Proton irradiation and characterization of additively manufactured 304L stainless steels

Journal of Nuclear Materials 531 (2020) 152007

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Proton irradiation and characterization of additively manufactured 304L stainless steels B.P. Eftink a, *, J.S. Weaver b, c, J.A. Valdez a, V. Livescu a, D. Chen a, d, Y. Wang a, C. Knapp e, N.A. Mara b, f, S.A. Maloy a, G.T. Gray III a a

Los Alamos National Laboratory: MST-8 Materials Under Irradiation and Dynamic Extremes, P.O. Box 1663, Los Alamos, NM, 87545, USA Los Alamos National Laboratory: Center for Integrated Nanotechnologies, P.O. Box 1663, Los Alamos, NM, 87545, USA National Institute of Standards and Technology: Engineering Laboratory, 100 Bureau Drive, Gaithersburg, MD, 20899, USA d University of Houston: Department of Physics, 3507 Cullen Blvd, Houston, TX, 77204, USA e Los Alamos National Laboratory: Sigma Division, P.O. Box 1663, Los Alamos, NM, 87545, USA f University of Minnesota-Twin Cities: Chemical Engineering and Materials Science, 421 Washington Ave. SE, Minneapolis, MN, 55455, USA b c

h i g h l i g h t s  Microstructures evolve similarly for irradiated AM and wrought 304L  Nanohardness values are similar for all samples after irradiation.  Proton irradiation of AM and wrought 304L

a r t i c l e i n f o

a b s t r a c t

Article history: Received 15 August 2019 Received in revised form 16 January 2020 Accepted 16 January 2020 Available online 21 January 2020

Irradiations were performed with 1.5 MeV protons to 0.6 dpa at 40e150  C on additively manufactured (AM) 304L stainless steel and the changes in microstructure and mechanical behavior after irradiation were compared to wrought 304L stainless steel. All microstructural and hardness results after irradiation suggest the samples evolve toward a similar state, despite significant differences in the unirradiated microstructures and hardness values. A TEM and nanoindentation-based investigation of before and after proton irradiation at 40e150  C is presented. Results are interpreted in terms of initial dislocation content, dislocation structures, and microstructural and chemical homogeneity. Published by Elsevier B.V.

Keywords: Additive manufacturing Dislocation structures Irradiated metals TEM X-ray diffraction

1. Introduction Materials utilized in extreme radiation environments are required to maintain a level of mechanical integrity as damage develops in the microstructure with irradiation exposure, at elevated temperature, and in a corrosive environment. To design advanced concept nuclear reactors with better power efficiency, it is important to design materials that can tolerate extreme temperature and irradiation regimes [1]. Additive manufacturing is one method of producing complex parts that, if the microstructure can

* Corresponding author. E-mail address: [email protected] (B.P. Eftink). https://doi.org/10.1016/j.jnucmat.2020.152007 0022-3115/Published by Elsevier B.V.

be adequately controlled during processing, offers a potentially promising pathway for manufacturing damage-tolerant reactor components. Such components not only have complex microstructures that differ from conventionally produced metals, but also microstructural variability within the same part [2e4]. Parameters used in the AM build process include power input, beam size and method (i.e., powder-bed, directed energy deposition, etc.), and have a profound impact on the resulting microstructures and therefore the properties of the component [5e8]. As an example, it is possible to vary the microstructural feature lengths, crystallographic texture, and porosity by simply varying the solidification rate during AM processing. Numerous publications have covered additively manufactured austenitic stainless steels in terms of build parameters,

