An insight into oversaturated deformation-induced sigma precipitation in Super304H austenitic stainless steel

Materials and Design 181 (2019) 108056

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An insight into oversaturated deformation-induced sigma precipitation in Super304H austenitic stainless steel Qingwen Zhou a, Jiangwen Liu a,b, Yan Gao a,b,⁎ a b

School of Materials Science and Engineering, South China University of Technology, Guangzhou 510641, PR China Key Laboratory of Advanced Energy Storage Materials of Guangdong Province, PR China

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• Sigma nucleates at GBs/ PBs/ TBs early before recrystallization in over deformed γ. • Cr segregation and multiple slipping may be the prophase of sigma transition. • Stress relief and enhanced Cr diffusion at REX interfaces cause sigma rapid growth. • Deformation promoted massive sigma precipitation appears only at moderate T(650 °C).

a r t i c l e

i n f o

Article history: Received 19 April 2019 Received in revised form 3 July 2019 Accepted 18 July 2019 Available online 19 July 2019 Keywords: Austenitic stainless steel Nanocrystalline Sigma phase Nucleation Recrystallization

a b s t r a c t A new mechanism of the deformation-promoted precipitation of the sigma phase during aging at 650 °C in nanocrystallized Super304H austenitic stainless steel was proposed for its extremely early nucleation and subsequently fast growth. Unlike the reports of sigma phase nucleation at the recrystallizing grain boundaries in other deformed steels, the sigma phase was found to nucleate before recrystallization during aging in the oversaturated shot-peened Super304H steel at stress concentrations with high distortion energy (nanograin boundaries, twinning intersections and interfaces of dispersion strengthening phases like Nb(C,N)). The high-distortion areas favored the segregation of chromium and consequently facilitated the structure transformation from facecentered cubic austenite to topologically close-packed sigma phase. Sigma phase particles were found to grow slowly in the early stage of aging due to the residual compressive stress field in the deformed austenite matrix, and began to grow abnormally quickly to be 1–2 μm in size at the recrystallizing boundaries when recrystallization commences, due to both the release of residual compressive stress at the recrystallized region and the fast chromium diffusion at the deformed nanostructure region. To avoid the occurrence of sigma phases, the degree of surface deformation should be controlled to be lower than the deformation saturation value. © 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/).

1. Introduction

⁎ Corresponding author at: School of Materials Science and Engineering, South China University of Technology, Guangzhou 510641, PR China. E-mail address: [email protected] (Y. Gao).

For excellent corrosion resistance and mechanical properties at high temperature, austenite stainless steels are widely used in components designed for high-temperature or corrosive environment applications, such as in the power, nuclear, oil and chemical industries [1,2].

https://doi.org/10.1016/j.matdes.2019.108056 0264-1275/© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

