Influence of Alpha/Gamma-Stabilizing Elements on the Hot Deformation Behaviour of Ferritic Stainless Steel

Abstract

This study investigates the hot deformation behaviour and microstructural evolution of two AISI 430 ferritic stainless steel variants: 0A (basic) and 1C (modified). These variants primarily differ in chemical composition, with 0A containing higher austenite-stabilizing elements (C, N) compared to 1C, which features lower interstitial content and slightly higher Si and Cr. This research aimed to optimize hot rolling conditions for enhanced forming properties. Uniaxial hot compression tests were conducted using a Gleeble thermo-mechanical system between 850 and 990 °C at a strain rate of 3.3 s⁻¹, simulating industrial finishing mill conditions. Analysis of flow curves, coupled with detailed microstructural characterization using electron backscatter diffraction, revealed distinct dynamic restoration mechanisms influencing each material’s response. Thermodynamic simulations confirmed significant austenite formation in both materials within the tested temperature range, notably affecting their deformation behaviour despite their initial ferritic state. Material 0A consistently exhibited a strong tendency towards dynamic recrystallization (DRX) across a wider temperature range, particularly at 850 °C. DRX led to a microstructure with a high concentration of low-angle grain boundaries and sharp deformation textures, actively reorienting grains towards energetically favourable configurations. However, under this condition, DRX did not fully complete the recrystallization process. In contrast, material 1C showed greater activity of both dynamic recovery and DRX, leading to a much more advanced state of grain refinement and recrystallization compared to 0A. This indicates that the composition of 1C helps mitigate the strong influence of the deformation temperature on the crystallographic texture, leading to a weaker texture overall than 0A.

1. Introduction

The processing of ferritic stainless steel (FSS) has garnered significant attention owing to their role in various industrial applications, particularly in manufacturing components that require enhanced formability and corrosion resistance [1]. The hot rolling of these materials is critical as it influences their microstructural evolution, which, in turn, affects their mechanical performance and formability. Laboratory-controlled hot deformation processes, such as compression testing at elevated temperatures allow researchers to evaluate how different parameters—like temperature and strain rate—impact the flow behaviour of the material, thus enabling the precise design of the finishing mill process for industrial steel production [2,3]. The underlying mechanisms, notably recrystallization and recovery, are crucial in determining the optimal hot rolling conditions necessary for enhancing ductility and the overall forming behaviour of the alloy [4]. Specifically, dynamic recrystallization (DRX) is a process that occurs during plastic deformation at high temperatures, where new dislocation-free grains are formed, leading to softening of the material. In contrast, dynamic recovery (DRV) is a more subtle process that occurs continuously during deformation, primarily driven by dislocation annihilation and their rearrangement into dislocation walls, forming subgrains [5,6].

Previous studies have demonstrated that fine-tuning hot rolling parameters can lead to favourable microstructural characteristics that significantly improve the material’s performance under deformation. For instance, Mehtonen et al. [7] have shown that controlled hot deformation can promote dynamic restoration mechanisms, which are key to refining grain structure in stabilized FSS. Additionally, their research indicates that the temperature range during the hot working significantly affects static recrystallization and texture evolution, ultimately leading to better drawability of the final product [8]. The synergy of these processes highlights the complexity of optimizing hot rolling conditions, specifically for AISI 430 ferritic stainless steel, where variations in chemical composition can introduce significant differences in material response [9]. Supporting this, Vaughan et al. [10] highlighted the influence of cold rolling on the Lankford coefficient—a measure of drawability—demonstrating that an enhanced understanding of rolling parameters leads to more effective processing routes.

The advancement of knowledge regarding the interplay between hot rolling conditions and the resulting microstructural characteristics is essential for the optimization of FSS. The continuous evolution of methodologies aimed at characterizing these relationships significantly contributes to the effective manufacturing of high-performance FSS components. In this regard, electron backscatter diffraction (EBSD), coupled with scanning electron microscopy (SEM), is indispensable for quantitatively analysing crystallographic orientation, grain size and shape distributions, grain boundary character (e.g., high-angle vs. low-angle boundaries), and texture evolution during hot rolling, which are critical to understanding how variations in hot rolling conditions affect material drawability [11]. In this regard, while Schneider et al. [12] investigated the influence of hot rolling conditions on the microstructure and texture of the hot band in Ferritic FeSi Steels, Kisko et al. [13] specifically observed how hot rolling and coiling temperatures affect the formation of structures and substructures (low-angle boundaries) to enhance drawability and resist ridging in FSS.

