Production of corrosion-resistant 316L stainless steel clads on carbon steel using powder bed fusion-selective laser melting

Journal of Materials Processing Tech. 273 (2019) 116243

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Production of corrosion-resistant 316L stainless steel clads on carbon steel using powder bed fusion-selective laser melting Pratik Murkutea, Somayeh Pasebanib,c, O. Burkan Isgord,

T

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a

Materials Science, School of Mechanical, Industrial and Manufacturing Engineering, Corvallis, OR, USA Manufacturing Engineering, School of Mechanical, Industrial and Manufacturing Engineering, Corvallis, OR, USA c Advanced Technology and Manufacturing Institute (ATAMI), Corvallis, OR, USA d School of Civil and Construction Engineering, Oregon State University, OR, USA b

A R T I C LE I N FO

A B S T R A C T

Associate Editor: A. Clare

Powder bed fusion-selective laser melting (PBF-SLM) was used to produce corrosion-resistant 316L stainless steel clads on 1018 carbon steel substrates. The PBF-SLM parameters such as laser power, laser-scanning speed, hatch spacing, and layer thickness were optimized to attain maximum clad density and a superior metallurgical bond to the substrate. Typical energy densities that are required for producing dense 3D parts (˜100 J/mm3) was inadequate for cladding operations, and a higher energy density (333–1333 J/mm3) was necessary to produce low-defect clads with good adherence to the substrate. A maximum clad thickness of 133.14 μm was achieved at the lowest tested scan speed of 100 mm/s after ten layers of powder melting. The clads had lower chromium content than the 316L powder due to evaporative losses experienced during laser melting process. However, chromium contents in the range of 13–15% were successfully achieved in all cladded specimens. Increasing laser scan speeds had a negative impact on the nanoindentation hardness of the clads; however, the clad hardness at all scan speeds was found to be higher than AISI 316L SS. Electrochemical tests showed that the corrosion properties of clads produced at low laser scan speeds were comparable to AISI 316L SS.

Keywords: Additive manufacturing Powder bed fusion selective laser melting 316L stainless steel 1018 low carbon steel Metal cladding Corrosion

1. Introduction Low carbon steel (LCS) has extensive applications in various industries (e.g., automotive, construction, and transportation) due to its high strength-to-cost ratio. However, its low corrosion resistance in neutral, acidic, or saline environments limits its utilization and service life. Corrosion resistant alloys, such as stainless steels (SS), provide superior corrosion resistance than LCS in such applications, albeit with a significantly higher material cost. Various sectors would benefit from a low-cost corrosion-resistant material. The cladding is one of the viable methods to manufacture a composite with wear and corrosion resistant surface layer on a low-cost substrate with desired mechanical properties, as in the case of 316L SS clad on LCS substrate. Conventional manufacturing techniques have been used to join/ fuse two dissimilar metals in the past; for example, Ishida (1991) conducted a metallurgical investigation of stainless steel clads formed on mild steel by three different welding techniques, namely, tungsten inert gas (TIG) welding, high current pulsed arc welding, and constricted arc welding. Deqing et al. (2007) focused their efforts on improving the bonding characteristics between 304 SS and Q235A steel

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using interlayer diffusion bonding method, where they sandwiched an Al-Cu-Mg interlayer between 304 SS and Q235A steel and heated the three-layered composite between 600° and 640 °C for various durations. Khara et al. (2016) developed a corrosion resistant chromium powder coating on mild steel by powder roll bonding followed by high-temperature annealing to promote Cr diffusion in steel substrate. Yoshimura (1987) made a unique attempt of hot rolling stainless steel tubes filled with carbon steel shavings, resulting in a composite with a core of carbon steel cladded with a metallurgically bonded stainless steel layer. Jing et al. (2014) explored a novel reduction bonding technique to join plain carbon steel plate (substrate) and 316L SS plate (clad). This method involved heating the clad and substrate plates to 1223–1323 K in a pure hydrogen atmosphere to reduce the oxide scales on the carbon steel, followed by conventional hot rolling. The key limitation with most clads produced with these conventional techniques is a distinct and abrupt metallurgical transition between SS and LCS at the clad-substrate interface. This sharp metallurgical change results in an inferior bonding between the clad and the substrate, often leading to cracking or delamination failures when subjected to manufacturing (e.g., rolling) or operational stresses,

Corresponding author. E-mail address: [email protected] (O.B. Isgor).

https://doi.org/10.1016/j.jmatprotec.2019.05.024 Received 31 October 2018; Received in revised form 29 March 2019; Accepted 20 May 2019 Available online 24 May 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.

Journal of Materials Processing Tech. 273 (2019) 116243

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making the composite unfit for field use. Furthermore, these conventional techniques often result in clads with defects such as crevices, pinholes, keyholes, and porosity, leading to reduced service life. In the past two decades, additive manufacturing (AM) techniques have been gaining impetus owing to their inherent capabilities and advantages, which include producing complex geometries, savings in material cost, reducing labor cost, lead time and controlling final product quality. In the past, several researchers have employed AM techniques for laser cladding. For example, Rashid et al. (2017) used direct metal deposition (DMD) technique to deposit up to 1 mm thick 316L SS clad on a mild steel substrate. Farren et al. (2007) employed laser engineered net shaping (LENS) technique to fabricate a transition joint between 1085 carbon steel and 316L SS. Majumdar et al. (2005) used a CO2 laser assisted, direct laser deposition (DLD) technique for 316L SS fabrication. The inherent advantage of these AM techniques lies in the powder dispensing method. In LENS, DMD and DLD techniques, the substrate is melted with a traveling heating source, and the clad material in powder form is blown on the melt pool through co-axial nozzles, allowing the cladding of very complex geometries but with low resolution and low dimensional accuracy. The powder bed fusion selective laser melting (PBF-SLM) technique differs from the abovementioned AM methods in that a thin metal powder layer that is spread uniformly on a flat substrate is locally melted using lasers. After melting of one layer, another thin layer of metal powder is spread on top of the previously fused metal, thereby producing three-dimensional (3D) geometry with repeated layers. The PBF-SLM technique has been conventionally used for printing complex 3D parts, and they have been extensively studied for properties. Bartolomeu et al. (2017) did a detailed microstructural, mechanical and wear behavior investigation of the 316L SS parts produced via three different techniques: SLM, hot pressing, and conventional casting. Chao et al. (2017) studied the corrosion properties of selective laser melted austenitic 316L SS and observed an increase in corrosion resistance as compared to the wrought alloy. Liverani et al. (2017) studied the effects of SLM process parameters on the mechanical properties of 316L SS parts, while Suryawanshi et al. (2017) focused on the effects on mechanical properties such as fatigue cracking due to various scan strategies. The SLM and laser cladding deposition processing techniques have been compared in regards to their effect on the metallurgical properties on 316L SS parts by Ma et al. (2017). Cherry et al. (2014) fine-tuned the SLM derived parameters like volumetric energy density (J/mm3) to achieve the highest part density. However, in producing 3D parts, the first few layers of fused metal in contact with the substrate are a part of the support structure. The purpose of the support structure is to support overhanging surfaces, maintain the structural integrity, and reduce distortion in 3D parts during the print, as noted by Samant et al. (2018). The support structure is machined off from the substrate and discarded after the print, and hence the bonding between the substrate and the deposited fused metal is often overlooked and not taken into consideration. On the contrary, the governing criteria for a successful cladding operation is a superior clad-substrate bond. Therefore, 3D printing and cladding operations are fundamentally different in the sense of bonding of materials. Although, the PBF-SLM technique has been extensively used to produce complex 3D parts, the potential for cladding or bonding two dissimilar metals has not been explored before. In this study, we investigate the feasibility of using the PBF-SLM technique for cladding 316L SS on and LCS substrate with the ultimate goal to increase the corrosion resistance of the substrate material while reducing the cost of producing corrosion resistant alloys. As shown in the previous studies that are focused on printing 3D components, the laser energy density has a significant effect on the properties of the final product even though many other factors (e.g., laser spot size, hatch orientation, hatch space, and powder layer thickness) affect the PBF-SLM process. Sufiiarov et al. (2017) observed that increasing layer thickness had an adverse effect on the strength of

