Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4V alloys

Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4V alloys

Author’s Accepted Manuscript Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4Valloys Aniruddha Samanta, Manjima Bhattacharya, Itishre...

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Author’s Accepted Manuscript Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4Valloys Aniruddha Samanta, Manjima Bhattacharya, Itishree Ratha, Himel Chakraborty, Susmit Datta, Jiten Ghosh, Sandip Bysakh, Monjoy Sreemany, Ramkrishna Rane, Alphonsa Joseph, Subroto Mukherjee, Biswanath Kundu, Mitun Das, Anoop K. Mukhopadhyay

PII: DOI: Reference:

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S1751-6161(17)30403-4 http://dx.doi.org/10.1016/j.jmbbm.2017.09.013 JMBBM2497

To appear in: Journal of the Mechanical Behavior of Biomedical Materials Received date: 10 February 2017 Revised date: 30 August 2017 Accepted date: 6 September 2017 Cite this article as: Aniruddha Samanta, Manjima Bhattacharya, Itishree Ratha, Himel Chakraborty, Susmit Datta, Jiten Ghosh, Sandip Bysakh, Monjoy Sreemany, Ramkrishna Rane, Alphonsa Joseph, Subroto Mukherjee, Biswanath Kundu, Mitun Das and Anoop K. Mukhopadhyay, Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4Valloys, Journal of the Mechanical Behavior of Biomedical Materials, http://dx.doi.org/10.1016/j.jmbbm.2017.09.013 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Nano- and micro-tribological behaviours of plasma nitrided Ti6Al4Valloys a

Aniruddha Samanta, bManjima Bhattacharya, cItishree Ratha, a,dHimel Chakraborty, cSusmit Datta, aJiten Ghosh, eSandip Bysakh, aMonjoy Sreemany, fRamkrishna Rane, fAlphonsa Joseph, f Subroto Mukherjee, cBiswanath Kundu, cMitunDas, and aAnoop K. Mukhopadhyay* a

Advanced Mechanical and Materials Characterization Division, Central Glass and Ceramic Research Institute, 196, Raja S C Mullik Road, Kolkata 700032, India, bDepartment of Chemical Sciences, Indian Institute of Science Education and Research, Kolkata, India, cBioceramics and Coating Division, CSIR-Central Glass & Ceramic Research Institute, 196 Raja S. C. Mullick Road, Kolkata 700032, India, eAdvanced Material Characterization Unit, Central Glass and Ceramic Research Institute, 196, Raja S C Mullik Road, Kolkata 700032, India, fFacilitation Centre for Industrial Plasma Technologies, Institute for Plasma Research, Gandhinagar 382428, India. Abstract Plasma nitriding of the Ti-6Al-4V alloy (TA) sample was carried out in a plasma reactor with a hot wall vacuum chamber. For ease of comparison these plasma nitrided samples were termed as TAPN. The TA and TAPN samples were characterized by XRD, Optical microscopy, FESEM, TEM, EDX, AFM, nanoindentation, micro scratch, nanotribology, sliding wear resistance evaluation and in vitro cytotoxicity evaluation techniques. The experimental results confirmed that the nanohardness, Young’s modulus, micro scratch wear resistance, nanowear resistance, sliding wear resistance of the TAPN samples were much better than those of the TA samples. Further, when the data are normalized with respect to those of the TA alloy, the TAPN sample showed cell viability about 11% higher than that of the TA alloy used in the present work. This happened due to the formation of a surface hardened embedded nitrided metallic alloy layer zone (ENMALZ) having a finer microstructure characterized by presence of hard ceramic Ti2N, TiN etc. phases in the TAPN samples, which could find enhanced application as a bioimplant material.

Keywords: Ti6Al4V; Plasma nitriding; Nanoindentation; Hardness; Scratch, Wear Corresponding Author: Tel.: +91 33 2473 3469/76/77/96; fax: +91 33 2473 0957; Present address: Central Glass and Ceramic Research Institute, 196, Raja S.C. Mullick Road, Kolkata-32, E-mail address: [email protected], [email protected] Authors E-mail [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected], [email protected] d

Presently working at National Institute for Locomotor Disabilities, Kolkata, India.

1. Introduction The global market size of metallic bioimplants including Ti-6Al-4V alloys (TA) is all set to touch ~19 billion US$ by 2017 with a further growth to ~ 27 billion US$ by 2022 (http://www.marketsandmarkets.com). During the last decade for a two year period from 2006 to 2008 the number of primary total hip replacement (THR) was 689,608 with a revision surgery

number of 79,231 in the developed countries alone (Labek et al., 2011). For the same period, the number of primary total knee replacement (TKR) was 377,462 with a revision surgery number of 19,752 in the same developed countries (Labek et al., 2011). Obviously, then the global number of primary and revision surgeries would be much more. This fact points towards a huge need for exceptionally large number bioimplants, which, in turn; justifies the large market size. Among the biomedical implants TA is unique with high biocompatibility, corrosion resistance, high specific strength but slightly inferior tribological properties (Balla et al., 2014; Geetha et al., 2009; Shum et al., 2007; Niinomi, 2003; Iannoet al., 1995; Luckey H. A., 1991). Therefore, many current efforts are directed to enhance its tribological performance further by various surface engineering techniques (Chan et al., 2016; Celik et al., 2016; Katahira et al., 2016; Kesik et al., 2016; Sivakumar et al., 2015; Xiang et al., 2015). Among these techniques, the plasma based surface engineering techniques emerged as the most promising one for commercial application (Samanta et al., 2017a; Silva et al., 2010; Ueda et al., 2010; Biswas et al., 2008; Yildiz et al., 2008; Fernandes, 2006; Fouquet et al., 2004; Yilbas et al., 1996). Further, among the plasma based techniques; the plasma nitriding (PN) process shows excellent initial prospect to enhance the tribological performance of TA (Ali and Raman; 2010; Fernandes, 2006; Molinari et al., 1997; Yilbas et al., 1996). With increment in percentage of nitrogen and duration of PN treatment both the mechanical as well as tribological properties are improved due to the formation of hard ceramic phases e.g., face centered cubic -TiN and tetragonal -Ti2N (Yildiz et al., 2008; Fernandes, 2006; Fouquet et al., 2004; Morita et al., 2001; Yilbas et al., 1996). For ease of discussion we prefer to term TA after the PN treatment as TAPN. The detailed representative literature survey on quasi-static/dynamic micromechanical properties of

surface modified Ti6Al4V is included in Table 1. The survey is not absolutely exhaustive but it is presented with the inclusion of those works which typically highlight the knowledge trend gathered over the period span of about two decades from 1996 to 2016. Based on the literature survey data presented in Table 1 a few important traits emerge out. The first is that in 9 (e.g., Yilbas et al., 1996; Molinari et al., 1997; Fouquet et al., 2004; Fernandes et al., 2006; Yildiz et al., 2008; Biswas et al., 2008; Yilmazer et al., 2009; Vasylyev et al., 2012; Mahdipoor et al., 2016) out of the 15 i.e., in 60% of the cases cited microhardness was measured; while microscratch was measured in a singular (Yildiz and Alsaran, 2010) occasion

(i.e.,

about

6.67%).

The

second

is

that

mainly mechanical

wear

and

corrosion/tribocorrosion studies were done in three (She et al., 2015; Kim et al., 2016; Zhao et al., 2016) i.e., in 20% out of the 15 cases cited. The third is that water erosion studies was done in a singular case (Batory et al., 2016) i.e., in about 6.67% of the 15 cases cited. The fourth is that there was again a singular attempt to study adherence strength (Cassar et al., 2012) i.e., about 6.67%, of the 15 cases cited. These information’s on plasma nitrided Ti6Al4V alloys provide one of the scopes e.g., study of the microscratch resistance, in the present work. Moreover, in cases where a very thin nitride layer formed (Yilbas et al., 1996; Molinari et al., 1997) microhardness in fact decreased from the surface towards the interior of the bulk material after plasma nitriding of Ti6Al4V alloys. But the surface microhardness was improved (Yilbas et al., 1996; Molinari et al., 1997) with respect to that of the pristine alloy. On the other hand, formation of a relatively thicker nitride layer along with the formations of the -TiN and Ti2N phases led to even 40% improvement in microhardness of plasma nitrided Ti6Al4V alloys (Fouquet et al., 2004; Fernandes et al., 2006; Yildiz et al., 2008). Microhardness had also increased due to formation of TiN phase with N/Ti ratio of 0.75 to 1.05 (Vasylyev et al., 2012).

Dendritic TiN phase formed during laser surface nitriding (Biswas et al., 2008) also improved Young’s modulus of Ti6Al4V alloys. Increment in surface microhardness was strongly sensitive to formation of the aforesaid phases which in turn was a strongly sensitive function of the nitriding technique, time and temperature (Yilbas et al., 1996; Molinari et al., 1997; Fouquet et al., 2004; Fernandes et al., 2006; Yildiz et al., 2008; Biswas et al., 2008; Yilmazer et al., 2009; Vasylyev et al., 2012; Mahdipoor et al., 2016). Not only microhardness, plasma nitriding also enhanced microscratch resistance, wear and erosion resistance, corrosion/tribocorrosion resistance as well as adherence of Ti6Al4V alloys along with the reduction of friction coefficient (Yildiz and Alsaran, 2010; She et al., 2015; Kim et al., 2016; Zhao et al., 2016; Batory et al., 2016; Cassar et al., 2012). Similarly the literature survey on quasi-static nanomechanical properties of surface modified Ti6Al4V is shown in Table 2. Here also the survey is not absolutely exhaustive but it is presented with the inclusion of those works which typically highlight the knowledge trend gathered over the period span of about one and a half decades from 2002 to 2016. Based on the literature survey data presented in Table 2 two important traits emerge out. The first is that in 6 (i.e., Barbieri et al., 2002; Vadiraj and Kamaraj, 2007; Biswas et al., 2008; Luo and Ge, 2009; Kim et al., 2016; Batory et al., 2016) out of the 6 studies i.e., in 100% cases nanoindentation was the preferred experimental techniques used for evaluation of the nanohardness and / or Young’s modulus of the plasma nitrided Ti6Al4V and Ti6Al7Nb alloys. The second one is that there has not been a single study reported yet on nanoscratch or nanotribology of plasma nitrided Ti6Al4V. This information, therefore, provides a major scope and justification of the present work.

