Microstructure and properties of high-strength C + N austenitic stainless steel processed by laser powder bed fusion

Journal Pre-proof Microstructure and properties of high-strength C+N austenitic stainless steel processed by laser powder bed fusion J. Boes (Conceptualization) (Data curation) (Formal analysis) (Investigation) (Methodology) (Validation) (Visualization) (Writing original draft) (Writing - review and editing), A. R¨ottger (Conceptualization) (Funding acquisition) (Methodology) (Project administration) (Supervision) (Writing - review and editing), W. Theisen (Conceptualization) (Funding acquisition) (Project administration) (Resources) (Supervision) (Writing - review and editing)

PII:

S2214-8604(19)31410-1

DOI:

https://doi.org/10.1016/j.addma.2020.101081

Reference:

ADDMA 101081

To appear in:

Additive Manufacturing

Received Date:

2 September 2019

Revised Date:

20 December 2019

Accepted Date:

19 January 2020

Please cite this article as: Boes J, R¨ottger A, Theisen W, Microstructure and properties of high-strength C+N austenitic stainless steel processed by laser powder bed fusion, Additive Manufacturing (2020), doi: https://doi.org/10.1016/j.addma.2020.101081

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Microstructure and properties of high-strength C+N austenitic stainless steel processed by laser powder bed fusion

J. Boes*, A. Röttger, W. Theisen

Chair for Materials Technology, Institute for Materials, Ruhr University Bochum, Germany

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Abstract In the developing field of laser powder bed fusion (L-PBF), austenitic stainless steels, such as AISI 316L, have gained great importance owing to their excellent processability. However, the moderate strength of these steels limits their applicability. This can be counteracted by the use of nitrogen as an alloying element to improve both strength and corrosion resistance.

Keywords:

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In this work, nitrogen-alloyed high-strength austenitic stainless steel X40MnCrMoN19-18-1 was processed by L-PBF, and the resulting microstructural and mechanical properties were investigated. The same material was also processed by hot isostatic pressing (HIP), which was used as a reference state. In the L-PBF process, argon and nitrogen were used as process gases to investigate the influence of process atmosphere on the microstructure and on changes in the chemical composition during processing. The results show a minor decrease in the nitrogen content of the steel after L-PBF, independently of the process gas, whereby argon resulted in a slightly higher specimen density. The microstructure after L-PBF processing contained small precipitates that could be removed by a short solution-annealing treatment. The tensile properties of the L-PBF-built steel are comparable to those of the steel produced by hot isostatic pressing in terms of ultimate tensile strength, but had lower elongation to fracture values. The ductility of the material was enhanced by solution annealing without significant impairment of the ultimate tensile strength. This work demonstrates that nitrogen-alloyed stainless steels can be processed by means of L-PBF and can extend the variety of appropriate steels towards applications with high requirements for the material strength and chemical resistance.

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additive manufacturing, laser powder bed fusion, C+N steel, high-nitrogen steel, austenite

Introduction

In the last decades, nitrogen has gained growing importance as an alloying element in various steel grades [1]. A beneficial effect of nitrogen on the mechanical properties of steel was first studied by Frehser et al., who found an increase in the yield strength induced by nitrogen additions without a pronounced reduction in material toughness [2]. In this context, Nilsson et al. confirmed a linear dependency of the yield strength of different austenitic stainless steels on the nitrogen content, which had already been predicted by Irvine et al. based on investigations of the influence of different alloying elements on the mechanical properties of various steels [3]. Besides the pronounced solidsolution strengthening induced by nitrogen, a positive effect of nitrogen on the cold-work hardening

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behavior of austenitic stainless steels was reported [4]. According to early studies of Uggowitzer et al., a combination of solid-solution strengthening and increased work-hardening coefficients lead to increased material strength. In addition to the influence of nitrogen on the mechanical properties, nitrogen was found to improve the corrosion resistance with regard to intercrystalline corrosion and pitting corrosion in several studies [5–8]. The pronounced influence of nitrogen on the pitting resistance of stainless steels is taken into account in the pitting resistance equivalent number (PREN), in which the impact of nitrogen is amplified by the highest factor compared to all other alloying elements [1]. The mechanisms of the described improvements of the mechanical and chemical properties of steels due to nitrogen addition were further investigated by Gavriljuk et al., Berns et al., and Shanina et al. using Mössbauer spectroscopy, electron spin resonance investigations, and ab initio calculations [1,9–13]. As an interstitial atom in the Fe lattice, nitrogen atoms strongly hinder the movement of dislocations and therefore contribute to the material’s strength. Simultaneously, nitrogen increases the electron state density at the Fermi level and thus donates free electrons to the lattice [13]. In consequence, nondirectional metallic bonding of atoms is promoted, and the material’s ductility is not impaired significantly by relatively high contents of interstitial atoms [14]. The increase in electron density at the Fermi level can be amplified by combinational alloying with nitrogen and carbon, leading to high material strength and simultaneously good ductility [9,15]. A further reason for increased strength induced by nitrogen additions can be found in the effect of nitrogen on the short-range ordering in the Fe lattice [9]. Due to specific interactions of nitrogen atoms with substitutional atoms, nitrogen favors a homogeneous distribution of substitutional atoms in the lattice. As a result, the impact of the substitutional atoms on the solid-solution strengthening is amplified indirectly by nitrogen.

