475°C Embrittlement of Duplex Stainless Steel - A Comprehensive Microstructure Characterization Study

C. Örnek1,2,3, M. G. Burke1,2, T. Hashimoto2, J.J.H. Lim1 & D. L. Engelberg1,2,3

1Materials Performance Centre,School of Materials, The University of Manchester,

Sackville Street, Manchester, M1 3AL, United Kingdom

2Corrosion and Protection Centre,School of Materials, The University of Manchester,

Sackville Street, Manchester, M1 3AL, United Kingdom

3Research Centre for Radwaste & DecommissioningSchool of Materials, The University of Manchester,

Sackville Street, Manchester, M1 3AL, United Kingdom

Email: cem.oernek@manchester.ac.uk, tel.: +44 161 306 4838 (corresponding author)Email: joven.lim@manchester.ac.uk, tel.: +44 161 306 4838

Email: M.G.Burke@manchester.ac.uk, tel.: +44 161 306-4858Email: T.Hashimoto@manchester.ac.uk, tel.: +44 161 306 5938Email: dirk.engelberg@manchester.ac.uk, tel: +44 161 306 5952

AbstractThe effect of 475°C embrittlement on microstructure development of grade 2205 duplex stainless steel has been investigated. Spinodal decomposition products and associated precipitates in ferrite, austenite, and at interphase boundaries were characterized using analytical Transmission Electron Microscopy (TEM) and Scanning Electron Microscopy (SEM) techniques. Micro-analyses confirmed the presence of Cr-enriched α and Cr-depleted α spinodal structures in the ferrite after 5 hours of aging at 475°C. Long-term aging for 255 hours resulted in heavily-faulted R-phase precipitates with sizes of ~50 – 400 nm, χ-phase, and ε-Cu in the ferrite, TiN and Cr2N precipitates in the austenite, and a continuous network of M23C6-carbides at interphase boundaries. A significant hardness increase was observed after 255 hours of aging, which was accompanied by a reduction of ferrite fraction. X-ray diffraction (XRD) stress measurements showed a general reduction of residual stresses in both ferrite and austenite with aging. Electron Backscatter Diffraction (EBSD) showed increased local misorientations, primarily close to precipitate interfaces within the ferrite, indicating the development of strain heterogeneities in the microstructure. The data presented provides a better understanding of 475°C embrittlement in duplex stainless steel, suggesting that not only the ferrite alone is responsible for embrittlement. A comprehensive microstructure characterization study has been provided and the explanation for 475°C embrittlement of duplex stainless steel has been discussed.

KeywordsDuplex stainless steel; 475°C embrittlement; Microstructure characterization; Transmission electron microscopy; Scanning electron microscopy; X-ray diffraction; Electron backscatter diffraction; Spinodal decomposition; Electron energy loss spectroscopy; Secondary phases

IntroductionDuplex stainless steels (DSS) have now been used for various industry applications, due to their excellent mechanical properties and good corrosion resistance [1-3]. The DSS

microstructure typically consists of a balanced ratio of ferrite (δ) and austenite (γ), providing mechanical strength, ductility, and fracture toughness, with the latter a result of the chemical composition and the small grain size, typically in the order of 5 – 10 µm. The ferrite to austenite ratio, therefore, has a significant influence on the mechanical behavior [4], with strength and creep resistance governed by the ferrite, whereas ductility and toughness are controlled by the austenite [5]. However, service temperatures between 250 – 550°C can limit the application of DSS due to microstructure embrittlement reactions [5-10]. The highest rate of embrittlement in 2205 DSS has been reported at 475°C; and this phenomenon is therefore referred to 475°C embrittlement or low-temperature embrittlement in the literature [5, 8, 9].

Low-temperature embrittlement in the 250 – 550°C temperature range is associated with microstructure transformations on sub-micron scale. This is often attributed to phase separation reactions occurring in the ferrite phase, and the volume fraction of ferrite is, therefore, believed to be the life-limiting factor [8-11]. Embrittlement is mostly ascribed to spinodal decomposition or the formation of α’’ precipitates after long-term aging treatments [8, 12-16]. However, additional phases can also form in the decomposed ferrite such as G-phase, carbides, nitrides, secondary austenite, χ-, π-, τ-, I-, S-, J-, Z-, R- and Laves-phases [6, 8-12, 17-27]. These may also have significant impact on mechanical properties. Formation of α’ + α’’ is typically associated with an increase of hardness, loss of magnetic permeability, and direct observation of such changes have now become routine practice [28-31]. Moreover, decomposition reactions and phase transformations in austenite similar to that of ferrite could also occur, due to an existing miscibility gap in the Fe-Ni binary system [32-36]. 475°C embrittlement may, therefore, be caused by a variety of phase transformations occurring in ferrite and in austenite to some extent, if any. However, compared to the austenite, the ferrite is more prone to phase transformation due to the more than 1oo-times faster diffusivity of alloying elements [37]. Most investigations to characterize 475°C embrittlement have therefore focused on phase transformation kinetics.

