# High temperature EBSD experiment versus crystallographic reconstruction to study phase transformation induced microstructures - Application examples for steels

Maissa Fekih, T. Martinez Ostormujof, L Germain, M. Piette, D. Sornin,
Nathalie Gey

► To cite this version:

Maissa Fekih, T. Martinez Ostormujof, L Germain, M. Piette, D. Sornin, et al. High temperature EBSD experiment versus crystallographic reconstruction to study phase transformation induced mi- crostructures - Application examples for steels. Materials Characterization, 2025, 228, pp.115427. 10.1016/j.matchar.2025.115427. hal-05221037

HAL Id: hal-05221037

https://hal.univ-lorraine.fr/hal-05221037v1

Submitted on 25 Aug 2025

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# High temperature EBSD experiment versus crystallographic reconstruction to study phase transformation induced microstructures - Application examples for steels

M. Fekih a,*, T. Martinez Ostormujof a, L. Germain a, M. Piette c, D. Sornin b, N. Gey a,*

## ARTICLE INFO

### Keywords:

- In-situ high temperature EBSD
- Phase transformation
- Crystallographic reconstruction
- Steels
- Microstructure
- In-situ observation

## ABSTRACT

This study explores the capabilities and limitations of high-temperature electron backscatter diffraction (HT-EBSD) and crystallographic reconstruction methods for analyzing microstructure evolution induced by phase transformations in steels. Using a state-of-the-art CMOS detector with infrared filtering enabled EBSD map acquisition of austenite in real time at temperatures up to 1000 °C with acquisition speed over 1000 points per second (pps) and high indexing rates. These direct observations revealed critical features of the austenite microstructure, including annealing twins and grain size heterogeneities, that reconstruction methods typically overlook. Moreover, it was possible to confirm austenite grain refinement through austenitization cycles in a medium carbon steel and to reveal for the first-time ultrafine parent austenite grains in 10 % chromium oxide dispersion strengthened (ODS) steel, a case where reconstruction methods failed. Furthermore, in-situ HT-EBSD provides ground truth validation for reconstruction methods. By monitoring in-situ the martensitic transformation on 10 %Cr steel, we confirmed both the high accuracy of reconstructed parent austenite orientations and validated the average orientation relationship (OR) established through reconstruction methods. Finally, in-situ HT-EBSD could also provide valuable data for refining reconstruction algorithms to overcome their limitations in capturing precise grain morphology and annealing twins. Despite its advantages, HT-EBSD faces challenges including temperature measurement uncertainties and potential surface artifacts. This study ultimately advocates for judicious combination of both methodologies to comprehensively understand microstructure evolution during steel phase transformations, including the “memory effect” during re-austenitization and identification of austenite nucleation sites.

## 1. Introduction

The mechanical properties of steels depend fundamentally on their microstructures induced during processing. The high-temperature (HT) austenite phase and its transformation during cooling play key roles in determining the final microstructures and properties. Complex multi-constituent microstructures including ferrite, bainite, and martensite can form as a result of varied thermal treatments and cooling rates, ultimately determining critical material characteristics such as strength, ductility, hardness, and toughness [1–4]. Understanding and controlling the phase transformation mechanisms, both during heating and cooling, is crucial for controlling and optimizing the final microstructure and properties.

Two complementary approaches have emerged to study the high temperature austenite phase and subsequent phase transformations. The first, indirect parent-phase reconstruction, calculates austenite grain orientations from the orientation data of the transformed phase, based on established orientation relationships (ORs). Specifically, on the orientation map of the transformed phase, one considers neighboring crystallographic domains, corresponding to different variants and checks if they unambiguously define a unique parent orientation in OR with all variants. Since the early seminal work [5–7], there has been an impressive development of complementary strategies for reconstructing crystallographic orientation maps of the parent phase [8–17]. The first reconstructions have been successfully applied on Titanium and Zirconium alloys for microstructures inherited from the $β → α$ phase

<footnote>a Université de Lorraine, CNRS, Arts et Métiers Institute of Technology, LEM3, F-57000 Metz, France</footnote>

<footnote>b Université Paris-Saclay, CEA, Service de Recherches en Matériaux et procédés Avancés, 91191 Gif-sur-Yvette, France</footnote>

<footnote>c Centre de Recherche Vallourec, 59620 Aulnoye Aymeries, Hauts-de-France, France</footnote>

<footnote>* Corresponding authors at: Université de Lorraine, CNRS, Arts et Métiers Institute of Technology, LEM3, F-57000 Metz, France.
E-mail addresses: maissa.fekih@univ-lorraine.fr (M. Fekih), nathalie.gey@univ-lorraine.fr (N. Gey).</footnote>
transformation with respect to the Burgers OR [18]. Cayron et al. pioneered the application to prior austenite grains reconstructions in steels [11]. Today, reconstruction algorithms have gain in maturity and have been integrated into mainstream electron backscatter diffraction (EBSD) data post-processing software [16,19,20]. They are widely used to understand microstructure inheritance induced by forging, additive manufacturing, welding [10,21,22]. While powerful, these reconstruction methods have some limitations - they are primarily only applicable to displacive transformations, and the accuracy of some reconstructed grain boundaries remains uncertain. This is particularly the case for neighboring parent grains/twins sharing common variants [11,12]. Although these methods can also refine the average OR between parent and daughter phases [23,24], direct experimental validation remains essential for a deeper understanding of microstructure formation mechanisms.

