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Additive Manufacturing

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Full Length Article

Additively manufactured 316L stainless steel with improved corrosionresistance and biological response for biomedical applications

M.J.K. Lodhia, K.M. Deenb, M.C. Greenlee-Wackerc, Waseem Haidera,⁎

a School of Engineering and Technology, Central Michigan University, Mt. Pleasant, MI, 48859, USAbDepartment of Materials Engineering, The University of British Columbia, Vancouver, BC, V6T 1Z4, Canadac Department of Biology, Central Michigan University, Mt. Pleasant, 48859, MI, USA

A R T I C L E I N F O

Keywords:Selective laser meltedStainless steelOxide filmOsteoblast

A B S T R A C T

Enhancing the corrosion resistance and improving the biological response to 316 L stainless steel is a long-standing and active area of biomedical research. Here, we analyzed the structure and corrosion tendency ofselective laser melted-additively manufactured (AM) 316 L stainless steel (AM 316L SS) and its wrought coun-terpart. SEM analysis showed a fine (500–800 nm) interconnected sub-granular structure for the AM 316L SS,but a polygonal coarse-grained structure for the wrought sample. Relative to the wrought sample, the AM 316LSS also exhibited a higher charge transfer resistance and higher breakdown potential (˜1000mV vs. SCE) whentested in biological electrolytes, which included human serum, PBS, and 0.9M NaCl. A higher pitting resistance(extended passive region) and improved stability of the AM 316L SS was attributed to its dense structure of oxidefilm and refined microstructure. Finally, material compatibility with pre-osteoblasts was analyzed. Large cyto-plasmic extension of osteoblast cells and retention of stiller morphology was observed when cells were culturedon the AM 316L SS as compared to its wrought counterpart, suggesting that the AM 316L SS was a bettersubstrate for cell spreading and differentiation. The differentiation of cultured cells was further validated bywestern blot for Runx2. Runx2, an anti–proliferative marker indicative of differentiation, was equivalent in cellscultured on either samples, but overall more cells were present on the AM 316L SS. Given its higher corrosionresistance and ability to support osteoblast adherence, spreading and differentiation, the AM 316L SS has po-tential for use in the biomedical industry.

1. Introduction

Austenitic stainless steel is one of the commonly used alloys forbiomedical applications, including surgical instruments, orthopedicimplants, fixtures, orthodontics and pharmaceutical equipment. Thewide spread use of stainless steel can be attributed to its reasonablecost, ease of fabrication, biocompatibility, adequate mechanicalstrength and corrosion resistant properties [1–4]. However, stainlesssteel also has a tendency to exhibit localized corrosion, hampering itsstructural applications, particularly within biomedical industry, and24% of the implant failures are caused by this corrosion phenomenon[5]. Localized corrosion of the metallic implants leads to release ofmetallic ions in the surrounding tissues that could initiate inflammatoryand adverse cellular reactions. The structural integrity of the implantsinfluences the quality of life, since implant failure may cause severepain and repetitive post-surgical operations [6–9]. In stainless steel,MnS inclusions and the resulting chromium–depleted region

surrounding these inclusions results in the heterogeneous structuredoxide film and provide the initiating points for localized corrosion[10,11]. The long-term stability and improved biocompatibility ofstainless steel is still a challenge and solutions to these problems con-tinue to be explored. Various approaches, including the application ofcoatings, surface modification, laser surface treatment and grain re-finement have been adopted to overcome these issues with stainlesssteel [12–16]. Coatings are considered as an ineffective approach tothese problems because of their limited mechanical stability and de-fective structure. Surface modification through ion implantation mayreinforce the stability of the surface oxide film and this method is foundto be an appropriate remedy to minimizes corrosion damage of thestainless steel implants [17,18]. However, the high-energy bombardedions could adversely affect the surface topographical features, whichmay promote excessive dissolution and premature failure of the mate-rial. Surface laser melting was posited to solutionize the MnS inclusionsby selectively melting the stainless steel surface and then rapidly

https://doi.org/10.1016/j.addma.2019.02.005Received 27 October 2018; Received in revised form 29 January 2019; Accepted 11 February 2019

⁎ Corresponding author.E-mail address: haide1w@cmich.edu (W. Haider).

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Available online 16 February 20192214-8604/ © 2019 Elsevier B.V. All rights reserved.

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quenching it [19]. Currently, one of the most promising approaches toaddress the limitations of stainless steel is surface texture modificationat the nano-scale. In support of this notion, improved corrosion re-sistance and enhanced in-vivo cellular growth has been reported byforming nano-texture features on the surface of stainless steel [20].Compared with other commercial materials, several reports show thatalloys produced by nanoscale modification (refined grain structure)exhibit an improved corrosion response and enhance metabolic activ-ities of cells in culture [21–28]. In stainless steel, refined grain structurecould effectively decrease the path for chromium diffusion at the sur-face and may result in thickening of barrier oxide film, leading to im-proved corrosion resistance [29]. It is also considered that variation insurface roughness, texture, hydrophilicity and grain structure couldplay an important role in modifying the cellular response to the implantmaterials [30]. The higher density of interphase boundaries (high andlow angle grain boundaries) present in the fine grained material pro-vides regions of high energy that favor the growth of cells and resultinto better cellular response [31]. In order to obtain nano to sub-micronsize grain structure in stainless steel different thermo-mechanical op-erations have been reported in the literature [32–36]..