Spinodal decomposition is a phase reaction mechanism describing the segregation of an entire phase into two distinct phases with distinguishable chemical composition and physical properties, but similar crystallographic parameter. Ferrite, with its high Cr content, can undergo phase separation and transformation reactions due to the miscibility gap in the Fe-Cr binary system. Ferrite is not stable within the miscibility gap and decomposes into two decomposition products, BCC-α’ and BCC-α’’. Other secondary phases can be formed, causing further decomposition of ferrite. Unfortunately, no common denotation for this phenomenon exists in the literature. Some authors stated the reaction as δ α + α’, while others stated as α α’ + α’’, or α α1 + α2. In both cases, α’ can easily be confused with martensite. To avoid misinterpretation, it is therefore important to define the structure and composition of these reaction products, in the form of Fe-rich α’ and Cr-rich α’’.

Amongst the large variety of phases identified that can be formed in the 475°C embrittlement temperature window, the intermetallic R-, G-, tau (τ)- and chi (χ)-phases, secondary austenite, carbides and nitrides are together with the formation of α’ mainly believed to be responsible for the embrittlement. G-phase (molybdenum silicide) precipitates are reported to be of globular morphology with sizes between 20 – 50 nm according to Mateo et al. [20], 2 – 7 nm according to Hamaoka et al. [38], 8 – 11 nm according to Danoix et al. [39], and <50 nm according to Shiao et al. [22] having a FCC-structure with Fm3m space group and lattice parameters with a = 1.09 nm, b = 1.12 nm, and c = 1.14 nm [40]. Frank-Kasper-R-phase precipitates belong to the tetrahedrally close-packed phases with rhombohedral unit cell (reverse hexagonal lattice). Their lattice parameters are reported to be a = 0.901 nm and α =

74.5°. They usually assume a disc-shaped or lenticular morphology with sizes between 50 – 400 nm and are enriched in Mo and Si [8, 12, 13, 19, 21, 24, 41, 42]. The precipitation of τ-phase occurs by a diffusional process and exhibits a needle-like appearance with Fmmm space group and lattice parameters with a = 0.4054 nm, b = 0.3436 nm, and c = 0.2867 nm.

Secondary austenite formation from ferrite is also reported to occur in a diffusion-less manner, via a double-shear nucleation process [12, 42], which can assume Widmannstätten morphology [22]. Shiao et al. distinguished three different mechanisms of secondary austenite formation (γ2), also called as ‘new austenite’ (γnew), which can have different morphological appearance with possibly different chemical composition [22]. One is formed on ferrite/ferrite grain boundaries via nucleation and growth by the consumption of ferrite with usually a band-like shape forming an entire austenite network on the ferritic grain boundary. Another is formed on the interface between ferrite and austenite phase boundaries by a direct reversion of ferrite to austenite due to elemental redistribution with usually Ni depletion around the newly formed austenite and a precipitation-free zone. In the ferrite grain-interior, another form of secondary austenite can be formed having Widmannstätten structure [22]. Widmannstätten structure is well-known in ferritic steels causing severe embrittlement [43].

The work reported in this paper provides a detailed microstructure analysis of ferrite and austenite in a grade 2205 duplex stainless steel with 475°C embrittlement heat treatment exposures. The observed microstructure characteristics have been correlated with hardness measurements, residual stress profiles, and EBSD measured local strain components.

ExperimentalA grade 2205 duplex stainless steel plate with a composition (in wt.-%) of 22.4Cr, 5.8Ni, 3.2Mo, 1.5Mn, 0.4Si, 0.016C, 0.18N and Fe (bal.) has been used in this investigation. The plate was in a solution-annealed condition, referred to ‘as-received’ in this work. Rectangular coupon specimens were cut from the as-received plate and then heat treated at 475 ± 5°C for 5, 20, 50, and 255 hours followed by a water-quench. The surface of the specimens was ground to 4000-grit using SiC paper, followed by fine polishing up to 0.1 µm using diamond paste, and finalized with an OP-S active oxide polishing suspension treatment.

Microstructure CharacterizationMicrostructure characterization consisted of optical microscopy, SEM, EBSD, and TEM analyses. One sample of each heat treatment condition was used for microstructure characterization and measurement of mechanical properties. Optical microscopy analysis was performed to obtain information about the DSS microstructure appearance, and was carried out on the as-received condition only. The sample was etched using Beraha II dye-etch solution, which consisted of 800 ml deionized water, 400 ml hydrochloric acid, 48 g ammonium hydro-difluoride, and 1 g potassium disulfide.

