# Microstructural evolution of super duplex stainless steel during final annealing: The importance of the grain structure

D. Tresallet, H.P. van Landeghem, B. Langelier, H.S. Zurob, F. Krajcarz, C.

Tassin, J.D. Mithieux, M. Veron

► To cite this version:

D. Tresallet, H.P. van Landeghem, B. Langelier, H.S. Zurob, F. Krajcarz, et al., Microstructural evolution of super duplex stainless steel during final annealing: The importance of the grain structure. Materials Today Communications, 2024, 39, pp.108899. 10.1016/j.mtcomm.2024.108899. hal-04727347

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# Microstructural evolution of super duplex stainless steel during final annealing: The importance of the grain structure

D. Tresallet<sup>a</sup>, H.P. Van Landeghem<sup>a, *</sup>, B. Langelier<sup>b</sup>, H.S. Zurob<sup>c</sup>, F. Krajcarz<sup>d</sup>, C. Tassin<sup>a</sup>, J.D. Mithieux<sup>d</sup>, M. Veron<sup>a</sup>

## ARTICLE INFO

### Keywords:

Duplex stainless steel

Microstructure

Annealing

Coarsening

Capillarity

## ABSTRACT

Flat duplex stainless steel acquires its characteristic banded microstructure during its processing sequence, which ends with a final annealing step aimed at fine tuning the phase ratio to equal parts ferrite and austenite. Besides phase fraction, the grain structures, shapes and morphologies of both phases evolve too during this final step. In the present, it was shown that recrystallization is mostly complete at the end of heating and phase equilibration shortly after. As a consequence, chemical and elastic driving forces are rapidly depleted while the sizes and shapes of phases keep evolving, driven solely by capillarity. Using quantitative measurements that enable decorrelation of coarsening and globularization, it is shown therein that the grain structure of the material as well as the annealing temperature play a crucial role in the intensity of those two phenomena. A high grain boundary density enhances the kinetics of both. A higher annealing temperature, which leads to faster mass transport, favors the increase of the characteristic sizes of the microstructure over its morphological evolution.

<figure><img src="image_2.png" /></figure>

## 1. Introduction

Duplex stainless steel grades are a group of stainless grades whose microstructure consists of a nearly equal-part mixture of body centered cubic ferrite and face centered cubic austenite. They display a combination of both excellent mechanical properties and high corrosion resistance. Duplex stainless steel flat products all feature a characteristic banded microstructure due to the simultaneous presence of both phases throughout the manufacturing process of the alloy. This process is concluded by a final annealing step aimed at adjusting the ductility of the product and the phase ratio. During this step, multiple phenomena occur and lead to changes in the microstructure, the most prominent of

which is the pronounced morphological evolution. Before annealing, the cold-rolled metal microstructure shows fine and sharp austenite bands embedded in a ferrite matrix. Within a few minutes of annealing, those bands become thicker and more rounded. Additionally, they become considerably shorter as a result of the bands breaking up into smaller fragments.

The morphological evolution and coarsening of the microstructure during annealing of superduplex flat products have been previously noted in the literature. Keichel et al. remarked in particular that austenite bands were subject to fragmentation past certain annealing durations [2]. Wessman et al. studied the evolution of austenite band spacing during annealing and attempted to model it using, amongst other, the classical LSW theory [3]. These modeling efforts, considering solely the shape of phases and disregarding their grain structures, met only limited success. The authors closed their article wondering as to how capillarity could drive microstructure coarsening at such a rate.

There are different driving forces coexisting at the beginning of the heat treatment. First, the phase fractions do not initially match the annealing temperature equilibrium fractions, resulting in a chemical

<footnote><sup>a</sup> Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMAP, Grenoble F-38000, France</footnote>

<footnote><sup>b</sup> Canadian Centre for Electron Microscopy, Hamilton, ON L8S 4M1, Canada</footnote>

<footnote><sup>c</sup> McMaster University, Faculty of Engineering, Hamilton, ON L8S 4L7, Canada</footnote>

<footnote><sup>d</sup> Aperam Stainless, Isbergues F-62330, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: hugo.van-landeghem@grenoble-inp.fr (H.P. Van Landeghem).</footnote>
driving force. Second, the microstructure from cold rolling features a high dislocation density, inducing an elastic driving force. Finally, cold rolling also leads to a fine microstructure with high density of phase interfaces and grain boundaries. This high density of interfaces and boundaries also represent energy stored in the material that can serve as a driving force for the evolution of the duplex microstructure during annealing.

