Characterisation of TCP precipitation sequences of superaustenitic stainless steel and correlation with electrochemical and mechanical properties

This study investigates the influence of precipitation sequences formed at critical high-temperature range for various ageing time of super austenitic stainless steel 654SMO against the evolution of its corrosive resistance performance through the microstructural characterization, hardness and electrochemical corrosion testing. The specific type of stainless-steel alloy is one of the most corrosion resistant along with exceptional mechanical properties, tailor-made for pressurized and erosive systems handling chlorinated sea water, plate heat exchangers, and fuel gas cleaning applications and an excellent case study of in-depth phase transformations-electrochemical properties characterization phenomena. The microstructure evolution is characterized by scanning electron microscopy (SEM) coupled with energy-dispersive spectroscopy (EDS), hardness testing and potentiodynamic polarization curves. The corresponding precipitation characteristics and sequences of secondary phases are discussed and correlated with electrochemical properties. The results demonstrate σ phase with various morphologies formed during isothermal aging and the formation of three more secondary topologically close pack (TCP) phases, such as chi phase (χ), laves phase (A2B), and Cr2N against corrosive resistance and mechanical performance, which are of primary importance for understanding electrochemical phenomena of different phases and useful for optimizing such alloys for various applications.


Introduction
Super-austenitic stainless steels (SASSs) demonstrate a remarkable potential for deployment within exceedingly demanding operational environments.These include uses in the petrochemical sector, plate heat exchangers, sea water desalination, and flue gas desulfurization [1].Among these materials, the super-austenitic stainless steel S32654 presents increased concentrations of important alloying additions such as chromium (Cr) (24.0 -25.0 wt.%), nickel (Ni) (20.0 -21.0 wt.%), molybdenum (Mo) (7.0 -8.0 wt.%), and nitrogen (N) (0.45 -0.55 wt.%).This particular formulation has the combined benefit of exhibiting exceptional resistance to corrosive environments and outstanding mechanical properties [2].However, the augmentation of alloy elements drives not only to enhance mechanical strength and corrosion resistance but also to catalyze the formation of diverse secondary phases, such as TCP phases, during exposure to elevated temperatures [3].These phases impair essentially the inherent corrosion resistance and the mechanical integrity of the studied steel [4].
In addition to the diverse array of precipitates encountered, the process of precipitation in S32654 alloy is characterized by a high degree of complexity.This complexity arises from the coexistence of competitive and simultaneous solid-state reactions within a defined temperature spectrum.As described by [5], a comprehensive analysis of the equilibrium secondary phases prevalent in S32654 within the temperature range of 700-1100 °C, indicates the prominent presence of the σ phase and Cr2N.Notably, the swift precipitation kinetics observed for the σ phase can be attributed to the elevated chromium and molybdenum concentrations inherent to the S32654 composition.Similarly, it is revealed that three distinct intermetallic phases, namely χ, σ, and Laves phases, in conjunction with the nitride Cr2N, are known to appear within S31254 and S32654 alloys [6].Furthermore, their research presents the correlation between cold deformation and the facilitation of secondary phase precipitation and coarsening processes [7].
According S. Zhang [3], increasing Cr and Mo contents significantly promote and participate to the formation of σ phase and inhibit the precipitation of other low temperature secondary phases such as R and π phase.The promotion effect of Cr and Mo on precipitation further expands the area of Crand Mo-depleted zones around various precipitates, enhancing the intergranular corrosion susceptibility of S32654.The study also revealed that Mo has a stronger promotion effect on the formation of σ phase than Cr.The effects of Cr and Mo on the precipitation behavior were mainly achieved by affecting the driving force for precipitation and the activity of elements.So, the motivation behind this research lies in the need to understand the alloy's behavior under thermal exposure and the correlation of the microscale microstructural changes to its overall corrosion resistance and mechanical integrity by a broad study consideration between the alloy's exceptional properties and the challenges posed by the formation of secondary phases.Finally, the understanding of the effect of precipitation sequences on electrochemical properties can lead to modified improved alloy design and highly corrosion resistant ferrous materials for the upcoming demanding applications.