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In addition to the advantageous influence of nitrogen in steel, the challenging production of nitrogen-alloyed steels due the very different solubility of nitrogen in the phases occurring in the FeC system should be mentioned [16]. Whereas the nitrogen solubility in the octahedral sites of fcciron is approx. 2.8 wt%, it is only approx. 0.1 wt% in δ-ferrite. In the case of primary δ-ferritic solidification, the nitrogen present in the melt (up to approx. 0.14 wt%) is forced out of the primary solid phase into the melt, where the maximum solubility is eventually exceeded and the atomic nitrogen recombines to molecular gaseous nitrogen [1]. In consequence, gas pores form that diminish the material's properties. The nitrogen solubility in the melt can be increased by alloying with Mn, Cr, Ti, and V [17]. However, the use of elements Ti and V is limited owing to their high tendency to form hard phases and thereby bind the nitrogen. The elements Mn and Cr are used to increase nitrogen solubility in -Fe, whereas carbon, silicon, and nickel decrease the nitrogen solubility in the melt. Owing to the stabilizing effect of Mn and also N on the austenitic phase, Ni contents can be substituted without decreasing the austenite stability [18]. In addition, this substitution of Ni in austenitic stainless steels by the element Mn can also be beneficial in terms of cost reduction and biocompatibility [19,20]. Therefore, austenitic CrMnN steels are possible alternative materials to conventional CrNi austenitic steels, such as AISI 316L, which are used in various applications and have recently found extraordinary interest in connection with laser additive manufacturing [21]. Especially in powder bed-based laser additive manufacturing (laser powder bed fusion, L-PBF), 316L has been studied extensively owing to its superior processability and its relatively good mechanical and chemical properties. During L-PBF, a focused laser beam selectively melts a metallic powder layer, creating a part layer by layer from 3D-CAD data. Due to the locally and temporally unsteady heat input induced by the moving laser spot and the layer-wise build-up of a part, high cooling rates and inhomogeneous temperature distributions occur during L-PBF [22,23]. These process specific conditions can be problematic in terms of crack initiation when manufacturing steels. Thus, mainly steels with a high ductility, such as the aforementioned austenitic stainless steel 316L, can be processed by means of L-

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PBF. Besides the enhanced risk for crack formation during L-PBF, the process conditions promote the formation of unique microstructures. In L-PBF-built steels, highly heterogeneous non-equilibrium microstructures are present [24,25]. For example, anisotropic grain structures, microsegregations and cellular sub-grain structures with high local dislocation density are commonly found attributes in L-PBF-densified steels [26–28]. Especially, the extraordinary distribution and interaction of solute atoms and dislocations in cellular sub-grain structures appears to result in increased strength and decreased ductility of L-PBF-built steels compared to the conventionally produced material [28,29]. Whereas the microstructure formation mechanisms are discussed controversially, the effect of the microstructure on the mechanical properties is generally reported to be beneficial [26,30]. Correspondingly, L-PBF-built 316L was found to have a higher strength compared to the same material in the cast and HIP state. Nonetheless the material strength and ductility of L-PBF-built 316L is still exceeded by FeCrMnN austenitic steels [25,31–33]. Thus, the good L-PBF processability of austenitic steels could be combined with improved mechanical and chemical properties by the use of CrMnN steels in L-PBF. In this study, the feasibility of manufacturing a high-nitrogen CrMnN steel by means of L-PBF and characterization of the resulting mechanical properties were investigated by processing X40MnCrMoN19-18-1 steel with L-PBF. As a result, the variety of steels that can be processed by L-PBF can be extended and the use of L-PBF for the production of parts for applications with high requirements for the mechanical and chemical properties of the material can be achieved. It is the aim of this work to gain general knowledge concerning the microstructure and the associated mechanical properties that result from the L-PBF process.

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Experiments Materials

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In this study, austenitic X40MnCrMoN19-18-1 steel is investigated in the hot isostatically pressed (HIP) condition and in the L-PBF-densified condition. The gas-atomized steel powder was provided by Materials Science and Engineering Werkstoffzentrum Clausthal UG (Clausthal-Zellerfeld, Germany) with particle sizes < 250 µm. The powder fraction was set to 20 - 63 µm by means of sieving to provide sufficient processability by L-PBF. The chemical composition of the powder, which was used for both L-PBF and HIP, was analyzed by energy dispersive X-ray spectrometry (EDS). The carbon content and the nitrogen content of the powder were determined by carrier-gas hot extraction using CS-800 and ONH-2000 analyzers, respectively, from ELTRA GmbH. The particle size distribution and the average particle size diameters d(0.1), d(0.5) ,and d(0.9) were measured by means of laser diffraction using a Mastersizer 2000 from Malvern Instruments Ltd. The measurements were performed in accordance with ISO 13320, and the average mean was calculated from four single measurements. To characterize the feeding behavior of the powder, the apparent density was measured according to DIN ISO 697 and the powder flow rate was estimated using a Hall flowmeter according to DIN EN ISO 4490. The chemical composition of the investigated material is listed in Table 1. Figure 1 b shows an SEM image of the investigated powder particles. The powder particles have spherical shapes with a small amount of satellites. However, there is a small number of elongated or irregularly shaped particles, which were formed due to suboptimal atomizing conditions and which favor an irregular alteration of the particle shapes and the collision of partly solidified droplets. Pronounced microsegregations in the microstructure of the particles can be observed (Figure 1 c), as characterized by the grey color scale in the single-phase material. The investigated steel powder reveals a unimodal particle size distribution with mean diameters d(0.1) = 19.68 µm, d(0.5) = 39.66 µm, and d(0.9) = 63.27 µm (Figure 1 a), which indicates a shift in the particle size distribution toward smaller particle sizes than the desired lower fraction boundary of 20 µm. Moreover, the d(0.9) value indicates that at least 10 % of the particles have a diameter greater than

63 µm. The described morphology of the particles and the particle size distribution result in a flow time of 18.1 s in the Hall flow test, a bulk density of 4.42 g/mm3, and a packing density of 0.58. These properties were sufficient for the use of the Realizer SLM 100 device.

Table 1:

Chemical composition of the investigated steel powder in wt%. N 0.39

Mn 19.47

Cr 18.23

Mo 1.09

Si 0.24

∑P+S 0.013

Fe Bal.

a) Particle size distribution of the investigated steel powder, b) SEM micrograph of lose powder particles, and c) SEM micrograph of the cross-section of a powder particle.

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Figure 1:

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X40MnCrMoN19-18-1 powder

C 0.39

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HIP-densified X40MnCrMoN19-18-1 steel made from the same powder as for L-PBF processing was used as a reference material. Figure 2 shows the SEM image of the microstructure of the material in the HIP condition as a reference. In this condition, the steel possesses a single-phase austenitic microstructure.

SEM micrograph of the microstructure of X40MnCrMoN19-18-1 steel in the HIP reference state.

