# Investigation of the influence of manganese on the dynamic strain ageing of ferritic steels – Part II: Industrial alloy

W. Mottay, D. Caillard, H. Idrissi, P. Maugis, A. Portavoce, F. Roch, C.

Perrin-Pellegrino, K. Hoummada

► To cite this version:

W. Mottay, D. Caillard, H. Idrissi, P. Maugis, A. Portavoce, et al. Investigation of the influence of manganese on the dynamic strain ageing of ferritic steels – Part II: Industrial alloy. Acta Materialia, 2025, 301, pp.121576. 10.1016/j.actamat.2025.121576. hal-05286241

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# Investigation of the influence of manganese on the dynamic strain ageing of ferritic steels – Part II: Industrial alloy

W. Mottay<sup>a, *</sup>, D. Caillard<sup>b, D</sup>, H. Idrissi<sup>c, D</sup>, P. Maugis<sup>a, D</sup>, A. Portavoce<sup>a, D</sup>, F. Roch<sup>d</sup>, C. Perrin-Pellegrino<sup>a, D</sup>, K. Hoummada<sup>a, D</sup>

## ARTICLE INFO

### Keywords:

- Plastic deformation

- *In-situ* TEM

- Dynamic strain-aging

- Steel

- Dislocation dynamics

## ABSTRACT

Dynamic strain aging (DSA) was studied in an as-welded C-Mn steel at multiple scales, using macroscopic tensile tests, *in-situ* TEM tensile tests and atom probe tomography (APT). At the macroscopic scale, DSA is related to an increase in Ultimate Tensile Strength (UTS) between 423 K and 623 K, compared to room temperature. PLC serrations are observed in a limited temperature range, 423 K and 473 K. At the nano-scale, DSA manifests as a change in dislocation dynamics, namely the occurrence of dislocation bursts. Interestingly, high-temperature diffusion-controlled (HTDC) Peierls glide suppresses PLC serrations, and its velocity controls the end of DSA. Moreover, the occurrence of HTDC-Peierls glide suggests that interstitial atoms control DSA in the C-Mn weld, that being confirmed by the extracted activation energy for HTDC-Peierls glide and by APT characterizations showing the absence of Mn on dislocations. Contrary to what was observed in the Fe-3Mn-C model in the first part of this study, Mn does not influence DSA in this industrial C-Mn weld. We propose that the reduced Mn content does not lead to a sufficient anchoring of mixed dislocations, thus allowing for the occurrence of HTDC-Peierls glide. Consequently, screw dislocations, which do not attract Mn atoms, become prominent as deformation proceeds, therefore suppressing the Mn-controlled DSA. Based on these findings, we discuss previously published results regarding the manifestations of DSA in steels as a function of their chemical composition.

## 1. Introduction

The study of dynamic strain aging (DSA) in industrial steels is critical for understanding the mechanical behavior of welds, such as those used in nuclear power plants. DSA is a phenomenon that significantly influences the mechanical properties of steels at the macroscopic scale, particularly their strength, ductility, and work hardening rate [1–5]. It is characterized, at the atomic scale, by the interaction between dislocations and solute atoms [6,7], both interstitial (carbon, nitrogen) and substitutional (manganese) ones, which reduces dislocation mobility [8, 9] and generally leads to a hardening of the material but also to a deleterious embrittlement [10–12]. To be more specific, DSA leads to an increased tensile strength, due to the rise in work hardening rate [13], but it may lead to strain localization [4,14], plastic instabilities such as Portevin-Le Châtelier serrations [15,16], and an increase in ductile-brittle transition temperature [11,17].

From an industrial point of view, the best way to avoid this embrittlement is to perform a post-weld heat treatment in order to precipitate most of the interstitials from the deposited metal [17,18]. In binary Fe-C alloys, it was shown [19] by *in-situ* tensile tests in a transmission electron microscope (TEM) that carbon interacts with dislocations during plastic deformation and anchors them, provided the temperature is high enough for carbon diffusion. As a result, deformation then proceeds by brutal dislocation bursts. At higher temperatures, a high-temperature diffusion-controlled (HTDC) Peierls glide mechanism, described in part 1, is observed, corresponding to screw dislocations being able to drag segregated carbon atoms in their core [20,21]. Subsequent researches focused on model ternary alloys [22], Fe-M-C (M = Ni, Si, Cr, Al) and their dislocation dynamics in the DSA temperature range. In these ternary model alloys, two behaviors were observed. First, in Fe-Ni-C and Fe-Si-C systems, the ternary alloy exhibited the same dislocation dynamics as in the binary Fe-C alloy. This was interpreted as

<footnote><sup>a</sup> Aix Marseille Université, CNRS, IM2NP, 13397 Marseille, France</footnote>

<footnote><sup>b</sup> CEMES-CNRS, Université de Toulouse, 31055 Toulouse, France</footnote>

<footnote><sup>c</sup> IMMC, IMAP, Université Catholique de Louvain, B-1348 Louvain-la-Neuve, Belgium</footnote>

<footnote><sup>d</sup> Framatome 92084 Paris la Défense Cedex, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: william.mottay@im2np.fr (W. Mottay).</footnote>
the substitutional solute atoms not interacting with dislocation nor with carbon. The Fe-Cr-C and Fe-Al-C systems were, however, more complex. In both cases, dislocation dynamics is modified by the addition of the substitutional atom, compared to Fe-C alloys. In these systems, DSA is observed from 423 K to 473 K via the occurrence of brutal dislocation bursts. The HTDC-Peierls glide is not observed at 523 K or higher, contrary to what was observed in Fe-C alloys. Finally, the DSA temperature range is extended to higher temperatures, around 773 K.