Experimental procedure
The investigated material is a super austenitic stainless steel S32654.The chemical composition (in wt.%) of the steel is 0.01C-3.50Mn-24Cr-22Ni-7.30Mo-0.50N-0.50Siand Fe balanced.Starting with 3 mm thick as received plate of S32654, samples with the dimension of 20 mm × 10 mm × 3 mm were cut from the hot rolled plate and solution treated at 1300 °C for 15 min, followed by water quenching to achieve a homogenized super saturated solid solution austenite matrix at ambient temperature.Consequently, the samples were isothermally aged at 750 o C and 850 o C for 30 min, 6 h, 24 h, 120 h and 1000 h and finally quenched in water.Exposure to atmospheric conditions was chosen in order to simulate service conditions of this steel.Metallographic preparation for scanning electron microscopy (SEM) examination of the specimens, employing a JSM6380 at 20 kV, was performed involving standardized metallographic preparation techniques.The metallographic samples were firstly ground (240 -3000 grit) and then polished using diamond paste (6 μm -1 μm).Subsequently, these samples were electrolytically etched in a aqueous solution of 10 % oxalic acid and deionized water at 2V for 0.5 s to 4 min.Energy dispersive X-ray microanalysis (EDS) was used in order to assess the elemental composition of the various secondary phases.Finally, potentiodynamic polarization testing through 3.5%wt.NaCl aqueous solution with 1mV/sec rate and -800 up to +1400 mV range, as well as hardness tests with 197 N force for 15 seconds were employed to evaluate the influence of TCP phases on microstructure-properties relationship.

Results
Figure 1a and 1b depict SEM micrographs of the as-received and solution-treated samples.The asreceived sample (Fig. 1a) consisted of a single austenite phase, twin crystals, with no secondary phases and an approximate grain size of 30-50μm.The solution heat-treated sample (Fig. 1b) also exhibited an austenite phase but with significantly larger grain sizes ranging from 300-400μm.However, the heat treatments of 15 minutes at 1300°C was insufficient for the complete dissolution of large carbides in the centreline segregation.Upon aging at 750°C for 30 minutes (Fig. 1c), minimal changes were observed, but at 850°C (Fig. 1d), continuous precipitation sequences began to nucleate and grow at grain boundaries with a width less than 1 μm.Increasing the aging time to 6 hours at 750°C (Fig. 1e) led to similar precipitations as the 850°C, 30-minute sample.However, they nucleated as a grid of connected precipitates rather than cohesive ones.Notably, at 850°C and 6 hours (Fig. 1f), some cellular precipitates nucleated from the grain boundaries and gradually grew preferentially towards the austenite grains.
As the aging time approached 24 hours (Fig. 2), both the volume fraction and mean size of the cellular precipitates significantly increased.Additionally, many intragranular constituents transformed into a needle-like shape, and two new types of precipitates became observable.It wasn't until 120 hours of aging at 850°C (Fig. 3b) that cellular precipitates ceased to increase in size.Moreover, when the needlelike phases grew to a certain size, they pinned the migration of cell boundaries and hindered the growth of cellular phases.Meanwhile, precipitates reached a critical size, making them easily detectable by EDS.Their size and volume fraction noticeably increased, with black contrast ones appearing in larger quantities and in various shapes compared to the white contrast ones.The nucleation points of these precipitates were closely related to the different precipitates existing in the steel.After long-term aging at 750°C for 120 hours (Fig. 3a), the precipitation behaviour differed significantly compared to 850°C.The presence of cellular phases was reduced, and intergranular deposits increased substantially.An alternation between black and white contrast precipitates suggested a different enrichment in containing alloying elements.Table 1 presents the chemical composition of all intermetallic precipitates and phases found in the current steel.Concerning the heat treatment at 1000 hours, three main differences were observed between the 850°C and 750°C sample (Fig. 3c, Fig. 3d).Firstly, the 750°C sample was much finer-grained than the 850°C sample.Secondly, it featured white and black precipitates that increased over time, while the 850°C sample lacked white contrast phases and had an abundance of black ones.Thirdly, the absence of needle-like black contrast phases was apparent in the 750°C sample.The corrosion characteristics of the sample at 750 °C (Fig. 4) exhibit a relatively consistent trend within the range of 2.5 to 5.6 μA/cm 2 , except for a notable deviation observed at the 6-hour aging mark, where the corrosion current increases to 10.96 μA/cm 2 , subsequently decreasing to 5.75 μA/cm 2 at the 24-hour mark.Conversely, at 850 °C (Fig. 5), significant disparities are observed in the corrosion current of the steel.Following a transient increase during the initial 30 min of aging, there is a subsequent decline in corrosion current until the 24-hour aging period.However, a significant spike is observed at the 30min-mark reaching nearly 25 μA/cm 2 , indicating a substantial corrosion current for the steel.Eventually, the corrosion current gradually decreases, until the 1000-hour aging mark, to 12 μA/cm 2 .