EBSD was employed to characterize the grain size, austenite and ferrite phase fraction, and local misorientation (LMO). An FEI Quanta 650 SEM interfaced with a Nordlys EBSD detector from Oxford Instruments with AZtec V2.2 software was used for data acquisition. The experimental parameters for EBSD analysis included a step size of 0.15 µm over an area of typically 856 x 746 µm² with an accelerating voltage of 15 – 20 kV. High resolution imaging and EBSD analysis was also performed in an FEI Magellan, interfaced with a Nordlys EBSD detector from Oxford Instruments. The experimental parameters for the latter

analyses included an accelerating voltage of 15 kV and a step size of 10 – 20 nm of an area covering fine ferrite grain-interiors (<5 µm). Data post-processing was performed using HKL Channel 5 software. High-Angle Grain Boundaries (HAGB’s) were defined with misorientation ≥15° and LAGB’s between >1° and <15°. The grain size was determined by the mean linear intercept method as the mean of the vertical and horizontal directions (twins disregarded). Local misorientation (LMO) maps were generated by using a 3x3 binning and a 5° sub-grain angle threshold. This analysis gives the average LMO for a misorientation below the pre-determined sub-grain angle threshold, and can be used to locate regions with higher concentrations of misorientation in microstructures. The latter is typically associated with local micro-deformation in the form of elastic and plastic strain, due to the presence of dislocations [44].

A detailed microstructure analysis was carried out with thin foil specimens in the TEM; by electropolishing 80 – 100 µm thick 3.0 mm diameter disc samples in Tenupol-5 twin jet polished (Struers, Denmark) with a Jubalo closed cycle refrigeration system. The samples were electropolished at 20 kV in an electrolyte of 20% perchloric acid in methanol at a temperature of -40°C. All TEM specimens were subsequently examined in an FEI Tecnai F30 300 kV Field Emission Gun analytical TEM equipped with an Oxford Instruments Xmax80 SDD EDX and an AZtec analysis system. Structural analysis of spinodal decomposition and precipitates were carried out by X-ray diffraction analysis in a 4k x 4k Ultrascan CCD camera.

The sample aged for 255 hours at 475°C was analyzed first in a Philips CM20 200 kV LaB6

TEM, and then in a FEI Tecnai F30 at 300 kV and in a spherical aberration-corrected FEI Titan G2 80-200 analytical TEM with Super X EDX (ChemiSTEM™) technology, operated at 200 kV and equipped with a GIF Quantum 956 electron energy loss spectroscope (EELS). Electron Energy Loss Spectroscopy (EELS) in STEM on the FEI Titan G2 have been used to map regions of interest to provide compositional analyses of spinodal decomposition and second phase precipitates.

Hardness TestingHardness testing (HV30) was conducted to assess the overall hardness change with aging at 475°C. Micro-hardness (HV0.1) and nano-hardness (GPa) measurements were also carried out to obtain information about hardness change in the austenite and ferrite, respectively. Hardness measurements were carried out on a Vickers hardness device (Georg Reicherter Briviskop 187.5) with a load of 30 kg (HV30). For each sample, ten hardness indentations were made, and the arithmetical mean with the standard deviation calculated. The tested samples had a 600-grit ground surface.

Micro-hardness measurements were conducted with the Struers Duramin micro-hardness tester (Ballerup, Denmark) with a Vickers indenter and a test load of 0.098 N on ¼ µm polished surfaces. A total of 50 hardness measurements were obtained from each specimen: 25 measurements in the austenite and 25 measurements in the ferrite phase. The arithmetic means of all micro-hardness results were calculated.

Nano-indentation measurements were performed on samples polished to 100 nm followed by an OP-S end-polish using the Nano Indenter XP from MTS Nano Instruments. A Berkovich indenter with three-sided pyramid geometry and a 75 nm radius of the tip was used to keep the projected area-to-depth ratio constant. 10 x 10 indentations with a penetration depth of 200 nm were made in a serpentine array on each specimen. From each indenting, local nano-

hardness values at maximum load were obtained. The tested area was analyzed by SEM and EBSD to ascribe each indenting result to either ferrite or austenite. Indentations which did not penetrate into the interphase regions were accepted only. In total, 30 – 40 indentations were obtained for each microstructure phase.

