# Effect of cooling rate on the microstructures of three low carbon alloys with different manganese and molybdenum contents

Anna Benarosch, Bernard Marini, Caroline Toffolon-Masclet, Zofia Trzaska, Estelle Meslin, Ivan Guillot

► To cite this version:

Anna Benarosch, Bernard Marini, Caroline Toffolon-Masclet, Zofia Trzaska, Estelle Meslin, et al. Effect of cooling rate on the microstructures of three low carbon alloys with different manganese and molybdenum contents. Metallurgical Research & Technology, 2022, 119 (520), pp.520. 10.1051/metal/2022059. cea-04604033

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# Effect of cooling rate on the microstructures of three low carbon alloys with different manganese and molybdenum contents

Anna Benarosch, Bernard Marini, Caroline Toffolon-Masclet, Zofia Trzaska, Estelle Meslin, Ivan Guillot

<figure><img src="image_1.png" /></figure>

► To cite this version:

Anna Benarosch, Bernard Marini, Caroline Toffolon-Masclet, Zofia Trzaska, Estelle Meslin, et al. Effect of cooling rate on the microstructures of three low carbon alloys with different manganese and molybdenum contents. Metallurgical Research & Technology, 2022, 119 (520), pp.520. 10.1051/metal/2022059. cea-04604033

HAL Id: cea-04604033

https://cea.hal.science/cea-04604033v1

Submitted on 6 Jun 2024

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
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# Influence of Mo and Mn on microstructures in low alloy steels at different cooling rates

Anna Benarosch<sup>a,b</sup>, B. Marini<sup>a</sup>, C. Toffolon-Masclet<sup>a</sup>, Z. Trzaska<sup>b</sup>, E. Meslin<sup>a</sup>, I. Guillot<sup>b</sup>

## Authors :

- Anna Benarosch, +33 1 69 08 21 66, anna.benarosch@cea.fr

- Bernard Marini, Dr., +33 1 69 08 56 13, bernard.marini@cea.fr

- Caroline Toffolon-Masclet, Dr., +33 1 69 08 53 57, caroline.toffolon@cea.fr

- Zofia Trzaska, Dr., +33 1 49 78 13 45, zofia.trzaska@univ-paris13.fr

- Estelle Meslin, Dr., +33 1 69 08 21 68, estelle.meslin@cea.fr

- Ivan Guillot, Pr., +33 1 49 78 13 45, guillot@glvt-cnrs.fr

## Abstract

Low-alloy 16 to 20MND5 steels are used for the fabrication of nuclear reactor components. During manufacturing, an austenitization step is followed by quenching: different types of microstructures can appear during this process. Understanding the impact on these microstructures of two of the most important alloying elements of 16 to 20MND5 steels, Mo and Mn, can help predict how mechanical properties will evolve during tempering and during aging. The impact of molybdenum and manganese, as well as the impact of cooling rate, are studied on microstructures of three model steels, FeCMo, FeCMn and FeCMoMn, using continuous cooling transformation (CCT) diagrams and electron

<footnote><sup>a</sup> CEA, Energy Division, Applied Sciences and Simulation for Low Carbon Energy, Department of Materials for Nuclear Applications, 91191 Gif sur Yvette cedex, France</footnote>

<footnote><sup>b</sup> Univ Paris Est, CNRS, ICMPE, UMR 7182, F-94320 Thiais, France</footnote>
# Influence of Mo and Mn on microstructures in low alloy steels at different cooling rates

Anna Benarosch<sup>a,b</sup>, B. Marini<sup>a</sup>, C. Toffolon-Masclet<sup>a</sup>, Z. Trzaska<sup>b</sup>, E. Meslin<sup>a</sup>, I. Guillot<sup>b</sup>

## Authors :

- Anna Benarosch, +33 1 69 08 21 66, anna.benarosch@cea.fr

- Bernard Marini, Dr., +33 1 69 08 56 13, bernard.marini@cea.fr

- Caroline Toffolon-Masclet, Dr., +33 1 69 08 53 57, caroline.toffolon@cea.fr

- Zofia Trzaska, Dr., +33 1 49 78 13 45, zofia.trzaska@univ-paris13.fr

- Estelle Meslin, Dr., +33 1 69 08 21 68, estelle.meslin@cea.fr

- Ivan Guillot, Pr., +33 1 49 78 13 45, guillot@glvt-cnrs.fr

## Abstract

Low-alloy 16 to 20MND5 steels are used for the fabrication of nuclear reactor components. During manufacturing, an austenitization step is followed by quenching: different types of microstructures can appear during this process. Understanding the impact on these microstructures of two of the most important alloying elements of 16 to 20MND5 steels, Mo and Mn, can help predict how mechanical properties will evolve during tempering and during aging. The impact of molybdenum and manganese, as well as the impact of cooling rate, are studied on microstructures of three model steels, FeCMo, FeCMn and FeCMoMn, using continuous cooling transformation (CCT) diagrams and electron

<footnote><sup>a</sup> CEA, Energy Division, Applied Sciences and Simulation for Low Carbon Energy, Department of Materials for Nuclear Applications, 91191 Gif sur Yvette cedex, France</footnote>

<footnote><sup>b</sup> Univ Paris Est, CNRS, ICMPE, UMR 7182, F-94320 Thiais, France</footnote>
backscattering diffraction characterizations. Considering CCT diagram and hardness, FeCMoMn is found to be a good model alloy for 16 to 20MND5 steels. The presence of molybdenum or manganese does not modify the misorientation characteristics of martensite. In bainitic microstructures however, the presence of Mn seems to favor the presence of block boundaries with a misorientation of about 59° around [433]. Investigations carried out in a prior austenite grain (PAG) shows a continuous evolution in the crystallographic features from martensite to slowly cooled bainite, with the presence of block and sub-block boundaries. The microstructure becomes coarser with decreasing cooling rate, with fewer different crystallographic orientations in one PAG: where in the highest cooling rate three texture components are present, some PAG grains in the slowest cooled microstructure only one or two texture components per PAG can be found.

## Keywords

low-alloy steel, continuous cooling transformation, electron backscattered diffraction (EBSD), phase transformation, misorientation

## 1. Introduction

Low-alloy steels such as 16 to 20MND5, similar to A508 steels, are commonly used in the fabrication of nuclear pressure vessels, pressurizers and steam generators. 16MND5 contains 0.16wt% of carbon while 20MND5 contain 0.20wt% of C. The bainitic microstructures of these steels possess a good combination of strength, fracture toughness and weldability. Molybdenum and manganese are, with nickel, the main alloying elements. Segregations of carbon and alloying elements are observed in large ingots, at a macro- to a microscopic scale, which can alter locally the expected properties [1-3]. Furthermore, while a bainitic microstructure is targeted during the elaboration of the components, due to the important thickness, a gradient of microstructures forms during quenching, from allotriomorphic ferrite to martensite, with, potentially, very different mechanical properties. Once the
backscattering diffraction characterizations. Considering CCT diagram and hardness, FeCMoMn is found to be a good model alloy for 16 to 20MND5 steels. The presence of molybdenum or manganese does not modify the misorientation characteristics of martensite. In bainitic microstructures however, the presence of Mn seems to favor the presence of block boundaries with a misorientation of about 59° around [433]. Investigations carried out in a prior austenite grain (PAG) shows a continuous evolution in the crystallographic features from martensite to slowly cooled bainite, with the presence of block and sub-block boundaries. The microstructure becomes coarser with decreasing cooling rate, with fewer different crystallographic orientations in one PAG: where in the highest cooling rate three texture components are present, some PAG grains in the slowest cooled microstructure only one or two texture components per PAG can be found.

## Keywords

low-alloy steel, continuous cooling transformation, electron backscattered diffraction (EBSD), phase transformation, misorientation

## 1. Introduction

Low-alloy steels such as 16 to 20MND5, similar to A508 steels, are commonly used in the fabrication of nuclear pressure vessels, pressurizers and steam generators. 16MND5 contains 0.16wt% of carbon while 20MND5 contain 0.20wt% of C. The bainitic microstructures of these steels possess a good combination of strength, fracture toughness and weldability. Molybdenum and manganese are, with nickel, the main alloying elements. Segregations of carbon and alloying elements are observed in large ingots, at a macro- to a microscopic scale, which can alter locally the expected properties [1–3]. Furthermore, while a bainitic microstructure is targeted during the elaboration of the components, due to the important thickness, a gradient of microstructures forms during quenching, from allotriomorphic ferrite to martensite, with, potentially, very different mechanical properties. Once the
as-quenched microstructure is obtained, the component is tempered at approximately 650°C for
several hours.

This heterogeneity in chemical composition and in microstructure can be problematic for predicting
how the material will evolve under tempering and during aging. A good knowledge of the impact of
the alloying elements on the microstructure and of the microstructure itself can in turn help predict
the expected mechanical properties, for assessing the durability of the component during the plant life
and its possible extension. Indeed, under service conditions, the component is subjected for decades
to thermo-mechanical solicitations and, for the vessel especially, to neutron irradiation. Due to these
conditions, solute alloy elements and impurities segregate and precipitate, and are the cause of
embrittlement mechanisms. The microstructure plays an important role here for two reasons. First,
these segregations and precipitations are microstructure-dependent i.e. the misorientation angle/axis
pair, as well as the boundary plane, have an influence on the segregations of impurities, such as
phosphorus [4]. Second, the component can undergo embrittlement by the weakening of the
interfaces, which is microstructure-dependent as well. While the grain size is usually considered to be
an important factor for this type of embrittlement, the initial microstructure also comes into
consideration. For instance, low alloy steels with a martensitic microstructure are more sensitive to
embrittlement due to thermal aging than the same steel with a bainitic microstructure [5, 6].
Furthermore, neutron irradiation and thermal aging can trigger intergranular fracture in coarse-
grained martensite but not in coarse-grained bainite [7]. However, the impact of the microstructure
on the embrittlement induced by neutron irradiation is still debated [8].

Previous attempts were made to understand the influence of alloying elements on bainite and
martensite in low-alloy steels. For instance, it was shown that an increasing amount of molybdenum
or manganese lead to a finer bainitic microstructure [9–11], this being explained by the fact that Mo
and Mn decreases the bainite transformation temperature [10–12]. Other studies demonstrated that
the martensite formation depends on the Mn content [13] and that an increasing amount of Mo or Mn
decreases the martensite formation temperature [10, 14]. A systematic study of the effect of several
as-quenched microstructure is obtained, the component is tempered at approximately 650°C for
several hours.

This heterogeneity in chemical composition and in microstructure can be problematic for predicting
how the material will evolve under tempering and during aging. A good knowledge of the impact of
the alloying elements on the microstructure and of the microstructure itself can in turn help predict
the expected mechanical properties, for assessing the durability of the component during the plant life
and its possible extension. Indeed, under service conditions, the component is subjected for decades
to thermo-mechanical solicitations and, for the vessel especially, to neutron irradiation. Due to these
conditions, solute alloy elements and impurities segregate and precipitate, and are the cause of
embrittlement mechanisms. The microstructure plays an important role here for two reasons. First,
these segregations and precipitations are microstructure-dependent i.e. the misorientation angle/axis
pair, as well as the boundary plane, have an influence on the segregations of impurities, such as
phosphorus [4]. Second, the component can undergo embrittlement by the weakening of the
interfaces, which is microstructure-dependent as well. While the grain size is usually considered to be
an important factor for this type of embrittlement, the initial microstructure also comes into
consideration. For instance, low alloy steels with a martensitic microstructure are more sensitive to
embrittlement due to thermal aging than the same steel with a bainitic microstructure [5, 6].
Furthermore, neutron irradiation and thermal aging can trigger intergranular fracture in coarse-
grained martensite but not in coarse-grained bainite [7]. However, the impact of the microstructure
on the embrittlement induced by neutron irradiation is still debated [8].