Recently, additive manufacturing has emerged as a net shape pro-duction process with advantages associated to this process. Although arelatively new field, additive manufacturing technology is finding awide range of applications in the biomedical industry. This technologyhas lowered the entry barrier for new medical devices, and allows moredesign flexibility. Additive manufacturing can be used to manufactureorthopedic implants with a porosity that mimics human bone and couldbe customized in size or shape to the patient’s anatomical data [37].Additive manufacturing has intrinsically a very high cooling rate in therange of 103 – 108 K/s that results into a very non-conventional sub-granular structure [38,39]. The variation in the processing parametersof the additively-manufactured materials can significantly alter thesurface and bulk properties of the materials [40,41]. The micro-structural studies of additively manufactured 316 L stainless steel andsignificance of grain refinement in terms of mechanical properties hasbeen well discussed in the literature [42–44]. Wang et al. recentlyproposed additive manufacturing, a solution to overcome the strength-ductility tradeoff paradigm for 316 L stainless steel [42]. However, thecorrosion response of the additively manufactured materials is a rela-tively new and very active area of research. Corrosion is a phenomenonthat is highly dependent on the microstructural variations and on thenature of environment. Recent studies demonstrate that differentthermal processes and environmental conditions could significantlyinfluence the corrosion properties of the additively manufactured 316 Lstainless steel [45,46]. Furthermore, the biological response of addi-tively manufactured 316 L stainless steel remains less well defined andthere is an urgent need to explore these properties for biomedical ap-plications. Therefore, in this study the corrosion response of selectivelaser-melted (SLM) 316 L stainless steel in various biological environ-ments and proliferation and differentiation of the pre–osteoblast cellswere assessed.

2. Experimental procedure

2.1. Sample preparation

The additively-manufactured 316 L stainless steel, called the AM316L SS in this report, was fabricated by selective laser melting method.AM-250 unit (RENISHAW) equipped with ytterbium continuous laserbeam (λ=1060 nm) operating at a power level of 200W was used forthe fabrication of the AM 316L SS. The high-energy beam was used tofuse powder particles (35–50 μm) and multilayers of materials wereproduced by keeping a single layer thickness of 30 μm. The hatch dis-tance of 100 μm was adjusted to fabricate the AM 316L SS. The sampleswere fabricated under optimized conditions that were preset to avoidporosity and anomalous microstructural variations, including, but not

limited to, lack of fusion, formation of metastable phases, micro-cracksand non-metallic inclusion. The AM 316L SS samples were fabricatedinto circular disks with diameter of 15.8 cm and thickness of 0.5 cm.The wrought 316 L stainless steel was purchased from “OnlineMetals”and used as-received for comparison. The 0.5 cm thick wrought sampleswere cut from a one foot long rod with a diameter of 15.8 cm using highspeed saw. For electrochemical testing, both the wrought and AM 316LSS samples were polished sequentially on silicon carbide (SiC) papersranging from 180 to 1200 grit size. Subsequently, the samples weredegreased ultra-sonically in acetone for 15min, rinsed in water, anddried in the nitrogen stream.

2.2. Surface characterization

A Rigaku MiniFlexII diffractometer was used to acquire the X-raydiffraction (XRD) patterns. The X-rays were produced from the Cu K α(λ=0.154 nm) radiation source. The scan was performed over 2θ from30⁰ to 90⁰ to identify the constituent phases in the wrought and AM SSsamples. The elemental compositional analysis and the microstructuralfeatures of these samples were investigated via scanning electron mi-croscope (Hitachi S-3400-II) equipped with EDX analyzer. The sampleswere polished in a sequence by using 180 to 1200 grit SiC papers fol-lowed by fine polishing using three and one μm diamond suspension formicrostructural analysis. Following fine polishing, the samples wereultra-sonicated in acetone for 15min and washed in water.Furthermore, the polished samples were etched in a solution of 10mlHNO3 (70%), 15ml HCl (37%), 10ml CH3OOH (99%) and three dropsof glycerol (ASTM E-407). Contact angle measurement was performedusing kyowa contact angle meter (DM-CE1, Japan). Surface roughnesswas analyzed by using atomic force microscopy (AFM), scanning20 μm×20 μm surface area operated under tapping mode.

X-ray photoelectron spectra (XPS) were obtained to analyze thesurface chemistry of the samples. The XPS analyses were performed in aKratos Axis ultra-photoelectron spectroscope equipped with a mono-chromatic aluminum excitation source. Photoelectrons were detected at45⁰ take off angle.

2.3. Electrochemical analysis

Electrochemical performance of the samples was evaluated by usingReference 1000E potentiostat (GAMRY Instruments). All the electro-chemical experiments were conducted in a three-electrode cell. Theworking electrodes were polished samples having a 1.26 cm2 exposedarea. In this cell, the counter and reference electrodes were Pt coil andsaturated calomel electrode (SCE) (0.240 V vs. SHE), respectively.Human serum, phosphate buffer saline (PBS) and 0.9M NaCl, wereused separately as electrolytes to study the corrosion behavior. Theserum was obtained from whole blood that was allowed to clot in aglass tube for 30min at 37 °C followed by 30min at 4 °C. The clot wasremoved by centrifugation at 2000 rpm for 10min, the liquid interfacewas collected, and samples were frozen at −80 °C. To prepare the PBSsolution, PBS tablets (Sigma Aldrich) were dissolved in deionized water,and the 0.9M NaCl solution was prepared by dissolving calculatedamount of NaCl (Sigma Aldrich) in deionized water.