There was enhancement in nanohardness, Young’s modulus, and fretting wear resistance of plasma nitrided Ti6Al4V and Ti6Al7Nb alloys due to the formations of hard -Ti2N, -TiN as well as dendritic TiN phases (Barbieri et al., 2002; Vadiraj and Kamaraj, 2007; Biswas et al., 2008; Luo and Ge, 2009; Kim et al., 2016; Batory et al., 2016). It may please be noted that as the human lifespan has been significantly extended by the considerable advancements in medical technology, the short lifespan of implanted prostheses has become one of the most critical issues in the use of e.g., hip/knee joint replacements. This is exactly the point that the present work addresses. Thus, in other words, the present work aims to develop better surface engineered bioimplant materials. It needs to be categorically understood in this perspective that the number of joint replacements is expected to increase as the ageing population grows. As of date, according to (http://www.healthline.com/health/total-knee-replacemnet-surgery/revison) in US alone 54, 000 knee implant revision surgery is done per year. Similarly, as of date, according to (http:// www.smith-nephew.com/patient/treatments/hip-treatments/hip-revision-surgery) in US alone 30,000

revision

surgery

is

done

per

year.

It

has

been

estimated

(http://

www.healthline.com/health/total-knee-replacemnet-surgery/revison) that more than about 50% of such revision surgeries is required within nearly two years of the initial application of such prosthetic bioimplants. The revision surgery is not only expensive but also painful for the patient (Liao et al., 2013). Increasing the longevity of primary implants is an important issue also from the viewpoint of increasing human life expectancy. Thus, it follows that in the US alone the number of revision surgery is about 84,000 per year. Globally this number would be at least 10 times more even on a conservative basis. Thus,

the annual global number of revision surgeries due to bioimplant failure would be at least about 10 million per year or even more. This is the main issue that propels the main objective of the present work towards better surface engineered bio-implant development since the demand for greater longevity of implants is ever increasing in order to prevent revision surgery. Now, coming to the context of the present work as mentioned earlier; due to excellent mechanical properties, high biocompatibility, specific strength and corrosion resistance, titanium alloys (Ti6Al4V) began to be widely used for joint replacements. But the poor wear resistance of titanium alloys has restricted their use for the articulating surfaces in total hip and knee joint implants (Khanna et al., 2016). Due to poor wear resistance, titanium alloys generate wear debris in load bearing implants, which lead to localized inflammation reaction causing pain and loosening of implants due to osteolysis (Khanna et al., 2016). A common assembly of an artificial hip joint consists of a Ti6Al4V femoral head articulating in a Ti6Al4V acetabular cup with a polymer liner (e.g., ultra high molecular weight polyethylene, UHMWPE) separating the two metallic parts (Endo et al., 2002; Galvin et al., 2007; Lomholt et al, 2011). Since a polymer liner is used in the standard design (Lomholt et al, 2011) it is expected that Ti6Al4V may wear the polymer a little bit (Qu et al., 2005). But the main objective of the present work was the development of the surface engineered material for the femoral head. The wear aspect of the polymer is therefore beyond the scope of the present work but will definitely constitute a major future research area in our next level of developmental research activity. Globally however, many other wear couple combinations are being tried upon (Yildiz et al., 2009; Majumdar et al., 2011; Vangolu et al, 2011; Geringer et al., 2013; Çelik et al., 2014;

Vora et al., 2014; Kao et al., 2015; Julian and Munoz, 2015; Bartolomeu et al., 2016; Yang and Hutchinson, 2016). These include not only metal (Ti6Al4V) - polymer (PTFE) but also metal (Ti6Al4V) - ceramic (Al2O3/Si3N4) and metal (Ti6Al4V) - metal (SS440C) combinations (Qu et al., 2005).

But the optimum wear couple combination is yet to be unequivocally

established (Geetha et al., 2009; Mantripragada et al., 2013; Geringer et al., 2013; Vora et al., 2014). This fact is further supported by literature survey on tribological behaviour of bioimplant materials as given in Table 3. From the data presented in Table 3 a few definitive trends emerge out. The first is that out of the 39 cases reported, in 24 cases (i.e., about 62% cases) sliding wear was the preferred test technique while the fretting wear tests were carried out in 3 cases (i.e., in only about 7% cases). The second is that only two reports were available for microscratch studies as mentioned earlier (Table 1). The third is that nanoindentation and microindentation were also conducted occasionally. The fourth and the most important one is that the sliding wear tests were carried out in dry conditions as well as under simulated body fluid (SBF), phosphate-buffered saline (PBS), bovine serum etc. environments. Finally, the fifth one is that a very wide variety of ball/pin materials e.g., Al2O3, Ruby, ZrO2, Si3N4, WC-Co, SS316L, SS440C, PTFE, UHMWPE etc. was utilized as the counter body against Ti6Al4V and other metallic bioimplant materials. These data confirmed further that due to the very complex nature of the problem a globally acceptable unequivocal choice of wear couple is yet to be established. Once this ultra-important issue of wear couple selection in a hipimplant design is globally standardized, the most important issue of wear of the counter body against Ti6Al4V can be finally critically examined definitely in future. In the present work, the choice of Al2O3 ball was deliberately done as a test of efficacy of the surface engineered TAPN

materials in a worst possible scenario wherein it is exposed against a material of much higher hardness. As alumina ball is much harder so it had such a small wear rate (e.g., 10-7mm3N-1m-1) which can be assumed as negligible for all practical considerations. One obvious way of solving this problem is to develop a plasma nitrided surface containing hard -Ti2N, -TiN etc. phases with excellent biocompatibility (Cho et al., 2013) so that such ceramic phase aid to achieve high wear resistance along with improvement of other mechanical properties as already shown in Tables 1 and 2, mentioned above. The same approach has been adopted in the present work. The final application of these Ti6Al4V materials is planned to be as a bioprosthesis or as bioimplants, especially in hip-joints. Based on the literature data provided in Tables 1, 2 and 3 and the details of the references cited therein it was found that detailed simultaneous attempt on Rietveld’s structure refinement analysis along with quantitative estimation of the amounts of different phases done in the cases of both TA and TAPN, both plan and cross sectional FE-SEM and TEM studies (Yilbas et al., 1996; Molinari et al., 1997; Yildiz et al., 2008; Kao et al., 2015; Ali and Raman, 2010; Yetim et al., 2009), microscratch and nanoindentation as well as nanotribological response measured on the plan and cross sections of TA and TAPN along with pin-on-disk sliding wear studies in SBF against a hard alumina ball and in-vitro cytotoxicity studies were not reported in earlier studies. Therefore, the scholarly merit of the present work lies in simultaneous studies of (a) Rietveld structure refinement along with quantitative estimation of the amounts of different phases in the cases of TA and TAPN materials, (b) detailed plan and cross sectional FE-SEM and TEM studies of the microstructures of TA and TAPN materials, (c) microscratch studies of TA and TAPN materials (d) detailed nanoindentation studies on plan and cross sections of both

TA and TAPN samples as a function of load from 10 to 1000 mN, (e) nanotribological evaluation of TA and TAPN materials for the first time ever in literature, (f) sliding wear studies of TA and TAPN in SBF solution using a hard alumina ball as the counter face material, (g) cytotoxicity studies using a mouse embryonic fibroblast cell line (NIH3T3) and finally (h) to have correlation of mechanical and bio-properties with physical structure and microstructure of TA and TAPN. To the best of our knowledge this is the first ever such attempt for TA and TAPN materials. Thus, the major objective of the present study was to examine the suitability and preliminary efficacy of these surface engineered TAPN materials for such application. Therefore, the major sub-objectives of the present work was thus to critically study the nanotribological responses of TA and TAPN through characterizations of (a) microstructures in plan and cross sections (b) nanomechanical properties e.g., nanohardness and Young modulus in the embedded nitride metallic alloy layer zone (ENMALZ) and the bulk metallic alloy layer zone (BMALZ) (c) micro- and nanotribological behaviour (d) sliding wear resistance of TA and TAPN against a hard alumina ball as the counter face material in the simulated body fluid (SBF) (e) investigation of in vitro cytotoxicity of the TA and TAPN samples by standard MTT assay against mouse embryonic fibroblast cell line (NIH3T3) and, finally to have (f) structure-property correlation. 2. Materials and methods 2.1 Plasma Nitriding To produce the corresponding golden colour TAPN samples, about 10 mm diameter and 5 mm thick flat parallel mirror finished greyish bio-grade (Table 4) TA samples were plasma nitride (Saikia et. al, 2013) at 800C for 6h in a 4N2:1H2 mixture and 500 Pa nitriding pressure.

The plasma nitriding (PN) process was carried out in a plasma reactor operated at 431V with a hot wall vacuum chamber. Both TA and TAPN samples were characterized by chemical analysis, X-ray diffraction (XRD), optical microscopy, field emission scanning electron microscopy (FESEM), transmission electron microscopy (TEM), energy dispersive X-ray analysis (EDX), EDX line scanning, nanoindentation, microscratch, nanoscratch and AFM based surface topographical analysis techniques. Finally, the sliding wear studies were carried out in SBF. Further the bioimplant application efficacy of the TAPN samples as compared to that of the TA samples was examined by in vitro cytotoxicity tests conducted using the standard MTT assay against mouse embryonic fibroblast cell line (NIH3T3). 2.2 Phase Analysis The X-ray diffraction patterns of the TA and TAPN samples were recorded in X’pert Pro MPD diffractometer (PANalytical) using a monochromator operating at 40kV and 30 mA using CuK1 radiation. The characteristic wavelength of the CuK1 line was measured as 1.54056 Å. The XRD data were recorded with a step size of 0.05˚ and a step time of 2 sec with 2values ranging from 20˚ to 90˚ for these samples. To gain better insight into crystal structure and to conduct quantitative phase analysis the well-known Rietveld’s structure refinement analysis technique was adopted (Kumar et al., 2013; Waje et al., 2010; Kapczinski et al., 2003; Rietveld, 1969). 2.3 Microstructural Analysis 2.3.1 Optical Microscopy The preliminary microstructural examinations of the TA and TAPN samples were conducted by a calibrated optical microscope (Olympus, GX 51, USA) with an in-built image

analysis software. The plan sections of both TA and TAPN samples were finely polished using sequentially diamond pastes of 30, 9, 6, 3, 1 and 0.25 m grit sizes for this purpose. 2.3.2 Field Emission Scanning Electron Microscopy The detailed microstructural examinations of the TA and TAPN samples were carried out by field emission scanning electron microscopy (FESEM, Supra, VP35, Carl Zeiss, Germany). The typical order of magnitude estimates for compositional details on different regions of the TAPN surface were obtained from the same machine using EDX, EDX genesis line map and EDX spot mapping techniques. 2.3.3 Transmission Electron Microscopy The finer details of TA and TAPN microstructures in plan and cross-sections were examined by transmission electron microscopy (TEM, Tecnai G2 30, S-Twin, 300 KV, FEI, The Netherlands). At first a 200 m thin slice of about 10 mm diameter was cut by a low-speed wafering machine (Isomet, Buehler, USA). Then, a ultrasonic cutter was used to obtain a 3 mm diameter disk of 200 m thickness. The disk thickness was reduced to about 100 m by a disc grinder (Gatan, USA). Further, using diamond polishing (3 m) and alumina slurry polishing (0.05 m) the thickness at the center was reduced to about 30 m by a dimple grinder (Gatan, USA). Prior to placement on copper coated carbon grid for insertion in the TEM chamber this disk was ion polished using a precision ion polishing system (Gatan, USA). To prepare samples for cross-sectional TEM examination two rectangular pieces of about 2.7 mm width, 10 mm length and 1.5 mm thickness were sliced out using the same low-speed wafering machine as mentioned above. Especially, in the case of TAPN samples the two pieces were joined by epoxy (G-1, Gatan, USA) in a way that the nitrided surfaces were kept face-toface. The assembly was cured at 125ºC. From the cured sample about 10 mm long cylindrical

sample of 2.6 mm diameter was made by grinding on emery papers. Using the same epoxy as mentioned above for fixation purpose, this sandwich cylinder was again inserted into a stainless steel tube of 2.7 mm inner diameter and 3 mm outer diameter. This final assembly was again cured at the same temperature as mentined above. At the next step this sandwich within stainless steel tube support was sectioned to obtain cross-sectional slices of about 3 mm diameter and 200 m thickness. The same low-speed wafering machine as mentioned above was used for this purpose. The rest of the steps were similar to those as mentioned above for the plan sections. 2.4

