinds of precipitates in stainless steels, sigma phase is one of the most harmful kinds which will cause great damage to the performance of the materials. A research of

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long-term evolution of 304H austenitic stainless steel showed that σ phase nucleated in the δ ferrite after aging for 18 years and grew up subsequently in the austenite [23]. The δ ferrite may transform directly to sigma phase in austenite stainless steel by short-range diffusion transformation, due to their close chemical composition and the higher diffusion rates of alloying elements in δ ferrite [24][25]. In ferritic/martensitic (F/M) steels, δ-ferrite may induce the sigma phase formation rapidly at δ-ferrite/martensite boundaries [26]. In some cases, carbide particles can act as sources of chromium for σ-phase formation [27]. In another way, sigma phase may also transform directly from deformed austenite matrix by nucleating at the boundaries between the recrystallized and deformed grains due to the minimized energy barrier at the recrystallizing interfaces resulting from the rearrangement of dislocations and sufficient mass transport of the solute atoms for growth by short circuit diffusion [28][29]. But up to now, there is no systematic research of sigma phase precipitation behavior in the novel Super 304H stainless steel, since it is relatively newly developed and the available long term creep specimen do not show the existence of sigma phase [30]. Moreover, there are more alloying elements in this material and its sigma phase precipitation should be of difference from the conventional 304 stainless steel and is worth investigating. In this study, an attempt was made to adjust the shot peening parameter to 0.5 MPa/12 min, which was between 0.5 MPa/8 min [14] and 0.5 MPa/20 min [18], with the aim to get a shorter desensitization time without the occurrence of sigma phase, so as to endow the desensitization process with a better industry application prospect. In the end, woefully, sigma phase was found largely in the 0.5 MPa/12 min shot peened Super304H specimen after aging at 650℃ for 48h,

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therefore the precipitation characters of sigma phase and its relationship with cold deformation degree in the shot peened Super304H austenitic stainless steel were explored in details by comparing with our previous work [14][18].

2. Experimental 2.1 Materials and treatments The Super304H stainless steel tube was used as the experimental material with the composition shown in Table 1. The tube was solution treated at 1150℃ for 30min and water quenched rapidly. Then the material was processed into 50×30×4mm samples by wire-electrode cutting and ground and polished for shot peening treatment. Shot peening was conducted in an air blasting machine (AMS-1212P). Stainless steel balls with a diameter of 0.5 mm were used with a 100% surface coverage. Parameters of shot peening and solid solution treatment are given in Table 2. The sample after shot peening was cut into small pieces of 10×10×4mm by wire-electrode cutting for subsequent aging. The aging temperature was selected as 650℃ which is the highest service temperature of Super304H stainless steel as superheater tubes in ultra supercritical units. The selected aging times were 0.1 h, 1 h, 2 h, 10 h, 24 h 48 h, 96 h, 144 h, respectively. 2.2 Microstructural examination Quantitative measurement of the precipitates in nano-scaled specimen was performed by using X-ray diffraction (XRD, Philips X’pert MPD) with Cu-Kα radiation at 40 kV and 40 mA. The XRD tests were repeated twice for each sample to ensure the reproducibility. All the solution treated specimens (S1) were electrolytically polished in 10% perchloric acid

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(H2Cl2O4 2H2O) at 1 A/cm2 for 60s to reveal the precipitates morphology. For the shot peened specimens (S2), the deformed microstructures were not revealed clearly by this electrolytical polishing so the S2 specimens were mechanically ground and polished using diamond paste to a finish of 2.5 μm and then etched using Vilella’s reagent (1 g picric acid+ 5 ml HCl + 100 ml ethanol) to reveal both the microstructure and precipitates morphology. A scanning electron microscope (FESEM Zeiss Supra-40) was used to observe the morphology using secondary electron imaging. Image-Pro-Plus 6.0 (IPP 6.0) software was used in this study to statistically measure the volume fraction (V(t)) of sigma phase based on the stereology theory. All the experimental data of V(t) was based on the average of three measured results by IPP 6.0 software from three random SE fields. 2.3 DL-EPR test The DL-EPR test was carried out to measure the intergranular corrosion sensibility of the aged specimens. The specimens to be tested were ground on emery papers and polished with a diamond paste of 2.5 μm and then dried with ethanol. The DL-EPR test was carried in a solution of 0.01 M H2SO4 + 20 ppm KSCN at 35℃ on AutoLab PGSTAT30, using a three electrode system with graphite electrode as counter electrode and saturated calomel electrode (SCE) as reference electrode. The samples were kept in the solution for 30–45 min to obtain a stable open circuit potential (OCP). It was then scanned from -50 mV to +300 mV versus SCE and then scanned back to the starting potential at a scan rate of 1.67 mV/s. The ratio of the peak area measured in the reactivation (backward) scan to that in the activation (forward) scan was taken as the degree of sensitization (DOS) value. For the morphology observation after DL-EPR test, all the specimens