X-ray Diffraction Stress MeasurementA Proto iXRD Combo X-ray Diffractometer (Proto Manufacturing Inc., Michigan, USA) was used to measure surface residual stress profiles, with typically ± 20 MPa error band for all data [45, 46]. The testing machine is equipped with a two detector system, and a schematic illustration of the measurement setup is given in reference [47]. Cr and Mn x-ray tubes were used to measure the strain in ferrite and austenite, respectively, with a penetration depth of approximately 10 – 15 μm (99% absorption by workpiece at Ψ = 0°) [46, 48, 49]. Sin-2Ψ measurements were performed at eleven Ψ tilt angles (β angles on the Proto iXRD) with a 3° oscillation between at each β angle. The reader is referred to reference [47] for a detailed procedure description and all further setup parameters.

Prior to the tests, the x-ray diffractometer was calibrated to determine the zero stress position. Stress-free and pre-stressed standard samples were used for calibration. The multiple exposure technique was used and the inter-planar d-spacing of the respective planes measured. Two x-ray measurement orientations, i.e. 0° and 90° φ (Phi) angles, in all process orientations of each specimen were chosen, which was in alignment with the microstructure orientation. Each measured orientation corresponds to the stress direction as specified in Figure 1.

Phase Fracture Analysis using X-ray Diffraction and Electron Backscatter DiffractionPhase fraction analysis was carried out on the as-received and all aged specimens using x-ray diffraction (XRD) and EBSD analysis. EBSD phase fraction analyses were performed on large areas covering at least 2000 grains. XRD analyses were performed on a Bruker D8 Discover. The scan range was 40.0-130.0° 2θ with a step size of 0.02° 2θ. Co-Kα1 x-ray source with 35 kV accelerating voltage and 40 mA current was used. From the obtained XRD profiles the required lattice parameters for ferrite and austenite were selected from the ICDD PDF cards 04-007-9753 (ferrite) and 04-002-3692 (austenite), which were then read into Topas V4.2.0.2 analysis software. For data refinement the lattice size, scale, peak fitting type, atomic coordinates, temperature factors, preferred orientation etc. were adjusted to achieve best fit. The Rietveld weight percent (RWP) value was kept as low as possible, and all results are to the best of ± 2% precision.

Results and Discussion

Microstructure Characterization

As-received MicrostructureThe as-received microstructure consisted of 46 ± 1% ferrite and 54 ± 1% austenite with an average grain size of 6-8 µm for both phases (Figure 1 & Figure 2). The grains were elongated along RD, shown in the 3D representation of the microstructure in Figure 1. No secondary phases were observed. The austenite contained some dislocation structures with stacking faults and annealing twins, whereas the ferrite had intense dislocation networks and sub-grains (Figure 3).

Microstructure aged for 5 hoursSpinodal decomposition in the ferrite was observed after 5 hours aging, while neither precipitation nor chemical segregation in the austenite was apparent (Figure 4). In ferrite, spinodal decomposition was seen, shown in Figure 4. A spinodal arm width of 1 – 3 nm was measured which defines the size of α’ and α’’. Similar observation on a 2205 DSS has been reported by Weng et al. who also observed spinodal decomposition in the ferrite already after 2 hours aging at 475°C [16]. In contrast, Sahu et al. [11, 17] and Iacoviello et al. [50] did not observe spinodal decomposition, but α’’ precipitation in a 2205 DSS, whereas Nilsson et al. [19] observed a mottled contrast in the ferrite associated with spinodal decomposition in a DSS with 8 wt.-% Mo. Seemingly, the ferrite of grade 2205 DSS can be susceptible to both spinodal decomposition or precipitation of α’ and α’’. This suggests that the chemical composition of grade 2205 DSS is at or close to the chemical spinodal line which defines the border line in the Fe-Cr binary [51]. Slight variation in the chemical composition can apparently cause a great impact on the phase separation mechanism in the ferrite.

Microstructure aged for 20 hoursA distinctive increase of the dislocation density in both the ferrite and austenite was observed after 20 hours aging, summarized in Figure 5a+c. The ferrite contained dislocation forests and loops, whereas annealing twins and numerous dislocations were seen decorating the austenite. The mottled contrast in the ferrite became more apparent in comparison to the 5 hours-aged condition, with average sizes of 2 – 4 nm measured for Fe-enriched α’ and Cr-enriched α’’ (Figure 5b).

Microstructure aged for 50 hoursSecondary phases were observed in ferritic regions and at interphase boundaries after 50 hours aging, as shown in the SEM micrographs in Figure 6. The size of these precipitates was in the order of 50 – 150 nm. No precipitates in the austenite were observed. Weng et al. [16] did not report any secondary phase after aging for 64 hours at 475°C, but showed spinodal decomposition with sizes of ~5 nm for α’ and α’’. Sahu et al. [11, 17] reported coarsened α’’ precipitates only after aging for 100 hours at 475°C, whereas Nilsson et al. [19] showed Cr-Mo-rich χ-phase precipitates in the ferrite on interphase boundaries after 100 hours aging at 400°C. Redjaimia et al. also reported χ in a 2205 DSS in the ferrite with similar morphology and composition, but for conditions aged for 650°C for 336 hours. The morphological appearance may suggest the particle in Figure 6b at the interphase boundary as χ.

