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# Post-process lasering improves strength-ductility tradeoff and fatigue limit of additively manufactured stainless steels

Juan Guillermo Santos Macías, Kewei Chen, Alexandre Tanguy, Maxime

Vallet, Louis Cornet, Vincent Michel, Manas Vijay Upadhyay

► To cite this version:

Juan Guillermo Santos Macías, Kewei Chen, Alexandre Tanguy, Maxime Vallet, Louis Cornet, et al. Post-process lasering improves strength-ductility tradeoff and fatigue limit of additively manufactured stainless steels. 2024. hal-04530203

HAL Id: hal-04530203

https://hal.science/hal-04530203

Preprint submitted on 3 Apr 2024

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# High-vacuum laser treatments enhance strength, ductility and fatigue limit of additively manufactured stainless steel

Juan Guillermo Santos Macías, Kewei Chen, Alexandre Tanguy, Nathalie Isac, Maxime Vallet, Louis Cornet, Vincent Michel, Manas Vijay Upadhyay

► To cite this version:

Juan Guillermo Santos Macías, Kewei Chen, Alexandre Tanguy, Nathalie Isac, Maxime Vallet, et al. High-vacuum laser treatments enhance strength, ductility and fatigue limit of additively manufactured stainless steel. 2025. hal-04530203v3

HAL Id: hal-04530203

https://hal.science/hal-04530203v3

Preprint submitted on 4 Apr 2025

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
recherche français ou étrangers, des laboratoires
publics ou privés.

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

Distributed under a Creative Commons Attribution 4.0 International License
# Post-process lasering improves strength-ductility tradeoff and fatigue limit of additively manufactured stainless steels

Juan Guillermo Santos Macías¹,*, Kewei Chen¹, Alexandre Tanguy¹, Maxime Vallet²,³, Louis Cornet², Vincent Michel⁴, Manas Vijay Upadhyay¹,*

¹ Laboratoire de Mécanique des Solides (LMS), CNRS, École Polytechnique, Institut Polytechnique de Paris, Route de Saclay, 91120 Palaiseau, France

² Université Paris-Saclay, CentraleSupélec, ENS Paris-Saclay, CNRS, LMPS - Laboratoire de Mécanique Paris-Saclay, 91190 Gif-sur-Yvette, France

³ Université Paris-Saclay, CentraleSupélec, CNRS, Laboratoire SPMS - Structures, Propriétés et Modélisation des Solides, 91190 Gif-sur-Yvette, France

⁴ Laboratoire de Procédés et Ingénierie en Mécanique et Matériaux (PIMM), CNRS UMR 8006, Arts et Metiers Institute of Technology, HESAM University, 151 Boulevard de l'Hôpital, 75013 Paris, France

* First corresponding author email: manas.upadhyay@polytechnique.edu,

+ Second corresponding author email: guillermo.santos@imdea.org

## Abstract

Although stainless steels fabricated via additive manufacturing exhibit higher strength in comparison to their conventionally manufactured counterparts, they often suffer from lower ductility and poorer fatigue limit. In this work, post-process micron-sized lasering is used to re-engineer additively manufactured stainless steel microstructures close to the sample surface in order to improve their overall performance. The key idea is to use smaller spot sizes and faster scan speeds during post-process lasering compared to the fabrication process. Using a novel coupling between a continuous-wave laser and a scanning electron microscope, single line scans were performed to optimise the lasering parameters. The optimised parameters were then used to treat the surface of the entire sample. The result is an intragranular structure refinement near the surface, exhibiting nearly an order of magnitude reduction in microsegregation cell and dislocation structure size, and increase in dislocation density. The ensuing microstructure demonstrates a higher yield strength without ductility loss and significant enhancement of the fatigue limit due to reduced surface roughness and intragranular structure refinement. The proposed technique has tremendous potential to improve the mechanical response of other alloys fabricated via additive manufacturing or any manufacturing technique.

## Keywords

Laser treatment, mechanical behaviour, electron microscopy, fatigue, 3D printing
# High-vacuum laser treatments enhance strength, ductility and fatigue limit of additively manufactured stainless steel

Juan Guillermo Santos Macías¹,*, Kewei Chen¹, Alexandre Tanguy¹, Nathalie Isac¹, Maxime Vallet²,³, Louis Cornet², Vincent Michel⁴, Manas Vijay Upadhyay¹,*

¹ Laboratoire de Mécanique des Solides (LMS), CNRS, École Polytechnique, Institut Polytechnique de Paris, Route de Saclay, 91120 Palaiseau, France

² Université Paris-Saclay, CentraleSupélec, ENS Paris-Saclay, CNRS, LMPS - Laboratoire de Mécanique Paris-Saclay, 91190 Gif-sur-Yvette, France

³ Université Paris-Saclay, CentraleSupélec, CNRS, Laboratoire SPMS - Structures, Propriétés et Modélisation des Solides, 91190 Gif-sur-Yvette, France

⁴ Laboratoire de Procédés et Ingénierie en Mécanique et Matériaux (PIMM), CNRS UMR 8006, Arts et Métiers Institute of Technology, HESAM University, 151 Boulevard de l'Hôpital, 75013 Paris, France

* Corresponding author email:

guillermo.santos@imdea.org
manas.upadhyay@polytechnique.edu

## Abstract

Post-process laser scanning under high vacuum is proposed as a non-isothermal heat treatment to simultaneously refine the intragranular microstructure near the surface and reduce surface roughness, while also preventing oxidation, in order to enhance the overall mechanical response of an alloy. This treatment is performed using laser spot sizes and scan speeds that produce higher temperature gradients and faster heating/cooling rates than those encountered during manufacturing. The effectiveness of this approach is demonstrated on laser-based direct energy deposited (LDED) 316L stainless steel using parameters similar to those used in laser-based powder bed fusion (LPBF). High vacuum ($< 0.1$ Pa) lasering is conducted inside a newly integrated continuous-wave laser and scanning electron microscope (CW Laser-SEM). The treatments result in an order-of-magnitude reduction in microsegregation cell and dislocation structure sizes, as well as in the surface roughness of LDED 316L. For a parameter set in which the laser penetrates 14% of the total depth (7% each on the two widest surfaces of dogbone-shaped samples), significant improvements are obtained in yield strength (31.11%), ductility (14.2%), and fatigue limit (25%). The proposed approach has tremendous potential to alter the microstructure and improve the mechanical response of both additively and conventionally manufactured alloys.

## Keywords

Laser treatment, mechanical behaviour, electron microscopy, fatigue, 3D printing, AM 316L

## 1. Introduction

a Current address: Instituto madrileño de estudios avanzados materiales (IMDEA materiales), Calle Eric Kandel 2, 28906 Getafe, Spain
# 1. Introduction

One of the most important goals of research on engineering of stainless steels is to design microstructures that exhibit higher strength, ductility and fatigue limit than currently possible while reducing energy and material consumption. The advent of metal additive manufacturing (AM), has opened up the possibility to design as built (AB) parts with microstructures that present unprecedentedly higher strengths in comparison to their conventionally manufactured counterparts [1]. However, this increase in strength is often accompanied by a decrease in ductility and a poorer fatigue response [2], which is affected by a combination of material strength, surface roughness, defects (porosities or voids) and the residual stress state.

At the origin of the strength-ductility trade-off lies the hierarchical microstructure arising from the highly non-equilibrium processes occurring during AM. Heat-matter interaction induced melt pool dynamics, rapid solidification and solid-state thermal cycling result in a microstructure exhibiting physical and chemical heterogeneities ranging from few tens of nanometres to several hundreds of micrometres [3,4]. The primary contribution to material strength arises from the smallest of these features, which in stainless steels are precipitates, microsegregation cells and dislocation structures [5]. These intragranular features determine the mean free path of plastic deformation accommodating defects such as mobile dislocations. Typically, the smaller the size and higher the density of these features, the higher the strength and lower the ductility [1]. A commonly used approach to improve ductility is annealing, which is an isothermal heat treatment that evolves the metastable AB microstructure towards equilibrium by minimizing the stored energy. However, this process inadvertently causes an increase in feature size and decrease in density, inevitably resulting in a decrease in strength.

Meanwhile, the fatigue response of dense (negligible amount of porosities/voids) AB parts rely on the material strength as well as surface roughness [6]. During AM, unmelted powder particles sinter to the surface and become the dominant contributors to surface roughness of AB parts. Under fatigue loading, failure mainly occurs from surface crack nucleation (if not already present) and propagation in dense alloys. The fatigue response of AB parts can be improved via post-process surface treatments, the most common of which are mechanical in nature e.g., shot peening, polishing, etc., that induce in-plane compressive stresses on the surface and reduce surface roughness.

In this work, a non-isothermal heat treatment for AB parts is proposed in the form of continuous-wave (CW) laser scanning that simultaneously alters the underlying microstructure, specifically significantly refines the intragranular structure, and improves surface quality. The key idea is to perform this laser scanning using spot sizes and scanning speeds that are respectively smaller and faster than those used to build the samples in order to simultaneously improve the surface roughness and alter the intragranular microstructure.

This idea has been motivated by the microstructural differences between AB stainless steels fabricated via two commonly used techniques: CW laser-based powder bed fusion (LPBF) and CW laser-based direct energy deposition (LDED). However, LPBF processes use smaller laser spot sizes and faster scanning speeds than LDED processes. Consequently, stainless steel microstructures fabricated via LPBF exhibit smaller and denser features, and hence higher strength and lower ductility, than the LDED ones [4]. In this work, LDED stainless steel samples were used as base material and their surfaces were laser scanned using spot sizes and scan speeds
One of the most important goals in stainless steel engineering is to design microstructures with superior strength, ductility and fatigue limit while minimizing energy and material consumption. The advent of metal additive manufacturing (AM) has enabled the fabrication of as-built (AB) parts with microstructures that exhibit unprecedented strength in comparison to their conventionally manufactured counterparts [1]. However, this increased strength is often accompanied by reduced ductility and poorer fatigue performance [2], which is affected by a combination of material strength, surface roughness, volumetric defects (porosities or voids) and the residual stress state.

The strength-ductility trade-off in AB alloys arises from the hierarchical microstructure formed due to the highly non-equilibrium processes occurring during AM. The complex heat-matter interactions – melt pool dynamics, rapid solidification and solid-state thermal cycling – generate microstructures exhibiting physical and chemical heterogeneities ranging from few tens of nanometres to several hundreds of micrometres in scale [3,4]. The smallest of these features such as precipitates, microsegregation cells and dislocation structures including low angle grain boundaries [5,6], are the primary contributors to material strength. These intragranular features determine the mean free path for plastic-deformation-accommodating defects, the most common of which are mobile dislocations. Typically, the smaller the size and higher the density of these features, the higher the strength and lower the ductility [1]. A commonly used approach to improve ductility is annealing, which is an isothermal heat treatment that reduces stored energy by evolving the metastable AB microstructure towards equilibrium. However, this process inadvertently causes an increase in feature size (e.g., precipitate and grain sizes) and decrease in feature density (e.g., dislocation and precipitate densities and grain boundary interface-to-volume ratio) [7], inevitably resulting in a trade-off between strength and ductility.