The purpose of the present paper is to investigate the interplay between all available driving forces and the microstructural changes observed during annealing of duplex stainless steel. In particular, quantitative measurements of the phase characteristic dimensions will be provided along both the rolling and the transverse directions. The detailed grain structure examination provided herein shows the importance of a comprehensive treatment of capillary forces to explain the observed microstructural changes.

## 2. Material and methods

### 2.1. Processing

A quinary Fe-Cr-Ni-Mo-N alloy was cast into a 50 kg ingot. Its composition is detailed in Table 1 and is representative of duplex grade 1.4410. This ingot was cut into billets with initial dimensions 80 × 84 × 140 mm³. These billets were then hot-rolled in two steps starting from 1553 K (1280 °C). The first step reduced the thickness to 15 mm in 4 passes during which the strip temperature dropped between 1323 K (1050 °C) and 1273 K (1000 °C). The strip was then water quenched and cut into 150 mm long sections. Those were then reheated to 1553 K (1280 °C), hot-rolled a second time in 4 passes down to a thickness of 3 mm. The temperature had dropped between 1173 K (900 °C) and 1073 K (800 °C) before it was water quenched again. This resulted in an 85% total reduction by hot rolling. The strips were solution annealed at 1373 K (1100 °C) for 2 min before one last water quench. After pickling and grinding, the strips were cold-rolled to a thickness of 1.1 mm in 21 passes, that being a 66% reduction by cold rolling.

Strips having the dimensions 120.0 (transverse direction, TD) x 20.0 (rolling direction, RD) x 1.1 mm³ (normal direction, ND) were taken from the cold-rolled sheets. Isothermal annealing treatments were carried using a Gleeble 3500-GT device (Dynamic Systems Incorporated) at temperatures ranging from 1333 K (1060 °C) to 1453 K (1180 °C). The strips were resistance-heated at a rate of 100 K.s⁻¹ in air and held isothermally for up to 300 s then water quenched. It was verified that the cooling rate was greater than 100 K.s⁻¹ across the entire temperature range and that room temperature was reached within 2 s from the quench start. The minimum duration of 0 s corresponds to immediate quench after the target temperature has been reached. The temperature was controlled using K-type thermocouples welded to the center of one of the surfaces of the strip. Additional thermocouples were attached to the sample to confirm that the strips were heat-treated at the target temperature ± 10 K over 10 mm around the control thermocouple (TD) across the full width (RD). Thus, 10.0 (TD) x 20.0 (RD) mm² samples were taken from the middle of this homogeneous area for characterization purposes.

### 2.2. Characterization

#### 2.2.1. Phase fraction

Ferrite mass fraction $w$ was determined as the ratio of the measured maximum specific saturation magnetization $\sigma_s$ to the theoretical specific

<table><caption>**Table 1**
Mean chemical composition of the Fe-Cr-Ni-Mo-N model superduplex stainless steel (wt%).</caption><thead><tr><th>Alloying element</th><th>Cr</th><th>Ni</th><th>Mo</th><th>N</th><th>Mn</th><th>Si</th><th>C</th></tr></thead><tbody><tr><td>Mean content</td><td>26.1</td><td>7.40</td><td>3.83</td><td>0.25</td><td>0.02</td><td>&lt;0.01</td><td>&lt;0.01</td></tr></tbody></table>

saturation magnetization of the ferrite $\sigma_0$ according to Eq. (1).

$$w = \sigma_s / \sigma_0 \qquad (1)$$

The maximum of the magnetic moment was measured using a Sigameter (SETARAM Instrumentation) [4]. Measurements were carried out twice for each heat-treated sample. Dividing the magnetic moment by the mass of the sample gave the specific saturation magnetization $\sigma_s$. The theoretical specific saturation magnetization $\sigma_0$ of ferrite was calculated according to Eq. (2) [5].

$$\sigma_0 = 40\pi[2.160 - 0.0275Cr - 0.033Ni - 0.026Mo - 0.026Mn - 0.060V - 0.010(C+N) \pm 0.030] \mu T.m^3.kg^{-1} \qquad (2)$$

After phase fraction measurement, heat-treated samples were further cut into smaller samples oriented along either RD or TD for microstructural analyses.