Previous attempts were made to understand the influence of alloying elements on bainite and
martensite in low-alloy steels. For instance, it was shown that an increasing amount of molybdenum
or manganese lead to a finer bainitic microstructure [9–11], this being explained by the fact that Mo
and Mn decreases the bainite transformation temperature [10–12]. Other studies demonstrated that
the martensite formation depends on the Mn content [13] and that an increasing amount of Mo or Mn
decreases the martensite formation temperature [10, 14]. A systematic study of the effect of several
alloying elements including Mo and Mn demonstrated that Mo nor Mn had an impact on the martensitic hardness [15]. Maki *et al.* [16] investigated the impact of Mo and Mn on lath martensite morphology using model ternary Fe-C-X (X=Mo or Mn) alloys. Based on optical observations only, they concluded that an addition of about 1wt% of Mo or Mn does not impact the martensite packet size. When a finer characterization of the bainitic or martensitic microstructures is carried out [17–20], the impact of Mo or Mn is not explored.

As for the influence of the cooling rate on microstructural features, little information is available in literature. Su *et al.* [21] combined dilatometric and electron backscattered diffraction (EBSD) measurements but only for phase identification. Graf *et al.* [22] investigated by EBSD, on a low-alloy steel, the influence of cooling rate on martensite. They found no significant dependence of the cooling rate on the block and packet size but nothing was mentioned about the misorientation angles nor axes.

This is why, to the authors’ knowledge, no study reports the evolution of important crystallographic features, such as the distribution of misorientation angles and axes, as a function of both cooling rate and alloying element, for 16 to 20MND5 type of steels.

To understand the role of Mo and Mn on the different microstructures, two ternary high-purity model alloys, FeCMo and FeCMn, and one quaternary high-purity model alloy, FeCMoMn were investigated.

To study the influence of the cooling rate on the microstructure of these alloys, continuous cooling (CCT) diagrams were established, and the obtained microstructures characterized by EBSD. All the alloys contain a targeted concentration in carbon of 0.2wt%. The concentration in manganese in the Mn-bearing alloys was kept close to the concentration in 16 to 20MND5. However, the concentration in Mo was raised to a targeted concentration of 1wt%, against 0.43-0.6wt% in the industrial alloy. This was done to shift the precipitation of ferrite towards lower cooling rates [23, 24], to be able to study fully bainitic microstructures for the slowest cooling rates as possible.
alloying elements including Mo and Mn demonstrated that Mo nor Mn had an impact on the martensitic hardness [15]. Maki *et al.* [16] investigated the impact of Mo and Mn on lath martensite morphology using model ternary Fe-C-X (X=Mo or Mn) alloys. Based on optical observations only, they concluded that an addition of about 1wt% of Mo or Mn does not impact the martensite packet size. When a finer characterization of the bainitic or martensitic microstructures is carried out [17–20], the impact of Mo or Mn is not explored.

As for the influence of the cooling rate on microstructural features, little information is available in literature. Su *et al.* [21] combined dilatometric and electron backscattered diffraction (EBSD) measurements but only for phase identification. Graf *et al.* [22] investigated by EBSD, on a low-alloy steel, the influence of cooling rate on martensite. They found no significant dependence of the cooling rate on the block and packet size but nothing was mentioned about the misorientation angles nor axes.

This is why, to the authors’ knowledge, no study reports the evolution of important crystallographic features, such as the distribution of misorientation angles and axes, as a function of both cooling rate and alloying element, for 16 to 20MND5 type of steels.

To understand the role of Mo and Mn on the different microstructures, two ternary high-purity model alloys, FeCMo and FeCMn, and one quaternary high-purity model alloy, FeCMoMn were investigated.

To study the influence of the cooling rate on the microstructure of these alloys, continuous cooling (CCT) diagrams were established, and the obtained microstructures characterized by EBSD. All the alloys contain a targeted concentration in carbon of 0.2wt%. The concentration in manganese in the Mn-bearing alloys was kept close to the concentration in 16 to 20MND5. However, the concentration in Mo was raised to a targeted concentration of 1wt%, against 0.43-0.6wt% in the industrial alloy. This was done to shift the precipitation of ferrite towards lower cooling rates [23, 24], to be able to study fully bainitic microstructures for the slowest cooling rates as possible.
## 2. Materials and methods

### 2.1. Production of the high-purity model alloys and characterization of the received material

Each chemical element is purified by melting them using induction heating under successive gas fluxes: first oxygen, then hydrogen and argon. The elements are then melted together by induction heating to produce an ingot, which is homogenized by melting it again two more times. The ingot is shaped with a drop hammer and the final rod-geometry is obtained using a rotary forging machine. A final step of recrystallization is carried out at 860°C for two hours under a dynamic flux of argon.

The composition of the manufactured alloys is controlled by chemical analysis, see table 1. Mo and P concentrations are measured using Induced Coupled Plasma coupled with Atomic Emission Spectroscopy (ICP-AES), while carbon and sulfur concentrations are estimated with Combustion Infrared Absorption. Reducing Melting Infrared Absorption is used to measure oxygen, while Reducing Melting Thermal Conductivity estimates the concentration in nitrogen. Manganese content is evaluated thanks to Optical Emission Spectrometry.

<table><caption>Table 1: Composition of the high-purity alloys FeCMo, FeCMn and FeCMoMn</caption><thead><tr><th></th><th>C (wt%)</th><th>Mo (wt%)</th><th>Mn (wt%)</th><th>P (wt%)</th><th>N (wt%)</th><th>O (wt%)</th><th>S (wt%)</th></tr></thead><tbody><tr><td>FeCMo</td><td>0.18±0.02%</td><td>1.0±0.08%</td><td>-</td><td>&lt;0.005%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr><tr><td>FeCMn</td><td>0.20±0.02%</td><td>-</td><td>1.47±0.08%</td><td>&lt;0.004%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr><tr><td>FeCMoMn</td><td>0.19±0.02%</td><td>1.14±0.05%</td><td>1.46±0.05%</td><td>&lt;0.004%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr></tbody></table>

Segregations in Mo and Mn exist in the materials in some very small areas. These segregations most likely appeared during the production of the ingot, after solidification. These segregations are quantified by microprobe analysis, see figure 1 in appendix for more details. The quantification of the heterogeneities in these areas by electron microprobe analysis gives the following results: for the FeCMn alloy, the enriched areas contain a Mn concentration between 2 and 2.5wt%, as in FeCMoMn.
## 2. Materials and methods

### 2.1. Production of the high-purity model alloys and characterization of the received material

Each chemical element is purified by melting them using induction heating under successive gas fluxes: first oxygen, then hydrogen and argon. The elements are then melted together by induction heating to produce an ingot, which is homogenized by melting it again two more times. The ingot is shaped with a drop hammer and the final rod-geometry is obtained using a rotary forging machine. A final step of recrystallization is carried out at 860°C for two hours under a dynamic flux of argon.

The composition of the manufactured alloys is controlled by chemical analysis, see table 1. Mo and P concentrations are measured using Induced Coupled Plasma coupled with Atomic Emission Spectroscopy (ICP-AES), while carbon and sulfur concentrations are estimated with Combustion Infrared Absorption. Reducing Melting Infrared Absorption is used to measure oxygen, while Reducing Melting Thermal Conductivity estimates the concentration in nitrogen. Manganese content is evaluated thanks to Optical Emission Spectrometry.

<table><caption>Table 1: Composition of the high-purity alloys FeCMo, FeCMn and FeCMoMn</caption><thead><tr><th></th><th>C (wt%)</th><th>Mo (wt%)</th><th>Mn (wt%)</th><th>P (wt%)</th><th>N (wt%)</th><th>O (wt%)</th><th>S (wt%)</th></tr></thead><tbody><tr><td>FeCMo</td><td>0.18±0.02%</td><td>1.0±0.08%</td><td>-</td><td>&lt;0.005%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr><tr><td>FeCMn</td><td>0.20±0.02%</td><td>-</td><td>1.47±0.08%</td><td>&lt;0.004%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr><tr><td>FeCMoMn</td><td>0.19±0.02%</td><td>1.14±0.05%</td><td>1.46±0.05%</td><td>&lt;0.004%</td><td>&lt;0.003%</td><td>&lt;0.001%</td><td>&lt;0.001%</td></tr></tbody></table>

Segregations in Mo and Mn exist in the materials in some very small areas. These segregations most likely appeared during the production of the ingot, after solidification. These segregations are quantified by microprobe analysis, see figure 1 in appendix for more details. The quantification of the heterogeneities in these areas by electron microprobe analysis gives the following results: for the FeCMn alloy, the enriched areas contain a Mn concentration between 2 and 2.5wt%, as in FeCMoMn.
As for FeCMo, the segregated areas contain between 1 and 3.5 wt% in Mo, while for FeCMoMn they contain in average 2.5wt% of Mo. Primary carbides are present in all the alloys.

## 2.2. Dilatometry measurements

The construction of the continuous cooling transformation (CCT) diagrams for each alloy is obtained using a DT-1000 dilatometer modified by AET Technologies. The samples are of cuboid shape with dimensions of 2x3x12 mm³ taken from the rods. The chamber is purged three times using helium before thermal treatment.

* Austenitisation

All dilatometry samples are subjected to the same austenitisation conditions. Austenitization is carried out in a static helium atmosphere to limit decarburization. Austenitisation consists in two steps:

1. The sample is heated at 1.5°C.s⁻¹ until reaching 1200°C. The sample is held at this temperature for 5min.

2. The sample is cooled at 300°C.s⁻¹ until reaching 890°C. The sample is held at this temperature for 5min.

Austenitisation conditions are schematically presented figure //. The step 1 aims at reducing the local chemical heterogeneities. A compromise is found with a treatment at 1200°C for 5min to homogeneise as much as possible the matrix while limiting austenite grain growth. The step 2 aims at having an austenitisation temperature close to the industrial one before quenching. The step 2 is held for 5min so that the temperature is homogeneous throughout the dilatometry sample.

Figure // in appendix shows microprobe analysis of a dilatometry sample for each alloy before and after austenitisation.

* Quenching

The targeted cooling rates are obtained by flowing helium in the chamber and, for cooling rates of 100°C.s⁻¹ or more, helium previously cooled by liquid nitrogen is injected.

* Determination of transformation temperatures
As for FeCMo, the segregated areas contain between 1 and 3.5 wt% in Mo, while for FeCMoMn they contain in average 2.5wt% of Mo. Primary carbides are present in all the alloys.

## 2.2. Dilatometry measurements

The construction of the continuous cooling transformation (CCT) diagrams for each alloy is obtained using a DT-1000 dilatometer modified by AET Technologies. The samples are of cuboid shape with dimensions of 2x3x12 mm³ taken from the rods. The chamber is purged three times using helium before thermal treatment.

* Austenitisation

All dilatometry samples are subjected to the same austenitisation conditions. Austenitization is carried out in a static helium atmosphere to limit decarburization. Austenitisation consists in two steps:

1. The sample is heated at 1.5°C.s⁻¹ until reaching 1200°C. The sample is held at this temperature for 5min.

2. The sample is cooled at 300°C.s⁻¹ until reaching 890°C. The sample is held at this temperature for 5min.

Austenitisation conditions are schematically presented figure //. The step 1 aims at reducing the local chemical heterogeneities. A compromise is found with a treatment at 1200°C for 5min to homogeneise as much as possible the matrix while limiting austenite grain growth. The step 2 aims at having an austenitisation temperature close to the industrial one before quenching. The step 2 is held for 5min so that the temperature is homogeneous throughout the dilatometry sample.

Figure // in appendix shows microprobe analysis of a dilatometry sample for each alloy before and after austenitisation.

* Quenching

The targeted cooling rates are obtained by flowing helium in the chamber and, for cooling rates of 100°C.s⁻¹ or more, helium previously cooled by liquid nitrogen is injected.

* Determination of transformation temperatures
Transformation temperatures are determined from the dilatometry data using the deviation from the straight line method.

* Characterizations

After testing, the dilatometric samples are polished and cut in half in order to observe their section by light optical microscopy. These sections are polished and etched with 4% nital. The Vickers hardness is measured on at least 15 prints for each sample with a load of 200gf held for 10 to 15s.

