Near‐Atomic Scale Characterization of Mn‐Gradient in Reverted Austenite in Hot‐Rolled Medium‐Mn Steels

Herein, the Fe–5Mn–2Al–0.1C (wt%) and micro‐alloyed Fe–5Mn–2Al–0.1C–0.1Nb–0.2V (wt%) medium Mn steels are processed by a novel flash‐austenite‐reverse‐transformation (ART) annealing. The core–shell Mn‐gradient in austenite is observed and compositionally characterized at near‐atomic scale by using atom probe tomography. Compared with the conventional ART annealing, flash‐ART process doubled the austenite fraction in both materials, based on electron backscattering diffraction measurement. The rapid austenite reversion in flash‐ART annealed MMnS is attributed to the negligible partitioning local equilibrium controlled austenite kinetics at the flash temperature. The formation of core–shell Mn‐gradient is due to the different austenite growth kinetics and interface mobilities. Microalloying elements are found to form nano‐carbides, which leads to decelerating austenite reversion because of depleted carbon content in the initial martensitic microstructure, and retarding austenite growth according to pinning effect.


Introduction
The demands of weight savings, emission reduction, and passenger safety in the automotive industry have promoted the concept of medium-Mn steels (MMnS) design to achieve an extraordinary combination of high strength and superior ductility. [1,2][5][6][7] The principle of ART annealing, i.e., intercritical annealing (IA), is to revert parent martensite (α') or ferrite (α) to austenite phase (γ) in the temperature range between A c1 and A c3 .The austenite reversion process is accompanied by elemental partitioning, where Mn and C diffuse from α/α' into reverted austenite, enhancing the stability of austenite.
[7][8][9][10][11][12] Hu et al. [5] and Ma et al. [7] reported the role of elemental partitioning and segregation at transforming austenite/ferrite interface in controlling yielding behaviors.Han [8] systematically reviewed the mechanical properties and cracking behaviors governed by microstructural morphology in MMnS.Ding et al. [9] and Sun et al. [12] discussed the effect of pre-existed austenite phase on microstructural evolution and mechanical properties.The pre-existed austenite adjusts the initial Mn distribution among microstructure and provide tailored heterogeneous retained austenite in the as-annealed microstructure.To understand the austenite reversion mechanism is essential to provide practical guidance for adjusting austenite fraction, morphology, and distribution in MMnS.
Hot-rolling and cold-rolling processes are used for manufacturing of MMnS.The lamellar structure achieved in hot-rolled (HR) MMnS and globular structure obtained in cold-rolled (CR) MMnS exhibit distinct mechanical behaviors, e.g., tensile properties, toughness, fatigue property, and hydrogen embrittlement resistance.Uniform grain size distribution is obtained with lamellar or globular morphologies after IA.Unlike the CR MMnS with the large number density of defects, the austenite reversion proceeds in a sluggish way in HR MMnS, especially in relatively lean C and Mn condition.Increasing C and Mn contents will accelerate the austenite reversion according to thermodynamics, which in reverse reduces weldability and increases alloying costs.Wang et al. [13] investigated the austenite reversion kinetics in a Fe-9Mn-3Ni-1.4Al-0.01C(wt%) MMnS at CR and HR states.The austenite fraction reached %0.4 after 1 h annealing at 600 °C in CR sample, while the austenite fraction only approached %0.35 after 8 h annealing in HR sample.The sluggish austenite reversion in HR MMnS was attributed to the reduced density of deformation-induced defects, e.g., martensite grain boundary, stored dislocation.Most recently, new alloying design strategy and processing route, e.g., preservation-deformation-annealing, [12] cyclic and flash annealing, [14][15][16] have been proposed to design chemical heterogeneity in austenite to enhance mechanical properties [17] and hydrogen embrittlement resistance, [18,19] which brings new opportunity to control austenite reversion for HR MMnS with balanced properties.Wan et al. [15] first proposed a flash-austenite reversed transformation (flash-ART) annealing to produce a compositional core-shell austenite phase in CR Fe-0.2C-7.8Mn-2Al(wt%) MMnS, which succeeded in increasing austenite fraction.The chemical heterogeneity was characterized using nano-auger electron spectroscopy equipped with electron backscatter diffraction (AES-EBSD). [15]However, the formation mechanism of the chemical heterogeneity remains unclear, due to the limited spatial resolution and high surface sensitivity of characterization techniques.In order to further tailor the heterogeneous chemical gradients for enhancing materials properties, it is essential to characterize the heterogeneous chemical distribution behaviors with higher spatial resolution and understand its formation mechanisms.
In the present work, the chemical heterogeneity produced via flash-ART in HR Fe-5Mn-2Al-0.1C(wt%) and HR microalloyed Fe-5Mn-2Al-0.1C-0.1Nb-0.2V(wt%) MMnS are investigated at near-atomic scale using the atom probe tomography (APT).The comparison of microstructural evolution in flash-ART processed and conventional ART processed samples is reported based on electron backscattering diffraction (EBSD) measurements.The effect of microalloying elements (Nb and V) on microstructural evolution and chemical heterogeneity is discussed, and the formation mechanisms of Mn gradients are further explained based on thermodynamics calculations and experimental observations.

