CoCrMoC coatings on implant grade 316LVM stainless steel

Journal Pre-proof Tribocorrosion response of duplex layered CoCrMoC/CrN and CrN/CoCrMoC coatings on implant grade 316LVM stainless steel

Raisa Chetcuti, Peter A. Dearnley, Antonino Mazzonello, Joseph Buhagiar, Bertram Mallia PII:

S0257-8972(19)31303-9

DOI:

https://doi.org/10.1016/j.surfcoat.2019.125313

Reference:

SCT 125313

To appear in:

Surface & Coatings Technology

Received date:

11 September 2019

Revised date:

2 December 2019

Accepted date:

25 December 2019

Please cite this article as: R. Chetcuti, P.A. Dearnley, A. Mazzonello, et al., Tribocorrosion response of duplex layered CoCrMoC/CrN and CrN/CoCrMoC coatings on implant grade 316LVM stainless steel, Surface & Coatings Technology (2019), https://doi.org/10.1016/ j.surfcoat.2019.125313

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© 2019 Published by Elsevier.

Journal Pre-proof Tribocorrosion response of duplex layered CoCrMoC/CrN and CrN/CoCrMoC coatings on implant grade 316LVM stainless steel Raisa Chetcuti1, Peter A. Dearnley2, Antonino Mazzonello1, Joseph Buhagiar1, Bertram

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Mallia1*

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Corresponding author. Tel.: +356 2340 2057. E-mail: [email protected] (Bertram Mallia)

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Department of Metallurgy and Materials Engineering, University of Malta, Msida, MSD 2080,

Malta 2 †

Boride Services Ltd, Blakedown DY10 3NE, United Kingdom Email: [email protected]

Abstract: The objective of this work was to improve the tribocorrosion performance of a biomedical grade 316LVM stainless steel via the application of duplex layered coatings. Two

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different PVD coatings were deposited via magnetron sputtering: (1) an underlying CoCrMo carbon S-phase layer followed by an outermost CrN layer designated as CrN/S and; (2) an

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underlying CrN layer followed by an outermost CoCrMo carbon S-phase designated as S/CrN.

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Tribocorrosion of the coated and uncoated 316LVM was studied against an alumina ball which

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was made to slide in a reciprocation motion in Ringer’s solution. The contributions of wear, corrosion and their synergy were elucidated by performing tests under open circuit, cathodic and

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anodic potential conditions. The resultant wear scars were analysed via SEM, EDS, optical

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microscopy and contact profilometry. Both coatings displayed superior tribocorrosion response to

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AISI 316LVM with CrN/S exhibiting the best performance. CrN/S displayed lowest coefficient of friction under all electrochemical conditions (~0.10-0.15) and a material loss under anodic condition with a 93% decrease when compared to 316LVM. The CrN/S scars exhibited a smooth morphology under all test conditions in contrast with the other two test materials which displayed shearing marks oriented along the sliding direction under OCP and anodic conditions. By considering the resultant scar morphologies, a hypothesis is being proposed to explain the tribocorrosion degradation mechanisms encountered.

Keywords: 316LVM; CrN; magnetron sputtering; S-phase; tribocorrosion; synergism

2

Journal Pre-proof 1. INTRODUCTION Joint disease is amongst the five most frequent illnesses in the USA which induce long term intense pain to the patient [1]. The high demand for orthopaedic implants and the associated financial burden imposed on healthcare systems necessitate the long durability and functionality of the implant. In hip replacements, austenitic stainless steels articulating against ultra-high molecular weight polyethylene such as the Exeter V40 compatible stainless steel heads are used

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[2, 3]. One cause of implant degradation is tribocorrosion which encapsulates the combined effect

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of mechanical wear between articulating components of an implant and the chemical attack

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governed by body fluids [4-6].

Austenitic stainless steel is characterised with a nanometric self-healing oxide film which

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naturally forms on its surface and protects the bulk material from corrosion [7, 8]. This oxide film

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is mechanically worn by the rubbing action of the femoral head against the acetabular component. The tailoring of implant surfaces using various techniques, including Physical

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Vapour Deposition (PVD), are used to mitigate this damage [9]. PVD allows the surface

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properties to be enhanced without compromising the mechanical characteristics of the substrate

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material. The versatility of PVD enables the synthesis of coatings with different chemistries and structures [10, 11]. Most of the reported studies have been oriented towards mono-layer PVD coatings. The full potential of PVD that enables the synthesis of different coatings architectures with possible enhanced performance to mono-layer coatings has yet to be explored [12, 13]. S-phase, also known as expanded austenite, is a precipitate free, supersaturated, metastable solid solution of nitrogen or carbon in a face centred cubic chromium-containing alloy. S-phase layers have generally been produced on austenitic stainless steel and CoCrMo alloys via low temperature diffusion treatments such as carburising and nitriding [9, 14, 15]. In this work, PVD is used to produce CoCrMo supersaturated with carbon atoms (S-phase) at temperatures far 3

Journal Pre-proof below the sensitisation limit [15, 16]. Supersaturated metallic systems were previously shown to undergo significant oxidative wear (Type I corrosion-wear) damage but display higher load support than the base alloy [11, 17]. PVD sputtered CoCrMo(C) S-phase has only received limited investigation for biomedical implants since previous studies [9, 17, 18] have focused mostly on the Fe-Cr-Ni S-phase produced on metallic alloys via thermochemical treatments. Previous works [9, 11, 12] have shown that CrN coatings displayed excellent resistance to

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cyclic passive film breakdown and regeneration (Type I corrosion-wear) resulting in mitigation

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of material losses via this mechanism. CrN coated metallic biomaterials were however

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susceptible to accelerated corrosion attack of the substrate or coating-substrate interface resulting in blistering damage on stainless steel and catastrophic failure of the coating (Type II corrosion-

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wear) [9, 11].

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In a recent investigation of CoCrMo metal-on-metal PVD TiNbN coated bearings [19], which were retrieved 53 months after implantation, it was determined that the coating exhibited third

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body wear and delamination. Failure at the coating-substrate interface might have been induced

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by a mild form of Type II corrosion-wear damage [9, 19]. A strategically designed dual layer

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coating to mitigate the presence formation of paths for the solution to reach the substrate-coating interface can alleviate Type II corrosion-wear while maintaining a high resistance to Type I corrosion-wear.

In this work, CrN on CoCrMo(C) S-phase were deposited via PVD to form a CrN/S dual layer on biomedical grade stainless steel. The CoCrMo(C) S-phase is intended to provide better load support to the top CrN layer with excellent resistance to Type I corrosion-wear damage. This underlayer shall mitigate micro-cracking of the CrN layer; a potential precursor for coatingsubstrate interfacial corrosion that causes Type II corrosion-wear damage. This dual-layered coating helps to address the current concern of metallic ion release into the blood stream and the 4

Journal Pre-proof generation of wear particles at the bearing surfaces of the metallic implants [20]. The tribocorrosion of the CoCrMo(C) S-phase layer and its mechanical properties were also studied on specimens by depositing the layers in reverse order. The study of the latter was included to get better scientific insight on the potential degradation mechanisms and enable a comprehensive study of the tribocorrosion performance of both coating layer combinations.