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were ultrasonic cleaned and then examined under scanning electron microscope to reveal the detailed states of intergranular corrosion. 2.4 Field emission transmission electron microscope TEM investigation was performed using a JEOL-2100 field emission transmission electron microscopy at an accelerating voltage of 200kV. The thin foil specimens were prepared from a slice cutting from the shot peened layer with the thickness of about 100 μm. Then this slice was produced into small disks with a diameter of 3 mm and mechanically ground to less than 50 μm to make sure the left part was the shot peened nanocrystallined surface. Finally the thin foiled specimen were further polished by an TenuPol-5 twin-jet machine in a solution of 10% perchloric acid (HClO4) + 90% ethanol (C2H5OH) at -20℃~ -15℃. 2.5 Electron backscattered diffraction examination Electron backscatter diffraction (EBSD) has the potential to study grain boundary character distribution (GBCD) and phase distribution of materials. The specimens for EBSD analysis were electrolytically polished in 10% perchloric acid (HClO4) + 90% ethanol (C2H5OH) solution. Typically 10 mm-10 mm areas were polished for EBSD scans, with beam and video conditions keeping identical. EBSD measurements were made at an operating voltage of 20keV at a step size (distance between the two measuring points) of about 0.1 μm. Channel 5 software was used for scan data analysis. The CSL values (Σ) is correlated with boundary energy. Boundaries with Σ value less than 29 are considered special because of their low energy, based on the Brandon's criterion Δθ = 15/Σ1/2 [31]. In this study, low energy boundaries (special boundaries) were taken to be those with Σ value lower than 11 since the Σ values of most boundaries are lower than 11 after

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aging. All the boundaries with Σ equal to 1 were ignored from this analysis.

3. Results 3.1 Microstructure of the shot peened specimen Fig.1 shows the TEM morphology at the top layer (less than 50μm) in shot peened S2. Nearequiaxed nanograins and fine twin bands are introduced into the material. Deformed twins and their intersections are verified to be the domination mechanism for grain refining in Super304h stainless steel [14]. It is worthwhile to note that some areas like grain boundaries and twins intersections are more severely deformed compared with its surroundings. For example, the deformation is more severe on the deformed twinning intersections and strain-induced martensite is introduced at this sites, as shown in Fig.2. 3.2 Precipitation evolution with aging treatment The precipitation behaviors by SEM of the solution annealed (S1) and shot peened (S2) specimens aged at 650℃ for various times are shown in Fig. 3 and Fig. 4. For the solution treatment specimen (S1), Nb(C,N) is the main precipitate (the white spherical precipitates in Fig.3(a)). During aging process, the M23C6 carbides (shown with the yellow arrows) begin to precipitate at grain boundaries, as shown in Fig. 3(b-d). The shot peened specimen (S2) has more nucleation sites like the deformed bands in Fig. 4 (a) and (b) and Nb(C,N) is the main precipitate before aging [30] (shown with the white arrows). During aging M23C6 carbides precipitate not only at grain boundaries but also at some deformed bands and sub grains inside the austenite grains, which are quite small in nano-scale size and need to be identified by TEM. The TEM observation of M23C6 is shown in Fig.5. We can see a precipitate at deformed twins in Fig. 5(a)

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and (b). The selected area diffraction pattern in Fig. 5(a) and the dark field image in Fig. 5(c) confirm the precipitate to be M23C6. In Fig. 4(c) and (d), micron sized precipitates of a new phase are found to distribute uniformly after 24 h aging, indicating a fast growth rate of this phase which was identified as sigma phase by the XRD analysis in Fig. 6. However the smaller particle size of sigma phase in the specimen aging for 24 h shown in Fig. 4(c) may explain why its XRD diffraction peaks in Fig. 6 are not obvious compared to those after 48h aging. Besides the strain-induced martensite is found to transform gradually to austenite during aging and its peaks disappear completely after 10 h of aging in the XRD spectrums in Fig. 6. The peaks of sigma phase (between 45~50 degrees) begin to appear after aging of more than 24 h, which will be further confirmed by the DL-EPR tests. From these results it is known that the fast precipitation of sigma phase has not been suppressed during aging when the shot peening parameter is adjusted to 0.5 MPa/12 min. Consequently its intergranullar corrosion would be likely affected to a certain extent. 3.3 Evaluation of degree of sensitization by DL-EPR test Fig. 7 displays the DL-EPR curves of S1 and S2 aged at 650◦ C for various times, with typical activation current (Ia) and reactivation current (Ir) peaks. In Fig. 7(a) of S1, only one reactivation current peak is observed, relating to the dissolution of Cr-depleted zones near M23C6 precipitates at grain boundaries [32], and the Ir value increases gradually with aging time, indicating a higher intergranular corrosion sensitivity due to the gradual increase of M23C6 precipitates with longer aging time (as shown in Fig. 3 (a) - (d)). In Fig. 7(b) of S2, there appears a second reactivation