Microstructure aged for 255 hoursThe microstructure after 255 hours aging is shown in Figure 7. Numerous precipitates could be observed with high-resolution SEM, with spinodal decomposition becoming also apparent. The spinodal structures had a size of 14 – 29 nm, with α’ showing enrichment in Fe, Ni, and Mn and depletion in Cr, whereas the opposite was detected for α’, shown in Figure 8 and Figure 9. Precipitation was heterogeneous with ferritic sub-grain boundaries and slip planes being preferential nucleation sites. Some dislocation structures and stacking faults were observed in the austenite (Figure 10b). There seemed to be one major precipitate decorating the ferrite grain-interiors which had a platelet-like or disk-shaped morphology (Figure 10, Figure 11, and Figure 12). Selected area diffraction patterns (SADP) were obtained from a ferritic area containing platelet-shaped precipitates and showed an orientation relationship with 3-times larger unit cell and a growth direction along <111>. EDX spectra over such particles revealed Mo and Si enrichment, suggesting this to be R-phase (Figure 13). EELS chemical analysis showed enrichment in Cr, Ni, Mn, and Si (Figure 17). The diffraction

pattern and chemical composition are in good agreement with the works from Redjaimia et al. [12] and Nilsson et al. [19] who confirmed the presence of R in duplex stainless steel. TEM dark-field analysis showed that these disk-shaped particles were heavily faulted. This was also in line with the work of Redjaimia et al. [12] who reported a heavily faulted internal structure of R.

A continuous layer of a phase reaction product was seen decorating all interphase boundaries as well as some ferrite-ferrite grain boundaries (Figure 10). Selected area diffraction (SAD) over such regions showed a 2 – 2.5-times larger unit cell size with a cubic structure (Figure14) indicating M23C6-type carbide. Chemical analysis revealed enrichment in Cr, Mo, Ni, and Si and depletion in Fe (Figure 15 and Figure 9) with a heavily-faulted structure, as shown High-angle annular dark field (HAADF) images in Figure 14. Further EELS analyses showed particles in ferrite containing Cu with very little alloying elements (Figure 17). Cu-precipitates in stainless steels are typically ε-Cu with face-centered-cubic crystal lattice structure, and are typically found in the ferrite [52]. Dyja et al. observed ε-Cu in a 25Cr-7Ni cast super duplex stainless steel after aging at 480°C for 4 hours and reported coherency with the host lattice for small sizes of ε-Cu particles (~5 nm) and incoherency for larger sizes (~10 nm) [53]. The morphology of ε-Cu particles have been reported as spherical and seemed to form in the ferrite only [52, 53]. The size of these particles was less than 5 nm suggesting coherency with the matrix.

Some precipitates in the austenite were also observed indicating two different secondary phases. Spherical or disk-shaped features were seen in austenite in the vicinity of interphase boundaries, as shown in Figure 8c-d. Their morphological appearance suggests chromium nitride (Cr2N). Horvarth et al. observed spherical Cr2N precipitation in austenite in a 2205 and 2507 duplex stainless steel after aging at 300 – 500°C with a distinctive micro-hardness increase [36]. There is a preferential partitioning of nitrogen in austenite reported which increases with aging, facilitating the precipitation of chromium nitrides [9, 18, 36, 54, 55]. Shi et al. also reported Cr2N and R-phase precipitates in a duplex stainless steel after aging at 500°C for 5 months which is in good agreement to those shown in Figure 8c-d [56]. Further pillar-shaped precipitates in austenite were observed, having a length of ~300 nm and a width of ~60 nm (Figure 16). TEM-EDX analysis showed enrichment in Ti and N and a major depletion in Fe, Ni, and Cr (Figure 16). Diffraction pattern was obtained from high-resolution TEM imaging and the Fast-Fourier Transform (FFT) revealed a face-centered-cubic lattice with (111) zone axis for this precipitate (Figure 16d). The parent austenite grain was also in zone axis with (111) orientation showing a cube-on-cube orientation relationship of the austenite with the precipitate. The crystal structure and chemical composition suggests this precipitate as titanium nitride in the type of TiN. Titanium has a significantly higher affinity of C and N in the stainless steel microstructure which is typically added to improve the resistance to sensitization of stainless steels [9, 57]. TiN precipitation in duplex stainless steel is a rare observance, and the fact that Ti was not an alloying element but can contain some traces of refractory metals showed clearly the high affinity of Ti to bound N. To summarize, the microstructure aged for 255 hours at 475°C contained spinodally decomposed α’ and α’’, R-phase, M23C6 carbide, χ-phase, and ε-Cu in the ferrite and Cr2N and TiN in the austenite, with M23C6 carbides decorating interphase boundaries.