#### 2.2.2. Phase morphology and mean band thickness (light microscopy)

Phase morphology and their characteristic dimensions were analyzed using light microscopy after surface polishing and color etching. SiC grinding papers and diamond suspensions down to 1 µm in particle size were used to obtain a mirror-finish surface. Beraha I chemical etchant [6] (100 mL water, 20 mL hydrochloric acid 37%, 0.25 g potassium metabisulfite) was used at room temperature to color ferrite while better preserving the flatness of samples than other etchants, including electrolytic ones. An optical microscope DSX510 (Olympus, Japan) was used to acquire micrographs with a magnification of x1000. The DSX-BSW v.3.1.9 software was used to perform automatic 2D stitching of fields with 20% overlap and obtain images of 1000 µm along ND (full sample thickness) and 2000 µm along RD or TD.

For phase quantification, a threshold operation was performed with the Fiji v1.52o software. Micrographs were turned into 8-bit greyscale then binarized into phase images using the automated routine “Minimum” [7]. According to ASTM E1245-03 [8], the austenite fraction in volume was determined and the relative accuracy was calculated on 20 neighboring fields measuring 1000 (ND) x 100 (RD or TD) µm² constitutive of each image. This indicator is an estimate of the error on the phase fraction measurement as influenced by the field-to-field variability. For all images, the relative accuracy is lower than 1%, which confirmed that the area analyzed is sufficient for quantitative image analysis.

The mean band thickness was measured using the linear intercept method according to ASTM E1268-01 [9] coded in an automated routine through the Aphelion Dev 4.3.2 software package (ADCIS, France). The number of intersections and the edge-to-edge spacing of the bands were measured on all lines of pixels along the normal direction for higher confidence. The formalism used for the normalization of intercept distribution can be found in the supplementary material, along with considerations regarding the interpretation of the results.

#### 2.2.3. Grain boundary density and misorientation (EBSD)

The interfaces/grain boundaries were analyzed using electron backscatter diffraction (EBSD). The samples were first prepared the same way as those for optical microscopy. As EBSD requires a surface free from deformation, an additional finishing stage was performed using a 0.04 µm silica suspension. A CCD Hikari Pro detector (EDAX, USA) mounted on an Ultra 55 scanning electron microscope (FEG-SEM) (Carl Zeiss, Germany) were used with the beam voltage and current of 20 kV and 4 nA, respectively. A hexagonal grid with a 0.15 µm step size and a 5 × 5 binning were chosen as collection parameters in the TSM OIM Data Collection™ v7.0 software package (EDAX) allowing an acquisition rate of 184 fps. For each sample, an area of 355 × 355 µm² was scanned at half-thickness, representing a minimum of 4800 grains obtained for the coarsest microstructures. The average confidence index was greater than 0.73 for annealed samples and 0.41 for the as cold-
rolled sample, which showed a highly deformed structure.

TSM OIM Analysis™ v7.3.0 (EDAX) was used for data analysis. A grain dilation routine consisting in a single iteration was first performed on grains defined by a boundary misorientation superior to 5° and a minimum size of 5 pixels. A maximum of 1.9% of the pixels were replaced by this routine through all analyses, which allowed a proper measurement of grain boundary lengths. The boundary lengths were corrected by a factor of 0.77 to account for the hexagonal grid. Grain boundaries were then sorted into groups based on the phase whose grains they separate: ferritic (δ/δ) or austenitic (γ/γ) grain boundaries, or phase boundaries (δ/γ). Twin boundaries Σ3 (<111>/60°) present in austenite phase were considered apart from general austenitic grain boundaries as they were found to evolve differently with annealing. For each group, boundaries were additionally sorted based on their misorientation with a 2θ binning. The cumulative boundaries length was calculated for each interval and used to obtain various boundary densities. Ferritic and austenitic grain boundary densities were calculated as the ratio of the grain boundary length to the area of the respective phase, whereas the phase boundary density was defined as the ratio of the phase boundary length to the whole area of analysis comprising both phases. This approach is similar to that used by Haghdadi et al. [10] in a previous study.