Nanotribological Studies

2.4.1 Static Contact: Nanoindentation Study Using the well-known Oliver and Pharr method (Oliver et al., 1992) the nanohardness (H) and Young’s modulus (E) were measured by the nanoindentation technique (Fischerscope H100XYP, Fischer, Switzerland) applied to both plan and cross sections of the polished TA and TAPN samples. All experiments were conducted at a temperature of about 30 ± 4oC and relative humidity of ~70 ± 5%. Both loading and unloading times were kept fixed at 30s each. To ensure statistically reliable and reproducible nanoindentation data of the samples, the machine was calibrated before each and every experiment with an independently certified (nanohardness, H ≈ 4.14 ± 0.1 GPa and Young’s modulus E ≈ 84.6 ±3.5 GPa) standard BK7 (Schott, Germany) reference glass block provided by the supplier of the machine. For both plan and cross sections at a given value of applied load always line arrays of several nanoindentations were made to take care of the statistical reproducibility issue in the embedded nitrided metallic alloy layer zone (ENMALZ) and the unnitrided bulk metallic alloy layer zone (BMALZ). The nanoindentation experiments on the plan sections were conducted at 10, 50, 100, 200, 500, 700 and 1000 mN loads while those on the cross sections were conducted

at 30, 40, 45, 60, 70, 80, 90 and 100 mN loads. The machine worked according to the DIN 50359-1 standard. It used a Berkovich indenter with a tip radius of 150 nm. The depth and force sensing resolutions of the machine were 1 nm and 0.2 N, respectively. The loading time was kept constant at 30 seconds. The unloading time was the same as the loading time. Once the prescribed maximum load was attained the sample was unloaded within the next 30 seconds. There was no hold time utilized at the maximum load on a given sample. 2.4.2 Dynamic Contact: Microscratch Study A scratch tester (TR-102-M3, Ducom, Bangalore, India) equipped with a Rockwell C diamond indenter of ~200 m tip radius was used to perform the scratch tests following work reported earlier (Yao et al., 2011). The scratch experiments were conducted at three different constant applied normal loads (PN) of 5, 10 and 15 N with a constant scratching speed (v) of 200 m.s-1. For this purpose, a compound force transducer which could measure both normal load (PN) and lateral (PL) forces with resolution of ± 0.01 N was attached to the scratching head of the scratch tester. From these data the coefficient of friction (COF=) was evaluated as (COF) or [ ⁄

], where

is the applied normal force as mentioned above and

is the tangential force

generated during the scratch experiments. Further, the scratches had a fixed length of 3 mm and they were kept 500 m apart from each other. The data reported here were obtained as the average data based on at least three scratch experiments conducted at a given applied load. The same optical microscope as mentioned above in sub-section 2.3.1 was utilized to measure the scratch width. Further, a non-contact profilometer (Contour GT-K, Bruker Corporation, USA) was utilized for confirmatory measurements of the width (W) and depth (D) of the scratch tracks. All

scratch experiments were performed in air at a temperature of about 30 ± 4oC and at a constant relative humidity of about 70±5 %. Thus, humidity effects; if any; was assumed to be constant. Therefore, while discussing the results the influence of this constant relative humidity was not specifically considered. Similarly, giving any special consideration of temperature in the discussion of the results was avoided because the changes in temperature during the present scratch experiments were not measured. Further, both low and high magnification FESEM photomicrographs of the microscratch grooves created in both TA and TAPN were taken by the same FESEM as mentioned above in the section 2.3.2 (FESEM, Supra, VP35, Carl Zeiss, Germany). 2.5 Surface Roughness Study Surface roughness parameters of both the TA and TAPN samples were measured by using a conventional AFM equipment (Veeco di CP-II instrument, lateral resolution - 0.1 nm, z resolution - 0.01 nm). All experiments were conducted in non-contact (NC) mode. A 1-10 Ohmcm Phosphorous (n) doped silicon tip with aluminium coating at the back side was used for this purpose. The data on root mean square (RMS) surface roughness (Rq), average surface roughness value (Ra), the maximum profile peak height of the surface (Rp), the maximum profile depth of valley region (Rv) and maximum peak to valley height difference (Rp-v) were measured by the AFM technique for the TA and TAPN sample surfaces. 2.5.1 Nanoscale Dynamic Contact: Nanoscratch Study Using a diamond sphero-conical indenter (SBA63) of ~2 μm tip radius and fitted inside a cantilever (serial number: HL-126) of ~6.68 mN.μm-1 stiffness the nanoscale dynamic contact experiments were carried out in a nanoscratch tester (NST: 50-133, CSM, USA) by making 1

mm long scratches induced at 1 mm.min-1 scratching speed at 100 mN normal load. The friction coefficient (COF) was evaluated as mentioned earlier. The surface topographical features, width and depth of the scratch groves were measured by non-contact mode surface profilometry using a A1-10 ohm-cm phosphorous (n) doped and back side Al coated Si tip attached to the same atomic force microscope (AFM, Veeco di CP-II instrument) as mentioned earlier in Section 2.5. The AFM had respectively, 0.1 nm and 0.01 nm of lateral and depth resolutions. 2.6 Sliding Wear Studies As mentioned earlier, the final application of these Ti6Al4V materials is planned to be as bio-prosthesis or as bio-implant materials. Such surface engineered materials are globally being developed as the load bearing surface of hip or knee implants (Wang et al., 2010). These materials will have to perform load bearing application in a human body environment and under relative movement condition. One of the preliminary ways to test the efficacy for such application is to simulate such an atmosphere during testing. That is why the sliding wear resistance tests were done in a simulated body fluid (SBF) which has composition identical to that of the inorganic part of blood plasma. The environment of the human body is buffered so that the pH is maintained at ~7.40 at 36.5 °C. Wear behavior of materials is therefore routinely studied using simulated body fluids (SBF) which simulates the inorganic part of blood plasma (see e.g., Table 3). These tests are focused on the examination of materials and provide information to evaluate their suitability for bioimplant applications. If they perform well it signifies that they should also perform well when actually placed in the body where a corrosive biological environment (Kokubo and Takadama et al., 2006; Wang et al., 2010) prevails. A comparison of nominal concentrations of ions in human

blood plasma and in simulated body fluid (SBF) at pH of 7.4 is given in Table 5 following (Mutlu and Oktay, 2013). The data of Table 5 confirmed that human blood plasma and the SBF used in the present work have identical compositions. The sliding wear resistance studies of the TA and TAPN disk samples were carried out in a ball-on-disc configuration according to ASTM G 99 - 95a standard by using a pin-on-disc tribometer (Nanovea, USA), a 3 mm diameter alumina ball, a constant normal load of 3N, fixed sliding speed of 40 mm.s-1, sliding distance of 300 m and wear track radius of 5 mm. Freshly prepared SBF solutions were utilized in all the sliding wear experiments conducted inside the test chamber kept at a constant temperature of 37°C. The volume of material worn was measured (Das et al., 2016) from the width and depth of wear tracks experimentally measured by contact profilometry (Talysurf PGI 2000S, Taylor Hobson, USA). The wear rate was calculated as the volume of material worn per unit load per unit sliding distance (Samanta et al., 2017a; Das et al., 2016).The images of sliding wear scars made in TA and TAPN samples were obtained using a table-top SEM (PhenomproX, Phenom-World B.V., The Netherlands). 2.7In vitro cytotoxicity studies The bioimplant application efficacy of the surface engineered TAPN as well as the control TA materials was felt as to be of paramount importance. Therefore, the in vitro cytotoxicity’s of the TA and TAPN samples were also assessed by standard MTT assay technique. The biocompatibility of the TA and TAPN samples were assessed in terms of in vitro cytotoxicity using MTT assay. The proliferation of cell morphology on the TA and TAPN samples was observed using the same table top SEM (PhenomproX, Phenom-World B.V., The Netherlands) as mentioned earlier in Section 2.6.

A mouse embryonic fibroblast cell line (NIH3T3) was used in the present work (Rucinska et al., 2008). All the cell culture experiments were performed on triplicate samples. Polished samples were cleaned by ultrasonication in ethanol and sterilization was done in an autoclave at 121°C for 20 min at 15 psi saturated vapour pressure. Each sample was placed in each well of 24-well plate and seeded with 1×104 cells/well. After allowing some time for attachment of cells to substrates, plate wells were filled with sufficient cell culture media (DMEM, Dulbecco’s modification of Eagle’s Medium, supplemented with 10% v/v Fetal Bovine Serum) and incubated in a humidified incubator at 37ºC and 5% CO2. Exhausted culture medium was replaced by fresh medium during experiment as required. MTT assay was performed after 3 and 5 days of culture, but the data of 5 days of culture has been included in this present study. The MTT (Sigma, St. Louis, MO) 3-(4,5dimethylthiazole-2-yl)-2,5-diphenyl tetrazolium bromide (MTT) stock solution was prepared by dissolving the MTT salt (5 mg/ml) in phosphate-buffered saline (0.01 M). After the stipulated time periods, the culture medium was removed from wells and 1 mg/ml MTT solution (200 µl MTT stock solutions + 800 µl DMEM) was added to each well in 24-well plates. Prior to that, the solution was homogenized by pipetting it in-out for several times. After 4 h of incubation (37 ºC , 5% CO2) to allow the formation of Formazan crystals, the MTT solution was removed from the well and 1 ml of solubilization solution (Di-Methyl SulphOxide, AR grade, Merck) was added to each well to dissolve the insoluble formazan crystals. Then 100 µl of the resulting purple colored solution was transferred into each well of a 96-well plate (at least three wells from each sample well) and optical density (Samanta et al,