The second approach leverages recent technological advances in compact hot stages and high-speed EBSD cameras, enabling direct in-situ observations of microstructural evolution at elevated temperatures. Early attempts at HT EBSD were hampered by limited camera sensitivity and pattern acquisition speeds [25–35]. For instance, Seward et al. [26] pioneered in-situ $\alpha \rightarrow \beta$ transformation studies in titanium using EBSD but were limited by the acquisition speed of the first generation of CCD cameras reaching only 10 points per second (pps). While continuous progress in EBSD detectors over the years has gradually accelerated data acquisition [30,36], most studies at HT still operated at lower acquisition rates to ensure a good signal quality [32]. These limitations often compromised real-time tracking of phase transformations as well as spatial resolution of the EBSD map, preventing the detection of fine microstructural features. However, modern high-speed EBSD systems have overcome these limitations. Notable achievements include Taylor et al.’s development of a fully automated system capable of measurements up to 1000 °C [36,37], and Taniyama et al.’s first observation of the austenite-to-martensite transformation in low carbon steel using a specially designed heating stage [38].

Despite these advances, a critical evaluation of in-situ HT EBSD for studying phase transformations, including both its capabilities and limitations, remains lacking. Furthermore, few studies have explored the synergistic potential of combining direct measurements with indirect reconstruction methods [36,37].

This paper addresses these gaps by providing a comprehensive state-of-the-art analysis of the latest generation HT EBSD experiments in comparison with crystallographic reconstruction methods for studying phase transformation induced microstructures in steels. The analysis was conducted through two representative case studies: The first examines austenite grain size evolution and its potential refinement through thermal cycling in the austenitic range. The second investigates a nuclear steel grade, with and without oxide dispersion strengthened (ODS), to demonstrate the acquisition of HT data that was previously inaccessible through indirect methods. Our aim is to highlight the strengths of HT EBSD measurements while acknowledging its limitations, demonstrating how these measurements can be effectively combined with reconstruction approaches to provide a more complete understanding of phase transformations in steels. This analysis serves as a valuable reference for materials scientists selecting appropriate methodologies for their specific research needs in phase transformation studies across various material systems.

# 2. Materials and experimental procedure

## 2.1. Materials

In-situ HT EBSD experiments were conducted on different steels. The first case concerns a medium-carbon steel used for Oil Country Tubular Goods applications (referenced here as OT-steel). Its composition and transformation temperatures are detailed in Table 1. Our purpose was to characterize the HT austenite grain size distribution by direct in-situ HT

<caption>**Table 1**</caption>

Nominal composition (weight %) and transformation temperatures of the OT-Steel, (Ac1, Ac3) at heating with a rate of 0.5 °C/s and (Ms, Mf) at quenching.

<table><thead><tr><th colspan="4">Nominal composition</th><th colspan="4">Transformation temperatures (°C)</th></tr><tr><th>C</th><th>Cr</th><th>Mo</th><th>V</th><th>Ac1</th><th>Ac3</th><th>Ms</th><th>Mf</th></tr></thead><tbody><tr><td>0.4</td><td>1</td><td>1.3</td><td>0.2</td><td>775</td><td>850</td><td>310</td><td>120</td></tr></tbody></table>

EBSD analysis. Further, our point was to investigate the effect of thermal cycling on austenite grain size refinement. The general framework aims to improve the sulfide stress cracking (SSC) performance with microstructure refinement [39,40]. Thus, two metallurgical states of OT-steel were considered: state A results from a single austenization treatment (called Q/T) and state B from two successive Q/Ts (each Q/T treatment was carried out at a temperature of 880 °C for 10 min followed by quenching to obtain martensite, this temperature corresponds to Ac3 + 30 °C – see table 1).

The second study investigates a 10 % wt chromium oxide dispersion-strengthened (ODS) steel, referenced as 10Cr-ODS, fabricated via powder metallurgy and hot extrusion at 1050 °C. This material represents a promising candidate for next-generation nuclear reactors, demonstrating enhanced creep resistance and neutron irradiation resilience. For comparative analysis, a non-ODS variant of the same steel, referenced as 10Cr, was also examined. The chemical composition and transformation temperatures are detailed in Table 2.

Our purpose was again to characterize the HT austenite grains of 10Cr-ODS steel in a challenging configuration where crystallographic reconstruction fails, as observed previously by Durand et al. [41,42]. In addition, we tackle in-situ analysis of martensite transformation on cooling in the reference grade 10Cr and evaluate the austenite/martensite OR by direct measurement in comparison to that deduced by indirect reconstruction.