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Super304H austenitic stainless steel is widely used in reheater and superheater tubes in thermal power plants, especially in ultrasupercritical units, due to its excellent high-temperature performance [3–5]. The addition of nitrogen and carbon and other alloying elements, such as copper and nioboium, in the steel forms dispersive Nb(C,N) and a nanoscale copper-rich phase for precipitation strengthening [6–8]. Shot peening (SP) is used in the manufacturing processes of the Super304H steel intended to enhance its steam oxidation resistance by grain refining. SP has also been demonstrated to enhance surface hardness and corrosion resistance [9,10]. To get rid of the high intergranular corrosion susceptibility caused by the high carbon content in Super304H steel [11–13], high-intensity SP has been attempted to produce a nanostructure that facilitates its quick desensitization by enhanced chromium diffusion, which in coarse-grained austenite is normally slow, during aging [14–16]. Unfortunately, the highly deformed microstructure of Super304H causes the fast precipitation of the sigma phase, a well-known, hard and brittle intermetallic compound, during aging [17]. Due to the negative effect of the sigma phase on the mechanical properties of steels [18,19], the investigation into sigma phase formation is of technological importance. In duplex stainless steel, the sigma phase may transform directly from δ-ferrite by short-range diffusion of ferrite due to their close chemical composition and the high diffusion rate of alloying elements in δ-ferrite [20–22]. In austenite stainless steel, the sigma phase usually precipitates only after long-term aging at a temperature in the range of 600 °C to 900 °C [23]. Moreover, in deformed austenite, the sigma phase may nucleate early at the recrystallizing interfaces between the recrystallized and deformed grains resulting from a minimized energy barrier by dislocation rearrangements at the recrystallizing interfaces [24,25]. Such accelerated sigma precipitation was also observed in the highly deformed Super304H steel in our previous studies [17,26]. We have proposed term oversaturated deformation to be the SP critical condition to trigger fast and mass sigma precipitation at uneven stress concentration sites with high energy [26], which corresponds to the concept of ‘steady state’ of severe plastic deformation that is believed to be a balance between deformation-induced grain refinement and dynamic recovery [27,28]. Previous studies focused on sigma particles that had already grown to micron-scale sizes and therefore did not consider the nucleation stage in early sigma precipitation. [17,24–26]. Whether the overdeformed microstructures first affect the recrystallization process and then influence sigma precipitation at recrystallizing interfaces or directly influence early sigma nucleation behavior remains unknown. Therefore, the present research is a continuation of our previous study [26] with a more detailed mechanistic analysis of the initial sigma phase nucleation sites and the relation of its fast growth with recrystallization. Some studies have been reported on the precipitation behaviors in the deformed microstructure [29,30], but they mainly relate to nonferrous metals or other alloys used in precipitation hardening states. Austenitic stainless steels are used in annealed or normalized states and there are very limited research reported regarding the aging behavior of its deformed state, so the mechanism of fast sigma precipitation in the deformed austenite is not well understood. Up to now, there has been no direct observation of the initial nucleation sites where nanoscale sigma nuclei form in deformed austenite in the early stage of aging. In this study, therefore, the overall conditions for deformationpromoted sigma precipitation, including temperature and deformation degree, were investigated. More importantly, the aim of the study was to directly demonstrate the early sigma phase nucleation sites in deformed austenite and the evolution of sigma phase during the recrystallization process. Microcosmic and intuitive evidence of nanoscale sigma phase precipitates in the early stage of aging in the severely shotpeened (SPed) Super304H stainless steel obtained by transmission

electron microscopy (TEM) (including high-resolution TEM (HR-TEM) and scanning transmission electron microscopy (STEM) which is used to gather high-angle annular dark field (HAADF) image and energy dispersive spectrometer(EDS) mapping) and electron backscatter diffraction (EBSD) is presented in this report and a new mechanism of deformation-promoted nucleation of the sigma phase before recrystallization is propounded, which is different from the previously reported mechanism [24,25]. In addition, the mechanism for the abnormal fast growth of the sigma phase enhanced by recrystallization is interpreted with an aim to avoid sigma occurrence in the surface-deformed austenitic stainless steel. 2. Method A Super304H stainless steel tube (Sumitomo Metal Industries of Japan) of the composition shown in Table 1 was used as the experimental material. The tube was solution treated at 1150 °C for 30 min and quenched rapidly in water. The tube was then processed into 50 mm × 30 mm × 4 mm samples by wire electrode cutting and the samples were subsequently ground and polished for SP. SP was conducted in an air blasting machine (AMS\\1212P) with 100% surface coverage. The stainless steel shots had a diameter of 0.5 mm. After SP, the samples were cut into 10 mm × 10 mm × 4 mm pieces by wire electrode cutting for subsequent aging. A variety of SP parameters (3 min, 6 min and 12 min under 0.5 MPa) were chosen to study the influence of deformation degree on the sigma phase precipitation behavior. Different temperatures (450 °C, 550 °C, 650 °C and 750 °C) were also used during the aging process to study the influence of temperature on the deformation-promoted precipitation (DPP) of the sigma transition process. For the detailed study of sigma phase nucleation sites, SP was carried out for 12 min (0.5 MPa/12 min), which, according to [26], is under the condition of oversaturated deformation, and the aging temperature was set to be 650 °C, which is the vulnerable temperature for sigma transition in SPed Super304H stainless steel. The selected aging times were 2 h and 168 h for early aging and late aging, respectively. For TEM, thin foil samples were prepared from an approximately 300-μm thick slice of the SP layer. Then this slice was cut into small disks with a diameter of 3 mm and mechanically ground to b50 μm to ensure that the characteristics was the SP top surface. Finally, the thin foil samples were electrochemically polished using a TenuPol-5 twinjet machine (Struers Inc.) in a solution of 10 v/v% perchloric acid and 90 v/v% ethanol at −20–-15 °C. This electrochemical polishing will dissolve the mechanically polished surface to eliminate possible stress introduction. TEM was performed using a JEOL-2100 field emission transmission electron microscope at an accelerating voltage of 200 kV. EBSD was used to study the grain orientation distribution and for phase identification. The samples for EBSD analysis were first ground and mechanically polished to a finish of 2.5 μm using a diamond paste and then electrolytically polished in a 10 v/v% perchloric acid and 90 v/v% methanol solution at a current of 1 A/cm2 for 60 s. EBSD measurements were made at an operating voltage of 10 keV and a step size of about 0.1 μm. The low operating voltage makes it possible to show deformed microstructures clearly. HKL CHANNEL5 software from Oxford Instruments was used for EBSD data analysis. 3. Results 3.1. The critical conditions to trigger fast sigma transition Oversaturated deformation has been speculated to be the condition for rapid sigma precipitation [26]. For phase transformation, temperature is also a crucial condition that influences the conversion rate and determines whether the transition may happen. In this study, the gradients of both temperature and deformation degree were taken into consideration. Fig. 1 shows the overall critical conditions to trigger sigma phase precipitation in Super304H steel with different deformation