2017b) of resultant solution was measured using a microplate reader (BioRad, USA) at 595 nm. The reference wavelength was 650 nm. As MTT salt is photosensitive, the MTT assay procedure was carefully performed to avoid exposure of treated cells to external light. Aseptic techniques were used and sterile environment was maintained during experiment. Statistical analysis was performed using the Student’s “t-test” on MTT assay data and p < 0.05 was considered statistically significant. The data is presented as mean ± standard deviation based on at least 5 to 10 independent experiments. 2.7.1 SEM Studies SEM analysis was carried out to observe how the cultured cells adhered and proliferated on samples. After 5 days of culture, the samples were taken out of well plates, rinsed 10 min each time for three times with 0.01 M PBS and fixed in 2% Paraformaldehyde-2% Glutaraldehyde solution (in 0.01 M PBS) at 4 ºC overnight. The dehydration of substrates was done using ethanol series, 30%, 50%, 70%, 95% and 100%. In each case 10 minute dehydration was done. The process was repeated three times for each of the substrates. This step was followed by another step of dehydration. It was done by ethanol-HMDS (Hexa-methyl disilazane, 98+% CAS Number 999-97-3, Alfa Aesar) series, 30%, 50%, 70%, 95% (HMDS, v/v concentration). Here also, in each case 10 minute dehydration was done. The process was repeated three times for each of the substrates. The HMDS dehydration was done inside a chemical fume hood and following safety protocols. Finally, the substrate was kept in 100% HMDS overnight for drying. Dehydrated substrates with cells fixed on it were sputter coated and observed under the SEM (PhenomproX, Phenom-World B.V., The Netherlands) as mentioned earlier in Section 2.6. 3. Results and discussion

3.1 XRD analysis The XRD patterns of TA and TAPN samples are shown in Fig. 1a and b respectively. The XRD patterns (Fig. 1a) showed that the TA sample contained hexagonal compact α-Ti and body centered cubic β-Ti phases. The peaks in the XRD pattern were accordingly indexed in Fig.1 (a). The TAPN sample contained the α-Ti (hexagonal compact) as the major phase, along with significant amount of hard ceramic phases e.g., Ti2N and TiN formed due to plasma nitriding of the TA samples. In a similar manner, the peaks in the XRD pattern were accordingly indexed in Fig.1 (b). The results of the Rietveld analysis for the TA and TAPN samples are shown in Fig. 2 (a) and (b), respectively. In addition, the quality of fitting was also assessed from the difference plot (Fig. 2a and b) as obtained by Reitveld technique (Kumar et al., 2013; Rietveld, 1969). The fitted curve matched well with the experimentally recorded XRD patterns (Fig. 2a and b). The data on the wt. % of phases, cell parameters, unit cell volumes, average crystallite sizes (d) and average lattice strains () were estimated form Rietveld analysis and are presented in Table 6 (Kumar et al., 2013; Waje et al., 2010; Rietveld, 1969). The data presented in Table 6 revealed that the TA sample contained mostly -Ti (e.g., ~86 wt.%) with a small amount (e.g., ~14 wt.%) of -Ti phases. However, the TAPN sample contained ~ 47.60, 34.70 and 17.70 wt.% of -Ti, Ti2N and TiN phases, respectively (Table 4). Possibly, the -Ti phase got stabilized after the incorporation of interstitial elements, especially nitrogen; during plasma nitriding (Edrisy and Farokhzadeh,2016). The -Ti phase in the TA sample had a strain () of 0.074 which reduced to 0.052 in the TAPN sample (Table 6). This fact implied that the crystal lattice was more relaxed because the crystallite size (d) of -Ti had increased from ~22 nm in the TA sample to ~28 nm in the TAPN

sample (Table 6). The relatively higher lattice strain in the -Ti phase of the TA sample could be due to rolling induced deformation during the production of the TA alloy. Further, the observed reduction in ( and increase in (d) was also most likely linked to the annealing of the TA sample kept for 6h at 800oC during the plasma nitriding process, as mentioned above. Accordingly, the lattice parameters for the α-Ti phases were e.g., (a=b=2.938 Å, c=4.681 Å) in the TA and (a=b=2.925 Å, c=4.684 Å) in TAPN (Table 6) samples, as expected. 3.2 Microstructure and morphology analysis 3.2.1 Optical Microscopy Study The surface morphology of the polished plan section of the TA sample appeared featureless (Fig. 3a). However, the TAPN sample surface (Fig. 3b) appeared to have a surface decorated with relatively uniform distribution of approximately round-shaped particles which had possibly formed during the PN process. 3.2.2 Field Emission Scanning Electron Microscopic study The microstructure of the plan section of the TA sample (Fig. 4a) shows a predominantly coarse acicular (plate like)  phase microstructure with very small amounts of  in the  phase boundaries. It also revealed a bimodal distribution of interconnected equiaxed primary  grains and lamellar (+ colonies which got transformed to at the corresponding phase boundaries. These results corroborated well with the XRD data (Fig. 1a). The corresponding EDX spot mapping data shown in Fig. 4b and c respectively for ( and (+ grains revealed the presence of Ti, Al and V; as expected. The relative magnitudes of Ti and Al were similar for both ( and (+ grains, however; compared to its intensity in the ( grains the peak corresponding to V was much stronger in the (+ grains.

The presence of nearly uniformly distributed hemispherical particles were visible from FE-SEM photomicrographs taken on plan section of TAPN at both low (Fig. 5a) and high (Fig. 5b) magnifications. These features corroborated well with that observed already in photomicrograph obtained by optical microscopy, Fig. 3b. Further, the significant presence of both N and Ti in the corresponding EDX spot mapping data (Fig. 5c) of TAPN confirmed the formation of a surface nitride layer (Vasylyev et al., 2012; Biswas et al., 2008; Yildiz et al., 2008; Fernandes et al., 2006; Fouquet et al., 2004; Molinari et al., 1997; Yilbas et al., 1996) comprising of TiN, Ti2N etc. phases as indicated already by the XRD data (Fig. 1). The cross-sectional FESEM image along with in-situ EDX line mapping (Fig. 5d) of the TAPN sample confirmed the presence of about 16 m embedded nitrided metallic alloy layer zone (ENMALZ). Beyond the ENMALZ of course the unnitrided bulk metallic alloy layer zone (BMALZ) prevailed up to about 30 m and even beyond although, the EDX signals were recorded up to a width of about 16 m (Fig. 5d). During the PN process nitrogen diffused through the TA surface. Therefore, the region at the immediate vicinity of the left most point the ENMALZ showed a slightly higher N concentration. At the ENMALZ the amounts of nitrogen and Al were maximum but the V content was slightly reduced compared to that of the BMALZ. These results implied that the precipitations of TiN, Ti2N etc. phases were indeed more favourable in the near surface region of the ENMALZ. Further, the lower V content implied that at least on a comparative scale the PN process effectively reduced its presence in the ENMALZ. 3.2.3 TEM Analysis of Microstructure The results of transmission electron microscopy (TEM) characterization of the TA and TAPN samples are presented in Fig.6 (a-d) and Fig.7 (a-d), respectively. The bright-field TEM

image of TA shown in Fig.6a clearly revealed the formation of predominantly elongated Ti grains with sizes of the order of about 1m where, the individual grains had relatively lower dislocation density. The local fine-scale variation of diffraction contrast in the BF image was indicative of a strained crystal lattice. This information corroborated well with that obtained from the Rietveld analysis (Table 6) of the XRD data. However, occasionally in some different regions mostly equiaxed -Ti grains were also observed to have accommodated a large number of dislocation lines (Fig. 6b). These grain localized dislocations were further confirmed by high magnification bright field (Fig. 6c) and weak beam dark field (WBDF) TEM images (Fig. 6d). The spreading of the diffraction spots in the characteristic SAD pattern shown as inset of Fig. 6(b) indicated that the disorientation of boundaries was fairly large. However, in some cases the grain boundaries were not so well defined, Fig. 6(a-c).These regions appeared to reveal some extinction contours inside the grains, Fig. 6(a-c), with high angle grain boundaries. These observations most likely implied that  phase grains had easily fragmented in the  phase boundaries. It happened possibly because compared to that of the phase; the  phase with thin strip morphology had a relatively lesser capability to accommodate severe plastic strains (Murr et al., 2009). The bright filed TEM image of Fig. 7a confirmed the presence of TiN and Ti2N precipitates (Fig. 7a) and formation of nanocrystals within the thin surface layer (Fig. 7b and c) of the TAPN sample. The spots in the SAD pattern, shown as inset of Fig. 7b, indeed confirmed the presence of Ti2N in the thin surface nitrided layer. In addition, a high density of nonequilibrium crystal defects formed possibly due to the release of the stored energy of the surfaces through preferential orientations of titanium nitride nucleation sites (Murr et al., 2009).

Localized surface melting during the PN process at 800C probably led to the formation of the nanocrystalline thin layer with the highly stochastic formation of fine droplets (Fig. 7d) in the TAPN sample. 3.2.5 Nanohardness and Young’s modulus of TA and TAPN The experimentally measured, smooth load (P) versus depth (h) plots for the polished plan sections of TA and TAPN samples are shown in Fig. 8 (a) and (b) respectively. The FESEM images of typical single nanoindents at a load of 700 mN in the cases of both TA and TAPN are shown in the inset of Fig. 8 (a) and (b), respectively. Compared to those of the TA sample (Fig. 8a) the P-h plots of TAPN samples (Fig. 8b) showed both smaller depth of penetration and higher stiffness. As a result, the TAPN samples had respectively nanohardness (H), Fig. 8 (c) and Young’s modulus (E), Fig. 8 (d) enhanced by about 120 and 60% as compared to those (e.g., H~3.3GPa, E~ 120 GPa) of the TA samples, in correspondence. This happened due to the fact that the TA sample had undergone more plastic deformation as shown in Fig. 8a (Ballarre et al., 2009; Kern et al., 2006) while the TAPN sample had undergone more of elasto-plastic deformation (Fig. 8b), as expected from the presence of relatively harder TiN, Ti2N etc. phases (Biswas et al., 2008). It is noted further from the data presented in Fig. 8 (b) that the maximum penetration depth was close to only about 1.55 m which was smaller than 10% of the ENMALZ thickness (e.g., about 16 m). Hence, the mechanical properties of the underlying BMALZ (Fig. 5d) did not affect the nanomechanical properties measured in the ENMALZ (Fig. 8 a-d). 3.2.5.1 Indentation Size Effect Nanohardness (H) of the TA as well as TAPN samples exhibited (Fig. 8c and d) indentation size effect (ISE) which is often explained in terms of geometrically necessary

dislocations (GNDs) induced by the nanoindentation experiments which impose highly steep strain gradients inside the nanoindentation cavities (Marteau et al., 2012; Asgari et al., 2011; Bouanis et al., 2011; Oliveira et al., 2009; Ohmura et al., 2005; Nix and Gao, 1998). However, out of numerous such efforts to explain ISE; one of the most utilized simple, empirical relationships to explain the ISE is Meyer’s Law (Shanholtz et al., 2013; Csehova et al., 2010; Wang et al., 2007; Xu et al., 2003) expressed as P = Chn, where h is nanoindentation depth and P is applied nanoindentation load, as mentioned earlier. Further, A and n are constants to be determined by the simple data fitting technique. In particular, n is known as the Meyer’s index and when n is <2, the presence of ISE is expected. Thus, as shown in Fig. 9 the fitting of the current experimental data (Fig. 8a,b) into the aforesaid relationship yielded n values of 1.98 and 1.42 for TA and TAPN samples, respectively. As the Meyer’s indices were < 2 for both the TA and TAPN samples (Fig. 9), it corroborated well with the experimental observation (Fig. 8c) of the presence of ISE. Although it is beyond the scope of the present work, a separate dedicated effort to understand the genesis of ISE would be the scope of our future work in these materials.