Hardness DevelopmentThe macro-hardness results are shown in Figure 18. The as-received hardness was 272 ± 6 HV30. The hardness increased slightly with aging to 286 ± 2 HV30 after 50 hours, and then

increased sharply after 255 hours to 342 ± 5 HV30. This rise in hardness is typically associated with embrittlement which is in good agreement with other work [11, 17, 58, 59]. Micro-hardness measurements in ferrite and austenite as a function of aging time are shown in Figure 19. The micro-hardness of ferrite and austenite were similar in their as-received conditions showing a larger scatter. The hardness of ferrite increased from 273 ± 10 HV0.01 to 336 ± 26 HV0.01 after 255 hours of aging, whereas no change occurred in austenite. Apparently, the increase of hardness of the ferrite phase is therefore responsible for the hardness increase of the bulk microstructure. This is due to the precipitates and spinodal decomposition formed in the ferrite phases impeding dislocation mobility. An increasing hardness of the ferrite with 475°C embrittlement is typical for duplex stainless steels, while the hardness of austenite often remains unchanged [56, 60-62]. Nano-hardness measurements are shown in Figure 20. The trend of the nano-hardness of ferrite and austenite was similar to the trend seen for the micro-hardness. A significant increase of hardness was observed in ferrite with aging. Ferrite is typically responsible for the hardness increase in duplex stainless steels with aging in the 475°C temperature window [61, 63, 64], and it has been shown that this is more pronounced for larger ferrite fractions [4]. The large scatter of the nano-hardness measurements is due to different grain orientations and precipitation heterogeneity, and therefore, the scatter in ferrite is larger than in austenite.

XRD Stress DevelopmentXRD stress measurements for austenite and ferrite are shown in Figure 21 and Figure 22, respectively. The stress in austenite and in ferrite in the as-received condition was compressive. Overall, the entire microstructure seemed to possess more compressive surface stress components. No significant change occurred after 5 hours aging. After 20 hours aging, large compressive stresses developed in both austenite and ferrite. The compressive stresses in austenite further increased with a likewise stress development in ferrite. It seemed that the entire microstructure became more compressively stressed which might be related to the spinodal decomposition reaction in the ferrite. Spinodal decomposition may have caused negative stress fields due to local chemical strain misfit arising from elemental concentration fluctuations leading to slight expansion of the entire lattice. Expansion of the ferritic matrix lattice may have also affected austenite grains leading to a compressive stress development in austenite. After 50 hours aging, both phases indicated changes in the residual stress profiles, with austenite showing more balanced stresses, but not in the TD plane. The compressive stresses in ferrite decreased significantly. However, other reactions in the austenite must have occurred simultaneously since the initial compressive nature from the mill-annealing process was reduced, possibly caused by decomposition products and/or dislocation reactions. All residual stresses in the ferrite and austenite were significantly reduced and balanced after 255 hours of aging. The nucleation of R-phase precipitates and further secondary phases seemed to have the largest effect on the stress development in the microstructure, since large stress variations occurred between 20 and 50 hours and 50 and 255 hours of aging.

XRD and EBSD Phase Fraction AnalysisThe phase fraction of ferrite and austenite in all aged specimens were analyzed by XRD and EBSD and compared to the as-received microstructure condition. The results are summarized in Figure 23. The austenite-to-ferrite ratio, measured through XRD, increased from 52:48 ± 2% to 59:41 ± 2% after 255 hours of aging, while EBSD results showed a similar increase from initially 54:46 ± 1% to 57:43 ± 1%. Thus, either secondary austenite must have formed or primary austenite must have grown on the expense of ferrite during aging. The formation

of secondary austenite in 20.9Cr-9.9Ni duplex stainless steel [22] and 308 austenitic stainless steel [65] has been reported, associated with decreasing volume fractions of ferrite during aging in the 475°C temperature window. The growth of primary austenite by the formation of new austenite was analyzed by Shiao et al. [22], reporting the consumption of secondary austenite by the growth of primary austenite grains [22]. Furthermore, Takizawa and co-workers [66] investigated the effect of ferrite content on stress corrosion cracking susceptibility on a 23Cr-1Mo-1.5Mn duplex stainless steel and observed a decrease of the ferrite content by 5 – 10% after aging at 475°C for 690 hours.