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Table 1 Chemical composition of Super304H stainless steel (mass fraction, %).

Sample GB5310-2008a a

C

Si

Mn

P

S

Cr

Ni

Nb

Cu

B

N

Al

Mo

0.09 0.07–0.13

0.21 ≤0.3

0.67 ≤1.0

0.030 ≤0.030

b0.005 ≤0.010

18.00 17.0–19.0

8.44 7.50–10.50

0.45 0.30–0.60

3.00 2.50–3.50

– 0.001–0.01

– 0.05–0.12

0.019 0.003–0.030

0.22 –

Chinese standard code case: GB 5310-2008, Seamless steel tubes and pipes for high pressure boiler.

degrees after aging at different temperatures for 168 h. The microstructures of the aged samples differed depending on the temperature used. As for the non-deformed samples, aging is a process of grain coarsening at different temperatures [16,31,32]. However, the deformed microstructures underwent recovery and recrystallization processes. At 450 °C and 550 °C, the deformed microstructures remained and recrystallization had not yet happened in all deformed samples as the supplied heat was not sufficient to trigger recrystallization. In contrast, full recrystallization (mentioned as Full Rec in Fig. 1) occurred at 750 °C in all deformed samples and the recrystallized microstructure consisted of austenite without sigma phase precipitation. The deformation-promoted sigma phase was only triggered at 650 °C after severe SP treatment (0.5 MPa/8 min) and no sigma phase precipitation was observed when the deformation degree is mild (0.5 MPa/3 min) [16]. In the present study, the critical SP duration to trigger saturated deformation was found to be 6 min at 0.5 MPa, a little shorter than the 8 min reported in our previous study [26], which may be due to different SP equipment being used. From Fig. 1, deformationpromoted sigma phase transition is clearly related to both deformation degree and aging temperature. To investigate the nucleation sites and abnormal growth of deformation-promoted sigma precipitation in detail, the overdeformed 0.5 MPa/12 min sample was chosen and aged at 650 °C for 2 h and 168 h. The as-deformed microstructures of the 0.5 MPa/12 min sample are given in Fig. S1 in the supplementary material. 3.2. The early sigma phase precipitation sites in over-deformed samples In our former study, the sigma phase in 0.5 MPa/12 min samples was only observed by SEM after aged at 650 °C above 24 h [26]. Due to the