3.2.5.2Nanohardness and Young’s Modulus in Cross Section of TAPN The typical illustrative load–depth (P-h) plots obtained from the nanoindentation experiments conducted at a load of 100 mN on the cross-section of TAPN are presented in Fig. 10 (a). The typical illustrative data on distribution of nanohardness (H) and Young’s modulus (E) measured at the same load of 100 mN are shown in Fig. 10 (b) and (c), respectively as a function of distance starting from the beginning of the ENMALZ. Due to precipitation of mainly the TiN,

Ti2N etc. phases during the PN process (Fig. 1) the nanohardness (H) and Young’s modulus (E) data of the ENMLAZ of about 16 m width (Fig. 5d) were much higher (e.g., H ~ 6.5 GPa,E~250 GPa) than those (e.g., H~ 3.3 GPa, E~150 GPa) of the BMALZ that extended from ~16 m up to about 50 m (Fig. 10b,c). Further, the lower (h) of ~0.52 m compared to much higher (h) of ~0.78 m in BMALZ (Fig. 10a) corroborated very well with the higher magnitude of nanohardness measured in ENMALZ (Fig. 10b). In addition, the FESEM photomicrographs of nanoindentation arrays proved that even though made at the same load of e.g., 80 mN, the nanoindentation cavity sizes in the ENMLAZ were much smaller than those observed in the BMALZ (Fig. 11). This evidence corroborated also very well with the higher nanohardness of ENMALZ (Fig. 10b). 3.6

Micro Scratch analysis The optical photomicrographs of scratch tracks made respectively in the TA and TAPN

samples under applied normal loads (

) of 5 N (Fig. 12a,d),10 N (Fig. 12b,e) and 15 N (Fig.

12c,f) indicated that the scratch widths in the TAPN samples were relatively much smaller than those obtained in TA samples. These indications were further confirmed by actual measurements conducted by image analysis and shown in Fig. 13.

3.6.1 Variation of COF A similar trend followed for the coefficient of friction () data measured for the TA and TAPN samples, Fig. 14 (a, b). Depending on the applied normal load (e.g., 5 to 15N) the data of COF obtained in the present work beyond the initial run-in period typically varied in the range from ~0.30 to 0.40 with an average of about 0.35. Depending on the applied normal load range

(5-15N) beyond the initial run-in period COF of TAPN samples varied in the range from ~0.10 to 0.35 with an average of about 0.23. The bibliographic references to support the reliability of the obtained values of COF for TA and TAPN are provided in Table 7. From the data presented in Table 7 it is clearly evident that the COF for bare TA alloy generally varied in the range from 0.25 to 0.50 with an average of about 0.38 (Filemonowicz et al., 2005; Yildiz et al., 2008; Yetim et al., 2010; Bartolomeu et al., 2016; Chan et al., 2017). Thus, the data obtained in the present work (~0.35) was of the same order of the average data (e.g., 0.38) as reported in literature (Filemonowicz et al., 2005; Yildizet al., 2008; Yetim et al., 2010; Bartolomeu et al., 2016; Chan et al., 2017) and hence, it was accepted as reliable. From the data presented in Table 7 it is also clearly evident that the COF for TAPN alloy generally varied in the range from 0.05 to 0.55 with an average of about 0.30 (Filemonowicz et al., 2005; Xiong et al., 2007; Yildiz et al., 2008; Yetim et al., 2010;Bartolomeu et al., 2016;Chan et al., 2017). Thus, the data obtained in the present work (~0.23) was of the same order of the average data (e.g., 0.30) as reported in literature (Filemonowicz et al., 2005; Xiong et al., 2007; Yildiz et al., 2008; Yetim et al., 2010; Bartolomeu et al., 2016; Chan et al., 2017) and hence, it was accepted as reliable. The average COF data (e.g., 0.23) of TAPN alloys was much smaller than that (e.g., 0.35) obtained for the TA alloys in the present work. Similar trend was reported by other researchers as well (Filemonowicz et al., 2005; Xiong et al., 2007; Sahasrabudhe et al., 2015; Chan et al., 2017). Others reported that either COF of TA and TAPN alloys were similar (Yetim et al., 2010) or there was a marginal increase in COF of TAPN (e.g., 0..42-0.55) in comparison to that (e.g., 0.42-0.46) of the corresponding TA alloy (Yildiz et al., 2008). The decrease in COF

of TAPN with respect to that of the TA alloys was possibly linked to the decrease of surface roughness in TAPN in comparison to that of the TA alloys (Yildiz et al., 2008). Similar observations have been made also by other researchers as well (Yildiz and Alsaran, 2010; She et al., 2015; Kim et al., 2016; Zhao et al., 2016; Batory et al., 2016; Cassar et al., 2012). From the data presented in Fig. 14b, it appears that there is some variation in COF of both TA and TAPN alloys. Similar observations have been made by other researchers (see e.g., Fig. 4, Yetim et al., 2010; Fig. 7a,b, Sahasrabudhe et al.,2015; Fig. 12, Bartolomeu et al., 2016). Therefore, it seems that such variation is not uncommon at all. Rather, it appears as a real characteristic feature of the COF data of TA and TAPN materials in general. However, it will be interesting to have some idea about the genesis of such variations in COF. There were two types of variations in the COF. One type of variation was clearly dependent on load. The other type of variation COF was distance dependent. These two aspects need to be discussed separately. All experiments were conducted at the same constant speed of 200 m.s-1. Even if it is assumed that this speed is high, then; it should have caused similar magnitudes of variations in the cases of both TA and TAPN alloys. But the experimentally measured data (Fig. 14a,b) suggested that this was not the case, per se. Therefore, these variations cannot be directly linked to the speed factor. The surfaces of both TA and TAPN alloys had very minor amount of microporosities (Fig. 11a,b) if at all and hence, their possible role in these variations of COF cannot be, strictly speaking, totally ruled out. Also, if this hypothesis is correct surface microporosities of TAPN might have been decreased due to TiN, Ti2N etc. hard ceramic phase formation (Fig. 1 and Table 6). Such a reduction in surface microporosities should then be reflected in the reduction of COF

which was indeed the case (Fig. 14a,b). However, it would definitely require further detailed studies, which are beyond the scope of the present work, to derive a solid supportable correlation between the amount of surface microporosities and localized variation in COF due only to that factor. Now, it may be noted from the data presented in Fig. 14 (a) and (b) that the total sliding distance was 3 mm in each case. In the case of the bare TA alloy, the run-in length was about 0.75 mm (Fig. 14a) while in the case of the TAPN alloy it was about 1 mm. Up to these run-in path lengths the noise is due to the settling in of the dynamic contact between the moving conical indenter and the static TA and TAPN alloy surfaces. The corresponding noises in the data may be attributed to the strong initial interaction between the surface asperities on TA and TAPN alloys on the one hand and those on the synthetic diamond indenter on the other. Thus, for a given load the COF variations characteristic of the TA alloy surface are for the path length of about 0.75 mm to 3.0 mm. Similarly, for a given load the COF variations characteristic of the TAPN alloy surface are thus for the path length of about 1.0 mm to 3.0 mm. In the context of load dependence, for the TA alloys the steady state COF at 5, 10 and 15N loads were e.g., 0.40, 0.38 and 0.30, respectively (Fig. 14a). Similarly, in the case of the TAPN alloys the steady state COF at 5, 10 and 15N loads were e.g., 0.35, 0.30 and 0.10, respectively (Fig. 14b). Thus, in both the cases the COF decreased with increase in load because the contact area increased leading to a decrease in contact stress. Further, as mentioned earlier (Yildiz et al., 2008), the decrease in COF of the TAPN alloys with respect to those of the TA alloys may be linked to reduction in surface roughness parameters of TAPN as compared to that of the TA alloys.

In the context of distance dependent variations in COF of TA (Fig. 14a) at a given load such variations could be attributed to possibly concurrent yet independent contributions from several factors. These are suggested to be linked to spatial fluctuations in the lateral force due to e.g., (a) statistically random distribution of asperities on the surface, (b) high magnitudes of all average surface roughness parameters, (Yildiz et al., 2008) (c) possibly adhesive wear characteristic of titanium metals due to the rapid make and break of asperity contact junctions created between the asperities on TA alloys and the synthetic diamond indenter (Qu et al., 2005) and (d) statistically random distribution of the  and  phases (Fig. 1 and Table 6) on the surface of TA alloys (Filemonowicz et al., 2005) which governs the spatial variations in the nanohardness of the Ta alloy. Further dedicated studies, which are beyond the scope of the present work, would be definitely needed in future to verify the validity of the suggestions made here. Similarly, in the context of distance dependent variations in COF of TAPN (Fig. 14b) at a given load such variations could be attributed to possibly independent yet concurrent contributions from several factors. These factors are suggested to be due to spatial fluctuations in the lateral force linked to e.g., (a) higher nanohardness of TAPN surface (b) statistically random spatial distribution of nanohardness on the TAPN surface and its obvious dependence on the statistically random distribution of TiN and Ti2N phases (Fig. 1 and Table 6) on the surface of TAPN alloys (Xiong et al., 2007), (c) probability of more asperity interactions because of highly reduced surface roughness which means that at a given load for the same contact area more asperities of TAPN surface will be interacting with those on the diamond indenter surface (d) reduction in adhesive wear of TAPN due to reduced surface roughness (Dong et al., 2000) and (e) possible formation of a transfer layer followed by the consequent periodic but localized

fracture of the transfer layer because of the higher nanohardness of the TAPN alloys. However, as mentioned earlier further detailed studies which are beyond the scope of the present work, would be definitely required to be conducted in future to check out the validity of the propositions made here. Further, the relative variations in COF beyond the run-in lengths of e.g., 0.75 mm (Fig. 14a) in the TA and 1.0 mm (Fig. 14b) in the TAPN alloys were actually reduced with increase in load, especially at higher loads. For instance, in the case of TA alloys (Fig. 14a) for path differences of 1 to 1.5 mm, the variations in COF (e.g., COF) were 0.003, 0.001 and 0.002 respectively, for applied normal loads of 5, 10 and 15N. Similarly, for the consecutive path difference of 1.5 to 2.0 mm, COF were 0.022, 0.028 and 0.001 for applied normal loads of 5, 10 and 15N, respectively. Further, for the consecutive path difference of 2.0 to 2.5 mm, COF were respectively 0.019, 0.010 and 0.001 for applied normal loads of 5, 10 and 15N. Similarly, for the consecutive path difference of 2.5 to 3.0 mm, COF were respectively 0.009, 0.006 and 0.001 for applied normal loads of 5, 10 and 15N (Fig. 14a). In the case of the TAPN alloys, in a similar manner for path differences of 1 to 1.5 mm, COF were 0.051, 0.015 and 0.010 respectively, for applied normal loads of 5, 10 and 15N (Fig. 14b). Similarly, for the consecutive path difference of 1.5 to 2.0 mm, COF were 0.010, 0.010 and 0.012 for applied normal loads of 5, 10 and 15N, respectively. Further, for the consecutive path difference of 2.0 to 2.5 mm, COF were respectively 0.015, 0.006 and 0.001 for applied normal loads of 5, 10 and 15N. Similarly, for the consecutive path difference of 2.5 to 3.0 mm,