EBSD Local Misorientation MappingA LMO map with the corresponding phase map of the as-received microstructure is shown in Figure 24. The as-received microstructure had only smooth LMO variations across the entire microstructure, despite of the presence of numerous grain boundaries. A few LMO hot-spots were detected mainly on ferrite LAGB’s and HAGB’s and on small austenite grain clusters deriving from the process history of the material. The mean LMO of the entire microstructure was low, with austenite showing slightly higher degree of misorientation than ferrite. Ferrite and austenite have different thermal expansion coefficients and deformation behavior. During quenching, austenite can distort easier than ferrite which may lead to localized strain development in austenite¸ particularly in small grain clustered regions. Ferrite cannot bear large thermal stresses during quenching, hence, small in-grain distortions can cause localized strain development at grain boundaries. Aging at 475°C for 255 hours caused an increase of the number of LMO hot-spots in ferrite. The LMO map and corresponding EBSD phase map for the aged microstructure are shown in Figure 25 and Figure 26. The mean LMO shifted towards higher values indicating an increase in local strain. The phase transformation products formed in austenite seemed to have promoted the formation of discrete LMO hot-spots at ferrite boundaries. Large LMO hot-spots in ferrite, indicative to strain, were detected at precipitates and at grain boundaries. The LMO from ferrite across the precipitates was 1 – 1.8°, while across spinodally decomposed area in ferrite 0.1 – 0.6° was measured. This showed clearly a larger effect of microstructure strain caused by precipitates than by the spinodal decomposition.

ConclusionAgeing treatments of 2205 duplex stainless steel at 475°C resulted in a number of microstructure changes, which directly affected mechanical properties.

(i) Spinodal decomposition, consisting of Fe-rich α’ and Cr-rich α’’, were observed in ferrite, with the spinodal structures showing an increasing width of Cr and Fe enriched zones with longer aging time.

(ii) ε-Cu particles, χ-phase, and a continuous Cr-Mo-Ni-Si-rich network of precipitates, proposed to be M23C6, decorating interface boundaries and ferrite-ferrite boundaries, were observed after 255 hours aging in the ferrite, with Cr2N and TiN precipitates found in the austenite.

(iii) R-phase precipitates with disc- and lenticular-shape were detected in ferrite after 50 hours aging, which grew further from ~50 nm to 200-400 nm after 255 hours of 475°C exposure.

(iv) EBSD and XRD analysis confirmed a reduction of ferrite content with longer aging exposures, accompanied by an increase in the austenite content.

(v) Macro-, micro-, and nano-indentation measurements showed hardness changes with ageing for 255 hours, as a result of the reported phase precipitation reactions in the ferrite.

(vi) XRD stress measurements revealed a reduction of compressive stresses in both ferrite and austenite, resulting in a more balanced microstructure stress condition after 255 hours of aging.

(vii) EBSD analysis indicated the presence of higher local misorientation distributions, indicative of plastic strain, in the vicinity of precipitate interfaces within the ferrite.

(viii) The data showed that phase reactions occur in ferrite, austenite, and at interphase regions, hence, suggesting that not only the ferrite alone is responsible for the occurrence of microstructure embrittlement.

AcknowledgementThe authors acknowledge Radioactive Waste Management, a wholly owned subsidiary of the Nuclear Decommissioning Authority (NPO004411A-EPS02), and EPSRC (EP/I036397/1) for financial support. The authors are grateful for the kind provision of Grade 2205 Duplex Stainless Steel plate by Rolled Alloys. Special thanks to Gary Harrison, School of Materials, for his support during XRD measurements and analysis. The authors are also thankful to Dr Arne Janssen, University of Manchester, for valuable discussion for the diffraction pattern analysis.

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Figures

Figure 1: EBSD phase map showing the as-received microstructure of 2205 DSS in all process orientations.

Figure 2: Optical micrographs of grade 2205 DSS show the microstructure in (a) rolling direction (RD) and (b) in normal direction (ND).

Figure 3: TEM bright-field images of the as-received microstructure: (a+b) dislocation structures in austenite, (c+d) dislocation networks in ferrite.

Figure 4: TEM bright-field images of the microstructure aged for 5 hours at 475°C: (a) dislocation networks in austenite, (b) spinodal decomposition of ferrite with image insert showing diffraction pattern with (110) zone axis.

Figure 5: TEM bright field images of the microstructure aged for 20 hours at 475°C: (a) dislocation network with dislocation loops in a ferritic subgrain, (b) spinodal decomposition of ferrite, (c) dislocation networks in austenite, (d) annealing twin boundary with dislocations in ferrite [(111) zone axis].

Figure 6: SEM images of the microstructure aged for 50 hours at 475°C: (a) overview, (b+c) some precipitates formed within the ferrite and at interphase boundaries.

Figure 7: SEM images of the microstructure aged for 255 hours at 475°C: (a) overview, (b-d) multiple precipitation occurred in ferrite and at interphase boundaries. Note the spinodally-decomposed ferritic structure which has as a mottled, intertwined appearance.