For electrochemical characterization, initially the open circuit po-tential for each sample was stabilized (0.1 mV/s). The electrochemicalimpedance spectroscopy (EIS) was recorded by applying a small ACperturbation (5mV) in the frequency range of 100 kHz to 1mHz. Toestimate the pitting tendency and stability of the passive film, po-tentiodynamic cyclic polarization tests were conducted by applying a0.5 mV/s scan rate. During forward anodic polarization, the reversescans were obtained by reaching either 1.5 V (vs. OCP) apex potentialor producing 10mA/cm2 current density, whichever is achieved first.

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2.4. Cell culture study

To investigate the cell’s interaction with the wrought and AM 316LSS samples, the murine pre-osteoblasts cell line MC3T3-E1 (subclone 4)was used. For cell culturing, α modified minimum essential mediumhaving 1% of penicillin streptomycin and 10% fetal bovine serum wasused, placed in the controlled humidified environment at 37 ⁰C and 5%CO2. The cells used in experiments were from the fourth passage.15,000 cells (measured using Countess II FL, life technologies) wereseeded on the circular disks of polished wrought and AM samples andincubated for 24 and 48 h. Following the incubation time, the cells werestained using NucBlue Live cell stain ready probes (Invitrogen Inc.) andwere fixed by using 4% paraformaldehyde. The cells were dehydratedby using methanol then acetone for 10min each and washed with PBSsolution. Furthermore, the cells were stained using ActinRed™ 555ready probes reagent for 30min at 37 °C to visualize the F-actin fila-ments of cytoskeleton. The imaging was done by using EVOS® FL cellimaging system (AFM4300, Invitrogen Inc.).

2.5. Western blot study

MC3T3 cells cultured on either wrought and AM samples were lysedin RIPA buffer. Protein (50 μg) from cell lysates was separated by SDSpolyacrylamide gel electrophoresis (PAGE) and transferred on to PVDFmembrane. The membranes were blocked with 5% non-fat dry milk intris-buffer saline (TBS) for 1 h at ambient temperature. Membraneswere then incubated in antisera raised against Runx2 at a 1:1000 di-lution at 4 °C overnight. After incubation, the membranes were washedthree times with tris-buffered saline containing 0.1% Tween20 (TBST)and incubated with a suitable near-Infrared fluorescent secondary an-tibody (LI−COR) at a 1:5000 dilution in 1% non-fat dry milk for 1 h indark. After washing three times, the immune complexes involving thetarget proteins were detected using Odyssey imaging system fromLI−COR. The membranes then stripped and reprobed for GAPDH. Theband intensities were quantified using Image Studio software andprotein expression levels were plotted relative to GAPDH from the samemembrane.

3. Results and discussion

3.1. Elemental composition and X-ray diffraction analysis

Energy dispersive X-ray spectroscopy (EDX) was used to determinethe elemental composition of major alloying elements. Although thistechnique provides some information about the composition of thefabricated material, the concentration of C, S, P, and O in the wroughtand AM 316L SS samples were below the detection limits of EDX ana-lyzer and a more sensitive assay was employed. Standard photo spec-troscopic procedure (ASTM-350-18) was adopted to confirm the com-position of these alloy samples and the wt. % of these alloying elementswere approximately within 2% margin of error. The wrought and AM316L SS samples showed a similar elemental composition, and corre-sponded to the elemental composition of the ASTM 316 L stainless steelgrade (Table 1).

The structures of the wrought and AM 316L SS samples were ex-amined using X-ray diffraction (XRD) (Fig. 1a). The diffraction patternsfor the wrought sample revealed three peaks originating from the (111),(200) and (220) at 43.8⁰, 50.9⁰ and 74.8⁰, respectively. These peaks

were associated with the austenitic matrix phase of conventional 316 Lstainless steel. The diffraction pattern matched with the JCPDS 31-0619and confirmed the austenitic phase [39]. In addition to the three aus-tenitic phase peaks, a very small shoulder peak at 44.2⁰ was observed inthe wrought sample. Although this peak could have originated from the(110) plane of δ-ferrite phase in the wrought sample’s microstructure,no other characteristics peak of δ-ferrite was evident in the diffractionpattern. Using the WRC-1992 equivalent formulas [47], a Creq/Nieqratio greater than 1.69 indicates the existence of δ-ferrite as primaryphase in the austenitic stainless steels. The low Creq/Nieq ratio of 1.54for the wrought samples and 1.49 for the AM 316L SS samples, sup-ported the conclusion that a negligible concentration of δ-ferrite waspresent in these alloys. Although δ-ferrite may form during the laserpulsed welding process, it is not the case with the existing samplesunder investigation [48]. Likewise, the XRD pattern of the AM 316L SSsample had only the three austenitic phase signature diffraction peaks.One other subtle difference was noticed between the wrought and AM316L SS samples: The peak associated with the (220) plane for the AM316L SS sample was relatively more intense than wrought sample, andsince no additional diffraction peak was observed for the AM 316L SSsample, the fabricated material is of high metallurgical purity.