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peak, different from the law in S1, which is attributed to the replenishment of chromium-depleted zones near M23C6 carbides in martensite during the reactivation scan, termed as “martensite-induced sensitization” [33] and this phenomenon also occurred in our former research on different shot peened Super304H steel [14]. The quick chromium diffusion in martensite results in a shallow depletion of chromium-depleted zones, which leads to an active move of the potential for IGC attack, presented as a second reactivation peak in the DL-EPR curve [33]. Fig. 7(c) is the partial enlargement of Fig. 7(b) to show the reactivation peaks more clearly. Since a large part of the two reactivation peaks in Fig. 7(c) is overlapped, the DOS results would be inaccurate if they were calculated by peak height (current density). Therefore in this paper, the DOS was determined by using the area (quantity of electricity) of the reactivation current peak and activation current peak in the following equation [34]: DOS = (Qr/Qa) ×100% Where Qr is the charge for the reactivation scan and Qa is the charge for the activation scan. Fig. 8 presents the DOS values calculated from Fig. 7 as a function of aging time for the two specimens aged at 650℃. The DOS of solution-treated specimen (S1) increases monotonically with aging time, indicating that S1 specimen is at the stage of sensitization even aged to the prolonged time of 144h. However, the DOS of S2 specimen increases first and reaches a maximum of 29.2% after aging at 650℃ for 0.1 h, and then decreases rapidly to reach a rather low value at 10 h. This quick healing process is similar but quicker to our previous work [14]. When the specimens are aged for more than 24 h, however, the DOS values increase again, indicating a second sensitization.

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Chromium depleted zones in austenite and the second reactivation peak is believed to be martensite-induced sensitization. For the reactivation potentials of the first peak, the figures are between -200~ -225 mV during 0.1h to 10h, but change to -181~ -190 mV during 24 h to 144 h. This phenomenon indicates a change of chromium depleted zones because the main chromium-rich precipitates has changed from M23C6 to sigma phase after 24h of aging. The dissolution potential of chromium depleted zones near the sigma phase (24 h~144 h) has a positive move (about 30 mV) comparing to M23C6 (0.1 h~10 h), indicating its easier dissolving. However the chromium content of sigma phase (27~28 wt.%) is much lower than that of M23C6 [18] and the reason for this positive move of the potential is believed to be the extremely large amount of sigma phase (Fig. 4(c) and (d)), which cause a more severe chromium depletion near the sigma phase. When the aging time is less than 10 h, the intergranular corrosion sensibility is caused by the M23C6 in austenite as well as in strain-induced martensite. The enhanced chromium diffusion in both nanocrystalline austenite and martensite heals the chromium depleted zones quickly, which accounts for the rapid drop of DOS values from 29.2% to 4.7% in 10 hours (0.1 h to 10 h). But during longer time aging from 24 h to 144 h, the strain-induced martensite transfers to austenite and most nanostructures also grow up, hence chromium diffusion rate will slow down. Under this situation, the chromium-depleted zones caused by the later sigma phase precipitation are difficult to heal and consequently the DOS values keep increasing gradually from 3.8% to 15.7%. Fig.9 shows the SEM corrosion morphologies after the DL-EPR tests of the two specimens. The corrosion attacks of S1 in Fig. 9(a)-(d) are mainly along grain boundaries, which corresponds

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to the dissolution of the chromium depleted zones near the M23C6 carbides along grain boundaries. In Fig. 9(e) to (g) of S2, corrosion attacks are not only along grain boundaries but also along twin boundaries and subgrain boundaries in the austenite grains. Due to the fast healing of chromium-depleted zones by enhanced chromium diffusion in deformed structure, the S2 specimen has already reached a desensitization state at 24 h of aging as shown in Fig. 9(g). But the XRD and DL-EPR results have shown that the main precipitates change to sigma phase at 24 h of aging. At this time sigma phase has not grown up large enough so its influence on the DOS is not obvious and the morphology after DL-EPR test can not reveal its existence. In Fig. 9(h), however, the main attacks become the spalling of the large amount of micrometer sized sigma phases. The chromium-depleted zones near the sigma phase formed after about 24 h of aging are difficult to heal due to the slowing down of chromium diffusion in the recrystallized microstructure. Hence there is no fast decrease of the second sensitization in the later aging ( 48 h to 144 h) compared with the early stage aging (0 h to 10 h). 3.4 The distribution of sigma phase Fig.10 shows the EBSD results of S2 sample after aging for 144 h to characterize sigma phase more directly and precisely. Fig. 10(a) is the original EBSD map and Fig. 10(b) is the partial magnification in Fig. 10(a) to reveal more detailed phase distribution. Fig. 10(c) is the fitting outcome for phase map and grain boundaries of Fig. 10(b) using Channel 5 software after a slight noise reduction treatment and Fig. 10(d) is the processed data for the distribution of CSL grain boundaries. The zero resolution areas in the EBSD results are the black areas in Fig. 10 (a) and (b) and the white areas in Fig. 10 (c), whose presence is attributed to residual stress but the resolution