## 2.2. Heating set-up for in-situ EBSD experiment

HT EBSD experiments were conducted on a ZEISS Auriga 40 FIB-SEM microscope equipped with the latest generation of Oxford Instruments CMOS Symmetry camera and a heating stage developed by Kammrath & Weiss. The EBSD detector allows for high-resolution data acquisition at elevated temperatures while maintaining a favorable signal-to-noise ratio and minimal distortion [35]. An infrared (IR) filter was installed to protect the camera from IR radiation emitted during heating cycles. The CMOS Symmetry camera can easily operate at sample temperatures up to 1000 °C with indexing speeds of 3000 pps for a 156 × 128 pixel resolution [43]. Under these conditions, an angular resolution of 0.5° can still be achieved, as evaluated from the Kernel Average Misorientation distribution obtained on fully recrystallized data set, considering the threshold angle for 95 % of the data. Furthermore, it is possible to acquire secondary electron (SE) images during the experiments, while the forescattered electron detector (FSD) becomes non-functional at elevated temperatures due to thermal signal interference.

The heating module (Fig. 1) features a compact design with optimized geometry for full integration in the SEM chamber at EBSD 70° tilt condition and a working distance of 15 mm. It is capable of reaching a maximum specimen temperature of 1050 °C. The heating system operates by applying a current to a ceramic heating element, indicated by number 2 in Fig. 1 (a), which is covered by a molybdenum (Mo) cap where the sample is fixed (number 3 in Fig. 1 (b)). To enhance thermal conductivity, a thin layer of graphite is positioned between the heating element and the Mo cap. A drilled hole is located in the bottom right of this Mo cap to allow for the insertion of a thermocouple (indicated by number 1 in Fig. 1 (a)). Two protective shields are mounted over the device to minimize radiation and prevent interference with the diffraction signal (indicated by number 4 and 5 in Fig. 1 (b)). The
<caption>**Table 2**</caption>

a. Chemical composition (in w%), b. Phase transformation temperatures (°C) measured by dilatometry (heating at 1 °C/s and cooling at 10 °C/s) for 10Cr and 10Cr-ODS steels [42].

<table><thead><tr><th>(a)</th><th>C</th><th>Cr</th><th>W</th><th>Mn</th><th>Mo</th><th>Ni</th><th>Ti</th><th>O</th><th>Al</th><th>Si</th><th>N</th><th>Y</th></tr></thead><tbody><tr><td>10Cr</td><td>0.104</td><td>9.97</td><td>0.49</td><td>0.26</td><td>0.3</td><td>0.3</td><td>0.1</td><td>0.03</td><td>0.005</td><td>0.019</td><td>0.006</td><td>-</td></tr><tr><td>10Cr ODS</td><td>0.11</td><td>9.7</td><td>0.88</td><td>0.26</td><td>0.45</td><td>0.31</td><td>0.22</td><td>0.12</td><td>0.01</td><td>0.06</td><td>0.004</td><td>0.21</td></tr><tr><td>(b)</td><td></td><td></td><td>Ac1</td><td></td><td></td><td>Ac3</td><td></td><td></td><td>Ms</td><td></td><td></td><td>Mf</td></tr><tr><td>10Cr</td><td></td><td></td><td>852</td><td></td><td></td><td>883</td><td></td><td></td><td>385</td><td></td><td></td><td>260</td></tr><tr><td>10Cr ODS</td><td></td><td></td><td>946.8</td><td></td><td></td><td>993.8</td><td></td><td></td><td>328.8</td><td></td><td></td><td>219.0</td></tr></tbody></table>

<figure><img src="image_2.png" /><figcaption>**Fig. 1.** Kammrath & Weiss heating module: (a) shows the thermocouple (1) and the ceramic heating element (2), (b) shows the sample (3), the radiation shields (4 and 5) and the 70 °C pre-tilted stage.</figcaption></figure>

module is positioned on a pre-tilted support at 70°.

Temperature is regulated with a PID controller offering two different modes and the option to program heating cycles. The first mode, called “PID control” allows a maximum heating rate of 3 °C/s, from RT up to 900 °C. The second mode, “Rapid heating” operates at maximum power with a maximum heating rate of 50 °C/s, though this can result in an overshoot of 20 °C and temperature oscillations. The fastest way to cool the specimen is to cut the power, enabling a cooling speed of 7 °C/s in both modes. The system includes active water cooling of the module body to prevent thermal transfer to the microscope components. The allowed sample size is 10 mm × 10 mm and 1 mm thick. To minimize the temperature gradient between the heated surface and the analyzed top face, it is recommended to limit the sample size and thickness under 1 mm.

### 2.3. Temperature measurement within the SEM chamber

To ensure safe conditions of EBSD analysis, preliminary heating experiments to 1050 °C were conducted with the hot stage to measure the temperatures of the components of the SEM chamber and the EBSD camera. Thermocouples have been attached at different locations: (1) on the SEM’s pole piece, (2) at the top of EBSD screen where the infrared radiation is maximal, (3) on the right side of the EBSD screen and finally (4) on the top of the heating module. For these experiments, the EBSD camera was inserted at 195 mm corresponding to a detector distance of 0.63 mm and an elevation of -4.7 mm.

None of the four thermocouples recorded a temperature above 78 °C (Fig. 2 (b)), confirming that the operating conditions inside the SEM are completely safe and secure.