The present two-part research is dedicated to Fe-Mn-C systems. Indeed, manganese is among the most commonly used alloying elements in steels and it is known to have attractive interactions with carbon, nitrogen and dislocations [23–28]. The effects of these interactions on DSA are studied using in-situ TEM tensile tests on a model Fe-3Mn-C alloy (first part [29]) and on an industrial C-Mn steel weld (present work). Surprisingly, the mechanisms by which Mn atoms influence DSA in steels are not that clear. On the one hand, it was recently demonstrated [30] that Mn addition might reduce DSA sensitivity by precipitating with carbon and nitrogen, to form $\epsilon$-(Fe,Mn)$_{2-3}$(C,N) carbonitrides or $\alpha^{\prime\prime}$-(Fe,Mn)$_{16}N_2$ nitrides. On the other hand, it is not known whether Mn additions around 1–2 %, as is commonly observed in C-Mn and low-alloy steels, might induce DSA. In this sense, model alloys such as the Fe-3Mn-C material studied in part I, provide valuable insights into the fundamental mechanisms of DSA. These simpler systems allow for a detailed study of dislocation dynamics and the interactions between solute atoms. However, although model alloys are essential for isolating and understanding the core mechanisms of DSA, their simplified microstructures make it difficult to extrapolate these findings directly to industrial steels. Industrial alloys are much more complex, featuring a wide range of alloying elements, varied grain structures, and precipitates that can significantly modify dislocation behavior and overall mechanical properties [31]. This complexity makes it challenging to directly apply the results from model alloys to the macroscopic mechanical properties of industrial materials. The present work aims at linking macroscopic mechanical properties of an industrial steel to the nano-scale dislocation dynamics.

The first part of this study [29] provided insights into dislocation dynamics in the DSA temperature range of a model Fe-3Mn-C alloy. We extracted activation energies for dislocation glide based on their velocities. This showed that viscous glide and DSA at elevated temperature are controlled by Mn diffusion. Surprisingly, we showed that manganese diffusion also controls DSA in the low-temperature range. More specifically, we observed that Mn addition hinders the HTDC-Peierls glide observed above 473 K in binary Fe-C alloys. We proposed that Mn is able to co-segregate with carbon on mixed dislocations at around 473 K with the help of pipe-diffusion. This Mn segregation on mixed dislocations is supposed to suppress the ability for dislocations to re-arrange as long straight screw dislocations and give rise to the HTDC-Peierls glide. This second part is extended to a more complex, industrially used C-Mn steel weld. By studying the dislocation dynamics in this industrial material, we aim to understand how the mechanisms unraveled the first part apply to a more complex alloy and to investigate how the interaction of interstitial atoms and manganese influences DSA in a material with a more “realistic” microstructure.

A key aspect of this study is the use of *in-situ* TEM tensile testing on the industrial steel, a technique that has not been applied to such materials in previous research [32,33]. This enables us to correlate the macroscopic tensile tests of the C-Mn steel with dislocation dynamics, providing a detailed understanding of how solutes like nitrogen, carbon and manganese affect dislocation motion and the DSA behavior in this industrially relevant material. In addition to *in-situ* TEM, we use atom probe tomography (APT) to characterize the nanostructure of the studied material. APT allows for mapping the 3D distribution of solute atoms with high spatial resolution [34,35]. By combining APT data with the *in-situ* TEM results, we gain a comprehensive picture of how the solute elements interact with dislocations and how these interactions influence dislocation motion, strain hardening, and the overall mechanical

behavior of this DSA-sensitive steel. By linking atomic-scale insights with the macroscopic mechanical properties of the steel, we aim to provide a deeper understanding of the mechanisms governing DSA in industrial materials and offer valuable information for the optimization of steel alloys used in dynamic loading applications, such as welding and high-strength structural components.

## 2. Experimental

The studied material is an industrial C-Mn steel weld containing 0.065C - 1.37Mn - 0.27Si - 0.013 V - 0.014 N - 0.006S (wt %) (0.3C - 1.38Mn - 0.54Si - 0.014 V - 0.054 N - 0.011S in at %). Aluminum cannot be added in welds to precipitate with N, as it vaporizes in the welding arc and would oxidize before reaching the melt pool. Welding was carried out by deposition of 36 consecutive passes using manual arc welding on 25 mm thick plates. The cooling time from 800 °C to 500 °C was estimated at 5 s for most welding passes. The microstructure of the weld metal is a mix of acicular ferrite and polygonal ferrite, which is consistent with that described in the literature [36,37]. The dislocation density is in the order of $10^{14}$ m⁻² [13,30,38]. The material was studied in the “as-welded” state, without any further heat treatment. Note that this condition is not fully representative of the industrial use, but it maximizes DSA sensitivity. Previous studies on this weld revealed its sensitivity to Dynamic Strain-Aging (DSA), caused by C and N segregation at dislocations during tensile tests [13]. Moreover, it was shown that C, N and Mn segregate on dislocations in this material [39]. The increase in mechanical strength and in work hardening rate caused by DSA is explained by the inhibition of dynamic recovery, which originates from the interstitial solutes-dislocations interaction during plastic deformation at temperatures close to 573 K.

Macroscopic tensile tests were performed on the base metal of the weld from 323 K to 623 K by increments of 50 K to study DSA-sensitivity on the whole DSA temperature range. *In-situ* straining experiments have been carried out in a JEOL 2010HC TEM, using a Gatan low-temperature straining holder, and a high-temperature straining holder designed in CEMES. The movies have been recorded by a Megaview III video camera working at 50 images per second and analyzed frame by frame. Micro-samples were rectangles with an electro-polished thin-edged hole, glued with a high-temperature cement on two rings adapted to the straining holder. Further details on the experimental setup can be found in references [19,22]. Frames of the video are presented, as well as difference images between two frames. In the difference image, the grey color indicates no difference between the two parent images, while a black or white color indicates a displacement, such as a moving dislocation. All the videos were obtained during a single experiment in the same grain of a sample deformed at different temperatures, summing up to 50 min of video. The grain is oriented in the (210) plane, while the tensile axis is [1-25], vertical in all figures. Observations were performed with *g* = [002], corresponding to *a* + 10° tilt angle.

Atom Probe Tomography (APT) experiments were carried out on a Cameca Leap 3000X-HR to investigate segregation on dislocations, in voltage mode at a base temperature of 50 K, 20 % pulse fraction and 200 kHz pulse rate. Data were analyzed using IVAS 3.6.14 software. Samples from the C-Mn weld were prepared using a dual-beam focused ion-beam and scanning electron microscope (FIB-SEM) Helios Nanolab 600 from Thermofischer, using the standard lift-out procedure [35].