The microstructures of the wrought and AM 316L SS samples wereanalyzed using scanning electron microscope and images are presentedin Fig. 1(b, c, d). The wrought sample (Fig. 1b) showed an equiaxedgrain structure having sharp grain boundaries and coarser grains. Theparallel lines observed within the grain structure represented the twinbands, and could be due to the prior mechanical treatment of thewrought sample. For the AM 316L SS, the localized melting, extremelyhigh solidification rate and large temperature gradient within a narrowregion, resulted in the development of non-conventional microstructure(Fig. 1c). In contrast to the conventional large austenitic grain structureof the wrought sample, fine, cellular and columnar shaped sub–grainsare evident in the AM 316L SS sample. The well-oriented and inter-connected cellular and columnar sub–grains structure formed colonies,as shown in Fig. 1c. From a wider view, individual colonies could berecognized as large grain, and pattern similar to the wrought samplecould be visualized. The size of the sub–grains were found to be be-tween 500–800 nm in the AM 316L SS sample as shown in Fig. 1d.These data are consistent with previous reports that affiliated manu-facturing conditions [42,49], results in a similar grain structure to theAM 316L SS. Overall, the fine microstructural features of the AM 316LSS is of particular interest, since the structure of this additively man-ufactured material could improve the electrochemical and biologicalperformance of stainless steel.

To determine the wettability behavior of the samples, a key variablepromoting cell spreading and proliferation [50], contact angle mea-surements were performed. In measures of surface wettability, a lowcontact angle indicates a highly hydrophilic surface, and favors proteinadsorption, cell spreading, and cell attachment. Cell spreading andgrowth occur on hydrophilic surfaces because they promote the de-position of extracellular matrix and attachment via adhesion moleculestethering to intracellular actin filaments. The average value of thecontact angle for a water droplet on the wrought samples was 86.22ᵒ,and was reduced to 78.69ᵒ for the AM 316L SS samples (Fig. 1e). Thesedata support the conclusion that the AM 316L SS sample has a morehydrophilic surface, which could be attributed to the refined micro-structure obtained in the AM 316L SS samples, and suggests that the AM316L SS sample may prove to be a better substrate for the adhesion andproliferation of pre-osteoblast cells. In addition to the microstructuralfeatures, the surface wettability and adhesion of cells over the metallicsurfaces also depend on the surface roughness, which is affected by thesurface preparation method and properties of materials [50]. In thiscase both the wrought and AM 316L SS samples were prepared simi-larly by polishing on the SiC paper upto 1200 grit size. It is therefore torule out the influence of surface roughness, it is important to identifythe topographical features of the surface. To address this feature, AFM

Table 1Chemical composition (in wt. %) of the wrought and AM 316L SS samples.

Elements C Mn Mo S P O Ni Cr Fe

Wrought 0.03 2.3 2.1 0.03 0.06 0.04 11.5 17.1 Bal.AM 0.03 2.2 2.3 0.03 0.05 0.05 11.2 17.5 Bal.

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analysis was performed and the results are tabulated in Fig. 2. Typicalgrinding marks were evident on both wrought and AM 316L SS sam-ples. The almost similar Ra and Rq values of both samples revealed thatmajor difference in the surface characteristics of these sample i.e.wettability, cell adhesion and growth and electrochemical behavior canbe attributed to the difference in the microstructural features, (as

shown in Fig. 1b–e) which are intrinsic to the manufacturing method.

3.2. XPS analyses for the surface characterization

X-ray photoelectron spectroscopy (XPS) was done to further analyzethe chemistry of the samples. Fig. 3 shows the survey spectra of the

Fig. 1. (a) X-ray diffraction pattern of the wrought and AM 316L SS samples microstructure of the (b) wrought and (c) AM 316L SS sample (d) magnified section ofsub–granular structure from the AM 316L SS sample (e) contact angle measurement (n=5) for the wrought and AM 316L SS samples.

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surfaces without any etching process. The spectra for both the wroughtand AM 316L SS samples indicated the presence of Fe, Cr, Ni, Mo, Mnand O at the surface, but no difference in the intensity or positioning ofthe peaks was seen, in the spectra of either samples. Interestingly, adominant oxygen peak in the spectra could indicate the presence of apassive oxide film on the surface of the samples.

We measured the oxidation state of the O, Cr and Fe in the passivefilm by obtaining the high-resolution XPS spectra. The deconvolutionresults of these high-resolution spectra are presented in Fig. 4. Decon-volution was done by using the Gaussian function with the backgroundsubtraction using Shirley method. The spectra for O 1 s indicates thepresence of both hydroxide (OH–) (531.8 eV) and oxide (O2–)(530.4 eV) species in the passive film formed at the surface of both thesamples. Next, the oxidation state of Cr was analyzed, and finally Fe.The Cr2P3/2 spectra were dissociated into the binary peaks affiliatedwith the metallic (Cr°) and ionic states (Cr+3). The broad peak asso-ciated with Cr3+ suggested that the formation of hydroxide(Cr3+−OH–) is more feasible than chromium oxide. Additionally, asmall doublet peak of Cr3+– O2– species originating at 576.2 eV sup-ports our conclusion that a hydroxide is more concentrated at thesurface. The deconvolution of Fe2P3/2 peak resulted in four peaks, as-sociated with the metallic iron Fe° (707.2 eV), Fe+2 (710.5 eV) andFe+3 (711.3, 713.5 eV) species. The broad peak at 713.5 eV corre-sponds to the Fe3+−OH– species, whereas the peak at 711.3 eV in-dicates the presence of Fe3+– O2– species. The relatively small collarpeak at 710.5 eV affiliated with the Fe2+– O2– species and these resultsconfirmed the presence of both Fe2+ and Fe3+ species in the iron oxidephase at the surface of both the wrought and AM 316L SS samples. The

separate collar peak at low binding energy (707.2 eV) corresponded toFe°. Since the prominent signals originated at high binding energy(713.5 eV) for Fe3+−OH–, it is likely that the oxide film is composed ofboth hydroxide and oxide phases. After this comprehensive analysis ofthe XPS spectra (Fig. 4), no differences in the surface chemistry of thepassive film were found between the wrought and AM 316L SS samples.