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2. MATERIALS AND METHODS

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2.1 Material Preparation

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The substrate material used was AISI 316LVM austenitic stainless steel alloy (L. Klein SA, Switzerland) in conformity with the standard ASTM F138 and with the nominal composition

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given in Table 1. Disc specimens of 25 mm diameter and 6 mm thickness were cut using a SiC

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grinding disc from an annealed rolled bar. Parallel, flat surfaces were obtained using a JonesShipman 540 precision grinder (UK) followed by fine grinding with an emery cloth of 1200 µm

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grit size. The disc specimens were subsequently polished to a mirror finish (Ra: 3.5 ± 0.4 nm)

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using polycrystalline diamond polishing with 9 µm, 6 µm and 3 µm diamond pastes on Struers

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LaboPol 25 equipment (Denmark). Prior to coating deposition, the disc specimens were ultrasonically cleaned in ethanol for fifteen minutes. The surface roughness of the test materials was measured using a NanoMap-500LS 2D contact profilometer (USA) equipped with a stylus having a tip radius of 1 μm. A scan distance of 1000 µm, a scan speed of 25 μms -1, a contact load of 15 mg and a sample frequency of 20 Hz were used. The coating deposition was carried out using a custom built laboratory scale unbalanced magnetron sputtering setup at Boride Services Ltd. (UK). Radio frequency (r.f.) sputter cleaning process of the 316LVM substrate prior to deposition was carried out in an Argon atmosphere at a pressure of 0.12 Pa and an induced bias of -50 V. A thin layer of CoCrMo was initially sputter 5

Journal Pre-proof deposited on the 316LVM substrate followed by a ~8 µm CoCrMo(C) S-phase layer using a CoCrMo alloy target. This designates the underlayer of the CrN/S coating system as shown in the schematic in Figure 1(a). The top layer constituted of a 2.3 µm thick CrN layer sputter deposited using a Cr target in a reactive atmosphere following the deposition of a thin Cr adhesion layer and with a surface roughness, Ra, of 9.1 ± 0.3 nm. The deposition parameters are summarised in Table 2. The S/CrN coating system holds the same architecture as CrN/S, but in reverse order. In

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this case, Cr interlayers were deposited at both substrate/underlayer interface and coating/coating

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interface as indicated in Figure 1(b). The surface roughness of the S/CrN coating system was

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measured to be 6.9 ± 0.5 nm. Although it was expected that the CrN/S would have superior tribocorrosion performance to S/CrN, the study of the latter was included to get better scientific

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insight on the potential degradation mechanisms.

2.2 Chemical composition and phase identification

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The crystalline phases of the uncoated and coated specimens were determined by Glancing

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Angle X-Ray Diffraction (GXRD) using a Rigaku Ultima IV (Texas, USA) equipped with a

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copper tube. In this study, the measured reflections for CrN/S were indexed using powder diffraction files (PDFs) from the International Centre of Diffraction Data (ICDD), while reflections for S/CrN, were manually indexed using an analytical technique [21]. In this approach, the angle for each reflection is divided by an integer, i.e. (sin2θ/2, sin2θ/3, sin2θ/4, etc...) and the first quotient, common for both reflections, K is determined. The angle for each reflection is divided by the quotient, K, which represents λ2/4a2. A list of integers corresponding to h2 +k2+l2 were obtained, and the structure of the surface material was identified to be FCC by identification of the Bravais lattice. The thickness of each coating layer in the duplex system was determined from cross-sectional analysis of fractured coated specimens using Zeiss Merlin Gemini (Germany) Scanning Electron Microscopy (SEM). Coating composition was qualitatively 6

Journal Pre-proof determined using an Apollo X Ametek (USA) Energy Dispersive Spectrometer (EDS) attached to the SEM.

2.3 Mechanical properties A Micromaterials NanoTest 600 (UK) Nanotester equipped with a Berkovich diamond indenter of a 150 nm tip radius was used to measure the surface hardness and reduced modulus of

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the uncoated 316 LVM and coated samples. Thirty indents were performed on each sample at

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24 °C, each at an indentation load of 30 mN and a loading/unloading rate of 0.5 mN/s. The

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Nanotester equipped with a Synton-MDP 600 conical diamond indenter of tip radius 10 µm was

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used to assess the deformation behaviour of U/T 316 LVM and coated samples. A constant load of 0.8 mN was applied for the initial topographic scan (pre-load scan), followed by a ramped

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scratch scan (load scan). During the latter scan, the load was held constant at 0.8 mN for the first

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50 µm but then gradually increased at a constant rate until a maximum load of 450 mN was reached at 350 µm from the beginning of the scratch track. A final topographic scan (post-load

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scan) was subsequently carried out at a constant load of 0.8 mN to determine the elastic recovery

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of the material. Each test was repeated for five times. 2.4 Corrosion and Tribocorrosion) testing Potentiodynamic polarisation tests were carried out in Ringer’s solution using a Gamry Reference 600TM potentiostat (USA) to study the corrosion behaviour of the test samples. The potentiostat was connected to a three-electrode cell consisting of the sample (working electrode), a saturated calomel reference electrode (SCE) as the reference electrode and a platinum-coated titanium rod acting as the counter-electrode. The scanning sequence involved 1 hour of open circuit potential for potential stabilisation followed by a potentiodynamic sweep at a scan rate of 0.167 mVs-1 between -100 mV versus OCP and 1000 mV versus SCE. 7

Journal Pre-proof The tribocorrosion behaviour of U/T 316LVM and the coated samples was investigated with a computer controlled Gamry Reference 1000TM potentiostat (USA) connected to a threeelectrode cell, similar to that used for static corrosion testing. All tests were conducted in Ringer’s solution held at 37 ± 1°C by a double walled glass heated water jacket. Tribocorrosion tests were performed using a custom-built reciprocating sliding tribometer connected to a potentiostat. The counter face material was a sintered alumina ball with a 7.9 mm diameter ball

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bearing (Spheric Trafalgar, UK) having a hardness of 16.7 GPa and a Young’s modulus of 365

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GPa. The use of an alumina ball enables better understanding of the tribocorrosion interactions on

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the test material since the recorded current would only be generated from the test materials. Prior to each test, the test sample was ultrasonically cleaned in acetone at 50°C for fifteen minutes and

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the tribocorrosion cell was flushed with deionised water. The test specimen was then mounted in

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the tribometer and Ringer’s solution was poured in the cell. Reciprocating sliding tests were conducted at a frequency of 2 Hz under a normal load of 1 N and a 7 mm stroke length. The

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normal load was chosen such that the resulting uniaxial and subsurface shear stresses were close

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to their respective yield strengths to avoid gross plastic deformation of the test samples [22].