resolution limitation in SEM, TEM was used in this study to observe the early sigma phase precipitation sites. The 0.5 MPa/12 min sample was aged at 650 °C for 2 h and then imaged using TEM to observe the early sigma phase nucleation sites in the oversaturated SPed Super304H stainless steel. 3.2.1. Grain boundaries Fig. 2 shows the microstructure of the sigma phase at grain boundaries (GBs) and triple junctions in the SPed sample aged at 650 °C for 2 h. In Fig. 2(a), both the sigma phase and M23C6 are located at the GBs where the unevenly deformed strain contrast microstructure shows the existence of locally high stress and strain concentration. The electron diffraction patterns of these two precipitates are shown in Fig. 2(b) and (c). In Fig. 2(c), two sets of diffraction patterns of the sigma phase are observed, which indicates the existence of two sigma particles. Unfortunately, the morphology is not sufficiently clear to distinguish them in Fig. 2(a) because one is located at the GB with high distortion. Fig. 2(d) shows three sigma phase particles located at the triple junction of austenite grains. The STEM image in Fig. 2(e) and corresponding EDS mapping of chromium in Fig. 2(f) support the identification of these particles as sigma phase. The surroundings around the sigma precipitates after tilting are shown in more detail in the supplementary materials (Fig. S2). In Fig. 3(a), the interplanar spacing of (111)A in the deformed austenite matrix was measured from its diffraction pattern to be about 0.207 nm, which was similar to the non-deformed value (0.208 nm) and indicated that the lattice parameters of the austenite grains remained stable although SP introduced residual compressive stress [33,34]. However, the interplanar spacing of (111)A was only 0.197 nm in the region where sigma transition occurred as in Fig. 3

Fig. 1. Overall critical conditions to trigger sigma phase precipitation in Super304H stainless steel with different deformation degrees after aging for 168 h at different temperatures.

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Fig. 2. TEM of the sigma phase microstructures in a SPed sample aged at 650 °C for 2 h; (a) sigma phase near M23C6 at GBs; (b) diffraction pattern of M23C6; (c) diffraction pattern of the sigma phase; (d) three sigma phase particles at the triple junction of GBs; (e) STEM image of sigma phase particles in (d); (f) EDS mapping of chromium in (e).

(b) and (c), which indicated a higher residual compressive strain here. This change in lattice parameter shown in Fig. 3 further proved that the sigma phases nucleated at sites with locally high strain. As these regions were nano- micro-sized, their contribution to the overall lattice parameter in the diffraction ring could not be determined. During plastic deformation, GBs inhibit slip movement and cause high residual stress and strain concentration at GBs, especially at triple junctions, which provide a high energy fluctuation favoring the nucleation of precipitates such as M23C6 and the sigma phase. Normally compressive stress fields exist after SP [33,34], which is not conducive to the growth of sigma nuclei (diffusional phase transformation); at oversaturated deformation sites, however, a complicated stress situation may disturb the regular distribution of compressive stress and generate some areas of tensile stress, which favors early sigma phase nucleation. For examples, the d(111)A near GBs in Fig. 2(b) was measured to be 0.212 nm, which is larger than its non-deformed value (0.208 nm). The d(200)A value was measured to be 0.184 nm, which was likewise larger than its non-deformed value (0.180 nm). Also, a widened interplanar distance in (111)A of 0.214 nm is found in Fig. 5 (b). These data support the existence of local tensile strain near GBs in samples with oversaturated deformation. A recent work also demonstrates an increase in tensile stress near GBs where dislocation motion was halted [35]. From the above results, it is clear that oversaturated SP causes an increased local stress and strain near GBs. Although overall residual compressive stress fields exist after SP [33,34], local tension and compression stresses coexist at GBs.

3.2.2. Deformation twins Deformation twinning is the main means of grain refinement in Super304H stainless steel [16]. After aging at 650 °C for 2 h, most of the deformed microstructures remained un-recrystallized as shown in Fig. 4. In the severely deformed sample, de-twinning and dislocations coexist and the straight twin boundaries (TBs) are fragmentized. Sigma phase particles were found at the intersection of fragmentized TBs, as shown in Fig. 4(a). Fig. 4(b) shows the magnification of sigma particles in the white square in Fig. 4(a) after a certain tilting to show the TBs clear and Fig. 4(c) shows the diffraction pattern of the area marked in Fig. 4(b). The deformation extent (strain contrast) around sigma particles at the intersection of TBs was also obviously higher than in the adjacent region in Fig. 4(b). Fig. 4(d) shows the EDS mapping of chromium in Fig. 4(b) to demonstrate the chromium enrichment in sigma particles. Fig. 4(e) shows the EDS line scanning results across one sigma particle (white line with arrow in Fig. 4(a)), where a copper-rich particle adjacent to the sigma phase is detected. In Super304H stainless steel, more alloying elements have been added and a rich variety of dispersion strengthening phases, such as Nb(C,N) [6,31] and copper-rich phases [7,8], exist to ensure its excellent high-temperature performance [6–8].