The composition and structure of the oxide film is very important, asit determines its electrochemical stability when it exposed to a specificenvironment. Fig. 5 illustrates the concentration profile of the majorconstituents (O, Fe, Cr and Ni) present in the oxide film. The highconcentration of oxygen is evident at the start of surface etching, andoxygen decays upon extended etching time. These data validated thatoxygen was enriched in a surface layer on both the samples. The con-centration of Fe, Cr and Ni are modestly increased during initial ˜200 sof etch time in each sample before becoming constant. However, therelatively high concentration of ‘Fe’ (˜11%), compared to other passivefilm constituents at the start of etching (0 s) indicates that the outerregion of the oxide film is primarily composed of oxide/hydroxide ofiron. Furthermore, a sharp increase in the concentration of chromiumafter 30 s etching indicates the enrichment of ‘Cr’ in the oxide film. Thisbehavior suggested the formation of duplex oxide film composed of ‘Fe’and ‘Cr’ oxides/hydroxides enriched outer and inner layers, respec-tively. After ˜600 s, the oxygen concentration drops to< 10% and be-comes constant which indicates the metal/oxide interface. These dataare consistent with other well established observations that duplexoxide film forms on surface of 316 L stainless steel [51–55].

3.3. Electrochemical behavior of the wrought and AM 316L SS samples inbiologically relevant electrolytes

Since an oxide film may form spontaneously on the surface of thewrought and AM 316L SS samples when exposed to air, EIS analysis wascarried out by placing the two samples in different electrolytes (serum,PBS and 0.9M NaCl). Nyquist plots are shown in Fig. 6 a–c and can beused to estimate the stability of the oxide film on the wrought and AM316L SS samples. The Nyquist plots of both the samples in all theelectrolytes showed one depressed semi-circle. The larger diameter ofthe Nyquist plot represents better stability of passive film and higherresistance to electrochemical dissolution. In order to simulate theelectrochemical processes at the interface of electrode/electrolyte, andquantify the charge transfer resistance and the charge distributionwithin the electrical double layer, an equivalent electrical circuit (EEC)(Randle circuit) model (Fig. 6d) was employed. The estimated seedvalues were inserted in the software (Echem analyst – 6.25) and iter-ated to achieve simulated curve.

Fig. 2. AFM roughness analysis for the wrought and AM 316L SS samples.

Fig. 3. The XPS survey spectra of the wrought and AM 316L SS samples.

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The high frequency impedance behavior is associated with anelectrolyte resistance and represented as ‘Rs’ in the EEC. The impedanceresponse at medium and low frequency is associated with the chargedistribution in the double layer (non–faradaic) and faradaic processesoccurring at the interface of electrode and electrolyte. In the EEC theseprocesses are quantified as constant phase element ‘Qo’ and chargetransfer resistance ‘Rct’, respectively. Due to the depressed semi–circlein the Nyquist plots, phase shift (> –90°) and power coefficient n< 1the ‘Qo’ instead of pure double layer capacitance was used. ‘Qo’ attri-butes to the non–homogeneous distribution of the charge at an oxide/electrolyte interface and it is related to the oxide film dielectric char-acteristics and nature of the ionic species present at the oxide/elec-trolyte interface. The stability of the oxide film is indicated as ‘Rct’ andrepresents the resistance to charge transport through the oxide film and

across the oxide/electrolyte interface.Fig. 7a represents the values of ‘Qo’ calculated from Nyquist plots

(Fig. 6a–c) using EEC model (Fig. 6d) in different electrolytes. The re-latively lower ‘Qo’ values (28.7 μS.sn-cm–2 for serum, 28.5 μS.sn-cm–2

for PBS and 51.8 μS.sn-cm–2 for 0.9M NaCl) exhibited by the AM 316LSS compared to the wrought samples and corresponded to the presenceof the least defective structure of oxide film on the AM 316L SS. Thesurficial charge distribution across the oxide/electrolyte interface de-pends on the dielectric characteristics of the oxide film, and is directlyrelated with majority charge carrier type and concentration. Similarly,the relatively large ‘Qo’ values of both AM and wrought samples in0.9 M NaCl solution corroborate the depolarizing tendency of Cl– spe-cies and interaction with the oxide film. This behavior is depicted inFig. 7b, and validated the depolarization tendency of Cl– ions by

Fig. 4. The high-resolution spectra of O 1 s, Cr 2p3/2, and Fe 2p3/2 species.

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illustrating relatively low values of ‘Rct’ by the AM and wrought 316L SSsamples in 0.9 M NaCl compared to ‘Rct’ in serum and PBS. Also, notablyhigher ‘Rct’ were registered by the AM 316L SS compared to wroughtsample (2.23 MΩ–cm2 vs. 1.66 MΩ–cm2 for serum, 2.21 MΩ–cm2 vs.1.42 MΩ–cm2 for PBS and 0.62 MΩ–cm2 vs. 0.26 MΩ–cm2 for 0.9 MNaCl) and confirmed the improved stability and barrier characteristicsof the oxide film formed on the AM 316L SS sample. Furthermore, thehighest ‘Rct’ offered by both samples was in serum, and could be asso-ciated with the specific adsorption of albumin and other proteins on thesurface. The isoelectric point for albumin and stainless steel is 4.7 and8.5, respectively [56]. Therefore, in serum (pH=7.4) the albumin andstainless steel carry opposite charges and that could facilitate

adsorption of albumin on the stainless steel surface [57]. The lowest‘Rct’ calculated was when both the samples were in 0.9 M NaCl elec-trolyte, and could be associated with the adverse interaction of Cl–species with the oxide film [58]. The presence of Cl– could acceleratelocalized dissolution reactions due to its strong depolarization tendencyand formation of water soluble complexes [59,60].