the component. Cherry et al. (2014) showed a proportional relationship between the laser power, energy density, and part density. Energy density is mainly determined by four parameters following:

E=

P vs * h * d

(1) 3

where E is the laser energy density (J/mm ), P is the laser power (W), vs is the scan speed (mm/s), h is the hatch distance (mm), and d is the powder layer thickness (mm) spread on the substrate. Cherry et al. (2014) observed that the optimum energy density needed for printing 3D components with minimum porosity and maximum density using the 316L SS powder was 104.52 J/mm3. However, the applicability of this criterion for producing corrosion resistance cladded surfaces on LCS is not known. In this study, we focus on the effect of the parameters that affect the laser energy density on the properties of the clad surfaces. We characterize various properties of the cladded systems using metallurgical, mechanical and corrosion tests. 2. Experimental procedure 2.1. Materials In the present study, gas atomized 316L SS powder with a particle size in the range of 15–45 μm (D50 = 32 μm), and an apparent density of 4.49 g/cm3 was used as a clad power feedstock. The secondary electron (SE) micrograph of the powder (Fig. 1(a)) shows a typical spherical morphology for gas-atomized powders. The 1018 LCS plate with dimensions 25.4 × 25.4 × 3.175 mm3 was used as a substrate material. The microstructure of the substrate comprising of α-ferrite (bright phase) and pearlite (α + Fe3C) − (dark phase) fine grain structure is presented in Fig. 1(b). Furthermore, wrought and annealed AISI 316L SS was also used in this study to benchmark and compare the mechanical and corrosion properties of the SS clad developed using PBF-SLM. The microstructure of the AISI 316L SS mainly consisted of annealed austenite grains as shown in Fig. 1(c). The nanoindentation hardness values for the LCS substrate and AISI 316L SS plate were 2.12 ± 0.09 GPa and 3.01 ± 0.25 GPa, respectively. The chemical compositions of 316L SS powder, LCS, and AISI 316L SS are presented in Table 1. The AM cladding process was divided into two sets based on input energy density values as shown in Table 2: low energy density (LED) clads and high energy density (HED) clads. In the first set (LED clads), only two parameters, power (P) and laser scan speeds (vs) were varied to achieve the energy densities that are close to 104.52 J/mm3, as recommended by Cherry et al. (2014). The resulting energy density tested was in the range of 83–250 J/mm3. Based on the outcomes of the LED clad set, further optimization of the PBF-SLM process was investigated by producing HED clads using higher energy density in the range of 333–1333 J/mm3 by keeping the P = 200 W and setting lower vs in a range between 100 and 400 mm/s. To investigate the effect of hatch overlap on HED clad properties an additional clad was produced at 100 mm/s with a hatch overlap of 50%. Hatch spacing is the distance between the centers of two melt tracks laid side by side. For LED and HED clads, other parameters such as hatch spacing (h), hatch orientation, layer thickness (d), and total thickness (T) were fixed, and their values are presented in Table 2. The PBF-SLM chamber schematic is illustrated in Fig. 2. 2.2. Characterization studies 2.2.1. Scanning electron microscopy The backscattered emission (BSE) imaging and energy dispersive Xray spectroscopy (EDS) studies were done using Thermo Scientific FEI quanta 600 FEG scanning electron microscope (SEM). For SEM analysis, the clad specimens were cut, and cold mounted in epoxy to expose the transverse face of the specimens to reveal the clad-substrate interface. 2

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sample preparation process was followed for AISI 316L SS and LCS specimens. The chromium contrast observed in BSE imaging mode was used to investigate the clad-substrate interface for clad thickness measurements. The EDS line and area scans were performed across the cladsubstrate interface for elemental analysis. 2.2.2. X-ray diffraction The Bruker-AXS D8 Discover, X-Ray diffractometer was used for XRD scans of the 316L SS powder, LCS, and the as-produced clad surfaces. The scans were performed with Cu-Kα radiations (λ = 1.54 Å) xray source, operating at 40 kV and 40 mA. The 2θ range for the x-ray scan was set 20°–110° with a step size of 0.05° and the scan speed of 4°/ min was set for all the x-ray scans. 2.2.3. Optical microscope For microstructural analysis, LCS was etched with 2 vol. % Nital (Ethanol + HNO3), whereas AISI 316L SS was etched with aquaregia (HCL:HNO3 = 3:1) for 10 s. However, to reveal the clad layer microstructures, the specimens had to be electro-etched in 10% oxalic acid at 15 V for 1 min. Surface preparation procedure for optical microscope studies was similar to those for SEM specimens. The microstructural analysis was done using a Leica DM2500 optical microscope. 2.2.4. Nano-indentation The mechanical properties of the clad, HAZ and the substrate were characterized using the nanoindentation technique. The tests were performed on Micromaterials-NanoTest Vantage setup, and the indentations were made using a Berkovich diamond tip indenter. A series of 15 indents were made across the clad-substrate interface, and the hardness in all three regions: Clad, HAZ, and the Substrate, was measured in triplicates as shown in Fig. 3(a). In addition, AISI 316L SS and LCS were separately tested for benchmarking the hardness values and the representative load-depth curves for AISI 316L SS and LCS (Fig. 3(b)). Only 40% of the post indentation data was curve fitted for hardness analysis as shown in Fig. 3(b) (shown by bold curve). The spacing between two successive indents was kept at 20 μm. All the tests were performed at loading and unloading rates of 2 mN/s, with a maximum load of 150 mN for clads and 250 mN for LCS and AISI 316L SS for a dwell period of 5 s at maximum load. 2.2.5. Corrosion testing Open circuit potential monitoring (OCP), electrochemical impedance spectroscopy (EIS), linear resistance polarization (LPR), and cyclic polarization (CP) tests were used to investigate the corrosion behavior of the clads, LCS, and AISI 316L SS. The corrosion tests were performed on as-printed specimens, and no additional surface preparations or treatments were done on the clad surface. Therefore, the inherent surface roughness and defects in the SS clad layer resulting from PBF-SLM process were accounted for in corrosion results. However, the LCS and AISI 316L SS specimens were polished up to 2000 grit silicon carbide paper and followed by cloth polishing up to 0.05 μm aqueous alumina solution. The test specimens were ultrasonically cleaned in isopropanol bath before starting the tests. A three-electrode electrochemical cell was used for the corrosion tests. A near neutral 3.56 wt.% NaCl aqueous solution was used as a test electrolyte. The cladded specimens were used as working electrodes; the graphite block and saturated calomel electrode (SCE) formed the counter electrode and the reference electrode, respectively. A circular surface area of 2.85 cm2 was exposed to the electrolyte during the tests. The corrosion tests were performed using a Gamry 3000 potentiostat/ galvanostat/ZRA. The corrosion testing sequence was set up on the Gamry Instruments Framework software. The sequence comprised of 1 h of OCP monitoring followed by EIS scan in the frequency range of 50,000-0.01 Hz at 5 mV r.m.s AC potential. After EIS, LPR test in the potential range of ± 15 mV vs. OCP were run, and lastly, CP scan with an initial and final potential