Furthermore, the use of thermodynamic simulation software provides crucial insights into phase transformations, equilibrium compositions, and precipitate formation at various temperatures, complementing experimental investigations by predicting microstructural constituents and stability during processing [14]. In this regard, Zhang et al. [15] predicted the formation of (Cr,Fe)₂₃C₆ and AlN precipitates in hot-formed ferritic stainless steel, which was later evidenced by advanced electron microscopy techniques. Nabiran et al. [16] demonstrated the use of CALPHAD modelling in designing FSS for high-temperature applications, specifically to stabilize the microstructure through solid-state precipitation of MX carbonitrides, by calculating phase equilibria and evaluating phase stabilities to determine the influence of alloying elements.

While considerable research exists on the hot deformation of stainless steels [17,18], a detailed comparative analysis focusing on the subtle influences of chemical composition within specific FSS grades, coupled with precise simulation of industrial hot rolling finishing mill conditions, remains an area requiring further investigation. This research aims to fill this gap by studying the hot deformation microstructure of two variations of the same AISI 430 FSS grade in terms of chemical composition and manufacturing process; the AISI 430 FSS grade was studied for materials after cold rolling in a preliminary research study [19]. Hot compression tests, conducted in the temperature range of 850–990 °C and at a strain rate of 3.3 s⁻¹, were performed to establish the most adequate conditions taking place during high-temperature deformation. The industrial finisher pass program consists of four passes at 990, 960, 930, and 850 °C, with an average reduction of 45% per pass and an average speed of 220 m/min. Each pass was simulated individually, not performed in a multi-pass way, at these same temperatures and conditions. The flow curves were examined to understand the deformation characteristics. Microstructural evolutions were studied using SEM and EBSD. We hope our findings contribute to the established body of knowledge regarding optimal processing parameters that enhance the performance of FSS in industrial applications.

2. Materials and Methods

Two samples of AISI 430 steel with different initial compositions, known as basic (identified as 0A) and modified (1C), were sourced from the daily production of Acerinox Europa S.A.U. (www.acerinox.es, accessed on 28 July 2025) at various stages of the processing route. The chemical composition of the samples was analyzed by X-ray fluorescence and spark optical emission spectrometry (S-OES), using instruments commonly found in a typical stainless steel factory. Specifically, a Panalytical Axios Fast X-ray fluorescence spectrometer (Malvern Panalytical, Malvern, UK, www.malvernpanalytical.com, accessed on 28 July 2025), an OBLF Qsn 750 (OBLF Spectrometry, Witten, Germany, www.oblf.de, accessed on 28 July 2025), and a Leco CS 600 and TC 600 analyzers (LECO Corporation, St. Joseph, MI, USA, www.leco.com, accessed on 28 July 2025) were used. The ThermoCalc software (version 2025a using the TCFE12 database) was used to calculate the phase diagrams [14].

Table 1. Composition of the as-supplied materials.

Elements (wt.%)
Sample ID	C	N	Si	Cr	Ni	Fe
0A (basic)	0.050	0.035	0.35	16.3	0.26	balance
1C (modified)	0.025	0.025	0.45	16.7	0.23	balance

summarizes the materials selected for this study, along with their global chemical composition and the type of annealing applied. The composition of sample 0A included a high content of austenite-stabilizing elements (C, N, Ni), while the composition of sample 1C was modified with a lower content of interstitial elements such as C and N.

The microstructure was determined after polishing and etching the samples using Vilella’s reagent for 30 to 75 s. Observation was carried out using an Olympus GX71 light optical microscope (Olympus, Tokyo, Japan, www.olympus.co.uk, accessed on 28 July 2025).