COF were respectively 0.028, 0.020 and 0.003 for applied normal loads of 5, 10 and 15N (Fig. 14b). If the hypothesis of adhesive junction formation is assumed to be valid, then more transfer layer formation following breakage of the adhesive junction layer would be expected to happen at especially the higher loads particularly at higher sliding distances which provided more time for such interaction to happen. This should reduce the magnitude of COF as had indeed happened. Depending on the extent of their periodic formation and fracture and consequent spatial distribution in between the moving diamond indenter and the TA or TAPN surface (Qu et al., 2005) the extent of relative spatial variation of the lateral force and hence COF would be determined. If this picture is correct, then the relative variations in COF beyond the run-in lengths should reduce with the increase in applied normal load as indeed appeared to have had happened in the experimental data plotted in Fig. 14 (a) and (b). While such a match between the experimental data and the speculative explanation of the same appears to be encouraging, it must be appreciated at the same time that fully dedicated further studies, which are beyond the scope of the present work, would be definitely desired in future to accept the mechanisms of relative small variation in COF of the TA and TAPN alloys as proposed in the present work.

3.6.2 Noncontact Profilometry Study The typical representative 2D optical images along with the confirmatory data on both width (W) and depth (D) of the microscratch grooves as obtained by noncontact profilometry of

TA and TAPN samples and the corresponding 3D views are shown in Figs. 15 (a, b and c), 16(a, b and c) and 17 (a, b and c) for the applied normal loads of 5, 10 and 15 N, respectively. The corresponding widths of the TA samples were respectively measured as 53m from Fig. 15 (a, b), 73 m from Fig. 16 (a, b) and 75 m from Fig. 17 (a, b) for the applied normal loads of 5, 10 and 15N in correspondence. Similarly, the corresponding depths for the applied normal loads of 5, 10 and 15N were measured respectively as 3.047 m from Fig. 15 (b, c), 3.832m from Fig. 16 (b, c) and 5.064m from Fig. 17 (b, c). The corresponding widths of the TAPN samples were in turn measured as 47 m from Fig. 15 (d, e), 59 m from Fig. 16 (d, e) and 62 m from Fig. 17 (d, e) for the applied normal loads of 5, 10 and 15N. In a similar manner, the corresponding depths for the applied normal loads of 5, 10 and 15N were measured respectively, as 3.048 m from Fig.15 (e, f), 3.202 m from Fig. 16 (e, f) and 3.706 m from Fig. 17 (e, f). These results further confirmed that most likely due to significant surface hardening caused by the presence of TiN, Ti2N etc. phases; thescratch widths (W) and depths (D) of the TAPN samples were much smaller than those of the TA samples. Since the scratch tester used a conical indenter of about 200 m tip radius, the contact zone may be assumed as hemispherical. So, the load bearing contact area (Ac) was approximated from Figs 15-17 as, (Ac) ~ reff2 where (reff) is given by (W/2).Thus, based on the scratch width (W) data presented for TA and TAPN samples in aforementioned Figs. 15-17; (Ac) were evaluated to be (~2200, 4183and 4416 m2) and (~1734, 2733 and 3018 m2) for applied normal loads of 5, 10 and 15N, respectively. Therefore, the corresponding contact stresses (e.g., c~P/Ac) developed due to the applied normal loads of 5, 10 and 15 N on the TAPN samples

(e.g., 2.88, 3.65 and 4.97 GPa) were much higher than those (e.g., 2.27, 2.39 and 3.39 GPa) developed on the TA samples due the same normal loads. Further, assuming the yield strength (y) of TA and TAPN as ~0.93 GPa (Zherebtov et al., 2005); the ratio (c/y) were (~2.44, 2.56 and 3.64) and (~3.09, 3.92 and 5.34) for the TA and TAPN samples respectively, for applied normal loads of 5, 10 and 15N. Hence; in general; there was localized yielding induced flow which could reduce the friction coefficient (COF). Since, the (c/y) ratios in TAPN were much higher than those of TA; the extent of localized yielding induced flow caused a larger reduction of (COF) as a function of load, Fig. 14. The higher the applied normal load, the higher was the ratio (c/y) and hence the higher was the extent of reduction in (COF) with load as shown in Fig. 14. Further, the presence of TiN, Ti2N etc. phases within ~16 m of the ENMALZ surface might have had reduced (COF) through their lubricating properties and non-textured nature (Patel et al., 2010; Hove et al., 2015). Again, it has been suggested that the generation of micro wear debris inside the scratch groove may be responsible for decrease in the friction coefficient of the TA samples (Blau, 1992). Therefore, as TAPN was much harder than TA; the possibility of the generation of more wear debris in TAPN than in TA samples (Fig. 12) could not be, per se, strictly ruled out. As a matter of fact it appears from the optical photomicrograph (Fig. 12) to be an independent yet parallel possibility, as suggested above. It follows thus from the aforesaid discussions based the set of data displayed in Figs. 12-17 that there was some minute amount of material loss that had happened in both TA and TAPN during the microscratch experiments. This aspect therefore deserves special attention. As the indenter had a large tip radius of about 200 m, for the purpose of simplification the scratch grooves were roughly approximated as shallow triangular zones having the width

( ) and depth ( ) of the actual scratches as measured by the non-contact profilometry technique mentioned above. Thus, the volumes (

) of material worn out of the TA and TAPN samples

were estimated using the following Equation (1) as: [(

where,

)

]

(1)

is the length of the scratch groove.

Subsequently, the wear rate at a given applied normal load of (

) producing a wear volume

was estimated using the following Equation (2) as: [

]

(2)

Therefore, using respectively the Equations (1) and (2) mentioned earlier, both microscratch induced wear volume and microscratch induced wear rate due to the applied normal loads of (

) 5, 10 and 15 N were evaluated for both TA and TAPN. It may be noted in this context that

for all three applied normal loads of 5, 10 and 15 N the width and depth of the scratch tracks in TAPN (Figs. 15-17) were indeed much smaller than that (Figs 15-17) of the TA sample. As a result, the average wear volumes of the TAPN samples (e.g., ~2.14 x 10-13, 2.83 x 10-13 and 3.44 x 10-13 m3 measured respectively at applied normal loads of 5, 10 and 15 N) were relatively much smaller than those (e.g., ~2.42 x 10-13, 4.2 x 10-13 and 5.7 x 10-13 m3 measured respectively at normal loads of 5, 10 and 15 N) of the TA samples. As mentioned above, all wear volumes were calculated by using Equation (1). These data are shown in Fig. 18 (a) as the double log plot of wear volume versus (

).

The interesting point is that the wear volumes of TA and TAPN exhibited empirical power law dependencies with a positive exponent on the applied normal load (

). The

corresponding result obtained for the TA sample by the empirical fitting is given in the following Equation (3) as: (

) [(

)

]

(3)

The high magnitude of the goodness of fit (e.g., R2~0.999) suggested that the empirical Equation (3) represented the trend of dependence of the experimentally obtained (

) data on (

)

properly. Similarly, the corresponding result obtained for TAPN by the empirical fitting is given in the following Equation (4) as: (

) [(

)

]

(4)

The high magnitude of the goodness of fit (e.g., R2~0.997) suggested that the empirical Equation (4) represented the trend of dependence of the experimentally obtained

data on (

)

properly. Based on the data plotted in Fig. 18(a) and by using Equation (2) the wear rates ( both TA and TAPN were measured. For TAPN, the magnitudes of (

) of

) were measured to be ~

1.43 x 10-11, 9.44 x 10-12 and 7.65 x 10-12 m3.N-1m-1 at applied normal loads (

) of 5, 10 and 15

N, respectively. Thus, the wear rate of TAPN decreased by about 50% as the applied normal load

was increased from 5 to 15N. This fact corroborated well with the fact that TAPN had much reduced values of the COF (i.e., ). Further, these wear rates were obviously much smaller than those e.g., ~1.61 x 10-11, 1.39 x 10-11 and 1.26 x 10-11 m3.N-1m-1 measured respectively for applied normal loads of 5, 10 and 15 N in the TA samples. Here also the wear rate of TA decreased but only by about 22% as the applied normal load was increased from 5 to 15N. These data are shown in Fig. 18 (b) as the double log plot of wear rate versus (

).