Fig. 1. (a) BSE mode image of gas atomized 316L SS powder showing spherical morphology (b) LCS microstructure showing fine grains of α- ferrite (bright phase) and pearlite (dark phase), and (c) AISI 316L stainless steel annealed austenite grain microstructure.

Table 1 Nominal chemical composition of 316L SS powder, LCS and AISI 316L SS (wt. %). Material

C

Cr

Ni

Mo

Mn

Si

P/S

Fe

316L SS - powder LCS AISI 316L SS

0.01 0.20 0.03

17.50 0.7 16-18

12.57 0.9-1.0 10-14

2.36 0.01 2-3

0.4 0.8 2.0

0.40 0.70 0.75

< 0.01 < 0.04 0.045

Balance

After cold mounting, the specimens were polished to 2000 grit silicon carbide paper, followed by cloth polishing up to 0.05 μm sized alumina suspension and finally ultrasonic cleaning in isopropyl alcohol bath. Before the SEM analysis, the samples were sputtered with Au-Pd coating for better electron grounding and improved imaging. A similar 3

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Table 2 Process parameters used for LED and HED clads*.

LED clads HED clads

Power (W)

Scan speed (mm/s)

Energy density (J/mm3)

150, 175, 200, 225 200

600, 800, 1000, 1200 100 (0%, 50% overlap), 200, 300, 400

83-250 333-1333

* Fixed parameters: h = 30 μm, d = 50 μm, hatch angle = 45°, T = 500 μm.

of −0.5 V vs. OCP and apex potential of +1 V vs. OCP was run. The apex current of 20 mA/cm2 and a potential scan rate of 0.166 mV/s, as specified by ASTM (2018), was set for all the CP tests. The EIS, LPR and CP data presented in this paper were area normalized. The equivalent circuit modeling for EIS data was performed in Gamry Echem Analyst software (V 6.04).

(1981). In-depth analysis and detailed reasoning for the occurrence of the balling phenomenon and its adverse effects of the final part properties are reported by Bidare et al. (2018) and Cacace et al. (2017); therefore, they are not discussed here. In addition to balling, geometric warping of the clad layer was observed in all cases. Fig. 4(b) presents the side view of the clads shown in Fig. 4(a) to demonstrate the extent of geometric warping in each case. It can be seen that the extent of geometric warping decreases with increasing power levels, i.e., most significant clad warpings were visible at 150 W, while the least warping was at 225 W. Moreover, clad delamination was observed in all cases, regardless of power and laser scan speed used. The extent of balling, geometric warping and clad delamination were further aggravated at higher scan speeds of 1000, and 1200 mm/s (not shown here) to the extent of total clad peel-off from the substrate. Warping and clad delamination are attributed to high residual and thermal stresses. The residual stresses arise because of incomplete melting of metal particles, at higher speeds or inadequate energy densities. Incomplete or partial melting of the powder resulted in the differential expansion of the melted and unmelted powder, thereby creating a mismatch at the melt-solid interface. This differential expansion resulted in the generation of residual stresses, as observed by Simson et al. (2017). Moreover, rapid cooling rates associated with radiative cooling in an inert nitrogen atmosphere, result in high thermal stresses in the clad layer. Residual stresses in conjunction with thermal stresses resulted in geometrical warping. Furthermore, due to the difference in coefficient of thermal expansion values of clad material (16.0 × 10−6 K−1) and substrate (11.7 × 10−6 K−1), the expansion mismatch on melting resulted in the poor bonding of clad materials on the

3. Results and discussion 3.1. Visual observations This section presents the visual observations made on the LED, and HED clads. Interpretation and reasoning of the defects observed in the clad produced are provided in the following sections. 3.1.1. LED clads The LED clads produced by PBF-SLM are presented in Fig. 4. As shown in Fig. 4(a), the clads were produced at vs. = 800 and 600 mm/s and with laser powers in the range of 150 W to 225 W. The balling phenomenon was observed at all scan speeds and power levels. However, lower power levels showed extensive balling, resulting in increased surface roughness. The increase in balling with decreasing power levels and increasing scans speeds was also observed in two separate research works by Li et al. (2011) and Sufiiarov et al. (2017), who reported similar trends. The occurrence of balling could be due to several reasons such as inadequate laser energy input resulting in little liquid content in melt pools, melt pool splashing under high scan speeds, capillary instability, and Plateau-Rayleigh instability. The details on the Plateau-Rayleigh phenomenon are given by Chandrasekhar

Fig. 2. PBF-SLM print chamber schematic. 4

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Fig. 3. Nano-indentation test schedule showing (a) the indents made across the clad-substrate interface (b) loading/unloading nano-indentation plots for AISI 316L SS and LCS and region selected for curve fitting.