## 3. Results

### 3.1. Macroscopic tensile tests

Macroscopic tensile tests were performed on the base metal of the weld at 293 K and from 323 K to 623 K by increments of 50 K. Defining the beginning of DSA as the first temperature from which mechanical strength becomes higher than at 293 K, it is observed that DSA macroscopically starts at 423 K. The maximum in UTS is reached at 523 K. UTS
value at 623 K is nearly equal to that at 293 K. Interestingly, serrations on the tensile curves are observed at 423 K and 473 K, corresponding to Portevin–Le Châtelier (PLC) instabilities [16,27,40]. The DSA temperature range observed in the present C-Mn weld base metal is in agreement with results from similar C-Mn steels [2,10,41]. Interestingly, the PLC serrations are observed in the same limited temperature range [10], around 423–473 K, while the maximum value for UTS is also reached at 523 K.

Regarding the C-Mn weld metal, tensile tests were performed solely at 293 K and 573 K, revealing its sensitivity to DSA, i.e. UTS increase, at 573 K, without PLC serrations [13]. At 293 K, the yield point phenomenon is highly marked. The stress drop is about 50 MPa and the Lüders strain is about 3 %, observed around 530 MPa. UTS reaches 664 MPa. At 573 K, the yield point phenomenon and Lüders strain are no more observed, as in the base metal. The yield stress is around 450 MPa while UTS increases to 792 MPa. Overall, the tensile tests at 293 K and 573 K exhibit the same features in the base and the weld metals. Yield point phenomenon is observed at 293 K and disappears at 573 K, while the work hardening rate and the flow stress increase at 573 K. The main difference between the weld and base metals relies on the difference in flow stress. The weld metal presents a more important strength than the base metal, at the expense of uniform elongation. This difference originates from the microstructural differences between the weld and base metals. The weld metal has a finer grain size and higher dislocation density, contributing to the higher flow stress.

The weld's DSA sensitivity was explained by a decrease in dynamic recovery, as suggested from the Kocks-Mecking model [42,43], most probably caused by interstitial atoms segregation on dislocations during the plastic deformation. The decreased dynamic recovery leads to an increased work-hardening rate, which overall suggests that dislocation dynamics is modified in the DSA temperature range. This point is further studied in the next section, using *in-situ* TEM tensile testing.

## 3.2. *In-situ* TEM tensile tests

*In-situ* straining experiments were performed at increasing temperatures in order to study the effect of DSA on dislocation dynamics. Experiments at 100 K and 200 K, available as supplementary materials, revealed a slow and viscous glide, characteristic of the low-temperature Peierls mechanism [8,44]. At 313 K, see Fig. 2, dislocations are curved and their movement is rather fast, in accordance with the athermal plasticity regime.

In Fig. 2, we observe individual movements of a near-screw dislocation from 1 to 1' and 1'''. Other similar near-screw dislocations are often anchored on macro-kinks. The slip traces are sinuous, as a result of intensive cross-slip. This mechanism corresponds to that already observed in all the iron samples containing high concentrations of substitutional atoms [45]. The first evidences of DSA are observed at 423 K, Fig. 3. At this temperature, dislocations glide at high speed over long distances, before being firmly anchored. Once the slip bands appear, in less than one frame of the video, no further individual, slow motion of dislocation is observed. This corresponds to the beginning of a low-intensity DSA, which is in rather good agreement with the beginning of DSA at the macroscopic scale, Fig. 1. It is worth noting the weak intensity of dislocation bursts, weaker than that observed in ternary Fe-Cr-C and Fe-Al-C alloys [22] and in the Fe-3Mn-C model studied in part I. However, whatever the added substitutional atom, the first occurrence of DSA is observed near 450 K. This regime of dislocation motion, by bursts or avalanches, characteristic of DSA, still goes on at 513 K, Fig. 4. Once again, high speed dislocation glide is observed, and appear in less than one frame of the video. These fast-gliding dislocations are then stopped, sometimes in the observed zone, as the dislocation marked (d) in Fig. 4. The Burgers vector is 1/2 [11–1] in the (101) plane. The screw direction is indicated in projection at tilt +10°. This dislocation is a screw one in agreement with the sinuous slip traces characteristic of frequent cross-slip. No individual slow movement of

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** True stress-true strain tensile curves obtained on the base metal of the weld at increasing temperatures, from 293 K to 623 K. DSA starts at 423 K and is nearly finished at 623 K. PLC serrations are observed at 423KC and 473 K.</figcaption></figure>

dislocation is observed.

At 553 K (Fig. 5), most of the plasticity is still ensured by dislocation bursts. However, some anchored dislocations begin to glide slowly and viscously. A straight screw dislocation, marked (1) in Fig. 5, moves slowly, at a velocity estimated as 0.3 nm.s⁻¹. The Burgers vector is 1/2 [11–1] in the (101) plane. Cross-slip from the (101) plane to a plane close to (112) is observed. This viscous glide of a straight screw dislocation around 550 K is the signature of the HTDC-Peierls mechanism, already studied in references [19,20,22]. Dislocation dynamics is similar at 613 K, Fig. 6. Once again, plasticity is mainly controlled by fast-gliding dislocations, but the viscous glide of straight screw dislocations is faster than at 553 K. The straight screw dislocation labeled 1 glides to position 1' within about 10 s, with an average velocity of 2.7 nm.s⁻¹, nearly an order of magnitude higher than that measured at 553 K.

At 673 K (Fig. 7), evidences of DSA can hardly be detected. Very few dislocation bursts remain and most of the plasticity is achieved by viscous glide of faster screw dislocations. The dislocation identified in Fig. 7 glides at an average velocity of 16 nm.s⁻¹. This indicates a progressive transition to homogeneous and steady glide plasticity, which is confirmed by straining at higher temperatures, available in supplementary materials.

The results from *in-situ* TEM tensile tests with increasing temperatures in the industrial C-Mn steel weld reveal different mechanisms of plasticity. From 100 K to approximately room temperature (RT), plasticity is lattice-controlled. Dislocations are straight in their screw orientation and glide in a slow, viscous manner, controlled by the HTDC-Peierls mechanism. Lattice-controlled plasticity progressively disappears around RT in favor of obstacle-controlled plasticity, characterized by faster and curved dislocations. These results, unpresented here, are available in supplementary materials.