3.2.3. Nb(C,N) phase boundaries Fig. 5(a) is the combination of HAADF image and the EDS mappings of chromium and niobium where several sigma particles (with high

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Fig. 3. SPed samples aged at 650 °C for 2 h; (a) TEM image of the austenite matrix with corresponding diffraction pattern; (b) TEM image of the sigma phase microstructure; (c) The magnified HR-TEM image of the boundary between the sigma phase and austenite.

chromium content marked in red) are located near the phase boundary of Nb(C,N) (with high niobium content marked in yellow) in the early stage of aging. Fig. 5(b) shows the magnified sigma particles (marked

by yellow arrows) and one of them is located at the Nb(C,N) phase boundary (marked by a red arrow). The dispersive Nb(C,N) phase results in the piling up of vacancies, dislocations and high stress

Fig. 4. TEM of the sigma-phase microstructures in a SPed sample aged at 650 °C for 2 h; (a) bright field image of two sigma-phase precipitates; (b) magnification of the square in (a); (c) corresponding (red circle) diffraction pattern of a sigma grain; (d) EDS mapping of (b); (e) EDS line scanning of the white line in (a).

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Fig. 5. (a) The combined EDS mapping of chromium and niobium in a SPed sample aged at 650 °C for 2 h; (b) The magnification of precipitates in the white square in (a) and the corresponding selected area diffraction pattern, which shows the presence of the sigma phase near Nb(C,N).

concentration at the phase boundaries during deformation, which contributes to the early sigma precipitation. 3.3. Crystallography of sigma phase transformation from deformed austenite 3.3.1. Chromium segregation as a prophase for sigma nucleation Chromium enrichment was observed at the phase boundary of Nb(C, N). In Fig. 6(a), two sigma phase precipitates at the interface between Nb(C,N) and austenite are identified by the EDS mapping results in Fig. 6(b) and (c), from which the location of the sigma phase particles can be determined to be at the phase boundary of the Nb(C,N) particle. Most importantly, there is a chromium-rich zone at this phase boundary marked as “segregation” in Fig. 6(a) and (c). In Fig. 6(d) the incoherent boundary in the HR-TEM image between this segregation zone and sigma nucleus proves that this chromium-rich zone had not completed its structure transition from austenite to sigma but only to the stage of being a chromium accumulation area. This phenomenon indicates that chromium segregation could be the prophase of sigma transition during

aging. Also, some fringes in this segregation region are visible in Fig. 7 (d), which may be a superlattice since chromium atoms may occupy specific lattice points in the face-centered cubic (FCC) structure to form an intermediate state. However, more evidence is needed to figure out the exact microstructure of this state. Under radiation treatment, α’ precipitation may be induced because of more point defects and dislocations are introduced by the radiation, iron–chromium alloys undergo α-α’ phase separation to form chromium-rich zones and the sigma phase may form subsequently [36–38]. As a diffusive transformation, the nucleation of chromiumrich sigma phases is controlled by chromium diffusion and chromium segregation. Chromium diffusion is by substitutional diffusion, which is a low process. However, the chromium diffusion process is accelerated by the high vacancy concentration in the over-deformed microstructure, which is similar to the radiation process. More HR-TEM images were recorded to show detailed information about the surrounding of chromium-rich segregation area adjacent to the sigma phase, as shown in Fig. 7. From the similar fast Fourier transform (FFT) pattern of square A (austenite) and square B (segregation

Fig. 6. TEM of sigma phase microstructures in a SPed sample aged at 650 °C for 2 h; (a) the bright field image of two sigma phase precipitates at the phase boundary between Nb(C,N) and austenite; (b) the EDS mapping of niobium; (c) the EDS mapping of chromium (in red) and iron (in blue); (d) A HR-TEM image of the boundary between a sigma nucleus and segregation area.