Fig. 4. LED clads: (a) produced at increasing power (top to bottom) and increasing speeds (right to left) (b) geometrical warping due to excessive thermal and residual stresses.

substrate. An additional observation was made for the clads produced at 600 mm/s (Fig. 4(a)); the clads were not entirely formed; the dashed white lines show the demarcation between the fully formed and incomplete clad layer. This is attributed to improper powder spreading resulting due to recoater blade damage. Recoater blade or roller damage due to interaction with metal balling, sharp edges and warping of the previously fused metal layer is a well-known phenomenon demonstrated by Li et al. (2018). The blade damage was caused by out of plane sharp edges and corners of the clad to the left side (800 mm/s), obstructing the movement of the coater blade that led to inefficient and hindered powder spreading. Subsequently, the complete cladding was done only for the first few layers, when the residual stresses were not enough to cause delamination and clad warping.

were melted, which formed the clad layer. On the contrary, in the conventional PBF-SLM printing, first 10 layers constitute the base of the support structure, where the bond strength is not the priority. Moreover, the contact area between the support structure and the baseplate is less than the clad areas in this study. Consequently, HED clads were produced with process parameters presented in Table 2. Fig. 5(a) shows the clads produced at 100 mm/s (top), and 200 mm/s (bottom) have a consistent and relatively smooth surface finish with minimal balling. The similar surface finish was observed in clads produced at 300 and 400 mm/s (not shown here). It was observed that the surface roughness increased with increasing scan speeds as previously observed in a LED clads. However, the balling phenomenon was alleviated, and the surface finish was considerably better in HED clads. For HED clads, the laser power of 200 W and low scan speeds (100–400 mm/s) resulted in higher melt pool temperatures due to more substantial dwell time. The previous work Li et al. (2011) supports the observations made in this paper. They showed that the higher melt pool temperatures result in the better wetting ability of the metal which in turn, alleviate the balling phenomenon and thereby promotes better bonding between the substrate and the clad layer. Therefore, HED clads show a superior clad quality as compared to LED clads. The investigation of the effect of hatch overlap on clad properties revealed that 50% overlap (100 mm/s) resulted in the better surface

3.1.2. HED clads After analyzing the LED clads, it was realized that the energy density input, reported by Cherry et al. (2014), would be inadequate to clad SS on LCS, and slower laser scans speeds would be necessary. In a PBFSLM process for manufacturing of 3D SS components, the input energy density value is dictated by the bonding of the new SS layer with the previously fused SS layer, which only accounts for cohesive forces. However, in this study, since two dissimilar metals, 316L SS and LCS, are bonded, adhesive forces are in play, and the physics and thermodynamics are different from bonding stainless steel with itself. Furthermore, the total cladding thickness was set at T = 500 μm and the layer thickness was d = 50 μm. Therefore, only 10 layers of powder 5

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Fig. 5. (a) HED clads produced at 100, and 200 mm/s (b) Chromium contrast observed in BSE imaging mode used for clad thickness measurements (left) and chromium EDS- area profile (right) (c) clad thickness variation with scan speeds.

finish, i.e., reduced surface roughness than 0% overlap (100 mm/s). Since HED clads had a better surface finish with minimal balling, all the geometrical, metallurgical, mechanical and electrochemical characterizations were performed only on HED clads and the results and presented hereafter.

Bidare et al. (2018) and Khairallah et al. (2016) increasingly higher amounts of powder is blown away from the melt pool with increasing scan speeds due to the more significant interaction between laserplasma plume and the powder bed causing more instability. Furthermore, on account of evaporation of metal vapor from the melt pool and net upward flux during laser melting, inward flow of the ambient inert gas towards the melt track is created due to the Bernoulli effect and similar observations were made by Gunenthiram et al. (2018). This inward flow of the inert ambient gas is sufficient to entrain powder particles from adjoining areas, which can become incorporated into the melt pool or ejected with the metal vapor. The Bernoulli effect is intensified at high scan speeds thereby leading to the formation of wider denudation zones. Where denudation zone is the depletion of metal powder particles in the zone immediately adjoining the solidified track. Consequently, the formation of wide denudation zones due to the Bernoulli effect results in less powder being available for melting to form the clad layer at higher scan speeds as shown in two different works by Khairallah and Anderson (2014) and Matthews et al. (2016). This resulted in melting of powder layers thinner than initially spread layer thickness of 50 μm. As shown in Fig. 5(c), the decrease in clad thickness with increasing scan speed, is in agreement with the theory of denudation zones and the Bernoulli effect. In this range of scan speeds, the clad thickness data follows a linear trend, although, more data points are needed for establishing a statistical relationship between scan speed and clad thickness.

3.2. Geometrical characterization The clad thickness measurements were done based on the chromium contrast observed in SEM-BSE imaging mode, and the measurements were confirmed by optical microscopy. Fig. 5(b)-left side, shows a representative BSE image of the clad-substrate interface showing chromium contrast in the clad region. This contrast is due to the difference of the elemental atomic weight in Cr-rich stainless steel clad layer, and Cr-deficient LCS substrate as shown in EDS area maps for Chromium of the same specimen (Fig. 5(b)-right side). The thickness measurements were done at multiple locations along the clad length, and the results were plotted as a function of vs. as shown in Fig. 5(c). It was observed that with increasing scan speeds the clad thickness decreased. The maximum average clad thickness of 133.14 μm was achieved at the lowest scan speed of 100 mm/s and 50% overlap, whereas the lowest average clad thickness of 76 μm was observed at 400 mm/s. This trend could be explained on the basis of input energy density (Eq. (1)), powder bed-laser interaction and the denudation zone theory during powder melting. As observed previously in two independent research studies by 6

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Fig. 6. (a) optical micrograph showing the heat affect zone (HAZ) underlying the clad layer (b) HAZ thickness variation with scan speeds (c) variation of HAZ thickness (white arrow) and microstructural evolution in HAZ with increasing scan speeds.

differential cooling rates experiences at various depths. Higher temperatures and rapid cooling rates at shallower depths of the HAZ resulted in a decarburized layer of fine δ-ferrite grains. Whereas relatively low temperatures and slower cooling at higher depths facilitated carbon diffusion from δ-ferrite grains to adjoining γ-ferrite (austenitic) areas, which transformed to pearlite grains on cooling below austenizing temperatures. Furthermore, the resulting HAZ microstructure, which shows formation of discrete pearlite islands in predominantly ferrite matrix is dictated by the formation second phase (cementite) formation and its wetting/de-wetting ability. In two separate works by Straumal et al. (2012) and Straumal et al. (2014), it was demonstrated that the ability of the second solid phase (γ) to completely or partially wet the grain boundaries of the primary phase (α) is dependent on the energy of α/α grain boundary in the α phase energy (σαα) and interfacial energy (2σαγ) two interfaces between the α and γ phases. If the σαα < 2σαγ, the α/α grain boundary can exist in equilibrium contact with the γ phase, however, if the 2σαγ < σαα, the α/α grain boundary should be replaced by the layer of the second solid phase γ. These wetting and de-wetting regimes are highly dependent on the system temperature and the associated cooling rates. Typically, the PBF-SLM process is associated with very high melt pool temperatures of 5000–7000 K and rapid cooling rates of 103–105 K/s, therefore, the dynamics of second phase grain boundary wetting phenomenon observed during the SLM melting process are significantly different that the one observed by Straumal et al. (2012) at equilibrium heating and cooling processes such as annealing or normalizing. However, the underlying principles of