### 2.2.4. Atom probe tomography

Atom probe specimens were extracted from cross-section samples by focused ion beam (FIB) lift out. The surface of the samples was prepared following the same procedure as for EBSD measurements.

Regions of interest 25 (RD) x 4 (ND) µm² in size were located in a Zeiss NVision 40 dual beam microscope. Those regions were first protected with a tungsten deposit before milling with Ga ion beam. Smaller sections were then attached to pre-sharpened silicon micro tip arrays. Annular sharpening of the specimens was carried out at 30 kV, 40 pA before a cleaning step at 10 kV, 80 pA to enhance the sharpness of the tip and minimize FIB damage. The specimens thus obtained were analyzed in a Cameca local electrode atom probe (LEAP) 4000 XHR at a base temperature of approximately 47.7 K in laser mode. The pulse energy of the UV laser was set between 35 and 45 pJ and its repetition rate at 200 kHz. The evaporation was conducted with a target detection rate of 0.01 ions/pulse (1%), controlled by steadily increasing the standing voltage from approximately 2 kV to 8 kV. The collected datasets were reconstructed and analyzed using CAMECA IVAS v 3.8.2 and established algorithms. Interfaces in the specimens were determined using iso-concentration surfaces (isosurfaces) of 6 at% Ni, as calculated using 1 nm cubic voxels and a delocalization of 3 x the voxel size in X and Y directions, and 1.5 x the voxel size in Z (the analysis direction). Concentration profiles across the interfaces were obtained by calculating proximity histograms (proxigrams) at set distance intervals based on those isosurfaces.

## 3. Results

### 3.1. Phase fractions, dislocations and interfaces

The evolution of ferrite fraction over time is shown in Fig. 1 for each heat-treatment temperature. Starting from 48 wt% in the as-cold-rolled sample, this fraction increases during the first ten seconds, regardless of temperature. The magnitude of this increase depends on temperature: the higher the temperature, the larger the increase. The ferrite fraction then keeps evolving until the 30 s mark and towards its thermodynamic equilibrium value calculated using the Thermo-Calc software package with the TCFE9 database (Thermo-Calc Software, Sweden).

The distribution of plastic deformation in the microstructures in the as-cold-rolled state and after a brief annealing treatment can be appraised using the inverse pole figures in Fig. 2a and Fig. 2b and the corresponding phase maps in Fig. 2c and in Fig. 2d. The cold-rolled state microstructure consists of deformed grains mostly oriented along the

<figure><img src="image_3.png" /><figcaption>Fig. 1. Ferrite fraction measured by magnetization at the cold-rolled state and after annealing at 1060 and 1180 °C. The evolution of fraction with tie is compared to the theoretical equilibrium value calculated with Thermo-Calc (TCFE9).</figcaption></figure>

rolling direction. Most of the grains are crossed by slip bands oriented at 45° from the rolling direction, a sign of severe plastic deformation. After a short annealing treatment at 1060 °C the deformation structure has entirely disappeared from ferrite grains and only traces of plastic deformation remain in a few austenitic grains. The grains constituting the bands are more equiaxed than in the as-rolled condition for both phases. For longer duration (10 s at 1060 °C) or higher temperature (0 s at 1180 °C) treatments, no signs of deformation can be observed in either phase.

Fig. 3 represents the absolute (a, c) and relative (b, d) evolution of the density of interfaces present in the microstructure with annealing at 1060 °C and 1180 °C. First, it can be noticed that a large fraction (about 65–70%) of the austenitic grain boundaries are twin boundaries at 0 s for both temperatures. At 1060 °C, austenitic/ferritic phase boundary density remains almost constant during the first ten seconds and decreases at longer annealing times. After a brief annealing, the grain boundary density of austenite is almost twice that of ferrite despite a quasi-equality in phase fractions because recrystallized austenite starts with a finer grain size than recovered ferrite. The densities of austenitic (excluding twin boundaries) and ferritic grain boundaries then decrease with annealing and converge towards the same value of about 0.10 µm⁻¹ after 300 s. This translates into a drop by about 20% for ferrite while it is over 60% for austenite. The density of twin boundaries in austenite remains more elevated and is about 3 times than that of non-special boundaries. At 1080 °C, the trend is generally the same although the values of densities drop faster to lower values. One difference can be found when considering the relative rates of decrease of non-special and twin boundary densities. At 1060 °C, the non-special boundary density shows faster relative decrease than twin boundaries, while at 1080 °C, they both decrease at the same relative rate.