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Specifically, the maximum Hertz contact pressure (Pmax) on uncoated 316LM was 0.61 GPa whilst the maximum subsurface shear stress (τmax) was 0.19 GPa, where the calculated Hertzian elastic contact radius (a) was 28 μm. The latter stresses compare to the uniaxial yield strength and shear yield strengths of 0.27 GPa and 0.15 GPa respectively for uncoated 316LVM. All values were calculated mathematically using the Hertzian elastic contact theory equations for a ball-on-flat contact. For the uncoated 316LVM, at the onset of sliding, the maximum shear stress (0.19 GPa) was slightly higher than its shear yield strength (0.15 GPa). However, the stresses are expected to become immediately lower as the contact area increases following the onset of

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Journal Pre-proof sliding. Using a load which is close to the shear yield strength of the substrate enables getting measurable wear in a reasonable testing time. Tribocorrosion tests were conducted under open circuit potential (OCP) conditions, cathodic potential (CP) conditions of -0.7 versus SCE and anodic potential (AP) conditions of +0.1 versus SCE. Under CP and AP, the current was recorded for ten minutes without sliding, followed by two hours of sliding and a further ten minutes without sliding. The same procedure was followed

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when testing under OCP conditions, during which the OCP potential was recorded. During

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tribocorrosion experiments, the electrochemical parameters (current and potential) and frictional

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forces were continuously monitored via a Gamry Reference 1000TM potentiostat (USA) connected to a PC running a dedicated LabviewTM program. All tribocorrosion tests were

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triplicated.

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The profiles across the scars on the uncoated 316LVM and coated samples following tribocorrosion testing were measured with a NanoMap-500LS (USA) contact profilometer

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equipped with a stylus having a tip radius of 1 μm. The scans were carried out in the middle of

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the wear scars as well as at around 2 mm away from each of the ends of the scar. The volumetric

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material loss was calculated by multiplying the average area of the corrosion-wear scar (based on three scars) by the scar length. The surface morphology and coating composition of the corrosion-wear tracks were assessed using SEM (Zeiss Merlin GeminiTM, Germany) equipped with an EDS (Apollo X Ametek, USA), while a Leica MicrosystemsTM light optical microscope (Heerbrugg, Switzerland) was used to view the scar on the alumina counter body material after tribocorrosion testing. A method developed by Diomidis et al. [23] to calculate the synergistic component (S) of the total volumetric material loss (TML) under OCP and AP conditions was used. This takes into

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Journal Pre-proof account the synergism (S) of wear aggravated corrosion, CW, i.e. primarily Type I corrosion-wear, and corrosion aggravated wear, WC. That is:

𝑇𝑀𝐿𝐴𝑃/𝑂𝐶𝑃 = 𝑊0 + 𝐶0 + 𝑆

(Equation 1)

Where W0 is the total volume of material lost due to mechanical wear only (obtained from potentiostatic tests carried out under CP) and C0 is the total volume of material lost due to

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corrosion only, which is considered to be practically zero for passive materials. And:

(Equation 2)

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𝑆 = 𝐶𝑊 + 𝑊𝐶

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Specifically considering OCP and AP conditions, Equation 1 can be rearranged to obtain S: (Equation 3)

𝑆𝐴𝑃 = 𝑊𝐶𝐴𝑃 + 𝐶𝑊𝐴𝑃 = 𝑇𝑀𝐿𝐴𝑃 − 𝑇𝑀𝐿𝐶𝑃

(Equation 4)

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𝑆𝑂𝐶𝑃 = 𝑊𝐶𝑂𝐶𝑃 + 𝐶𝑊𝑂𝐶𝑃 = 𝑇𝑀𝐿𝑂𝐶𝑃 − 𝑇𝑀𝐿𝐶𝑃

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Here, the volumetric quantities TMLOCP , TMLAP and TMLCP are directly obtained from the

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cross-sectional area of the tribocorrosion tracks determined from contacting profilometry

3. RESULTS

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(produced under OCP, AP and CP test conditions), multiplied by the track length.

3.1 Coating crystal structure The GXRD plots, shown in Figure 2, only gave structural information of the outermost coating layer; insufficient X-ray penetration took place to allow detection of the adjacent sublayers or the 316LVM substrate. For the CrN/S coated 316LVM, characteristic CrN (111), (200), (220), (311) and (222) reflections were detected (Figure 2a) in agreement with the ICDD card number 04-015-0334 (cubic CrN with a = 0.41137 nm) whilst for S/CrN coated 316LVM two broad reflections at ~44° and ~80° of 2θ (Cu Kα), respectively corresponding to the (111) and 10

Journal Pre-proof (220) inter-planer spacings of the cubic lattice of CoCrMo matrix carbon S-phase [24], were detected (Figure 2b). The intensity of the S-phase (111) reflection was about 10 times higher than the peak attributed to the (220) inter-planar spacing (Figure 2b). 3.2 Coating composition, hardness and modulus Chemical composition of the top (exterior) layers of the duplex coatings are shown in Table 3. The exterior CrN coating of the CrN/S coated 316L material was close to the theoretical

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stoichiometric maximum of 50Cr:50N, being 48 at.% N and 52 at.% Cr. On the other hand, the

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S-phase solid solution top layer of the S/CrN coated 316LVM contained 16 at.% C which is

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similar to reports of carbon S-phase contents produced by low temperature carburising of CoCrMo alloys [25].

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The nano-indentation hardness of the duplex CrN/S and S/CrN duplex coated 316LVM steels

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was 21.6 ± 1.3 and 12.7 ± 0.3 GPa, respectively, compared to 3.6 ± 0.2 GPa for the uncoated substrate (Table 3). The values for the duplex coatings are close to those reported for single

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layered PVD-CrN and PVD S-phase [12, 13] coatings, indicating that the nano-hardness values

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were not significantly affected by the different hardness values of their adjacent (neighbouring)

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sub-layers. The nano-indentation reduced elastic moduli values for the duplex CrN/S and S/CrN coatings 316LVM were very similar, being 221 ± 5 and 209 ± 4 GPa respectively; both, however, were higher than that of uncoated 316LVM (179 ± 10 GPa).

3.3 Scratch testing These tests revealed the surface plastic deformation and fracture responses of the coated and uncoated materials. No delamination of the coated materials took place during these tests, indicating excellent coating adherence to the substrates. The uncoated 316LVM exhibited a roughened scar due to ploughing of material ahead of the indenter (Figure 3a) as supported 11

Journal Pre-proof further by the post load surface profile scan (Figure 4a). Evidence of plastic flow and plastic fracture of the external S-phase coating of the S/CrN coated sample was evident from flattening of the original hemi-spherical coating asperities within the scratch track and slight coating material detachment along the sides of the scratch (Figure 3b). Further, ripples in the final region of the load surface profile scan (Figure 4b) were revealed that indicated there had been a material build-up in front of the scratch indenter just before scratching was stopped. A different response

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was shown by the CrN/S coated 316LVM. Here, the initial hemi-spherical coating topography of

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the CrN/S sample surface was only changed slightly within the scratch track (Figure 3c) – the

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asperities have been made smooth – indicating superficial plastic flow, without any fracture taking place. It was also apparent from the ‘load’ and ‘post-load’ profilometer traces (Figure 4c)

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that only slight plastic flow of the adjacent substrate had taken place during scratching. This

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demonstrated that the CrN/S coated 316LM material was very resistant to mechanical damage

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under the selected nano-scratch test conditions.