A new regime of dislocation motion emerges near 420 K. From this temperature, no slow dislocation motion is observed. Plastic deformation relies on bursts of individual or small groups of dislocations. This is characteristic of a weak DSA regime, where dislocations are firmly locked by segregated solute atoms. Around 550 K, some straight dislocations start to glide slowly, in agreement with a HTDC-Peierls mechanism. This was already observed by Caillard & Bonneville [19] in Fe-C model alloys and attributed to the segregation of carbon to screw dislocation cores [20]. The temperature being sufficiently high, dislocations are able to drag the segregated carbon. In the present work, it cannot be stated whether carbon or nitrogen are responsible for DSA. Both dislocation bursts and HTDC-Peierls mechanism co-exist between 550 K and 670 K. DSA vanishes at about 670 K in favor of the sole
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** "Athermal" glide of dislocations at 313 K. The motion is called athermal because it is highly discontinuous with strongly varying instantaneous velocities, independently of the dislocation character (edge/screw). Locking and unlocking events from various obstacles are frequently observed. The slip traces at the two foil surfaces are noted "tr".</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** Rapid slip traces (noted tr) appearing in less than one frame of the video (0.04 s) at 423 K. No further dislocation motion is observed during the following 6 s.</figcaption></figure>

viscous (or steady) glide regime. The observation of the HTDC-Peierls mechanism and the absence of high temperature DSA (above 700 K) is rather surprising. Even though the present weld contains Mn as an alloying element, its effect on dislocation dynamics is not observed, in contrast with the Fe-3Mn-C model alloy studied in part I.

### 3.3. APT characterizations

*In-situ* TEM tensile tests from the previous sections suggest that

dislocation motion in the DSA regime is mostly controlled by interstitial atoms segregation on dislocations. Such feature is further confirmed by APT performed in the C-Mn steel weld. An APT volume from the C-Mn steel is presented in Fig. 8a. Each colored dot represents an atom. Iso-concentrations of N with a 1 % threshold value are displayed in green. They reveal α'' carbonitrides – (Fe,Mn)₁₆N₂ – which were previously reported in this material [30]. A 3 % isoconcentration of C is presented, in brown. This C-enrichment is rather linear and cylindrical, as can be seen from the top and axial views, and is thus interpreted as carbon
<figure><img src="image_6.png" /><figcaption>**Fig. 4.** Rapid slip event (in <0.04 s) and locking of dislocation d at 513 K. The dislocation is locked in the pure screw orientation. The slip traces at the two foil surfaces are noted "tr".</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Fig. 5.** Rapid slip and cross-slip event (in <0.16 s) of a dislocation coming from (or stopping at) a grain boundary noted GB, at 553 K. Note that the two slip traces noted "tr" can be deduced from each other by a translation parallel to the Burgers vector. During the 23 s following this event, no other motion can be observed, except for a very slow and steady motion of the straight screw dislocation noted 1 (0.3nm/s).</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>**Fig. 6.** Slow and steady motions of two screw dislocations at 613 K. The screw segments noted 1 moves to 1' within about 10 s at the velocity 2.7 nm/s.</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Fig. 7. Steady motion of a screw dislocation from 1 to 1' and 1'' at 673 K (16nm/s).</figcaption></figure>

atoms segregation on a dislocation. The 1D concentration profile across the dislocation (radial direction) evidences the segregation of both carbon and a few nitrogen on the dislocation, but no Mn-enrichment is detected.

APT characterizations were also performed on the deformed C-Mn weld, after failure during tensile testing at 573 K, Fig. 8b In this volume, a planar 1 % carbon isoconcentration (brown) corresponds to a C-segregated grain boundary. Besides, a 3 % Mn isoconcentration (yellow) reveals a network of parallel lines, corresponding to a Mn-segregated low-angle grain boundary [46]. The top view of the APT volume shows that these two planes are tilted relative to each other. Four concentration profiles, orthogonal to the Mn-segregated plane, are plotted. The first one is located somewhat across the rotation axis, the fourth one is located where the distance between the C-segregated and the Mn-segregated plane is maximal. These profiles reveal that, in position 1, C and Mn peaks are superimposed. In positions 2, 3 and 4, the C and Mn peaks are shifted by a value which increases while the distance between the two planes increase. No carbon peak is observed at the location of the Mn peak, meaning that carbon is not segregated with Mn on the yellow plane. However, there might be a small peak of Mn correlated with the brown plane.

The following scenario is proposed to explain these observations. We suggest that the yellow plane was a Mn-segregated low-angle grain boundary (LAGB), as it consists of a dislocation network, which was present in the volume prior to the tensile tests at 573 K. During plastic deformation, this planar defect rotated. From a crystallographic point of view, there is no more planar defect at the yellow plane. The grain boundary is now located at the brown plane. The Mn-enriched lines would only correspond to segregated Mn atoms which remain stuck in their initial place, as Mn is not mobile at 573 K [47,48]. However, since carbon diffuses fast enough at the tensile temperature, it could segregate to the new location of the grain boundary. Then, this APT volume reveals three interesting elements. First, it confirms that carbon is indeed mobile at 573 K. Second, it evidences that Mn is able to segregate on edge dislocations, which networks form the LAGBs. This result is in agreement with the ones obtained in the first part of this research, as well as results from the literature evidencing Mn segregation on LAGBs or edge dislocations [26,46,49]. Finally, manganese bulk diffusion is

negligible at 573 K, at is does not able to diffuse to the new location of the grain boundary.