Uniaxial hot compression tests were conducted using a Gleeble 3500-GTC thermo-mechanical testing system to physically model the hot rolling finishing mill (FM) process conditions for each material. Rectangular specimens with overall dimensions of 20 × 15 × 10 mm, featuring a compressed surface of 20 × 5 mm with the compression axis aligned with the normal direction, were taken from the industrial process transfer bar. The specimens were heated at a rate of 15 °C/s to the test temperature and held for 5 s prior to one-hit compression. The temperature–time schematic diagram is shown in Figure 1. The single-hit hot compression tests were carried out under varying temperatures ($T_{FM}$)—those corresponding to the rolling passes of the FM (990, 960, 930, and 850 °C) and at a constant strain rate of 3.3 s⁻¹. The analysis of the resultant stress–strain curves provides insights into the DRX behaviour.

EBSD was performed with a Zeiss Ultra 55 field emission gun scanning electron microscope (FEG-SEM) (ZEISS, Oberkochen, Germany, www.zeiss.es, accessed on 28 July 2025) equipped with a Channel 5 EBSD system from Oxford Instruments (Abingdon, UK, www.oxinst.com, accessed on 28 July 2025), using polished longitudinal sections of the samples. The sample sections were prepared as per conventional metallographic procedures, with a final polish using colloidal silica suspension. EBSD maps were acquired at 20 kV, a working distance of 16 mm, and a 0.45 µm step size, using the post-processing software Tango and Salsa. For the grain boundary component, Tango produces a correlated misorientation angle distribution that makes it possible to obtain the content of grains, subgrains, and twin boundaries. The Tango recrystallized fraction component detects deformed and recrystallized grains and, therefore, the content of these fractions in the analysed area. Salsa is the tool used to generate an orientation distribution function, an approach to texture measurements.

3. Results and Discussion

Figure 2 displays the stress–strain curves obtained from thermo-mechanical compression tests on 0A and 1C ferritic steels at four different temperatures. For the 0A sample, at 850 °C, the curve showed typical high-temperature yielding behaviour, with significant plastic deformation before failure. Softening as a result of DRX can be observed. This suggests that new grains are formed while the deformed grains are consumed [20]. At 930 °C, steady-state flow is observed without a pronounced peak. This can indicate that the softening is attributable to a combination of DRV and DRX [21]. At this temperature, DRV seems to be the dominant mechanism, while DRX probably exhibits different kinetics than at 850 °C. Finally, at 960 °C and 990 °C, the curves showed a smooth peak and lower strength compared to the previous two temperatures, indicating that DRX may be taking place continuously and more effectively [22]. Work hardening appears to be compensated by DRV, and therefore, softening was almost constant at these two temperatures. The stress–strain curves observed (Figure 2) are characteristic of a coexistence of dynamic recrystallization and dynamic recovery, as described by Poliak and Jonas [23].

In Figure 2b, related to the 1C samples, the effect of DRX and DRV is also evident. It can be observed that the flow curves initially increased due to DRX and then started to gently drop as a result of DRV. Based on the curve shapes, DRX appeared to be highly active and efficient at all four testing temperatures for the 1C sample. In addition, softening seemed to be more abrupt in 1C than in 0A.

The most significant discrepancies between samples 0A and 1C are observed in the results of the 850 °C and 960 °C tests. At 850 °C, the softening appeared to be less pronounced in 1C than in 0A. This indicates that work hardening was less significant or was more effectively counteracted by softening mechanisms in 1C in the pre-peak region. At 960 °C, the curve for sample 1C reaches notably higher flow stresses and exhibits more pronounced post-peak softening compared to sample 0A. This suggests that, for the 1C sample at 960 °C, the initial work hardening rate might be higher, or the efficiency of DRV is lower, leading to a greater accumulation of dislocations before DRX is able to dramatically reduce the flow stress.

Despite the initial ferritic state of the samples, the obtained flow stress curves (Figure 2) exhibited characteristics typical of austenitic stainless steel, notably indicating the occurrence of discontinuous DRX. This behaviour suggests a significant influence of the $\gamma$-phase (austenite) during the high-temperature deformation. This is corroborated by the phase diagram in Figure 3, which illustrates the temperature-dependent phase transformations. Specifically, Figure 3 confirms that the ferrite (BCC, body-centred cubic) to austenite (FCC, face-centred cubic) transformation occurs in 0A and 1C alloys when heated from 850 °C to 930 °C, where there is a significant volume fraction of austenite. In this diagram, carbides (K) have been omitted because they represent a small conversion, although they are involved in the ferrite to austenite transformation.