Meanwhile, fatigue resistance in dense (i.e., with negligible porosity) AB components relies primarily on material strength and surface roughness [8]. During AM, unmelted powder particles sinter to the surface, becoming dominant contributors to surface roughness [1]. Under cyclic loading, failure mainly occurs from surface crack nucleation (if not already present) and propagation. Thus, surface roughness negatively affects the fatigue response. Post-process surface treatments, including mechanical methods (e.g., shot peening [9], polishing [10,11], preloading [12], etc.) and isothermal heat treatments (e.g., annealing for 30 minutes at 900°C [13]) can enhance fatigue performance by reducing tensile residual stresses and surface roughness [14].

In this work, we propose a non-isothermal, post-process surface heat treatment in the form of continuous-wave (CW) laser scanning under high vacuum (< 0.1 Pa). The key strategy and novelty of this approach is two-fold: (i) Optimized laser parameters – using smaller laser spot sizes and higher scanning speeds than those used during fabrication to obtain large temperature gradients and faster heating/cooling rates, and (ii) High-vacuum processing – preventing surface oxidation, which can embrittle the surface and negatively impact fatigue performance. The strategy is expected to simultaneously refine the intragranular microstructure near the surface and reduce surface roughness, thereby enhancing strength, ductility, and the fatigue limit of an alloy. We demonstrate its effectiveness for an additively manufactured 316L stainless steel, a reference alloy in AM research and applications. However, the proposed approach can be applied to any AM or conventionally manufactured alloy.
corresponding to LPBF. Note that while post-process laser-based surface treatments have been to
polish surfaces of AM alloys [7,8] i.e., to reduce their surface roughness, it has not been used to
simultaneously refine the intragranular structure and improve surface roughness, both of which are
important to improve the overall mechanical response of a material.

Identifying appropriate laser scanning parameters to obtain a desired microstructure requires performing a parametric study in conjunction with microstructure characterisation before and after lasering. To facilitate such studies, a recently designed novel coupling between a CW laser and an environmental scanning electron microscope (CW laser-SEM (Fig. 1a,b)) [9] is used. While there are precedents for coupling between pulsed lasers and SEM [10–13], none exist for a CW laser-SEM. An important advantage of this coupling is its ability to perform laser scanning under secondary vacuum (pressure < 10⁻⁴ mbar) or under a controlled environment (e.g., inert gas environment) inside the environmental SEM. It prevents surface oxidation during lasering and facilitates a one-to-one comparison before and after scanning without additional surface preparation using different detectors available inside the SEM. The CW laser-SEM device is used to first perform a parametric study to optimise the lasering parameters.

<figure><img src="image_3.png" /><figcaption>**Fig. 1:** CW laser-SEM coupling; illustration of the CW laser-SEM coupling (a) and image of the experimental setup (b). (c) Plot quantifying the laser spot sizes (1.7 times the full width at half maximum) in the *x* and *y* directions (in (a)) as a function of optical *z* (focusing distance controller).</figcaption></figure>
The first part of the proposed strategy, i.e. to use smaller laser spot sizes and faster scan speeds than those used during fabrication, has been motivated by the microstructural differences between AB alloys fabricated via two commonly used techniques: CW laser-based powder bed fusion (LPBF) and CW laser-based direct energy deposition (LDED). LPBF employs smaller laser spot sizes and faster scanning speeds than LDED. Consequently, LPBF microstructures exhibit smaller and denser features, and hence higher strength and lower ductility, than the LDED ones [1]. In this work, LDED stainless steel samples are used as the base material and their surfaces are subjected to post-process laser scanning using LPBF-like spot sizes and scan speeds.

The second part of the proposed strategy, i.e. to perform laser scanning inside high vacuum, is important because a surface oxide layer can degrade mechanical performance, particularly under cyclic loading; oxide layers can embrittle surfaces and accelerate crack nucleation under cyclic loading making them detrimental to fatigue life, especially at elevated temperatures. To achieve a high-vacuum environment, laser scanning is performed inside a newly designed CW laser-scanning electron microscope (CW laser-SEM) system (Fig. 1a, b) [15]. While previous studies have coupled pulsed lasers with SEM [16–18], this is the first CW laser-SEM system, developed by some of the co-authors [15]. This unique setup allows not only to perform oxidation-free laser treatments but also enables direct, one-to-one microstructure comparisons before and after lasering without intermediate mechanical polishing. It has already been successfully used to provide microstructural input and validation for a thermomechanical polycrystal model [19] and a coupled computational thermo-fluid dynamics and phase-field solidification model [20] applied to 316L.

Most of the previous studies on laser surface treatments of AM 316L [21–23] have focused on surface roughness reduction for improved corrosion resistance and hardness. The only prior work examining enhancement of mechanical response is by Chen et al. [24], which reported slight improvements in tensile properties of LPBF 316L when subjected to low-vacuum laser treatment (10 Pa) with LPBF-like parameters. However, while their approach reduced surface roughness, it did not significantly refine the microstructure to obtain an overall improvement in strength-ductility tradeoff and fatigue limit. In contrast, we demonstrate that our high-vacuum laser treatment with LPBF-type parameters significantly improves the strength, ductility and fatigue limit of LDED 316L. The enhanced surface quality is also expected to improve the corrosion and wear behaviour.

The remainder of the article is divided as follows: the experimental procedure and methods, including a detailed presentation of the CW laser-SEM, are given in section 2. Section 3 presents results and discussion around the parametric study, and tensile and fatigue testing. Finally, the conclusions of the study are presented.

## 2. Material and methods

### 2.1 Material and additive manufacturing via LDED

Wrought 316L is used to manufacture 316L powder via inert gas atomisation by Oerlikon GMBH (Germany). Granulometry reveals powder particle size range in between 45 µm and 106 µm ($d50_{316L} = 57.86$ µm).

Thin walls with dimensions of 100 (x) × 12 (y) × 0.8 (z) mm³ (for tensile testing) and 160 (x) × 30 (y) × 0.8 (z) mm³ (for fatigue testing) were printed using a BeAM Modulo 400 LDED
The article is divided as follows: the experimental procedure and methods, including the CW laser-SEM, are described in section 2. Section 3 presents results and discussion around the parametric study, and tensile and fatigue testing. Finally, the conclusions of the study are presented.

## 2. Material and methods

### 2.1 Material

The materials studied in this work are two 316L stainless steels with different Si content. Henceforth, they are called 316L and 316L-Si. Wrought alloys are used to manufacture 316L and 316L-Si powders via inert gas atomisation by Oerlikon GMBH (Germany). The 316L powder has the following composition in weight percent (wt. %): Fe – balance, Cr – 17.34, Ni – 12.55, Mo – 2.34, Mn – 1.4, Si – 0.49, N – 0.08, Cu – 0.04, O – 0.03, P – 0.01, C < 0.1 and S < 0.1. Meanwhile, 316L-Si has the following composition in wt. %: Fe – balance, Cr – 17, Ni – 12, Mo – 2.5, Mn – 1.0, Si – 2.3, P < 0.1, C – 0.03 and S < 0.1. Granulometry reveals powder particle size range in between 45 µm and 90 µm.

### 2.2 Additive manufacturing via LDED

Thin walls with dimensions of 100 (x) × 12 (y) × 0.8 (z) mm³ (for tensile testing) and 160 (x) × 30 (y) × 0.8 (z) mm³ (for fatigue testing) were printed using a BeAM Modulo 400 LDED machine equipped with a 500W YLR-fibre laser. The printing process employed a bidirectional scanning strategy with a laser power of 250 W, scanning speed of 2000 mm/min, a powder flow rate of 6.8 g/min, a vertical displacement of focusing head of 0.2 mm after each layer deposition and a spot size of 0.7 mm.

### 2.3 CW laser-SEM coupling and sample polishing

The key equipment used in this work is the CW laser-SEM [9]. It is comprised of a coupling between a 1070 nm wavelength SPI QUBE 200 W air-cooled fibre laser (Industrial Laser Systems, France) and a FEI Quanta 600 environmental-SEM for characterisation (see Fig. 1a,b). The laser power can be varied from 9 to 209 W, the scanning speed from 0.01 to 20000 mm/s and the spot size from 45 to 500 µm with the help of the incorporated Scanlab varioscan<sub>de</sub>20i type 133 (optical z). Virtually full 2D scanning liberty in a relatively large area (100 × 100 mm²) is possible with the equipped Scanlab intelliscan III 20 scanner. The laser spot sizes as a function of optical z (see Fig. 1c) and the power have been calibrated using the femto easy BP 13.9 beam profiler and gentec UP55N-40S-H9-D0 calorimeter.

Some AB 316L and 316L-Si thin-walls (excluding those used for fatigue testing) were polished using 320, 600, 1200, 2400 and 4000 grit SiC paper followed by 3 µm and 1 µm diamond paste polishing. The final polishing step involved active oxide polishing (OPA) suspension polishing for two and a half minutes. Next, the samples (except the ones used for fatigue) were subjected to electrochemical etching using 10% vol. oxalic acid aqueous solution to reveal the microsegregations during backscatter electron (BSE) imaging. Electron backscatter diffraction
machine equipped with a 500W YLR-fibre laser. The printing process employed a bidirectional scanning strategy with a laser power of 250 W, scanning speed of 2000 mm/min, a powder flow rate of 6.8 g/min, a vertical displacement of printing head of 0.2 mm after each layer deposition and a spot size of 0.7 mm. The volumetric energy density of manufacturing, calculated as the laser power divided by the product of the scanning speed, vertical displacement of printing head and laser spot size, is 53.57 J/mm³. The chemical composition of the AB material measured at FiLAB, France using ICP spectroscopy and total organic carbon analysis is given in table 1 [25].

<table><caption>Table 1. Chemical composition (wt.%) with standard deviation (SD) of LDED 316L from [25]</caption><thead><tr><th></th><th>Fe</th><th>Cr</th><th>Ni</th><th>Mo</th><th>Mn</th><th>Si</th><th>N</th><th>Cu</th><th>O</th><th>P</th><th>C</th><th>S</th><th>Co</th></tr></thead><tbody><tr><td>316L</td><td>Bal.</td><td>18.7</td><td>13.3</td><td>2.4</td><td>1.4</td><td>0.73</td><td>0.083</td><td>0.063</td><td>0.024</td><td>0.012</td><td>0.018</td><td>0.005</td><td>0.026</td></tr><tr><td>SD</td><td></td><td>0.438</td><td>0.312</td><td>0.01</td><td>0.005</td><td>0.005</td><td>0.001</td><td>0.001</td><td>0.000</td><td>0.000</td><td>0.000</td><td>0.000</td><td>0.000</td></tr></tbody></table>

As shall be seen in the results section 3, post process laser scans of interest are those that result in conduction mode melting with much smaller melt pools than those occurring during AM. Therefore, as well acknowledged in the literature (see for example [20,26,27]), negligible variations in chemical composition are expected due to post-process laser scanning.