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3.4 Corrosion and Tribocorrosion tests

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3.4.1 Potentiodynamic tests

The potentiodynamic polarisation curves for the test materials are shown in Figure 5. The uncoated 316LVM material displayed a wide passive region and low passive current density. Fluctuations in current between 0.18 – 0.83 V versus SCE show the occurrence of metastable pitting events. The S/CrN and CrN/S coatings exhibited a similar passive current density to that of uncoated 316LVM. The anodic current of the coated samples started to increase rapidly at ~0.6 V versus SCE, signifying the onset of transpassive dissolution. No pitting or crevice corrosion on the coated and uncoated test areas were observed after potentiodynamic testing. Following testing, of the coated specimens it was observed that the appearance of the test surfaces changed from shiny to a matt finish as a result of transpassive dissolution of chromium. This change in 12

Journal Pre-proof test area appearance was not evident for the uncoated 316LVM which exhibited a low current density till the end of the test. CrN/S had a passive current density which was slightly higher (smaller than ~0.13 µAcm -2) than that of the uncoated 316LVM over the entire passive potential range. This result can be explained by comparing the surface topographies of CrN/S coated and uncoated 316LVM. In the case of CrN/S the topography comprises dome-tipped asperities and had an Ra of 9.1 nm while

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the uncoated 316LVM had a featureless topography with an Ra of 3.5 nm. As a result, CrN/S is

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expected to have a higher surface area for the same projected area under test compared to the

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uncoated 316LVM contributing to a higher measured current.

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3.4.2 Potential and coefficient of friction (COF) variation under open circuit potential (OCP) conditions

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The evolution of the open circuit potential (OCP) and the coefficient of friction (COF) during

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tribocorrosion testing is shown in Figure 6. During the first few hundred seconds of exposure to

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the saline electrolyte, the OCP of the coated and uncoated test pieces stabilised at ~ -50 mV vs. SCE, indicating the onset of passivation via the formation of a protective passive oxide layer. On the onset of sliding against the Al2O3 counterface, a sudden drop in OCP of ~ -300 mV vs. SCE was observed for the S/CrN coated and uncoated 316LVM samples (Figure 6a); in contrast only a slight initial fall of about 10 mV vs. SCE was observed for the CrN/S coated 316LVM. Significant oscillations of several tens of mV vs. SCE were noted during sliding contact of the alumina balls on the S/CrN coated and uncoated 316LVM test-pieces, whereas, no similar fluctuations were observed for the CrN/S coated 316LVM test piece (Figure 6a). Here, only a slow alteration in OCP took place throughout the entire test, dropping from an initial value of ~ -50 mV vs. SCE to ~ -100 mV by the end of the test. The fluctuations in OCP (during sliding 13

Journal Pre-proof contact) observed for the S/CrN coated and uncoated 316LVM test-pieces are similar to those observed previously by other workers and can be attributed to the cyclic removal and regeneration of the passive film formed on metal alloy surfaces [9, 13, 26, 27]; this leads to material loss via Type I corrosion-wear [9]. Beyond the initial ~4000 s of testing, the amplitude of fluctuation diminished. Once sliding against the Al2O3 counterface was stopped after 7800s, the OCP returned towards its initial value before testing.

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Figure 6a also shows that the dynamic OCP is largely unchanged over sliding time for the

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uncoated 316LVM. On the other hand, the dynamic OCP shifts gradually to more negative values

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with time of sliding for S/CrN until it stabilised after circa 5000s from test start. This indicates that the uncoated 316LVM reached a steady-state condition before S/CrN. An increased overall

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damage to the passive film is expected to shift lower the open circuit potential to balance the

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anodic and cathodic reactions. A longer time in reaching a steady-state is also reflected in the dynamic friction coefficient which increased at a steeper rate with time of sliding for S/CrN albeit

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exhibiting a lower steady-state value (~0.4) than that of uncoated 316LVM (~0.6) as shown in

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Figure 6b. The behaviour of CrN/S coated 316LVM was profoundly different. Here, the COF

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was stable and relatively low (~0.15) throughout the test duration (Figure 6b). The evolution of the OCP during reciprocating tests was considered in determining the cathodic and anodic potentials used for potentiostatic tribocorrosion testing. A cathodic potential of -700 mV versus SCE was selected by shifting the lowest dynamic OCP recorded, which in this study was exhibited by S/CrN (~450 mV) by 250 mV in the cathodic direction. The same cathodic potential was used for all test materials. No hydrogen evolution was observed during these tests. The anodic potential was chosen by consulting Figures 5 and 6a . The potential was selected to be 150 mV higher than that of the highest static OCP recorded of -50mV for CrN/S. This resulted in 14

Journal Pre-proof a static potential of +100 mV versus SCE which was within the passive zone of all test materials (Figure 5) 3.4.2 Current and coefficient of friction variation (COF) under cathodic potential (CP) conditions The application of an external cathodic potential in principle suppresses dissolution of the surface metallic ions (via the anodic reaction). Under such conditions the mechanical component

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of tribocorrosion dominates material loss since anodic ion dissolution is suppressed [24].

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Tribocorrosion tests performed with an externally applied cathodic voltage of -700 mV vs. SCE

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revealed the variations in cathodic current and COF, given in Figure 7. The uncoated 316LVM experienced a rapid (in under 10 seconds) decrease in cathodic current magnitude from an initial

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value of ~ -100 μA to a value of ~ -600 μA when sliding contact with the Al2O3 ball was started.

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Similar behaviour took place for the coated test-pieces although the magnitude of the cathodic current fall was lower. In all cases the cathodic current value returned to near its original value

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once sliding was stopped (Figure 7a). Regarding the changes in COF (Figure 7b), relatively high

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values were noted for the S/CrN coated 316LVM (> 0.3) and uncoated 316LVM (~0.35-0.45),

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whilst the CrN/S coated 316LVM displayed a very low value of ~0.15 throughout the test.