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microstructures and properties [9e22]. For example, AM austenitic stainless steels commonly have cellular microstructural features and these features have been attributed to increased tensile strength and equivalent or increased ductility [19,23]. It has also been shown that a cellular microstructure in a 316L stainless steel can have segregation of Mo to the cell boundaries likely resulting in the reduced corrosion resistance of AM 316L compared to wrought 316L stainless steel [12]. Processing parameters have even been shown capable of tuning the mechanical properties, for example, heat input for AM 304L has shown an impact on yield strength, ultimate tensile strength, as well as ductility [20]. Mechanical testing of neutron irradiated 316L has also been reported, however, without a comparison to wrought material [24]. Microstructural response after ion irradiation of AM 316L above 300  C has been investigated, showing instability of cellular microstructures [25] and modestly increased swelling compared to wrought material [26]. Some properties can be improved in AM austenitic steels, such as yield strength, while others, such as corrosion resistance, may be less favorable compared to conventionally prepared wrought material. For wrought austenitic stainless steels, extensive research has been performed on how initial microstructure impacts the evolution of properties caused by irradiation [27,28]. One microstructural feature known to affect a materials response to radiation is initial dislocation density [29]. Annealed compared to cold worked 316 stainless steel results in a different rate of irradiation hardening/ softening to the final, equilibrium, hardness dependent on irradiation temperature [29]. At 427  C, annealed 316 stainless steel hardens at a faster rate with neutron fluence than 20% cold-worked 316 stainless steel. At a higher irradiation temperature, 538  C, 20% cold-worked 316 stainless steel softens with neutron fluence and annealed 316 stainless steel hardens to eventually approach the same yield strength. Grain size is also an important consideration and austenitic stainless steels with smaller grains have been shown to have lower swelling rates [30,31]. Additively manufactured 304L austenitic stainless steel has microstructural features including increased lattice dislocation content, and fine grain or subgrain size that could be beneficial to microstructural stability when irradiated. Wrought material with a high dislocation density was chosen to provide a comparison between AM and traditionally produced material with high number densities of irradiation defect sinks. This work investigates the effect of proton irradiation at 40e150  C on several microstructures of AM 304L stainless steel as compared to wrought. This temperature regime is relevant to certain reactor components, for example, the primary coolant supply adapter [32]. The focus of this study is to investigate 1) the effects of low dose and temperature irradiation (using protons) on the microstructures in AM 304L stainless steel and 2) mechanical properties through nanohardness measurements. From prior studies, it should be expected that the microstructures and hardness should evolve to a similar equilibrium despite differences in the unirradiated material. X-ray diffraction was also used to compare the overall phase distribution within each 304L stainless steel sample before and after irradiation. 2. Experimental procedures Additively manufactured materials were produced using 304L stainless steel powders obtained from Carpenter Powder Products in Sweden. Two powder types were used, a pedigreed with low sulfur (S) and a commercial off the shelf (COTS) with higher S content powder, both Micro-Melt 304L SS. The compositions are provided in Table 1. Two machines, a laser directed-energydeposition (DED) Optomec LENS MR-7 and laser powder-bed Electro-Optical Systems (EOS) M280, were employed to produce

the AM samples. Both high and low S powders were used in the production of components using the DED machine whereas only the low S powder was used to manufacture samples in the powderbed machine. Powder size for the DED and powder-bed machines were 44e106 and 15e45 mm, respectively. The AM build with the EOS M280 powder-bed machine was on a 50.5 mm thick AISI 304L baseplate while the baseplate thickness for the Optomec LENS MR-7 builds was 12.7 mm. The EOS machine parameters used were the EOS proprietary PH-1 setting with layer heights of 20 mm. Parameters such as the beam power and raster speed are unknown, however, the beam raster uses a rotational rectilinear hatching. Settings on the Optomec LENS MR-7 are known, using a 1070 nm fiber laser, with a raster speed of 1.016 m/ min, power of 800 W, 9.5 mm overfocus, z-step of 0.76 mm, hatch spacing of 1.02 mm, hatch pattern of 45 e135 -225 e315 , mass feed of 33.7 g/min, powder efficiency of 20%, deposition rate of 400 g/h, ~40% hatch overlay, and in an Argon environment with <5 ppm O2. Compositions, post build, were determined through outside vendors and presented in Table 2. The wrought material was processed by high energy rate forging [33]. Ion irradiations were performed at the Ion Beam Materials Lab at Los Alamos National Laboratory. To prepare the samples for ion irradiation, the surfaces were mechanically polished with a final step of 0.01 mm colloidal silica. The samples were irradiated with 1.5 MeV protons to an irradiation dose of 0.6 displacements per atom (dpa) at a sample depth of 5 mm which is in the plateau region (peak dose was 12 dpa). The irradiation dose was calculated with SRIM using the Kinchin and Pease model on a pure Fe target (assuming a Fe threshold displacement energy of 40 eV) [34]. Irradiation was conducted on all samples at the same time by defocusing the beam over a 6 by 6 mm area. A small-step beam wobbling was also used during the irradiation to improve the irradiation uniformity. The samples were attached to a Cu mount using Ag paint, and temperature was measured using a thermocouple on the Cu mount. The Cu mount was water cooled. The chamber vacuum steadily improved from 6:2  10 7 to 7:0  10 8 torr during the roughly 76.5 h long irradiation. The beam current density was increased from 0.03 to 0.18 A m 2 during the irradiation, and the shape of the beam was monitored. Temperature was dictated by beam heating, and variations occurred due to the variability of the beam current. The temperature was measured during the first 36 h between 55 and 95  C, the next 24 h between 120 and 153  C and the last 16.5 h between 70 and 120  C. However, all irradiated specimens experienced the same temperature variation, thus allowing comparison between them. Fluence for the irradiation was 1:46  1019 ions per cm2. Fig. 1 shows the calculated damage distribution as a function of distance from the sample surface. The dose reached a maximum value of 12 dpa at a depth of 12 mm into the samples; however, this region of the samples was not considered in the study because of the high dose sensitivity to depth as shown in Fig. 1 and the high injected interstitial content. Electron backscatter diffraction (EBSD) crystal orientation maps were acquired using a FEI Inspect SEM with an EDAX detector using TSL software. Samples were prepared for EBSD by mechanical polishing before an electropolish etch in 40% nitric acid in water, first for 6 s at 6 V and next 4 s at 3 V, at room temperature. For the DED samples and powder-bed sample, step sizes of 5 and 2 mm were used, respectively. A step size of 2 mm was used for the wrought material EBSD scan. Sample preparation for transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) was performed by focused ion beam (FIB) milling using a FEI Helios 600i. FIB milling was chosen to extract samples before irradiation and then to extract samples after irradiation from an adjacent area. This was done because in the AM samples, microstructure can change significantly depending on the