3.3. Metallurgical characterization 3.3.1. Microstructural observations In separate studies by Bidare et al. (2018) and Huang et al. (2016b) showed via FEM modeling that in PBF-SLM process, the melt pool temperatures could reach as high as 7000 K; such high temperatures coupled with high cooling rates in the order of 103–105 K/s in nitrogen atmosphere result in the formation of heat affected zone (HAZ) in substrate region. HAZ is the unmelted region of the substrate, which undergoes microstructural changes influenced by high temperatures and rapid cooling rates. Fig. 6(a) shows the HAZ underlying the clad layer revealed by chemical etching by 2% nital solution. The HAZ thickness measurements were made using optical microscopy. As shown in Fig. 6(b), the HAZ thickness increased with decreasing scan speeds. Since P, h and d were fixed, increasing scan speeds from 100 mm/s to 400 mm/s, resulted in a 4-fold decrease in energy density value from 1333.33 J/mm3 to 333.33 J/mm3 respectively. Consequently, at high energy densities, the depths to which the temperature increased above the austenitizing temperature (A3 temperature line in iron-iron carbide diagram) was deeper. As a result, for slow scan speeds, the microstructural changes were observed at higher depths, and vice versa. On further microscopic investigation of the HAZ microstructures, it was observed that areas immediately adjacent to clad layer showed very fine decarburized δ-ferrite microstructure and at higher depths, coarse pearlite grains in the form of carbon clustering was observed. Assuming, HAZ temperatures reached well above austenizing temperatures, these microstructural observations could be attributed to 7

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interfacial energy and contact angles remains unchanged and the results obtained in the present study has close semblance to the previous studies. In the HAZ, the areas close to the clad-substrate interface showed δferrite grains completely separated by cementite, in other words, Fe3C phase completely wets the δ-ferrite grain boundaries, resulting in contact angles close to 0°. This could be attributed to higher temperatures and lower cooling rates associated with shallow HAZ depths, allowing lowering of interfacial energy (2σδ/Fe3C) than the δ-ferrite grain boundary energy (σδ/δ). However, at higher HAZ depths, discrete pearlite islands were observed, which implies a de-wetting regime, which is due to relatively lower temperatures and rapid cooling rates. The primary solidification mechanism coupled with wetting/de-wetting of δ-ferrite grain boundaries by second solid phase Fe3C, results in HAZ microstructure as shown in Fig. 6(c). Furthermore, since the heat penetration depths decrease with increasing laser scan speed, the completely wetted and partially wetted areas decrease in size as shown in Fig. 6(b). Fig. 6(c) presents the HAZ for all scan speeds, and it was observed that the coarse pearlite grains (carbon clusters) are formed closer to the clad layer (at shallower depths) with increasing laser scan speeds. This is due to increasingly shallower depths of heat penetration at higher laser scan speeds. For the microstructural investigation, the cladded specimens were electroetched according to the procedure mentioned in the earlier section. Fig. 7 presents a characteristic microstructure of the 316L SS clads deposited via the PBF-SLM process. It is evident from the figure that the PBF-SLM process resulted in columnar grains in the stainless steel clad layer. The columnar grain structure is typically associated with rapid unidirectional cooling. In this case, unidirectional cooling was an outcome of heat dissipation from the clad layer to the LCS substrate, which acted as a heat sink. It was observed that the columnar grain size was nearly the same with changing scan speeds, albeit, the grain size observed in clads was smaller than that in AISI 316L SS (Fig. 1(c)). Owing to rapid cooling rates of the melt pools, grain coarsening or grain growth could not be facilitated. The oblique black line markings (shown by white arrows) visible on the clad layer (Fig. 7) show the melt track morphologies. For all scan speed, the top portion of the clad exhibits widely spaced melt tracks, whereas the clad region near clad-substrate interface showed narrow spacing between melt tracks and shallower melt pools. This is due to changing scan orientation by 45° after each layer melting, thereby changing effective cross-section of the melt pool exposed to the face being electroetched and remelting of the melt tracks in subsequent passes. It is interesting to note that the columnar grain growth is not limited to the melt pool boundary, but all the grains grow across several

Fig. 8. EDS line profile for chromium content across the clad-substrate interface at different scan speeds.

melt pools along the thickness of the clad. 3.3.2. Elemental mapping The EDS line profile scans for chromium element across clad-substrate interface are shown in Fig. 8. It was observed that the chromium content in all clad layers was less than initially present in 316L SS powder (Table 1). Bidare et al. (2018) observed in their FEM simulations that, based on the laser power and scan speeds, the localized melt pool temperatures could reach as high as 3500–7000 K. Since these local temperatures are well above the vaporization temperatures of iron (3134 K) and chromium (2944 K), the Fe and Cr metal losses were experienced due to evaporation in study done by Huang et al. (2016a). The Cr losses were more dominant at a low scan speed of 100 mm/s due to higher melt pool temperatures and more considerable laser-dwell time. The dotted lines denote the clad-substrate interface marked by a sharp decrease in chromium content. The clad thickness measurements (Fig. 5(c)) and receding clad-substrate interface as shown in Fig. 8 are in good agreement. 3.3.3. X-ray diffraction The clad surfaces, 316L SS powder, and LCS were subjected to X-ray diffraction for phase analysis, and the XRD spectra are shown in Fig. 9. As illustrated in Fig. 9(a), LCS exhibited a typical α-ferrite (ferrite) and Fe3C (cementite) peaks, whereas 316L SS powder showed characteristic austenite peaks. However, clads produced at slowest scan speeds of 100 mm/s (0% and 50% overlap) showed a good mix of austenite, αferrite and cementite peaks from pearlite. On the contrary, clads produced at higher scan speeds (Fig. 9(b)) showed only austenitic peaks which is a characteristic of pure stainless steel powder. Cullity (1978) and Weymouth and John (2005) in their texts on principles of X-ray diffraction have demonstrated that the penetration depth of X-rays in stainless steels is of few tens of microns. Considering the clad thicknesses produced via SLM, the possibility of an X-ray signal

Fig. 7. The optical micrographs showing melt pool boundaries (white arrows) and columnar grain at 50× magnification. 8

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Fig. 9. XRD spectra of the HED clads produced at (a) 100 mm/s at overlap 0%, 50% and (b) 200–400 mm/s compared with LCS and AISI 316L SS.

Fig. 10. Nano-indentation test results showing (a) hardness plots for 100 mm/s at 0% and 50% (b) hardness plots of scan speeds of 200, 300 and 400 mm/s.

coming from the substrate layer can be ruled out. Furthermore, the thinner clads at high scan speeds (200–400 mm/s) showed only austenitic peaks, which is suggestive of the output signal being received only from the clad layer. Therefore, it could be said with certainty that, high energy density input resulted in the considerable mixing of substrate and clad materials which would imply better bonding between the 316L SS clad layer and LCS substrate.

microstructure adjoining the coarse columnar grained SS clad. The high clad hardness and even higher HAZ hardness imply an in-service superior wear resistance than AISI 316L SS, thereby extending the service life of the cladded components. Additionally, it was observed that the increasing scan speed adversely affected the clad hardness, with 100 mm/s clads showing the maximum hardness of 6.09 GPa (0% overlap) and 5.01 GPa (50% overlap), whereas the 400 mm/s clads had the lowest hardness value of 3.85 GPa. The decreasing hardness with increasing scan speed relates to decreasing melt density and increasing defect, i.e., porosity, keyholes and balling. Inadequate energy density leading to the inefficient melting of powder resulted in non-ideal clad density thereby reducing hardness at higher scan speeds.