### 3.2. Microstructure size and morphology

#### 3.2.1. Quantitative analysis of band thickness

The coarsening of the bands can be appreciated by tracking the mean band thickness of either or both phases. In Fig. 4, it can be seen that the average band thickness changes little during heating by comparing the mean thickness in the cold rolled state to that of the sample annealed for 0 s at 1060 °C. It can further be observed that the general evolution of the thickness is the same at both temperatures. However, the higher the temperature, the higher the rate of increase. The bands also get thicker at higher temperature. After 300 s, the mean band thickness has doubled at 1060 °C whereas it has quadrupled at 1180 °C. Fig. 5a and b reveal that the bands of both phases are responsible for this trend almost equally, with ferrite bands thickening slightly more than austenite bands. This minor difference could be owed to the ferrite fraction
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Inverse pole figures (a, b) and corresponding phase maps (c, d) indexed from EBSD analyses in the cold-rolled state (a, c) and after brief heating (0 s) at 1060 °C (b, d). Rolling direction (RD) is horizontal and normal direction (ND) is vertical.</figcaption></figure>

increase especially at higher temperature, where this increase is more pronounced.

### 3.2.2. Qualitative observation of the morphological evolution

The morphology of phases shows only a minor evolution during heating as seen in Fig. 2c and d. Major changes are only visible after longer annealing times. They will be summarized below by considering the microstructures after annealing for 0 s and 300 s at 1060 °C. The trend is the same at 1180 °C but changes happen faster and in a more pronounced fashion.

Fig. 6a and c show the two-phase microstructure of the alloy observed after 0 s at 1060 °C in the RD / ND plane and the TD / ND plane, respectively. The austenite bands are dispersed in the ferrite matrix. This is consistent with the ferritic solidification mode of the alloy and its subsequent hot and cold rolling [11]. However, it should be noted that given the phase fraction ratio close to unity, austenite can locally act as the matrix in which ferrite bands are embedded. Those local instances are highlighted by red circles in Fig. 6a and in c. These bands can be sorted into two populations. The thicker austenitic bands measure a few microns in thickness and can be considered as plates as they measure up to a hundred microns along RD and between 20 and 60 µm along TD. The thinner austenitic bands can be considered as ellipsoidal particles as they measure less than 30 µm along RD and less than 20 µm along TD. Some early stage of band breakup can also be seen but only in locations where they are the thinnest, as highlighted by the yellow dashed circles in Fig. 6a and c.

Fig. 6b and d show the two-phase microstructure of the alloy observed after 300 s at 1060 °C in the RD / ND plane and the TD / ND plane, respectively. After 300 s at 1060 °C, the morphology and size of austenite bands have evolved considerably. The interfaces between the bands show features where one phase forms small dents into the other. Examination of Fig. 7 reveals that those dents are located at the triple junction of two-phase interfaces and a grain boundary and are the

results of either phase wetting the grain boundaries of their complementary phase. All grain boundaries seem to be equally affected except for austenite Σ3 twin boundaries. In several locations, it can be observed that this wetting is about to lead to austenite band breakup as highlighted by the yellow dashed circles in Fig. 6b and d. Where the phenomenon is more advanced, aligned fragments of austenite bands remain after such breakup as pointed out by the yellow arrows. Some break up of ferrite bands can also be spotted as indicated by the red circles, although it is a lot less frequent. Additionally, it cannot be distinguished whether those sites correspond to austenite wetting ferrite grain boundaries parallel to ND or ferrite wetting austenite grain boundaries parallel to RD or TD.

### 3.2.3. Statistical analysis of the morphology

Fig. 8a and b show the evolution of the distributions of intercept lengths in austenite normalized by their average at 1060 °C and 1180 °C respectively. These plots permit to identify a morphological change in a banded microstructure independently of size changes as detailed in Appendix A in the Supplementary Material. The plots show that band breakup and globularization happen steadily along the 300 s treatment at 1060 °C and could possibly continue further. In contrast, most of the morphological evolution of the bands is complete after 10 s at 1180°C with little further evolution for the remaining 290 s.