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rate of this sample is over 80% so the result is relatively reliable. The fitting outcome in Fig. 10(c) is found to be in good consistency with the original data. From the results we can see that most sigma phases locate at grain boundaries especially at triple junctions and the size of most sigma phase is up to 1μm which indicates a rapid growth of this phase during aging. Although most of the Nb(C,N) particles locate uniformly [29] while most of the M23C6 carbides are at grain boundaries [8], some of them are found to be wrapped by sigma phase, as shown in Fig. 10(b) and (c), which indicates that these precipitates locate at the same area in the early stage of aging but afterwards sigma phases grow far quicker than other precipitates. The content of sigma phase (characterized as FeCr in EBSD) is 3.48% which is higher than those of Nb(C,N) and M23C6. But most zero resolution areas are near the sigma phase hence the real content of sigma phase may be even higher. What’s more, the stress concentration around sigma phase will make the phase boundaries high-energy as well as more sensitive to corrosion. From Fig. 10(d) we can see that most grain boundaries in the recrystallized austenite are low-energy and low-Σ coincidence site lattice (CSL). This microstructure has been proved to exhibit higher intergranular corrosion resistance [35]. If no sigma phase had appeared, the recrystallized microstructure of austenite would have performed a good intergranular corrosion resistance [14]. The chromium-rich sigma phase has very high interfacial energy and also results in chromium depletion zones around it [36], all of which will lead to the preferential corrosion along the phase boundaries. Combining with the result in Fig. 9(h), we can conclude that the phase boundaries between austenite and sigma phase is the preferential sites for intergranular corrosion and the appearance of sigma phase has deteriorated the intergranular corrosion resistance of shot

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peened Super304H stainless steel. 3.5 The relationship between sigma phase precipitation and the degree of shot peening Compared with the specimens without cold deformation [30] or with low intensity shot peening [14] where no sigma phase precipitated during long time aging, the shot peening intensity in this paper is greater and as a result there are more strain-induced martensite, dislocations and residual stress in this specimen. These differences may help promote the precipitation of sigma phase during aging. The microstructures of different shot peening specimens in our research after aging at 650℃ for long time (168 h) are summarized in Fig. 11, to show the effect of cold deformation degree on the precipitation of sigma phase. In Fig. 11(a) and (b), the recrystallized grains are large and some nano-scale precipitates distribute at grain boundaries and in grains. There are still some deformed twins in the grains shown in Fig. 11(a). Due to the large interval between 8-12 min, we further studied the SP process of 10 min under 0.5 MPa and its microstructure aging for 168 h is put in Fig. 11(c), where a limited number of micro-scaled sigma particles (marked by red arrows) are found. So the critical SP time of triggering sigma phase precipitation should be between 8-10 min. Considering the time control roughness of SP process, we choose 2 min as an interval so we may take 10 min under 0.5 MPa pressure as the critical time of triggering sigma phase precipitation during aging. To conclude, when the shot peening time increases to 10 min, the sigma phase suddenly appears and grows quickly to large size during 650oC/168 h aging, proving that the cold deformation degree is the main reason for the promoted precipitation of sigma phase. Fig. 12(a) gives a summary of the sigma phase content evolution during aging at 650℃ of

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Super304H steel with different shot peening time, with some data cited from our previous work [14][18]. We find out the different change trends of sigma phase content during aging in the specimens of different shot peening intensity. When the shot peening time is 3 min and 8 min, there is no sigma phase precipitation during aging. When the shot peening time increases to 10 min, the sigma phase suddenly appears, meaning that the critical shot peening time under 0.5 MPa to trigger sigma phase precipitation is 10 min. The content of sigma phase in SP 0.5 MPa/10 min specimen aging at 650oC for 168 h was calculated (3.02%) and is listed in Fig. 12(a). Because the sigma content is minute and its distribution is quite uneven in this specimen, its content evolution during aging has not been calculated due to calculation accuracy. Further increase of shot peening time to 12 min and 20 min brings an increase in both velocity and amount of sigma precipitates and a quicker reaching of stable state at 144h and 96 h respectively. This phenomenon indicates that larger cold deformation degree brings quicker sigma phase precipitation. Fig. 12(b) cites the results of average grain size and content of strain-induced martensite calculated from XRD results and Fig. 12(c) cites the depth of deformed layer in different SP specimens in our previous work [18], to show the effect of cold deformation degree on grain refinement, martensite formation and deformed layer. It is obvious that the increase of shot peening time brings the fast change of grain size, martensite content and deformation layer. From the former results [14], it is known that there is a faster increasing rate of intergranular corrosion susceptibility (DOS value) in specimens with smaller grain size since more grain boundaries offer more precipitation sites for M23C6 and bring its quicker precipitation. Moreover, the healing rate of the Cr-depleted zones is also quicker due to the massive diffusion channels caused by more