2.3. Electron backscattered diffraction

Electron backscattered diffraction (EBSD) enables further characterization of the microstructures. The section of the dilatometric samples is polished using a diamond suspension down to 3µm grade and then polished with a colloidal silica suspension. The EBSD analyses are done with a MEB FEG JEOL JSM and a MEB FEG Zeiss Sigma HD, both equipped with a Bruker e-Flash HR EBSD detector. Several mappings are done for each sample so that the total area analyzed would be 2.1 mm² with a step of 1µm. The data is acquired with the Bruker Esprit software. To measure the average size of prior austenite grains (PAG), a reconstruction is done using the methodology developed by Cayron *et al.* [25] on the software ARPGE 1.2. EBSD data are treated with HKL Channel 5 provided by Oxford Instruments.

The misorientation angle distributions are the result of the average of all the mappings. Distributions with a minimum of 5° or 45° are presented, meaning that misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. The misorientation axis distributions in inverse pole figures presented in this paper are taken from one of the mappings but, for a given sample, the misorientation axis distributions in inverse pole figures were the same for every mapping. The misorientation axis distributions give, for an angular interval, the axis along which is the misorientation. The axes are given in the crystal coordinates. The densities are expressed relative to that from a random distribution [26], they are calculated in terms of multiple uniform density (mud).

EBSD analysis of microstructure is used to derive the misorientation between neighbouring grains in terms of a pair of angle and of rotation axis. These pairs are compared to the expected pairs related to
Transformation temperatures are determined from the dilatometry data using the deviation from the straight line method.

* Characterizations

After testing, the dilatometric samples are polished and cut in half in order to observe their section by light optical microscopy. These sections are polished and etched with 4% nital. The Vickers hardness is measured on at least 15 prints for each sample with a load of 200gf held for 10 to 15s.

2.3. Electron backscattered diffraction

Electron backscattered diffraction (EBSD) enables further characterization of the microstructures. The section of the dilatometric samples is polished using a diamond suspension down to 3µm grade and then polished with a colloidal silica suspension. The EBSD analyses are done with a MEB FEG JEOL JSM and a MEB FEG Zeiss Sigma HD, both equipped with a Bruker e-Flash HR EBSD detector. Several mappings are done for each sample so that the total area analyzed would be 2.1 mm² with a step of 1µm. The data is acquired with the Bruker Esprit software. To measure the average size of prior austenite grains (PAG), a reconstruction is done using the methodology developed by Cayron *et al.* [25] on the software ARPGE 1.2. EBSD data are treated with HKL Channel 5 provided by Oxford Instruments.

The misorientation angle distributions are the result of the average of all the mappings. Distributions with a minimum of 5° or 45° are presented, meaning that misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. The misorientation axis distributions in inverse pole figures presented in this paper are taken from one of the mappings but, for a given sample, the misorientation axis distributions in inverse pole figures were the same for every mapping. The misorientation axis distributions give, for an angular interval, the axis along which is the misorientation. The axes are given in the crystal coordinates. The densities are expressed relative to that from a random distribution [26], they are calculated in terms of multiple uniform density (mud).

EBSD analysis of microstructure is used to derive the misorientation between neighbouring grains in terms of a pair of angle and of rotation axis. These pairs are compared to the expected pairs related to
phase transformation models. The considered microstructures can have with austenite a Kurdjumov-Sachs (KS) [27] orientation relationship (OR), a Nishiyama-Wasserman (NW) [28] OR or a Greninger-Troiano (GT) [29] orientation relationship. Misorientation can be calculated between variants bearing one of these OR. The misorientation angle/axis pairs for KS/KS, NW/NW, KS/NW and GT/GT misorientation relationships are listed in supplementary material, in table S1.

The prior austenite grains examined in detail in the present work are selected using the boundaries with a misorientation between 20 and 49° and chosen to have a diameter close to the average size of prior austenite grains.

## 3. Results

### 3.1. Phase transformation and CCT diagrams

The CCT diagrams of the model alloys and of an industrial 20MND5 [7] steel, figure 1, cover cooling rates from 300°C.s⁻¹ to 0.5°C.s⁻¹ for ternary alloys and to 0.1°C.s⁻¹ for FeCMoMn. The size of the prior austenite grains of model alloys was reconstructed from EBSD measurements on the martensitic microstructure cooled at 300°C.s⁻¹. All three alloys presented a prior austenite grain size with an average between 85 and 96µm.
phase transformation models. The considered microstructures can have with austenite a Kurdjumov-Sachs (KS) [27] orientation relationship (OR), a Nishiyama-Wasserman (NW) [28] OR or a Greninger-Troiano (GT) [29] orientation relationship. Misorientation can be calculated between variants bearing one of these OR. The misorientation angle/axis pairs for KS/KS, NW/NW, KS/NW and GT/GT misorientation relationships are listed in supplementary material, in table S1.

The prior austenite grains examined in detail in the present work are selected using the boundaries with a misorientation between 20 and 49° and chosen to have a diameter close to the average size of prior austenite grains.

## 3. Results

### 3.1. Phase transformation and CCT diagrams

The CCT diagrams of the model alloys and of an industrial 20MND5 [7] steel, figure 1, cover cooling rates from 300°C.s⁻¹ to 0.5°C.s⁻¹ for ternary alloys and to 0.1°C.s⁻¹ for FeCMoMn. The size of the prior austenite grains of model alloys was reconstructed from EBSD measurements on the martensitic microstructure cooled at 300°C.s⁻¹. All three alloys presented a prior austenite grain size with an average between 85 and 96µm.
<figure><img src="image_2.png" /><figcaption>Figure 1: CCT diagrams for a) FeCMo, b) FeCMn and c) FeCMoMn and d) 20MND5 [7]: B=Bainite, M=Martensite, A+W F=allotriomorphic and Widmanstätten ferrite, P=pearlite, F=ferrite. (color necessary for print)</figcaption></figure>
<figure><img src="image_3.png" /><figcaption>Figure 1: CCT diagrams for a) FeCMo, b) FeCMn and c) FeCMoMn and d) 20MND5 [7]: B=Bainite, M=Martensite, A+W F=allotriomorphic and Widmanstätten ferrite, P=pearlite, F=ferrite. (color necessary for print)</figcaption></figure>
<table><caption>Figure 2: Micrographs for different cooling speeds for FeCMo (a,b,c), FeCMn (d,e,f)</caption><thead><tr><th></th><th>a) 300°C/s</th><th>b) 5°C/s</th><th>c) 0.5°C/s</th></tr></thead><tbody><tr><td>FeCMo</td><td></td><td></td><td></td></tr><tr><td>FeCMn</td><td></td><td></td><td></td></tr><tr><td>FeCMoMn</td><td></td><td></td><td></td></tr></tbody></table>
<table><caption>Figure 2: Micrographs for different cooling speeds for FeCMo (a,b,c), FeCMn (d,e,f)</caption><thead><tr><th></th><th>a) 300°C/s</th><th>b) 5°C/s</th><th>c) 0.5°C/s</th></tr></thead><tbody><tr><td>FeCMo</td><td></td><td></td><td></td></tr><tr><td>FeCMn</td><td></td><td></td><td></td></tr><tr><td>FeCMoMn</td><td></td><td></td><td></td></tr></tbody></table>
The alloys containing molybdenum have similar CCT diagrams, figures 1a) and 1c). These diagrams are split in two clear domains: one fully bainitic with a rather constant formation temperature, and one fully martensitic. However, for the FeCMn alloy, while there is a fully martensitic domain, there is no fully bainitic domain (figure 1b). For cooling rates between 20°C.s⁻¹ and 10°C.s⁻¹, EBSD measurements showed that the fraction of martensite increases with the cooling rate, which is coherent with the evolution of the Vickers hardness. The figure 2e) shows that martensite is present in rather small proportions for a cooling rate of 10°C/s. From cooling rates of 5°C.s⁻¹ and less, the FeCMn microstructure is a mix of allotriomorphic and Widmanstätten ferrite with bainite at 5°C.s⁻¹ or pearlite for lower cooling rates, see figure 2f). Observing the presence of allotriomorphic and Widmanstätten ferrite together is not surprising since the Widmanstätten ferritic laths nucleate on allotriomorphic ferrite grains, with a low misorientation angle between them [30, 31]. Furthermore, the formation of Widmanstätten ferrite in the present study is favored by the large size of the prior austenite grains.

In the fully martensitic domain, the formation temperature (Mₛ) is close for all the alloys, with an average value of 430°C, and the microstructures are similar, see figures 2a), 2d) and 2g). The experimental Mₛ value is comparable to what Andrews empirical equation predicts: 451°C for FeCMo, 414°C for FeCMn and 406°C for FeCMoMn. In that domain, the Vickers hardness is rather constant with cooling rate, with close values for all compositions.

The presence of Mn increases the martensitic hardenability of the steel, as reported before in literature [24], and martensitic microstructures can be seen until cooling rates of 30°C.s⁻¹ or even less for the FeCMn alloy, which is not the case for the FeCMo composition.

Combining molybdenum and manganese decreases the bainite formation temperature (Bₛ): it is around 650°C in the ternary alloys, while it is 50°C lower for FeCMoMn, with an average of 585°C, see figures 1a) to 1c). Wang *et al.* also reported that both Mn and Mo contribute to decrease Bₛ, after compiling several CCT diagrams through an artificial neural network model [32]. Combining both elements also decreases the martensite finish temperature (Mf): in the quaternary alloy Mf is in average 290°C while for the ternary alloys the average Mf is between 320°C and 340°C.
The alloys containing molybdenum have similar CCT diagrams, figures 1a) and 1c). These diagrams are split in two clear domains: one fully bainitic with a rather constant formation temperature, and one fully martensitic. However, for the FeCMn alloy, while there is a fully martensitic domain, there is no fully bainitic domain (figure 1b). For cooling rates between 20°C.s⁻¹ and 10°C.s⁻¹, EBSD measurements showed that the fraction of martensite increases with the cooling rate, which is coherent with the evolution of the Vickers hardness. The figure 2e) shows that martensite is present in rather small proportions for a cooling rate of 10°C/s. From cooling rates of 5°C.s⁻¹ and less, the FeCMn microstructure is a mix of allotriomorphic and Widmanstätten ferrite with bainite at 5°C.s⁻¹ or pearlite for lower cooling rates, see figure 2f). Observing the presence of allotriomorphic and Widmanstätten ferrite together is not surprising since the Widmanstätten ferritic laths nucleate on allotriomorphic ferrite grains, with a low misorientation angle between them [30, 31]. Furthermore, the formation of Widmanstätten ferrite in the present study is favored by the large size of the prior austenite grains.

In the fully martensitic domain, the formation temperature (Mₛ) is close for all the alloys, with an average value of 430°C, and the microstructures are similar, see figures 2a), 2d) and 2g). The experimental Mₛ value is comparable to what Andrews empirical equation predicts: 451°C for FeCMo, 414°C for FeCMn and 406°C for FeCMoMn. In that domain, the Vickers hardness is rather constant with cooling rate, with close values for all compositions.

The presence of Mn increases the martensitic hardenability of the steel, as reported before in literature [24], and martensitic microstructures can be seen until cooling rates of 30°C.s⁻¹ or even less for the FeCMn alloy, which is not the case for the FeCMo composition.

Combining molybdenum and manganese decreases the bainite formation temperature (Bₛ): it is around 650°C in the ternary alloys, while it is 50°C lower for FeCMoMn, with an average of 585°C, see figures 1a) to 1c). Wang *et al.* also reported that both Mn and Mo contribute to decrease Bₛ, after compiling several CCT diagrams through an artificial neural network model [32]. Combining both elements also decreases the martensite finish temperature (Mf): in the quaternary alloy Mf is in average 290°C while for the ternary alloys the average Mf is between 320°C and 340°C.
In the FeCMo and FeCMoMn systems, the bainitic microstructure becomes coarser with decreasing cooling rates according to optical observations, see figures 2b), 2c) and 2h), 2i). In these alloys, bainite is the main phase even for cooling rates as low as 0.5°C.s⁻¹ for FeCMo and 0.1°C.s⁻¹ for FeCMoMn. Nevertheless, ferrite at grain boundaries precipitates for FeCMo at cooling rates equal to 1 and 0.5°C.s⁻¹ (figure 2c). However, the formation of ferrite could not be detected by dilatometry, see figure 1a).