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Figure 2:

Thermodynamic equilibrium calculations

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Thermodynamic equilibrium calculations were performed using ThermoCalc® software. Simulations were used to gather information on the phase stability as a function of the temperature and the phase compositions. Thermodynamic data for the simulations was taken from the TCFe9 database. The conducted simulations consider the phases LIQUID (liquid phase), FCC_A1#1 (γ-Fe), FCC_A1#2 (M(C,N) carbonitrides), BCC_A2 (α-Fe, δ-Fe), M23C6, HCP (M2(C,N) carbonitrides), and gaseous phase. The system size and pressure were set to n = 1 mol and 101325 Pa, respectively.

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Laser powder bed fusion

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For the densification of the investigated X40MnCrMoN19-18-1 steel with L-PBF, a Realizer SLM 100 device was used. This device is equipped with an ytterbium fiber laser with a nominal output power of 100 W, a wavelength of ≈ 1064 nm, and a focal diameter of ≈ 90 µm. The effective output power, measured by the manufacturer, amounted to 77.4 W. The experiments were performed under an N2 or Ar atmosphere to ensure an oxygen content of less than 0.3 vol% in the process chamber. Due to the lower density of N2 compared to air or Ar, the gas filling system of the L-PBF machine was modified by creating an additional gas outlet at the bottom of the build chamber. To evaluate suitable process parameters for the buildup of dense and crack-free specimens, the laser parameters exposure time tE and point distance pd were varied, whereas the slice thickness (30 µm) and the laser output power (77.4 W) remained constant. Thus, the effective input energy per unit length El (Eq. 1) ranged between 232.2 J/m and 451.5 J/m. A hatch distance of 120 µm was used. The respective laser parameters were chosen based on the results of previous findings on densification of AISI 316L stainless steel and several tool steels [34–36]. It was assumed that the densification behavior of the AISI 316L and the investigated X40MnCrMoN19-18-1 steel were similar, and the chosen parameters were identified as promising. An x/y interlayer stagger strategy was used for the buildup. 𝐸𝑙 =

𝑃𝑒𝑓𝑓 𝑝𝑑 𝑡𝐸

(1)

For microscopic investigations of the resulting microstructure of the X40MnCrMoN19-18-1 steel, only small cubic samples with an edge length 5 mm x 5 mm x 2.55 mm (length x width x height) were built. Furthermore, tensile specimens (Figure 3) were produced with the most favorable set of laser

parameters in order to characterize the mechanical properties of the L-PBF-built X40MnCrMoN1-181 steel compared to the same steel in the HIPed condition. Hot isostatic pressing and heat treatment

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In addition to L-PBF-densified X40MnCrMo19-18-1 steel in the as-built condition (L-PBF), L-PBF-built and solution-annealed specimens (L-PBF annealed) were investigated. Solution annealing was performed at a temperature of 1180 °C for 30 min in an inert gas atmosphere. The same powder was used for producing specimens by hot isostatic pressing as a reference state. Hot isostatic pressing of encapsulated steel powder was conducted using a QIH-9 URQ hot isostatic press from Quintus Technologies (Sweden), which allows rapid quenching of the material by gas quenching immediately after densification so that a further solution-annealing heat treatment is not necessary. During the HIP densification process, an Ar pressure of 150 MPa and a temperature of 1150 °C were applied for 3 h. Subsequently, Ar was used for gas quenching of the encapsulated and HIP-densified material inside the HIP to suppress precipitation processes. Microscopy

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Microscopic investigations were performed on specimens taken from both L-PBF-built and HIPed X40MnCrMoN19-18-1 steel. Specimens of the reference material for microscopic investigations were produced by cutting, grinding (SiC abrasive paper with mesh sizes of 320, 500, and 1000), polishing (diamond suspension with average particle diameters of 6 µm, 3 µm, and 1 µm), and etching (NITAL solution) of the specimen's cross-section.

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Microscopic investigations of the microstructure were performed using an MIRA 3 SEM from Tescan. The SEM images were obtained with a working distance of 15 mm and an acceleration voltage of 15 keV. The local chemical composition of microstructural constituents was determined by energy dispersive X-ray spectroscopy (EDS) using an AZtec Energy Advanced system from Oxford Instruments with a working distance of 15 mm and an acceleration voltage of 20 keV.

X-ray diffraction

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The porosity of the L-PBF-built specimens was measured by image analysis of four binarized OM micrographs for each investigated parameter set using image analysis software ImageJ (version 1.49v), and the percentage of black pixels was evaluated for each image. Porosity values are given as mean values with a standard deviation.

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In order to characterize the phase composition of the investigated steel in the L-PBF-built condition as well as after a subsequent solution annealing step, X-ray diffraction (XRD) measurements were conducted. For this purpose, a Bruker D8-Advanced device was used with a Bragg-Brentano set-up and CuKα radiation (wavelength: 1.5406 Å). The X-ray diffraction pattern was recorded in the range of 30 – 80 ° 2θ with a step size of 0.01 and an acquisition time of 10 s per step. The obtained diffraction patterns were analyzed using the software DIFFRAC.EVA V3.0. The preparation of the investigated specimens was performed as described in the previous section. Tensile testing

Tensile tests were conducted to compare the mechanical properties of the X40MnCrMo19-18-1 steel in the different conditions. The specimen geometry is shown in Figure 3. Specimens were built with L-PBF or produced from bulk material (HIP) by electrodischarge machining. The specimens produced by L-PBF were oriented with their tensile direction perpendicular to the build direction. The gauge length of the specimens was 10 mm at a width of 2 mm and a specimen thickness of 2 mm. The total length of the specimens was 26 mm. Tensile experiments were performed in accordance with DIN EN

Geometry of the investigated tensile specimens [34]. Dimensions are given in mm.

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Figure 3:

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ISO 6892-1:2016 [37] using a ZWICK-Roell Z100 tensile testing machine. A crosshead speed of 0.5 mm/min was applied in the tests, which were performed until the specimen ruptured. Specimen elongation was recorded by an extensometer and by crosshead display. Additionally, strain gauges were used for determining the Young’s modulus.