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grain boundaries in smaller grain size specimens. These influences of fine grains will bring a faster sensitization and desensitization of IGCS. However, after 8 min of shot peening, those deformation variables of grain size, martensite content and deformation layer all reach a relatively stable state. These phenomena indicate that the deforming energy before the critical point of 8 min brings about uniform deformation, and after 8 min of shot peening the deforming energy fails to distribute uniformly in the saturated uniform deformed layer but concentrate in some more deformable sites like triple junctions of grain boundaries, phase boundaries and strain-induced martensite, and the results of uniform deformation like grain size and deformation layer remain nearly stable. Fig.13 presents a schematic illustration of microstructure evolution with the increase of shot peening intensity. The stage before 8 min is the early stage of uniform nanocrystalline layer formation with the obvious increase of layer depth (Fig. 13(b)). Further increasing shot peening time only makes the deformation take place at some more deformable sites. As a result, those over deformed sites with a more severe stress concentration (like the red points in Fig. 13(c)) become the preferential sites for sigma phase precipitation. These sites have low energy barrier for sigma phase nucleation and also promote its growth because of the enhanced chromium diffusion in the over deformed microstructure, especially the strain-induced martensite. The mechanism of the fast sigma phase precipitation in the over deformed Super304H steel at the very early stage of aging needs further detailed study since there has been no research report to explain it up to now. To conclude, the extra high energy introduced by high intensity cold deformation greatly favors the nucleation and growth of sigma phase in the Super304H steel. 4. Discussion 16

Comparing with the 0.5 MPa/8 min specimen, the 0.5 MPa/12 min specimen reaches the desensitization state more quickly (650 ℃ /10 h), meaning that its intergranular corrosion sensitivity caused by the chromium depleted zones near M23C6 carbides has been healed in 10 h of aging at 650℃. Unfortunately chromium-rich sigma phase precipitates massively after long time aging (more than 24 h), which results in chromium depleted zones along its phase boundaries and brings about the second sensitization in the 0.5 MPa/12 min specimen. Since the fast diffusion channels in deformed microstructure gradually disappear during aging through stress recovering and recrystallization, the chromium depleted zones near sigma phase are not easy to be healed in the later aging so the degree of second sensitization keeps on increasing with aging time. Although the chromium ratio of sigma phase is much lower than that of M23C6, its massive content has induced an accumulated deeper chromium depletion for a higher IGCS and brought a positive move of its reactivation potential compared with that of Cr23C6 induced sensitization. However, fast sigma phase precipitation is not common in austenitic stainless steels when serving as boiler components. For the conventional 304H steel, sigma phase only appears after a very long time service（18 years）under high temperature service [23] and there is no sigma phase precipitation in the Super 304H steel in the long-term (4578h) creep test at 650℃ [30]. In the shot peened S2 specimens, more high energy sites are [37] introduced during shot peening, which favor the nucleation of M23C6 [14] as well as of sigma phase by high energy fluctuation and the chromium diffusion enhancement from austenite to this phase. Meanwhile, dislocation pile-up leads to the formation of strain-induced martensite, which also promotes chromium diffusion and may transfer to sigma phase more easily due to its smaller microstructure difference with sigma

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phase than austenite matrix. Researches have shown that the sigma phases nucleate at the interfaces between recrystallized matrix and deformed matrix by short circuit diffusion and then left behind in the recrystallized grains in the cold rolled Fe-10Cr-30Mn [28] [29]. However, the sigma phases locate mainly at triple junctions in our study. Obviously the mechanism of sigma phase precipitation in the shot peened Super304H stainless steel is different from its precipitation behavior in other deformed materials. From Fig. 12 and Fig. 13, we may speculate that the nucleation and fast growth of sigma phase is due to the uneven stress distribution after over saturated shot peening. When increasing the shot peening time, the evolution of nanocrystallization experiences two stages. The first stage is the uniform transfer from as-received matrix to deformed one where the grain refining degree, martensite content [38] as well as the depth of deformed layer increase in a quick rate with shot peening time. When the deformation reaches a saturate state (8 min/0.5 MPa in this research), further increase in shot peening time does not bring overall deformation but local severe stress and strain concentration in some sites, representing with a smooth line of deformation degree with shot peening time (Fig. 12(b) and (c)). At those sites, the uneven concentrated energy is high enough to overcome the energy barrier of sigma transition and triggers sigma phase nucleation at the early stage of aging before recrystallization. Meanwhile the deformed microstructure will promote the chromium diffusion which may explain the fast growth of sigma phase in high density shot peened Super 304H stainless steel. Up to now, there is no research concerning the fast sigma phase precipitation and its mechanism in this material after high intensity cold deformation and further work needs to be done to understand the stress concentration induced sigma phase precipitation.