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Table 1 Chemical compositions in weight percent of the two 304L stainless steel powders used for additively manufacturing the samples.

Low S High S

C

Si

Mn

P

S

Cr

Ni

Mo

Cu

N

O

Fe

0.015 0.020

0.53 0.77

1.5 1.5

0.012 0.008

0.003 0.009

18.4 18.5

9.8 9.8

0.0 0.0

0.0 0.0

0.05 0.08

0.019 0.023

Bal. Bal.

Table 2 Chemical compositions in weight percent of the three 304L stainless steel samples additively manufactured, and the wrought sample.

Wrought Low S DED High S DED Low S Powder-Bed

C

Si

Mn

P

S

Cr

Ni

Mo

Cu

N

O

Fe

0.015 0.012 0.018

0.57 0.68 0.84

1.43 1.59 1.53

0.019 0.010 0.017

0.002 0.005 0.007

18.47 18.10 18.46

10.02 9.57 9.30

0.11 0.008 0.036

0.18 0.011 0.017

0.05 0.051 0.074

<0.01 0.022 0.013

Bal. Bal. Bal.

0.020

0.63

1.43

0.007

0.004

18.32

10.03

0.015

0.008

0.05

0.03

Bal.

nanoindentation. 3. Results 3.1. Initial microstructures

Fig. 1. Dose in dpa as a function of depth into the sample.

position in the sample. Focused ion beam milling was started at 30 keV then 5 keV and final milling at 2 keV. The samples were then milled further using a Gatan PIPS II at 1 keV and cryogenic temperatures to minimize sample preparation artifacts. TEM and STEM were performed on either a FEI Tecnai F-30 or FEI Titan, both operating at 300 keV. Nanoindentation was performed on a Keysight G200 Nanoindenter. Indents with a diamond Berkovich (pyramidal) tip were performed using the continuous stiffness measurement (CSM) at 2 nm displacement amplitude and frequency of 45 Hz to a final displacement of 1000 nm and a constant strain rate (loading rate over the load) of 0.05 s 1. A minimum of ten indents for the wrought samples and 22 for additive samples were used to determine the average Young’s modulus and Hardness according to the Oliver-Pharr method [35]. The modulus remained relatively constant with indentation depth for all samples, and the average was determined over a displacement range of 200e1000 nm. The average hardness was determined over a displacement range of 900e1000 nm which approximately corresponds to a plastic zone depth of 4e10 mm [36]. For the irradiated samples, nanoindentation was performed on the irradiated surface. It is assumed that the plastic zone depth is closer to 4 mm in the irradiated material and closer to 10 mm in the unirradiated material based on observations in similar ion irradiated materials [37,38]. Thus, we expect the irradiated microstructures of TEM foils taken between 4 and 6 mm from the surface to be representative of the material probed by