3.4.3 Corrosion current and coefficient of friction (COF) variation under anodic potential (AP) conditions The application of an external anodic potential accelerates the dissolution of surface metallic ions [28] and a more rapid formation of a protective passive oxide film; this leads to a more rapid material loss via Type I corrosion-wear, compared to that which takes place under OCP conditions [9]. Further, testing under AP conditions, allows the monitoring of the resulting anodic current during sliding contact, as shown in Figure 8a. 15

Journal Pre-proof The overall currents were relatively high for both the S/CrN coated 316LVM (~15 to 40 μA) and uncoated test pieces (~25 to 70 μA). In contrast, for the CrN/S coated 316LVM the anodic current was much lower (< 5 μA). Large fluctuations (several tens of micro-amps) in the dynamic current were recorded for the uncoated and S/CrN coated 316LVM test-pieces during sliding contact with the alumina counterface material, whilst negligible fluctuations in this parameter were noted for the CrN/S coated 316LVM (Figure 8a). Here, the corrosion current was relatively

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smooth with respect to time. The fluctuating current for the uncoated and S/CrN coated stainless

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steel can be attributed to the cyclic mechanical removal and reformation of the passive film/oxide

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scale during tribocorrosion testing. For all test materials, the cessation of sliding contact with the Al2O3 test ball, caused the anodic current to decrease to values slightly in excess of zero as shown

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after sliding contact was stopped.

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in Figure 8a; the latter effect indicates minor electrochemical activity continued to take place

The recorded COF (Figure 8b), showed high values for the uncoated (~0.4) and S/CrN coated

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316LVM (~0.25) compared to values of ~0.1 recorded for the CrN/S coated 316LVM. The COF

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values for the S/CrN [17, 29] and CrN/S coated 316LVM [12, 17, 26, 29] are in agreement with

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work published by others.

3.4.4 Morphology of the corrosion-wear tracks SEM images of representative wear scars for all the tribocorrosion tested samples under OCP, CP and AP conditions are shown in Figure 9, while light optical micrographs of the corresponding worn Al2O3 surface are shown in Figure 10. Specifically, the tribocorrosion tracks appeared as follows:Under OCP conditions both S/CrN and uncoated 316LVM displayed plastic shearing marks oriented in the direction of reciprocation contact with the Al2O3 ball counter-face. This was far 16

Journal Pre-proof more pronounced for the uncoated 316LVM test-piece (Figure 9a) than the S/CrN coated 316LVM (Figure 9b). In contrast the CrN/S coated 316LVM was comparatively smooth (Figure 9c) indicating an absence of plastic shearing. Some wear debris and/or oxide scale formation was noted at the sides of the corrosion-wear tracks produced on the S/CrN coated and uncoated 316LVM test-pieces whilst no similar deposits were seen on the tested CrN/S coated 316LVM. No micro-cracking or coating blistering of any of the coated materials was observed.

f

Under CP conditions all test materials displayed smooth surfaces (Figure 9d, e, f) with minor

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occurrence of plastic shearing in the direction of sliding only for uncoated (Figure 9d) and S/CrN

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coated 316LVM (Figure 9e). No wear debris and/or oxide scale formation was noted at the sides

micro-cracking of the coatings.

e-

of the corrosion-wear tracks produced under these conditions and there was no evidence of

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Under AP conditions the surfaces of the corrosion-wear tracks resembled those produced under OCP conditions. Hence, both uncoated 316LVM (Figure 9g) and S/CrN coated 316LVM

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(Figure 9h) displayed plastic shearing marks oriented parallel with the direction of reciprocation

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sliding contact whilst the CrN/S coated 316LVM track was very smooth (Figure 9i).

No

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evidence of micro-cracking of the coatings was observed in the tracks. Wear debris and/or oxide scale formation at the sides and the ends of the test tracks of the uncoated 316LVM (Figure 9g) and S/CrN (Figure 9h) coated 316LVM was more extensive than that observed under OCP conditions (Figure 9a, b). By contrast, no such similar debris/deposit formation was noted on the CrN/S coated 316LVM (Figure 9i), which was similar to that observed under OCP conditions (Figure 9c). Both CrN/S and S/CrN coated 316LVM were resistant to material loss via microcracking and/or coating blistering (Type II corrosion-wear).

3.4.5 Morphology of the alumina (Al2O3) ball counter-face contact areas 17

Journal Pre-proof The alumina (Al2O3) ball counter-face contact areas, after sliding against most coated and uncoated 316LVM test samples under OCP, CP and AP test conditions, retained their original spherical shape, suffering no significant damage (Figure 10). The elliptical shape of the contact zone on the alumina balls (Figure 10), shows that the ball is conforming with the scar on the disc specimen rather than exhibiting wear. The absence of abrasive wear of the alumina ball was expected when sliding against the much softer uncoated 316LVM. Sliding against CrN/S which

f

had a somewhat higher hardness than the alumina ball still did not result in observable wear

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damage on the ball. Furthermore, no evidence of alumina grain pull-out was observed in any of

pr

the conducted experiments. Varying degrees of coverage of the ball contact areas with debris transferred from the coated and uncoated test-plate tracks was visible. There was no clear

e-

correlation between the amount of such coverage and the type of coated or uncoated test-plates

Pr

tested against the alumina balls nor the electrochemical conditions that had been deployed in each

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3.4.6 Material losses

al

specific case.

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Total volumetric material loss data for all the coated and uncoated test materials, after tribocorrosion sliding contact tests against Al2O3 balls under OCP, CP and AP conditions are summarised in Figure 11. Under all electrochemical conditions the TML was highest for the uncoated 316LVM test-plates, whilst the lowest TML recorded was for the CrN/S coated 316LVM. Whilst less good than the latter, the S/CrN coated 316LVM was consistently superior to uncoated 316LVM. The use of AP conditions (Figure 11) accelerated the TML rate for all the coated and uncoated 316LVM test plates compared to that produced under OCP conditions, whilst conversely, the application of CP conditions caused the TML to be markedly reduced; here the TML was very low (< 1x10-12 m3) for all test-plates. Figure 12 illustrates the volumetric total 18

Journal Pre-proof material losses due to the synergy components (WC + CW) under AP and OCP conditions, determined via Equations 3 and 4. Using OCP conditions the percentage material loss due to synergism (S OCP) was ~82% of the TML for uncoated 316LVM; ~72% of the TML for S/CrN coated 316LVM and ~59% of the TML for CrN/S coated 316LVM. Under AP conditions the percentage material loss due to synergism (SAP) was ~92% of the TML for uncoated 316LVM; ~95% of the TML for S/CrN

f

coated 316LVM and ~72% of the TML for CrN/S coated 316LVM. Hence, the effect of applying

oo

a positive potential during tribo-corrosion testing was to increase the %S (WC + CW) of all the

pr

coated and uncoated test-pieces compared that produced under OCP conditions.

e-

4. DISCUSSION

Pr

Direct evidence of corrosion activity during tribo-corrosion testing was three-fold: (1) The fall in potential under OCP conditions (Figure 6a); (2) The accumulation of wear debris at the

al

sides and ends of some of the tribocorrosion tracks (Figure 7a, b and Figure 7g, h) and; (3) the

rn

detection of current under AP (Figure 8a) testing conditions. Prior EDS analysis of similar wear