These results suggested that formation of a less defective and morecompact oxide film on the surface of the AM 316L SS sample was as-sociated to its fine sub–grain structure. Relative to serum and PBS, therelatively higher ‘Qo’ (51.8 μS.sn–cm–2) registered by the AM 316L SSsample in 0.9M NaCl solution reflected an increase in defective siteswithin the oxide film and was attributed to the interaction with Cl–

Fig. 5. The concentration profiles of ‘O’, ‘Fe’, ‘Cr’ and ‘Ni’ determined from the XPS depth profile analysis of the (a) the wrought and (b) AM 316L SS samples.

Fig. 6. Nyquist plots of the wrought and AM 316L SS samples in (a) Serum, (b) PBS, and (c) 0.9M NaCl. (d) The EEC model used to simulate the impedance spectra.

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ions. However, even in the presence of Cl– ions, the higher ‘Rct’ (0.62 MΩ–cm2) and lower ‘Qo’ (51.8 μS.sn–cm–2) offered by the AM 316L SSwas superior to values recorded using the wrought sample (‘Rct’=0.26MΩ–cm2, ‘Qo’=64.9 μS.sn–cm–2) and further affirmed the improvedcorrosion resistant properties of the AM 316L SS. The improved stabi-lity of the AM 316L SS sample against the corrosion phenomenon is inline with the other studies [22,61,62], and suggested the higher re-sistance to corrosion for stainless steel is a result of refined grainstructure. The fine grain structure could significantly enhance the dif-fusion of cationic species towards oxide/electrolyte interface and torapidly form the dense and least defective oxide film [63].

To investigate the localized corrosion tendency of the wrought andAM 316L SS samples, cyclic potentiodynamic polarization (CPP) scanswere obtained. The stability of the oxide film in the presence of dif-ferent ionic/non–ionic species can also be estimated from these CPPscans (Fig. 8a– c). The cathodic polarization represents the combinedeffect of possible H2O or dissolved oxygen reduction processes. In ad-dition, these curves also depict the kinetics of oxide film dissolution dueto thermodynamic un-stability of oxides at such low potentials. Duringanodic scan, the reformation of oxide film, stability and resistance tolocalized corrosion can be estimated from CPP. The influence of ionicspecies toward localized corrosion is determined from the hysteresisduring reverse anodic polarization. Analysis of cathodic polarizationtrends of the wrought and AM 316L SS samples in each electrolyteshowed similar charge transport characteristic at the oxide/electrolyteinterface; H2O and/or dissolved oxygen reduction according to reaction1 and 2, respectively.

2H2O + 2e– → H2 + 2OH– (1)

O2 + 2H2O+ 4e– → 4OH– (2)

The corrosion current density (icorr) obtained from the interpolationof both the cathodic and anodic curves within Tafel region are pre-sented in Table 2. Comparatively, lower icorr estimated for the AM 316LSS samples is supported with the impedance trends and suggests thefabricated material possesses better resistance to dissolution in eachelectrolyte. The Ecorr of both the wrought (−299mV vs. SCE) and AM(–331mV vs. SCE) 316L SS samples in serum was found to be slightlymore negative than observed in PBS and 0.9M NaCl, possibly asso-ciated with the tendency of the protein adsorption at the surface. On theother hand, the relatively positive Ecorr and lower icorr of the wrought(−203mV, 0.269 μA/cm2) and AM (−179mV, 0.212 μA/cm2) samplesin 0.9M NaCl indicated appreciable change in the nature of oxide filmduring cathodic polarization. In 0.9 M NaCl solution there is a strongtendency for the dissociation of water into H+ and OH– and the spe-cifically adsorbed OH– and Cl– species could affect the kinetics of filmdissolution as indicated by anodic scans. Under applied conditions, theanodic polarization trends of both the wrought and AM 316L SS

samples were associated with the intrinsic barrier characteristics of theoxide film as evident in Fig. 8.

In contrast to the characteristic passive behavior of the wrought316L SS, the AM 316L SS sample showed a much higher breakdownpotential (Ebd) in the different electrolytes. The ‘Ebd’ was defined as theinflection point for the rapid increase in the current density with a smallchange in potential. Compared to the wrought sample, the higher ‘Ebd’measured for the AM 316L SS sample in all electrolytes, confirmed thehigher stability of the passive oxide film, which developed on thisfabricated material. Furthermore, no fluctuation in the passive region ofthe current was observed. This data suggests that a homogenous andcompact oxide film forms on AM 316L SS. This higher stability con-ferred by passive oxide film on the AM 316L SS sample correlated to thefine interconnected sub–granular structure resulting from the additivemanufacturing process. Consistent with the results of this study, pre-vious studies also reported that fine grain structure is associated withenhanced stability of the passive film which developed on stainless steel[62,64,65]. Furthermore, the additive manufacturing process con-tributes to rapid solidification and reduces or completely removes thenucleation of manganese sulphide (MnS) [66]. As described earlier, theMnS inclusions and chromium depleted region in the vicinity are theinitiating points for localized corrosion of stainless steel [10].