Different processing methods (wrought, DED, and powder-bed) of the 304L were found to result in variations in the microstructure. The overall microstructures are shown in the EBSD orientation maps presented in Fig. 2. Both DED samples, Fig. 2(a) and (b), displayed a combination of larger and elongated grains hundreds of mm long and smaller equiaxed grains tens of mm in diameter. A much finer structure is observed in the powder-bed material, Fig. 2(c). For a direct comparison to the DED material, the inset in Fig. 2(c) has the same scale as the DED orientation maps. The powder-bed material has a combination of larger and smaller grains, tens of mm and several mm in diameter, respectively. Wrought produced 304L is presented in Fig. 2(d) and shows equiaxed grains on the order of 100 mm in diameter. A detailed TEM characterization correlating to specific regions in the EBSD maps is not within the scope of this manuscript. Regions for TEM observation were chosen independently of the EBSD results and are presented below. Each AM sample had distinct microstructural features despite the fact that all were built with 304L powder. It was also common for an individual sample to exhibit multiple microstructures, necessitating extraction of the TEM foils from the same region before and after irradiation. To confirm the same microstructures were observed in the unirradiated and irradiated samples, the irradiated TEM samples were produced to include material past the irradiation depth, Supplementary Fig. S1. The three samples include: high S DED, low S DED, and low S powder-bed. A control sample of the conventionally processed wrought material was also investigated and is shown in the bright-field TEM micrograph in Fig. 3. The wrought material was characterized by dislocation structures including subgrains on the order of several hundred nanometers across. Austenite was the only phase observed in the wrought TEM samples. 3.1.1. 304L high S additively manufactured directed-energydeposition The microstructure of the DED processed sample using the 304L high S powder is an austenitic matrix with ferritic ellipsoids on the order of 400 nm in diameter, marked by arrows in Fig. 4. The ferritic ellipsoids are extended in the TEM foil normal direction as shown in Fig. 4(a) and (b) which is why they appear circular. The diffraction pattern in Fig. 4(c) is a selected area diffraction pattern of the

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Fig. 2. EBSD orientation maps for (a) DED high S, (b) DED low S, (c) powder-bed low S and (d) wrought 304L. Insets of (c) and (d) are the same scale as (a) and (b) for direct comparison.

BCC ferrite marked by the arrow in Fig. 4(b). Chemical mapping using STEM/EDX of the region imaged using STEM bright-field in Fig. 5(a) is presented in Fig. 5(b)-(f). Fig. 5(c), reveals Ni depletion in and surrounding the regions of ferrite. It is also observed that the ferrite is primarily Cr rich, but also to a lesser extent Mn and O rich compared to the austenite matrix, Fig. 5(d)e(f). In the austenitic matrix, dislocations are evenly distributed. Grain/subgrain size was not determined as the entire electron transparent region, 12 by 12 mm, contained a single crystallographic orientation. The austenite matrix contains planar features which are found near the ferrite in the Ni depleted regions, the planar features are marked by arrowheads in Fig. 4. The planar features have diffraction contrast consistent with stacking faults. Decreasing Ni content, as is the case surrounding the ferrite, decreases the stacking fault energy in austenitic stainless steels [39], in addition to decreasing the stability of austenite.

Fig. 3. Bright-field TEM micrograph of wrought 304L stainless steel produced by conventional methods.

3.1.2. 304L Low S additively manufactured directed-energydeposition Fig. 6(a) presents a representative TEM bright-field micrograph of the microstructure observed in the DED produced sample using the low S powder. This image shows irregularly shaped sub-grains on the order of microns with boundaries consisting of high densities of dislocations. Second phases observed in this sample include ferrite as well as Mn, Si and O rich precipitates. The regions

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Fig. 4. (a) and (b) Bright-field transmission electron micrographs of additively manufactured 304L stainless steel, DED high S. (c) Selected area diffraction pattern of ferrite phase marked with arrow from (b). Arrows mark ferrite and arrowheads mark planar features.