When comparing the FeCMoMn CCT diagram with a CCT diagram of the 20MND5 steel [7] with a similar prior austenite grain size, figures 1c) and 1d), many similarities are found. For instance, the martensite and bainite temperature formation are the same and the cooling rate at which the bainitic domain starts (10°C.s⁻¹) is also similar. However, the industrial steel diagram has a range of cooling rates for which the martensitic and bainitic domain overlap, which is not the case in the quaternary model alloy. Furthermore, ferrite precipitates from 0.3°C.s⁻¹ in the industrial diagram, while no ferrite at grain boundaries was present for FeCMoMn at the lowest cooling rates. This is due to the higher concentration in molybdenum added in the model alloy compared to the industrial 20MND5.

Despite these differences, the FeCMoMn and the 20MND5 CCT diagrams are almost the same. It means that Mo and Mn here play a great role in explaining the phase transformations in the industrial alloy, and that the other alloying elements only have a small contribution to explaining the diagram.

## 3.2. Distributions of misorientation angle

Figure 3 gives the misorientation angle distributions for several cooling rates investigated by EBSD, for FeCMo. Results for the two other compositions can be found in the supplementary material figures S1a) to S1d).
In the FeCMo and FeCMoMn systems, the bainitic microstructure becomes coarser with decreasing cooling rates according to optical observations, see figures 2b), 2c) and 2h), 2i). In these alloys, bainite is the main phase even for cooling rates as low as 0.5°C.s⁻¹ for FeCMo and 0.1°C.s⁻¹ for FeCMoMn. Nevertheless, ferrite at grain boundaries precipitates for FeCMo at cooling rates equal to 1 and 0.5°C.s⁻¹ (figure 2c). However, the formation of ferrite could not be detected by dilatometry, see figure 1a).

When comparing the FeCMoMn CCT diagram with a CCT diagram of the 20MND5 steel [7] with a similar prior austenite grain size, figures 1c) and 1d), many similarities are found. For instance, the martensite and bainite temperature formation are the same and the cooling rate at which the bainitic domain starts (10°C.s⁻¹) is also similar. However, the industrial steel diagram has a range of cooling rates for which the martensitic and bainitic domain overlap, which is not the case in the quaternary model alloy. Furthermore, ferrite precipitates from 0.3°C.s⁻¹ in the industrial diagram, while no ferrite at grain boundaries was present for FeCMoMn at the lowest cooling rates. This is due to the higher concentration in molybdenum added in the model alloy compared to the industrial 20MND5.

Despite these differences, the FeCMoMn and the 20MND5 CCT diagrams are almost the same. It means that Mo and Mn here play a great role in explaining the phase transformations in the industrial alloy, and that the other alloying elements only have a small contribution to explaining the diagram.

## 3.2. Distributions of misorientation angle

Figure 3 gives the misorientation angle distributions for several cooling rates investigated by EBSD, for FeCMo. Results for the two other compositions can be found in the supplementary material figures S1a) to S1d).
<figure><img src="image_3.png" /><figcaption>Figure 3: Distributions of misorientation angle for the FeCMo composition with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)</figcaption></figure>

In the FeCMo alloy, it is possible to split the distributions into two groups, see figures 3a) and 3b). The first group of distributions, containing the cooling rates from 300°C.s⁻¹ to 30°C.s⁻¹, has a high frequency of boundaries with a misorientation around 59°. The second group, containing the cooling rates below 10°C.s⁻¹, has boundaries with a misorientation around 53° and 59° of equal frequency. Most of the time, this group also presents boundaries with a misorientation around 8°, which is not the case for the second group. This classification is also applicable to the FeCMo CCT diagram. The first group corresponds to the martensitic microstructures, except for the microstructure cooled at 30°C.s⁻¹, while the second group corresponds to the bainitic domain. The microstructure cooled at 30°C.s⁻¹, although in the bainitic domain, has a similar misorientation angle distribution to martensitic microstructures. It is the only cooling rate for which the bainitic temperature formation is lower than for the other bainitic microstructures. This is in agreement with the distinction made by Gourgues et al. [17] between microstructures formed at high temperature versus microstructures formed at low temperatures.
<figure><img src="image_4.png" /><figcaption>Figure 3: Distributions of misorientation angle for the FeCMo composition with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)</figcaption></figure>

In the FeCMo alloy, it is possible to split the distributions into two groups, see figures 3a) and 3b). The first group of distributions, containing the cooling rates from 300°C.s⁻¹ to 30°C.s⁻¹, has a high frequency of boundaries with a misorientation around 59°. The second group, containing the cooling rates below 10°C.s⁻¹, has boundaries with a misorientation around 53° and 59° of equal frequency. Most of the time, this group also presents boundaries with a misorientation around 8°, which is not the case for the second group. This classification is also applicable to the FeCMo CCT diagram. The first group corresponds to the martensitic microstructures, except for the microstructure cooled at 30°C.s⁻¹, while the second group corresponds to the bainitic domain. The microstructure cooled at 30°C.s⁻¹, although in the bainitic domain, has a similar misorientation angle distribution to martensitic microstructures. It is the only cooling rate for which the bainitic temperature formation is lower than for the other bainitic microstructures. This is in agreement with the distinction made by Gourgues et al. [17] between microstructures formed at high temperature versus microstructures formed at low temperatures.
For FeCMn however, see supplementary material figures S1a) and S1b), it is not possible to regroup
the microstructures according to their misorientation angle distributions since these distributions
rather present a continuum. The more the cooling rate decreases, the fewer boundaries with a
misorientation around 59° there are and the more boundaries with a misorientation around 53° are
present. Furthermore, the higher the cooling rate, the lower is the misorientation frequency around
8°. The microstructure cooled at 0.5°C.s⁻¹ even tends towards a random (McKenzie) distribution [33].

Even though the Widmanstätten ferrite is known to have a Kurdjumov-Sachs (KS) relationship [27] with
the austenite phase [34], the large amount of pearlite and allotriomorphic ferrite in the sample, see
figure 2f), explains why the distribution tends toward a McKenzie distribution [26].

Although FeCMoMn and FeCMo have similar CCT diagrams, the evolutions of the misorientation angle distributions with cooling rate is different, see supplementary material figures S1c) and S1d). For microstructures from 300°C.s⁻¹ to 2°C.s⁻¹, the frequency of boundaries with a misorientation around 59° increases with decreasing cooling rate, where for FeCMo and FeCMn it is the opposite. This is due to a higher proportion of boundaries with a low misorientation and of boundaries with a misorientation between 50 and 57° for the fastest cooling rates.

Nevertheless, the distributions for microstructures with a cooling rate lower than 2°C.s⁻¹ tend to have similarities with the FeCMo bainitic distributions: the most frequent misorientation for boundaries are around 53°, 59° or 8°.

Therefore, the FeCMoMn bainitic microstructures could be sorted in a group with microstructures
from 10°C.s⁻¹ to 2°C.s⁻¹ and another one with microstructures cooled from 1°C.s⁻¹ to 0.1°C.s⁻¹. However,
this classification could hardly be explained by what is reported in the CCT diagram since all these
microstructures have the same formation temperature.

The peak observed at 30° is due to an error in indexation of the Kikuchi patterns [17].
For FeCMn however, see supplementary material figures S1a) and S1b), it is not possible to regroup
the microstructures according to their misorientation angle distributions since these distributions
rather present a continuum. The more the cooling rate decreases, the fewer boundaries with a
misorientation around 59° there are and the more boundaries with a misorientation around 53° are
present. Furthermore, the higher the cooling rate, the lower is the misorientation frequency around
8°. The microstructure cooled at 0.5°C.s⁻¹ even tends towards a random (McKenzie) distribution [33].

Even though the Widmanstätten ferrite is known to have a Kurdjumov-Sachs (KS) relationship [27] with
the austenite phase [34], the large amount of pearlite and allotriomorphic ferrite in the sample, see
figure 2f), explains why the distribution tends toward a McKenzie distribution [26].

Although FeCMoMn and FeCMo have similar CCT diagrams, the evolutions of the misorientation angle distributions with cooling rate is different, see supplementary material figures S1c) and S1d). For microstructures from 300°C.s⁻¹ to 2°C.s⁻¹, the frequency of boundaries with a misorientation around 59° increases with decreasing cooling rate, where for FeCMo and FeCMn it is the opposite. This is due to a higher proportion of boundaries with a low misorientation and of boundaries with a misorientation between 50 and 57° for the fastest cooling rates.

Nevertheless, the distributions for microstructures with a cooling rate lower than 2°C.s⁻¹ tend to have similarities with the FeCMo bainitic distributions: the most frequent misorientation for boundaries are around 53°, 59° or 8°.

Therefore, the FeCMoMn bainitic microstructures could be sorted in a group with microstructures
from 10°C.s⁻¹ to 2°C.s⁻¹ and another one with microstructures cooled from 1°C.s⁻¹ to 0.1°C.s⁻¹. However,
this classification could hardly be explained by what is reported in the CCT diagram since all these
microstructures have the same formation temperature.

The peak observed at 30° is due to an error in indexation of the Kikuchi patterns [17].
### 3.3. Hardness vs cooling rates

The evolution of the Vickers hardness with the cooling rate is given figure 4. Despite different CCT diagrams, the ternary alloys have similar hardness variations with cooling rate and close hardness values. It means that Mo and Mn have the same effect on hardness value for a given cooling rate, despite different microstructures. For example, the FeCMn mix of allotriomorphic and Widmanstätten ferrite with pearlite cooled at 2°C.s⁻¹ has the same hardness value as the FeCMo bainite cooled at 2°C.s⁻¹. For bainitic microstructures and for compositions around 1wt% of Mo or Mn, Irvine and Pickering have reported that the hardness is expected to be similar [35].

<figure><img src="image_4.png" /><figcaption>Figure 4: Microhardness for each composition depending on the cooling rate (color necessary for print)</figcaption></figure>

In the quaternary alloy, Vickers hardness variation with cooling rate is the same as for the ternary alloys, but shifted towards lower cooling rates by a factor of ten. Furthermore, the hardness values and evolution for the quaternary alloy are the same as for the 20MND5 steel [36], see figure 4. Similar
### 3.3. Hardness vs cooling rates

The evolution of the Vickers hardness with the cooling rate is given figure 4. Despite different CCT diagrams, the ternary alloys have similar hardness variations with cooling rate and close hardness values. It means that Mo and Mn have the same effect on hardness value for a given cooling rate, despite different microstructures. For example, the FeCMn mix of allotriomorphic and Widmanstätten ferrite with pearlite cooled at 2°C.s⁻¹ has the same hardness value as the FeCMo bainite cooled at 2°C.s⁻¹. For bainitic microstructures and for compositions around 1wt% of Mo or Mn, Irvine and Pickering have reported that the hardness is expected to be similar [35].

<figure><img src="image_5.png" /><figcaption>Figure 4: Microhardness for each composition depending on the cooling rate (color necessary for print)</figcaption></figure>

In the quaternary alloy, Vickers hardness variation with cooling rate is the same as for the ternary alloys, but shifted towards lower cooling rates by a factor of ten. Furthermore, the hardness values and evolution for the quaternary alloy are the same as for the 20MND5 steel [36], see figure 4. Similar
CCT diagrams and hardness values indicate that the FeCMoMn alloy can be used as a fair model for studying microstructures of 16 to 20MND5 type of steels.

For all three model alloys, the Vickers hardness evolution for the model alloys can be divided in three domains: two domains where hardness is constant and one where hardness increases with increasing cooling rate.

The first domain goes from 5°C.s⁻¹ and less for the ternary alloys and 1°C.s⁻¹ and less for the quaternary composition. Here, the hardness value is constant. The second hardness domain goes from 10°C.s⁻¹ to a 100°C.s⁻¹ for the ternary alloys, and from 2°C.s⁻¹ until 30°C.s⁻¹ for FeCMoMn. For each kind of alloy, the hardness increases with the cooling rate. In the FeCMn alloy, this can be explained by the increasing proportion of martensite until reaching a fully martensitic microstructure at 100°C.s⁻¹. For FeCMo and FeCMoMn however, the microstructure is fully bainitic. The increase in hardness can be explained here by a microstructure becoming finer with increasing cooling rate. The last hardness domain corresponds to the martensitic domain. It goes from 100°C.s⁻¹ and above for the ternary alloys and above 30°C.s⁻¹ for the quaternary one. The hardness in the martensitic domain is independent of the cooling rate.