Results and discussion

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Porosity and crack density

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The results of the parameter study are presented in Table 2. As described before, the investigated parameter sets were chosen from previous studies of the densification of a stainless steel with L-PBF [25,35,36] and based on the assumption of similar densification characteristics of both N-alloyed and N-free austenitic stainless steels [25]. The obtained results confirm the aforementioned assumption showing sample densities of more than 99.8 % and 99.5 %, depending on the respective inert gas atmosphere. In accordance with literature on basic densification phenomena in L-PBF [24], there is an increase in the density with increasing energy input that is due to the increasing size of the melt pool [43]. However, if a critical input energy is exceeded, the sample density decreases rapidly due to pronounced material evaporation, balling phenomena, and unstable fluid kinetics in the melt pool. Therefore, an exposure time of 110 µs and a point distance of 30 µm were evaluated as optimal parameters for the production of dense samples of the investigated steel. However, the inert gas atmosphere had a notable impact on the sample density. The results display a decreasing density when Ar is replaced with N2 as the inert process gas. The larger number of pores and binding defects inside the samples was also detected in OM images, as depicted in Figure 4. Interestingly, the microstructure of the samples built under an N2 atmosphere revealed a larger number of regularly distributed binding defects, indicating a lower energy input during the L-PBF process in an N2 atmosphere compared to an Ar atmosphere. Various theories on the influence of the process gas on the melt pool stability and evaporation processes are discussed in the literature [44,45]. Dai et al. found lower densities in L-PBF-built parts consisting of TiC/AlSi10Mg composite when using an N2 atmosphere compared to an Ar atmosphere. Combining their experimental findings with numerical simulations of the L-PBF process, they concluded that, in contrast to N2, Ar stabilizes the melt pool and therefore provides better part quality [45]. Moreover, the efficiency of the inert gas flow in

removing byproducts (smoke) from the interaction zone of the laser and the powder bed during melting strongly influences the ratio between the energy that is brought into the material and the energy that is absorbed above the powder bed/melt pool by the aforementioned byproducts [46]. Therefore, the efficiency of the inert gas flow in removing byproducts from the interaction zone limits the melt pool size for a given laser energy. In this context, the different densities of the Ar and N2 gases might be the reason for the different inert gas flow characteristics inside the build chamber, leading to different energy inputs by the laser radiation into the powder bed. The higher porosity of the samples built under an N2 atmosphere might therefore be a result of inefficient removal of particles from the laser/powder interaction zone by the inert gas flow. Nevertheless, the attained sample densities can be considered as adequate for testing the material’s properties in the L-PBFbuilt condition.

30

100 110 150 175 150 175

Energy input in J/m 258.0 283.8 387.0 451.5 232.2 270.9

Porosity in vol% Ar 0.20 ± 0.05 0.17 ± 0.04 0.20 ± 0.04 0.68 ± 0.32 0.35 ± 0.20 0.28 ± 0.06

N 0.68 ± 0.06 0.44 ± 0.05 1.64 ± 0.27 2.51 ± 0.03 0.52 ± 0.31 0.40 ± 0.33

OM micrographs of the microstructure of the investigated steel manufactured by LPBF using a) a nitrogen atmosphere and b) an argon atmosphere

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tE in µs

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Pd in µm

Laser parameters and resulting specimen porosity values of the L-PBF-built X40MnCrMoN19-18-1 steel.

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Table 2:

Microstructure

SEM micrographs of the microstructure of the L-PBF-built steel manufactured under a N2 atmosphere and under an Ar atmosphere are depicted in Figure 5 a-d. Independent of the used process gas, the steel has a fine-grained microstructure. Similar to various other L-PBF-built materials, there is a cellular subgrain structure that consists of equiaxed to columnar dendrites with microsegregations in the interdendritic spaces [26,29,47]. As indicated by EDS measurements, mainly heavier elements such as Cr and Mo segregate into the interdendritic regions (Figure 6). This is in accordance with literature concerning more commonly processed stainless CrNi-austenites containing Cr and Mo [28]. In addition to Cr, Mo and Si, also Mn can be found in higher concentrations in the interdendritic

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spaces compared to the dendritic areas in the case of X40MnCrMoN19-18-1 (Figure 6). Micropores with a diameter of several nm are also present in the microstructure. Whereas the aforementioned microstructural elements are commonly found in L-PBF-built austenitic stainless steels and numerous other materials, the investigated steel also contains finely dispersed precipitates having a spherical to blocky shape (Figure 5 c,d). Using X-ray diffraction analysis (XRD), three different phases besides the austenitic matrix can be identified (Figure 7). The measurements indicate the presence of Cr- and Mo-rich (carbo-) nitrides of M2N and M2(C,N) type in the L-PBF-built steel. Furthermore, Cr-rich M23C6 carbides can be found. The XRD measurements are in good agreement with results from thermodynamic equilibrium calculations (Figure 8), which suggest the formation of Cr- and Mo-rich carbides and carbo-nitrides. Furthermore, the identified phases of the investigated steel in L-PBFbuilt condition are similar to those found in conventionally produced high-nitrogen CrMn steels in terms of type and chemical composition [48,49]. Combining the results of the phase analysis and the SEM investigations, it can be concluded, that the microstructure of L-PBF-built X40MnCrMoN19-18-1 steel consist of an austenitic matrix with finely dispersed Cr- and Mo-rich (carbo-) nitrides. According to solidification simulations, precipitates are expected to form at the end of the solidification sequence, which implies that they are formed in the interdendritic regions. As shown in Figure 5 c,d, precipitates are found in both the dendritic and interdendritic regions of the steel. Thus, consideration must also be given to formation of precipitates from the solid austenitic matrix. It is possible that the heat input by the laser during build-up might catalyze the precipitation process.

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SEM images of the microstructure of the L-PBF-built steel manufactured under a N2 atmosphere (a, c, e) and under an Ar atmosphere (b, d, f) in the as-built condition (ad) and after subsequent solution annealing (e, f).

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Local chemical composition in the interdendritic (EDS spectrum 1) and dendritic spaces (EDS spectrum 2) of the investigated steel in L-PBF-built condition measured by EDS.