## 2.2 CW laser-SEM coupling and scanning electron microscopy

The CW laser-SEM is a coupling between a 1070 nm wavelength SPI QUBE 200 W air-cooled fibre laser (Industrial Laser Systems, France) and a FEI Quanta 600 environmental-SEM for characterisation (see Fig. 1a, b). The laser power can be varied from 9 to 209 W, the scanning speed from 0.01 to 20000 mm/s and the spot size from 45 to 500 µm with the help of the incorporated Scanlab varioscan<sub>de</sub>20i type 133 (optical z). Virtually full 2D scanning liberty in a relatively large area (100 × 100 mm²) is possible with the equipped Scanlab intelliscan III 20 scanner. The laser spot sizes as a function of optical z (see Fig. 1c) and the power have been calibrated using the femto easy BP 13.9 beam profiler and gentec UP55N-40S-H9-D0 calorimeter. As seen in Fig. 1a, the laser beam is oriented horizontally because the electron source blocks the vertical alignment of the laser. The samples have to be placed in such a way that laser scanning is performed on the surface normal to the incident laser beam. Such a placement could raise a concern regarding the effect of gravity on melt pool dynamics. However, recent computational thermo-fluid dynamics simulations [20] have shown that this effect is negligible, at least for the case of 316L for a wide range of scanning parameters.

Single line scan parametric studies are conducted in sections 3.1 - 3.3 and on the basis of these studies, optimized laser scanning parameters are identified to laser treat the entire surfaces. Since the parametric studies themselves form an important part of the results in this work, the obtained parameters are directly presented in the results and discussion section 3.

SEM analysis is first performed on some mechanically polished AB samples. Then, the same samples are subjected to laser scanning inside the CW Laser-SEM and restudied via SEM without any additional mechanical polishing; some of these samples are used for dogbone testing as described in section 2.3 but none of these mechanically polished samples are used for fatigue testing described in 2.6. Some AB thin-walls (excluding those used for fatigue testing) were polished using 320, 600, 1200, 2400 and 4000 grit SiC paper followed by 3 µm and 1 µm diamond paste polishing. The final polishing step involved active oxide polishing (OPA) suspension polishing for two and a half minutes. Next, the samples (except the ones used for fatigue testing)
(EBSD) was performed with 1 µm step size using a Symmetry® detector (Oxford Instruments).
BSE imaging is performed using an angular back scatter detector from Thermofisher.

In addition to SEM, OM is performed using the Keyence VHX-100 microscope with a VH-
Z100R objective.

**2.4 Uniaxial tensile testing**

Rectangular plates with dimensions of 35 × 12 × 0.8 mm³ were machined from the 316L and 316L-Si walls using HERMLE U 630 T, perpendicular to the DED build direction. After that, they were polished using the procedure described in section 2.3. Some of the polished plates were clamped in a small region (~7 × 3 mm²) on one side and subjected to laser scanning. Dogbone specimens were extracted from AB and lasered 316L and 316L-Si samples using Hermle U 630 T; sample sides were mirror polished before machining (and before lasering for lasered samples) to avoid surface roughness influence on the mechanical testing. The gauge region had the dimensions 8 × 4 × 0.5 mm³ ensuring a minimum length to width ratio of 2 that corresponds to the JIS Z 2241 standard; respecting this condition is sufficient for a comparative study between AB and lasered samples. Two specimens per condition were tested on a microtest stage at a (quasi-static) loading rate of 2 µm/s. The use of a Sill Optics correctal camera and digital image correlation helped to obtain the mechanical behaviour of the samples. All samples failed in the gauge region. After failure, fractography was performed on the specimens using SE imaging.

**2.5 Scanning transmission electron microscopy and lamellae preparation**

Thin-film lamellae were extracted inside a FEI Helios Nanolab 660 dualbeam SEM microscope equipped with a focused ion beam. The lamellae have a thickness of ~100 nm. They were studied in the Titan³ G2 TEM microscope, which is equipped with a Cs-probe corrector, STEM detectors and a SuperX detector. It also has the capability to perform energy dispersive X-ray spectroscopy in order to characterise the chemical composition of the lamellae. Spectroscopy was performed with an acquisition time of 1 hour under operating conditions of 20000 counts per second and 300 keV. The spectra were deconvoluted and quantified as described in [14].

**2.6 Residual stress measurement – X-ray diffraction**

The sin²Ψ X-ray diffraction method was used to determine the stresses in the austenitic phase. This method is based on the measurement of Bragg peak positions for a given *hkl* reflection and for various directions of the scattering vector with respect to sample. Due to a significant absorption of X-rays by metallic alloys, the method allows measurements on surfaces. In the case of austenitic steels, the average depth penetrated by X-rays at Ψ = 0 is ~7.2 µm; the average measurement depth is defined as the depth at which 67% of the incoming intensity is reflected. Hence, the surface residual stress measures are average measures over this depth.

The residual stresses were evaluated on the surface of both sides of the specimen using the Ψ tilt method and an *in situ* diffractometer X-RAYBOT (manufactured by MRX, France). To statistically optimise the number of grains in diffraction conditions, an oscillation of +/- 5° on Φ and Ψ angles was applied during the peak's acquisitions. A collimator with a circular irradiated area with a diameter of 2.5 mm was used. The diffraction conditions were chosen in accordance with the NF EN 15305-2009 standard: Mn-Kα radiation at 0.2290 nm wavelength, voltage 20 kV,
were subjected to electrochemical etching using 10% vol. oxalic acid aqueous solution to reveal
the microsegregations during backscatter electron (BSE) imaging. Electron backscatter diffraction
(EBSD) was performed with 0.5, 0.1 and 1 µm step size for Figs. 3, 4 and 6b, respectively, using
a Symmetry® detector (Oxford Instruments). BSE imaging is performed using an angular back
scatter detector from Thermofisher.

<figure><img src="image_2.png" /><figcaption>Fig. 1. CW laser-SEM coupling; illustration of the CW laser-SEM coupling (a) and image of the experimental setup (b). (c) Plot quantifying the laser spot sizes (1.7 times the full width at half maximum) in the x and y directions (in (a)) as a function of optical z (focusing distance controller).</figcaption></figure>

In addition to SEM, optical microscopy (OM) is performed using the Keyence VHX-100 microscope with a VH-Z100R objective.

**2.3 Uniaxial tensile testing**

All uniaxial tensile tests are conducted on mechanically polished AB and lasered samples.
To that end, rectangular plates with dimensions of 35 × 12 × 0.8 mm³ were machined from AB
and lasered 316L walls using Hermle U 630 T, perpendicular to the DED build direction. After
current 1 mA, 311 (hkl) lattice plane, diffraction angle $2\theta = 152^{\circ}$, 21 $\Psi$ angles in the range [-37.27$^{\circ}$, 39.23$^{\circ}$] and $\pm 5^{\circ}$ oscillations in $\Phi/\Psi$.

## 2.7 Fatigue testing

Rectangular plates with dimensions of 70 × 14 × 0.8 mm³ were machined from the 316L printed walls using Hermle U 630 T, perpendicular to the DED build direction. After that, the plates were clamped into a customised fixture and laser scanned under secondary vacuum with scanning speed of 100 mm/s, power of 70 W, spot size of 60 µm and scan overlap of 50%, within the CW Laser-SEM chamber. As a result, a 60 × 12 mm² area in the centre of the plates was laser treated. Subsequently, fatigue samples with 8 × 4 × 0.8 mm³ gauge volume and 32 mm blending fillet (adhering to the ASTM E466-21 standard) were machined from the processed area using Hermle U 630 T. Simultaneously, the AB fatigue samples were directly machined from the unprocessed rectangular plates.

Constant amplitude uniaxial tensile fatigue test was conducted on both AB and L70 samples, using a 10 kN load cell on a Material Testing System (MTS 810), employing a sinusoidal stress waveform with a stress ratio (R) 0.1 and a constant frequency of 40 Hz, at room temperature. The loading direction was set perpendicular to the DED build direction. In order to determine the fatigue strength ($\sigma_f$), the staircase method was employed. Here, the starting stresses ($\sigma_{max}$) are chosen below yield strength of AB and L70 316L samples, at 275 and 325 MPa. Note that, since the difference between number of cycles corresponding to 300 MPa and 275 MPa is large, a median stress of 287.5 MPa is used for the fatigue testing. The testing was interrupted at $3 \times 10^6$ cycles (run-outs) for samples that do not fail until that point. Fracture in all the failed samples occurred in the gauge region.

## 2.8 Surface roughness measurement

Surface roughness measurements for unpolished AB 316L and unpolished + lasered 316L samples were conducted following the ISO 25178 norm using Bruker Alicona optical system with a ring light illuminator to improve image quality. Surface texture images were captured using 20x objective lens with vertical resolution less than 13 nm, Lambda C ($L_c$) filter of 80 µm, and stitched together. MeasureSuite of Alicona was used to obtain $S_a$.

# 3 Results and discussion

The reference AB samples used are 316L and 316L-Si stainless steel thin walls manufactured via LDED on hot-rolled 316L substrates using a single-pass-per-layer bidirectional scanning strategy with P 250 W, $v$ 33.33 mm/s and $s$ 0.7 mm. These parameters result in a surface energy density $E_s = P/(v \cdot s) = 10.71$ J/mm². Electron backscatter diffraction (EBSD) imaging after mirror polishing and etching the AB sample reveals a weak texture and a bimodal grain size distribution due to long columnar grains along layer height with a tilt along the printing direction, and smaller equiaxed grains in between layers (Fig. 2).
that, they were mechanically polished using the procedure described in section 2.3. Some of the polished plates were clamped in a small region (~7 × 3 mm²) on one side and subjected to laser scanning. Dogbone specimens were extracted from AB and lasered rectangular plates also using Hermle U 630 T (see Fig. 2); sample sides were mirror polished before machining (and before lasering for lasered samples) to avoid surface roughness influence on the mechanical testing. The gauge region had the dimensions 8 × 4 × 0.5 mm³ ensuring a minimum length to width ratio of 2 that corresponds to the JIS Z 2241 standard; respecting this condition is sufficient for a comparative study between AB and lasered samples. Three specimens per condition were tested in displacement control employing a microtesting stage at a (quasi-static) displacement rate of 2 µm/s. The use of a Sill Optics correctal camera and digital image correlation helped to obtain the mechanical behaviour of the samples. All samples failed in the gauge region. After failure, fractography was performed on the specimens using secondary electron (SE) imaging.

<figure><img src="image_3.png" /><figcaption>Fig. 2. Schematic of wall pieces with laser scanned region (depicted with red lines) and tensile (top) and fatigue (bottom) sample geometries. The sample thickness values are preceded by a “t”.</figcaption></figure>

## 2.4 Scanning transmission electron microscopy and lamellae preparation

Thin-film lamellae (~100 nm thick) were extracted from AB and lasered samples used for uniaxial tensile testing. A FEI Helios Nanolab 660 dualbeam SEM microscope equipped with a focused ion beam was used to extract these lamellae. They were studied in the Titan³ G2 TEM microscope, which is equipped with a Cs-probe corrector, scanning transmission electron microscopy (STEM) detectors and a SuperX detector. It also has the capability to perform energy dispersive X-ray spectroscopy (EDS) in order to characterise the chemical composition of the lamellae. Spectroscopy was performed with an acquisition time of 1 hour under operating conditions of 20000 counts per second and 300 keV. The spectra were deconvoluted and quantified as described in [28].