Fig. 5. (a) Bright-field STEM electron micrograph and (b)e(f) x-ray analysis chemical maps of DED high S AM sample.

of ferrite were on the order of 400 nm while the Mn and Si rich precipitates were on the order of 150 nm. The precipitates are likely MnSiO3 [40]. Fig. 6(b) is a high angle annular dark field (HAADF) micrograph which shows the different phases, ferrite is marked with arrowheads while the Mn, Si and O rich precipitates are marked with arrows. The second phases where characterized using x-ray chemical maps presented in Fig. 7, which corresponds to the same sample region as Fig. 6(b). Ferrite was characteristically depleted of Ni similar to the observation in Fig. 5. Additionally, the Mn, Si and O rich precipitates have S content localized on an edge of the precipitate, which is marked by the arrowheads in Fig. 7, similar to prior work [41]. 3.1.3. 304L Low S additively manufactured powder-bed Two different microstructures were observed in the sample manufactured using the low S 304L powder in the powder-bed machine. The first is a dislocation cell structure with a cell

diameter of 400 nm contained within grains whose diameter is on the order of microns, shown in Fig. 8(a). The dislocation cells are either spherical or elongated, however, the smallest dimension is still the 400 nm diameter. The other microstructure observed in the sample is more disordered, and contains a relatively higher though less organized dislocation density, shown in Fig. 8(b). For the cell structure, at the cell boundaries, there are precipitates, cavities or both. Spherical features corresponding to either the precipitates or cavities are present in the TEM brightfield images in Fig. 9(a) and (b). The features appear to be cavities from the contrast change observed with TEM and over- and underfocusing in Fig. 9(a) and (b) [42,43], however, chemical segregation is observed to some of the larger features as shown in Fig. 9(c)e(i). Black circles are placed on regions of elevated Si and Mn content. The chemical segregation of Si, Mn, and possibly S suggests the features are likely precipitates, or possibly cavities with chemical segregation on the surface. The smaller features were not

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Fig. 6. (a) Bright-field electron micrograph of additively manufactured 304L stainless steel, DED low S. (b) HAADF image of additively manufactured 304L stainless steel, DED low S, of a different region than (a). Arrowheads mark ferrite and arrows mark Mn and Si rich precipitates.

Fig. 7. X-ray analysis chemical maps of DED low S AM sample. Arrowheads mark location of S rich regions in the SieK and SeK maps. Arrowheads mark S rich regions in (e)e(g) showing the location of S with respect to the Si and Mn rich oxides.

characterized by clear chemical segregation, likely due to the limits in the x-ray analysis, though are also likely precipitates. Average feature size, from a sample of 100 features, is 10 nm with a standard deviation of 5 nm. A histogram of feature size is shown in Fig. 10. Volume fraction of features is around 0.1% as calculated from the three-dimensional model presented in Supplementary Video 1 which was created with stereographic methods [44e48]. Supplementary video related to this article can be found at https://doi.org/10.1016/j.jnucmat.2020.152007 The interiors of the dislocation cells are relatively free of dislocations. Additionally, differences in composition were found using x-ray chemical analysis between the cell interiors and boundaries, with the boundaries enriched in Cr, Mn and Ni while depleted in Fe, which has been previously reported for additively manufactured 316L stainless steel [23].

3.2. Irradiated microstructures Irradiation with protons to 0.6 dpa at 40e150  C had the effect of reducing the lattice dislocation density in all of the samples, as shown in the bright-field TEM micrographs in Fig. 11. Bright-field TEM micrographs in Fig. 11 show a comparison of the (a)-(d) irradiated and (e)-(h) unirradiated microstructures. Each sample after irradiation was observed to contain irradiation induced dislocation loops, apparent as black dot damage. The dislocation cell structure is observed to be unstable in the low S powder-bed sample. Lattice dislocation content is observed to be lower in the high S DED, low S DED and wrought samples. Chemical depletion of Ni surrounding the ferrite was still observed as well as the planar features in the high S DED sample. In the low S DED and wrought samples, the retained lattice dislocation content was highest. The dislocation cell

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Fig. 8. (a) and (b) Bright-field TEM electron micrograph of additively manufactured 304L stainless steel, powder-bed low S.

Fig. 9. (a)(b) bright-field TEM and (c-i) x-ray chemical maps from the same region in low S powder-bed additively manufactured material. Over-under focus bright-field TEM imaging is shown in (a) and (b). Black circles indicate regions enriched in Si and Mn.