<figure><img src="image_3.png" /><figcaption>**Fig. 2.** (a) Thermocouples positioned on various components within the SEM chamber prior to the installation of the heating module, (b) temperature curves recorded for the components relative to the heating temperature with the heating module in position (Detector Distance = 0.65 mm).</figcaption></figure>
# 3. Results and discussion

## 3.1. High temperature parent phase analysis: Potential and limitation

Fig. 3 shows the result of a large field of view EBSD acquisition at 840 °C on sample OT-steel-B in austenite phase (see section 2.1). The map was acquired with 0.5 µm step size over 290 × 210 µm area and acquisition speed of 1000 pps, completing the entire acquisition in less than 4 min. The fraction of reliable indexed points of the map exceeded 88 %.

This result demonstrates the capabilities of the latest EBSD technology in characterizing parent phase microstructures at HT. High-speed acquisition with reliable indexing is clearly evidenced and opens up the possibility of capturing dynamic changes in the microstructure. In this case, we observed limited grain growth at the plateau temperature.

Interestingly, this large EBSD acquisition revealed austenite grain size heterogeneities organized in vertical bands approximately 10–20 µm thick, spaced about 100 µm apart. These bands contain very small austenitic grains, likely induced by localized variations in chemical composition.

One drawback of such HT acquisition is related to sample temperature control. On our hot stage, the temperature is controlled with a thermocouple inserted in a hole drilled in the bottom right of the Mo cap. The temperature difference between the thermocouple and sample surface is suspected to reach several tens of degrees. At least, we often observed a difference between the temperatures start/finish of the transformation as determined by HT EBSD and by dilatometry for close heating rates. For samples of similar metallurgical state heated identically, the temperature for fully austenitic transformation varied from one experiment to the other. This temperature variability can be reduced by ensuring good contact with the heating element and measuring the EBSD map as close as possible to the thermocouple. Of course, an alternative is to weld an additional thermocouple on the sample to better control the temperature. This is technically possible but with the risk of unexcluded local changes in the sample composition and uncontrolled in-situ thermal cycle. In our in-situ heating experiments, the temperature required for complete transformation of OT-steel-B into austenite varied from 810 °C to 840 °C from one sample to another.

To overcome these limitations in the analysis of microstructures inherited by phase transformation, we have conducted our HT experiments by step wise heating to detect precisely the transformation start and finish (Ac1/Ac3) temperature and set the desired temperature from these reference temperatures (instead of considering the absolute temperature measured by the thermocouple).

Using this strategy, we were able to evaluate the grain size refinement with austenization cycles. Fig. 4 (a-b) shows the EBSD maps acquired for sample OT-steel-A and sample OT-steel-B at Ac3 + 30 °C after fifteen minutes (t₁₅) holding time (the EBSD analysis has been performed outside of the vertical band seen in Fig. 3). The average grain size value and the grain size distribution for each sample is indicated respectively in Fig. 4c. The grain size is given by the equivalent diameter of the detected grains weighted by their surface fraction (grain detection was obtained with a threshold angle of 5° ignoring twin boundaries). Notice that the EBSD data have been acquired with a step size of 0.3 µm, a map size of 180 × 130 µm and an acquisition speed of 970 pps resulting in a total acquisition time per map of 4 min. The average indexing rate of the maps was above 91 %. The grain size analysis was performed on cleaned and extrapolated data up to 100 % (the cleaning step consisted in correcting the limited indexing errors induced by pseudo-symmetries and in replacing “wrongly” indexed pixels identified as small grains of 1 to 3 pixels by the orientations of the nearest neighbors).

The result shows that the additional austenization cycle seen by sample B reduced the austenite grain size from 7.1 µm to 5.1 µm, confirming austenite grain refinement through austenization. Similar results were obtained by indirect parent map reconstruction [44].

## 3.2. Complementarities between in-situ measurements and indirect reconstruction

In the past decade, researchers have developed various indirect methods for reconstructing parent microstructures from orientation maps acquired at room temperature (RT) after displacive phase transformations. Whatever the reconstruction algorithm [14,45,46], the approach relies on some assumptions and the final result is often influenced by parameter selection, thereby raising significant questions about the reliability of the reconstructed parent contours. Nowadays in-situ measurements give access to the ground truth parent map and offer the possibility of validating the predictive capabilities of reconstruction algorithms.

Fig. 5 shows a comparison on sample OT-steel-B between in-situ measurement (Fig. 5 (a)) and reconstruction (Fig. 5 (c)) from EBSD data acquired at RT after fast cooling at 3 °C/s to form a lath like microstructure (Fig. 5 (a)). The reconstruction was performed using Merengue, our homemade reconstruction tool [9]. A very good agreement was observed between both maps in terms of austenite orientations, grains size and shape. However, for specific configurations, the reconstruction encountered specific challenges: (1) Certain twins remained undetected in a limited number of grains (indicated by white stars), (2) some regions could not be reconstructed (appearing black in