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Fig. 7. (a) A HR-TEM image of the intersection between the Nb(C,N) particle, austenite matrix and chromium segregation area with the corresponding FFT patterns; (b) a HR-TEM image of the junction of austenite and the segregation area; (c) schematic diagram showing the formation of chromium segregation areas.

region) in Fig. 7(a), we concluded that the segregation zone (B) had the same crystal structure as austenite (A). In Fig. 7(b), the interplanar crystal spacing (d(111) = 0.223 nm) of the segregation zone (B) is larger than that of the austenite matrix (d(111) = 0.207 nm). A stacking fault is observed to show the severe deformation degree that caused a microstructure with a high lattice disorder, such as vacancies. The vacancies at the over-deformed sites offer more positions for chromium atoms to occupy. In the early stage of aging, there are mainly three elements (iron, chromium and nickel) in the austenite matrix. Iron is the solvent atom and the radius of nickel is larger than that of chromium so the diffusion ability of chromium is higher than that of nickel and this atom is prone to occupy the vacancies at over-deformed sites to form chromium segregation regions. Fig. 7(c) shows a schematic diagram of chromium enrichment, as modelled by the equation of the net flux of chromium (JCr) in radiation-induced-segregation (RIS) [39] where the segregation process is attributed to the association of solute atoms with the irradiation-induced-vacancies and the original concentration gradient. There is no original concentration gradient of chromium (∇0Cr = 0) compared to the adjacent matrix after SP, but more vacancies are formed at the severely deformed zones and the chromium atoms are prone to occupy the vacancy through atom transition. Deformation-promoted segregation at overdeformed sites is attributed to the association of solute atoms with deformation-induced point defects, similar to the vacancy mechanism of RIS. 3.3.2. Lattice transformation from austenite to sigma phase With a magnification of the precipitates in Fig. 5, those sigma particles and one Nb(C,N) particle are distinguished more clearly, as shown in Fig. 8(a). In Fig. 8(b), the HR-TEM image of the interface between the sigma and Nb(C,N) phases (red square in (a)) is found to be

incoherent, so the parent phase of sigma transformation is not Nb(C, N). Fig. 8(c) shows the HR-TEM images of the interfaces between the austenite and three sigma particles (yellow square in (a)) with the corresponding FFT patterns. A coherent interface was observed between the sigma phase and austenite and (330)σ was clearly parallel to (111)A in M1 and M2. Furthermore, the FFT diffraction pattern P1 of the boundary region between sigma 2 and austenite shows the weak diffraction spots of (110)σ and the strong coincident spots of (330)σ and (111)A. This region (in M1) is therefore a mixture of the sigma phase and austenite. The FFT diffraction patterns P2, P3 and P4 of the sigma particles in Fig. 8(c) are different from each other and there is a certain angular deviation between the zone axes in P3 and P4 in spite of the same (111)σ spots being present. The contrast of HR-TEM image is always influenced by the atomic number of the elements in the matrix. In M1, all the parallel atom planes have similar contrast. In M2 on the other hand, the (330)σ atomic planes have different contrast from the (110)σ atomic planes. Based on recent research on the crystallography of topologically closepacked (TCP) phases [40], the atom arrangement along the [001] direction in the sigma unit cells is derived as shown in Fig. 5 (M2). In planes A and D (corresponding to (110)σ), the ratio n of chromium to iron is two, while in planes B and C (corresponding to (330)σ), the ratio n is one, hence the different contrasts of (110)σ and (330)σ crystal planes are introduced in their high-resolution images. Under most conditions, a new phase is likely to form on a certain crystallographic plane (habit plane) of the parent phase to reduce the interfacial energy [41]. From the above results it becomes clear that the FCC austenite matrix transforms to the TCP sigma phase. The crystallographic orientation between the sigma phase and austenite have been widely studied in other materials, and different results have been found

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Fig. 8. (a) The magnified three sigma particles at the phase boundary between Nb(C,N) and austenite shown in Fig. 4; (b) HR-TEM image of the red square (the interface of sigma 1 and Nb (C,N)) in (a); (c) HR-TEM image of the yellow square (the interfaces of three sigma grains) in (a) with some detailed magnifications and FFT patterns.