# 4. Discussion

## 4.1. Correlation between DSA at the macro-scale and nano-scale

The results presented in Sections 3.1. and 3.2. allow for establishing a direct correlation between DSA at the macro and nano scales in the C-Mn weld. In other words, it is possible to directly observe the modification of dislocation dynamics in the DSA temperature range and to analyze its impact on the mechanical response. The UTS values from Fig. 1 are plotted as a function of tensile temperature in Fig. 9. Based on the variations of UTS, 3 macro-scale plasticity regimes are proposed:

1. $T \le 373$ K, absence of DSA. UTS decreases while temperature increases, dislocation motion is athermal;

2. $423 K \le T \le 623$ K, DSA regime revealed by UTS values larger than those at 293 K. Dislocation motion consist mainly of bursts. HTDC-Peierls glide can be observed;

3. $T > 623$ K, end of DSA. UTS decreases once again. Dislocation bursts are no more observed, HTDC-Peierls glide controls plasticity;

The exact temperatures for the transition between each regime are unknown. As an example, DSA starts between 373 K (no DSA observed) and 423 K (DSA observed). Besides, the DSA regime is split in two sub-domains, one where PLC serrations are observed [16,40,50], at 423 K and 473 K, and another one where PLC serrations are no more observed, from 523 K to the end of the DSA regime. Finally, dislocations dynamics in each domain are indicated. Interestingly, dislocations dynamics change between each domain, and the transition is observed at a temperature in excellent agreement with the macroscopic scale observations. As expected, the athermal regime for dislocation motion corresponds to the absence of DSA at the macroscale. Then, the beginning of DSA, as well as PLC serrations, are correlated with the appearance of dislocation bursts. At the authors' best knowledge, this is the first time that dislocation bursts are directly observed and correlated with PLC serrations. The occurrence of HTDC-Peierls glide marks the end of
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** (a) APT volume from the as-welded C-Mn steel in which a dislocation is observed. C and N are segregated on the dislocation (b) APT volume from the as-welded C-Mn steel after tensile tests at 573 K. Prior to the deformation, a Mn-segregated LAGB was present. It moved under the applied stress to form a new GB, which nature is unknown, segregated only by carbon.</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Summary of the results regarding mechanical properties. UTS is plotted as a function of tensile temperatures. The variations in UTS values allow for the determination of 4 domains, which are linked to the regime of dislocation motion. Domain 1: No DSA, UTS decreases while temperature increases, dislocation motion is athermal. Domain 2: DSA + PLC, UTS increases, dislocations move by brutal bursts. Domain 3: DSA, UTS decreases but stays higher than at room temperature. Dislocation bursts are still present but HTDC-Peierls glide is also observed. Domain 4: End of DSA. UTS is lower than at room temperature. HTDC-Peierls glide controls plasticity.</figcaption></figure>

PLC serrations, but DSA is still present. The end of DSA is explained by the disappearance of dislocation bursts, at 623 K. These results, obtained for the C-Mn weld, could be extended to the model alloys from part I, Fe-C and Fe-3Mn-C.

Regarding Fe-C, the same 4 domains are expected. The beginning of DSA (transition between domain 1 and 2), is a function of the carbon content. For the Fe-C 16 ppm.at alloy, DSA starts around 400 K. The transition from domain 2 to domain 3 relies on the appearance of HTDC- Peierls glide. As was shown by Caillard [22] and confirmed in the two parts of this research, HTDC-Peierls glide starts once the couple (dislo- cations + segregated carbon atoms) are mobile enough to be dragged. At the strain rate used in these works, this happens around 500 K. A higher strain rate would shift the appearance of HTDC-Peierls glide to higher temperatures and vice versa. Dislocation bursts and HTDC-Peierls glide were not observed simultaneously in the Fe-C alloy, compared to what is observed in the C-Mn weld. This might be explained by the complexity of the industrial C-Mn weld compared to the model Fe-C alloy (slight Mn segregation delaying the HTDC-Peierls glide, smaller grain size, higher dislocation density). Finally, the end of DSA, marked by the disap- pearance of dislocation bursts, relies on the HTDC-Peierls glide being fast enough to accommodate the imposed strain rate. Then, the tem- perature for the end of DSA (about 625 K in the present work) is once again a function of the imposed strain rate. A higher strain rate would shift the end of DSA to higher temperatures.

The case of the model Fe-3Mn-C alloy from part I is much more complex. DSA starts at the same temperature, about 400 K, as the carbon content and the imposed strain rate are identical to the two other ma- terials. However, it was shown that HTDC-Peierls glide was suppressed in the Fe-3Mn-C model material. The temperature range where dislo- cation bursts are observed is much more important, somewhat up to 750 K. This suggests that the second domain of DSA (UTS increase + PLC) is extended to higher temperatures, up to 750 K. Based on this interpre- tation, it can be expected that PLC serrations on macroscopic tensile curves of the Fe-3Mn-C alloys would be observed between 450 K and 750 K. The third domain would rely on the appearance of viscous glide of mixed dislocations rather than on the appearance of HTDC-Peierls glide. Once again, this is also a function of the imposed strain rate. Finally, the end of DSA would be observed once the viscous glide be- comes fast enough to accommodate the imposed strain rate.

This correlation between macro and nano scale reveals interesting insights. First of all, DSA regime is clearly linked to dislocation bursts. This can be understood in terms of solute / dislocation interactions. At a given temperature, which depends on the considered solute diffusion and the applied strain rate, solute atoms are able to segregate on dis- locations temporarily arrested, via pipe-diffusion [42,43]. These dislo- cations are then anchored and cannot be put back into movement, leading to a mechanical paradox: a deformation is imposed but the carriers of the deformation, namely the dislocations, are immobilized. Plastic flow then becomes unsa, dislocations bursts are emitted in various areas of the samples, most probably as a result of stress con- centrations. Ultimately, the emission of avalanches of dislocations leads to macroscopic stress drops, known as PLC serrations. At about 523 K, interstitial solutes diffusion appears fast enough for HTDC-Peierls glide to occur. Even though the material is still in the DSA regime, this slow viscous movement of screw dislocations suppresses PLC serrations. We suggest that HTDC-Peierls glide tends to homogenize plastic flow, making it smoother. Dislocations bursts are still observed, but they do not transfer as stress drops on the macroscopic tensile curves. Finally, upon increasing again the temperature, HTDC-Peierls glide is acceler- ated, as measured in Fig. 5, 6 and 7. We suppose that DSA disappears once HTDC-Peierls glide is fast enough to accommodate the imposed strain rate, as was proposed in the first part of this research. The link between the end of DSA and the disappearance of dislocations bursts is thus a simple consequence of HTDC-Peierls glide being fast enough.