<figure><img src="image_4.png" /><figcaption>Fig. 3. HT austenite map acquired at 840 °C on sample OT-steel-B: (a) Pattern Quality map, (b) Grain boundary map (misorientation angles between 5 and 15° in white, >15° in black, >45° in blue, and > 58° in red). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Fig. 4. HT misorientation map for austenite: EBSD data acquired at $Ac_3 + 30$ °C on sample OT-steel-A (a), OT-steel-B (b) and the corresponding grain size distribution (c).</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>Fig. 5. EBSD analysis of OT-Steel-B (a) IPF-Z map of austenite acquired in-situ at $910$ °C, (b) Band Contrast map of the inherited microstructure at RT after in-situ cooling at $3$ °C/s and (c) IPF-Z map of reconstructed austenite. (Z: normal to the acquisition plane).</figcaption></figure>

the map) and (3) Small grains were occasionally not detected and covered by surrounding grains. These limitations stem mainly from an insufficient number of local crystallographic variants to determine a unique, high-confidence austenite orientation. Despite those minor discrepancies, the reconstruction yields reliable and realistic results that can be confidently interpreted. Similar comparisons between in-situ measurement and reconstruction are also available in [36].

However, one has to keep in mind that the quality of the reconstruction significantly depends on the existence of an OR between parent and daughter phases; the stricter this relationship, the better the reconstruction. As a consequence, for microstructure inherited from a more diffusive process, like upper bainite or ferrite, indirect reconstructions are at best limited or even out of purpose. Notice also that reconstruction methods are also able to refine the OR between parent and inherited phases. This point is discussed further in section 3.3 with a comparison between reconstruction and direct measurement.

Finally, combining reconstruction and in-situ measurement offers a powerful approach to understanding microstructure evolution during thermal cycling. Fig. 6 illustrates an original strategy schemed in Fig. 6 (a): starting with sample OT-steel-A, obtained after a single austenization treatment (Fig. 6 (b)), the prior austenite map was reconstructed (Fig. 6 (c)). Subsequent in-situ reheating of sample OT-steel-A to $870$ °C revealed critical microstructural dynamics on the HT austenite map (Fig. 6 (d)). Some prior austenite grains are partially recovered during heating. Inherited variants within those grains transform by developing back their common parent orientation (examples are marked by stars in Fig. 6). Residual austenite may contribute to this “memory effect”. Significant new austenite nucleation occurs also during heating, explaining grain refinement from the austenization cycle.

These key observations confirm grain refinement during

austenization and provides insights to analyze austenite nucleation sites during thermal heating. This example demonstrates the complementary power of reconstruction and direct in-situ measurement in unveiling complex microstructural transformation mechanisms.

### 3.3. In-situ EBSD addressing challenges beyond traditional reconstruction methods

This section presents two results obtained on the 10Cr/10Cr-ODS steels highlighting the capabilities of in-situ EBSD to address challenges beyond traditional reconstruction methods. The first result demonstrates a case where it was impossible to reconstruct the HT austenite microstructure, revealing the limitations of indirect reconstruction methods and the unique insights offered by in-situ analysis. The second result exemplifies the monitoring of the martensitic transformation during cooling, coupled with a comparative OR evaluation from direct and indirect data analysis.

#### 3.3.1. Parent austenite grains of ODS steels

This analysis builds upon previous work [41] that highlights the limitations of reconstruction methods in assessing the prior austenite microstructure of a 10Cr ODS steel austenitized at $1100$ °C and quenched at $10$ °C/s. The same steel was subjected to in-situ heat treatment within a scanning electron microscope (SEM). Fig. 7 presents EBSD maps of the austenite microstructure acquired at $1000$ °C with a step size of 75 nm over an area of $48 \times 36$ µm², at an acquisition speed of 970 pps. Obtaining an EBSD map at this elevated temperature with such high spatial resolution represents a significant technical achievement. The HT microstructure, characterized by ultrafine austenitic grains with random orientations, is reported for the first time using high-
<figure><img src="image_7.png" /><figcaption>**Fig. 6.** (a) Schematic of the applied thermal cycle, (b) Martensite map of OT-steel-A obtained after a single heat treatment, (c) Prior austenite map obtained by reconstruction from (b), (d) Austenitic map acquired in-situ at 870 °C (all maps are plotted according to the IPF//Z colour code).</figcaption></figure>

temperature EBSD. However, in-situ high-temperature measurements in ODS steels present certain challenges; the sample surface developed significant localized depressions (similar to highly localized thermal etching), which create shadowing on the EBSD patterns and prevent their indexation. Consequently, the indexation rate of the map does not exceed 65 % (non-indexed pixels are shown in black in Fig. 7). These surface irregularities were attributed to the presence of ODS particles, as they were not observed in non-ODS 10Cr steel (as evidenced in the next section).

Despite this limitation, the austenite microstructure is clearly revealed. The extremely fine austenite grains, with an average grain size of approximately 0.5 µm, are probably the core issue limiting indirect reconstruction. Indeed, with such small austenite grains, the resulting martensitic microstructure tends to adopt a more granular morphology (as observed in [41]), with fewer crystallographic variants per austenite grain. This hinders accurate reconstruction of the austenite grains.