## 2.5 Residual stress measurement – X-ray diffraction
<figure><img src="image_4.png" /><figcaption>**Fig. 2:** EBSD IPF z images of the plane formed by building (−y) and printing (x) directions of the polished AB LDED 316L thin-wall before and after laser scanning using s 60 µm, P 24 W and varying v (first parametric study). The noise (black dots) in the AB region of the EBSD maps is a consequence of surface undulations caused by chemical etching to facilitate BSE imaging of Cr-Mo microsegregations in Fig. 3.</figcaption></figure>
The sin²Ψ X-ray diffraction method was used to determine the stresses in the austenitic phase. This method is based on the measurement of Bragg peak positions for a given hkl reflection and for various directions of the scattering vector with respect to sample. Due to a significant absorption of X-rays by metals, the method only allows surface measurements. In the case of austenitic stainless steels, the average depth penetrated by X-rays at Ψ = 0 is ~7.2 µm; the average measurement depth is defined as the depth at which 67% of the incoming intensity is reflected. Hence, the surface residual stress measures are average measures over this depth.

The residual stresses were evaluated on the surface of both sides of the specimen using the Ψ tilt method and an in situ diffractometer X-RAYBOT (manufactured by MRX, France). To statistically optimise the number of grains in diffraction conditions, an oscillation of +/- 5° on Φ and Ψ angles was applied during the peak's acquisitions. A collimator with a circular irradiated area with a diameter of 2.5 mm was used. The diffraction conditions were chosen in accordance with the NF EN 15305-2009 standard: Mn-Kα radiation at 0.2290 nm wavelength, voltage 20 kV, current 1 mA, 311 (hkl) lattice plane, diffraction angle 2θ = 152°, 21 Ψ angles in the range [-37.27°, 39.23°] and ±5° oscillations in Φ/Ψ.

## 2.6 Fatigue testing

Unlike the dogbone tensile tests, all fatigue tests were performed on AB and lasered samples that were not mechanically polished. Rectangular plates with dimensions of 70 × 14 × 0.8 mm³ were machined from AB walls using Hermle U 630 T, perpendicular to the LDED build direction (Fig. 2 bottom). After that, the plates were clamped into a customised fixture and laser scanned under high vacuum with scanning speed of 100 mm/s, power of 70 W, spot size of 60 µm and scan overlap of 50%, within the CW laser-SEM chamber. As a result, a 60 × 12 mm² area in the centre of the plates was laser treated. Subsequently, fatigue samples with 8 × 4 × 0.8 mm³ gauge volume and 32 mm blending fillet (adhering to the ASTM E466-21 standard) were machined from the processed area using Hermle U 630 T. Simultaneously, the AB fatigue samples were directly machined from the unprocessed rectangular plates.

Constant amplitude uniaxial tensile fatigue test was conducted on both AB and lasered samples, using a 10 kN load cell on a Material Testing System (MTS 810), employing a sinusoidal stress waveform with a stress ratio (R) 0.1 and a constant frequency of 40 Hz, at room temperature. The loading direction was set perpendicular to the DED build direction. In order to determine the fatigue strength (σf), the staircase method was employed. Here, the starting stresses (σmax) are chosen below yield strength of AB and lasered samples, at 275 and 325 MPa. Note that since the difference between the number of cycles corresponding to 300 MPa and 275 MPa is large, a median stress of 287.5 MPa is used for the fatigue testing. The testing was interrupted at 3 × 10⁶ cycles (run-outs) for samples that do not fail until that point. Fracture in all the failed samples occurred in the gauge region.

## 2.7 Surface roughness measurement

Surface roughness measurements for unpolished AB and lasered samples were conducted following the ISO 25178 norm using Bruker Alicona optical system with a ring light illuminator to improve image quality. Surface texture images were captured using 20x objective lens with
## 3.1 Parametric post-process laser scanning studies

Two parametric laser scanning studies were first conducted with 316L inside the CW Laser-SEM under secondary vacuum. For both studies, a constant spot size $s$ 60 µm was used; it is an order of magnitude smaller than the one used during LDED.

<figure><img src="image_5.png" /><figcaption>**Fig. 3:** BSE and EBSD IPF z images of the cross-sections across laser tracks studied in Fig. 2. The noise (black dots) in the AB region of the EBSD maps is a consequence of surface undulations caused by chemical etching to facilitate BSE imaging of Cr-Mo microsegregations. BSE and EBSD images share the same scale.</figcaption></figure>

### 3.1.1 Varying scan speed, constant laser power and laser spot size

The first parametric study was performed to understand the role of varying $E_s$ by varying $v$ on microstructure alteration while keeping $P$, $s$ and scan strategy constant. $P$ was set at 10% of the maximum laser power, 24 W, which is the smallest stable value for the laser used. $s$ was set to 60 µm, which is in the range of spot sizes used during LPBF. The scan strategy was a single line
vertical resolution less than 13 nm, Lambda C ($L_c$) filter of 80 µm, and stitched together.
MeasureSuite of Alicona was used to obtain $S_α$.

### 3 Results and discussion

<figure><img src="image_4.png" /></figure>
scan along the y direction (Figs. 2,3). The laser scans were parallel to each other and separated by a distance of 0.8 mm from their neighbours as seen in Fig. 2 to avoid any influence of one scan on the other. The $v$ were chosen to correspond to five different $E_s$ lower than the one used to manufacture the samples; since the CW laser-SEM operates under vacuum and does not use any powder feedstock, there are no convection-related heat losses or powder-related laser reflection losses; hence, a lower $E_s$ should be needed than during the LMD process to melt the material. The five different $E_s$ ($v$) were: 8 J/mm² (50 mm/s), 4 J/mm² (100 mm/s), 1.6 J/mm² (250 mm/s), 0.8 J/mm² (500 mm/s) and 0.4 J/mm² (1000 mm/s). Laser scanning at the four highest $E_s$ (four lowest $v$) results in melting, solidification and rapid cooling of the material, whereas laser scanning at the lowest $E_s$ (highest $v$) results in heating and cooling without any melting (Figs. 2,3).

The melt pool sizes of the four remelted and solidified tracks (Fig. 3) expectedly decrease with decreasing $E_s$ (increasing $v$) and their shapes reveal that conduction mode melting has occurred. The Cr and Mo microsegregation cells (a signature of rapid solidification in 316LSS [15,16]) in the lasered regions (Fig. 3) are at least an order of magnitude smaller than the ones in the AB material. Their size does not significantly decrease with increasing $v$, which was confirmed by comparing their sizes in the 50 mm/s and 500 mm/s tracks. The cell size difference between AB and lasered samples strongly indicates that solidification rates occurring during laser scanning are faster than those occurring during the LMD process, and they increase with increasing $v$. For the slowest laser scan (50 mm/s), the clearly visible microsegregation cells have an internal diameter that is ~300 nm on average. Assuming this value to be the primary cellular arm spacing ($\lambda$) during solidification, and using the empirical formula $\lambda = 80/\dot{T}^{1/3}$ for stainless steels [17], the deduced solidification rate ($\dot{T}$) in the slowest laser scanned region is ~1.9 × 10⁷ K/s; in comparison, the microsegregation cells in the AB 316L have a size of ~2.2 µm and the corresponding $\dot{T}$ is slower by three orders of magnitude. At such large solidification rates, and corresponding high temperature gradients, the $\gamma$-austenite grains epitaxially grow from the melt pool boundary via planar or cellular solidification mode and bend along the laser scanning direction [18], which is evidenced in the EBSD images in Figs. 2,3.

Following solidification, internal and sample constraints result in the formation of thermal stresses in the heat-affected solid. These stresses are sufficiently high to cause some of the grains to undergo plastic deformation e.g., via dislocation dynamics, which can be deduced from the presence of misorientation bands observed in some grains in the remelted zone (Fig. 2).

After lasering, a remarkably cleaner surface is obtained along all tracks. The cleaner surfaces are primarily a consequence of remelting and subsequent refinement of the microsegregation structure. A refined surface is also visible for the fastest laser scan, indicating that the surface could have slightly remelted even though it is not discernible in the BSE and EBSD images of the cross-section (Fig. 3).

### 3.1.2 Varying laser power, constant scan velocity and laser spot size

The next parametric study was performed to investigate laser penetration depth as a function of $E_s$ but this time varying $P$ and keeping $v$ (50 mm/s), $s$ (60 µm) and scan strategy (single line scan along the building direction) constant. Two different $E_s$ were tested: 15 J/mm² (45 W – 20% of maximum $P$) and 22 J/mm² (66 W – 30% of maximum $P$) and compared with the laser
<figure><img src="image_5.png" /><figcaption>Fig. 4. BSE and EBSD IPF z images of the cross-sections across laser tracks studied in Fig. 3. The noise (black dots) in the AB region of the EBSD maps is a consequence of surface undulations caused by chemical etching to facilitate BSE imaging of Cr-Mo microsegregations. BSE and EBSD images share the same scale. Similar to Fig. 3, the noise (black non-indexed points) in the EBSD maps is caused by chemical etching to reveal the intragranular structure in BSE images. This noise is lower in the lasered zones due to the microsegregation structures being smaller resulting in a cleaner (high confidence index) surface.</figcaption></figure>

The reference AB samples used are 316L stainless steel thin walls manufactured via LDED on hot-rolled 316L substrates using a single-pass-per-layer bidirectional scanning strategy with P 250 W, v 33.33 mm/s and s 0.7 mm. These parameters result in a surface energy density E_s = P/(v·s) = 10.71 J/mm². EBSD imaging after mirror polishing and etching an AB sample reveals
scan already performed at 8 J/mm² (24 W – 10% of maximum P). Based on the melt pool shapes (Fig. 4), the 8 J/mm² track resulted in conduction mode melting, whereas the 15 J/mm² and 22 J/mm² tracks resulted in keyhole mode melting. A pore can be observed close to the bottom of the melt pool in the 22 J/mm² laser track (black feature in the figure for 66 W), which is a signature of unstable keyhole mode melting. Such porosity is not visible in the 15 J/mm² track, however, the possibility of pores occurring elsewhere along the laser track cannot be eliminated.