3.4. Mechanical characterization Fig. 10 shows the nanoindentation hardness profiles for all the clads (Fig. 10(a) and (b)). As previously mentioned, the average hardness of AISI 316L SS and LCS was 3.01 GPa and 2.12 GPa, is marked by the horizontal solid lines in figures. For all the scan speeds, the SS clad hardness was found to be higher than AISI 316L SS. However the clad hardness decreased with increasing scan speed. Smaller austenitic columnar grain size due to rapid cooling rates associated SLM process resulted in enhanced clad hardness. Additionally, it was observed that the hardness for 50% hatch overlap was lower than 0% overlap, and this could be attributed to the remelting of the melt tracks that resulted in the tempering effect, i.e., stress relaxation and recrystallization. Remelting of the melt tracks led to a reduction of residual stresses, and minimal grain growth thereby lowering the hardness of the clads. For all scan speeds, the HAZ region adjoining clads were found to have a higher hardness that AISI 316L SS, furthermore, for higher scan speeds (200–400 mm/s), the hardness was higher than the clad itself. This high HAZ hardness resulted from a very fine δ-ferrite

3.5. Electrochemical characterization 3.5.1. Open circuit potential (OCP) At the start of corrosion testing, OCP was monitored for one hour. The OCP curves for all tested specimens are plotted in Fig. 11. AISI 316L SS specimens showed high OCP (≲−0.2 V SCE) that increased further with time due to the formation of a corrosion-resistant chromium oxide surface film, which reduced corrosion rates to a passive state. LCS exhibited the most negative potentials, which is the typical characteristics of an actively corroding metal. OCP of all cladded specimens were higher than that of LCS, which indicated that cladded surfaces were more corrosion resistant than LCS. The OCP values of clads that were prepared using 100 mm/s laser scan speed were similar 9

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LCS, and the clads. High-frequency impedance values in the impedance Bode plot (e.g., at 104 Hz) indicate the solution resistance (Rs), which was around 17.5 ohm cm2 for the 3.5 wt. % NaCl electrolyte. The impedance values at low-frequencies (e.g., at 0.01 Hz) reflects the sum of charge transfer resistance (Rct) and solution resistance (Rs). Charge transfer resistance, Rct, is the electrical resistance to the charge movement offered by the electrical double layer formed on the alloy surface. Typically, protective films of corrosion resistant alloys exhibit high charge transfer resistance. The observed Rs + Rct value for the AISI 316L SS (4.68 × 105 ohm .cm2) was three orders of magnitude higher than that of LCS (7.46 × 102 ohm.cm2), confirming the superior corrosion resistance of stainless steel over carbon steel. For the clads produced at the slowest scan speed of 100 mm/s, the impedance values at 0.01 Hz were comparable to AISI 316L SS (1.49 × 105 ohm .cm2 and 1.21 × 105 ohm. cm2 for 50% and 0% hatch overlap, respectively. The clads with 50% hatch overlap shows marginally higher impedance than 0% overlap; this behavior is likely because of partial remelting of the melt track during the overlapping scan, which resulted in additional input energy, higher energy density, higher melt pool temperatures, and volumes. Increased melt pool volume leads to slower cooling rates of the melt tracks, which promotes grain coarsening in the clad layer, relieves stresses, and alleviates surface roughness as noted by Li et al. (2011). Lower laser scan speeds also cause slower cooling rates, which explains the lower impedance values observed at 0.01 Hz of clads produced at higher laser scan speeds, as shown in Fig. 12(a). The clads produced at higher scan speeds show two orders of magnitude lower impedance (at 0.01 Hz) than clads produced at 100 mm/s scan rate. As illustrated in Fig. 12(b), phase angle for AISI 316L SS was highly negative, in the range of −80° to −85° for a wide range of frequencies,

Fig. 11. Open circuit potential curves for AISI 316L SS, LCS and clads produced at 100–400 mm/s scan speeds.

to that of AISI 316L SS. However, at higher laser scan speeds, the OCP values were progressively more negative. This trend suggests that the corrosion resistance of clads decreases with increasing scan speeds. The interpretation for this behavior is provided in subsequent sections while discussing EIS, CP and LPR results.

3.5.2. Electrochemical impedance spectroscopy (EIS) Fig. 12, shows the EIS data presented in the form of impedance Bode plot (a), phase angle Bode plot (b) and Nyquist plot (c) for AISI 316L SS,

Fig. 12. EIS data showing (a) impedance bode plot (b) phase angle bode plot (c) Nyquist plot for AISI 316L SS, LCS, and the HED clads at various scan speeds. 10

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suggesting a near ideal capacitive behavior (indicated by a phase angle of −90°). On the other hand, LCS showed higher phase angles, in the range of −60° to −63° in a very short range of frequencies (at around 5–10 Hz), which is a characteristic of unprotective oxide film formation. The 100 mm/s clads showed phase angles similar to AISI 316L SS; however, with increasing scan speeds, the phase angle increased, and eventually, the clad produced at 400 mm/s laser scan rate showed comparable values to LCS. The observations that are presented in the Bode plots for the impedance and phase angle are also confirmed in the Nyquist plot, as shown in Fig. 12(c). The EIS data were further analyzed using equivalent circuit modeling by fitting the data to a Randles circuit, which consisted of a resistor representing solution resistance, Rs, in series with a resistor (Rct) and constant phase element (CPE) that are connected in parallel to represent the electrochemical processes that take place on the clad surface. This simplified equivalent circuit model has been used by Blanco et al. (2006) and Liu et al. (2017) for passivating stainless steel in high pH concrete pore solution and corroding metal surfaces to quantify assumed circuit elements and their corresponding physical representations. Constant phase element represents the imperfect capacitive behavior of the metal/electrolyte interface, and it is quantified by a capacitance term (CCPE), and a constant (α), a similar representation of the electrical elements were used by Orazem and Tribollet (2017). The results of equivalent circuit modeling are presented in Table 3. The CPE parameter α was found to be between 0.9 and 1 for AISI 316L SS and clads produced at slow laser scan speeds (100 mm/s and 200 mm/s). Note that α values close to 1 indicate passivity of the metal surface. The α values for clad produced at 400 mm/s and LCS were similar (˜0.86). The highest Rct value was observed for AISI 316L SS, followed by clads that were produced at 100 mm/s laser scan speed. The charge transfer resistance drops significantly for higher scan speeds, which is an indication that they are not as protective as the clads produced were 100 mm/s laser scan speed. 3.5.3. Cyclic polarization (CP) Fig. 13 illustrates the CP curves for clads, LCS and AISI 316L SS. As shown in Fig. 13(a), the AISI 316L SS had the highest pitting potential (Ep = 324 mV vs. SCE), whereas LCS exhibited the lowest pitting potential (Ep = −650 mV vs. SCE). Higher pitting potentials indicate stronger resistance to pitting. As shown in Fig. 13(a) and (b), the pitting potentials for the clads were in between these two cases and were considerably higher than those of LCS. As shown in Figs. 13(b) and 14 , pitting potentials were generally correlated inversely with laser scan speed. For the clads that were produced at 100 mm/s laser scan speed, the clad with 50% hatch overlap had a higher Ep (73.44 mV vs. SCE) than clad with no hatch overlap (-26.2 mV vs. SCE). The difference of Ep between LCS and chromium containing specimens (i.e., clads and AISI 316L SS) can be mainly attributed to the presence of chromium in the latter. in their text on iron oxides showed that oxides that form on the LCS in neutral conditions are thermodynamically unprotective. The difference in Ep between clads and AISI 316L SS could be attributed to various factors such as chemical composition, surface roughness, and microstructural differences arising from different processing techniques. AISI 316L SS had higher