Table 2. Phase transformations taking place and the resulting theoretical austenite and ferrite percentages in alloys 0A and 1C at the temperatures for which the thermo-mechanical compression tests were performed.

Sample	Temperature (°C)	990	960	930	850
0A	Phase trans. reaction	BCC + K → FCC	BCC + K → FCC	BCC + K ← FCC	Finished
	Austenite (v/v)	37	39	38	0
	Ferrite (v/v)	63	61	62	99
1C	Phase trans. reaction	BCC + K → FCC	BCC + K ↔ FCC	BCC + K ← FCC	Finished
	Austenite (v/v)	25	27	25	0
	Ferrite (v/v)	75	73	75	99

presents a comparison of equilibrium phase transformations and the resulting theoretical austenite and ferrite percentages in 0A and 1C alloys at various temperatures, derived from the equilibrium phase diagram. A key point to highlight is that at 960 °C, the phase transformation reactions differ between both alloys: while in 0A the transformation is from ferrite to austenite (BCC + K → FCC), in 1C, the reverse transformation is observed to start (BCC + K ← FCC).

While flow stress curves suggest discontinuous dynamic recrystallization (DDRX), microstructural analysis is essential for definitive confirmation. DRX is a complex process influenced by temperature, strain rate, and material composition. Although test temperatures can promote DDRX, the initial microstructure of the samples may also play a significant role.

To investigate the microstructural evolution during phase transformation, EBSD analysis was performed on samples 0A and 1C at 850 °C and 960 °C. As illustrated in the equilibrium phase diagrams (Figure 3), 850 °C marks the onset of the austenite phase transformation, while 960 °C represents a temperature where a significant proportion of austenite is predicted to form for both alloys within the studied temperature range, approaching their respective maximum volume fractions. Furthermore, Figure 2 showed that the softening mechanisms differ notably between the two samples (0A and 1C) at these two temperatures.

Figure 4 presents the band contrast, grain boundary (GB), and coincidence-site lattice (CSL) maps of these four conditions, along with a histogram displaying the grain size distributions. The white lines/regions most likely represent grain boundaries or subgrain boundaries where the local strain field or atomic arrangement is such that the Kikuchi pattern quality is degraded, leading to low confidence in the indexing or outright non-indexing. This is a common feature in EBSD maps and often indicates areas of high misorientation, lattice distortion, or possibly very fine, unresolved features [24].

The microstructure of sample 0A (Figure 4a) reveals a state characterized by elongated ferrite grains with a varied range of grain sizes and shapes, indicative of the prior hot-deformation process. Some diffuse coloured bands, interpreted as regions containing martensite according to our previous research, were found [25]. The formation of this martensite is attributed to the austenite formation during hot rolling at elevated temperatures, followed by rapid cooling. This formation mechanism was studied in detail in reference [25], where it was found that for the grain size analysis by EBSD (Figure 4e), CSLs were considered special boundaries and not as regular grain boundaries, specially the well-known $\Sigma 3$ twin boundary in face-centred cubic materials, which appeared as concentrated red lines in sample 0A (Figure 4a,b) and represented the martensite fraction. At 850 °C (Figure 4a), the bands interpreted as martensite do not exhibit a clearly preferred crystallographic orientation relative to the ferrite matrix. However, at 960 °C (Figure 4b), these bands appear to align preferentially along the rolling direction (RD), and the ferrite grains seem to have undergone some refinement, appearing smaller and less uniform in size compared to the microstructure at 850 °C.

In contrast, the 850 °C microstructure of sample 1C (Figure 4c) shows a more heterogeneous ferrite grain structure, with a wider distribution of grain sizes and shapes and regions exhibiting very fine grains. The presence of highly-deformed bands is less pronounced in this state of 1C compared to 0A. Notably, these bands almost completely disappear upon heating to 960 °C (Figure 4d).