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5. Conclusions The relationship between the second sensitization and sigma phase precipitation in the Super 304H stainless steel with high intensity shot peening is investigated. The shot peened specimens aged at 650℃ for different time were examined focusing mainly on the evolution of its intergranular corrosion sensitivity. The main results obtained are as follows. (1) The increase of shot peen time from 8 min to 12 min makes it possible to achieve desensitization state in a shorter time (10 h) at 650℃ but second sensitization presents after long time aging (more than 24 h) due to the appearance of sigma phase. The rapid growth of sigma phase brings an increase of IGCS as well as a positive move of the reactivation current peak. (2) The critical shot peening time to stimulate sigma phase precipitation is between 8~12 min under 0.5 MPa in the shot peened Super304H. For shot peening before 8 min, the austenite matrix transforms to uniform deformed layer but after 8 min, the deformation energy distributes unevenly at some easier deforming sites like triple junctions of grain boundaries, phase boundaries and strain-induced martensite and causes severe local stress /strain concentration. (3) The local stress concentration site with higher density of dislocations and vacancies after over saturated cold deformation favors the fast nucleation of sigma phase. Meanwhile the deformed microstructure promotes the chromium diffusion which accounts for the fast growth of sigma phase in the high density shot peened Super 304H steel. (4) Most of the grain boundaries in the recrystallized austenite are low Σ CSL boundaries. So the intergranular corrosion after long time aging is more prone to happen along the phase

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boundaries between austenite and sigma phases where there exist both chromium depletion and stress concentration.

Acknowledgement The authors would acknowledge greatly the financial support from the National Nature Science Foundation of China (51471072) and the Key Laboratory of Advanced Energy Storage Materials of Guangdong Province.

Reference [1] R. Viswanathan, J.S.A.J., Boiler materials for ultra-supercritical coal power plantssteam-side oxidation. Journal of Materials Engineering and Performance. 15(2006) 255-274. [2] Bakker, R.V.A.W., Materials for ultra supercritical coal power plants -boiler materials: Part 1. Journal of Materials Engineering and Performance. 10(2001) 81-95. [3] MV Kumar，V Balasubramanian，AG Rao, Hot tensile properties and strain hardening behaviour of Super 304HCu stainless steel, Journal of Materials Research & Technology. 6(2016) 116-122 [4] Xin-mei Li, Y.Z.Z.Z., Microstructure evolution of a novel super304h steel aged at high temperatures. Materials Transactions. 51(2010) 305-309. [5] A. Abou-Elazm, R.A.I.E., Correlation between the degree of sensitization and stress corrosion cracking susceptibility of type 304H stainless steel. Corrosion Science. 51(2009) 203-208. [6] P. Lin, G.P.U.E., Influence of grain boundary character distribution on sensitization and intergranular corrosion of alloy 600. Scripta Metallurgical Materialia. 33(1995) 1387-1392.

20

[7] Kenji Kaneko, T.F.K.Y., Formation of M23C6-type precipitates and chromium-depleted zones in austenite stainless steel. Scripta Materialia. 65(2011) 509-512. [8] Yan Gao, C.Z.X.X., Intergranular corrosion susceptibility of a novel Super304H stainless steel. Engineering Failure Analysis. 24(2012) 26-32. [9] Armijo J.S., Intergranular stress corrosion cracking of austenitic stainless steels in oxygenated high temperature water. Corrosion-Nace. 24(1968) 319-325. [10] Indrani Sen, E.A.N.S., Microstructure and mechanical properties of annealed SUS 304H austenitic stainless steel with copper. Materials Science and Engineering A. 528(2011) 4491-4499. [11] R. Peraldi, B.A.P., Effect of Cr and Ni contents on the oxidation behavior of ferritic and austenitic model alloys in air with water vapor. Oxidation of Metals. 61(2004) 463-483. [12] K. K. Alanemea, S.M.H.I., Effect of copper addition on the fracture and fatigue crack growth behavior of solution heat-treated SUS 304H austenitic steel. Materials Science and Engineering A. 527(2010) 4600-4604. [13] Sunil Kumar B., B.S.P.V., Methods for making alloy 600 resistant to sensitization and intergranular corrosion. Corrosion Science. 70(2013) 55-61. [14] Ruikun Wang, Z.Z. Q.Z., Effect of surface nanocrystallization on the sensitization and desensitization behavior of Super304H stainless steel. Corrosion Science. 111(2016) 728-741. [15] T. S. Wang, B.L.M.Z., Nanocrystallization and martensite formation in the surface layer of medium-manganese austenitic wear-resistant steel caused by shot peening. Materials Science