## 3.3. Composition across the interface

Fig. 9 shows the tomograms reconstructed from atom probe data for specimens as rolled and annealed for 0 s and 300 s. Fig. 10 shows typical composition profiles found across austenite/ferrite interfaces at various stages during annealing at 1060 °C for all the alloying elements present in the model alloy. The interface can be located on these plots using the segregation peaks formed by N, Cr and Mo. The maxima of these peaks are used to define the location of the interface (i.e. x = 0). Clear
<figure><img src="image_5.png" /><figcaption>Fig. 3. Absolute (a, b) and relative (c, d) boundary densities with annealing time at 1060 °C (a, c) or 1180 °C (b, d).</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>Fig. 4. Average thickness of all bands (ferrite + austenite) as a function of annealing time at either 1060 °C or 1180 °C.</figcaption></figure>

partitioning of all elements can be seen on either side of the interface for all cases to concentrations close to the calculated equilibrium concentrations. Given the displacement of the interface observed through other techniques, these composition profiles suggest that conditions close to full equilibrium are maintained during the movement of the interface.

# 4. Discussion

## 4.1. Microstructural evolution and associated driving forces

The above results show that multiple changes happen during annealing of the as-cold-rolled alloy. The first change is the recrystallization of austenite and ferrite. These transformations happen almost entirely upon heating, with only a small fraction of austenite not yet recrystallized after 0 s at 1060 °C. This suggests that the elastic energy stored in the material in the form of dislocations during prior plastic deformation is mostly dissipated upon heating. The second is the change of the ferrite fraction towards its equilibrium value as a function of temperature. A large proportion of the fraction variation occurs upon heating and its value is almost stable after 30 s. This indicates that the chemical driving force available for phase transformation is largely dissipated.

The last change that is noticed is the size and morphology change of the microstructure. Contrary to the other two phenomena, this modification does not start during heating, as the initial band sizes of both phases are practically identical to the as-rolled condition for all temperatures. Once the treatment temperature is reached, thickening of the bands begins rapidly. While it slows down as the treatment progresses, it is still ongoing even for the longer time investigated here, 300 s, at all temperatures. This thickening of bands is logically accompanied by a reduction of the density of phase boundaries in the material, leading to a decrease of the amount of interfacial energy. This decrease is amplified by the decrease of grain boundary densities in both austenite and ferrite.
<figure><img src="image_7.png" /><figcaption>Fig. 5. Average thickness of (a) ferrite and (b) austenite bands as a function of annealing time at either 1060°C or 1180°C.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>Fig. 6. Optical micrographs after color etching of samples heated at 1060 °C for (a, c) 0 s and (b, d) 300 s. Horizontal corresponds to the rolling direction (RD) on micrographs (a, b) and to the transverse direction (TD) on micrographs (c, d). Vertical is the normal direction (ND).</figcaption></figure>

This shows that the size and morphological evolution of the micro-
structure can happen in the absence of chemical and elastic driving
forces and that it is primarily driven by a reduction of interfacial energy.

4.2. *Importance of the grain structure in morphological evolution and coarsening*

As Keichel et al. noticed [2], the breakup of the bands is primarily initiated at grain boundaries of one phase, where it is seen that the other phase wets those boundaries, a phenomenon termed grooving [12,13].
<figure><img src="image_9.png" /><figcaption>Fig. 7. Phase map after annealing at 1060 °C for 300 s. Grains boundaries, phase boundaries and twin boundaries are also represented. Rolling direction (RD) is horizontal and normal direction (ND) is vertical.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>Fig. 8. Distribution of normalized intercept lengths in austenite along the normal direction (ND) during annealing at a) 1060 °C and b) 1180 °C.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>Fig. 9. Atom probe tomograms of austenite/ferrite interfaces a) as cold-rolled and after annealing at 1060°C for b) 0 s and c) 300 s.</figcaption></figure>