Figure 4f presents the normalized grain size distributions for the four tested conditions. A prominent feature across all samples is a common peak in frequency at approximately 5 µm, which also corresponds to the mode for 0A at 960 °C and 0A at 850 °C. However, the detailed distributions and the overall fineness of the grains vary significantly between samples and temperatures. For sample 0A, the statistical analysis reveals a generally finer grain structure, with an average grain size of 9.34 µm at 850 °C and a significantly finer grain size of 4.53 µm at 960 °C. This confirms that DRX was effective in producing a significantly refined microstructure. In 0A, at 850 °C, the presence of some larger grains evidences that DRX was not completely dominant in all regions. The mode for 0A at both temperatures is around 1.6–1.7 µm, which aligns with the initial sharp rise in frequency in the histogram. At 960 °C, the distribution for 0A becomes notably sharper and more refined, as indicated by a lower standard deviation (10.46 µm vs. 13.55 µm at 850 °C), a lower median (2.83 µm vs. 4.71 µm), and a strong dominant peak around 5 µm in the histogram, suggesting a more uniform and finer microstructure. For sample 0A at 960 °C, the grain size distribution exhibits a sharp peak and a long tail towards larger grain sizes. This is quantitatively supported by the high kurtosis value of 1504 and a skewness value of 32.58, respectively, even though these characteristics are less visually apparent in the normalized frequency plot. Furthermore, Figure 4a,b revealed a high concentration of low-angle GBs (LAGBs) (those between 1 and 10°) in the microstructure of the 0A sample, particularly at 960 °C. This elevated concentration of LAGBs suggests that the material is primarily undergoing DRV and potentially ongoing continuous DRX (CDRX), where these LAGBs progressively evolve into high-angle GB (those with angles higher than 10°) [26]. This microstructural evidence is in good agreement with Figure 2.

In contrast, sample 1C generally exhibits coarser average grain sizes (mean GS of 18.43 µm at 850 °C and 15.43 µm at 960 °C) and a broader overall grain structure compared to 0A due to the absence of martensite. However, Figure 4c,d present a high concentration of finer grains, with no observable LAGBs. This microstructural characteristic provided clear evidence that the CDRX mechanism was highly active and complete under these specific conditions, effectively consuming the prior recovered substructure. At 850 °C, 1C displays a more dispersed grain size distribution, with a mode around 1.68 µm and a high standard deviation (16.74 µm), indicating a wider range of grain sizes. Increasing the temperature to 960 °C for sample 1C leads to a slight decrease in average grain size (15.43 µm), but the distribution remains broad, as suggested by similar standard deviations and medians to the 850 °C condition. This behaviour, where higher temperatures do not universally lead to coarser grains across all samples, may indicate complex interactions among grain growth, dynamic recrystallization, and phase transformation kinetics in these alloys.

From these results, a correlation can be established between the DRX/DRV behavior and the composition of these materials. The composition of material 0A strongly favours dynamic recovery and potentially CDRX, where LAGBs evolve in some areas and finer grains form in others. On the other hand, the composition of material 1C appears to be active in dynamic recovery and recrystallization, leading to a higher refinement of the recrystallized grains. Overall, in the 1C sample, DRX was more active at 960 °C, while in the 0A sample, DRX activity was more pronounced at 850 °C. Grain size refinement is a direct result of active DRX, in line with the principles of dynamic recrystallization [27,28].

Figure 5 shows $\left\{100\right\}$, $\left\{110\right\}$, and $\left\{111\right\}$ pole figures of 0A and 1C after hot compression at the selected temperatures 850 °C and 960 °C. The colour scales are presented in multiples of a uniform distribution (mud), where higher intensity (yellow to red regions) indicates stronger preferred crystallographic orientations (texture components), and lower intensity (blue to green) suggests more random orientations. All pole figures display some degree of texture, confirming that hot compression induces crystallographic alignment. For sample 0A, when hot-compressed at 850 °C (Figure 5a), the texture exhibits a maximum intensity of 2.84 mud. The $\left\{100\right\}$ pole figure shows relatively low intensity, while the $\left\{110\right\}$ pole figure has some concentrations near the centre. The $\left\{111\right\}$ pole figure displays more distinct partial rings. This weak and randomized texture is consistent with a recrystallization process where DRX was started and promoted dispersed orientations. Upon hot compression at 960 °C (Figure 5b), sample 0A develops a significantly sharper texture, as evidenced by a substantial increase in maximum intensity to 7.70 mud. The pole figures show more concentrated high-intensity regions, particularly in the $\left\{100\right\}$ figure, indicating a strong development of specific preferred orientations at this higher temperature, which indicates that recrystallization has not been completed.