under different situations [24,42–44]. The sigma phase prefers to transfer at some special sites where misorientation angle distribution is of shallow peak in a range of 35–45° [43] and follows Kurdjumov–Sachs, Nishiyama–Wasserman, Greninger–Troiano, and Pitsch relationships [42]. In the oversaturated SPed Super304H stainless steel, the coincident spots of (330)σ and (111)A are observed. The similar interplanar spacing between them may be advantageous for the phase transformation from austenite to sigma. The transition process of (111)A to (110)σ was chosen to study the lattice transformation based on the above results and its schematic transition diagram is shown in Fig. 9. In austenite stainless steel, the iron, chromium and nickel occupancies are not fixed so grey dots are used to represent the lattice points that may be occupied by these atoms. Chromium enrichment is the first step of early sigma nucleation when some lattice positions in the FCC structure are occupied by chromium atoms, which is facilitated by chromium diffusion enhancement in the deformed matrix. The slip system of FCC is (111)[110], so slipping along this direction is eased during deformation. The lattice transition from FCC to primary TCP in step 2 is facilitated by the high energy concentration at the overdeformed sites, which may overcome the energy barrier of lattice transformation. As TCP is a tetragonal crystal system, the lattice parameters a and c are different, and c is smaller than a in the sigma phase unit cell with limited space for atoms to occupy. As a result, the atoms at the c side transport spontaneously to other sides with higher capacity to reach a stable sigma unit cell (step 3). Because of the higher energy state at unevenly deformed sites, multiple slips can occur simultaneously in austenite and there may also be multiple crystal orientation relationship between the sigma phase and austenite. All in all, the oversaturated SPed microstructure features, such as the intersections of GBs and TBs with a high density of dislocations and vacancies, greatly accelerate sigma nucleation at the early stage of aging from both chromium segregation and a facilitated lattice transition.

3.4. The interpretation of the abnormal rapid growth of sigma phase In former studies, both the fast nucleation and growth of the sigma phase have been attributed to the recrystallization of deformed austenite [24,25]. The present study has proven that sigma nucleation occurs before the recrystallization of deformed austenite, but it remains unknown whether recrystallization influences the fast growth of the sigma phase as it does in other kinds of discontinuous precipitation reactions [29,30]. Fig. 10 shows the distribution of micron-sized sigma particles in SPed Super304H stainless steel after aging at 650 °C for 168 h determined by EBSD. In Fig. 10(a), some austenite grains (red) have grown to micron size (fully recrystallized) with large sigma particles (blue) whose size is up to 1 μm at GBs or GB triple junctions. However, the identification of sigma particles in the deformed microstructure may be influenced by residual stress. EDS mapping was therefore used to

Fig. 9. Schematic diagram of solid transformation from austenite to sigma phase.

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distinguish them since sigma particles have a higher chromium content and lower nickel and iron content than austenite. The results show that micron-sized sigma particles only appear in the recrystallized area in Fig. 10(b). Moreover some sigma particles located adjacent to the Nb (C,N) particles verified that the phase boundaries were the preferential sites for sigma nucleation. The residual stress in the deformed area influenced the resolution of EBSD and the small size of the sigma particles here made them difficult to identify by EDS. As a result, the nano-sized sigma particles are not observed in Fig. 10. These phenomena indicate that a rapid growth of sigma phase only occurred in the recrystallized area. Since micron-sized sigma particles only appear in the recrystallized area, as shown in Fig. 10(b), the residual compressive stress field [33,34] in the deformed austenite is assumed to have impeded the growth of the sigma phase; however, the sigma phase may undergo rapid growth at the recrystallizing boundaries, as schematically shown in Fig. 11. In Fig. 11(a), sigma particles (blue) are shown to have nucleated at unevenly high distortion sites in the oversaturated austenite before recrystallization and the residual compressive stress field after SP (green circle) existed around them, which hinders their growth. When sigma particles met the recrystallizing boundary, they started to grow quickly towards the recrystallized side because of the relief in residual compressive stress. Additionally, the fast chromium diffusion in the deformed side facilitated the growth of the FeCr intermetallic sigma phase, as shown in Fig. 11(b). As recrystallization progressed, the sigma particles grow to be micron sized and were gradually surrounded by a recrystallized microstructure (Fig. 11(c)). When completely surrounded by the recrystallized grains, sigma phase growth slowed down significantly due to the disappearance of enhanced chromium diffusion. Since sigma particles have a pinning effect on the movement of the recrystallizing interface, all those micronsized sigma particles (in Fig. 11(d)) were located at the boundaries instead of in the austenite grains after recrystallization. Moreover, austenite grains adjacent to the large sigma particles were smaller due to the pinning effect, as shown in Fig. 10(a). 4. Discussion From the above results and analysis, it is known that the recrystallization process contributes to the rapid growth of the sigma phase, but