Table 2 gives the average hardness values for different types of microstructures, relying on CCT diagrams. The average martensitic hardness is the same for all the model alloys, meaning it is independent of the type of alloying element, as already reported in [15]. However, the average bainitic hardness is higher in the quaternary alloy than in the ternary alloys, which both have a bainitic hardness of 240, despite the presence of martensite in the FeCMn composition. Since the bainitic microstructure has a lower temperature formation in the quaternary alloy, this increase in bainitic hardness is due to a microstructure refinement coming from combining the two alloying elements.
CCT diagrams and hardness values indicate that the FeCMoMn alloy can be used as a fair model for studying microstructures of 16 to 20MND5 type of steels.

For all three model alloys, the Vickers hardness evolution for the model alloys can be divided in three domains: two domains where hardness is constant and one where hardness increases with increasing cooling rate.

The first domain goes from 5°C.s⁻¹ and less for the ternary alloys and 1°C.s⁻¹ and less for the quaternary composition. Here, the hardness value is constant. The second hardness domain goes from 10°C.s⁻¹ to a 100°C.s⁻¹ for the ternary alloys, and from 2°C.s⁻¹ until 30°C.s⁻¹ for FeCMoMn. For each kind of alloy, the hardness increases with the cooling rate. In the FeCMn alloy, this can be explained by the increasing proportion of martensite until reaching a fully martensitic microstructure at 100°C.s⁻¹. For FeCMo and FeCMoMn however, the microstructure is fully bainitic. The increase in hardness can be explained here by a microstructure becoming finer with increasing cooling rate. The last hardness domain corresponds to the martensitic domain. It goes from 100°C.s⁻¹ and above for the ternary alloys and above 30°C.s⁻¹ for the quaternary one. The hardness in the martensitic domain is independent of the cooling rate.

Table 2 gives the average hardness values for different types of microstructures, relying on CCT diagrams. The average martensitic hardness is the same for all the model alloys, meaning it is independent of the type of alloying element, as already reported in [15]. However, the average bainitic hardness is higher in the quaternary alloy than in the ternary alloys, which both have a bainitic hardness of 240, despite the presence of martensite in the FeCMn composition. Since the bainitic microstructure has a lower temperature formation in the quaternary alloy, this increase in bainitic hardness is due to a microstructure refinement coming from combining the two alloying elements.
<table><caption>Table 2: Average hardness according to microstructure for the model alloys, using the CCT diagrams.</caption><thead><tr><th></th><th>FeCMn</th><th>FeCMo</th><th>FeCMoMn</th></tr></thead><tbody><tr><td colspan="4"><strong>Martensite</strong></td></tr><tr><td>Average hardness</td><td>433±7</td><td>443±4</td><td>448±3</td></tr><tr><td>Cooling rate interval</td><td>100°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td><td>100°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td><td>30°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td></tr><tr><td colspan="4"><strong>Bainite</strong></td></tr><tr><td>Average hardness</td><td>240±18*</td><td>240±3</td><td>305±3</td></tr><tr><td>Cooling rate interval</td><td>10°C.s<sup>-1</sup></td><td>2°C.s<sup>-1</sup> to 30°C.s<sup>-1</sup></td><td>0.1°C.s<sup>-1</sup> to 10°C.s<sup>-1</sup></td></tr><tr><td colspan="4"><strong>Ferrite + Pearlite</strong></td></tr><tr><td>Average hardness</td><td>179±4</td><td>-</td><td>-</td></tr><tr><td>Cooling rate interval</td><td>0.5°C.s<sup>-1</sup> to 5°C.s<sup>-1</sup></td><td>-</td><td>-</td></tr></tbody></table>

*The value given for the FeCMn bainite is the hardness of the microstructure cooled at 10°C/s, with the least possible martensite.

# 4. Discussion

## 4.1. The martensitic domain

Martensitic microstructures cooled at 300°C.s⁻¹ are chosen for each alloy to understand if Mo and Mn have an impact on this kind of microstructure. Figures 5a) and 5b) show that all three alloys have similar misorientation angle distributions. Figure 5c) gives the misorientation axis distribution in inverse pole figures for the quaternary alloy, for angles superior to 49°; see supplementary material figures S2a) and S2b) for the misorientation axis distributions of the ternary alloys.

All three compositions have similar misorientation axis/angle pairs. The main misorientation pairs are of 53-57° around [101] and of 59°-61° around [433]. This is in compliance with the misorientation angle distributions, figure 5b). These two angle/axis pairs are found for a GT/GT misorientation relationship
<table><caption>Table 2: Average hardness according to microstructure for the model alloys, using the CCT diagrams.</caption><thead><tr><th></th><th>FeCMn</th><th>FeCMo</th><th>FeCMoMn</th></tr></thead><tbody><tr><td colspan="4"><strong>Martensite</strong></td></tr><tr><td>Average hardness</td><td>433±7</td><td>443±4</td><td>448±3</td></tr><tr><td>Cooling rate interval</td><td>100°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td><td>100°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td><td>30°C.s<sup>-1</sup> to 300°C.s<sup>-1</sup></td></tr><tr><td colspan="4"><strong>Bainite</strong></td></tr><tr><td>Average hardness</td><td>240±18*</td><td>240±3</td><td>305±3</td></tr><tr><td>Cooling rate interval</td><td>10°C.s<sup>-1</sup></td><td>2°C.s<sup>-1</sup> to 30°C.s<sup>-1</sup></td><td>0.1°C.s<sup>-1</sup> to 10°C.s<sup>-1</sup></td></tr><tr><td colspan="4"><strong>Ferrite + Pearlite</strong></td></tr><tr><td>Average hardness</td><td>179±4</td><td>-</td><td>-</td></tr><tr><td>Cooling rate interval</td><td>0.5°C.s<sup>-1</sup> to 5°C.s<sup>-1</sup></td><td>-</td><td>-</td></tr></tbody></table>

*The value given for the FeCMn bainite is the hardness of the microstructure cooled at 10°C/s, with the least possible martensite.

# 4. Discussion

## 4.1. The martensitic domain

Martensitic microstructures cooled at 300°C.s⁻¹ are chosen for each alloy to understand if Mo and Mn have an impact on this kind of microstructure. Figures 5a) and 5b) show that all three alloys have similar misorientation angle distributions. Figure 5c) gives the misorientation axis distribution in inverse pole figures for the quaternary alloy, for angles superior to 49°; see supplementary material figures S2a) and S2b) for the misorientation axis distributions of the ternary alloys.

All three compositions have similar misorientation axis/angle pairs. The main misorientation pairs are of 53-57° around [101] and of 59°-61° around [433]. This is in compliance with the misorientation angle distributions, figure 5b). These two angle/axis pairs are found for a GT/GT misorientation relationship
(MR), see table 3, and in KS/NW MR [29]. However, they do not exist in KS/KS or NW/NW, see table S1 in supplementary material [17, 28].

A study in a prior austenite grain (PAG) is carried out in each of the three alloys and gives the same results for all three compositions. Figures 6b) to 6d) is the result of the investigations made in a FeCMoMn PAG selected from the EBSD map, figure 6a). Three texture components could be identified on the pole figure, figure 6b), and each of these texture components present between them a misorientation superior to 49°, as illustrated in figure 6c). Figure 6d) shows the misorientation profile associated with the white arrow on figure 6c). Two types of boundaries can be identified: boundaries with a misorientation of more than 50° and boundaries with a misorientation around 10°. According to Morito's definition [37], the first type of boundaries are block boundaries while the low angle boundaries are sub-block boundaries. Inside the blocks and sub-blocks, the misorientation is quite low, inferior to 4°, figure 6c). In the martensitic microstructure, according to figures 6d) and 5b), the most probable block boundary misorientation is around 59°.

<table><caption>Table 3: High angle/axis misorientation pairs measured for the martensite of the three model alloys compared to the closest misorientation pairs from GT/GT and KS/NW MR. The pair marked with (*) is less present than the others. The angle value given in parenthesis is the deviation from the specified axis.</caption><thead><tr><th>Martensite</th><th colspan="3">Model Alloys</th><th colspan="2">GT/GT</th><th colspan="2">KS/NW</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2*</td><td>55° ±2</td><td>59° ±2</td><td>54.3°</td><td>60.2°</td><td>54.7°</td><td>60.2°</td></tr><tr><th>Axis<br>(Offset<br>angle)</th><td>443*</td><td>101</td><td>433</td><td>101<br>(0°)</td><td>433<br>(3.25°)</td><td>101<br>(3.08°)</td><td>433<br>(2.80°)</td></tr></tbody></table>
(MR), see table 3, and in KS/NW MR [29]. However, they do not exist in KS/KS or NW/NW, see table S1 in supplementary material [17, 28].

A study in a prior austenite grain (PAG) is carried out in each of the three alloys and gives the same results for all three compositions. Figures 6b) to 6d) is the result of the investigations made in a FeCMoMn PAG selected from the EBSD map, figure 6a). Three texture components could be identified on the pole figure, figure 6b), and each of these texture components present between them a misorientation superior to 49°, as illustrated in figure 6c). Figure 6d) shows the misorientation profile associated with the white arrow on figure 6c). Two types of boundaries can be identified: boundaries with a misorientation of more than 50° and boundaries with a misorientation around 10°. According to Morito's definition [37], the first type of boundaries are block boundaries while the low angle boundaries are sub-block boundaries. Inside the blocks and sub-blocks, the misorientation is quite low, inferior to 4°, figure 6c). In the martensitic microstructure, according to figures 6d) and 5b), the most probable block boundary misorientation is around 59°.

<table><caption>Table 3: High angle/axis misorientation pairs measured for the martensite of the three model alloys compared to the closest misorientation pairs from GT/GT and KS/NW MR. The pair marked with (*) is less present than the others. The angle value given in parenthesis is the deviation from the specified axis.</caption><thead><tr><th>Martensite</th><th colspan="3">Model Alloys</th><th colspan="2">GT/GT</th><th colspan="2">KS/NW</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2*</td><td>55° ±2</td><td>59° ±2</td><td>54.3°</td><td>60.2°</td><td>54.7°</td><td>60.2°</td></tr><tr><th>Axis<br>(Offset<br>angle)</th><td>443*</td><td>101</td><td>433</td><td>101<br>(0°)</td><td>433<br>(3.25°)</td><td>101<br>(3.08°)</td><td>433<br>(2.80°)</td></tr></tbody></table>
<figure><img src="image_5.png" /></figure>

Figure 5: EBSD on the martensitic domain: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c) is the misorientation axis distribution for the FeCMoMn composition cooled at 300°C/s. (color necessary for print)
<figure><img src="image_6.png" /></figure>

Figure 5: EBSD on the martensitic domain: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c) is the misorientation axis distribution for the FeCMoMn composition cooled at 300°C/s. (color necessary for print)
<figure><img src="image_6.png" /><figcaption>Figure 6: Study in a prior austenite grain for FeCMoMn 300°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Figure 6: Study in a prior austenite grain for FeCMoMn 300°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
All three compositions present the same martensitic crystallographic features, as well as similar micrographs and hardness, meaning that Mo nor Mn have an impact on the martensitic microstructure.

## 4.2. The bainitic domain

To study the effect of Mo and Mn on the bainitic domain, the microstructures cooled at 5°C.s⁻¹ for FeCMo and cooled at 1°C.s⁻¹ for FeCMoMn were chosen because they correspond to cooling rates where the bainite transformation temperature is independent of the cooling rate. For FeCMn, the microstructure with the least martensite and without allotriomorphic and Widmanstätten ferrite was selected: the microstructure cooled at 10°C.s⁻¹ is the closest to a fully bainitic microstructure.

In that domain, the misorientation angle distributions presented figures 7a) and 7b) differ from one composition to another. The Mn-containing alloys have similar distributions with a higher proportion of boundaries with a misorientation around 59° than what can be found in the FeCMo bainite. In the FeCMo microstructure cooled at 5°C.s⁻¹, there are as many boundaries with a misorientation around 53° as there are around 59°. Furthermore, there is a higher tendency to form boundaries with a misorientation peak below 10° than for the Mn-bearing alloys.