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Figure 6:

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To study the influence of the typical microstructure of the L-PBF-built steel, the L-PBF-densified steel was also investigated after subsequent solution annealing. As shown in Figure 5 e,f, the microstructure of the steel in the solution-annealed condition does not show precipitates or microsegregations. The results of the SEM investigations are supported by XRD measurements, which confirm a single-phase austenitic microstructure without a detectable amount of other phases (Figure 7). The chosen solution temperature of 1180 °C and the annealing time of 30 minutes are thus sufficient for the dissolution of both precipitates and segregations. Probably, the small size of both aforementioned microstructural constituents and the associated short diffusion distances, which are necessary for chemical homogenization during the heat treatment, allow successful solution annealing in a comparatively short time. The heat treatment of conventionally produced austenitic CrMn-HNS that contains large precipitates often requires longer annealing times, enhancing the risk of decarburization or nitriding of the respective part during heat treatment [1]. Moreover, the high solution-annealing temperature can initiate considerable grain growth, especially during a long heat treatment at high temperatures. Thus, long solution-annealing times would diminish the overall properties of the steel. In contrast, the L-PBF-processed steel with its finegrained microstructure allows a relatively simple and short heat treatment without a pronounced risk of significant changes in the chemical composition or the grain size of the steel during annealing.

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Figure 7:

XRD diffraction pattern of the investigated steel in the L-PBF-built condition and after a subsequent solution annealing.

Chemical composition

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The investigated steel was developed with the aim of an increased material strength and a simultaneously high ductility and toughness, combined with superior corrosion resistance. Based on the findings of Gavriljuk et al., a combination of nitrogen and carbon as interstitial elements was chosen in order to attain the aforementioned properties by increasing the electron state density at the Fermi level and to improve the short-range atomic order [10,11,50]. The required solubility of the liquid phase of the steel for a nitrogen content of ~0.4 wt% is provided by large amounts of the elements Mn and Cr. Jiang et al. showed that combinational alloying with Cr and Mn at about 18 wt% each can significantly increase the nitrogen solubility in the Fe-Cr-Mn-N system at a given nitrogen pressure [17]. The comparatively large amounts of Cr and Mn in the investigated steel were chosen in order to use a pressureless conventional production. After L-PBF-densification, steel samples have a slightly decreased nitrogen content compared to the nitrogen content of the starting powder (Figure 9). The differences in the nitrogen content reveal depletion of nitrogen during L-PBF. It should be mentioned in this context that the decreasing nitrogen solubility of iron in the liquid state with increasing temperature also applies to steels [17,51]. To characterize the influence of temperature on the nitrogen solubility of the investigated steel in the liquid state, the nitrogen content of the melt was calculated using ThermoCalc software. The calculations were carried out for the temperature regime in which the solid phase, liquid phase, and/or gaseous phase are thermodynamically stable (temperatures indicated in Figure 8). The results are given in Table 3. They reveal a notable decrease in the nitrogen content of the melt at temperatures higher than 1760 °C. This is accompanied by the formation of a gaseous phase in the investigated system, which mainly consists of nitrogen-rich molecules. Since this effect takes place at temperatures significantly higher than Tliq, the effect is not problematic in casting processes, in which the steel is melted at temperatures much closer to Tliq. In contrast to this, temperatures in the melt pool of around 2000 °C can occur in L-PBF processes [22,23,44,52]. Thus, the higher temperatures during the L-PBF process might induce a drop in the nitrogen solubility of the investigated steel. In consequence, this may lead to nitrogen degassing that would lower the nitrogen content of the L-PBF-built sample compared to the starting powder. Degassing of nitrogen due to the formation of primary δ-ferrite, which is suggested by the calculated phase-temperature diagram (Figure 8) is unlikely because there was no nitrogen loss measured when comparing the chemical composition of the steel before atomization and in the gas-atomized condition.

Table 3:

Nitrogen content and volume fractions of liquid and gaseous phase obtained by thermodynamic equilibrium calculations.

Temperature in °C

Volume fraction of liquid phase 0.011 1.000 1.000 0.473 0.095

Volume fraction of gas phase 0.000 0.000 0.000 0.527 0.905

Thermodynamically stable phases and respective volume fractions as functions of the temperature obtained by thermodynamic equilibrium calculations.

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Figure 8:

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1275 °C 1400 °C 1600 °C 1760 °C 1800 °C

N content of liquid phase in wt% 0.465 0.391 0.390 0.387 0.368

Figure 9:

N and C content of the investigated steel in the gas-atomized condition and the LPBF-built condition (given in wt%).

Tensile experiments To investigate the influence of the L-PBF process on the material behavior under a quasistatic load, tensile experiments were conducted comparing materials produced by L-PBF and by hot isostatic pressing. L-PBF-built and subsequently solution-annealed specimens were also tested. Only specimens built under an Ar atmosphere were investigated in the tensile experiments because they

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had a slightly lower porosity compared to the specimens built under an N2 atmosphere. The microstructure and chemical composition of both batches do not show any other significant differences.

Representative stress-strain curves for X40MnCrMoN19-18-1 stainless steel in different processing conditions and for 316L in L-PBF-built condition [53].