<figure><img src="image_6.png" /><figcaption>Fig. 4: BSE images showing the melt pool depths at the sample edge after lasering with s 60 µm, v 50 mm/s and varying P (second parametric study).</figcaption></figure>

**3.2 Summarising the parametric studies**

In general, the melt pool depth increases with increasing E_s. It is more sensitive to P than to v; e.g., the depth triples when P is increased by a factor of 1.875 from 24 W, whereas it is halved when v is increased five-fold from 50 mm/s. A transition from conduction to keyhole mode melting occurs between 8 J/mm² and 15 J/mm². For all the scanning parameters used, cellular or planar grain growth occurs from the melt pool boundary during solidification; this deduction is supported by the crystallographic orientations and shapes of grains (Figs. 2, 3) and phase-field simulations aimed at understanding solidification during LPBF of 316L [18]. The newly formed grains have sizes similar to the unmelted ones in the base material and, as expected [18], they are bent along the direction of laser scanning. Following solidification, tensile residual stresses are generated along the laser track during cooldown as a consequence of the temperature gradient mechanism [19]. The local stresses can be sufficiently large to trigger dislocation dynamics causing local plastic deformation, as reported in a recent *in situ* synchrotron X-ray diffraction study [20]; this argument is further supported by the misorientation bands observed in the 50 mm/s track (Fig. 2). All laser scans result in a smoother surface, which can be deduced from the disappearance of noise in the EBSD maps caused by the reduction in surface undulations induced during etching.

Based on these studies, two key hypotheses can be derived. The smallest dislocation structures observed in both LDED and LPBF 316L typically either coincide with Cr-Mo
a weak texture (see Fig. S1) and a bimodal grain size distribution (see Fig. S2) due to long columnar grains along layer height with a tilt along the printing direction, and smaller equiaxed grains in between layers (see upper half of Fig. 3a). The AB microstructure is found to be nearly fully austenitic (greater than 99.9% γ-austenite), as verified from the analysis of the phase map (see Fig. S3) obtained from EBSD. A wide range of grain boundary misorientations are observed with a distinct high peak at 60° disorientation (Fig. S2) indicating an important presence of Σ3 twin boundaries.

The scratches in the EBSD maps in Fig. 3a are a consequence of polishing defects. The noisy (non-indexed) regions in Fig. 3a and 4 are caused by chemical etching of the sample before lasering to reveal the underlying intragranular microstructure of the AB material (Fig. 4). These scratches and noisy regions are deliberately kept in the images in order to highlight the surface improvement caused by laser scanning, which is discussed below.

Two parametric laser scanning studies were first conducted inside the CW laser-SEM under high vacuum. For both studies, a constant spot size *s* 60 µm was used; it is an order of magnitude smaller than the one used during LDED.

## 3.1 Parametric post-process laser scanning studies: Varying scan speed, constant laser power and laser spot size

The first parametric study was performed to understand the role of varying *E*<sub>s</sub> by varying *v* on microstructure alteration while keeping *P*, *s* and scan strategy constant. *P* was set at 10% of the maximum laser power, 24 W, which is the smallest stable value for the laser used. *s* was set to 60 µm, which is in the range of spot sizes used during LPBF. The scan strategy was a single line scan along the *y* direction (Figs. 3, 4, S4 (SE image of laser tracks), S5 (kernel average misorientation map of laser tracks)). The laser scans were parallel to each other and separated by a distance of 0.8 mm from their neighbours as seen in Fig. 3 (and Figs. S4, S5) to avoid any influence of one scan on the other. The *v* were chosen to correspond to five different *E*<sub>s</sub> lower than the one used to manufacture the samples; since the CW laser-SEM operates under vacuum and does not use any powder feedstock, there are no convection-related heat losses or powder-related laser reflection losses; hence, a lower *E*<sub>s</sub> should be needed than during the LDED process to melt the material. The five different *E*<sub>s</sub> (*v*) were: 8 J/mm<sup>2</sup> (50 mm/s), 4 J/mm<sup>2</sup> (100 mm/s), 1.6 J/mm<sup>2</sup> (250 mm/s), 0.8 J/mm<sup>2</sup> (500 mm/s) and 0.4 J/mm<sup>2</sup> (1000 mm/s). Laser scanning at the four highest *E*<sub>s</sub> (four lowest *v*) results in melting, solidification and rapid cooling of the material, whereas laser scanning at the lowest *E*<sub>s</sub> (highest *v*) results in heating and cooling without any apparent melting (Figs. 3,4). The SE image (Fig. S4) demonstrates the impact of the laser scans on the polished sample through the darker contrast along the laser tracks; the darker contrast in the laser tracks should be due to the topological changes brought about by melt-pool dynamics [20]; however, these changes do not affect the acquisition of EBSD data (Figs. 3, S5 and S6).

The melt pool sizes of the four remelted and solidified tracks (Fig. 4) expectedly decrease with decreasing *E*<sub>s</sub> (increasing *v*) and their shapes reveal that conduction mode melting has occurred. The Cr and Mo microsegregation cells (a signature of fast solidification in 316L [28–30]) in the lasered regions (Fig. 4) are at least an order of magnitude smaller than the ones in the AB material. Their size does not significantly decrease with increasing *v*, which was confirmed
microsegregation cells [3] or they are smaller than these cells [21–23]. Furthermore, smaller the size of the microsegregation cells, smaller the dislocation structures and higher the dislocation density [3,21–23]. It is well accepted that smaller (tens of nanometres) and denser feature sizes strengthen the material. The first hypothesis (hypothesis 1) is that the material in the lasered zone should exhibit smaller and denser dislocation structures, and consequently higher strength, than the AB material. Meanwhile, the observed surface refinement after lasering suggests that laser scanning unpolished samples should reduce the surface roughness if the lasering causes local melting. The second hypothesis (hypothesis 2) is that, together with an increase in strength, these samples should exhibit an improved fatigue limit.

### 3.3 Tensile testing polished and lasered 316L and 316L-Si – improving strength-ductility tradeoff

Hypothesis 1 was tested by performing the following experiment. Six 35 mm-long pieces were machined from multiple AB 316L walls and polished on their largest surfaces (x and y in Fig. 2); the polishing was done to perform EBSD mapping before laser scanning and compare with the EBSD maps after laser scanning. In addition, four pieces of the same size were machined from LDED 316L-Si walls. Two polished 316L and 316L-Si walls were lasered on both sides (one side in its entirety before the other) using a bidirectional scan strategy with a 50% overlap (30 µm step) between two consecutive scans and an $E_s$ of 11.67 J/mm² using P 35 W, s 60 µm, v 50 mm/s; these set of parameters, henceforth denominated L35, result in conduction mode melting, which is preferred from a manufacturing and service standpoint as it avoids creating keyhole porosities. Another two polished 316L walls were lasered on both sides with the same strategy and same $E_s$ (11.67 J/mm²), but double P and v i.e., with P 70 W, s 60 µm, v 100 mm/s (henceforth denominated L70). All laser scans were conducted under secondary vacuum. Dogbone shaped samples were extracted from all walls. Optical microscopy (OM), SEM, and scanning transmission electron microscopy (STEM) were performed before and after laser scanning. For an L70 316L sample, residual stresses were measured before and after lasering as well as after extracting the dogbone.

Polishing AB walls induces compressive stresses on their surfaces along both in-plane directions; e.g., a polished AB 316L wall exhibits $\sigma_{xx} \approx \sigma_{yy} \approx -75 \pm 20$ MPa. However, subsequent lasering on one face followed by the other results in bending and unbending of the walls and a significant alteration of the stresses; the same sample exhibits $\sigma_{xx} = 0 \pm 11$ MPa and $\sigma_{yy} = 271 \pm 39$ MPa on the first lasered surface and $\sigma_{xx} = 189 \pm 18$ MPa and $\sigma_{yy} = 350 \pm 41$ MPa on the second lasered surface. This difference in residual stress states before and after lasering implies that the sample has undergone plastic deformation during lasering. The residual stress state is further altered during dogbone extraction; the aforementioned sample exhibits $\sigma_{xx} = -57 \pm 21$ MPa and $\sigma_{yy} = 94 \pm 18$ MPa on the first surface and $\sigma_{xx} = 142 \pm 15$ MPa and $\sigma_{yy} = -34 \pm 20$ MPa on the second surface demonstrating a remarkable ~135° rotation of the in-plane stress vector.

Continuing to focus on 316L, the melt pool shapes in the lasered zones reveal that conduction mode melting had indeed occurred (Fig. 5a). The L35 and L70 scans have respectively penetrated 8% and 14% of the total thickness of the AB walls. Both scans result in the formation
by comparing their sizes in the 50 mm/s and 500 mm/s tracks. The cell size difference between
AB and lasered samples strongly indicates that solidification rates occurring during laser scanning
are faster than those occurring during the LDED process, and they increase with increasing $v$. For
the slowest laser scan (50 mm/s), the clearly visible microsegregation cells have an internal
diameter that is ~300 nm on average. Assuming this value to be the primary cellular arm spacing
($\lambda$) during solidification, and using the empirical formula (originally designed for secondary
dendritic arm spacing) $\lambda = 80/\dot{T}^{1/3}$ for stainless steels [31], the deduced solidification rate ($\dot{T}$) in
the slowest laser scanned region is ~1.9 × 10⁷ K/s; in comparison, the microsegregation cells in
the AB have a size of ~2.2 µm and the corresponding $\dot{T}$ is slower by three orders of magnitude. At
such large solidification rates, and corresponding high temperature gradients, the $\gamma$-austenite
grains epitaxially grow from the melt pool boundary via either cellular or dendritic (without
secondary dendrites) solidification mode and bend along the laser scanning direction [20,32],
which is evidenced in the EBSD images in Figs. 3, 4. In a recent work by Chen et al. [25] involving
the same LDED 316L used in this work, it has been argued that either primary austenite or primary
austenite with secondary ferrite mode occur during AM, which strongly supports the occurrence
of cellular solidification.

Following solidification, internal and sample constraints result in the formation of thermal
stresses in the heat-affected solid. These stresses are sufficiently high to cause some of the grains
to undergo plastic deformation e.g., via dislocation dynamics, which can be deduced from the
presence of misorientation bands observed in some grains in the remelted zone (see kernel average
misorientation map in Fig. S5, S6). Interestingly, the kernel average misorientation is overall lower
in the regions that melted during the laser scan than in the unmelted surrounding regions (Fig. S5),
which indicates that either the AB and/or the mechanical polishing prior to laser scanning induces
significant deformation that results in the manifestation of these larger misorientations.

After laser scanning, remarkably cleaner (higher indexation) EBSD images (Figs. 3b, 4)
are obtained along all the tracks; recall that no surface mechanical polishing was performed after
laser scanning. These cleaner EBSD images are a direct consequence of remelting (under vacuum,
resulting in a smoother and oxide-free surface) and subsequent refinement of the microsegregation
cells. A cleaner track is also visible in the EBSD image of the fastest laser scan, indicating that the
surface could have slightly remelted even though it is not discernible in the BSE and EBSD images
of the cross-section (Fig. 4).