Akin to the case of wear volume, the wear rate also; exhibited empirical power law dependence on the applied normal load but with a negative exponent. The negative exponent physically means a reduction in (

) with increase in(

). The corresponding result obtained

for the TA sample by the empirical fitting is given in the following Equation (5) as: (

) [(

)

]

(5)

The high magnitude of the goodness of fit (e.g., R2~0.997) suggested that the empirical Equation (5) represented the trend of dependence of the experimentally obtained

data on (

)

properly. Similarly, the corresponding result obtained for the TAPN sample by the empirical fitting is given in the following Equation (6) as: (

) [(

)

]

(6)

The high magnitude of the goodness of fit (e.g., R2~0.998) suggested that the empirical Equation (6) represented the trend of dependence of the experimentally obtained properly.

data on (

)

It is interesting to note that the pre-exponent term in Equation (6) was twice as high as that noted in Equation (5) obtained for the TA sample. Further, the magnitude of the negative exponent in Equation (6) derived for TAPN was nearly thrice as that obtained in Equation (5) derived for the TA sample. Thus, even at the highest applied normal load of 15 N, both wear volume and wear rate of the TA sample were about 40% higher than those of the TAPN samples. Therefore, the rate of change of wear volume with applied normal and rate of change of wear rate with applied normal loads was expected to be much smaller in case of TAPN samples as compared to those of the TA samples. This is exactly what had happened indeed. The rate of change of wear volume of TA with applied normal load was 3.27 x 10-14 m3.N-1. The magnitude for the rate of change of wear rate of TA with applied normal load was evaluated to be 3.49 x 1013

(m3.N-1.m-1).N-1. On the other hand, the rate of change of wear volume of TAPN with applied

normal load was 6.67 x 10-13 m3.N-1 which was much smaller than that of the TA sample. Further, the magnitude for the rate of change of wear rate of TAPN with applied normal load was evaluated to be 1.29 x 10-14 (m3.N-1m-1).N-1 which, in turn, was much smaller than that of the TA sample. In this context it may be noted that since the synthetic diamond indenter was used in the microscratch experiments, it had hardness (e.g., about 70-150 GPa) much higher than those of the TA (e.g., ~3 GPa) and TAPN (e.g., ~8 GPa) alloys. Therefore, only mild abrasive wear with localized plastic ploughing as the mechanism of very small wear volume (e.g., ~10-13 m3, Fig. 18a) and wear rate (e.g., ~10-11-10-12 m3N-1m-1, Fig. 18b) was expected to happen. In addition, due to plastic ploughing process in metals at least the materials are usually pushed out towards the edges of the plastically deformed scratch groove. This process subsequently leads to formation of pile ups at the edges of the scratch grove. FESEM based

evidences of both plastic ploughing and pile up formations have been provided later in the present work. But it needs to be understood that for localized ploughing to occur, the precursor stage is the localized occurrence of plastic flow induced deformation. The localized plastic flow deformation can happen in turn, if and only if the localized tensile stress active at the wake of the moving indenter that caused the genesis of the microscratch is higher than the yield stress of the given material. Even in every static indentation process a tensile stress is generated at the back of the indenter. When the indenter is in motion e.g., as in the present micro-scratch experiments, this tensile stress is significantly enhanced (Lawn, 1967; Lawn and Frank, 1967). Since the indenter tip had a radius of 200 m, the maximum tensile stress due to static contacts (

) under the normal load (PN) of 5, 10 and 15 N in the present experiments were

therefore calculated from Hertzian contact mechanics using the following Equation (5) as (Lawn, 1967; Lawn and Frank, 1967; Bandyopadhyay et al., 2012): ) ⁄

(

(7)

is the unit of stress, the corresponding contact radius is ‘as’ and s is the

where,

Poisson’s ratio of the sample. The magnitudes of half of the width of the corresponding scratch grooves made at applied normal loads of 5, 10 and 15 N in TA and TAPN samples and measured by non-contact profilometry (Figs. 15-17) were taken as “as” for calculating the magnitude of

,

the unit of stress. Since the indenter tip had a radius of 200 m, the maximum tensile stress due to static contacts (

) under a normal load

(5, 10 and 15 N) in the present experiments were

therefore again calculated from the Hertzian contact mechanics using the following Equation (8) as (Lawn, 1967; Lawn and Frank, 1967; Bandyopadhyay et al., 2012): ( where,

)

(8)

is the COF and other terms are already defined. Thus, the maximum tensile stresses created at the wake of the indenter due to both static

(

) and dynamic (

) contacts made at the applied normal loads of 5, 10 and 15 N for the

TA and TAPN samples are shown respectively in Fig. 19 (a) and (b). It is particularly interesting to note that in both TA and TAPN the magnitudes of ( (

) were much higher than those of

), as expected. Moreover, depending on the applied load the magnitudes of (

) varied

in the range 6.43 to 7.25 GPa in the case of TA which had a relatively higher (  value. Similarly, depending on the applied load, the magnitudes of (

) varied in the range 4.70 to

7.42 GPa in the case of TAPN which had a relatively lower ( value. Nevertheless, the magnitudes of these maximum tensile stresses (Fig. 19a,b) were much higher (e.g., ~4.7 to 7.5 GPa) than the yield strength (e.g., y ~0.93 GPa) as mentioned earlier (Zherebtov et al., 2005). Therefore, as (

) was » y, localized plastic flow induced deformation was most

expected to happen. It follows as a consequence, that physically this process would have to occur through localized plastic ploughing. This is exactly what is observed from both low and high magnification FESEM photomicrographs presented below for the microscratch grooves made at 5, 10 and 15 N in both TA (Fig. 20 a-f) and TAPN (Fig. 21 a-f) samples. The low magnification FESEM photomicrographs presented respectively in Fig. 20 (a), (b) and (c) showed the extensive plastic ploughing in the scratch grooves created by the applied normal loads of 5, 10 and 15 N in the TA sample. The corresponding high magnification FESEM

photomicrographs presented in Fig. 20 (d), (e) and (f) confirmed the observation made in the low magnification FESEM photomicrographs (Fig. 20 a-c). Further, as a result of the plastic ploughing process the displaced TA material was piled up just beside the edge of the corresponding scratch groves; as indicated by the white arrows in Fig. 20 (a), (b) and (c). The low magnification FESEM photomicrographs presented in Fig. 21 (a), (b) and (c) respectively showed the extensive plastic ploughing in the scratch grooves created respectively by the applied normal loads of 5, 10 and 15 N in the TAPN sample. The corresponding high magnification FESEM photomicrographs presented in Fig. 21 (d), (e) and (f) confirmed the observation made in the low magnification FESEM photomicrographs (Fig. 21 a-c). In addition, some very minor microcracking appeared to be present in the scratch groove made at especially very high load of e.g., 15 N in TAPN (Fig. 21c,f). The occurrence of such slight microcracking could be linked to the high nanohardness of the TAPN layer. However, there was no large scale delamination or fracture observed in any of the scratch grooves made at 5, 10 and 15 N in the TAPN samples. Further, as a result of the plastic ploughing process; the displaced TAPN material was piled up just beside the edge of the corresponding scratch groves as indicated by the white arrows in Fig. 21 (a), (b) and (c). But, their obvious presence was relatively lesser noticed compared to the case of the TA samples (Fig. 20 a-c), presumably due to much higher nanohardness of the TAPN sample (Fig. 8c) as compared to that of the TA sample. However, further dedicated work; which is beyond the scope of the present work, will be necessary to look into ultimate details of the underlying complex deformation and damage mechanisms particularly in a surface hardened layer comprising of nanocrystalline Ti2N, TiN etc. hard

ceramic phases dispersed in the ductile TA matrix as was the case of the present TAPN samples. The aforesaid discussion has provided both theoretical justification and experimental evidence in support of the occurrence of plastic deformation in both TA and TAPN samples. Thus, all the aforesaid discussion proved beyond doubt that localized plastic deformation induced ploughing was the major operative mechanism behind the ultralow wear volumes of both TA and TAPN. In this connection it may be noted that scratch hardness (

) was also evaluated for both

TA and TAPN samples using the following Equation (9) as (Bellemare et al., 2010): (

(

)

where, (

) ⁄(

(9)

)

) is the applied normal load utilized in the microscratch experiments and as before, the

corresponding contact radius is ‘as’. Here also, as mentioned earlier; the magnitudes of half of the width of the corresponding scratch grooves made at 5, 10 and 15 N applied normal loads in TA and TAPN samples and measured by non-contact profilometry (Figs. 15-17); were taken as “as” for calculating the magnitude of

, the scratch hardness.

These data are plotted as a function of the applied normal load for both TA and TAPN samples in Fig. 22. The results showed that (

) increased respectively from about 4 to 6.5 GPa

and about 6 to 9 GPa for TA and TAPN as the load was increased from 5 to 15 N. As conceptually scratch hardness means the intrinsic resistance of a given materials against scratch formation and related abrasive processes of material removal, these data confirmed again that TAPN was much more scratch resistant than the TA material.

3.7

Nanotribology of TA and TAPN samples For the applied normal load of 100 mN, the nanotribological data of the TA and TAPN

samples are shown in Fig. 23 (a, b). The corresponding AFM images are shown in Fig. 24 (a-d). The experimentally measured data on nanoscratch depth and (COF) of the TA and TAPN samples are shown respectively; in Fig. 23 (a) and (b) as functions of the sliding distance (e.g., 1 mm). The average scratch depth was about 9 m (Fig. 23a) for the TA sample but it reduced to about 7.5m (Fig. 23b) in the case of the TAPN sample. Similarly, the average magnitude of (COF) data was about ~ 0.4 (Fig. 23a) for the TA sample but it reduced by about 50% to about ~ 0.2 (Fig. 23b) for the TAPN sample. The reductions in nanoscratch depth and (COF) data happened due to significant surface hardening caused by the presence of TiN, Ti2N etc. phases (Fig. 1) as mentioned earlier; while explaining the results of the microscratch experiments (Figs.12-20). The typical topographical AFM photomicrographs of the nanoscratches made on the TA and TAPN samples are shown in Fig. 24 (a) and (b), respectively. The corresponding 3D AFM photomicrographs of the nanoscratches made on the TA and TAPN samples are shown respectively in Fig. 24 (c) and (d). For the TAPN sample the width of plastically deformed nanoscratch groove (Fig. 24b) was also much smaller than that of the TA sample (Fig. 24a). The 3D views confirmed further that both width and depth of the plastically deformed nanoscratch groove in the TAPN sample (Fig. 24d) were much smaller than those of the TA sample (Fig. 24c). These data (Figs. 23 and 24) implied that due to the significant surface

hardening provided by the PN process, the TAPN samples had nanowear resistance much higher than that of the TA sample. 3.7.1 Surface Roughness Studies The data on root mean square (RMS) surface roughness (Rq), average surface roughness value (Ra), the maximum profile peak height of the surface (Rp), the maximum profile depth of valley region (Rv) and maximum peak to valley height difference (Rp-v) as measured by the AFM technique for the TA and TAPN sample surfaces respectively, are shown in Tables 8 and 9. The grand average values of the root mean square (RMS) surface roughness (R q), average surface roughness value (Ra), the maximum profile peak height of the surface (Rp), the maximum profile depth of valley region (Rv) and maximum peak to valley height difference (Rp-v) of the TAPN samples were about 70% smaller than those of the TA sample. Thus, these data corroborated nicely with the suggestions made above in Section 3.6.1 above. In addition, this information confirmed that the surface of the TAPN sample was much smoother as well as finer compared to that of the TA sample. The surface topographies of the TA and TAPN samples were examined by the same AFM technique, as mentioned earlier. The results are shown in Fig. 24 (a-d). These data indeed confirmed the formation of a relatively finer surface microstructure (Fig. 24b, d) in the TAPN sample as compared to that (Fig. 24a, c) of the TA sample. 3.8 Sliding wear studies The data related to the results obtained from the sliding wear experiments in SBF solution are given in Table 10.The corresponding optical microscopy based images and FESEM images of the sliding wear tracks are given in Figs. 25 and 26, respectively. The surface mode contact profilometry measurements showed that the average cross sectional area of the wear tracks on