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debris formed on uncoated and S-phase coated 316LVM, under similar test conditions, has shown these to contain >20 at.% oxygen [11, 17]. This probably points to the formation of fine oxidised debris which have accumulated at the edges of the wear tracks over the cycles. Such deposits, comprising of very fine particles, only formed after testing the exterior carbon S-phase coating (of the duplex S/CrN) and uncoated 316LVM stainless steel; it was not detected on the exterior CrN coating (of the duplex CrN/S) after testing (Figure 9c, f, i). The latter observation is in keeping with the hypothesis that corrosion activity was very low for the CrN/S, which was clear under AP conditions when a very low corrosion (oxidation) current was recorded compared to uncoated 316LVM and S/CrN (Figure 8a). 19

Journal Pre-proof Evidence of mechanical wear during tribo-corrosion testing comprised the formation of plastic shearing marks oriented in the direction of reciprocation contact with the Al2O3 ball counter-face. This was only observed under AP and OCP test conditions and was more pronounced for uncoated 316LVM (Figure 9a, g) compared to S/CrN (Figure 9b, h). In contrast, the exterior CrN coating (of the duplex CrN/S) appeared smooth (Figure 9c, i). Under CP conditions, no evidence of plastic shearing was observed after tribocorrosion testing uncoated

f

316LVM and S/CrN coated 316LVM (Figure 9d, e). A hypothesis for this observation will be

oo

proposed in this discussion.

pr

For the test conditions used in this work, what appears to be simple micro-asperity shearing [30] is in fact corrosion aggravated wear (WC), i.e., mechanical wear that requires prior corrosion

e-

to take place. This mechanism, depicted schematically in Figure 13 requires that the test surface

Pr

exhibits uneven material loss leading to roughening and associated increased asperity contact pressure. Conducting passive test surfaces immersed in Ringer’s solution under OCP conditions

al

will become activated in the contact region during sliding. The active region becomes polarised

rn

with respect to the undisturbed region outside the activated zone and as a consequence the OCP

Jo u

shifts to lower values (Figure 6a). Initial contact with the alumina counterface will happen at the surface micro-asperities (Figure 13a). Shear stress exerted by the Al2O3 ball counterface, plastically deforms the topographical ‘hills’ of metallic materials in the reciprocating directions of sliding contact provided that their shear yield strength (τy) is exceeded. This localised shearing of the surface will result in deformed asperities, plastic fracture leading to debris generation, damage to the passive film and an increased dislocation density in the material [31, 32]. The electrochemical reactivity of passive metals will depend on the degree of plastic work. This phenomenon has been recently observed by Acharyya et al. [33] who performed scanning electrochemical microscopy on polished and ground austenitic stainless steel. The deformed 20

Journal Pre-proof asperity regions will exhibit increased metal dissolution due to the higher damage to the passive film and higher surface reactivity when compared to the neighbouring regions in the active wear scar (Figure 13b). This tribocorrosion material loss mechanism comprise mechanical wear aggravated corrosion (CW) – also termed Type I corrosion-wear [5, 34]. Due to the resultant uneven dissolution, former topographical ‘hills’ will be removed and replaced by new ‘valleys’ as shown schematically in Figure 13b,c. The new topographic hills will subsequently exhibit

f

increased plastic shearing and mechanical loss giving rise to corrosion aggravated wear (WC) as

oo

shown in Figure 13c. Overall, for a given position on the wear track, material loss will take place

pr

by the actions of CW and WC synergies.

e-

A similar scenario will take place under AP conditions. Here, all the surface of the specimen is

Pr

polarised at the same potential of 100 mV vs SCE, however the dissolution rate within different regions within the contact zone is dependent on the amount of plastic work they contain. The area

al

outside of the wear scar is protected by the passive film and exhibits low currents and hence

rn

minimal dissolution rates. However, within the wear scar, the passive film is being damaged or

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removed and contributes to the measured anodic current which is related to material dissolution – also known type I tribocorrosion mechanism (Figure 8a). The degree of dissolution within the active zone will depend on the degree of prior deformation which results in uneven dissolution such that former topographical ‘hills’ will be removed and replaced by new ‘valleys’ (Figure 13 b,c). Hence, the material losses within the wear scar are expected to be via WC and CW synergies similar to tests under OCP but more severe. In fact, this resulted in greater total material losses (Figure 12) and increased roughening under AP conditions when compared to those under OCP conditions (Figure 9).

21

Journal Pre-proof CW is promoted by the ease of dissolution of surface ions into the electrolyte (Ringer’s solution) and WC is promoted by the ease of plastic flow (once a ‘hill’ and ‘valley’ topography has been created – Figure 13b, c – shear stresses being highest on ‘hills’). Hence, high values of τy (shear yield strength) suppress WC activity, whereas, CW is suppressed by passive materials that undergo limited anodic dissolution and form little (if any) oxide scale/films that are vulnerable to removal by an applied shear stress. Values of τy calculated from nano-indentation hardness (via

f

the Tresca yield criterion (τy = H/6)) are shown in Table 3. These confirm that the τy of CrN/S is

oo

greater than that of S/CrN which is greater than uncoated 316LVM. The relative ease of dynamic

pr

ion dissolution can be gleaned from Figure 8a. This shows that the resistance to anodic ion dissolution of CrN/S is greater than S/CrN which is greater than uncoated 316LVM. Hence, the

e-

total material losses due to tribocorrosion, shown by the various test materials in Figure 11, is due

Pr

to a combination of material properties. The outer CrN ceramic coating of the CrN/S coated 316LVM is outstandingly the best due to a combination of high hardness and its inability to

al

plastically deform avoiding the formation of zones with increased dislocation density having

rn

different reactivity together with its high resistance to oxidation (slow ion dissolution and little

Jo u

oxide film/scale formation) in Ringer’s solution. The outer S-phase coating of the S/CrN coated 316LVM has an intermediate hardness/yield strength and a metallic character which allows deformation by slip leading to the formation of zones of different reactivity with intermediate resistance to synergic tribocorrosion losses (CW + WC) placing it second in terms of tribocorrosion resistance. On the other hand, uncoated 316LVM being of the lowest hardness/yield strength and lowest resistance to synergic tribocorrosion losses (CW + WC) resulting in the worst tribocorrosion resistance of this study. A further point that reinforces the aforementioned hypothesis (Figure 13) is that testing under CP conditions prevents ion dissolution – only cathodic current was observed (Figure 7a). It also suppresses the formation of passive film (Figure 9d, e, f) and prevents micro22