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the extremely early nucleation of sigma phase is mainly attributed to the unevenly distributed distortion energy in the oversaturated SPed Super304H stainless steel. This new mechanism is partially different from the previously reported one [24,25] in which sigma nucleation was believed to occur simultaneously with recrystallization. Normally, the sigma phase will not appear in Super304H stainless steel until after long-term aging [6]. Even after a relatively mild SP treatment, the sigma phase does not precipitate during aging, which indicates that the structure and composition barrier for sigma transformation in Super304H stainless steel is quite high [26]. The oversaturated SP helps to overcome the energy barrier and trigger sigma phase transformation, with chromium segregation and multiple slips as the prophase at some stress concentration sites during aging. A coherent phase boundary is observed between the sigma phase and austenite with coherent diffraction spots of (330)σ and (111)A. A complete deformation-promoted sigma transformation process is schematically illustrated in the supplementary materials (Fig. S3). For oversaturated SPed austenite, both the high energy input (including deformation energy and heat energy) and the evolution of deformed microstructure during aging influenced the sigma transformation process. As a result, the precipitation behavior of the SPed austenite was different under different temperatures. The energy evolution with aging time at different aging temperatures is illustrated in Fig. 12. Low-temperature aging mainly resulted in a slow stress relief, and the deformed microstructures remained during aging without sigma nucleation due to the insufficient energy input (the red, horizontal, dotted line in Fig. 12 represents the critical energy barrier of sigma precipitation). In contrast, high-temperature aging caused in rapid stress relief and recrystallization, which happened much earlier than the sigma phase nucleation (vertical dashed line), and resulted in the fully recrystallized austenite without sigma phase precipitation. Only at moderate temperatures does stress relief occur in moderation and the heat energy and deformation effect function synergistically to trigger sigma nucleation. Although surface nanocrystallization is a promising way to solve the problem of the high intergranular corrosion sensitivity of Super304H stainless steel [16] and its desensitization time decreases with a higher deformation degree [17], its service temperature (600–650 °C) in ultrasupercritical units is in the sensitive region to trigger DPP of the sigma

Fig. 10. Elemental maps obtained by EBSD data of a SPed sample after aging at 650 °C for 168 h; (a) phase map with legends and a partial magnification of micron-sized sigma particles; (b) corresponding EDS mappings of iron, chromium, nickel and niobium.

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Fig. 11. Schematic illustration of the rapid sigma phase growth during recrystallization in the over-saturated SPed specimen.

phase, which must be inhibited in practical applications because of its deteriorative effect on the sample performance. The above results reveal that sigma nucleation happens at over-deformed sites; therefore, the surface SP deformation degree must be controlled to be lower than the saturation value, so as to avoid deformation-promoted sigmaphase precipitation. 5. Conclusions (1) Due to the uneven distribution of stress and strain after oversaturated deformation, the sigma phase nucleates at high stress concentration sites, such as triple junctions of GBs, TBs intersections, and phase boundaries of Nb(C,N) and copper-rich phases, extremely early during aging. (2) The sigma phase is transformed directly from deformed austenite with chromium segregation as the prophase before recrystallization. A coherent phase boundary between the sigma phase and austenite is observed at the early stage of phase transformation. (3) Sigma particles remain nano-sized in the deformed microstructure because of the constraint of residual compressive stress around them. However, they start to grow rapidly to micronsizes at the recrystallizing interfaces due to both the stress relief at the recrystallized side and the enhanced chromium diffusion at the deformed side. (4) The surface SP degree must be controlled to be lower than the deformation saturation value, so as to avoid deformationpromoted, fast sigma-phase precipitation in austenitic stainless steel.

Fig. 12. Schematic energy evolution of oversaturated SPed samples with aging time at different aging temperatures.

CRediT authorship contribution statement Qingwen Zhou: Investigation, Methodology, Visualization, Writing original draft. Jiangwen Liu: Formal analysis, Data curation, Investigation. Yan Gao: Conceptualization, Supervision, Resources, Writing - review & editing. Declaration of Competing Interest None. Acknowledgement This work was supported by the National Natural Science Foundation of China [grant number 51471072] and the Key Laboratory of Advanced Energy Storage Materials of Guangdong Province. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.matdes.2019.108056.

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