The FeCMo, FeCMn and FeCMoMn misorientation axis distribution for high angles is presented in figure 7c), d) and e) respectively.

When analyzing the misorientation angle/axis pairs, the same differences as for the misorientation angle distributions can be noted. FeCMo has two main misorientation angle/axis pairs: 49-53° around [221] and of 59-61° around [111], which is a Σ3 boundary. These pairs can come from a KS/KS misorientation relationship, see table 4a), but not from a KS/NW, a NW/NW or a GT/GT MR [17, 28], (table S1 in supplementary material). For the FeCMn alloy, despite being mainly bainitic, similar features to martensite exist with the same misorientation of 59-61° around [433]. As previously stated, this angle/axis pair can come from a KS/NW or GT/GT misorientation relationship, but not from a KS/KS nor NW/NW misorientation relationship [17]. This means that the orientation relationship of the bainitic microstructure with parent austenite may be changing depending on the addition of Mo or
All three compositions present the same martensitic crystallographic features, as well as similar micrographs and hardness, meaning that Mo nor Mn have an impact on the martensitic microstructure.

## 4.2. The bainitic domain

To study the effect of Mo and Mn on the bainitic domain, the microstructures cooled at 5°C.s⁻¹ for FeCMo and cooled at 1°C.s⁻¹ for FeCMoMn were chosen because they correspond to cooling rates where the bainite transformation temperature is independent of the cooling rate. For FeCMn, the microstructure with the least martensite and without allotriomorphic and Widmanstätten ferrite was selected: the microstructure cooled at 10°C.s⁻¹ is the closest to a fully bainitic microstructure.

In that domain, the misorientation angle distributions presented figures 7a) and 7b) differ from one composition to another. The Mn-containing alloys have similar distributions with a higher proportion of boundaries with a misorientation around 59° than what can be found in the FeCMo bainite. In the FeCMo microstructure cooled at 5°C.s⁻¹, there are as many boundaries with a misorientation around 53° as there are around 59°. Furthermore, there is a higher tendency to form boundaries with a misorientation peak below 10° than for the Mn-bearing alloys.

The FeCMo, FeCMn and FeCMoMn misorientation axis distribution for high angles is presented in figure 7c), d) and e) respectively.

When analyzing the misorientation angle/axis pairs, the same differences as for the misorientation angle distributions can be noted. FeCMo has two main misorientation angle/axis pairs: 49-53° around [221] and of 59-61° around [111], which is a Σ3 boundary. These pairs can come from a KS/KS misorientation relationship, see table 4a), but not from a KS/NW, a NW/NW or a GT/GT MR [17, 28], (table S1 in supplementary material). For the FeCMn alloy, despite being mainly bainitic, similar features to martensite exist with the same misorientation of 59-61° around [433]. As previously stated, this angle/axis pair can come from a KS/NW or GT/GT misorientation relationship, but not from a KS/KS nor NW/NW misorientation relationship [17]. This means that the orientation relationship of the bainitic microstructure with parent austenite may be changing depending on the addition of Mo or
<figure><img src="image_7.png" /><figcaption>Figure 7: EBSD on the bainitic domain: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c), (d) and (e) are the corresponding misorientation axis distribution. (color necessary for print)</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>Figure 7: EBSD on the bainitic domain: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c), (d) and (e) are the corresponding misorientation axis distribution. (color necessary for print)</figcaption></figure>
Mn. Furthermore, since the 59-61° [433] pair is also present in the FeCMoMn bainitic microstructure, but absent from the FeCMo bainitic one, it suggests that the presence of manganese here in the bainitic microstructure favors this specific angle/axis pair. Here, the FeCMn and FeCMoMn microstructures share many common misorientation angle/axis pairs which can be found in GT/GT MR, see table 4b). The pair 49-53° around [443] in FeCMn and FeCMoMn can be considered close to the GT/GT pair 50.8° [332] because [443] and [332] are close.

a)

<table><thead><tr><th>Bainite</th><th colspan="3">FeCMo</th><th colspan="3">KS/KS</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2</td><td>55° ±2</td><td>59° ±2</td><td>49.5°</td><td>51.7°</td><td>60°</td></tr><tr><th>Axis<br>(Offset angle)</th><td>221</td><td>221</td><td>111</td><td>111<br>(0°)</td><td>221<br>(1.80°)</td><td>111<br>(0°)</td></tr></tbody></table>

b)

<table><caption>Table 4: a) High angle/axis misorientation pairs measured for the bainite of the FeCMo alloy compared to the closest misorientation pairs from KS orientation relationship (OR), and b) high angle/axis misorientation pairs measured for the FeCMo alloy compared to the closest misorientation pairs from KS OR. The pairs marked with (*) are less present than the others.</caption><thead><tr><th>Bainite</th><th colspan="3">FeCMn and FeCMoMn</th><th colspan="4">GT/GT</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2*</td><td>55° ±2*</td><td>59° ±2</td><td>50.8°</td><td>54.3°</td><td>60°</td><td>60.2°</td></tr><tr><th>Axis<br>(Offset angle)</th><td>221* and 443*</td><td>101*</td><td>433</td><td>332<br>(0.94°)</td><td>101<br>(0°)</td><td>101<br>(0°)</td><td>433<br>(3.25°)</td></tr></tbody></table>

Figures 8a) and b) presents the misorientation of boundaries in a PAG taken from the FeCMoMn bainitic microstructure. More details can be found in supplementary material, figure S3. A similar investigation was carried out in bainitic PAGs in FeCMn and FeCMo and gave similar results.

The bainitic microstructure has here many crystallographic features in common with martensite. The microstructure is organized as well in blocks, separated by boundaries with a misorientation of 50° or
Mn. Furthermore, since the 59-61° [433] pair is also present in the FeCMoMn bainitic microstructure, but absent from the FeCMo bainitic one, it suggests that the presence of manganese here in the bainitic microstructure favors this specific angle/axis pair. Here, the FeCMn and FeCMoMn microstructures share many common misorientation angle/axis pairs which can be found in GT/GT MR, see table 4b). The pair 49-53° around [443] in FeCMn and FeCMoMn can be considered close to the GT/GT pair 50.8° [332] because [443] and [332] are close.

a)

<table><thead><tr><th>Bainite</th><th colspan="3">FeCMo</th><th colspan="3">KS/KS</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2</td><td>55° ±2</td><td>59° ±2</td><td>49.5°</td><td>51.7°</td><td>60°</td></tr><tr><th>Axis<br>(Offset angle)</th><td>221</td><td>221</td><td>111</td><td>111<br>(0°)</td><td>221<br>(1.80°)</td><td>111<br>(0°)</td></tr></tbody></table>

b)

<table><caption>Table 4: a) High angle/axis misorientation pairs measured for the bainite of the FeCMo alloy compared to the closest misorientation pairs from KS orientation relationship (OR), and b) high angle/axis misorientation pairs measured for the FeCMo alloy compared to the closest misorientation pairs from KS OR. The pairs marked with (*) are less present than the others.</caption><thead><tr><th>Bainite</th><th colspan="3">FeCMn and FeCMoMn</th><th colspan="4">GT/GT</th></tr></thead><tbody><tr><th>Angle</th><td>51° ±2*</td><td>55° ±2*</td><td>59° ±2</td><td>50.8°</td><td>54.3°</td><td>60°</td><td>60.2°</td></tr><tr><th>Axis<br>(Offset angle)</th><td>221* and 443*</td><td>101*</td><td>433</td><td>332<br>(0.94°)</td><td>101<br>(0°)</td><td>101<br>(0°)</td><td>433<br>(3.25°)</td></tr></tbody></table>

Figures 8a) and b) presents the misorientation of boundaries in a PAG taken from the FeCMoMn bainitic microstructure. More details can be found in supplementary material, figure S3. A similar investigation was carried out in bainitic PAGs in FeCMn and FeCMo and gave similar results.

The bainitic microstructure has here many crystallographic features in common with martensite. The microstructure is organized as well in blocks, separated by boundaries with a misorientation of 50° or
<figure><img src="image_8.png" /><figcaption>Figure 8: Study in a prior austenite grain for FeCMoMn 1°C/s: a) EBSD map of the misorientation angles in the PAG with a profile in white, and the misorientation profile in PAG (b). (color necessary for print)</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Figure 8: Study in a prior austenite grain for FeCMoMn 1°C/s: a) EBSD map of the misorientation angles in the PAG with a profile in white, and the misorientation profile in PAG (b). (color necessary for print)</figcaption></figure>
more, and sub-blocks, with a misorientation from 10° to 20°, see figures 8a) and 8b). Since the FeCMo
alloy has a higher frequency of misorientation below 20° than the Mn-bearing alloys, figure 7a), it
means that sub-block boundaries are more frequent in the FeCMo alloy than block boundaries. This
can have an impact on the resulting mechanical properties since high angle boundaries have the ability
to arrest or deviate cracks during cleavage fracture [17, 39] while low angle boundaries only show a
low angle change in the fissure propagation path inducing a lower dissipation of energy during fracture
[39]. As for martensite, the pole figure shows, see supplementary figure S3b), the presence of three
texture components, with a misorientation of 49° at least between each other.

In the bainitic domain, the alloying elements have an impact on the bainitic microstructure, unlike for martensite, despite common crystallographic features at the PAG level.

4.3. Lowest cooling rates domain

In this domain, the microstructures cooled at 0.5°C.s⁻¹ for FeCMo and FeCMn, and the one cooled at 0.1°C.s⁻¹ for FeCMoMn were selected for comparison. The molybdenum-containing alloys have misorientation angle distributions comparable and close to what is observed for the FeCMo cooled at 5°C.s⁻¹, see figures 9a), 9b) and 7a), 7b). For FeCMn, the presence of an important proportion of allotriomorphic ferrite and pearlite induces a larger proportion of boundaries with a misorientation between 20 and 49°. Furthermore, the misorientation of 59-61° around [433] has disappeared in the FeCMn microstructure, see supplementary figure S4b) as well as in the FeCMoMn bainitic microstructure cooled at 0.1°C.s⁻¹, see figure 9c). Instead, it has been replaced by Σ3 boundaries in the Mn-bearing alloys. In the case of FeCMn, the presence of Σ3 boundaries can come from a KS/KS MR, since Widmanstätten ferrite has a KS OR.

Despite different microstructures, FeCMn and FeCMoMn share this common twin boundary and Mo
seems to have little impact on the selection of misorientation angle/axis pairs in the FeCMoMn alloy,
compared to Mn. In the FeCMo microstructure, supplementary figure S4a), the boundaries present at
more, and sub-blocks, with a misorientation from 10° to 20°, see figures 8a) and 8b). Since the FeCMo
alloy has a higher frequency of misorientation below 20° than the Mn-bearing alloys, figure 7a), it
means that sub-block boundaries are more frequent in the FeCMo alloy than block boundaries. This
can have an impact on the resulting mechanical properties since high angle boundaries have the ability
to arrest or deviate cracks during cleavage fracture [17, 39] while low angle boundaries only show a
low angle change in the fissure propagation path inducing a lower dissipation of energy during fracture
[39]. As for martensite, the pole figure shows, see supplementary figure S3b), the presence of three
texture components, with a misorientation of 49° at least between each other.

In the bainitic domain, the alloying elements have an impact on the bainitic microstructure, unlike for martensite, despite common crystallographic features at the PAG level.

4.3. Lowest cooling rates domain

In this domain, the microstructures cooled at 0.5°C.s⁻¹ for FeCMo and FeCMn, and the one cooled at 0.1°C.s⁻¹ for FeCMoMn were selected for comparison. The molybdenum-containing alloys have misorientation angle distributions comparable and close to what is observed for the FeCMo cooled at 5°C.s⁻¹, see figures 9a), 9b) and 7a), 7b). For FeCMn, the presence of an important proportion of allotriomorphic ferrite and pearlite induces a larger proportion of boundaries with a misorientation between 20 and 49°. Furthermore, the misorientation of 59-61° around [433] has disappeared in the FeCMn microstructure, see supplementary figure S4b) as well as in the FeCMoMn bainitic microstructure cooled at 0.1°C.s⁻¹, see figure 9c). Instead, it has been replaced by Σ3 boundaries in the Mn-bearing alloys. In the case of FeCMn, the presence of Σ3 boundaries can come from a KS/KS MR, since Widmanstätten ferrite has a KS OR.