### 3.3.2. Monitoring the martensitic transformation during cooling

In-situ EBSD analysis was performed on 10Cr steel to capture the progression of the phase transformation by tracking the same region of interest (ROI) during cooling from 950 °C to a plateau at 380 °C with a controlled cooling rate of 3 °C/s and finally to RT. Martensite was detected at 390 °C, which is 5 °C higher than the martensite start temperature determined by dilatometry (385 °C). Fig. 8 illustrates the cooling scheme with two key EBSD acquisitions: an almost fully austenite map at 390 °C and a partially transformed at 380 °C. Data were collected with a step size of 0.1 µm over a map area of approximately 115 × 85 µm² with acquisition speed of 970 pps. At 380 °C, the transformation into martensite was incomplete with retained blocky and interlath austenite accounting for 25 % of the microstructure.

Fig. 9 focuses on a single prior austenite grain extracted from the overall microstructure shown in Fig. 8 (yellow boxed region). Figs. 9 (a-d) illustrate the dynamic transformation of the microstructure throughout the cooling process of the analyzed grain from 390 °C to RT. Each panel displays the IPF-Z map and the {100} pole figure of either austenite or martensite. Starting from a single austenite orientation (Fig. 9 (a)) excluding annealing twins measured at HT), the progression of the martensitic transformation induced an orientation gradient within the retained austenite phase, resulting from strain-induced transformation plasticity (Fig. 9 (b)). As transformation progresses towards RT, this orientation gradient was progressively transferred to the martensitic variants, resulting in a significant spread in the martensite orientations at RT (Fig. 9 (c)). Notably, the strict orientations inherited from the initial austenitic orientation were not distinctly visible. Instead, an orientation spread was observed, that continuously decorates the Bain domains. This type of orientation spread has been frequently reported in steels [47,48]. The parent austenite grain was finally reconstructed from the RT map in Fig. 9 (c). Interestingly, the reconstruction from the RT orientations using a strict OR (discussed in the next section) restored austenite orientations, exhibiting a spread closely resembling that observed at 380 °C (Fig. 9(d)). Averaging these reconstructed austenite orientations yielded a grain orientation remarkably close to the orientation measured prior to the start of transformation (represented by crosses in the pole figure). The directly measured orientation at 390 °C was characterized by the Euler angles: $\phi_1 = 198.7^\circ$, $\Phi = 31.3^\circ$, and $\phi_2 = 18.0^\circ$. The averaged reconstructed orientation was determined to be: $\phi_1 = 198.4^\circ$, $\Phi = 30.5^\circ$, and $\phi_2 = 18.2^\circ$. These orientations differ by only 0.82°, demonstrating high consistency and accuracy in the reconstruction process.

Beyond the comparison between measured and reconstructed orientations, it is also important to examine the OR itself. This is particularly crucial in the case of steels where several ORs, often close to one another, have been reported. Classical ORs such as Kurdjumov-Sachs (K-S), Nishiyama-Wassermann (N-W) or Greninger-Troiano (G-T) differ by only a few degrees, which makes their distinction difficult but fundamental for the interpretation of the transformation mechanisms
<figure><img src="image_8.png" /><figcaption>Fig. 7. EBSD maps acquired on 10Cr ODS steel at 1000 °C: (a) Band contrast map with boundary angles, showing low-angle grain boundaries (3° to 15°) in yellow and high-angle boundaries (>15°) in black; (b) Austenite IPF-Z map. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>Fig. 8. In-situ EBSD temperature profile during cooling of 10Cr steel over the same ROI: Map at 390 °C presents the fully austenitic phase, map at 380 °C displays band contrast and retained austenite in IPF-Z coloring.</figcaption></figure>

and also for the accuracy of reconstructions [12].

Fig. 10 plots the different OR obtained on the 10Cr steel through both direct measurements and indirect reconstructions. The OR is referenced by its deviation from KS-OR, considering the angle $\theta_2$ between the KS coincidence directions $[1\bar{1}0]_{fcc}$ and $[\bar{1}11]_{bcc}$ as a function of the angle $\theta_1$ between the KS coincidence planes $(111)_{fcc}$ and $(110)_{bcc}$. Blue and green colors distinguish results obtained by direct measurements and indirect reconstructions, respectively, considering either averaged ORs obtained over the whole map (see triangles) or averaged ORs obtained over the specific austenite grain highlighted in Fig. 9 (see crosses).

The averaged OR was determined through a two-step reconstruction process. First, clusters of the parent grain (with high confidence indices) were identified using G-T OR. Second, the OR was progressively refined using the method implemented in Merengue software [9]. This approach establishes an average OR that minimizes the deviation between daughter and reconstructed parent clusters across the entire dataset. The OR refinement process works iteratively, alternating between reconstructing austenite using the current OR and refining this OR until no further improvement in reconstructed austenite fraction is achieved.

The OR determined directly from the EBSD map at 380 °C was close to that obtained through indirect reconstruction, with a difference of less

than 0.5°. Similar deviations were observed when considering only a specific grain. This is a step forwards to validate the average OR deduced by reconstruction. Moreover, the analysis of the HT map shows that the OR along different austenite/martensite interfaces varied about 3° (even inside the same grain). This possibly reflects that ORs in steels often exhibit systematic deviations from a theoretical unique OR and can vary according to the local stress state during transformation progress.