Let us discuss now the mechanisms underlying these correlations between macro and nano scale. The end of DSA relies on a HTDC-Peierls glide being fast enough to accommodate plastic deformation. This means that the end of DSA, for a defined material, depends solely on the imposed strain rate. Besides, it was previously shown that the activation energy for the HTDC-Peierls glide is close to that of carbon (or nitrogen) diffusion [19,20]. It results that, for a given strain rate, the first occur- rences of HTDC-Peierls glide will always be observed at the same tem- perature. Most of the tensile tests being performed at a strain rate of about 10⁻⁴ s⁻¹ [10,51,52], HTDC-Peierls glide should be mostly observed around 473 K and 523 K in steels. As HTDC-Peierls glide marks the end of PLC serrations, we can thus state that PLC disappears once carbon or nitrogen atoms are sufficiently mobile to be dragged by dislocations. Finally, the beginning of DSA relies on the anchoring of temporarily arrested dislocations on obstacles of the microstructure, including forest dislocations. Once again, this relies on the diffusion of interstitial atoms. However, in this case, it implies short-circuit diffusions, such as pipe-diffusion along dislocation cores, these mechanisms being much faster than bulk diffusion [53,54]. This is somehow in accordance with results available in the literature, showing that the activation energy for the end of DSA is close to the activation energy for the bulk diffusion of the solute atom, while the activation energy for the beginning of DSA is about half of the bulk diffusion activation energy [15,28,55,56].

In other words, DSA starts when pipe-diffusion of solute atoms is fast enough to anchor temporarily arrested dislocations. From this point, plastic flow becomes unstable and relies on dislocation bursts, up to a temperature where it is kinetically possible for HTDC-Peierls glide to be activated, i.e., a temperature where bulk diffusion of the solute becomes non-negligible. This viscous movement of screw dislocations stabilizes the plastic flow and suppresses the PLC serrations, even though dislo- cation bursts are still required to accommodate the imposed strain rate. Finally, the end of DSA is reached once HTDC-Peierls glide is fast enough to fully accommodate plastic deformation.

## 4.2. Solute atoms – dislocations interactions

As in the first part of this research, we measured the velocity of dislocation glide at different temperatures in order to extract activation energies of viscous dislocation glide, which values allow determining the solute responsible for DSA. Indeed, it was previously shown that activation energy for HTDC-Peierls glide corresponds to that for C or N
diffusion plus the energy for the formation of a critical kink pair [20]. In the case of viscous glide of mixed and edge dislocations (as is the case in the Fe-3Mn-C model alloy from part I), the activation energy corresponds directly to that for the diffusion of the segregated solute atom.

Fig. 10 shows the logarithm of dislocations glide velocity $\ln(v)$ as a function of 1000/T for the studied C-Mn weld (red squares), the Fe-3Mn-C model alloy (yellow circles) studied in part I and a model Fe-C (black triangles) which data are taken from Caillard and collaborators [19,22]. All experimental data are correctly fitted by lines, as expected from thermally activated processes. DSA temperature ranges are plotted as horizontal dotted lines for each material (yellow for Fe-3Mn-C, red for the C-Mn weld and black for Fe-C). As was observed in part I, DSA vanishes once dislocation glide velocity reaches $2-3 \text{ nm.s}^{-1}$ ($\ln(v)=1$), a value sufficient to ensure smooth plastic deformation, without dislocation bursts, in these experimental conditions. This observation confirms the scenario proposed in Section 4.1., the end of DSA is reached when HTDC-Peierls glide becomes fast enough to ensure a homogeneous plastic deformation. It might seem surprising that the difference in dislocation density between the different samples does not impact the dislocation glide velocity required to reach the end of DSA. However, the mobile dislocation density must be considered for each sample, rather than the total dislocation density, including mobile and immobile (forest) dislocations. It is generally considered that the mobile dislocation density is one to two orders of magnitude lower than the total dislocation density (i.e. $1-10\%$ of dislocations are mobile during plastic deformation). We then suggest that the mobile dislocation density is nearly identical in all the tested samples. Interestingly, the lower temperature of DSA is about the same for each material, around $450 \pm 30 \text{ K}$. As was shown by Caillard [22], the lower temperature of DSA, at a similar strain rate, in Fe-C alloys depends solely on carbon content. Then, as expected from their identical carbon content and the identical imposed strain rate, Fe-C (16 ppm.at) and Fe-3Mn-C (16 ppm.at C) model systems have about the same temperature for the beginning of DSA. Surprisingly, the industrial C-Mn weld exhibits a similar beginning of DSA, around 420 K, suggesting that its interstitial content in solid solution (carbon and nitrogen) is also close to 20 ppm.at. It might seem surprising that the weld exhibits such a low C and N content in solid solution. However, as was previously shown, on the same C-Mn weld [30], solute C and N are mostly precipitated as metastable $\epsilon$.

<figure><img src="image_12.png" /><figcaption>Fig. 10. Logarithm of the high-temperature viscous dislocation velocity, $\ln(v)$, ($v$ expressed in nm/s) as a function of 1000/T. Experimental data from the as-welded C-Mn steel are in red. The values for the model Fe-C (black) are taken from Caillard & Bonneville [19]. Values for the Fe-3Mn-C model alloy from part I are also reported (yellow). Dashed lines are the linear fits of these data. Dotted horizontal lines are the DSA temperature ranges for each material.</figcaption></figure>

iron-manganese carbonitrides, in accordance with the rather low C and N content in solid solution estimated in the present work. Interestingly, these metastable precipitates are dissolved upon heating, at temperatures as low as 573 K. Then, it could be expected that C and N content in the matrix of the C-Mn weld progressively increases as the testing temperature increases. This might explain why the DSA starting temperature of the C-Mn weld seems slightly lower.

