**1. Introduction**

Duplex stainless steels (DSSs), and consequently, Super Duplex stainless steels (SDSSs) consist of austenite and ferrite phases. The resultant microstructure exhibits good combinations of strength, ductility, and corrosion resistance, since it takes advantages of the single-phase counterparts [1], with the main difference from more common austenitic stainless steels [2,3]. The steels not only inherit the mechanical properties of the completely ferritic or completely austenitic alloys, but they also exceed them. A factor of economic importance is the low content of expensive nickel, usually 4–7% compared with 10% or more in austenitic grades, as a result of which the life cycle cost of the DSSs is the lowest in many applications [4–7]. In DSSs, the two structure components b.c.c α-ferrite and f.c.c. γ-austenite lies as crystals of the same size statistically distributed next to each other [8]. However, adjustment of the two-phase microstructure of duplex stainless steels is complicated because a balanced phase ratio does not only depend on alloy components [9]. Indeed, the decomposition of ferrite to austenite can occur over a wide temperature range. This phenomenon can be understood on the basis that the duplex structure is quenched from a higher temperature, at which the equilibrium fraction of α-ferrite is higher. There appear to be three mechanisms by which austenite can precipitate within α-ferrite grains: By the eutectoid reaction, as Widmannstätten precipitates, and via a martensitic shear

process [9–11]. Martensitic transformation in solids provides an unusual mechanical behavior ranging from the superelastic behavior typical of shape-memory alloys to non-thermoelastic behavior where the transformation induced plasticity (TRIP) phenomenon allows the development of steels with a good compromise between ductility and toughness [12–15]. This typical property of TRIP-aided steels results from the strong couplings between plasticity by dislocation motion and martensitic phase transformation through the internal stresses generated by both inelastic processes. The TRIP mechanism is based on deformation-stimulated displacive transformation of a metastable former phase. In such materials, the overall behavior depends deeply on the so-called chemical energy, which leads the martensitic shear transformation. The motivation for that is twofold. First, the former phase should be sufficiently unstable such that a transformation-induced plasticity effect is initiated upon loading. Second, the former phase should be sufficiently stable that the TRIP effect occurs over a wide strain regime, specifically at high strains, where strain-hardening reserves are usually more desirable than at low strains. In order to obtain this phenomenon in DSS steels, different studies have been recently performed; designing new alloys compositions [16–18]. The aim of this work is to achieve the occurrence of the TRIP effect in SDSSs commercial alloys, exploiting the martensitic shear transformation of the γ-austenite within α-ferritic grains, just by tuning the proper thermal treatment.

#### **2. Materials and Methods**

This study has been performed on a commercial F55-UNS S32760 super duplex stainless steel, which had a chemical composition, as designed by standards and measured via optical emission spectroscopy (OES), is reported in Table 1. Samples have been drawn from a bulk ingot and thermally treated. First, a solution thermal treatment (STT) at 1573 K (1300 ◦C) for 145 s/mm has been executed, in order to erase the previous thermal and stress history of the specimens and to provide the super-saturation of the α-ferritic matrix with γ-formers elements, such as Ni, Mn, Cu, and N. Then, an annealing thermal treatment (ATT) has been performed at 1353 K (1080 ◦C) for different holding times: 36, 72, 215, 355, 710, and 1135 s/mm (Figure 1). The purpose of this heat treatment is the supply of the energy needed to trigger the γ-austenite precipitation within the α-matrix. This temperature range grants to avoid precipitation of embrittling phases within α-ferritic grains, such as σ (Fe-Cr-Mo), χ (Fe36Cr12Mo10), and nitrides (CrN and Cr2N). Both thermal treatments have been followed by water quench [19]. Afterwards, specimens have been machined into Round Tension Test Specimen shapes and tensile tests have been performed; both these operations have been executed, following the ASTM E8/E8M standard. The specimens have been treated for the microstructural examination using Beraha's tint etching (5 mL H2O, 1 mL HCl, 0.06 g K2S2O5, 0.06 g NH4FHF). Subsequently, the samples have been analyzed via stereoscopy, optical microscopy, and electron microscopy. The volume fraction of γ-austenite within the material has been calculated through automatic image analysis of 10 micrographs measuring 1 mm<sup>2</sup> randomly taken on the samples, following ASTM E1245 standard. Electron microscopy has been used in order to obtain morphological information via Secondary Electron (SEM/SE) imaging, chemical composition data through Energy-Dispersive X-ray Spectroscopy (SEM/EDS), and crystallographic data by Electron Backscatter Diffraction (SEM/EBSD) analysis. The beam spot has 1 μm radius. The chemical composition data have been obtained, averaging five measures. The Electron Backscatter Diffraction (SEM/EBSD) analysis has been calculated through the software INCA provided by Oxford Instruments (INCA Oxford Instrument, Oxford, UK). The crystallographic data have been used to highlight the influence of grain boundary distribution on the onset of secondary recrystallization. The presence of special Coincidence Site Lattice (CSL) boundaries between primary and secondary grains the development of recrystallization.

**Table 1.** UNS S32760 super duplex stainless steel chemical composition (expressed in wt.%), as designed by standards and measured via optical emission spectroscopy (OES).

**Figure 1.** Diagram representing the thermal treatments parameters.