Fig. 10. Histogram of cavity or precipitate feature size in powder-bed low S 304L stainless steel.

structure in the low S powder-bed sample was observed to be unstable during irradiation and was eliminated. On the other hand, the precipitates were retained in the low S powder-bed sample. The microstructures observed after irradiation were confirmed to be similar to those presented of the unirradiated microstructures by observations past the irradiation depth in the TEM foils, Supplementary Fig. S1. Dislocation loops were seen to be generated from the irradiation in all samples. The dislocation loop density was measured by imaging at the [001] zone axis using bright-field STEM to capture all possible loop types, Fig. 12. Additively manufactured samples high S DED, low S DED, and low S powder-bed are presented in Fig. 13 (a), (b) and (c), respectively. The irradiated wrought 304L is presented in Fig. 12 (d). Dislocation loop type and size were not determined as the loops are too small to do so accurately. Foil

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Fig. 11. Bright-field TEM micrographs of irradiated 304L samples. (a)e(d) irradiated and (e)e(h) unirradiated. (a) and (e) DED high S, (b) and (f) DED low S, (c) and (g) powder-bed low S, (d) and (h) wrought. All images are in the <111> diffraction condition. The micrographs between the images of the same sample type are of the same scale.

Fig. 12. On-zone [001] bright-field STEM imaging of irradiated (a) DED high S, (b) DED low S, (c) powder-bed low S, and (d) wrought 304L stainless steel. Images taken between 4 and 6 mm from the sample surfaces.

thicknesses were measured using stereomicroscopy [44,48,49]. Loop number density will be impacted by surface effects and FIB preparation artifacts, however, the relative number density between the samples is sufficient for comparison. The loop number densities are presented in Table 3.

content in the powder, in order of lowest to highest hardness: DED, powder-bed, and wrought. Despite the difference between the hardening (the hardening was 5.5 times greater for the low S DED samples compared to the wrought), all of the samples had similar hardness after irradiation. All of the samples displayed hardness values within error of each other after irradiation.

3.3. Nanohardness 3.4. X-ray diffraction Nanohardness testing was performed on the samples before and after irradiation. On the irradiated samples this was conducted on the irradiated surface. As mentioned in the experimental procedures, the plastic zone depth is around 10 mm and 4 mm for the unirradiated and irradiated specimens, respectively. The average hardness values before and after irradiation, the change in average hardness values, as well as the number of indentations for each sample are presented in Table 4. Before irradiation there are differences in hardness. After irradiation the hardness values for all samples are similar. The initial hardness was observed to be dependent mostly on the manufacturing method independent of S

X-ray diffraction allows a macroscale analysis of the phase distribution in the samples before and after irradiation that is not as susceptible as TEM and nanohardness is to local microstructural variations. X-ray diffraction data is shown in Fig. 13 for all of the samples. In Table 5 values for peak locations and full-width at half max (FWHM) are presented. One peak, (110)a, corresponds to ferrite, and is marked by the arrows in Fig. 13, the peak is not observed for the wrought sample. The ferrite peak is not observed in the low S powder-bed sample before irradiation either, despite its presence after irradiation. Ferrite has previously been observed

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Fig. 13. X-ray diffraction before and after irradiation.

Table 3 Dislocation loop number density of the different 304L stainless steel samples irradiated to 0.6 dpa at 50e140  C with protons. High S DED Loops/m

3

6:22±:6 

Low S DED

1022

1:29±:1 

Low S powder-bed

1022

2:09±:2 

Wrought

1022

4:01±:4  1022

Table 4 Nanohardness values both before and after irradiation for the various 304L stainless steel samples. Values are the mean ± standard deviation. The number of indents is listed in parenthesis.

Initial hardness (GPa) Irradiated hardness (GPa) Change in hardness (GPa) Percent increase in hardness

High S DED

Low S DED

Low S powder-bed

Wrought

2.63 ± 0.13 (24) 4.27 ± 0.10 (24) 1.64 ± 0.23 62%

2.76 ± 0.30 (22) 4.48 ± 0.22 (23) 1.72 ± 0.52 62%

3.06 ± 0.23 (23) 4.19 ± 0.16 (22) 1.13 ± 0.39 37%

3.91 ± 0.16 (13) 4.22 ± 0.21 (10) 0.31 ± 0.37 8%

Table 5 Quantitative values for peak location and full-width at half max (FWHM), in degrees, for each measured peak of each sample both before and after irradiation. Error in the instrument is ±0.01. Wrought

DED high S

As-received

(111)g (110)a (200)g (220)g

Irradiated

As-received

FWHM

Peak (degrees)

FWHM

Peak (degrees)

FWHM

Peak (degrees)