In 0.9 M NaCl, both the wrought (200mV vs. SCE) and AM 316L SSsamples (920mV vs. SCE) showed lower Ebd than in serum and PBS,demonstrating that stainless steel is prone to Cl−-dependent pittingcorrosion. The migration of Cl– toward surface during anodic polar-ization and their affinity to react with the oxide film can be estimatedfrom these results. The mechanism of passive film breakdown can bedescribed using the point deflect model [67]. Burstein et al. [68] ex-plained the mechanism for the breakdown of passive film in the pre-sence of Cl–, and according to this model, Cl–may adsorb on the surfaceoxide film and penetrate to form soluble metal chlorides. When theconcentration of metallic chlorides in the oxide film increases relativeto the metal oxide, there may be a rupture of passive oxide film and thestable pit is nucleated. There are many other reports also describing thepenetration of Cl– in the oxide film [69–72]. Overall, the AM 316L SSsamples had a much higher breakdown potential in 0.9M NaCl whencompared to the wrought 316L SS sample.

To further estimate the localized corrosion behavior of the samples,a reverse anodic scan was used. The development of large positive loopduring reverse anodic scan and intersection with the forward scan atthe protection potential (Ep) dictate the passive film breakdown abovewhich the pits may form and continue to grow. In order to quantify thepitting resistance (PR = Ebd – Ecorr) and protection tendency (PT = Ep -Ecorr) of the samples and to compare their localized corrosion tendencyin each electrolyte, a summary of the data is presented in Table 2.Significantly higher PR (> 1000mV) conferred by the AM 316L SSsample indicated high stability of the passive oxide film. However, in

Fig. 7. Histograms showing the variation in (a) Qo and (b) Rct of the wrought and AM 316L SS samples in serum, PBS and 0.9M NaCl.

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0.9M NaCl solution, both the wrought and AM 316L SS samples re-gistered negative PT values, suggesting that once the pre-existing pas-sive oxide film is locally deteriorated, the severe localized corrosionreactions may progress uninterruptedly. The high stability of the pas-sive oxide film formed on the AM 316L SS sample caused a noticeableimprovement in the localized corrosion resistance in natural and si-mulated biological environments and could overcome the longstandingchallenge which restricts the use of stainless steel (wrought) in manybiomedical applications. Importantly the improvement in the Ebd valuesafter grain refinement of the wrought stainless steel by various me-chanical means has already been reported repeatedly in the literature[62,64,65]. S. V. Muley et al. reported the increase in Ebd from 200 to389, 419 and 436mV under the application of effective cumulativestrain (εeff) of 0, 1.4, 2.8 and 4.2 respectively, tested in 3.5 wt. % NaClsolution [62]. However, our study showed the Ebd of 200 and 920mVfor the wrought and AM SS 316 L samples, respectively, when tested in0.9 M NaCl (equivalent to 3.5 wt.% NaCl). The results of this studyshowed that the AM 316L SS samples registered significantly higher Ebd

values, and suggested that the improved stability of the oxide filmformed on AM 316 L stainless steel samples could minimize the possi-bility of 316 L stainless steel localized corrosion in biological environ-ments.

3.4. Cell morphological analysis

In order to investigate the in–vitro biological response of thewrought and AM 316L SS, cell proliferation and morphology of theMC3T3 pre–osteoblast cells were assessed after 24 and 48 h of in-cubation period. Cell proliferation and spreading on the implant surfaceare the initial steps in a cascade of cell-implant interaction and it pre-dicts the comparable in–vivo response of the implant in terms of bio–-functionality. The surface texture and topographical features of theimplant material are the critical factors which may affect the growthtendency of the cells, and consequently describes the implant/tissuesinteraction [73,74]. Cell proliferation after 24 and 48-hours incubationtimes on the wrought and AM 316L SS samples are shown in Fig. 9. On

Fig. 8. Cyclic Polarization plots for the wrought and AM 316L SS samples in (a) serum, (b) PBS and (c) 0.9M NaCl. (d) Histogram representing breakdown potential.

Table 2Quantitative information about the electrochemical properties of wrought and AM 316L SS in various biological solutions as evaluated from the cyclic polarization.

Electrolyte Sample Ecorr (mV vs. SCE) Icorr (μA-cm−2) Ebd (mV) Ep (mV) PR (mV) PT (mV)

Serum Wrought −299 0.475 409 −218.0 708 81AM −331 0.457 984 −138.4 1315 193

PBS Wrought −261 0.599 403 −200.0 664 61AM −229 0.569 1038 −96.2 1267 133

0.9M NaCl Wrought −203 0.269 200 −238.0 403 −35AM −179 0.212 920 −252.3 1099 −73

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both the samples, the spherical and compact morphology of the pre–-osteoblast cells was observed after 24 h of incubation as shown inFig. 9a and b. Nearly twice as many cells were observed on the AM 316LSS sample compared to the cell growth on the wrought sample during24 h of incubation. For example, 77 cells adhered on AM, whereas 45cells adhered on wrought. The higher number of cells on the AM 316LSS sample suggests that the cell attachment on the surface of the spe-cimen is influenced by the grain size of this sample and not to the celladaption or development over time. The development of filopodia wasnot pronounced in the initial 24 h of incubation and that resulted inpoorly extended, concentric geometries, indicating the initial attach-ment of the cells on the surface. However, with increasing time thesurface characteristics starts to influence the cell proliferation and celladhesion. Therefore, we also assessed the cells proliferation after 48 hof incubation period. As shown in Fig. 9c and d, both materials sup-ported cell spreading and the appearance of stiller morphology after48 h incubation, indicating the cellular compatibility of both the sam-ples. Although both samples supported spreading of cytoplasmic ex-tensions, cell numbers remained increased on the surface of AM 316LSS sample (68 cells on the AM 316L SS vs. 39 cells on the wroughtsample) after 48-hour incubation, suggesting the enhanced cell survivaland pronounced osteoblast differentiation. This behavior suggested thebetter growth rates and superior cell attachment characteristics on theAM 316L SS sample.