As was previously shown, the activation energy for dislocation glide is 1.0 eV in Fe-C alloys, corresponding to HTDC-Peierls glide, and 2.8 eV in the Fe-3Mn-C, corresponding to viscous glide of Mn-segregated dislocations. Interestingly, the activation energy in the industrial C-Mn weld is 1.1 eV, in accordance with the observed HTDC-Peierls glide and the absence of segregated-Mn on dislocations, in APT observations. This means that DSA in the C-Mn weld is controlled by carbon or nitrogen core segregation to screw dislocations, in agreement with the macroscopic results [13,57,58]. It however appears that the dislocation glide velocity is shifted to higher temperatures in the industrial C-Mn weld with respect to the binary Fe-C alloy. This is most probably explained by hardening from microstructural features, such as the higher dislocation density or the presence of precipitates, as the ones observed in Fig. 8a, and alloying elements [31]. Then, the addition of 1.3% manganese does not change the dislocation glide mechanism taking place in binary Fe-C alloys. This is unexpected considering that manganese addition did modify the dislocation dynamics in the Fe-3Mn-C model alloy. Such behavior deserves special attention.

For the Fe-3Mn-C alloy, we proposed in the first part a mechanism for dislocations bursts, relying on mixed dislocations being first anchored by carbon atoms, at the beginning of DSA. Then, Mn could segregate on the already C-pinned dislocations, via pipe-diffusion. As a result, mixed dislocations are favored as they cannot rearrange as straight screw dislocations. The multiplication of mixed dislocations during deformation thus leads to a high temperature range of DSA, up to about 800 K. Plastic deformation proceeds by dislocation bursts up to a sufficient temperature for dislocations to drag their manganese atmosphere. However, some scarce unlocking events of dislocations were observed in the DSA temperature range of the model Fe-3Mn-C, giving rise to the HTDC-Peierls glide. These rare unlocking and relocking events of dislocations suggest that some dislocations are able to escape from segregated Mn atoms but still drag segregated C atoms. Considering a Gaussian distribution of Mn segregation on dislocations, the unlocked dislocations are those that are the least anchored by Mn atoms. If manganese concentration is reduced, as is the case in the C-Mn weld compared to the Fe-3Mn-C, manganese segregation is also reduced [39,59], thus increasing the proportion of dislocations able to escape from Mn segregation.

Therefore, we propose that the reduced Mn content in the C-Mn weld leads to an insufficient anchoring of mixed dislocations by Mn atoms, as sketched in Fig. 11. Mixed dislocations are temporarily pinned by Mn atoms. Once they are able to unlock, they re-arrange as straight screw dislocations because of interstitials core segregation, finally allowing HTDC-Peierls glide to be observed and leading to two major consequences. First, the high temperature HTDC-Peierls glide will stabilize plastic flow, suppress PLC serrations at the macro-scale and reduce overall DSA sensitivity. Second, plasticity becomes controlled by HTDC-Peierls glide of screw dislocations. As a result, they multiply as deformation proceeds and the screw character of dislocations becomes prominent in the microstructure. Since screw dislocations do not produce hydrostatic strain field [9,60], substitutional atoms do not interact with them, thus explaining the absence of high-temperature range of DSA in the C-Mn weld. It could be argued that Mn atoms, and more generally substitutional solutes, could nevertheless interact with screw dislocation cores. However, the absence of DSA in the high temperature range and the non-interaction between substitutional atoms and screw dislocations strain field proves that Mn atoms do not segregate in screw dislocation cores.

Overall, the results presented in the two parts of this research allow
<figure><img src="image_13.png" /><figcaption>**Fig. 11.** Schematic representation of the effect of Mn on DSA in the C-Mn weld. C segregates on mixed dislocations at the beginning of DSA. Mn might co-segregate with C, but the anchoring is not strong enough. Mixed dislocations are able to unlock from their Mn atmosphere, thus giving rise to the HTDC-Peierls glide of screw dislocations.</figcaption></figure>

discussing the optimal Mn content desirable in steels to reduce DSA sensitivity. Based on the present work and on the article from Landes [30], it appears that the 1.37 % Mn in the C-Mn weld does not induce DSA but is already sufficient to partly precipitate C and N from the solid solution and thus reduce DSA sensitivity. As was shown in the first part, 2.8 % Mn does induce DSA, but some scarce unlocking events of dislocations were observed, suggesting that dislocations anchoring by Mn was not that strong. Based on these results, we propose that the optimal Mn content to reduce DSA sensitivity might be around 2 % Mn. However, it must be noted that this value is proposed for model alloys or C-Mn steels. It has not yet been determined whether the addition of multiple substitutional elements (Mn, Cr, Al ...) leads to additive effects regarding DSA sensitivity.

This comparative study between a model Fe-3Mn-C alloy (part I) and an industrial C-Mn weld (part II) provides interesting insights on the elementary mechanisms controlling DSA sensitivity and the observation

of PLC effects in various steels. Seminal works from Leslie et al. [27,56] revealed the complex correlations between a given alloy with a specific chemical composition and the observation of PLC serrations. Building on Caillard's research [22], the present article might provide a better understanding of the link between alloy composition and PLC serrations. Four parameters have to be considered, which are summarized in Fig. 12:

1. The very first consideration relies on the interstitial content in the studied steel. The total absence of carbon / nitrogen atoms in solid solution should lead to an absence of low-temperature DSA (423–623 K). The athermal motion of mixed dislocations would then control plasticity. However, residual traces of interstitial atoms (as low as 1 ppm.at [19,22]) are sufficient to favor HTDC-Peierls glide of screw dislocations and associated DSA.

<figure><img src="image_14.png" /><figcaption>**Fig. 12.** Decision tree summarizing the possible observations of DSA as a function of the alloy chemical composition.</figcaption></figure>
2. The second consideration concerns the nature of the substitutional element added in the steel. We can consider two possibilities. First, the substitutional atom does not interact strongly with dislocations, as it seems to be the case for Ni or Si [22]. Then, HTDC-Peierls is allowed, and the mechanical behavior depends solely on interstitial atoms content as in case 1. Second, the substitutional elements interact with dislocations and can inhibit the HTDC Peierls. A third point must then be considered.