Discussion
The sigma phase is regarded as the most common and significant intermetallic phase in SASSs.It is observed based on the change of hardness testing results (Fig. 6) that the rich in Cr and Mo σ phase transforms the studied alloy to a harder but also more brittle microstructure.Nevertheless, the precipitation behaviour of this phase varies, depending on many factors, such as local chemical composition of austenite grain and grain boundaries, initial grain size, pre-existing stress-strain state and ageing conditions [5][6][7][8].In this study, three types of σ phases were detected and studied; intergranular, cellular and intragranular ones.The first type of σ phase started to form at 850 o C and 30 min of ageing.It formed as a continuous precipitate sequence on grain boundaries with thickness about 1 μm.After long-term ageing the shape of this intergranular σ phase changed, with many chromium nitrides interrupting the continuous precipitate and forcing it to take a slightly different tile shape.Following the same pattern at 750 o C but with longer periods of ageing, approximately 6 h, refined intergranular sigma phase started to form again on boundaries.It seems that also promotes the nucleation and growth of other precipitates next to it.Finally, at 1000 h grain boundaries have a completely discontinued formation of precipitates.This phenomenon is more significant at 750 o C as the precipitates are finer than 850 o C. The second type of σ phase, the cellular formation, is nucleated on the grain boundaries and towards the austenite grain.Concerning the precipitation temperature, it is widely observed that the cellular σ phase only formed at the temperature bigger than 800 °C and no cellular precipitation of σ phase occurred when the aging temperature was 750 °C.This phenomenon can be attributed to the fact that ageing at lower temperature (i.e.750 °C) presents too low diffusion rate of elements to meet the nucleation and growth kinetics of a eutectoid transformation to cellular type σ phase and secondary austenite.Regarding the intragranular σ phase formation, its nucleation and growth rates were much lower than those of intergranular σ phase.,The nucleation energy required for intragranular nucleation is higher than that for intergranular nucleation.Furthermore, the diffusion rate of σ phase forming elements (Cr and Mo, etc.) within the austenite grain are much lower than that along grain boundaries [8].The high Cr and Mo contents in S32654 greatly promote the formation of σ phase, and its nucleation and growth consumed large amounts of these elements and left Ni and N in surrounding matrix.As a result, Cr-and Mo-depleted and N-rich zones formed adjacent to σ phases leads to austenite γ2 with significantly less Cr, Mo compounds and much more Ni in its lattice.Theoretically, the first type of precipitates that nucleated was the CrN.Chromium nitride is an interstitial compound, in which nitrogen atoms reside in the octahedral spaces between the chromium atoms in a fcc lattice, making the compound susceptible to stoichiometry deficiency.This results in the formation of metallic dichromium nitride (Cr2N), a second interstitial compound.Cr2N, on the other hand, crystallizes in a hexagonal crystal structure in contrast to CrN, and is typically described as a secondary phase in either CrN hard coatings or steel.[9].It is observed that it started to form after 24 h of ageing at 850 o C, whereas the nucleation started a little before 120 h at 750 o C. At first they seemed as black spots at grain boundaries with irregular shape.As the ageing process continued, they grew next to existing sigma, chi and laves phase precipitates.The most interesting about Cr2N is the one formed within the needle-like σ phases.More precisely, it is more like to be transformed by local needle-like σ phases.As mentioned above, the σ phase is rich in Cr and Mo and its surrounding matrix is rich in N, while the Cr2N and chi phase are Cr-rich and Mo-rich phase, respectively.Thus, during the transformation process from σ phase to Cr2N, the σ phase could provide Cr and the matrix surrounding σ phase could supply N for the nucleation and growth of Cr2N, between their interphase boundary.As a consequence, the Mo in the transformed σ phase was left and supplied for the formation of, also Morich, laves phase.The transformation between needle-like σ phase and Cr2N was a very slow process that is the reason that it started to show up as needles after 120 h.Regarding the Laves phase shows a remarkably higher content in molybdenum than any other phase present.Due to the fact that it is the brightest phase, is thus richer in the heavy element molybdenum than σ phase, it may be clearly recognized on the basis of BSE (back scattered electrons) contrast in SEM.Although the amount of Laves phase is unquestionably bigger, sigma phase and laves phase develop together at the start of the process.Even after an initial increase in the volume fraction, it decreased again, suggesting that laves phase is metastable in this steel.
Last but not least the observation of potentiodynamic polarization curves, as well as of potentiodynamic parameters (Ecorr and icorr) suggests that the presence of precipitates affects the corrosion behavior of aged specimens.Furthermore, it is demonstrated that, depending on the aging time and temperature, potentiodynamic parameters may shift to higher or lower values, a fact that is related to multiple microstructural characteristics such as the volume fraction, size, spatial distribution geometric morphology and chemical composition of precipitates and depleted zones.The presence of sigma phase has been shown to deteriorate corrosion resistance, since related to Cr and Mo depletion, especially at grain boundaries [10,11].Furthermore, it has been demonstrated that Cr2N is also responsible for corrosion properties degradation, especially for the nucleation of pitting sites [12,13].Regarding the corrosion potential (Ecorr), a trend to higher values, thus improved corrosion resistance, is noted in samples aged at 850 o C. By contrast, lower Ecorr values are noted in samples aged at 750 o C, which may be attributed to the presence of only intergranular deposits of sample aged at 750 o C for 120 h, while the presence of more and new inter-and intragranular precipitates constitute a reason for corrosion susceptibility of sample aged at 750 o for 1000 h.Finally, the presence of intragranular precipitates observed in the sample aged at 850 o for 24 h explains the variation in corrosion potential, which was found lower compared to that of the sample aged at 750 o C for the same time.