Microstructure Characterization
The EBSD measurements of Fe5Mn and Fe5Mn-NbV steels processed by ART and flash-ART routes provide the quantitative information of microstructure in terms of austenite fraction, grain size, as displayed in Figure 1.The materials and experimental details are provided in Section 4. The direct influence of flash-ART annealing on austenite reversion is revealed.According to the calculation based thermal history, flash-ART annealed samples were heated from 680 to 800 °C within 1.2 s and in turn cooled down within 2.4 s, which sums as 3.6 s.Compared with the conventional ART process, flash-ART route only differs in a 3.6-second process before isothermal holding.The flash-ART process significantly increases the austenite fraction.
In Fe5Mn steel, the austenite percent increased from 5.9 vol% in ART annealed condition (Figure 1a) to 12.8 vol% in flash-ART annealed condition (Figure 1b).For Fe5Mn-NbV steel, the doubled austenite percent of 3.7 vol% was observed in flash-ART annealed sample, compared with ART annealed sample with 1.6 vol% austenite.Unlike the CR MMnS with large density of deformation-induced defects as fast diffusion path, [15] the flash-ART annealing route is proven to be identically effective in accelerating austenite reversion kinetics in HR MMnS with less deformation-stored energy.
The additions of microalloying elements reduced austenite fraction and average grain size, both in ART and flash-ART annealed samples.Although the flash-ART process doubled the austenite fraction in Fe5Mn-NbV sample, the resultant austenite fraction was merely around one-fourth of that in flash-ART annealed Fe5Mn sample.Average austenite grain size decreased with the additions of microalloying elements, where it drops from 0.35 AE 0.15 to 0.24 AE 0.08 μm in ART annealed samples and similarly from 0.34 AE 0.12 to 0.27 AE 0.10 μm in flash-ART annealed samples.The mechanisms of flash-ART process and effect of microalloying addition on austenite transformation will be investigated using near-atomic scale characterization technique with high chemical, spatial resolutions, i.e., atom probe.