## 3.2 Parametric post-process laser scanning studies: Varying laser power, constant scan velocity and laser spot size

The next parametric study was performed to investigate laser penetration depth as a
function of $E_s$ but this time varying $P$ and keeping $v$ (50 mm/s), $s$ (60 µm) and scan strategy (single
line scan along the building direction) constant. Two different $E_s$ were tested: 15 J/mm² (45 W –
20% of maximum $P$) and 22 J/mm² (66 W – 30% of maximum $P$) and compared with the laser
scan already performed at 8 J/mm² (24 W – 10% of maximum $P$). Based on the melt pool shapes
(Fig. 5), the 8 J/mm² track resulted in conduction mode melting, whereas the 15 J/mm² and 22
J/mm² tracks resulted in keyhole mode melting. A pore can be observed close to the bottom of the
melt pool in the 22 J/mm² laser track (black feature in the figure for 66 W), which is a signature
of zig zag grains in accordance to the bidirectional scanning strategy used. L35 results in a weak
texture, similar to the AB samples, but smaller grain sizes than the columnar grains of the AB
samples (Fig. 5b). In contrast, L70 results in a strong <100> texture with grains traversing multiple
tracks normal to the direction of lasering. Dislocation structures and density in the lasered region
are respectively smaller and higher than in the AB samples, and dislocation structures mainly
coincide with the microsegregations cells (Fig. 5c, d). While similar co-occurrence has been
reported for LPBF 316L [3], microsegregation cells after L70 are richer in Ni-Cr-Mo-Mn-Si-O as
opposed to only Cr-Mo ones reported in [3]. Furthermore, microsegregation cells in L70 samples
are decorated with Mn-Si-rich oxide precipitates that are smaller in size and more evenly
distributed than those occurring in LDED 316L.

<figure><img src="image_7.png" /><figcaption>**Fig. 5:** (a) OM (inset BSE) images of the cross-sections of L35 and L70 316L samples. (b) EBSD images of the surface of 316L before and after L35 (A) and L70 (B). (c) Bright-field STEM images of lamellae extracted from AB and L70 316L. (d) Energy dispersive X-ray spectroscopy images from L70 revealing chemical content in the region indicated with a black box in (c).</figcaption></figure>

Tensile testing reveals that AB 316L and 316L-Si exhibit a mean yield strength of 360 MPa and 414 MPa, respectively (Fig. 6a). Focusing on 316L, L35 and L70 result in an improvement of the mean yield strength by 11.67% and 31.11%. Using the rule of mixtures, the mean yield strengths of the lasered zones in L35 and L70 are found to be 884 MPa and 1157 MPa, respectively. To the best of our knowledge, such high yield strengths have never been reported for 316L fabricated either via LDED or LPBF.
of unstable keyhole mode melting. Such porosity is not visible in the 15 J/mm² track, however, the possibility of pores occurring elsewhere along the laser track cannot be eliminated.

<figure><img src="image_6.png" /><figcaption>Fig. 5. BSE images showing the melt pool depths at the sample edge after lasering with s 60 µm, v 50 mm/s and varying P (second parametric study).</figcaption></figure>

### 3.3 Summarising the parametric studies

In general, the melt pool depth increases with increasing $E_s$ (see Figs. 4, 5, S7). The parametric study with velocity variations reveals that the depth varies as the natural logarithm of $E_s$ (Fig. S7). However, the depth is more sensitive to P than to v; e.g., it triples when P is increased by a factor of 1.875 from 24 W, whereas it is halved when v is increased five-fold from 50 mm/s. A transition from conduction to keyhole mode melting occurs between 8 J/mm² and 15 J/mm². For all the scanning parameters used, cellular (or primary dendritic) grain growth occurs from the melt pool boundary during solidification; this deduction is supported by the crystallographic orientations and shapes of grains (Figs. 3,4) and phase-field simulations aimed at understanding solidification during laser scanning of 316L [20,32]. The newly formed grains have sizes similar to the unmelted ones in the base material and, as expected [20,32], they are bent along the direction of laser scanning. Following solidification, tensile residual stresses are generated along the laser track during cooldown as a consequence of the temperature gradient mechanism [33]. The local stresses can be sufficiently large to trigger dislocation dynamics causing local plastic deformation, as reported in a recent *in situ* synchrotron X-ray diffraction study [34,35] and predicted via crystal thermo-elasto-viscoplasticity simulations in [19]; this argument is further supported by the misorientation bands observed in the 50 mm/s track (Fig. 3). All laser scans result in a smoother surface, which can be deduced from the disappearance of noise in the EBSD maps caused by the reduction in surface undulations induced during etching.
Meanwhile, L35 316L-Si exhibits a mean yield strength of 469 MPa (an increase by 13.29%). These results together with the smaller denser dislocation structures (Fig. 5c) in the lasered zones validate hypothesis 1.

<figure><img src="image_8.png" /><figcaption>**Fig. 6:** (a) True stress-strain curves for AB, L35 and L70 316L and AB and L70 316L-Si. The circles represent the stresses and strains at failure of each sample. The small dotted lines are linear connectors between the end of strain measurement and the sample failure. The inset shows a zoomed in view of the boxed region. Post-deformation dark-field STEM images of lamellae taken from (b) AB and (c) lasered zone of L70 316L samples. The 316L AB-1 and L35-1 samples have been extracted from one wall and AB-2 and L35-2 from another wall; thus, while there is a difference in strength between the two AB samples of these walls, yet L35 results in a significant increase in strength for both.</figcaption></figure>

All samples fulfil the Considère criterion (Fig. 7), which means all microstructures achieve their full deformation capability prior to failure. This result is further corroborated by fractography studies (Fig. 8) showing the perfect bonding between lasered and AB regions; interestingly, the lasered region in L70 shows smaller voids, caused due to ductile damage, than in the AB region underneath as well as in the AB sample. This implies that post-process laser scanning has not adversely affected the ductility of 316L and 316L-Si. In fact, L70 has resulted in a slight (9.4%) improvement in ductility of 316L (Fig. 6a). Therefore, laser scanning has resulted in a net increase in the strength-ductility tradeoff for both materials.
Based on these studies, two logical outcomes can be expected. It has already been reported that the smallest dislocation structures observed in both LDED and LPBF 316L typically either coincide with Cr-Mo microsegregation cells [3] or they are smaller than these cells [36–38]. Furthermore, smaller the size of the microsegregation cells, smaller the dislocation structures and higher the dislocation density [3,36–38]. It is well accepted that smaller (tens of nanometres) and denser feature sizes strengthen the material [39,40]. The first expected outcome is that the material in the lasered zone should exhibit smaller and denser dislocation structures, and consequently higher strength, than the AB material. Meanwhile, the observed surface refinement after lasering suggests that laser scanning unpolished samples should reduce the surface roughness if the lasering causes local melting. The second expected outcome is that the intragranular structure refinement together with a reduced surface roughness should result in an improved fatigue limit.

To test these two outcomes, two optimized parameter sets with the same $E_s = 11.67 \text{ J/mm}^2$ are identified to increase melt-pool depth while remaining in conduction mode melting, which is desirable from a manufacturing standpoint to avoid porosities and uneven laser penetration. The first parameter set corresponds to P 35 W, $s$ 60 µm, $v$ 50 mm/s and it is henceforth denominated as L35. The second set denominated as L70 corresponds to P 70 W, $s$ 60 µm, $v$ 100 mm/s.

### 3.4 Tensile testing mechanically polished AB and lasered samples – improving strength-ductility tradeoff

The first expected outcome discussed in section 3.3 was tested by performing the following experiment. Six 35 mm-long pieces were machined from multiple AB walls and polished on their largest surfaces (x and y in Fig. 3); the polishing was done to perform EBSD mapping before laser scanning and compare with the EBSD maps after laser scanning. Three polished AB walls were lasered on both sides (one side in its entirety before the other) using a bidirectional scan strategy with a 50% overlap (30 µm step) between two consecutive scans and the L35 parameter set (see end of section 3.3). Another three polished AB walls were lasered on both sides with the same strategy and the L70 parameter set. All laser scans were conducted under high vacuum. Dogbone shaped samples (Fig. 2 top) were extracted from all walls. OM, SEM, and STEM were performed before and after laser scanning. For an L70 sample, residual stresses were measured before and after lasering as well as after extracting the dogbone.

Polishing AB walls induces compressive stresses on their surfaces along both in-plane directions; e.g., a polished AB wall exhibits $\sigma_{xx} \approx \sigma_{yy} \approx -75 \pm 20 \text{ MPa}$. However, subsequent lasering on one face followed by the other results in bending and unbending of the walls and a significant alteration of stresses; the same sample exhibits $\sigma_{xx} = 0 \pm 11 \text{ MPa}$ and $\sigma_{yy} = 271 \pm 39 \text{ MPa}$ on the first lasered surface and $\sigma_{xx} = 189 \pm 18 \text{ MPa}$ and $\sigma_{yy} = 350 \pm 41 \text{ MPa}$ on the second lasered surface. This difference in residual stress states before and after lasering implies that the sample has undergone plastic deformation during lasering. The residual stress state is further altered during dogbone extraction; the sample then exhibits $\sigma_{xx} = -57 \pm 21 \text{ MPa}$ and $\sigma_{yy} = 94 \pm 18 \text{ MPa}$ on the first surface and $\sigma_{xx} = 142 \pm 15 \text{ MPa}$ and $\sigma_{yy} = -34 \pm 20 \text{ MPa}$ on the second surface demonstrating a remarkable ~135° rotation of the in-plane stress vector.
<figure><img src="image_9.png" /><figcaption>**Figure 7:** Tangent modulus of the true stress-strain curve as a function of the true stress for all the 316L tensile testing samples. The black line indicates the Considère criterion.</figcaption></figure>

In order to better understand the origin of no change (and even a slight improvement) in ductility, TEM lamellae extracted post-mortem from AB and L70 316L samples were analysed via scanning TEM. The analysis revealed that while permanent deformation in the AB 316L sample was accommodated via dislocation structure evolution, significant deformation twinning along with dislocation structure evolution had occurred in the L70 316L sample (Fig. 6b, c for AB and L70 316L, respectively); these results were confirmed with selective area diffraction maps (not shown here) and the observation of newer and smaller (than the microsegregation) dislocation cells in both AB and L70 316L samples. EDS analysis revealed that the microsegregations were not affected indicating that in the L70 sample, the finer cells do not prevent twin propagation and dislocation structure evolution. The fact that both deformation twinning and dislocation structure evolution occur in the L70 sample explain why ductility is not adversely affected after lasering.