Journal Pre-proof asperity shearing. This indicates that the entire surface was rendered cathodic when testing under CP conditions and no anodic zones were created. Hence, the formation of a ‘hill’ and ‘valley’ topography was prevented and degradation by WC and CW synergies could not take place. These parameters were therefore effectively zero so the overall synergy under CP conditions was zero. The observed volumetric total material losses (TML) under CP (determined via profilometry of the tribocorrosion test tracks), as can be seen in Figure 11, was only slight. Effectively, this

f

represents the sum of the volumetric material losses under pure corrosion (C0) and pure wear

oo

(W0) conditions alone. Moreover, given that C0 is also practically zero due to suppression of

pr

material dissolution under CP conditions, the volumetric material loss under CP conditions can be attributed entirely to W0. This latter material loss is probably a combination of superficial

Pr

e-

micro-asperity shearing and light abrasion by the Al2O3 ball counterface.

al

5. CONCLUSIONS

rn

The tribocorrosion responses of the coated and uncoated steel test-pieces was determined under

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reciprocation sliding contact with alumina (Al2O3) whilst immersed in Ringer’s solution at 37 ± 1°C. The latter were applied under open circuit potential (OCP), anodic potential (AP) and cathodic potential (CP) conditions. The following are the main conclusions from this investigation:

1. Total material losses due to tribocorrosion under AP, OCP and CP electrochemical test conditions

increased

in

the

sequence:

CrN/S

coated

316LVM

→S/CrN

coated

316LVM→uncoated 316LVM; that is, CrN/S coated 316LVM was the most tribocorrosion resistant material. 23

Journal Pre-proof

2. CrN/S coated 316LVM which has CrN as the external layer is outstandingly the best compared to the other tested materials. This response is attributed to its high hardness (21.6 ± 1.3 GPa); very good adhesion to the substrate and good load support of the CrN external layer by the underlying S-phase layer; resistance to catastrophic failure via Type II tribocorrosion damage; its ceramic crystal structure (cubic CrN with a =0.41137nm) which precludes plastic deformation

f

and associated changes is surface reactivity; and its very high resistance to material dissolution

oo

under sliding conditions. The high hardness together with the low and uniform dissolution

pr

resulted in relatively small amount of wear and synergistic (CW +WC) tribocorrosion losses and

e-

the formation of smooth scars under both OCP and anodic test conditions.

Pr

3. S/CrN coated 316LVM was resistant against catastrophic and Type II tribocorrosion damage but showed relatively poor resistance to metal dissolution during sliding when compared

al

to CrN/S. Despite its high hardness (12.7 ± 0.3 GPa) it still exhibited micro-asperity shearing via

rn

plastic deformation under both OCP and anodic test conditions. This resulted in fine debris

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generation and the formation of regions with different reactivity within the wear scar. These contributed to an increased and uneven dissolution and roughening which aggravated material losses via the degradation by WC and CW synergies.

4. Uncoated 316LVM exhibited the highest material losses of the test materials investigated under all electrochemical test potential conditions. Its low hardness (3.6 ± 0.2 GPa) and low resistance to metal dissolution during sliding resulted in roughening of the scar and relatively high material losses via the degradation by WC and CW synergies for tests under both OCP and anodic test conditions. 24

Journal Pre-proof

5. The effect of applying a positive potential during tribocorrosion testing was to increase drastically the synergistic component of the material loss for S/CrN and uncoated 316LVM testpieces when compared to that produced under OCP conditions. This behaviour was not evident for CrN/S which still displayed low material losses and a smooth scar.

ACKNOWLEDGEMENTS

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This research was funded by the Endeavour Scholarship Scheme (Malta). This scholarship is part financed by the European Union European Social Fund (ESF) under Operational Programme

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II – Cohesion Policy 2014-2020, “Investing in human capital to create more opportunities and

e-

promote the well-being of society”. The authors would also like to thank ERDF (Malta) for

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financing the testing equipment through the project: “Developing an Interdisciplinary Material

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Testing and Rapid Prototyping R&D Facility (Ref. no. 012)”.

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None.

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DECLERATIONS OF INTEREST

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A. J. Smith, P. Dieppe, K. Vernon, M. Porter, A. W. Blom, Failure rates of stemmed metal-onmetal hip replacements: analysis of data from the National Joint Registry of England and Wales, The Lancet. 379 (2012) 1199-1204.

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Y. Niu, J. Wei, Z. Yu, Microstructure and tribological behavior of multilayered CrN coating by arc ion plating, Surf Coat. Technol. 275 (2015) 332-340.

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B. Mallia, P. A. Dearnley, Exploring new W–B coating materials for the aqueous corrosion–wear protection of austenitic stainless steel, Thin Solid Films. 549 (2013) 204-215.

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B. Zhang, J. Wang, Y. Zhang, G. Han, F. Yan, Comparison of tribocorrosion behavior between 304 austenitic and 410 martensitic stainless steels in artificial seawater, RSC Advances, 6 (2016) 107933-107941.

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D. H. Buckley, Wear, in Surface effects in adhesion, friction, wear, and lubrication. ed: Elsevier, 5 (1981) 429-509.

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P. A. Dearnley, K. Dahm, H. Çimenoǧlu, The corrosion–wear behaviour of thermally oxidised CP-Ti and Ti–6Al–4V, Wear, 256 (2004) 469-479.

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rn

al

Pr

e-

pr

oo

f

[25]

27

Journal Pre-proof TABLES

Table I: Chemical composition of medical grade 316LVM supplied by L. Klein SA Chemical Composition (wt %) Element

Cr

Ni

Mo

Mn

Si

N

Cu

P

C

S

Fe

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rn

al

Pr

e-

pr

oo

f

316LVM 17.33 14.56 2.78 1.69 0.25 0.088 0.03 0.016 0.011 0.003 Bal.

28

Journal Pre-proof Table II: Deposition parameters for S/CrN and CrN/S coatings performed at an induced substrate bias of -50 V and chamber pressure of 0.27 Pa. Coating Deposition Parameters Target Material

Target Current (A)

Chamber Atmosphere

Gas flow (SCCM)

Duration (hrs)

Interlayer

Functional Coating

Under layer (S-phase)

CoCrMo

0.47

Argon

10

0.5

CoCrMo

-

CoCrMo

0.47

Argon + CH4

Ar = 10 CH4 = 3

4.5

-

CoCrMo(C) (S-phase)

Cr

0.47

Argon

10

0.5

Cr

-

Cr

0.5

Argon + N2

Ar = 10 N2 = 3

4.5

-

CrN

Cr

0.5

Argon

10

0.5

Cr

-

Cr

0.5

Argon + N2

Ar = 10 N2 = 3

4.5

-

CrN

Cr

0.5

Argon

10

0.5

Cr

-

CoCrMo

0.47

Argon + CH4

Ar = 10 CH4 = 3

4.5

-

CoCrMo(C) (S-phase)

CrN/S Top layer (CrN) Under layer (CrN)

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rn

al

Pr

e-

Top layer (S-phase)

pr

S/CrN

f

Deposition Run

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Coating

29

Journal Pre-proof Table III: Characteristics of the uncoated 316LVM, CrN/S and S/CrN. Errors indicate the standard deviation from the mean.