Figure 5c, for sample 1C when hot-compressed at 850 °C, shows a maximum intensity of 3.59 mud. Its pole figures generally exhibit lower overall intensity and more dispersed concentrations compared to 0A at 960 °C, suggesting a weaker or more randomized texture promoted by a more active DRX. When hot-compressed at 960 °C (d), sample 1C also shows a smooth decrease in texture intensity compared to 1C at 850 °C, reaching a maximum of 3.34 mud. More distinct high-intensity regions emerge in the pole figures, indicating texture development. However, the maximum intensity for 1C at 960 °C remains lower than that for 0A at the same temperature, suggesting that while texture develops due to temperature, it does not reach the same strength or sharpness as in sample 0A. At both temperatures for the 1C material, DRX is already active and has a notable impact on the texture, leading it to be relatively weak and diffuse.

The detailed volume fractions of texture components (shown in the bar chart of Figure 5, along with the pole figure observations, demonstrated that the hot compression temperature significantly influenced the resulting crystallographic texture in both stainless steel types (0A and 1C). The specific dominant components and their distributions were notably altered by the hot compression temperature, reflecting the distinct deformation and recrystallization mechanisms active at each temperature. When hot-compressed at 850 °C, sample 0A developed a texture primarily characterized by the $\langle 210\rangle \left|\right|X0$ component. The presence of multiple components with significant volume fractions indicates a relatively complex texture not dominated by a single component. This is consistent with the activation of DRX, which modifies the typical deformation texture. However, when hot-compressed at the higher temperature of 960 °C, the resulting texture was significantly different, with the $\langle 101\rangle \left|\right|Y0$ component becoming overwhelmingly dominant. This indicates that the hot compression temperature played a decisive role in determining the final crystallographic texture in material 0A, leading to distinct and stronger preferred orientations at higher deformation temperatures.

At 850 °C, sample 1C exhibits a dispersed texture with several prominent components, notably $\langle 311\rangle \left|\right|Z0$, $\langle 001\rangle \left|\right|X0$, and $\langle 210\rangle \left|\right|X0$, which strongly suggests that deformation is influencing the orientation of the grains. This is consistent with a more pronounced DRV, although the presence of DRX is already confirmed in the process. When hot-compressed at 960 °C, the texture remained similarly dispersed, but there is a shift in the specific dominant components. The $\langle 310\rangle \left|\right|Z0$ component emerged as the most prominent, alongside strong contributions from $\langle 210\rangle \left|\right|X0$, $\langle 211\rangle \left|\right|Z0$, and $\langle 210\rangle \left|\right|Z0$. This suggests that, for material 1C, hot compression at the higher temperature altered the favoured crystallographic orientations, even if the overall degree of texture dispersion, remained similar. This indicates that the composition of 1C helps mitigate the strong influence of the deformation temperature on the crystallographic texture, leading to more random orientations or a weaker texture overall than 0A.

Figure 6 displays the normal-projected (Z0) inverse pole figure (IPF) orientation map for the 0A and 1C samples after hot compression at the selected temperatures of 850 °C and 960 °C. The microstructure of material 0A at 850 °C (a) displays a mix of equiaxed and somewhat elongated or irregular grains, along with some fine and larger grains with a mix of colours, suggesting a highly varied and diverse distribution of crystallographic orientations. Visually, there is a significant presence of red ($\langle 001\rangle$ oriented) and green ($\langle 101\rangle$ oriented) grains, along with blue ($\langle 111\rangle$ oriented) and other colours. The overall appearance suggests a largely randomized orientation that is highly consistent with very active DRX. In Figure 6b, material 0A at 960 °C presents a highly-concentrated LAGBs microstructure and reveals a clear predominance of $\langle 101\rangle$, forming large clusters and bands. This orientation offers an energetic advantage for nucleation and growth during the hot compression process [22], despite the fact that DRX was not fully completed under these conditions. This is likely due to the fact that DRV is promoted at these temperatures, potentially competing with DRX.