It should be noted that not all materials will have the possibility to accommodate deformation via multiple micromechanisms such as those occurring in 316L and therefore such an improvement in the strength-ductility tradeoff should not be a priori assumed for other materials.
<figure><img src="image_7.png" /><figcaption>Fig. 6. (a) OM (inset BSE) images of the cross-sections of L35 and L70 samples. (b) EBSD images of the surface of the sample before and after L35 (A) and L70 (B). (c) Bright-field STEM images of lamellae extracted from AB and L70. (d) Energy dispersive X-ray spectroscopy images from L70 revealing chemical content in the region indicated with a black box in (c).</figcaption></figure>

The melt pool shapes in the lasered zones reveal that conduction mode melting had indeed occurred (Fig. 6a). The L35 and L70 scans have respectively penetrated 8% and 14% of the total thickness of the AB walls without any decohesion of the interfaces between the refined and original untreated zones. Thus, the bonding between the lasered and AB zones can be considered to be perfect. Then, the resulting microstructure of L35 and L70 can be considered as a sandwich microstructure (a type of heterogeneous microstructure). Both scans result in the formation of zig zag grains in the remelted region in accordance to the bidirectional scanning strategy used. Furthermore, the distinct peak at 60° in the grain boundary misorientation significantly reduces for L35 and even more for L70 (Fig. S2). L35 results in a weak texture, similar to the AB samples, but smaller grain sizes than the columnar grains of the AB samples (Figs. 6b, S2). In contrast, L70 results in a strong <100> texture (Fig. S2) with grains traversing multiple tracks normal to the direction of lasering; some grains are larger than those encountered in AB and L35 (Fig. S2). A phase map analysis (not shown) reveals that the microstructure remains nearly fully austenitic for both L35 and L70; ~99.5% γ-austenite for both. Dislocation structures and density in the lasered region are respectively smaller and higher than in the AB samples, and dislocation structures mainly coincide with the microsegregations cells (Figs. 6c, d, S8). While similar co-occurrence has been reported for LPBF 316L [3], microsegregation cells after L70 are richer in Ni-Cr-Mo-
<figure><img src="image_10.png" /><figcaption>**Figure 8:** Fractography of AB (Top) and L70 (bottom) 316L samples. The scale for both samples is shown in the bottom image.</figcaption></figure>

**3.4 Fatigue testing unpolished and lasered samples – fatigue limit enhancement**

Hypothesis 2 is validated with the following experiment. Multiple 70-mm long unpolished AB 316L walls were laser scanned under secondary vacuum on both surfaces via L70. The surface roughness improved considerably with the mean surface height ($S_a$) decreasing from 16.593 µm on AB samples to 0.902 µm after L70. Dogbone samples were extracted from unpolished AB and lasered 316L walls and subjected to tension-tension fatigue at varying stress amplitudes and mean values. Despite the presence of tensile residual stresses due to lasering, L70 results in a significant (~25%) increase in the fatigue limit ($\Delta\sigma$) from 181.94 MPa for the AB sample to 227.24 MPa (Fig. 9). This improvement is a combined consequence of intragranular structure refinement (and the accompanying yield strength increase) along with the reduction in surface roughness.

<figure><img src="image_11.png" /></figure>
Mn-Si-O as opposed to only Cr-Mo ones reported in [3]. Furthermore, microsegregation cells in L70 samples are decorated with Mn-Si-rich oxide precipitates that are smaller in size, higher in number (increased number density) and more evenly distributed than those occurring in LDED 316L. Note that EDS analysis performed on the TEM lamellae of both samples did not reveal a significant change in the chemical composition of the material after lasering (see table S1). Despite the fact that EDS analysis on TEM lamellae is not statistically representative of the chemical composition of the whole material, the results of this analysis support the well-known fact that conduction mode melting does not result in significant change in chemical composition.

Tensile testing reveals that AB exhibits a mean yield strength of 360 MPa (Figs. 7a, S9, S10). L35 and L70 result in an improvement of the mean yield strength by 11.67% (402 MPa) and 31.11% (472 MPa), respectively. Since the tensile tests have been conducted under displacement control, and since the bonding between the lasered and the untreated zones is perfect for both L35 and L70, the macroscopic total strain experienced by both these zones is the same but the macroscopic stresses are distributed between the refined and untreated zones. Due to the (nearly) planar nature of the interface between the lasered and the original untreated zones, a binary rule of mixtures can be used as a zeroth order approximation to determine the yield strengths of the lasered region assuming that the untreated regions of L35 and L70 have the same yield strength as the AB sample. The following formula is used to compute the yield strength of the lasered zone based on the rule of mixtures: $\sigma^{\text{lasered}} = (\sigma^{\text{sample}} - (1-f)\sigma^{\text{AB}})/f$, where $\sigma^{\text{lasered}}$, $\sigma^{\text{sample}}$ and $\sigma^{\text{AB}}$ are the mean yield strength of the lasered zone, entire sample (L35 or L70) and the AB zone, respectively, and $f$ is the volume fraction of the lasered zone that is equal to 0.08 for L35 and 0.14 for L70; note that the length and width of the gauge region of the dogbone specimens for both samples are the same, hence, $f$ varies between samples only based on laser penetration depth. Then, the mean yield strengths of the lasered zones in L35 and L70 are found to be 885 MPa and 1160 MPa, respectively. To the best of our knowledge, such high yield strengths have never been reported for 316L fabricated either via LDED or LPBF. This increase in strength in comparison to the AB sample is primarily caused by refinement of the intragranular structure in the lasered zones of L35 and L70: microsegregation cell refinement, smaller dislocation structures and increase in dislocation density, higher precipitate size and density, and more solute trapping (due to higher cooling rates than those occurring during LDED) resulting in higher solid solution strengthening.

The lasered region in the L70 sample also exhibits a higher yield strength in comparison to some existing LPBF 316L in the literature [3,5,41–46], which is a result of several factors that have a compounded effect. The key reason is the higher cooling rate that leads to smaller microstructural features, larger dislocation density and greater solid solution without the occurrence of volumetric defects such as porosity or voids. The solidification rate, calculated using the empirical formula $\lambda = 80/\dot{T}^{1/3}$ [34], is ~1.9 × 10⁷ K/s for the slowest laser scan (50 mm/s). This cooling rate is three orders of magnitude higher than the one for LDED, and it is also higher than the typical LPBF rate [41,42,47]. Even when the LPBF cooling rate is similar [41], judging by the cell size, there tends to be insufficient densification in LPBF, leading to poorer performance [41,43]. The cell size, which contributes very significantly to strength through a Hall-Petch like relationship [3,43], is closer to the lower end of the range for our treated material (~300 nm) compared to LPBF (~200 nm - ~2 µm [3,41–44,47]). The dislocation density in the cell walls appears to be higher based on the TEM data. The microstructural features also include small, dense
**Fig. 9:** Secondary electron (SE) images of the (a, b) surfaces and (c, d) cross-sections of (a, c) unpolished AB and (b, d) L70 316L walls. (e) Wöhler (S-N) curve with double the stress amplitude (Δσ) vs number of cycles to failure (N) for tension-tension fatigue cycling of unpolished AB and lasered 316L samples. The trendlines are plotted for both samples and the arrows indicate fatigue runout cut-off at 3 million cycles.

# 4 Conclusions and perspectives

In this work, 316L and 316L-Si stainless steel thin-walls fabricated via the LDED process were subjected to post-process laser scanning inside the secondary vacuum of a newly designed system coupling a continuous-wave laser and an environmental SEM, called CW laser-SEM. The two widest surfaces of the LDED thin-walls were laser scanned using parameters similar to those used in LPBF processes. The results clearly demonstrate that lasering surfaces of as built (AB) stainless steels, using smaller spot sizes and faster scan speeds than those used during fabrication, can simultaneously significantly improve their strength-ductility trade-off and the fatigue limit. Notably,

* Laser scanning resulted in the formation of sandwiched microstructures exhibiting an intragranular structure refinement in the laser melted region in comparison to the base as built (AB) material. In fact, an order of magnitude reduction in the intragranular segregation cells and dislocation cell structure was achieved in the laser scanned region, which resulted in an increase in the overall yield strength of the sample. In one case, an overall increase in the yield strength by 31.11% over the 360 MPa of the AB material was obtained after laser scanning. For this sample, the laser penetration depth (where melting had occurred) was only 14% of the total depth of the sample. Consequently, through the rule of mixtures, a yield strength of ~1.15 GPa was obtained in the lasered region, which is significantly higher than the typical yield strengths achieved via LPBF of 316L.

* This increase in yield strength occurs without any ductility loss (and even a slight improvement in one case) indicating an overall improvement in the strength-ductility tradeoff. A deeper investigation in the case of lasered 316L reveals that this improvement occurs because plastic deformation is accommodated via dislocation structure evolution as well as deformation twinning in the lasered region.

* The strength enhancement, together with the surface roughness reduction due to laser scanning, translates into a significant improvement in fatigue life. An increase in fatigue limit by 25% was obtained for 316L.

Along with these results, a plethora of new interesting findings are presented, which merit deeper investigations in the future. These include (i) generation of residual stresses with strongly varying principal stresses across the sample thickness, (ii) grain structure and texture formation as a function of scan strategies keeping the energy density constant, (iii) coincidence of Ni-Mn-Si-O-rich segregations with the expected Cr-Mo segregations during solidification, (iv) reduction in the strain hardening effect after laser scanning, (v) the precise contributions of surface roughness reduction and intragranular structure refinement on fatigue limit improvement, etc.

The post-process CW lasering approach has promising prospects. For example, it can be integrated within an AM process, potentially minimising or eliminating the need for extensive
and more evenly spread precipitates (~15 - ~45 nm; see Fig. 6d) that will further strengthen the material. Furthermore, the high cooling rate leads to higher solid solution strengthening. In addition, during the laser treatment, plastic deformation also occurs [19], which further contributes to strengthening through work hardening.

All the aforementioned factors virtuously combine in the laser treated samples to produce a strengthening effect that surpasses the typical LPBF case. A comparison with other works also demonstrates how the laser treatment improves the mechanical response in comparison to LPBF. For instance, Yin et al. [45] present a material with less coherent cell walls, lesser precipitates and longer dislocation/microsegregation cells. The shorter cells and smaller grain size for the laser treated material could also be influenced by the smaller size of the treated region, which positively contributes towards improving the strength of the material. In a similar manner, Eres-Castellanos et al. [44] present LPBF 316L with larger and longer dislocation cells and a corresponding yield strength between 760 and 985 MPa. Similarly, Dryepondt et al. [46] also present LPBF 316L with larger cells. These results together with the smaller denser dislocation structures (Fig. 6c) in the lasered zones validate the first expected outcome.

<figure><img src="image_8.png" /><figcaption>Fig. 7. (a) True stress-strain representative curves for AB, L35 and L70. The circles represent the stresses and strains at failure of each sample. The small dotted lines are linear connectors between the end of strain measurement and the sample failure. The inset shows a zoomed in view of the boxed region. Post-deformation dark-field STEM images of lamellae taken from (b) AB and (c) lasered zone of L70 samples.</figcaption></figure>

All conditions fulfil the Considère criterion (Fig. 8), which means all microstructures achieve their full deformation capability prior to failure. This result is further corroborated by fractography studies (Fig. 9) showing the perfect bonding between lasered and AB regions; interestingly, the lasered region in L70 shows smaller voids, caused due to ductile damage, than in the AB region underneath as well as in the AB sample. This implies that post-process laser scanning has not adversely affected the ductility of 316L. In fact, L35 and L70 have respectively resulted in a 7.4% and 14.2% improvement in ductility of 316L (Fig. 7a). Therefore, laser scanning has resulted in an overall increase in the strength-ductility tradeoff.
post-processing steps and enabling the production of net-shaped parts with superior mechanical properties. Furthermore, it creates new avenues for optimising performance of not only AM stainless steels but also other alloys, including those manufactured using other techniques.

## Acknowledgments

The authors acknowledge the support of the following people: Pierre-Louis Hanappier (LMS) for printing the 316L samples, Nathalie Isac (LMS) for printing the 316L-Si samples, Simon Hallais (LMS) for discussions regarding SEM, Hind Meissera for assisting with the fatigue testing during their internship at LMS, and Eva Héripré (PIMM) for discussions on TEM lamellae preparation. JGSM, KC and MVU are grateful to the European Research Council (ERC) for their support through the European Union's Horizon 2020 – EXCELLENT SCIENCE – research and innovation program (grant agreement number 946959). All authors would like to thank LMS for providing financial support to build the CW Laser-SEM system.