FWHM

43.6 e 50.8 74.7

0.36 e 0.46 0.55

43.6 e 50.7 74.7

0.31 e 0.38 0.41

43.7 44.7 50.8 74.7

0.25 0.59 0.27 0.38

43.7 44.7 50.9 74.8

0.24 0.18 0.24 0.32

DED low S As-received

(111)g (110)a (200)g (220)g

Irradiated

Peak (degrees)

Powder-bed low S Irradiated

As-received

Irradiated

Peak (degrees)

FWHM

Peak (degrees)

FWHM

Peak (degrees)

FWHM

Peak (degrees)

FWHM

43.7 44.6 50.8 74.7

0.30 0.65 0.33 0.44

43.6 44.6 50.8 74.8

0.25 0.82 0.29 0.38

43.7 e 50.9 74.8

0.25 e 0.31 0.35

43.6 44.5 50.9 74.7

0.26 1.18 0.32 0.40

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in additively manufactured 316L stainless steel [11]. The ferrite was retained after irradiation. The FWHM decreases for each peak after irradiation except for all peaks in the powder-bed low S sample, and the (110)a peak corresponding to ferrite in the DED low S sample which show an increase in FWHM. Peak sharpening after irradiation is indicative of a decrease in dislocation density. With TEM, lattice dislocation density was observed to decrease after irradiation while irradiation induced dislocation loops increased. Peak location was consistent in all samples both before and after irradiation. 4. Discussion 4.1. Microstructure evolution under irradiation Austenitic stainless steels irradiated below 300  C exhibit microstructural changes including the formation of small, less than 5 nm diameter, dislocation loops and decreased lattice dislocation density [50,51]. The small dislocation loops, or black dot damage, are likely Frank loops [51]. At the temperature regime below 300  C, the microstructures in austenitic stainless steels reach an equilibrium in terms of irradiation induced dislocation loops and lattice dislocations at doses as low as 0.1 dpa [50]. The microstructural observations in this study are in agreement with the rapid evolution of microstructures. However, at 0.6 dpa irradiated below 150  C, each of the microstructures from the different processing methods had similar but not identical microstructures. For example, the dislocation loop number density varied from 1.29  1022 to 6.22  1022 loops per cubic meter. The equilibrium loop number density in 304 and 316 stainless steel has previously been found to be 1023 loops per cubic meter after irradiation [51]. Differences observed here may be explained by initial lattice dislocation density, which is qualitatively lowest for the microstructure shown for the High S DED and highest for the Low S DED. In the case of the lowest lattice dislocation microstructural region, the highest irradiation induced dislocation loop number density was found and the opposite is also true. This suggests initial lattice dislocation content has an impact on the dose required for microstructural equilibrium in the low temperature, less than 300  C, irradiation regime which is consistent with prior studies [29]. Another possible source in the differences of the irradiated microstructures could be due to chemical composition including of trace elements. Prior studies of austenitic stainless steels have revealed trace elements, particularly P, impact the transient swelling regime. This is thought to be due to vacancy diffusivity and precipitate evolution [52]. Due to the additive manufacturing processes, exact compositions are difficult to control which is observable from the differences in composition of the powder and the builds, Tables 1 and 2, respectively. Instability of the lattice dislocations to low temperature irradiation can be attributed to creation of interstitials and vacancies that increase the disorder in the matrix and increase the atomic diffusion rates. This can explain the instability of the dislocation cells in the powder-bed sample. Dislocation cells are low energy structures [53], however, cells observed in this study lacked stability under proton irradiation at low temperature. Annealing of AM Ale12Si has shown instability of the cells [54], similar to the effect of irradiation. Proton irradiated AM 316L stainless steel to 2.5 dpa at 360  C has also shown instability of the dislocation cell structures and Song et al. provides a diffusion based explanation of the instability [25]. Using the same approach, with a damage rate of 2:18  10 6 dpa=s and cell size of 400 nm, irradiation diffusion of 2:9  10 20 m s is calculated, equivalent to a thermal diffusion rate at around 717  C. Similar to the conclusion of Song et al., 717  C is higher than the recovery temperature, suggesting irradiation-