In the present, it was confirmed that grooving occurs at grain boundaries of both ferrite and austenite regardless of their orientation with the sheet directions, with the exception of Σ3 twin boundaries in austenite. This is consistent with non-special grain boundaries in ferrite and austenite presenting an average energy of about 0.8 J.m⁻² [14], higher than the average energy of the austenite/ferrite interface of 0.56 J.m⁻² [14], while the energy of Σ3 twin boundaries is reported below 0.02 J.m⁻² in austenitic stainless steel [15]. The grooving process, driven by capillary forces, is a contributor to the decrease in the grain boundary density in both phases. It induces local curvature differences in the same band and between neighboring bands through the mechanism schematically outlined in Fig. 11. High curvature elevates the free energy of the phase,

leading to its local dissolution. Conversely, low curvature lowers its free energy, resulting in local growth. The bands that compose the micro-structure of super duplex stainless steel flat products comprise tens to hundreds of grains. Thus, the presence of grain boundaries in the bands increases the density of sites where capillary driven evolutions take place. Within the same band, this promotes band breakup and glob-ularization, whereas between two different bands/particles, it enhances coarsening. This reinforces the conclusion of Wessman et al. who noted that phase interface curvature only is insufficient to drive the observed coarsening [3]. This also implies that correct modeling of morphology and size changes necessitates accounting for all interfaces, including grain boundaries and phase interfaces, and their associated energies.
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** Composition profiles a,b,c) for chromium and nickel and d,e,f) for molybdenum and nitrogen measured by atom probe tomography, across the austenite/ferrite interfaces shown in Fig. 9, a,d) as cold-rolled, after annealing at 1060 °C b,e) for 0 s and c,f) for 300 s. The dashed lines indicate the equilibrium concentrations in either phase calculated using Thermo Calc with the TCFE9 database. Equilibrium concentrations for the as-rolled was calculated at 1100°C, which corresponds to the temperature of the solution annealing step that preceded cold rolling.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 11.** Schematic representation of the grain boundary grooving process. a) After cold rolling, grain boundaries intersect the interface, which is practically flat at the local scale. b) During the early stages of annealing, the angles formed by the surfaces at the triple junction will evolve towards equilibrium values, causing the complementary phase (δ here) to wet the grain boundary. c) The local curvature thus created causes differences in solubility that promote globularization.</figcaption></figure>

### 4.3. Process parameters and microstructural evolution trajectory

It was observed here that the annealing temperature exerts a strong influence on the competition between globularization and coarsening. At 1060 °C, the degree of band subdivision deduced from Fig. 8 observed after 300 s is similar to that observed at 1180 °C after the same duration, even though the rate at which it occurs is very different. However, band thickening is a lot more subdued at the lower temperature. The

composition measurements near the interface in Fig. 10 show that the interface moves with partitioning of all alloying elements. Thus, any interface displacement is associated with some mass transport, whose characteristic diffusion distance is affected by annealing temperature. Globularization and band breakup require this transport to happen over reduced distances compared to coarsening. In the former case, it can additionally happen along the phase interface where diffusion is enhanced. Therefore, globularization, resulting from band subdivision
and subsequent spheroidization of the resulting particles, can occur as readily as 1060 °C as it does at 1180 °C, albeit more slowly. Thus, the lower thermal energy available for diffusion activation limits the amount of thickening seen at 1060 °C compared to 1180 °C.

It was also demonstrated that grain boundaries play a prominent role in the capillarity-driven microstructural evolution. However, it is seen that Σ3 twin boundaries are an exception and contrary to general boundaries do not initiate grooving. It is known that the grain boundary distribution can be engineered in austenitic stainless steel by tuning the thermomechanical schedule [16–18]. Consequently, the results of this study suggest that tailoring the ratio of Σ3 twin boundaries over general grain boundaries, for instance by adjusting the rolling schedule, could afford a degree of control over the microstructural evolution during final annealing.