3. The third element to consider is the substitutional atoms content, if this element interacts with dislocations, as is the case for Mn or Cr [22]. If the C content is null, a sufficient substitutional content can lead to a high temperature DSA, once the diffusion of the latter is fast enough. If the C content is not null, an insufficient substitutional content leads to a C-controlled DSA, i.e. a low temperature DSA. The high temperature DSA is suppressed, because of the prominence of screw dislocations. Finally, a sufficient substitutional content can extend the DSA temperature range. A low temperature range exists as C content is not null, but the substitutional atom is able to co-segregate even at low temperatures, thus hindering the HTDC-Peierls glide. Dislocations keep their mixed characters, which induces high temperature DSA.

4. Finally, the exact temperature ranges of DSA and serrations on the tensile curves depend on other parameters, such as the applied strain rate and the dislocation density. Indeed, a lower strain rate tends to increase the arrest time of dislocations on obstacles [60], thus increasing the probability of dislocations pinning by solute atoms. Besides, a higher dislocation density provides more diffusion short-circuits, which increases the apparent diffusion coefficient.

The reasons for which some substitutional solutes interact with dislocations and some do not are still not that clear. Excluding dislocation cores, it is not expected that any substitutional solutes segregate around screw dislocations. Considering edge and mixed dislocations, the segregation energy is a function atomic size difference between the solute element and iron atoms. That being said, Si would then be expected to segregate on dislocations and modify dislocation dynamics, which was not observed experimentally. It could be argued that this relies on Si-C repulsive interactions. In other words, the segregation of exclusively of C atoms on dislocations could be energetically more favorable than the segregation of both Si and C, because of the repulsive interactions between the two elements.

To summarize, this study sheds new light on the complex mechanisms underlying DSA, especially the solute – solute and solute – dislocations interactions. Some alloying elements such as Ni and Si have nearly no effect on DSA [22]. Indeed, they do not interact with carbon nor with dislocations. However, some other elements, such as Mn or Cr exhibit a much more complex behavior. On the one hand, they are able to precipitate with C or N in solid solution, and thus reduce DSA sensitivity [30]. On the other hand, a too high addition of these elements might lead to DSA sensitivity. Besides, the HTDC-Peierls glide appears especially beneficial for structural applications, as it suppresses PLC serrations and its deleterious consequences, such as strain localization. This suggests that, for each type of alloying elements, there exists an optimal concentration to reduce DSA sensitivity (i.e. precipitate C or N in solid solution), without inducing a "substitutional-controlled" DSA. Moreover, this optimal content is a function of the applied strain rate and microstructural features such as the dislocation density.

## 5. Conclusion

We studied the DSA sensitivity in two Fe-Mn-C systems, a model Fe-3Mn-C alloy in the first part and an industrial as-welded C-Mn steel in the present work. By correlating macroscopic tensile tests with in-situ TEM tensile tests, we directly observe the changes in dislocation dynamics at the nano scale, responsible for DSA at the macro scale. We directly show that PLC serrations are caused by the appearance of

dislocation bursts in the material. The occurrence of HTDC-Peierls glide caused by C segregation on screw dislocation cores suppresses PLC serrations, probably because it homogenizes plastic flow.

The observations of HTDC-Peierls glide, the extraction of the activation energies for screw dislocations glide as well as APT characterization all indicate that carbon and nitrogen are responsible for DSA in the industrial C-Mn weld. Contrary to the model Fe-3Mn-C, manganese does not influence DSA in this industrial steel. We suggest that the reduced Mn content does not allow for a sufficient anchoring of mixed dislocations, which are thus able to unlock and re-arrange as screw dislocations. Plastic deformation becomes controlled by the HTDC-Peierls glide of screw dislocations, the latter becoming prominent in the microstructure. As substitutional atoms do not interact with screw dislocations, there is no Mn-induced DSA in the industrial weld. In some aspects, HTDC-Peierls glide prevents high temperature DSA sensitivity. It can then be concluded that manganese influence on DSA is closely linked to its concentration in the alloy. A 3 % Mn addition is sufficient to induce DSA and have deleterious effects on mechanical properties, but a 1,3 % Mn addition does not induce DSA and is already sufficient to partly precipitate C and N in solid solution, thus reducing DSA sensitivity. Additional work is needed to determine the optimum Mn concentration.

The combined analyses of the model Fe-3Mn-C alloy and the industrial weld are used to discuss the effect of the chemical composition of the steel on DSA. The manifestations of DSA at the macroscopic scale are influenced by the interstitial atoms content (i.e. the occurrence of HTDC-Peierls glide), the nature of added substitutional atoms (do they interact with dislocations?) and the content of substitutional atoms (sufficient to hinder HTDC-Peierls glide?). Further research is required to consolidate these conclusions, as well as determining the effect of the applied strain rate on the temperature range of DSA.

## Data availability

Videos of the in-situ TEM tensile tests are available as supplementary materials.

## CRediT authorship contribution statement

W. Mottay: Writing – review & editing, Writing – original draft, Visualization, Methodology, Investigation, Conceptualization. D. Cail-lard: Writing – review & editing, Writing – original draft, Validation, Investigation, Data curation. H. Idrissi: Writing – review & editing, Validation, Methodology, Conceptualization. P. Maugis: Writing – review & editing, Validation, Supervision, Project administration, Conceptualization. A. Portavoce: Writing – review & editing, Validation, Supervision, Resources, Funding acquisition, Conceptualization. F. Roch: Writing – review & editing, Supervision, Resources, Funding acquisition. C. Perrin-Pellegrino: Writing – review & editing, Supervision, Resources, Project administration, Methodology, Conceptualization. K. Hoummada: Writing – review & editing, Validation, Supervision, Resources, Project administration, Funding acquisition, Conceptualization.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgements

This project was supported by the National Association for Research and Technology, France (ANRT Project n° 2020/0883 and ANRT Project n° 2022/0152). This work was also supported by the European Union through the Program PIIEC ME/CT (project SCIRIUS, No. DOS0217953).
H. Idrissi is mandated by the Belgian National Fund for Scientific Research (FSR-FNRS). The authors acknowledge financial support from the CNRS-CEA "METSA" French network (FR CNRS 3507) on the platforms CEMES and PACA.

## Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.actamat.2025.121576.

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