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As can be seen from the results of the tensile tests, the investigated steel has high strength and high elongation at fracture values in all investigated conditions (Table 4, Figure 10). For comparison, results from tensile tests on L-PBF-built 316L stainless steel are given in Table 4 [53]. These results were obtained using the same L-PBF machine and the same tensile testing setup as in this work and therefore possess very good comparability. Compared to 316L, which is commonly used in L-PBF, the high strength of X40MnCrMoN19-18-1 steel is a result of the strong solid-solution strengthening with a simultaneously highly metallic atomic bonding character induced by C and N. In as-built condition, the investigated steel possesses significantly higher strength than 316L with nearly the same elongation at fracture value. However, the mechanical properties of the investigated steel strongly vary depending on the manufacturing route. The conducted HIP process enables production of dense samples that have only small micropores along the former particle boundaries (Figure 2). Furthermore, the chosen temperature of 1150 °C and the gas quenching inside the HIP furnace lead to the formation of a single-phase austenitic microstructure. For this reason, the HIP-densified steel has a high tensile strength and a moderate yield strength at a simultaneously high elongation at fracture. In contrast to the HIP state, the L-PBF process includes complete re-melting of the powder followed by rapid cooling of the small melt pools that are created by the focused laser radiation [22,54]. Thus, the microstructure of the L-PBF-built specimens is characterized by very fine grains, in which a subgrain structure is found (Figure 5). Various studies on the microstructure of L-PBF-built austenitic stainless steels report that the subgrain structure consists of dendrites with microsegregations in the interdendritic spaces [26,47]. In addition to local enrichment in alloying elements, a high dislocation density is found in these regions [55]. Due to the high dislocation density in the L-PBF-built condition, the ductility of the steel is reduced (Table 4). Simultaneously, the yield strength is considerably increased. Since the dislocation network inside the interdendritic regions, which separate dendrites with slightly different crystallographic orientations, is similar to small-angle grain boundaries, the high yield strength might be correlated to the Hall-Petch relation [47]. This assumption is supported by the decrease in the yield strength after solution annealing of the L-PBFbuilt steel. During the heat treatment, the dislocation network as well as the microsegregations vanish due to the material's recovery, recrystallization, diffusion, and defect annihilation phenomena. In accordance with the Hall-Petch relation, the yield strength decreases due to the

lowered density of interfaces in the material. These findings are in line with the results from Wang et al., who report a strong influence of the subgrain structure in L-PBF-built 316L austenitic stainless steel on the material’s strength [28]. Moreover, the precipitates present in the as-built condition are dissolved at the solution-annealing temperature, which contributes to the ductility of the material. After solution annealing, the investigated steel shows both high yield strength and high tensile strength, as well as high elongation at fracture values. Due to the shorter time at elevated temperatures in the solution annealing step (30 min, 1180 °C) compared to the HIP step (3 h, 1150 °C), the solution annealed material reveals smaller grains in large areas than the HIPed material. For this reason, the yield strength of the investigated steel is higher in the L-PBF-built and solution annealed condition than in the HIP condition. The mechanical behavior of the L-PBF-built and heattreated steel can be understood as a compromise between the as-built condition and the HIPdensified material.

Results of tensile tests performed on the investigated steel in different tested conditions.

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Table 4:

Yield strength Rp0.2 in MPa

Tensile strength Rm in MPa

X40MnCrMoN19-18-1 L-PBF Ar X40MnCrMoN19-18-1 L-PBF + soln. ann. X40MnCrMoN19-18-1 HIP 316L [53]

689 ± 68 587 ± 8

Elongation at fracture in % 29.4 ± 8.2 67.2 ± 1.3

1014 ± 0 573 ± 19

69.9 ± 3.5 31.9 ± 4.3

79.6 ± 3.8 38.5 ± 4.4

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Conclusions

568 ± 0 440 ± 22

985 ± 107 1057 ± 36

Uniform elongation in % 27.1 ± 7.5 64.4 ± 3.4

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Material

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In this work, austenitic X40MnCrMoN19-18-1 steel was processed by L-PBF, and the resulting microstructure and associated mechanical properties were compared to those of hot isostatically pressed X40MnCrMoN19-18-1. Several conclusions can be drawn from the results:

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1. High-strength austenitic X40MnCrMoN19-18-1 steel containing C and N can be processed by means of L-PBF. It was shown that it is possible to densify the steel up to approx. 99.8 %. Use of nitrogen instead of argon as the process gas leads to a slight increase in porosity, which might be due to different flow characteristics of the gases and the differences associated with the removal of contaminants from the laser/powder interaction zone. 2. Regardless which process gas is used, the nitrogen content of the steel is reduced slightly during the L-PBF process. It is possible that high melt pool temperatures, as reported in literature, induce a decrease in the solubility of nitrogen in the melt and cause it to degas. 3. The microstructure of the L-PBF-built steel consists of fine grains with a cellular subgrain structure with microsegregations in the interdendritic regions. In contrast to conventional LPBF-built austenitic steels, the investigated steel further reveals fine precipitates of Cr2N, Mo2(C,N) and Cr23C6 type. The precipitates, as well as the microsegregations are removed by a short solution-annealing treatment. In future investigations, the precipitates should be investigated by means of high-resolution investigation techniques, such as transmission electron microscopy (TEM). 4. The investigated steel has a high strength in both the as-built and in the subsequently solution-annealed condition. The elongation at fracture value is lowest for the as-built

material and increases after solution annealing. The strength of the L-PBF-built steel is comparable to the material produced by HIP, but it has lower elongation at fracture values. Author Contribution Statement Article: Microstructure and properties of high-strength C+N austenitic stainless steel processed by laser powder bed fusion Journal: Additive Manufacturing Authors: Boes, Johannes; Röttger, Arne; Theisen, Werner

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Boes, Johannes: Conceptualization; Data curation; Formal analysis; Investigation; Methodology; Validation; Visualization; Writing – Original Draft; Writing – review & editing Röttger, Arne: Conceptualization; Funding acquisition; Methodology; Project administration; Supervision; Writing – review & editing Theisen, Werner: Conceptualization; Funding acquisition; Project administration; Resources; Supervision; Writing – review & editing

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Acknowledgements

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We gratefully acknowledge financial support provided by the German Science Foundation (DFG) within the framework of the research project “TH-531/20-1 - Mechanism-based assessment of the influence of powder production and process parameters on the microstructure and the deformation behavior of SLM-compacted C + N steels in air and in corrosive environments”.