## 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.

## CRediT authorship contribution statement

**Juan Guillermo Santos Macías:** Conceptualisation, Data Curation, Formal analysis, Investigation, Methodology, Validation, Visualisation, Writing – original draft, Writing – review & editing.

**Kewei Chen:** Formal analysis, Investigation, Validation, Visualisation, Writing – review & editing.

**Alexandre Tanguy:** Conceptualisation, Formal analysis, Investigation, Validation, Methodology, Writing – review & editing.

**Maxime Vallet:** Investigation, Visualisation, Writing – review & editing.

**Louis Cornet:** Investigation, Visualisation, Writing – review & editing.

**Vincent Michel:** Investigation, Visualisation, Writing – review & editing.

**Upadhyay Manas Vijay:** Conceptualisation, Formal analysis, Funding acquisition, Investigation, Methodology, Project administration, Resources, Supervision, Writing – original draft, Writing – review & editing, Validation, Visualisation.

**Data availability:** Data sets generated during the current study are available from the corresponding author on reasonable request.

**Code availability:** No code has been developed for this work.

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<figure><img src="image_9.png" /><figcaption>Fig. 8. Tangent modulus of the true stress-strain representative curve as a function of the true stress for AB, L35 and L70. The black line indicates the Considère criterion</figcaption></figure>

In order to better understand the origin of no change (and even a slight improvement) in ductility, thin-film lamellae extracted post-mortem from AB and the lasered zone of L70 were analysed via STEM. The analysis revealed that while permanent deformation in the AB sample was accommodated via dislocation structure evolution, significant deformation twinning along with dislocation structure evolution had occurred in the L70 sample (Fig. 7b, c for AB and L70, respectively); these results were confirmed with selective area diffraction maps (not shown here) and the observation of newer and smaller dislocation cells (than the microsegregation cells) in both AB and the lasered zone of L70 samples. The deformation twins in the lasered zone of L70 are thinner but higher in number density than the ones in AB. A similar response as the one reported for AB samples is expected in the untreated portion of L70 that has the same microstructure as the AB sample. Note that a thorough analysis of the deformation mechanisms occurring during tensile testing in the lasered and untreated zones of L70 requires performing *in situ* high resolution digital image correlation in the SEM during tensile testing, which is beyond the scope of this work. EDS analysis revealed that the microsegregations in the lasered zone of L70 were not affected, indicating that in the L70 sample, the finer cells do not prevent twin propagation and dislocation structure evolution. The fact that both deformation twinning and dislocation structure evolution also occur in the L70 sample explain why ductility is not adversely affected after lasering.

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<figure><img src="image_10.png" /><figcaption>Fig. 9. Fractography of AB (Top) and L70 (bottom) samples. The scale for both samples is shown in the bottom image.</figcaption></figure>

**3.5 Fatigue testing unpolished AB and lasered samples – fatigue limit enhancement**

The second expected outcome, discussed in section 3.3, is tested with the following experiment. Multiple 70-mm long unpolished AB walls were laser scanned under high vacuum on both surfaces via L70. The surface roughness improved considerably with the mean surface height (S<sub>a</sub>) decreasing from 16.593 µm on AB samples to 0.902 µm after L70. Dogbone samples were extracted from unpolished AB and lasered 316L walls and subjected to tension-tension fatigue at varying stress amplitudes and mean values. Despite the presence of tensile residual stresses due to lasering, L70 results in a significant (~25%) increase in the fatigue limit (Δσ) from 181.94 MPa for the AB sample to 227.24 MPa (Fig. 10).
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<figure><img src="image_11.png" /><figcaption>Fig. 10. SE images of the (a, b) surfaces and (c, d) cross-sections of (a, c) unpolished AB and (b, d) L70 walls. (e) Wöhler (S-N) curve with double the stress amplitude (Δσ) vs number of cycles to failure (N) for tension-tension fatigue cycling of unpolished AB and L70 samples. The trendlines are plotted for both samples and the arrows indicate fatigue runout cut-off at 3 million cycles.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>Fig. 11. SE images of fatigue fracture surfaces for (a) AB and (b) L70 samples under Δσ = 247.5 MPa. (a1-b1), (a2-b2) and (a3-b3) show SE images from the crack initiation region, crack propagation region, and final fracture region, respectively, for the (a1-3) as-built and (b1-3) L70 samples. The scale bars for a1-3 and b1-3 are shown in a3 and b3, respectively.</figcaption></figure>
This improvement is a combined consequence of intragranular structure refinement (and the accompanying yield strength increase) along with the reduction in surface roughness. The strengthening will increase the fatigue yield threshold of the material, while the surface smoothening will remove regions of stress concentration that would allow preferential fatigue crack nucleation and growth.

Fig. 11 shows the fracture surfaces of AB and L70 samples subjected to $\Delta\sigma = 247.5$ MPa. For the AB sample, cracks initiate from a large unmelted powder on the surface, whereas for the L70 sample, cracks originate from a region beneath the surface that appears to contain unmelted powder and pores. This indicates that the laser scanning process effectively eliminates unmelted powders and pores on the surface, as well as a thin layer beneath the surface, making it more difficult for cracks to originate from these areas. In the crack initiation region, brittle transgranular (cleavage) fracture predominates in both samples, as evidenced by the numerous crystalline cleavage facets and river-like patterns. In contrast, the final fracture region of both samples exhibits numerous dimples, suggesting a ductile fracture mechanism in this region.

It is worth commenting that the significant microstructural refinement and chemical heterogeneity reduction brought about by the treatment, together with the surface smoothening effect, should improve corrosion and wear resistance.

## 4 Conclusions and perspectives

In this work, 316L stainless steel thin-walls fabricated via the LDED process were subjected to post-process laser scanning inside the high vacuum of a newly designed system coupling a continuous-wave laser and an environmental SEM, called CW laser-SEM. The two widest surfaces of the LDED thin-walls were laser scanned using parameters similar to those used in LPBF processes. The results clearly demonstrate that lasering surfaces of as built (AB) stainless steels, using smaller spot sizes and faster scan speeds than those used during fabrication, can simultaneously significantly improve their strength-ductility trade-off and the fatigue limit. Notably,

* Laser scanning resulted in the formation of sandwiched microstructures exhibiting an intragranular structure refinement in the laser melted region in comparison to the base AB material. In fact, an order of magnitude reduction in the intragranular segregation cells and dislocation cell structure was achieved in the laser scanned region, which resulted in an increase in the overall yield strength of the sample. In one case (power 70 W, scan speed 100 mm/s and laser spot size 60 µm), an overall increase in the yield strength by 31.11% over the 360 MPa of the AB material was obtained after laser scanning. For this sample, the laser penetration depth (where melting had occurred) was only 14% of the total depth of the sample. Consequently, through the rule of mixtures, a yield strength of ~1.16 GPa was obtained in the lasered region, which is significantly higher than the typical yield strengths achieved via LPBF of 316L. The microstructural and chemical heterogeneity refinement coupled with the evened surface can also be positive for the corrosion behaviour, which is particularly interesting for the studied alloys.

* This increase in yield strength occurs with an improvement in ductility (14.2% for the aforementioned case) indicating an overall improvement in the strength-ductility tradeoff. A deeper investigation in the case of lasered 316L reveals that this improvement occurs
because plastic deformation is accommodated via dislocation structure evolution as well as deformation twinning in the lasered region.

* The strength enhancement, together with the surface roughness reduction due to laser scanning, translates into a significant improvement in fatigue life. A 25% increase in fatigue limit was obtained for the aforementioned case.

Along with these results, a plethora of new interesting findings are presented, which merit deeper investigations in the future. These include (i) generation of residual stresses with strongly varying principal stresses across the sample thickness, (ii) grain structure and texture formation as a function of scan strategies keeping the energy density constant, (iii) coincidence of Ni-Mn-Si-O-rich segregations with the expected Cr-Mo segregations during solidification, (iv) reduction in the strain hardening effect after laser scanning, (v) the precise contributions of surface roughness reduction and intragranular structure refinement on fatigue limit improvement, etc.

The post-process CW lasering approach has promising prospects. For example, it can be integrated within an AM process, potentially minimising or eliminating the need for extensive post-processing steps and enabling the production of net-shaped parts with superior mechanical properties.

Furthermore, it creates new avenues for optimising performance of not only AM stainless steels but also other alloys, including those manufactured using other (including conventional) techniques. For example, using the L70 parameters, laser scanning a warm rolled 316L sample with the same chemical composition as the one studied in this work, a yield strength of 280 MPa [48], the same sample thickness as the one used in Fig. 2 and the same laser penetration depth as shown in Fig. 6a would result in an increase in yield strength by ~44%. Therefore, the scope for mechanical property improvement is even higher for conventionally manufactured alloys.

Finally, the proposed post-processing approach is readily applicable to external surfaces (flat or curved) of AB samples. However, for sample geometries having internal surfaces, e.g., metamaterials, the same approach applied in-process (during fabrication, possibly as frequently as after the addition of each layer) would be more suitable.

## Acknowledgments

The authors acknowledge the support of the following people: Pierre-Louis Hanappier (LMS) for printing the 316L samples, Simon Hallais (LMS) for discussions regarding SEM, Hind Meissera for assisting with the fatigue testing during their internship at LMS, and Eva Héripré (PIMM) for discussions on TEM lamellae preparation. JGSM, KC and MVU are grateful to the European Research Council (ERC) for their support through the European Union's Horizon 2020 - EXCELLENT SCIENCE - research and innovation program (grant agreement number 946959). All authors would like to thank LMS for providing financial support to build the CW laser-SEM system and IMDEA Materiales for their support in performing granulometry on the AM powder.

## 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.
**CRediT authorship contribution statement**

**Juan Guillermo Santos Macías:** Conceptualisation, Data Curation, Formal analysis, Investigation, Methodology, Validation, Visualisation, Writing – original draft, Writing – review & editing.

**Kewei Chen:** Formal analysis, Investigation, Validation, Visualisation, Writing – review & editing.

**Alexandre Tanguy:** Conceptualisation, Formal analysis, Investigation, Validation, Methodology, Writing – review & editing.

**Nathalie Isac:** Investigation, Visualisation, Writing – review & editing.

**Maxime Vallet:** Investigation, Visualisation, Writing – review & editing.

**Louis Cornet:** Investigation, Visualisation, Writing – review & editing.

**Vincent Michel:** Investigation, Visualisation, Writing – review & editing.

**Manas Vijay Upadhyay:** Conceptualisation, Formal analysis, Funding acquisition, Investigation, Methodology, Project administration, Resources, Supervision, Writing – original draft, Writing – review & editing, Validation, Visualisation.

**Data availability:** Data sets generated during the current study are available from the corresponding author on reasonable request.

**Code availability:** No code has been developed for this work.

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