induced diffusion could explain the instability of the dislocation cells. All samples appear to approach the same final grain interior microstructure. It is expected that using additive manufacturing processing conditions to optimize microstructural stability in austenitic stainless steels for certain irradiation environments is likely less effective than has been shown for optimizing mechanical properties with additive manufacturing processing conditions [20]. An exception is the effect of processing on grain size and morphology. Grain size and morphology are more stable to irradiations than lattice dislocations and will impact the bulk mechanical properties of the material. It is expected that even after the grain interior microstructures reach an equilibrium state, differences in bulk mechanical response will be observed due to grain size and morphology. In the orientation maps of Fig. 2, grain size distribution is much smaller for the powder-bed compared to the DED produced samples, which should result in an increased yield strength for the powder-bed material from the Hall-Petch relationship, assuming similar grain interior microstructures. Neutron irradiation would be required to confirm differences in bulk mechanical behavior. 4.2. Hardness before and after irradiation Nanoindentation with a minimum of 10 and 22 indents for wrought and AM samples, respectively, was conducted for each sample in the irradiated and unirradiated conditions. The nanoindentation values were averaged to obtain the values in Table 4. Due to the values of several indents being averaged over a large area, sensitivity is decreased to both crystallographic orientation and local variations in microstructure including dislocation density. As a result, the average values do not correlate exactly to the distinct microstructures presented in the TEM micrographs. Additionally, grain size effects on strength are not apparent in the nanoindentation results. The results specifically gauge the relative ease of plasticity within a localized volume. The smallest grains in the powder-bed sample were several microns in diameter, resulting in each nanoindentation likely localized to a single grain. Therefore a comparison of average nanoindentation hardness values between the powder-bed, DED, and wrought produced material should correlate with differences of the in-grain structures (e.g., dislocation density, precipitates, etc.). Smith et al. investigated the individual contributions of microstructural features to yield strength for AM DED 304L and found dislocation density and compositional segregation to be significant, with precipitates contributing less due to their size and number density [55]. Hardness values of the unirradiated samples varied depending on processing method, with the wrought sample exhibiting the highest value of hardness followed by the low S powder-bed, the low S DED, and high S DED samples. The wrought sample contains a significant amount of cold-work evident from the EBSD and TEM in Figs. 2(d) and 3, respectively, and this results in a high initial hardness. The hardness trends in the unirradiated material are predicted to be a result of differences in dislocation density. The XRD FWHM results support this argument with the wrought having the highest FWHM and the AM samples having lower FWHM. Larger FWHM indicates higher dislocation density. The agreement is not perfect for the ranking of the AM samples with the DED low S having slightly higher FWHM values than similarly FWHM valued DED high S and powder-bed low S samples. Initial hardness shows the powder-bed low S with a slightly higher hardness than the DED high S and low S. After irradiation the different samples, significantly, exhibited similar hardness values. This can be justified by the microstructures in the grain interiors evolving to similar states. Prior studies on austenitic stainless steels support a transition during low temperature

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irradiation to similar microstructures and mechanical responses even at low doses [50,51]. 5. Conclusions  Using 304L powder with varying levels of sulfur and produced by different additive manufacturing methods resulted in significant microstructural differences in builds. These microstructural differences highlight the importance of precursor material characterization on qualifying additively manufactured materials for engineering applications.  The microstructures responded to proton irradiation to 0.6 dpa below 150  C by a decrease in lattice dislocation content and introduction of small, several nm diameter, irradiation-induced dislocation loops. The results are consistent with previous studies of conventionally (i.e. non-AM) prepared austenitic stainless steels.  All samples, including wrought, after irradiation had average nanohardness values within error of each other despite varying initial nanohardness values. Consistent with austenitic stainless steels approaching equilibrium mechanical properties, despite different initial lattice dislocation content, when irradiated at low temperatures.

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 B.P. Eftink: Conceptualization, Methodology, Formal analysis, Investigation, Writing - original draft, Visualization. J.S. Weaver: Methodology, Formal analysis, Investigation, Writing - original draft. J.A. Valdez: Investigation. V. Livescu: Investigation. D. Chen: Investigation. Y. Wang: Supervision. C. Knapp: Resources. N.A. Mara: Supervision. S.A. Maloy: Project administration. G.T. Gray: Project administration, Funding acquisition. Acknowledgements This work was performed at Los Alamos National Laboratory. Los Alamos National Laboratory, an affirmative action equal opportunity employer, is operated by Triad National Security, LLC, for the National Nuclear Security Administration of the U.S. Department of Energy. Work was supported by the LANL Dynamic Materials Properties Program, United States. This work was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science. Dr. Saryu Fensin is acknowledged for fruitful discussions. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2020.152007. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.

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