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References

Jo

ur

na

[1] V.G. Gavriljuk, H. Berns, High Nitrogen Steels: Structure, Properties, Manufacture, Applications, Springer, Berlin, Heidelberg, 1999. [2] J. Frehser, C. Kubisch, Metallurgie und Eigenschaften unter hohem Druck erschmolzender stickstoffhaltiger Stähle, BHM (1963) 369–380. [3] K.J. Irvine, T. Gladman, F.B. Pickering, The strength of austenitic stainless steels, J Iron Steel I (1969) 1017–1028. [4] P.J. Uggowitzer, M. Harzenmoser, Strengthening of austenitic stainless steels by nitrogen, High Nitrogen Steels—HNS 88 (1988) 174–179. [5] J.R. Kearns, The effect of nitrogen on the corrosion resistance of austenitic stainless alloys containing molybdenum, JMES 7 (1985) 16–26. [6] H.J. Grabke, High Nitrogen Steels. The Role of Nitrogen in the Corrosion of Iron and Steels, ISIJ International 36 (1996) 777–786. [7] C.L. Briant, Effects of nitrogen and cold work on the sensitization of austenitic stainless steels Final report, United States, 1982. [8] C. Escher, H. Berns, Druckaufgestickte Stähle für Verbrennungsmotoren: Schlußbericht ; BMBFProjekt (FKZ 03N3003B5) ; [Abschlußdatum des Vorhabens: 28.02.1999], Universitätsbibliothek u. Technische Informationsbibliothek; Ruhr-Univ. Inst. für Werkstoffe Lehrstuhl Werkstofftechnik, Hannover, Bochum, 2000. [9] V.G. Gavriljuk, B.D. Shanina, H. Berns, On the correlation between electron structure and short range atomic order in iron-based alloys, Acta Materialia 48 (2000) 3879–3893.

Jo

ur

na

lP

re

-p

ro of

[10] V.G. Gavriljuk, B.D. Shanina, H. Berns, Ab initio development of a high-strength corrosionresistant austenitic steel, Acta Materialia 56 (2008) 5071–5082. [11] V.G. Gavriljuk, V.N. Shivanyuk, B.D. Shanina, Change in the electron structure caused by C, N and H atoms in iron and its effect on their interaction with dislocations, Acta Materialia 53 (2005) 5017–5024. [12] V.G. Gavriljuk, S.P. Efimenko, Y.E. Smuk, S.Y. Smuk, B.D. Shanina, N.P. Baran, V.M. Maksimenko, Electron-spin-resonance study of electron properties in nitrogen and carbon austenites, Phys. Rev. B 48 (1993) 3224–3231. [13] B.D. Shanina, V.G. Gavrilyuk, A.A. Konchits, S.P. Kolesnik, A.V. Tarasenko, Exchange interaction between electron subsystems in iron-based F.C.C. alloys doped by nitrogen or carbon, physica status solidi (a) 149 (1995) 711–722. [14] V.G. Gavriljuk, J. Rawers, B.D. Shanina, H. Berns, Nitrogen and Carbon in Austenitic and Martensitic Steels, MSF 426-432 (2003) 943–950. [15] B.D. Shanina, V.G. Gavriljuk, H. Berns, F. Schmalt, Concept of a new high-strength austenitic stainless steel, Steel Research 73 (2002) 105–113. [16] H. Zitter, L. Habel, Zur Löslichkeit des Stickstoffs in Reineisen und austenitischen Chrom-NickelStählen, Archiv für das Eisenhüttenwesen 44 (1973) 181–188. [17] Z. Jiang, H. Li, Z. Chen, Z. Huang, D. Zou, L. Liang, The Nitrogen Solubility in Molten Stainless Steel, steel research international 76 (2005) 740–745. [18] H. Berns, W. Theisen, Eisenwerkstoffe: Stahl und Gusseisen, 4th ed., Springer-Verlag, s.l., 2008. [19] M. Talha, C.K. Behera, O.P. Sinha, A review on nickel-free nitrogen containing austenitic stainless steels for biomedical applications, Materials science & engineering. C, Materials for biological applications 33 (2013) 3563–3575. [20] K.H. Lo, C.H. Shek, J.K.L. Lai, Recent developments in stainless steels, Materials Science and Engineering: R: Reports 65 (2009) 39–104. [21] H. Fayazfar, M. Salarian, A. Rogalsky, D. Sarker, P. Russo, V. Paserin, E. Toyserkani, A critical review of powder-based additive manufacturing of ferrous alloys: Process parameters, microstructure and mechanical properties, Materials & Design 144 (2018) 98–128. [22] P.A. Hooper, Melt pool temperature and cooling rates in laser powder bed fusion, Additive Manufacturing 22 (2018) 548–559. [23] Du Yang, X. You, F. Qiao, L. Guo, Z. Liu, A model for predicting the temperature field during selective laser melting, Results in Physics 12 (2019) 52–60. [24] E. Liverani, S. Toschi, L. Ceschini, A. Fortunato, Effect of selective laser melting (SLM) process parameters on microstructure and mechanical properties of 316L austenitic stainless steel, Journal of Materials Processing Technology 249 (2017) 255–263. [25] A. Röttger, K. Geenen, M. Windmann, F. Binner, W. Theisen, Comparison of microstructure and mechanical properties of 316L austenitic steel processed by selective laser melting with hotisostatic pressed and cast material, Materials Science and Engineering: A 678 (2016) 365–376. [26] K.G. Prashanth, J. Eckert, Formation of metastable cellular microstructures in selective laser melted alloys, Journal of Alloys and Compounds 707 (2017) 27–34. [27] Di Wang, C. Song, Y. Yang, Y. Bai, Investigation of crystal growth mechanism during selective laser melting and mechanical property characterization of 316L stainless steel parts, Materials & Design 100 (2016) 291–299. [28] Y.M. Wang, T. Voisin, J.T. McKeown, J. Ye, N.P. Calta, Z. Li, Z. Zeng, Y. Zhang, W. Chen, T.T. Roehling, R.T. Ott, M.K. Santala, P.J. Depond, M.J. Matthews, A.V. Hamza, T. Zhu, Additively manufactured hierarchical stainless steels with high strength and ductility, Nature Materials 17 (2017) 63 EP -. [29] K. Saeidi, X. Gao, Y. Zhong, Z.J. Shen, Hardened austenite steel with columnar sub-grain structure formed by laser melting, Materials Science and Engineering: A 625 (2015) 221–229.