Conclusions
Overall, the main points of this research can be summarized as follows:  The main equilibrium secondary phases of S32654 at 750-850 o C are the σ phase and Cr2N. The intergranular σ phase is nucleated much earlier than the intragranular one due to the lower nucleation energy and higher element diffusion rate on grain boundary region. When aged at 850 °C, the cellular σ phase keeps rapid growth until 120 h and then gradually tend to be constant at the stage of long-term aging. The σ phase provides the necessary Cr and the secondary γ2 matrix supplies N for the formation of Cr2N, between the interphase boundary.
 Aging increases the steel's hardness but reduces its corrosion resistance, both due to precipitation sequences occurred.Ageing at 850°C leads to faster corrosion due to microstructural characteristics, while steel aged at 750°C maintains reasonable corrosion resistance. Hardness changes act as an indicator of overall microstructural changes, which finally affect in a more complex way the local electrochemical characteristics and finally the overall corrosion resistance.

Figure 2 .
Figure 2. Aged sample at 850 o C for 24 h.

Figure 3 .
Figure 3. Microstructure evolution at 120 h and 1000 h: (a) 750 o C for 120 h, (b) 850 o C for 120 h , (c) 750 o C for 1000 h, (d) 850 o C for 1000 h.

Figure 4 .
Figure 4. Potentiodynamic polarization curves of the heat-treated samples at 750 °C and the Ecorr comparison of all studied samples.

Figure 5 .
Figure 5. Potentiodynamic polarization curves of the heat-treated samples at 850 °C and the Icorr comparison of all studied samples.

7thFigure 6 .
Figure 6.Hardness testing evolution versus the ageing time of the studied samples.

Table 1 .
Chemical composition (in wt.%) of various phases found at the studied samples via EDS microanalysis (ND=Not Detected).