Near-Atomic Scale Characterization of Mn Gradients
Figure 2 shows the nano-sized austenite (γ)-martensite (α') duplex microstructure of flash-ART annealed Fe5Mn and Fe5Mn-NbV samples by three-dimensional (3D) APT technique.In flash-ART annealed Fe5Mn sample (Figure 2a), austenite and martensite with distinct Mn contents are separated by Mn iso-concentration surfaces, i.e., 3.5 at% iso-surface in yellow and 8.5 at% iso-surface in purple.A cylindrical region of interest   (ROI) was placed across the whole γ phase for compositional examination.The two-dimensional (2D) Mn density contour map and concentration profile along the z-axis of ROI are shown in Figure 2b,c respectively.The clear Mn heterogeneity can be observed, forming a core-shell structure.The shell regions with the thickness of %30 nm refer to the areas with over 10 at% Mn content, while the inner core region only revealed a low Mn concentration of %7 at%.This compositional core-shell structure has been reported to stabilize the reverted γ phase, where the martensitic transformation of core austenite will be suppressed according to the "shielding effect". [15,17]It is interesting to observe the asymmetric Mn concentration profiles from the core region to the opposite parallel α'/γ phase boundary, which is attributed to the nucleation of austenite at the martensite lath boundary and altered interface mobility during austenite growth. [20]In HR MMnS, austenite nuclei and cementite initially form at martensite lath/packet/block boundaries during ART annealing. [8]The existing austenite nuclei grow, and cementite may act as nucleation sites for the new austenite phase, [21,22] exhibiting lamellar α'-γ matrix after conventional annealing.In the proposed flash-ART process, the cementite precipitation window is rapidly passed by due to a high heating rate until 800 °C.The cementite precipitation during flash heating is thus ignored, and the austenite is mainly obtained by martensite-to-austenite transformation.The asymmetric Mn concentration profiles in Figure 2c were produced when austenite grew from boundary into neighboring martensite grain and encountered distinct crystalline structures, e.g., nano carbides, substructure.The α'/γ interface with a wider heterogeneous Mn gradient consequently showed a high migration rate during austenite reversion. [23]In the microstructure of MMnS, there existed a large amount of elemental inhomogeneity.Although no cementite is expected to form during the rapid heating process, cementite can be formed during intermediate cooling or isothermal holding in Fe5Mn steel.Moreover, C/Mn enrichments at various nanostructures (substructure, dislocation etc.) may also exhibit similar effects as carbides. [23]ifferent from the core-shell structure characterized by AES-EBSD and energy-dispersive X-ray analysis (EDX)transmission electron microscopy (TEM) in flash-ART annealed CR Fe-0.2C-7.8Mn-2Al(wt%) MMnS, [15] a continuous decreasing Mn concentration can be clearly observed in Figure 2c.The sharp concentration drops in ref. [15] may be attributed to 1) the surface sensitivity of the characterization technique due to the small interaction volume of auger electrons and 2) possible subgrained austenite with high Mn content and inclined EDX measurement line in respect to α'/γ interface.Mn gradient can be explicitly determined by APT measurement in the present work, providing hint to elucidate the austenite reversion mechanism in terms of Mn diffusion and interface migration, which will be followed up in discussion.
Figure 2d shows the reconstructed APT tip of flash-ART annealed Fe5Mn-NbV MMnS.The green 7 at% Mn iso-concentration surface separates the austenite (upper part) and martensite (bottom part) phases, and the red 4 at% C iso-concentration surfaces illustrate the oval-shaped carbides.A cylindrical ROI is placed perpendicular to the α'/γ interface and corresponding Mn density mapping, and atomic concentration profiles are shown in Figure 2e,f.A similar core-shellstructured austenite phase can be observed.Compared to the Fe5Mn alloy, the width of high-Mn shell of Fe5Mn-NbV was restricted to only %10 nm, and the peak Mn concentration (%11.5 at%) was slightly higher than that in Fe5Mn alloy (%10 at%).The core region is Fe5Mn-NbV alloy is not completely captured, but showing a relatively low Mn content of %7 at%.Nano-sized carbides appeared at α'/γ interface and within the austenite phase.These carbides are identified as (V,Nb)C carbides with an approximate stoichiometry of V 0.9 Nb 0.1 C, which is determined by proxigram method.The rapid heating rate leads to negligible diffusion of substitutional alloying elements.The (V,Nb)C carbides were thus mainly formed during hot rolling process and during isothermal holding.The (V,Nb)C carbide at α'/γ interface reveals a radius size of 4 nm, while those carbides inside the austenite grain are much smaller (1-2 nm).The growth of austenite phase is accompanied by the dissolution of metastable carbides [24,25] and precipitation of face-centered cubic structured (V,Nb)C carbides.In the investigated microalloyed MMnS, the austenite grew and the α'/γ interface continuously encountered (V,Nb)C carbides.When the α'/γ interface migrated forward, the (V,Nb)C carbides will be dissolved, which corresponds to the small carbides inside the austenite grain in Figure 2d.The (V,Nb)C carbide at α'/γ interface (blue box in Figure 2d) maintained a large size and can exhibit drag effect on α'/γ interface, [26] continuously retarding austenite reversion kinetics.Besides the pinning effect of interface-localized carbides, the local C concentrations identically influence the austenite reversion kinetics.The initial stage of austenite growth is reported to be controlled under negligible partitioning local equilibrium (NPLE) mode. [15,17,23]Due to the presence of carbides that locally with high C contents, the initial martensitic microstructure of Fe5Mn-NbV steel contained less C content compared with Fe5Mn steel, resulting in retarded austenite growth kinetics and small amount of austenite in annealed Fe5Mn-NbV samples.The significantly reduced width of high-Mn shell was also resulted from the slow mobility of α'/γ interface.