### 3.4. Discussion

The results presented in this study provide a comprehensive comparison between direct EBSD measurements at HT and indirect reconstruction methods for characterizing phase transformations in steels. This comparison reveals important complementarities and limitations worth discussing in detail.

#### 3.4.1. HT- EBSD capabilities

The different examples given in this contribution confirms that HT-EBSD enables successful direct observation of microstructural evolution induced by phase transformations in steels. As shown in Fig. 3 and Fig. 9, critical features of the HT austenite microstructure can easily be revealed. It includes the detection of annealing twins and grain size heterogeneities that reconstruction methods typically overlook. For the first-time, ultrafine parent austenite grains in 10 % chromium oxide dispersion-strengthened (ODS) steel were observed at HT (Fig. 7). Moreover, austenite grain refinement through austenitization cycles in a medium carbon steel could be confirmed (Figs. 4 and 6).

The impressive progress in HT-EBSD is due to (1) the development of compact heating stages with optimized design within the SEM chamber at 70° tilt, capable of reaching temperatures up to 1100 °C and (2) the new generation of detectors with infrared filtering offering high acquisition speed over 1000 pps with no compromise on pattern quality. The speed of the EBSD system is extremely important for in situ experiments to allow dynamic microstructural changes to be characterized. Of course, heating and cooling rates depend on the performance of the stage. Notably, some stages use laser heating, offering the highest heating rates [49]. However, acquisition during heating or cooling is inevitably limited by the thermal drift. In practice, acquisitions are only possible at a plateau temperature after the whole set-up has stabilized for a certain period; this period increases with the rate of heating or cooling. Automated acquisition as proposed with the TANIST platform [36] is a promising way to optimize in situ HT-EBSD acquisitions, especially to capture dynamic events during heating or cooling.

These advances make HT-EBSD a promising technique, with various applications, including in-situ observation of nucleation sites [50,51],
<figure><img src="image_10.png" /><figcaption>**Fig. 9.** EBSD maps and {100} pole figures illustrating the microstructural evolution of a single prior austenite grain during cooling from 390 °C to RT: (a) Austenite microstructure at 390 °C, (b) Retained austenite at 380 °C, (c) Martensitic microstructure at RT, (d) Reconstructed austenite from (e) Black crosses on the pole figure show the averaged orientations of the reconstructed austenite.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**Fig. 10.** θ₂ angle between |**1**1**1**|<sub>bcc</sub> and |**1**0**0**|<sub>fcc</sub> as a function of θ₁ angle between (110)<sub>bcc</sub> and (111)<sub>fcc</sub>, measured within the analyzed grain and across the entire map at 380 °C. Triangles and crosses represent the ORs averaged across the entire map and within the austenite grain analyzed in Fig. 9, respectively. Blue symbols indicate ORs obtained through direct measurement at austenite/martensite interfaces, while green symbols denote those determined by reconstruction.</figcaption></figure>

tracking of growth mechanisms during solid-state transformations [52], monitoring of crystallographic ORs between parent and product phases (see Fig. 10), and interface migration during diffusional and diffusionless transformations. Furthermore, combining HT-EBSD with EDS enables simultaneous chemical mapping [53], providing additional insights into phase transformation mechanisms.

### 3.4.2. HT-EBSD limitations

Despite its advantages, HT-EBSD faces several challenges. A significant limitation relies on accurate temperature control and measurement at the precise location of observation. Temperature differences between the thermocouple and the observed surface can exceed several tens of degrees, affecting transformation temperature measurements. Our step-wise heating approach mitigates but does not eliminate this issue.

Another inherent limitation is that direct observations of orientation changes during in-situ experiments can be subjected to free surface effects. Special care must be given in case of displacive transformation where free surfaces may affect variant selections. In our study, the

martensitic transformation in 10Cr steel was analyzed after cooling to RT. Fig. 11 presents the overlapped band contrast and misorientation boundaries maps measured at the top surface and in the bulk, along with the corresponding misorientation angle distribution. The analysis in the bulk was performed after additional mechanical polishing of the sample, until half of the initial thickness was removed (final thickness: 0.5 mm). Similar results were observed on other samples polished by removing the top surface (~9 µm), confirming the reproducibility of the observation.

A difference in the frequency of high and low misorientation boundaries is observed between the two regions. The observed differences can be attributed to free surface effects on the transformation process. At the surface, elastic strain energy is more easily relieved compared to the bulk, where mechanical constraints influence the selection of martensitic variants. The difference could also be due to a difference in cooling rate between the top surface and the bulk.

The top surface analysis under vacuum may also alter the surface composition due to preferential evaporation of volatile elements. On very low carbon steel, top surface decarburization has been observed [44]. In addition, any changes to the surface during the experiment can have deleterious effects on the pattern quality and indexing rates, as observed in our 10Cr-ODS steel experiments (see Fig. 7). Also, other surface irregularities like thermal etching of grain boundaries could affect nucleation sites.

According to these surface effects, care must be taken to obtain complementary information from the bulk measurements to confirm any conclusions drawn from the surface measurements.