21

and Engineering A. 458(2007) 249–252. [16] M. A. Arafin, J.A.S., A new understanding of intergranular stress corrosion cracking resistance of pipeline steel through grain boundary character and crystallographic texture studies. Corrosion Science. 51(2009) 119–128. [17] Fei Xie, WH Zhu, ZJ Zheng, Yan Gao. Influence of high temperature softening and solid solution on intergranular corrosion susceptibility of Super304H stainless steel, Corrosion Engineering, Science and Technology. 52(2017) 302-310. [18] Ruikun Wang. Effect of surface grain refining by shot peening on the intergranular corrosion susceptibility and desensitization characteristics of Super304H stainless steel [D]. Guangzhou: South China University of Technology, 2016. [19] Lo K H, Shek C H, Lai J K L. Recent developments in stainless steels. Materials Science & Engineering R. 65(2009) 39-104. [20] Sara Bagherifard, S.S.I.F. Nanoscale surface modification of AISI 316L stainless steel by severe shot peening. Materials and Design. 102(2016) 68–77. [21] M. Schwind, J. Kallqvist, J.-O. Nilsson, J. Agren, H.-O. Andren. Evolution of Secondary Phases Formed upon Solidification of a Ni-Based Alloy. Acta Materialia. 48 (2000) 2473. [22] Y.S. Sato, H. Kokawa. The Mechanism of Grain Coarsening in Friction-Stir-Welded AA5083 after Heat Treatment. Scripta Materialia. 40 (1999) 659. [23] Young-Su Ji, Jihye Park. Long-term evolution of σ phase in 304H austenitic stainless steel: Experimental and computational investigation. Materials Characterization. 12(2017) 23-29 [24] C.C. Tseng, Y. Shen. Fracture and the Formation of Sigma Phase, M23C6, and Austenite from

22

Delta-Ferrite in an AISI 304L Stainless Steel. Metallurgical and Materials Transactions A. 25A(1994) 1147-1158 [25] Xavier Ledoux, Francois Buy. Kinetics of sigma phase precipitation in niobium-stabilized austenitic stainless steel and effect on the mechanical properties. Materials Science Forum. 783-786(2014) 848-853 [26] Yinzhong Shen, Xiaoling Zhou. Sigma phases in an 11%Cr ferritic/martensitic steel with the normalized and tempered condition. Materials Characterization. 122(2016) 113-123 [27] G. RestrepoGarces, J. Le Coze. σ-phase precipitation in two heat-resistant steels––influence of carbides and microstructure. Scripta Materialis. 50(2004) 651-654 [28] Kazumitsu Shinohara, Toshihiro Seo. Recrystallization and sigma phase formation as concurrent and interacting phenomena in 25%Cr-20%Ni steel. Material Transactions Jim. 20(2002) 713-723 [29] F. Abe, H. Araki. Discontinuous precipitation of σ-phase during recrystallisation in cold

rolled Fe–10Cr–30Mn austenite. Materials Science and Technology. 4(1988) 885-893 [30] Xiang Huang, Qingwen Zhou. Microstructure and property evolutions of a novel Super304H steel during high temperature creeping. Materials at High Temperature, Doi: 10.1080/ 09603409. 2017. 1376831 [31] T. Watanabe, H. Fujji, H. Oikawa, K.-I. Arai. Grain boundaries in rapidly solidified and annealed Fe-6.5 mass% Si polycrystalline ribbons with high ductility. Acta Metallurgica. 37(1989) 941-952 [32] Kenji Kaneko, Tatsuya Fukunaga, Kazuhiro Yamada. Formation of M 23C6-type precipitates

23

and chromium-depleted zones in austenite stainless steel. Scripta Materialia.65(2011) 509–512 [33] V. Kain, K. Chandra, K. N. Adhe, P. K. De, Detecting classical and martensite-induced sensitization using the electrochemical potentiokinetic reactivation test, Corrosion. 61(2005) 587–593. [34] K.S. de Assis, F.V.V. de Sousa, M. Miranda, I.C.P. Margarit-Mattos, V. Vivier, O.R.Mattos, Assessment of electrochemical methods used on corrosion of super duplex stainless steel, Corrosion Science. 59(2012) 71–80 [35] Shigeaki Kobayashi, Ryosuke Kobayashi, Tadao Watanbe, Control of grain boundary connectivity based on fractal analysis for improvement of intergranular corrosion resistance on SUS316L austenitic stainless steel, Acta Materialia. 102(2016) 397-405 [36] A.Perron, C.Toffolon-Masclet, X.Ledoux. Understanding sigma-phase precipitation in a stabilized austenitic stainless steel (316Nb) through complementary CALPHAD-based and experimental investigations, Acta Materialia . 79(2014) 16-29 [37] Z.B.Wang, N.R.Tao, W.P.Tong. Diffusion of chromium in nanocrystalline iron produced by means of surface mechanical attrition treatment, Acta Materialia. 51(2003) 4319-4329 [38] Jiabin Liu, Chenxu Chen, Qiong Feng, et.al. Dislocation activities at the martensite phase transformation interface in metastable austenitic stainless steel: An in-situ TEM study, Materials Science and Engineering A. 703(2017) 236-243