Thermodynamics Calculations
Thermodynamic calculation is complemented to experimental observation, in order to deliver the deep understanding of austenite reversion mechanism during flash-ART annealing.In the previous discussion and literature, [15,17,23] the fast austenite formation is attributed to the NPLE-controlled growth and is influenced by C content in the parent martensitic microstructure.The thermodynamic calculations provided equilibrium Mn concentrations in austenite phase at 680 and 800 °C in two investigated materials, which is shown in Figure 3a.The Mn concentration in the core region of austenite is close to the equilibrium Mn concentration in austenite at 800 °C, and the Mn concentration in the shell region is close to the equilibrium Mn concentration at 680 °C (Figure 3b,c).This phenomenon is significant in understanding austenite growth kinetics, and the Mn heterogeneity can thus be tailored by adjusting flash-ART annealing parameters.The flash heating to the elevated ART temperature led to initial austenite growth under NPLE mode and already formed the large amount of the reverted austenite phase in Fe5Mn steel when reaching 800 °C.However, the martensitic microstructure with less C content due to carbide existence showed much less austenite fraction in annealed Fe5Mn-NbV steel.Following austenite growth kinetics is dominated by partitioning local equilibrium (PLE), where Mn diffusion and α'/γ interface mobility are sluggish.The decreasing Mn gradient from shell to core is expected to form when sample cooled from 800 to 680 °C and subsequent isothermal holding.Meanwhile, the width of high-Mn shell is determined by α'/γ interface mobility.The austenite/ferrite interface reveals varied interface mobility during the transformation. [23]When the migrating interface encounters regions of crystal containing some precipitates or microstructural defects, the mobility changes due to the altered dissipation energy. [23]The increased mobility thus leads to fast austenite growth.The isothermal holding at 680 °C contributed further to the austenite reversion, maintaining the equilibrium Mn concentration at α'/γ interface as well as triggering internal shell-to-core Mn diffusion.The reasons for the near-equilibrium Mn content (at 800 °C) in the core parts are 1) preexisting high-Mn region assisting early austenite nucleation and locally provide Mn solutes and 2) thermodynamics and internal Mn diffusion in austenite from shell region to core region.The local equilibrium state can be achieved in most austenite formed during flash annealing.As a result, the large amount of stabilized austenite phase can be obtained in HR MMnS without extending processing duration.The formation mechanism of heterogeneous austenite in MMnS will contribute to the MMnS deign with balanced mechanical properties [12,16,17] and enhanced hydrogen resistance. [18,19]

Conclusions
In summary, the flash-ART annealing benefits the austenite reversion by forming a core-shell Mn gradient in the austenite phase, lowering the overall Mn content in austenite.The core-shell Mn gradient in the investigated MMnS is formed owing to different austenite growth modes and changed α'/γ interface mobility.The initial high austenite growth kinetics is due to NPLE mode and followed by sluggish kinetics under PLE modes.Microalloying carbide on the one hand decreases the C content in parent martensite microstructure, resulting in a significant reduction in austenite fraction.On the other hand, austenite/martensite interface may continuously encounter microalloying carbides and exhibit a retarded mobility.The understanding of these Mn distribution behaviors is beneficial for tailoring the chemistry heterogeneity, enabling further novel design of metastable nanostructures in advanced high-strength steels.