Figure 6c shows that the microstructure of material 1C at 850 °C is dominated by elongated, fine, and aligned grains in bands, which is characteristic of DRX. Fine grains have a predominant orientation along the $\langle 001\rangle$ component and some of the $\langle 101\rangle$ component. This represents a highly preferential distribution aligned with the deformation direction, not a random one. Although DRV/DRX mechanisms are present and refinement has been completed, recrystallization was not fully reached. In Figure 6d, material 1C at 960 °C exhibits a microstructure characteristic of more equiaxed and finer grains, and the crystallographic orientations are more widely distributed compared to 1C at 850 °C, indicating more complete recrystallization of the microstructure. Nevertheless, certain alignments or bands of predominant colours, particularly red ($\langle 001\rangle$), are still somewhat discernible, though less pronounced than in 1C at 850 °C. The proportion of other intermediate colours is higher, contributing to a lower overall texture intensity. In summary, Figure 6 confirmed that DRV and DRX are significantly more active in sample 1C compared to sample 0A, with the latter showing a predominance of DRV. Overall, grain refinement and recrystallization state are much more advanced in sample 1C.

4. Conclusions

This study investigated the thermo-mechanical behaviour of two AISI 430 steel variants (0A and 1C) under plain stress at 850 °C and 960 °C, revealing significant differences in dynamic recrystallization (DRX) and dynamic recovery (DRV) mechanisms influenced by their different chemical compositions. The following conclusions are drawn:

1.

The stress–strain curves extracted from the plane stress compression test, performed to reproduce the hot rolling conditions, exhibited a peak stress followed by a clear softening stage. In most tested conditions, a steady-state flow was observed, providing unequivocal evidence of DRX occurrence. In other cases, the absence of a pronounced peak is attributable to a combination of DRX and DRV. Generally, the curves displayed characteristics of austenitic stainless steel, indicating the occurrence of discontinuous DRX.

2.

The difference in the chemical composition of the two ferritic steels under study significantly influenced the kinetics and extent of the DRX and DRV mechanisms:

• Material 0A, which corresponds to a more standard AISI 430 composition, showed that DRX is highly dependent on phase transformation. This is likely attributed to the higher content of interstitial elements (C, N), which can influence stored strain energy and DRX nucleation.
• Material 1C, a modified composition with lower C, N, and higher Cr and Si contents, showed more efficient DRX highly dependent on temperature. The lower C and N content, along with higher Cr and Si, appears to encourage grain refinement and texture control compared to 0A.

3.

The analysis of the EBSD results gave the following specific conclusions per material regarding their behaviour at the two studied temperatures:

• Material 0A, at 850 °C, showed an active DRX but one that was perhaps not completely homogeneous across the material, with a mixed microstructure with a variety of orientations in a weak and diffuse texture.
• Material 0A, at 960 °C, primarily underwent DRV and potentially ongoing CDRX, where the microstructure evolved from a low- to high-angle microstructure, leading to a very strong and sharp recrystallization texture, demonstrating that DRX has not been completed.
• Material 1C, at 850 °C, developed a heterogenous microstructure based on the bands of fine grains, with a sharp texture, and other areas with larger grains. This evidences that DRV and DRX were active, but only for the refinement of the grains.
• Material 1C, at 960 °C, shows a more homogenous microstructure with a greater refinement of grains. DRV and DRX led to a weaker overall texture.

4.

The results of this article demonstrated that the composition of material 1C improves the plastic deformation response, leading to significant grain refinement and efficient DRX at hot rolling temperatures. For material 0A, DRX is almost entirely inhibited due to a higher fraction of transformed austenite. This causes the mechanical response to work hardening to differ from that of material 1C, notably at lower temperatures of 850 and 930 °C. At these temperatures, the mobility of dislocations is observed to be lower in 0A, with a reduced tendency for recovery and recrystallisation. In contrast, for material 1C, DRX continues to be favoured despite the decrease in temperature. All these phenomena can be clearly seen in the flow curves derived from the high-temperature compression tests. The control of the temperature and the strain level is crucial in 0A for managing the final texture and grain size, directly influencing mechanical properties and product performance.