### 3.4.3. In-situ HT-EBSD versus indirect reconstruction

HT-EBSD is ideal for revealing transformation dynamics and intermediate states during thermal processing, providing essential insights into transformation mechanisms and kinetics. Meanwhile, reconstruction methods offer efficient post-mortem analysis of the final microstructure to recover the parent microtexture. It excels in revealing parent microstructures in case of industrial processes like additive manufacturing [54] or thermomechanical treatments [10], where HT-EBSD is impractical.

Of course, reconstruction methods have their own limitations that are already well identified [9,10,16]. Fundamentally, sufficient variants are required to accurately determine parent orientations. As
<figure><img src="image_12.png" /><figcaption>Fig. 11. Band contrast with misorientation boundaries map of 10Cr steel after cooling to RT: (a) top surface, (b) bulk and (c) misorientation angle histogram.</figcaption></figure>

demonstrated in our 10Cr-ODS steel with ultrafine grains, limited variant formation per austenite grain renders reconstruction ineffective [42]. HT-EBSD faces no such limitation, providing direct measurement regardless of variant selection. Finally, the quality of the reconstruction depends on the used OR; its knowledge is a pre-requisite. Indirect reconstruction performs well for displacive transformations that respect a ‘strict’ OR (like martensite). In contrast, HT-EBSD performance remains consistent regardless of transformation type.

HT EBSD can also be a valuable tool to establish ground truth validation for indirect reconstruction. In-situ HT-EBSD provides direct measurement of parent phase orientations and morphologies, against which reconstruction algorithms can be validated and refined.

The different comparisons made in this study demonstrate that in-
direct reconstruction is able to accurately determine of parent orienta-
tion provided sufficient variants are available, as also reported in [36].
OR selection is also important and impacts the quality of the recon-
struction [17]. HT EBSD enables the evaluation of OR during ongoing
transformation, allowing for comparison with ORs determined through
reconstruction methods. Such validation is critical for establishing
confidence in reconstruction methods, particularly when applied to
materials or conditions where direct measurement is impractical.

In future, in-situ HT-EBSD could also provide valuable data for
refining reconstruction algorithms to overcome their actual limitations
in particular to better recover grain morphologies and annealing twins.

Finally, we have shown examples of judicious combination of both
methodologies (see Fig. 6) This is the most efficient way to develop a
comprehensive understanding of phase transformation induced micro-
structures in steels.

4. Conclusion

This study has reviewed both the capabilities and complementarities
of HT-EBSD and crystallographic reconstruction techniques for
analyzing phase transformations in steels. Our different application
examples highlight several key points:

* A modern HT-EBSD detector coupled with a compact heating stage is highly efficient for direct observation of austenite microstructures at high temperatures, with fast acquisition speed (up to 1000 pps) and good indexing reliability. This enables detailed parent-phase characterization with high spatial resolution. HT EBSD opens up new potential for studying phase transformations like the possibility to quantify austenite grain refinement during thermal cycling, identify austenite nucleation sites and track martensitic transformations in real time.

* Certain limitations persist for HT-EBSD, including challenges such as achieving precise temperature control across the analyzed sample area, managing SEM image drift, and mitigating surface effects

pronounced for some materials and transformation conditions (for
example in ODS steels and for martensitic transformations).

* Parent austenite reconstruction algorithms perform well for displacive transformations following known ORs, but face limitations in cases with insufficient variants per parent grain (as evidenced in the ultrafine-grained 10Cr-ODS steel) or less strict ORs. Precise grain morphology restoration and twin identification remain challenging. In this context, HT-EBSD not only provides more accurate austenite microstructures but also serves as ground-truth validation for future refinement of reconstruction algorithms.

* Direct measurement of ORs between austenite and martensite confirmed the validity of ORs refined through reconstruction methods, with deviations <0.5°. However, local OR variations up to 3° within the same austenite grain highlight the influence of local stress conditions during transformation.

In summary, while crystallographic reconstruction remains a valu-
able approach for analyzing quenched microstructures inherited from
industrial processes, HT-EBSD offers complementary advantages in
challenging scenarios. The strategic combination of both techniques, as
demonstrated in this study, enables researchers to leverage their
respective strengths, for a more comprehensive understanding of
microstructural evolution during thermal processing.

CRediT authorship contribution statement

M. Fekih: Writing – original draft, Software, Funding acquisition, Methodology, Formal analysis, Validation, Investigation, Data curation.

T. Martinez Ostromujof: Writing – review & editing, Investigation, Methodology, Data curation, Software, Formal analysis. L. Germain:

Supervision, Funding acquisition, Writing – review & editing, Method-
ology, Data curation. M. Piette: Resources. D. Sornin: Resources. N.
Gey: Supervision, Data curation, Writing – original draft, Methodology,
Funding acquisition, Resources, Investigation, Conceptualization.

Declaration of competing interest

The authors 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

Data will be made available on request.

Acknowledgements

We acknowledge the experimental facilities of the MicroMat
platform from LEM3 (Université de Lorraine - CNRS UMR 7239). We also thank Kammrath & Weiss for fruitful discussions and their support in temperature measurements in the SEM chamber under EBSD conditions.

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