Experimental Section
The chemical compositions of the investigated MMnS are Fe-5Mn-2Al-0.1C(wt%), labeled as Fe5Mn, and Fe-5Mn-2Al-0.1C-0.1Nb-0.2V(wt%), labeled as Fe5Mn-NbV.The alloy was melted in a 50 kg vacuum induction furnace and then cast into ingots, which were then hot forged to 60 mm-thick plates.The forged steel plates were homogenized at 1200 °C for 2 h and then hot rolled to 4.5 mm-thick strips with the start and finish rolling temperatures of 1100 and 850 °C, respectively, followed by the oil quenching to room temperature.Fully martensitic microstructure (>99%) was obtained at HR state according to high-energy synchrotron X-ray diffraction measurements.No obvious Mn segregation band at micron scale was detected by EDX analysis mapping in a 50 μm Â 50 μm region.According to the Thermo-Calc calculations, the Ac 1 and Ac 3 of Fe5Mn steel are 476 and 953 and 500 and 970 °C for Fe5Mn-NbV steel, respectively.
Two heat treatment routes were performed on Fe5Mn, Fe5Mn-NbV samples, respectively, using Bähr 805A/D dilatometer: 1) for flash-ART process, the samples were heated to 800 °C with a heating rate of 100 °C s À1 and cooled immediately to 680 °C with a cooling rate of 50 °C s À1 and isothermally annealed at 680 °C for 1 h.The samples were finally quenched to room temperature at a cooling rate of 200 °C s À1 and 2) for the reference ART process, samples were heated to 680 °C at a heating rate of 100 °C s À1 isothermally annealed for 1 h.The samples were finally quenched to room temperature at a cooling rate of 200 °C s À1 .
The microstructures were characterized by EBSD using Zeiss SIGMA field emission scanning electron microscope (SEM) equipped with the EBSD detector (Oxford Instrument).The EBSD samples were ground with SiC abrasive paper with grit sizes of P320, P500, P800, P1200, P2400, and P4000 and then mechanically polished using 3 and 1 μm diamond paste.The final polishing was performed using oxide polishing suspension.The EBSD measurements were operated at an accelerating voltage of 20 kV with the step size of 50 nm.The acquired EBSD data were analyzed using the software Aztec Crystal (Oxford Instrument).The needle-shaped APT samples were fabricated by the site-specific lift-out and annular milling method using FEI Helios Nanolab 660 dual beam SEM-focused ion beam system.APT measurements were ran using a Local Electron Atom Probe 4000X HR system (CAMECA Instrument Inc.).Laser-pulse mode (wavelength of 355 nm, frequency of 200 kHz, and laser energy of 30 pJ) was selected for data collection.The base temperature in the analysis chamber was kept at 60 K during the measurement.The collected data were reconstructed and analyzed using the IVAS module in AP suite software 6.1 (CAMECA Instruments Inc.).The thermodynamic calculations were performed using Thermo-Calc software (2020a) combined with TCFE7 databases.

Figure 1 .
Figure 1.EBSD measurement of ART, flash-ART annealed Fe5Mn and Fe5Mn-NbV steels.a) Phase map of Fe5Mn ART sample; b) phase map of Fe5Mn flash-ART sample; c) phase map of Fe5Mn-NbV ART sample; d) phase map of Fe5Mn-NbV flash-ART sample; e) quantitative microstructural information in terms of austenite fraction and austenite grain size in 4 above samples.

Figure 2 .
Figure 2. APT analysis of flash-ART annealed samples.a) Fe5Mn sample rendered by 3.5 at% Mn iso-concentration surface (yellow) and 8.5 at% Mn iso-concentration surface (purple) showing duplex nanostructure; b) atom map and corresponding 2D Mn density contour map of the selected ROI (green box in (a)); c) corresponding Mn and C concentration profiles along long axis of the ROI in (b); d) Fe5Mn-NbV sample rendered by 7.0 at% Mn iso-concentration surface (green) and 4.0 at% C iso-concentration surface (red) showing duplex nanostructure; e) atom map and corresponding 2D Mn density contour map of the selected ROI (orange box in (d)); f ) corresponding Mn and C concentration profiles along long axis of the ROI in (e).

Figure 3 .
Figure 3.Comparison of Mn concentration obtained from APT analysis and equilibrium concentration derived from thermodynamics calculations: a) Mn equilibrium concentration calculated by ThermoCalc of Fe5Mn and Fe5Mn-NbV steels; b) APT determined Mn gradient of Fe5Mn steel; c) APT determined Mn gradient of Fe5Mn-NbV steel.