Overall, this data suggested the influence of grain structure and theresulting wettability characteristic of the sample on the cell prolifera-tion and adhesion, since the chemical composition and surface topo-graphy (polished) of both the samples were very similar. The hydro-philic nature of the surface enhanced pre-osteoblast attachment andspreading. The lower contact angle of the AM 316L SS sample likelycontributes to an increase in surface energy and sites for absorption thatsupports better cell proliferation and adhesion [31]. Furthermore, thefine microstructure also promotes the activity of alkaline phosphatesand osteocalin, and could results in improved osteointegration [75,76].

On the basis of these results, the fine grain structure of the AM 316L SSsample may provided more compatible sites for the cell adhesion andmay have caused faster adhesion and cell proliferation as shown inFig. 9b, d. These findings suggest that faster and more efficient os-teointegration would occur because the faster colonization on the im-plant surface promotes bone tissue interaction [27].

The finest biocompatible materials not only support cell prolifera-tion, but also provide optimum environment for specific genes in cellsthat improve their ability to differentiate into osteoblastic cells [77].Runt-related transcription factor 2 (Runx2) is an anti–proliferative genethat minimizes cell proliferation once pre–osteoblast cells starts term-inal differentiation; therefore, we used this protein as a biomarker forimmature osteoblasts [78]. To determine the effect of microstructure ondifferentiation, expression of Runx2 was assessed by western blot inMC3T3 cells cultured on the wrought and AM 316L SS samples for 72 h(Fig. 10). Consistent with steady cell numbers after 24 h and initiationof osteoblast differentiation, Runx2 was equivalent in cells cultured oneither samples. These data indicated that both materials supporteddifferentiation of osteoblasts. As discussed above, the refined structureof the AM 316L SS sample most likely provided better surface for celladhesion and migration by promoting the deposition of extra cellularmatrix and cells during the differentiation process. Thus, the Runx2expression in an increased number of cells on the AM 316L SS samplecompared to the wrought sample suggests better biocompatibility of theAM 316L SS samples [27].

4. Conclusions

In this study, we investigated the corrosion behavior of selectivelaser melted-additively manufactured (AM) and wrought 316 L stainlesssteel in serum, PBS and 0.9M NaCl. The MC3T3 pre–osteoblast cellsproliferation and differentiation on these steel samples were also de-termined. The EDX and XRD analysis confirmed the metallurgicalpurity and development of austenitic phase in both the samples. The

Fig. 9. Fluorescence micrographs of MC3T3 cell proliferation on the wrought and AM 316L SS samples, stained with Nuc-blue (nucleus) and Actin-red (actinfilaments of cytoskeleton) (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).

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microstructural analysis of the wrought sample revealed the presence oftypical polygonal shape coarse austenitic grains, whereas, fine inter-connected sub–grain colonies were formed in the AM 316L SS sample.The contact angle measurement indicated the enhanced hydrophiliccharacteristic of the AM 316L SS sample. The XPS analysis demon-strated the presence of mixed Fe and Cr oxide/hydroxide layers in thepassive film in which the top surface layer was composed of Fe–oxide/hydroxide enriched phase on both the samples. The charge transferresistance (Rct) of the AM 316L SS sample was found to be higher thanthe wrought SS sample in all the electrolytes. The highest Rct registeredby the AM sample (2.23 MΩ–cm2) in serum was attributed to the spe-cific adsorption of proteins on the dense oxide film. The relatively lowerRct and higher Qo of the wrought and AM 316L SS samples in 0.9 MNaCl than in serum and PBS indicated the generation of defects in theoxide film by the Cl– species. However, the overall better barriercharacteristics of the AM 316L SS relative to the wrought sample cor-responded to the dense oxide film structure. The oxide film formed onthe AM 316L SS sample represented improved stability compared towrought sample and is confirmed form the CPP scans. The larger Ebd ˜1000mV vs. SCE and relatively high PR (pitting resistance) of the AM316L SS sample depicted the barrier characteristics of the oxide filmpossibly due to its dense structure and to the absence of MnS inclusions.In 0.9 M NaCl, the larger positive loop in the CPP scan and negative Ep(protection potential) divulged that once the passive oxide film is dis-rupted, the localized corrosion of both the wrought and AM 316L SSsamples would proceed in an uncontrolled manner. The dense coloniesof star shape osteoblasts cells and their improved adhesion on the AM316L SS sample indicated better proliferation or attachment of osteo-blast cells on the AM 316L SS, compared to wrought counterpart. Theenhanced biocompatibility of the AM 316L SS sample was also con-firmed through expression of osteogenic biomarker Runx2 in the cellscultured on both materials. Since twice as many cells undergo differ-entiation on the AM 316L SS compared to wrought sample, the AM316L SS would be preferred over wrought stainless steel for biomedicalapplications. Overall, the high corrosion resistance, and improvedability to support cell adherence and differentiation, indicates that theAM 316L SS eliminates the known limitations of wrought stainless steelfor biomedical usage.

Acknowledgements

The Authors would like to acknowledge Zia ur Rahman and Dr.Naveen Mekala for their help in doing cell proliferation and westernblots, respectively.

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