Despite different microstructures, FeCMn and FeCMoMn share this common twin boundary and Mo
seems to have little impact on the selection of misorientation angle/axis pairs in the FeCMoMn alloy,
compared to Mn. In the FeCMo microstructure, supplementary figure S4a), the boundaries present at
<figure><img src="image_9.png" /><figcaption>Figure 9: EBSD on the slowest cooling speeds: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c) is the misorientation axis distribution for FeCMoMn composition cooled at 0.1°C/s (c). (color necessary for print)</figcaption></figure>

5°C.s⁻¹, with a misorientation of 49-53° around [221] and Σ3 boundaries, tend to disappear to the benefit of boundaries with a misorientation of 53-57° around [331], which can from a KS/NW and a GT/GT MR [17], see table S1 in supplementary material.

Investigation in a FeCMoMn PAG and in a FeCMo bainitic PAG gave similar results. The result for the FeCMoMn PAG can be found in supplementary material figures S5a) to S5d). For FeCMn, however, due to the microstructure, it was not possible to study the bainitic misorientation in a PAG.
<figure><img src="image_10.png" /><figcaption>Figure 9: EBSD on the slowest cooling speeds: distribution of the misorientation angles with a minimum at 5° (a) and a minimum at 45° (b). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (c) is the misorientation axis distribution for FeCMoMn composition cooled at 0.1°C/s (c). (color necessary for print)</figcaption></figure>

5°C.s⁻¹, with a misorientation of 49-53° around [221] and Σ3 boundaries, tend to disappear to the benefit of boundaries with a misorientation of 53-57° around [331], which can from a KS/NW and a GT/GT MR [17], see table S1 in supplementary material.

Investigation in a FeCMoMn PAG and in a FeCMo bainitic PAG gave similar results. The result for the FeCMoMn PAG can be found in supplementary material figures S5a) to S5d). For FeCMn, however, due to the microstructure, it was not possible to study the bainitic misorientation in a PAG.
The bainitic microstructure in the Mo-bearing alloys have three texture components, supplementary figure S5b), but some PAG were found with only one or two texture components. As the cooling rate decreases, a narrower range of variants is present per PAG. The supplementary figure S5d) shows that, as for the martensite and the bainite cooled at 1°C.s⁻¹, block boundaries have a misorientation superior to 50° while sub-blocks have a misorientation of 10° or less. In that sense, even if the misorientation angle/axis pairs change from a cooling rate to another, it appears that, on the PAG level, the microstructure evolves continuously from martensite to slowly cooled bainite, conserving many crystallographic similarities. As can be seen by comparing supplementary figure S5c) with figure 8a) and 6c), the microstructure becomes coarser with decreasing cooling rate, with fewer high angle boundaries, meaning blocks are getting larger. In these large blocks, the density of sub-blocks boundaries is not increased with decreasing cooling rate, as shows figure 8a) and supplementary figure S5c).

### 4.4. Model and industrial alloys

The density of the misorientation angle for 16MND5 bainite and martensite [36] compared with the model alloys are plotted figure 10. Density means the frequency is divided by the bin size. It is necessary here to compare the frequency distributions because the bin values are different for the 16MND5 and the model alloys distributions. The misorientation angle distributions of the model alloy and of the 16MND5 bainite, figures 10a) and 10b), are similar. The 16MND5 bainite, figure 10a), presents a high proportion of low angle boundaries around 8°, like the FeCMo bainite. Regarding the high angle boundaries, that is to say block boundaries, there is a nearly equal proportion of boundaries with a misorientation of 53° and another of 59°, figure 10b). The 16MND5 bainite presents a higher proportion of boundaries around 53° than for the Mn-bearing model alloys.

As for the martensitic microstructures, the distributions of the industrial steel and of the model alloys are quite similar, figures 10c) and 10d). Boundaries with a misorientation around 60° are the most
The bainitic microstructure in the Mo-bearing alloys have three texture components, supplementary figure S5b), but some PAG were found with only one or two texture components. As the cooling rate decreases, a narrower range of variants is present per PAG. The supplementary figure S5d) shows that, as for the martensite and the bainite cooled at 1°C.s⁻¹, block boundaries have a misorientation superior to 50° while sub-blocks have a misorientation of 10° or less. In that sense, even if the misorientation angle/axis pairs change from a cooling rate to another, it appears that, on the PAG level, the microstructure evolves continuously from martensite to slowly cooled bainite, conserving many crystallographic similarities. As can be seen by comparing supplementary figure S5c) with figure 8a) and 6c), the microstructure becomes coarser with decreasing cooling rate, with fewer high angle boundaries, meaning blocks are getting larger. In these large blocks, the density of sub-blocks boundaries is not increased with decreasing cooling rate, as shows figure 8a) and supplementary figure S5c).

### 4.4. Model and industrial alloys

The density of the misorientation angle for 16MND5 bainite and martensite [36] compared with the model alloys are plotted figure 10. Density means the frequency is divided by the bin size. It is necessary here to compare the frequency distributions because the bin values are different for the 16MND5 and the model alloys distributions. The misorientation angle distributions of the model alloy and of the 16MND5 bainite, figures 10a) and 10b), are similar. The 16MND5 bainite, figure 10a), presents a high proportion of low angle boundaries around 8°, like the FeCMo bainite. Regarding the high angle boundaries, that is to say block boundaries, there is a nearly equal proportion of boundaries with a misorientation of 53° and another of 59°, figure 10b). The 16MND5 bainite presents a higher proportion of boundaries around 53° than for the Mn-bearing model alloys.

As for the martensitic microstructures, the distributions of the industrial steel and of the model alloys are quite similar, figures 10c) and 10d). Boundaries with a misorientation around 60° are the most
frequent for the 16MND5 martensite, as it is in the model alloys. The boundaries around 53°, 55° and 60° are present in similar proportion in the 16MND5 martensite as in the model alloys (figure 10d).

Taking into consideration the misorientation angle distributions for both bainite and martensite, the ternary and quaternary alloys investigated in the present study are good models for the 16 to 20MND5 steels.

<figure><img src="image_10.png" /></figure>

Figure 10: Comparison with industrial 16MND5: Density of the misorientation angles with a minimum at 25° and 45° for bainitic microstructures (a and b) and for martensitic microstructures (c and d). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)
frequent for the 16MND5 martensite, as it is in the model alloys. The boundaries around 53°, 55° and 60° are present in similar proportion in the 16MND5 martensite as in the model alloys (figure 10d).

Taking into consideration the misorientation angle distributions for both bainite and martensite, the ternary and quaternary alloys investigated in the present study are good models for the 16 to 20MND5 steels.

<figure><img src="image_11.png" /></figure>

Figure 10: Comparison with industrial 16MND5: Density of the misorientation angles with a minimum at 25° and 45° for bainitic microstructures (a and b) and for martensitic microstructures (c and d). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)
## 5. Conclusions

The microstructures of tree model alloys were studied: two ternary FeCMo and FeCMn and one quaternary FeCMoMn. The influence of the cooling rate and of the alloying elements were investigated thanks to CCT diagrams and EBSD. This led to the following conclusions:

1. The presence of Mo and Mn greatly explains the phase transformations observed in the 16 to 20MND5 alloys, since the FeCMoMn and the industrial 20MND5 CCT diagrams are almost the same. Furthermore, the model quaternary and industrial 20MND5 alloys have similar hardness values and hardness evolutions with cooling rate.

2. The average bainitic hardness in higher when combining Mo and Mn. The martensitic hardness is independent of cooling rate and of the type of alloying element.

3. While the FeCMo microstructures can be sorted out between martensitic and bainitic ones based on their misorientation angle distributions, the FeCMn microstructures shows a rather continuous evolution. Finally, for the FeCMoMn alloy, the evolution of the frequency of boundaries with a misorientation of 59° increases with decreasing cooling rate, from $300°C.s^{-1}$ to $2°C.s^{-1}$, which is the opposite for the ternary alloys.

4. In the martensitic microstructures, no effect of Mo or Mn was detected on misorientation angle nor misorientation axis distributions.

5. In the bainitic microstructures however, the presence of Mn seems to favor the presence of block boundaries with a misorientation of about 59° around [433].

6. On the PAG level, a continuum of microstructure can be observed from martensite to slowly cooled bainitic microstructure: the amount of block boundaries with a misorientation of 49° or more decreases with decreasing cooling rate and the presence of sub-block boundaries, with a misorientation between 4 and 20°, remains present for all cooling rates investigated.

7. The microstructure becomes coarser with decreasing cooling rate, with fewer different crystallographic orientations in one PAG: where in the highest cooling rate three texture components
5. Conclusions

The microstructures of tree model alloys were studied: two ternary FeCMo and FeCMn and one quaternary FeCMoMn. The influence of the cooling rate and of the alloying elements were investigated thanks to CCT diagrams and EBSD. This led to the following conclusions:

1. The presence of Mo and Mn greatly explains the phase transformations observed in the 16 to 20MND5 alloys, since the FeCMoMn and the industrial 20MND5 CCT diagrams are almost the same. Furthermore, the model quaternary and industrial 20MND5 alloys have similar hardness values and hardness evolutions with cooling rate.

2. The average bainitic hardness in higher when combining Mo and Mn. The martensitic hardness is independent of cooling rate and of the type of alloying element.

3. While the FeCMo microstructures can be sorted out between martensitic and bainitic ones based on their misorientation angle distributions, the FeCMn microstructures shows a rather continuous evolution. Finally, for the FeCMoMn alloy, the evolution of the frequency of boundaries with a misorientation of 59° increases with decreasing cooling rate, from 300°C.s⁻¹ to 2°C.s⁻¹, which is the opposite for the ternary alloys.

4. In the martensitic microstructures, no effect of Mo or Mn was detected on misorientation angle nor misorientation axis distributions.

5. In the bainitic microstructures however, the presence of Mn seems to favor the presence of block boundaries with a misorientation of about 59° around [433].

6. On the PAG level, a continuum of microstructure can be observed from martensite to slowly cooled bainitic microstructure: the amount of block boundaries with a misorientation of 49° or more decreases with decreasing cooling rate and the presence of sub-block boundaries, with a misorientation between 4 and 20°, remains present for all cooling rates investigated.