Cr

N

C

Co

Mo

-

-

-

-

-

-

-

41.9 ± 2.8 26.8 ± 0.8

34.9 ± 2.0

-

-

-

-

15.3 ± 2.0

45.5 ± 1.7

2.6 ± 0.1

7.8 ± 0.2

179 ± 10

3.60 ± 0.06

221 ± 5

2.12 ± 0.02

209 ± 4

pr

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6.8 ± 0.10

0.60 ± 0.06

e-

2.3 ± 0.01 2.3 ± 0.10

3.60 ± 0.20 21.60 ± 1.30 12.70 ± 0.30

f

CoCrMo(C)

Pr

S/CrN

CrN

Reduced Modulus, Er (GPa)

al

CrN/S

Hardness, H (GPa)

Top Layer Composition (at%)

rn

316LVM

Layer Thickness (µm)

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Material

Calculated shear yield strength τy (GPa)

30

Journal Pre-proof LIST OF FIGURE CAPTIONS

Figure 1: Designed coating structures deposited by PVD; (a) CrN deposited on CoCrMo(C) Sphase designated a CrN/S and (b) CoCrMo(C) S-phase deposited on CrN, designated as S/CrN. Figure 2: Glancing angle XRD results using Cu-Kα and an incident angle of 3o for (a) CrN/S (Fm-3m:225 space group, representing the top CrN layer), and (b) S/CrN (Fm-3m:225 space group, representing the top CoCrMo(C) layer).

oo

f

Figure 3: SEM micrographs depicting nano-scratch scars for (a) uncoated 316LVM, (b) S/CrN, and (c) CrN/S. Nano-scratching was performed with a conical diamond indenter having a 10 µm tip radius, under a ramped load. The micrographs show the end of the scratch track which is where the maximum applied load was reached. The arrow indicates the direction of sliding of the diamond nano-indenter.

Pr

e-

pr

Figure 4: Nano-scratch profiles for (a) uncoated 316LVM, (b) S/CrN, and (c) CrN/S. SS refers to commencement of the scratch test by the nano-indenter (with a ‘pre-load’ contact force of 0.8 mN); SS-Load refers to commencement of a ramped load of 1.5 mN/µm up to 450 mN; S F refers to the position where the scratch finishes. Figure 5: Potentiodynamic polarisation curves for uncoated 316 LVM, S/CrN and CrN/S in Ringer’s solution.

rn

al

Figure 6: (a) Dynamic OCP vs. time and (b) Dynamic friction vs. time during reciprocating corrosion-wear tests of uncoated 316LVM, S/CrN and CrN/S against an inert alumina ball under 1 N normal load and a frequency of 2 Hz in Ringer’s solution. SS refers to the sliding start whereas SF refers to the sliding finish.

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Figure 7: (a) Dynamic current vs. time and (b) Dynamic friction vs. time during reciprocating corrosion-wear tests of uncoated 316LVM, S/CrN and CrN/S against an inert alumina ball under 1 N normal load and a frequency of 2 Hz in Ringer’s solution under cathodic protection (CP) conditions. SS refers to the sliding start whereas SF refers to the sliding finish. Figure 8: (a) Dynamic current vs. time and (b) Dynamic friction vs. time during reciprocating corrosion-wear tests of uncoated 316LVM, S/CrN and CrN/S against an inert alumina ball under 1 N normal load and a frequency of 2 Hz in Ringer’s solution under anodic (AP) conditions. SS refers to the sliding start whereas SF refers to the sliding finish. Figure 9: Secondary Electron SEM images of the scars on (a,d,g) uncoated 316LVM, (b,e,h) S/CrN, (c,f,i) CrN/S following sliding, against an Al2O3 ball in Ringer’s Solution under (a-c) OCP, (d-f) cathodic and (g-i) anodic conditions. The direction of sliding contact during tribocorrosion testing was from top to bottom. Figure 10: Light optical microscopy images of the counterface material (Al2O3) after corrosionwear testing under a 1 N normal load in Ringer’s solution for (a,d,g) uncoated 316LVM, (b,e,h) 31

Journal Pre-proof CrN/S, (c.f.i) S/CrN under (a-c) OCP, (d-f) cathodic, (g-i) anodic potential conditions. The direction of sliding during corrosion-wear testing was from left to right. Figure 11: Total material volume losses for uncoated 316LVM, S/CrN and CrN/S following corrosion-wear testing under anodic, OCP and cathodic conditions. Error bars indicate the maximum and minimum error of the three repeats which were performed. Figure 12: Total material volume losses for uncoated 316LVM, S/CrN and CrN/S following corrosion-wear testing under OCP and AP conditions, clearly illustrating the corrosion-wear synergy components. The synergy at both OCP and AP conditions is composed of wear aggravated corrosion (CW) and corrosion aggravated wear (WC). Error bars indicate the maximum and minimum error of the three repeats which were performed.

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Figure 13: Proposed tribocorrosion synergisms for processes under OCP and AP conditions. (a) Before contact occurs; metal ion dissolution and passive film formation; (b) During contact; plastic shearing, fracture of hills (WC), passive film damage and wear debris generation; (c) During contact; Passive film regeneration and damage via corrosion due to wear (CW), formation of new ‘hills’ and ‘valleys’. Black areas represent zones with a higher surface reactivity and metal ion dissolution when compared to the white areas.

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FIGURES

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Figure 9

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Author contribution statement

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Raisa Chectuti: Performed the experiments; Analysed and interpreted the data; Drafted the paper. Peter Dearnley: Analysed and interpreted the data; Contributed the coatings; Conceived and designed the experiments. Antonino Mazzonello: Analysed and interpreted the data; Performed the experiments; Edited the paper Joseph Buhagiar: Conceived and designed the experiments; Analysed and interpreted the data. Bertram Mallia: Conceived and designed the experiments; Analysed and interpreted the data; Wrote the paper

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Journal Pre-proof FACULTY OF ENGINEERING Department of Metallurgy and Materials Engineering Prof. Inġ. Bertram Mallia, Associate Professor Room 222 Faculty of Engineering University of Malta Msida MSD 2080, Malta Office: (+356) 2340 2057

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www.um.edu.mt/eng/mme

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28th October 2019

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Re: Manuscript SURFCOAT-D-19-02971

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The authors of the manuscript SURFCOAT-D-19-02971 declare no conflict of interest.

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Sincerely,

Prof. Inġ. Bertram Mallia Associate Professor Department of Metallurgy and Materials Engineering Faculty of Engineering, University of Malta

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Journal Pre-proof CrN/CoCrMoC hard coated 316LVM displayed high scratch damage resistance. CoCrMoC/CrN is catastrophic damage resistant but exhibits synergistic tribocorrosion. Catastrophic damage resistant CrN/CoCrMoC displayed high tribocorrosion resistance. A model was proposed to describe wear due to corrosion and corrosion due to wear.

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Resistance to dissolution and plastic deformation supress synergistic tribocorrosion.

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