## 4.4. Segregation and interfacial mobility

Composition measurements across the interface between austenite and ferrite reveal segregation of the solute elements as shown in Fig. 10. The segregation behavior of each solute is consistent with the one found in the literature for austenite/ferrite interfaces: molybdenum is found to segregate strongly in all conditions reflecting the strong intrinsic attractive interaction it has with the interface [19–21]. Nickel, on the other hand, is found to segregate weakly only in the cold rolled sample and not at all in the other conditions, in line with the negligible interaction it is reported to have with the interface [19,20]. Chromium and nitrogen both tend to segregate to the interface but two distinct configurations can be seen: in the cold rolled specimen, they both show little to no segregation at the interface, while they both segregate strongly at the interface in both annealed conditions. This would be consistent with them only having a modest intrinsic attractive interaction with the interface [20,22,23] and a stronger attractive interaction with each other at the interface, as they do in both bulk phases [24,25].

The difference in segregation behavior between cold-rolled and annealed samples cannot be definitely explained given the information available. It could be related to the quench that the materials underwent last. Both annealed samples were quenched with precision at the end of the heat treatment as detailed in 2.1, while the cold rolled sample was quenched after solution annealing in a much less controlled manner. Consequently, it is possible that the interface started moving during this quench, potentially at high velocity. This hypothesis is consistent with the ferrite fraction found in the cold rolled sample, which is lower than expected at the solution annealing temperature. Alternatively, it could be related to the crystallographic nature of the interfaces, which was not checked before sampling the atom probe specimens in the present work. Haghdadi et al. have shown that rational interfaces (i.e. those separating phase units with specific orientation relationships) show little to no segregation as is the case for the cold-rolled specimen here, while irrational interfaces display a segregation behavior similar to that of the annealed samples [26].

The levels of segregation observed here reveal that alloying elements interact with, at least some of, the interfaces in the conditions of super duplex final annealing. This interaction can lead to an extra energy dissipation in the form of solute drag [27], which could slow down the movement of the interface. The amplitude of this dissipation has a complex dependence on the velocity of the interface and the transport kinetics within [28]. Further work is thus needed to evaluate whether this contribution needs to be accounted for to model accurately the present microstructural evolution. The observed segregation behaviors of the various alloying elements suggest that even though the majority of the data from literature is concerned with lower temperatures, it would likely constitute an adequate starting point.

## 5. Conclusion

The microstructure of super duplex stainless steel flat products

undergoes significant evolution during final annealing. The displacement of the interface between the two constitutive phases proved to occur with partitioning of all alloying elements. Alloying elements also segregate to the interface with a potential solute drag effect that would need to be considered in future modeling work. The observed microstructural changes can be sorted according to the nature of their driving force: chemical, elastic or capillary. The first induces an adjustment of the phase fractions towards their equilibrium values. The second leads to recrystallization of austenite and ferrite. The last is responsible for globularization and coarsening of the microstructure. While the first two start dissipating during heating and are mostly exhausted within the first seconds at the target temperature, the alterations caused by the third only begin once the treatment temperature is reached and still proceed after several minutes. In addition to the phase interface, grain boundaries in both phases play an important role in surface energy reduction through grooving, with the notable exception of Σ3 twin boundaries in austenite. Grooving induces local curvature differences that contribute to both band subdivision into globular particles and coarsening. The competition between the two phenomena is affected by annealing temperature due to their different dependence on atomic mobility. As austenite Σ3 twin boundaries were seen not to participate to capillarity-driven reactions, the microstructural evolution seen during final annealing can likely be tuned by grain boundary engineering.

## Declaration of Competing Interest

The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: Muriel Véron reports financial support and equipment, drugs, or supplies were provided by Aperam Stainless France. Florent Karjcarz and Jean Denis Mithieux report a relationship with Aperam Stainless France that includes: employment. If there are other authors, they declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Data availability

A Ph.D. thesis [1] covering more details on this study and complementary analyses is also available for reading at https://tel.archives-ouvertes.fr/tel-02510845. The raw data required to reproduce these findings are available to download from https://zenodo.org/records/10362513. The processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.

## Acknowledgments

This work was performed with the support of the Center of Excellence of Multifunctional Architecture Materials “CEMAM” No. ANR-10-LABX-44-01 funded by the “Investment for the Future” Program operated by the National Research Agency (ANR). APT work was carried out at the Canadian Centre for Electron Microscopy, a national facility supported by the Canada Foundation for Innovation under the Major Science Initiative program, NSERC and McMaster University.

## Appendix A. Supporting information

Supplementary data associated with this article can be found in the online version at doi:10.1016/j.mtcomm.2024.108899.

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