Jo

ur

na

lP

re

-p

ro of

[30] A.J. Birnbaum, J.C. Steuben, E.J. Barrick, A.P. Iliopoulos, J.G. Michopoulos, Intrinsic strain aging, Σ3 boundaries, and origins of cellular substructure in additively manufactured 316L, Additive Manufacturing 29 (2019) 100784. [31] N.P. Lavery, J. Cherry, S. Mehmood, H. Davies, B. Girling, E. Sackett, S.G.R. Brown, J. Sienz, Effects of hot isostatic pressing on the elastic modulus and tensile properties of 316L parts made by powder bed laser fusion, Materials Science and Engineering: A 693 (2017) 186–213. [32] S. Riedner, Höchstfeste nichtrostende austenitische CrMn-Stähle mit (C+N), 2011. [33] H. Berns, V. Gavriljuk, B. Shanina, Intensive Interstitial Strengthening of Stainless Steels, Adv. Eng. Mater. 10 (2008) 1083–1093. [34] A. Röttger, K. Geenen, M. Windmann, F. Binner, W. Theisen, Comparison of microstructure and mechanical properties of 316 L austenitic steel processed by selective laser melting with hotisostatic pressed and cast material, Materials Science & Engineering A (2016) 365–376. [35] J. Boes, A. Röttger, C. Mutke, C. Escher, W. Theisen, Microstructure and mechanical properties of X65MoCrWV3-2 cold-work tool steel produced by selective laser melting, Additive Manufacturing 23 (2018) 170–180. [36] J. Krell, A. Röttger, K. Geenen, W. Theisen, General investigations on processing tool steel X40CrMoV5-1 with selective laser melting, Journal of Materials Processing Technology 255 (2018) 679–688. [37] Metallic materials - Tensile testing - Part 1: Method of test at room temperature (ISO 68921:2016); German version EN ISO 6892-1:2016, Beuth Verlag GmbH, Berlin. [38] F. Nilsén, I. Aaltio, Y. Ge, T. Lindroos, S.P. Hannula, Characterization of Gas Atomized Ni-Mn-Ga Powders, Materials Today: Proceedings 2 (2015) S879 - S882. [39] P.W. Lee (Ed.), Powder metal technologies and applications, 10th ed., ASM International, Materials Park, Ohio, 2002. [40] S. Dietrich, M. Wunderer, A. Huissel, M.F. Zaeh, A New Approach for a Flexible Powder Production for Additive Manufacturing, Procedia Manufacturing 6 (2016) 88–95. [41] E.O. Olakanmi, Effect of mixing time on the bed density, and microstructure of selective laser sintered (sls) aluminium powders, Mat. Res. 15 (2012) 167–176. [42] P.W. Cleary, M.L. Sawley, DEM modelling of industrial granular flows: 3D case studies and the effect of particle shape on hopper discharge, Applied Mathematical Modelling 26 (2002) 89– 111. [43] I. Yadroitsev, P. Krakhmalev, I. Yadroitsava, S. Johansson, I. Smurov, Energy input effect on morphology and microstructure of selective laser melting single track from metallic powder, Journal of Materials Processing Technology 213 (2013) 606–613. [44] P. Bidare, I. Bitharas, R.M. Ward, M.M. Attallah, A.J. Moore, Fluid and particle dynamics in laser powder bed fusion, Acta Materialia 142 (2018) 107–120. [45] D. Dai, D. Gu, Effect of metal vaporization behavior on keyhole-mode surface morphology of selective laser melted composites using different protective atmospheres, Applied Surface Science 355 (2015) 310–319. [46] Di Wang, S. Wu, F. Fu, S. Mai, Y. Yang, Y. Liu, C. Song, Mechanisms and characteristics of spatter generation in SLM processing and its effect on the properties, Materials & Design 117 (2017) 121–130. [47] Y. Zhong, L. Liu, S. Wikman, D. Cui, Z. Shen, Intragranular cellular segregation network structure strengthening 316L stainless steel prepared by selective laser melting, Journal of Nuclear Materials 470 (2016) 170–178. [48] F. Vanderschaeve, R. Taillard, J. Foct, Discontinuous precipitation of Cr2N in a high nitrogen, chromium-manganese austenitic stainless steel, Journal of Materials Science 30 (1995) 6035– 6046.

Jo

ur

na

lP

re

-p

ro of

[49] N. Krasokha, S. Weber, S. Huth, K. Zumsande, W. Theisen, Gas‐Solid Interactions During Nonisothermal Heat Treatment of a High-Strength CrMnCN Austenitic Steel Powder: Influence of Atmospheric Conditions and Heating Rate on the Densification Behavior, Metallurgical and Materials Transactions A 43 (2012) 4237–4246. [50] V.G. Gavriljuk, A.L. Sozinov, A.G. Balanyuk, S.V. Grigoriev, O.A. Gubin, G.P. Kopitsa, A.I. Okorokov, V.V. Runov, Effect of carbon and nitrogen on chemical homogeneity of fcc ironbased alloys, Metallurgical and Materials Transactions A 28 (1997) 2195–2199. [51] A. Sieverts, G. Zapf, H. Moritz, Die Löslichkeit von Wasserstoff, Deuterium und Stickstoff in Eisen, Zeitschrift für physikalische Chemie (1938) 19–37. [52] S.A. Khairallah, A.T. Anderson, A. Rubenchik, W.E. King, Laser powder-bed fusion additive manufacturing: Physics of complex melt flow and formation mechanisms of pores, spatter, and denudation zones, Acta Materialia 108 (2016) 36–45. [53] K. Geenen, Werkstofftechnische Charakterisierung austenitischer und martensitischer Stähle nach dem selektiven Laserschmelzen. Dissertation, 1st ed., 2018. [54] U.S. Bertoli, G. Guss, S. Wu, M.J. Matthews, J.M. Schoenung, In-situ characterization of laserpowder interaction and cooling rates through high-speed imaging of powder bed fusion additive manufacturing, Materials & Design 135 (2017) 385–396. [55] S. Bahl, S. Mishra, K.U. Yazar, I.R. Kola, K. Chatterjee, S. Suwas, Non-equilibrium microstructure, crystallographic texture and morphological texture synergistically result in unusual mechanical properties of 3D printed 316L stainless steel, Additive Manufacturing 28 (2019) 65–77.