7. The microstructure becomes coarser with decreasing cooling rate, with fewer different crystallographic orientations in one PAG: where in the highest cooling rate three texture components
are present, some PAG grains in the slowest cooled microstructure only one or two texture
components per PAG can be found.
are present, some PAG grains in the slowest cooled microstructure only one or two texture
components per PAG can be found.
Supplementary material
Supplementary material
<table><thead><tr><th colspan="16">KS/KS [17]</th></tr><tr><th>Angle (°)</th><th>10.5</th><th>10.5</th><th>14.9</th><th>20.6</th><th>20.6</th><th>21.1</th><th>47.1</th><th>49.5</th><th>49.5</th><th>50.5</th><th>50.5</th><th>51.7</th><th>57.2</th><th>57.2</th><th>60</th><th>60</th></tr><tr><th>Low index axis</th><td>101</td><td>111</td><td>310</td><td>310</td><td>221</td><td>210</td><td>221</td><td>111</td><td>101</td><td>431</td><td>322</td><td>221</td><td>432</td><td>331</td><td>111</td><td>101</td></tr><tr><th>Offset from axis (°)</th><td>0</td><td>0</td><td>4.41</td><td>1.21</td><td>1.80</td><td>2.36</td><td>4.22</td><td>0</td><td>0</td><td>1.90</td><td>1.47</td><td>1.80</td><td>1.90</td><td>4.60</td><td>0</td><td>0</td></tr><tr><th colspan="16">NW/NW [17]</th></tr><tr><th>Angle (°)</th><td>13.8</td><td>19.5</td><td>50.0</td><td>53.7</td><td>60</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><th>Low index axis</th><td>101</td><td>100</td><td>443</td><td>331</td><td>101</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><th>Offset from axis (°)</th><td>3.45</td><td>3.45</td><td>0.13</td><td>0.76</td><td>0</td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td><td></td></tr><tr><th colspan="16">KS/NW [17]</th></tr><tr><th>Angle (°)</th><td>5.25</td><td>11.8</td><td>17.2</td><td>20.2</td><td>48.6</td><td>49.7</td><td>50.9</td><td>52.1</td><td>54.7</td><td>55.7</td><td>58.8</td><td>60.2</td><td></td><td></td><td></td><td></td></tr><tr><th>Low index axis</th><td>101</td><td>421</td><td>321</td><td>411</td><td>331</td><td>433</td><td>432</td><td>331</td><td>101</td><td>431</td><td>331</td><td>433</td><td></td><td></td><td></td><td></td></tr><tr><th>Offset from axis (°)</th><td>0</td><td>3.00</td><td>1.48</td><td>5.14</td><td>1.60</td><td>1.89</td><td>3.61</td><td>4.89</td><td>0</td><td>3.08</td><td>1.44</td><td>2.80</td><td></td><td></td><td></td><td></td></tr><tr><th colspan="16">GT/GT</th></tr><tr><th>Angle (°)</th><td>5.72</td><td>11.3</td><td>14.1</td><td>17.3</td><td>19.8</td><td>19.9</td><td>49.6</td><td>50.0</td><td>50.2</td><td>50.8</td><td>51.9</td><td>54.3</td><td>55.6</td><td>57.5</td><td>60</td><td>60.2</td></tr><tr><th>Low index axis</th><td>101</td><td>331</td><td>320</td><td>331</td><td>410</td><td>401</td><td>111</td><td>331</td><td>433</td><td>332</td><td>331</td><td>101</td><td>331</td><td>441</td><td>101</td><td>433</td></tr><tr><th>Offset from axis (°)</th><td>0</td><td>1.61</td><td>4.21</td><td>1.76</td><td>4.39</td><td>0.43</td><td>3.31</td><td>2.60</td><td>3.46</td><td>0.94</td><td>3.88</td><td>0</td><td>2.24</td><td>2.87</td><td>0</td><td>3.25</td></tr></thead></table>

Table S1: Theoretical angle/axis pairs for KS/KS, NW/NW, KS/NW [17] and GT/GT misorientation relationships. Some rotation axes having irrational indexes, they were approximated by low index crystallographic axes. The offset angle between the actual axes and the low-index axes is specified.
<table><caption>Table S1: Theoretical angle/axis pairs for KS/KS, NW/NW, KS/NW [17] and GT/GT misorientation relationships. Some rotation axes having irrational indexes, they were approximated by low index crystallographic axes. The offset angle between the actual axes and the low-index axes is specified.</caption><thead><tr><th colspan="15">KS/KS [17]</th></tr><tr><th>Angle (°)</th><th>10.5</th><th>10.5</th><th>14.9</th><th>20.6</th><th>20.6</th><th>21.1</th><th>47.1</th><th>49.5</th><th>49.5</th><th>50.5</th><th>50.5</th><th>51.7</th><th>57.2</th><th>57.2</th><th>60</th><th>60</th></tr></thead><tbody><tr><th>Low index axis</th><td>101</td><td>111</td><td>310</td><td>310</td><td>221</td><td>210</td><td>221</td><td>111</td><td>101</td><td>431</td><td>322</td><td>221</td><td>432</td><td>331</td><td>111</td><td>101</td></tr><tr><th>Offset from axis (°)</th><td>0</td><td>0</td><td>4.41</td><td>1.21</td><td>1.80</td><td>2.36</td><td>4.22</td><td>0</td><td>0</td><td>1.90</td><td>1.47</td><td>1.80</td><td>1.90</td><td>4.60</td><td>0</td><td>0</td></tr></tbody><thead><tr><th colspan="15">NW/NW [17]</th></tr><tr><th>Angle (°)</th><th>13.8</th><th>19.5</th><th>50.0</th><th>53.7</th><th>60</th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th></tr></thead><tbody><tr><th>Low index axis</th><td>101</td><td>100</td><td>443</td><td>331</td><td>101</td><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th></tr><tr><th>Offset from axis (°)</th><td>3.45</td><td>3.45</td><td>0.13</td><td>0.76</td><td>0</td><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th><th></th></tr></tbody><thead><tr><th colspan="15">KS/NW [17]</th></tr><tr><th>Angle (°)</th><th>5.25</th><th>11.8</th><th>17.2</th><th>20.2</th><th>48.6</th><th>49.7</th><th>50.9</th><th>52.1</th><th>54.7</th><th>55.7</th><th>58.8</th><th>60.2</th><th></th><th></th><th></th><th></th></tr></thead><tbody><tr><th>Low index axis</th><td>101</td><td>421</td><td>321</td><td>411</td><td>331</td><td>433</td><td>432</td><td>331</td><td>101</td><td>431</td><td>331</td><td>433</td><th></th><th></th><th></th><th></th></tr><tr><th>Offset from axis (°)</th><td>0</td><td>3.00</td><td>1.48</td><td>5.14</td><td>1.60</td><td>1.89</td><td>3.61</td><td>4.89</td><td>0</td><td>3.08</td><td>1.44</td><td>2.80</td><th></th><th></th><th></th><th></th></tr></tbody><thead><tr><th colspan="15">GT/GT</th></tr><tr><th>Angle (°)</th><th>5.72</th><th>11.3</th><th>14.1</th><th>17.3</th><th>19.8</th><th>19.9</th><th>49.6</th><th>50.0</th><th>50.2</th><th>50.8</th><th>51.9</th><th>54.3</th><th>55.6</th><th>57.5</th><th>60</th><th>60.2</th></tr></thead><tbody><tr><th>Low index axis</th><td>101</td><td>331</td><td>320</td><td>331</td><td>410</td><td>401</td><td>111</td><td>331</td><td>433</td><td>332</td><td>331</td><td>101</td><td>331</td><td>441</td><td>101</td><td>433</td></tr><tr><th>Offset from axis (°)</th><td>0</td><td>1.61</td><td>4.21</td><td>1.76</td><td>4.39</td><td>0.43</td><td>3.31</td><td>2.60</td><td>3.46</td><td>0.94</td><td>3.88</td><td>0</td><td>2.24</td><td>2.87</td><td>0</td><td>3.25</td></tr></tbody></table>
<figure><img src="image_11.png" /><figcaption>Figure S1: a) and c) Distributions of misorientation angle for FeCMn and FeCMoMn compositions with a minimum at 5° and a minimum at 45° (b and d). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>Figure S1: a) and c) Distributions of misorientation angle for FeCMn and FeCMoMn compositions with a minimum at 5° and a minimum at 45° (b and d). Misorientation angles with a value inferior to the minimum are disregarded for the frequency calculation. (color necessary for print)</figcaption></figure>
<table><thead><tr><th></th><th></th><th></th><th></th><th></th></tr></thead><tbody><tr><td>FeCMo<br>300°C/s</td><td>a)</td><td>49-51° (14054pts)<br>55-57° (43482pts)</td><td>51-53° (27395pts)<br>57-59° (79144pts)</td><td>53-55° (33121pts)<br>59-61° (149432pts)</td></tr><tr><td>FeCMn<br>300°C/s</td><td>b)</td><td>49-51° (13679pts)<br>55-57° (52580pts)</td><td>51-53° (29218pts)<br>57-59° (112558pts)</td><td>53-55° (39385pts)<br>59-61° (186906pts)</td></tr></tbody></table>

Figure S2: EBSD on the martensitic domain: misorientation axis distributions for the FeCMo composition (a) and FeCMn composition (b). (color necessary for print)
<table><thead><tr><th></th><th></th><th></th><th></th><th></th></tr></thead><tbody><tr><td>FeCMo<br>300°C/s</td><td>a)</td><td>49-51° (14054pts)<br>55-57° (43482pts)</td><td>51-53° (27395pts)<br>57-59° (79144pts)</td><td>53-55° (33121pts)<br>59-61° (149432pts)</td></tr><tr><td>FeCMn<br>300°C/s</td><td>b)</td><td>49-51° (13679pts)<br>55-57° (52580pts)</td><td>51-53° (29218pts)<br>57-59° (112558pts)</td><td>53-55° (39385pts)<br>59-61° (186906pts)</td></tr></tbody></table>

Figure S2: EBSD on the martensitic domain: misorientation axis distributions for the FeCMo composition (a) and FeCMn composition (b). (color necessary for print)
<figure><img src="image_12.png" /><figcaption>Figure S3: Study in a prior austenite grain for FeCMoMn 1°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
<figure><img src="image_13.png" /><figcaption>Figure S3: Study in a prior austenite grain for FeCMoMn 1°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
<figure><img src="image_13.png" /><figcaption>Figure S4: EBSD for the slowest cooling speeds: misorientation axis distribution for the FeCMo composition (a) and FeCMn composition (b). (color necessary for print)</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>Figure S4: EBSD for the slowest cooling speeds: misorientation axis distribution for the FeCMo composition (a) and FeCMn composition (b). (color necessary for print)</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>Figure S5: Study in a prior austenite grain for FeCMoMn 0.1°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>Figure S5: Study in a prior austenite grain for FeCMoMn 0.1°C/s: a) EBSD map with a prior austenite grain (PAG) in red, b) {100}, {110} and {111} pole figures of the PAG (stereographic projection), c) map of the misorientation angles in the PAG with a profile in white, d) misorientation profile in PAG. (color necessary for print)</figcaption></figure>
## Acknowledgements

The authors would like to acknowledge the Commissariat à l'Énergie Atomique et aux Énergies Alternatives (CEA) for funding. They wish to thank Jean-Baptiste Delattre from the Commissariat à l'Énergie Atomique et aux Énergies Alternatives (CEA) for providing microhardness data on 20MND5.

Thanks are also due to Françoise Barcelo, Amélie Gangloff and Élodie Rouesne from CEA for their help with EBSD measurements. The authors would like to acknowledge Thomas Gilbert for his help with dilatometry, as well as Remy Pires and Julie Bourgon from the ICMPE for their precious help and advice.

The authors would like to thank Anna Fraczkiewicz and her team from the Georges Friedel Laboratory of the Mines Saint-Étienne (UMR 5307) for the production of the model alloys.

## Data availability

The raw and processed data required to reproduce these findings cannot be shared at this time due to legal or ethical reasons.

## References

[1] S. Saillet, N. Rupa, and C. Benhamou, Impact of large forging macro-segregations on the reactor pressure vessel surveillance program (2006) in Fontevraud 6, SFEN, Fontevraud.

[2] E. J. Pickering and H. K. D. H. Bhadeshia, The consequences of macroscopic segregation on the transformation behavior of a pressure-vessel steel, *J. Press. Vessel Technol. Trans. ASME, **136*** (2014) doi: 10.1115/1.4026448.

[3] F. Barcelo, B. Marini, S. Saillet, and P. Todeschini, Metallurgical characterization of micro-heterogeneities in a 16MND5 forging (2018) in Fontevraud 9, SFEN, Avignon.

[4] L. Zhang, B. Radiguet, P. Todeschini, C. Domain, Y. Shen, and P. Pareige, Investigation of solute segregation behavior using a correlative EBSD/TKD/APT methodology in a 16MND5 weld, *J. Nucl. Mater.*, **523** (2019) 434-443, doi: 10.1016/j.jnucmat.2019.06.002.

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## Acknowledgements

The authors would like to acknowledge the Commissariat à l'Énergie Atomique et aux Énergies Alternatives (CEA) for funding. They wish to thank Jean-Baptiste Delattre from the Commissariat à l'Énergie Atomique et aux Énergies Alternatives (CEA) for providing microhardness data on 20MND5.

Thanks are also due to Françoise Barcelo, Amélie Gangloff and Élodie Rouesne from CEA for their help with EBSD measurements. The authors would like to acknowledge Thomas Gilbert for his help with dilatometry, as well as Remy Pires and Julie Bourgon from the ICMPE for their precious help and advice.

The authors would like to thank Anna Fraczkiewicz and her team from the Georges Friedel Laboratory of the Mines Saint-Étienne (UMR 5307) for the production of the model alloys.

## Data availability

The raw and processed data required to reproduce these findings cannot be shared at this time due to legal or ethical reasons.

## References

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