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Figure Captions Fig. 1 TEM morphology of the shot peened S2, (a) a field with equiaxed nanograins; (b) another field with deformation twins. Fig. 2 TEM morphology of strain-induced martensite on the deformed twinning intersection in the shot peened S2, (a) bright field image; (b) diffraction pattern of the selected area in (a); (d) the dark field image of the selected area in (a) showing the strain-induced martensite. Fig. 3 SEM morphologies of the S1 specimen aging at 650℃ for various times, (a) 0h, (b) 10h, (c) 96h, (d) 144h. Fig. 4 SEM morphologies of the S2 specimen aging at 650℃ for various times, (a) 0h, (b) 10h, (c) 24h, (d)144h. Fig.5 TEM morphology of M 23C6 in the S2 specimen aged at 650oC for 10h, (a) bright field image with selected area diffraction pattern; (b) clearer M23C6 after tilting; (c) corresponding dark field image of spot 1 in yellow dashed square in the diffraction pattern of (a).

Fig. 6 X-ray diffraction spectrums of the aged S2specimens Fig. 7 DL-EPR curves of the specimens aged at 650℃ for various times, (a) S1; (b) S2; (c) the partial enlargement to show the reactivation peaks in (b). Fig. 8 Dependence of DOS on aging time of the two specimens aged at 650℃ Fig. 9 SEM morphologies after DL-EPR tests of the S1 and S2 specimens aged at 650℃ for various times Fig. 10(a) the original EBSD map of S2 sample; (b) the partial magnification in (a); (c) Processed data for phase map and grain boundaries in channel 5 from the data in (b);(d) Processed data for 25

the distribution of CSL grain boundaries Fig. 11 SEM morphologies of specimens shot peened at 0.5MPa for different times and then aged at 650℃ for 168h, (a) 3min[18]; (b) 8min[18]; (c) 10min; (d) 12min; (e) 20min[18] Fig. 12 (a) Dependence of sigma phase content on aging time at 650℃ in the Super304H steel with different shot peening parameters; (b) Dependence of average grain size and content of strain-induced martensite on shot peening time under 0.5MPa

[18]

; (c) Dependence of the depth of

deformed layer in different SP specimens Fig. 13 Schematic illustration of stress concentration evolution with shot peening time in the shot peening specimen, (a) matrix before shot peening; (b) uniform deformation state (saturated point); (c) over saturated deformation state ( red points indicate the stress concentration sites).

26

Tables Table 1 Chemical composition of the Super304H stainless steel（mass fraction,%） C

Si

Mn

P

S

Cr

Ni

Nb

Cu

B

N

Al

Mo

As-received

0.09

0.21

0.67

0.030

<0.005

18.00

8.44

0.45

3.00

—

—

0.019

0.22

GB5310

0.07~

≤0.3

≤1.0

≤0.030

≤0.010

17.0~

7.50~

0.30~

2.50~

0.001~

0.05~

0.003~

—

-2008

0.13

19.0

10.50

0.60

3.50

0.01

0.12

0.030

27

Table 2 The parameters of solution treatment and shot peening treatment Specimen

Solution treatment

S1

1150℃ 30min+water cooling

S2

1150℃ 30min+ water cooling

28

Shot Peening 0.5MPa/12min

Table 3 The reactivation potential of S2 under different aging time Aging time/ h

0

0.1

1

2

10

24

48

96

144

Reactivation

First peak

/

-200

-219

-225

-220

-187

-190

-181

-189

Potential/ mV

Second peak

/

-275

-267

-290

-283

/

/

/

/

0

29.2

28.8

13.4

4.7

3.8

5.7

12.7

15.7

DOS/ %

29

Graphical abstract

30

Highlights

 Over saturated shot peening (SP) triggers fast sigma precipitation in Super304H.  Massive sigma phase brings a positive move of the reactivation current peak.  Before critical SP time, no sigma phase appears in deformed surface during aging.  Longer SP time results in severe local stress concentration (uneven deformation).  Uneven deformation sites favor the fast nucleation and growth of sigma phase.

31