Fig. 13. Cyclic polarization curves for (a) AISI 316L SS, LCS and clads at 100 mm/s (b) clads produced at scan speeds 100–400 mm/s. The curves are marked with the notations “F” and “B,” denoting forward and backward scan directions, respectively.

chromium content than clads (18% vs. 13–15%) as shown in Table 1 and Fig. 8. However, chromium content alone cannot be used to explain decreasing Ep for clads with increasing laser scan speeds because all clads had relatively large chromium content (˜13–15%) that is higher than typically required for good corrosion resistance. For that explanation, microstructural differences and surface roughness need to be considered. The AISI 316L SS was produced via conventional route of casting, resulting in large austenitic grains as presented in Fig. 1(c). However, all clads produced by PBF-SLM show columnar grains because of rapid unidirectional cooling. On comparing grain sizes, LCS had the smallest grains, followed by relatively coarser columnar grains in clads and the largest grain size in AISI 316L SS. The grain boundary area was larger in clads than AISI 316L SS, and they increased progressively with laser

Table 3 Equivalent circuit modeling parameters after EIS data fitting. Material

Avg. Rs (ohm cm.2)

Rct × 103 (ohm cm.2)

CCPE (×10−6) ohm−1 cm−2 sα

α

Goodness of Fit

AISI 100 100 200 400 LCS

17.5

505.1 162.2 104.7 6.46 7.24 0.796

27.74 34.31 32.89 35.39 37.00 318.2

0.935 0.933 0.928 0.908 0.860 0.858

1.197 × 10−3 1.508 × 10−3 2.836 × 10−3 2.489 × 10−3 1.000 × 10−3 2.526 × 10−3

316L SS mm/s Ov = 50% mm/s mm/s mm/s

11

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2

3

4

5

6 Fig. 14. Pitting potentials as obtained from CP curves and corrosion rates calculated from LPR data.

scan speeds. These grain boundaries are high-defect and high-energy areas, often serving as pit nucleation sites, and could partially explain the lower Ep of clads than of AISI 316L SS. Furthermore, AISI 316L SS specimens had highly polished surfaces, which drastically reduces the number of pit nucleation sites apart from the grain boundary effect. On the contrary, the clad surfaces have inherently higher surface roughness due to non-uniform melt pool morphology associated with the PBF-SLM melting process. As previously discussed, increasing scan speeds causes balling, thereby increasing surface roughness. High clad surface roughness increases pit nucleation sites, and thereby lowering pitting potential below AISI 316L SS. Increasing surface roughness and decreasing grain size with scan speeds result in decreasing pitting potential as shown in Fig. 14.

333–1333 J/mm3 had to be employed for cladding operation to achieve better bonding and minimize the described issues. A maximum clad thickness of 133.14 μm was achieved at the lowest scan speed (100 mm/s) and increasing laser scan speeds reduced the clad layer thickness. The 316L SS clads had lower chromium content than the SS powder due to evaporative losses experienced during laser melting process. However, chromium contents in the range of 13–15% were successfully achieved in all clads. Metallurgical characterization revealed the formation of fine δ-ferrite and coarse perlite phases in HAZ beneath the clad layer. In addition, the HAZ thickness decreased with increasing scan speeds. Increasing laser scan speeds had a negative impact on the nanoindentation hardness of the clads; however, the clad hardness at all scan speeds was found to be higher than AISI 316L SS. The corrosion properties of the clads produced at 100 mm/s were comparable to AISI 316L stainless steel. Furthermore, the 50% hatch overlap resulted in better corrosion resistance as compared to 0% overlap. With increasing scan speeds, the corrosion resistance of the clads decreased owing to increased surface roughness and decreasing relative grain size.

Further improvements in service life and corrosion resistance of the clads could be achieved by increasing clad thickness and using heat treatment methods. Acknowledgments The authors would like to acknowledge the support from Reyixiati Repukaiti for nano-indentation hardness measurements and OSU ATAMI facility and staff. References

3.5.4. Linear polarization resistance (LPR) The corrosion rates for clads, AISI 316L SS and LCS, were measured by the LPR tests. Fig. 14 illustrates the variation of corrosion rates (μm/ year) with increasing scan speeds. Clads produced at 100 mm/s laser scan speeds and with 50% hatch overlap showed about a five-fold higher corrosion rates than AISI 316L SS, but these rates were two orders of magnitude lower than those for LCS. The increasing trend of the corrosion rates with scan speeds reaffirms the observations made in OCP, EIS and CP tests. For clads produced at 100 mm/s and with 50% hatch overlap, considering the clad thickness (133.14 μm) and the corrosion rate (1.97 μm/year), the estimated service life of the component would be 67.6 years, which is a considerable improvement as compared to the service life of low carbon steel. During this period (i.e., 67.6 years), LCS would lose 19.99 mm thickness, on average, assuming the measured corrosion rate of 295.66 μm/year. Additional works to increase the quality, corrosion resistance, and service life of the clads produced via PBF-SLM are in progress. These works involve improved processes to increase clad thickness, surface polishing, and heat treatment. Furthermore, the effects of different PBFSLM process parameters like h, d, hatch orientation, and remelting will be investigated in the future.