Figure 8: TEM bright-field analysis of the microstructure aged for 255 hours at 475°C: (a) spinodal decomposition of ferrite (mottled contrast) with platelet-shaped precipitates, (b) diffraction pattern with (100) zone axis, (c) two precipitates formed in the austenite with the image insert showing the diffraction pattern of the austenite with (110) zone axis, (d) higher magnified view of the precipitate in the austenite as pointed by the arrow in (c).

Figure 9: EELS reconstructed spectrum images of the Fe, Cr, Ni, and Mn measured on the microstructure aged for 255 hours at 475°C showing Cr-enriched α’’ in the ferrite with depletion in Fe, Ni, and Mn. Chemical segregation of Cr, Ni, and Mn on the interphase boundary can also be seen.

Figure 10: TEM bright-field imaging of the microstructure aged for 255 hours at 475°C: (a) overview, (b) dislocation structures and stacking faults in austenite with multiple precipitation occurred in ferrite and interphase boundaries, (c) numerous precipitates (R and τ) arrayed preferentially along subgrain boundaries and slip planes, (d) selected area diffraction pattern within the ferrite containing R and τ phase.

Figure 11: TEM analysis of the microstructure aged for 255 hours at 475°C: (a) bright-field and (b) dark-field image of a ferrite grain, (c) convergent beam diffraction pattern and (d) selected area diffraction pattern of ferrite with (110) zone axis.

Figure 12: Detailed TEM analysis of precipitation in the ferrite of the microstructure aged for 255 hours at 475°C: (a) overview bright-field image showing a ferrite grain with (b) the corresponding dark-field image, (c) bright-field image of region 1 and 2, (d) SADP of region 1, (e) SADP of region 2, (f) SADP of region 3, (g) bright-field image of region 4, (h) SADP of region 4, (i) bright-field image of region 5, (j) SADP of region 5.

Figure 13: TEM-EDX analysis of the precipitates in the microstructure aged for 255 hours at 475°C showing Mo and some Si enrichment.

Figure 14: TEM analysis of the microstructure aged for 255 hours at 475°C: (a) bright-field and (b) dark-field images showing precipitates in the ferrite and at interphase boundaries, (c) selected area diffraction pattern of the ferrite, (d) selected area diffraction pattern of the interphase boundary.

Figure 15: TEM-EDX analysis on a ferrite-ferrite grain boundary in the microstructure aged for 255 hours at 475°C showing Cr, Mo, Ni, and Si enrichment.

Figure 16: TEM analysis in the austenite of the microstructure aged for 255 hours at 475°C: (a) overview showing a pillar-shaped TiN precipitate in the austenite (bright-field image) with a diffraction pattern of the austenite with (111) zone axis with (c) showing a corresponding dark-field image, (d) high-resolution image of the TiN with the image insert showing an FFT diffraction pattern.

Figure 17: EELS reconstructed spectra obtained from a ferrite grain of the microstructure aged for 255 hours at 475°C: Cr-Ni-Mn-Si enrichment of the platelet-shaped precipitate with a few nano particles enriched in Cu and Mn can be seen. Fe-enriched α’ and Cr-enriched α’’ can also be seen.

Figure 18: Macro-hardness test results.

Figure 19: Micro-hardness test results.

Figure 20: Nano-hardness test results.

Figure 21: XRD surface stress results of austenite as a function of process orientation and aging time.

Figure 22: XRD surface stress results of ferrite as a function of process orientation and aging time.

Figure 23: Phase fraction analysis as a function of aging time at 475°C: (a) XRD and (b) EBSD.

Figure 24: (a) EBSD phase map of the as-received microstructure. Austenite is shown in blue and ferrite in ferrite, (b) corresponding LMO map. A few weak strain heterogeneities can be seen located mainly on ferritic grain boundaries and on fine austenite grain cluster regions.

Figure 25: (a) EBSD phase map of the microstructure after 255 hours of aging: Austenite is shown in blue and ferrite in red, (b) corresponding LMO map: There are numerous LMO hot-spots visible primarily in ferrite with a lesser extent in fine austenite grains.

Figure 26: (a) EBSD phase map of the microstructure after 255 hours of aging: Austenite is shown in blue and ferrite in red, (b) corresponding LMO map: Large LMO hot-spots, indicative to strain, detected on precipitates and on grain boundaries. LMO from ferrite across the precipitates was 1 – 1.8° while across spinodally decomposed area in ferrite this was 0.1 – 0.6°. Step size was 10 nm, (c) band contrast map: Shows the pattern quality. Note the faint contrast difference in ferrite showing spinodal decomposition products, α’ and α’’.