ASTM, 2018. ASTM G61: Standard Test Method for Conducting Cyclic Potentiodynamic Polarization Measurements for Localized Corrosion Susceptibility of Iron-, Nickel-, or Cobalt-Based Alloys. ASTM. Bartolomeu, F., Buciumeanu, M., Pinto, E., Alves, N., Carvalho, O., Silva, F.S., Miranda, G., 2017. 316L stainless steel mechanical and tribological behavior—a comparison between selective laser melting, hot pressing and conventional casting. Addit. Manuf. 16, 81–89. Bidare, P., Bitharas, I., Ward, R.M., Attallah, M.M., Moore, A.J., 2018. Fluid and particle dynamics in laser powder bed fusion. Acta Mater. 142, 107–120. Blanco, G., Bautista, A., Takenouti, H., 2006. EIS study of passivation of austenitic and duplex stainless steels reinforcements in simulated pore solutions. Cem. Concr. Compos. 28, 212–219. Cacace, S., Demir, A.G., Semeraro, Q., 2017. Densification mechanism for different types of stainless steel powders in selective laser melting. Proc. CIRP 62, 475–480. Chandrasekhar, S., 1981. Hydrodynamic and Hydromagnetic Stability. Dover Publications. Chao, Q., Cruz, V., Thomas, S., Birbilis, N., Collins, P., Taylor, A., Hodgson, P.D., Fabijanic, D., 2017. On the enhanced corrosion resistance of a selective laser melted austenitic stainless steel. Scripta Mater. 141, 94–98. Cherry, J.A., Davies, H.M., Mehmood, S., Lavery, N.P., Brown, S.G.R., Sienz, J., 2014. Investigation into the effect of process parameters on microstructural and physical properties of 316L stainless steel parts by selective laser melting. Int. J. Adv. Manuf. Technol. 76, 869–879. Cullity, B.D., 1978. Elements of X-ray Diffraction. Addison-Wesley Publishing Company. Deqing, W., Ziyuan, S., Ruobin, Q., 2007. Cladding of stainless steel on aluminum and carbon steel by interlayer diffusion bonding. Scr. Mater. 56, 369–372. Farren, J.D., Dupont, J.N., Noecker II, F.F., 2007. Fabrication of a carbon steel to stainless steel transition joint using direct laser deposition- a feasibility study. Weld. J. 55s–61s. Gunenthiram, V., Peyre, P., Schneider, M., Dal, M., Coste, F., Koutiri, I., Fabbro, R., 2018. Experimental analysis of spatter generation and melt-pool behavior during the powder bed laser beam melting process. J. Mater. Process. Technol. 251, 376–386. Huang, Y., Khamesee, M.B., Toyserkani, E., 2016a. A comprehensive analytical model for laser powder-fed additive manufacturing. Addit. Manuf. 12, 90–99. Huang, Y., Yang, L.J., Du, X.Z., Yang, Y.P., 2016b. Finite element analysis of thermal behavior of metal powder during selective laser melting. Int. J. Therm. Sci. 104, 146–157. Ishida, T., 1991. Formation of stainless steel layer on mild steel by welding arc cladding. J. Mater. Sci. 6431–6435. Jing, Y.-a., Qin, Y., Zang, X., Shang, Q., Hua, S., 2014. A novel reduction-bonding process to fabricate stainless steel clad plate. J. Alloys Compd. 617, 688–698.

4. Conclusions The PBF-SLM technique was successfully used for cladding corrosion-resistant 316L stainless steel layer on 1018 carbon steel substrate. The following specific conclusions were drawn: 1 The reported energy density of E = 104.52 J/mm3 for producing 3D parts proved to be inadequate for the bonding of two dissimilar metals and resulted in extensive balling, geometric warping, and clad delamination. A higher energy density in the range of 12

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2016. Denudation of metal powder layers in laser powder bed fusion processes. Acta Mater. 114, 33–42. Orazem, M.E., Tribollet, B., 2017. Electrochemical Impedance Spectroscopy. Wiley. Rashid, R.R.A., Abaspour, S., Palanisamy, S., Matthews, N., Dargusch, M.S., 2017. Metallurgical and geometrical characterisation of the 316L stainless steel clad deposited on a mild steel substrate. Surf. Coat. Technol. 327, 174–184. Samant, R., Ranjan, R., Mhapsekar, K., Anand, S., 2018. Octree data structure for support accessibility and removal analysis in additive manufacturing. Addit. Manuf. 22, 618–633. Simson, T., Emmel, A., Dwars, A., Böhm, J., 2017. Residual stress measurements on AISI 316L samples manufactured by selective laser melting. Addit. Manuf. 17, 183–189. Straumal, B.B., Kogtenkova, O.A., Kolesnikova, K.I., Straumal, A.B., Bulatov, M.F., Nekrasov, A.N., 2014. Reversible “Wetting” of grain boundaries by the second solid phase in the Cu-In system. JETP Lett. 100, 535–539. Straumal, B.B., Kucheev, Y.O., Efron, L.I., Petelin, A.L., Majumdar, J.D., Manna, I., 2012. Complete and incomplete wetting of ferrite grain boundaries by austenite in the lowalloyed ferritic steel. J. Mater. Eng. Perform. 21, 667–670. Sufiiarov, V.S., Popovich, A.A., Borisov, E.V., Polozov, I.A., Masaylo, D.V., Orlov, A.V., 2017. The effect of layer thickness at selective laser melting. Procedia Eng. 174, 126–134. Suryawanshi, J., Prashanth, K.G., Ramamurty, U., 2017. Mechanical behavior of selective laser melted 316L stainless steel. Mater. Sci. Eng. A 696, 113–121. Weymouth, B.D.C., John, W., 2005. Elements of X-Ray Diffraction. Yoshimura, T., 1987. In: Office, E.P (Ed.), Corrosion-Resistant Clad Steel and Method for Producing the Same.

Khairallah, S.A., Anderson, A., 2014. Mesoscopic simulation model of selective laser melting of stainless steel powder. J. Mater. Process. Technol. 214, 2627–2636. Khairallah, S.A., Anderson, A.T., Rubenchik, A., King, W.E., 2016. Laser powder-bed fusion additive manufacturing: physics of complex melt flow and formation mechanisms of pores, spatter, and denudation zones. Acta Mater. 108, 36–45. Khara, S., Choudhary, S., Sangal, S., Mondal, K., 2016. Corrosion resistant Cr-coating on mild steel by powder roll bonding. Surf. Coat. Technol. 296, 203–210. Li, C., Liu, Z.Y., Fang, X.Y., Guo, Y.B., 2018. Residual stress in metal additive manufacturing. Procedia CIRP 71, 348–353. Li, R., Liu, J., Shi, Y., Wang, L., Jiang, W., 2011. Balling behavior of stainless steel and nickel powder during selective laser melting process. Int. J. Adv. Manuf. Technol. 59, 1025–1035. Liu, G., Zhang, Y., Wu, M., Huang, R., 2017. Study of depassivation of carbon steel in simulated concrete pore solution using different equivalent circuits. Constr. Build. Mater. 157, 357–362. Liverani, E., Toschi, S., Ceschini, L., Fortunato, A., 2017. Effect of selective laser melting (SLM) process parameters on microstructure and mechanical properties of 316L austenitic stainless steel. J. Mater. Process. Technol. 249, 255–263. Ma, M., Wang, Z., Zeng, X., 2017. A comparison on metallurgical behaviors of 316L stainless steel by selective laser melting and laser cladding deposition. Mater. Sci. Eng. A 685, 265–273. Majumdar, J.D., Pinkerton, A., Liu, Z., Manna, I., Li, L., 2005. Mechanical and electrochemical properties of multiple-layer diode laser cladding of 316L stainless steel. Appl. Surf. Sci. 247, 373–377. Matthews, M.J., Guss, G., Khairallah, S.A., Rubenchik, A.M., Depond, P.J., King, W.E.,

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