TAPN sample was only about 0.178 x 10-3 mm2 and it was about 99% smaller than that (e.g., 14.7 x 10-3 mm2) measured for the TA sample (Table 10). It happened because the width of the wear track in TAPN (Fig. 25b) was much narrower than that (Fig. 25a) of the TA sample. As a result, the average wear volume of the TAPN samples (e.g., ~5.432 x 10-3 mm3) was about 99% smaller than that (e.g., ~4.642 x10-1 mm3) of the TA samples (Table 10). The most interesting observation was that even when sliding against a hard alumina ball in SBF, the wear rate of TAPN (e.g., ~6.2 x 10-6 mm3.N-1m-1) samples was about 99% smaller than that (e.g., 5.158 x 10-4 mm3.N-1m-1) measured for the TA samples (Table 10). It was found further that as alumina ball was much harder so it had such a small wear rate (e.g., 10-7mm3.N-1m-1) which can be assumed as negligible for all practical considerations. It is further interesting to note that in spite of the differences in the experimental conditions, the present wear rates matched well with literature data (Yildiz et al., 2008; Kao et al., 2015). The wear volume and hence, the wear rate of the TAPN samples was much smaller as it

had nanohardness about 120% higher than that measured on the plan section of the corresponding TA sample (Fig.8c).Even in the cross section of TAPN, the nanohardness of ENMALZ was about 196% higher than that of the BMALZ (Fig. 10b). These results had also corroborated very well with results obtained from both microscratch (Section 3.6) and nanoscratch studies (Section 3.7). The SEM images of the sliding wear track portions are shown in Fig. 26 (a) and (b) for the TA and in Fig. 26 (c) and (d) for the TAPN samples. The white arrows in Fig. 26 (a, b, c and d) represent the respective directions of sliding. The sliding wear track of the TA sample showed huge plastic deformation and the occasional presence of plate like micro-wear debris due to adhesive wear (Fig. 26a). The magnified view of the localized shear deformation induced

microcrack seen in Fig. 26 (a) is shown in Fig. 26 (b). From the photomicrograph presented in Fig. 26 (b) the micro-crack appeared to be limited within the heavily plastically deformed thin surface layer of the TA sample. The sliding wear track of the TAPN sample did not show any signature of huge plastic deformation, Fig. 26 (c). It happened because plastic deformation and adhesive wear were minimized due to high nanohardness prevalent at the surface of the TAPN sample. Further, it may be noted from Fig. 26 (c) that there was very insignificant amount of surface region where the thin layer on the surface of TAPN sample suffered localized micro-fracture during the sliding wear experiments in SBF. The magnified view of the localized micro-fracture zone seen in Fig. 26 (c) is shown in Fig. 26 (d). From the photomicrograph presented in Fig. 26 (d) some microabrasion grooves appeared to be present. It happened possibly due to occasional localized microbreakage of hard TiN, Ti2N etc. phases present in the ENMALZ region of the TAPN sample. 3.9

In vitro cytotoxicity studies The cell viability of the TA and TAPN were determined over 5 days of culture, as

mentioned above by optical density method. A higher cell viability value of TAPN over TA especially after 5 days of culture suggests that the substrate is sufficiently biocompatible and favors adhesion/proliferation of cells. It is well documented that TA has a high level of biocompatibility (Cho et al., 2013). When the data are normalized with respect to those of the TA alloy, the TAPN sample showed cell viability about 11% higher than that of the TA alloy (Fig. 27) used in the present work. The corresponding SEM images are shown in Fig. 28 (a, b). The SEM photomicrograph confirmed that cultured NIH3T3 cells on TA had negligible proliferation (Fig. 28a) after 5 days of culture while proliferation of the NIH3T3 cells on the TAPN substrates was

abundant (Fig. 28b) after 5 days culture. These results confirmed our findings from MTT assay that TAPN developed in the present work was sufficiently biocompatible and favored cell adhesion/ proliferation of NIH3T3 cells. 4. Conclusions The major conclusions of the present work are: (a) The TAPN sample contained only ~ 47.60 wt.% of -Ti phase which was nearly half of that present in the TA alloy but it contained a substantial amount of hard e.g., ~34.70 of Ti2N and ~17.70 wt.% of the TiN phases. Further, due to about 30% reduction in lattice strain (e.g., from 0.074 in TA to 0.052 in TAPN) it had slightly relaxed lattice structure and hence a highly smooth surface because it had all surface roughness parameters smaller by about 70% as compared to those of the TA sample. (b) The experimental results confirmed that due to the formation of adherent internally embedded surface nitrided layer comprised of the dispersed nanocrystals of TiN, Ti2N etc. hard ceramic phases and the high density of non-equilibrium crystal defects formed possibly due to the release of the stored energy of the surfaces through preferential orientations of titanium nitride nucleation sites in the TAPN sample, the nanohardness (H) and Young’s modulus (E) of TAPN were respectively about 120% and 60% higher than those of the TA alloy. (c) The combination of highly reduced surface roughness,

presence of the adherent

internally embedded surface nitrided layer comprised of the dispersed nanocrystals of TiN, Ti2N etc. hard ceramic phases as well as the relatively higher magnitudes of H and E in TAPN, caused about 40% reductions in microscratch induced wear volume (e.g., ~3.45 x 10-13 m3 in TAPN versus ~5.70 x 10-13 m3 in TA) and hence, in wear rate (e.g., ~7.66 x 10-12 m3N-1m-1 in TAPN versus ~1.27 x 10-11 m3N-1m-1 in TA) of TAPN; even at the highest applied load of 15 N. In

addition, the same factors of TAPN led not only to about 20% decrease in the nanoscratch depth of TAPN compared to that of the TA sample but also to about 99% reduction in sliding wear rate (e.g., ~6.2 x 10-6 mm3.N-1m-1 in TAPN vs. ~5.16 x 10-4 mm3.N-1m-1 in TA) of TAPN along with a concurrent reduction of about 52% in the COF of TAPN compared to that of the TA sample; when sliding against a hard alumina ball in SBF. (d) Due to the formation of the highly biocompatible Ti2N and TiN phases, after 5 days of culture the proliferation of the NIH3T3 cells on the TAPN surface was abundant as compared to negligible proliferation on the TA sample and hence, the cell viability of TAPN was about 11% higher than that of the TA sample. These results finally proved that in addition to having extraordinary micro and nanotribological properties as well as high nanohardness and Young’s modulus; the TAPN material developed in the present work was sufficiently biocompatible and favored adequate adhesion/ proliferation of the NIH3T3 cells.

Acknowledgments The authors are grateful to Dr. K. Muraleedharan, Director, CSIR-Central Glass and Ceramic Research Institute (CGCRI), Kolkata for his kind permission to publish this paper. The authors appreciate the infrastructural supports received from all colleagues and particularly; those received from the colleagues of the Advanced Mechanical and Materials Characterization Division, (AMMCD). The authors appreciate the experimental supports received from Dr. V K Balla of CSIR-CGCRI. Finally, the authors gratefully acknowledge the financial supports received from the Department of Science and Technology, Govt. Of India through sponsored Project No. GAP 0240.

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List of Figures Ti2N

20

30

40

+

50

60



 50

60

2 (in Degree) Fig.1: XRD patterns of (a) TA and (b) TAPN.

90



 40

80

TA 

30

TiN

70



20

Ti2N Ti2N



Ti2N



TiN

Ti2N+

Ti2N Ti2N+ TiN



TAPN

TiN

Intensity (a.u.)

 = -Ti, = -Ti

70

 80

90

Counts

Counts Ti6Al4V (R) 86.1% a Ti Titanium 86.1 % ß Ti Titanium 13.9 % 13.9%

(a)

TA

300

400

Plasma spray-ni-Ti6Al4v 86.1% a Ti Titanium 47.6 % TiNTiN 17.717.7% % Ti2N Ti34.8 2N %

(b)

TAPN

300 200

200

100

100

0

0 30

40

50

60

70

80

20

90

30

Position [°2Theta] (Copper (Cu)) 40 20 0 -20 -40

40

50 60 70 Position [°2Theta] (Copper (Cu))

80

100 50 0 -50 -100

Fig. 2: Line profile analysis of different samples done by the conventional Rietveld technique(a) TA and (b) TAPN.

(b)

(a)

100 m

Fig. 3: Optical micrographs of plan sections (a) TA and (b) TAPN

100 m

2 m

2 m

Fig. 4: FE-SEM micrographs of plan section of (a) TA (b, c) EDX analysis of the top surface of TA.

TiKa

(d)

(c)

0 m

NKa

Distance

16 m

TiKb

Fig. 5: FE-SEM micrographs of plan section of (a, b) TAPN (c) EDX analysis of the top surface of TAPN (d) Cross sectional FE-SEM micrographs along with line EDX analysis of TAPN.

(c)

Fig. 6: Transmission Electron Microscopy (TEM) of TA(a) Bright-field (BF) TEM image shows elongated Ti grains(b) -Ti grains were also observed to have accommodated a large number of dislocation lines. SAD pattern shown as inset of Fig. 6(b) indicated that the disorientation of boundaries was fairly large (c) High magnification bright field TEM confirms grain localized dislocations (d)Weak beam dark field (WBDF) TEM image also confirms the same.

(a)

(b)

(d)

Fig. 7: Transmission Electron Microscopy (TEM) of TAPN (a) Low magnification bright-field (BF) TEM image shows the presence of TiN and Ti2N precipitates (b, c) High magnification bright field TEM confirms formation of nanocrystals within the thin surface layer. SAD pattern shown as inset of Fig.7 (b) indicated the same (d) Nanocrystalline thin layer formed with stochastic formation of fine droplets.

10 mN 50 mN 100 mN 200 mN 500 mN 700 mN 1000 mN

(a)

Load (mN)

800 600

700 mN

1000

TA

(b)

800

Load (mN)

1000

400

600

700 mN

TAPN

400

300 nm

300 nm

200

200

0.5

1.0

1.5

2.0

Depth (m)

2.5

3.0

TA TAPN

(c) 10 8 6 4 2 0 200

400

600

Load,P (mN)

0.5

1.0

1.5

2.0

2.5

Depth (m)

12

0

0 0.0

3.5

800

1000

Young's modulus, E (GPa)

0 0.0

Nanohardness, H (GPa)

10 mN 50 mN 100 mN 200 mN 500 mN 700 mN 1000 mN

300

TA TAPN

(d)

250 200 150 100 50 0 0

200

400

600

800

1000

Load, P (mN)

Fig. 8: Load-depth plot analysis of (a) TA, (b) TAPN, (c) nanohardness and (d) young’s modulus vs. load of both TA and TAPN.

TA TAPN 3

P (mN)

10

n=1.42 n=1.98

2

10

-1

10

0

h (m)

10

Fig. 9: Fitting of the current experimental data according to Meyer’s Law in case of both TA and TAPN.

Load, P (mN)

80

(a) ENMALZ BMALZ

60

40

20

0 0.0

0.2

0.4

0.6

0.8

1.0

Depth, h(m) 10

280

(b)

6 5

3 2 1

BMALZ

0

240 220 200 180

ENMALZ

Young's modulus (GPa)

7

4

(c)

260

8

ENMALZ

Nanohardness (GPa)

9

160 140 120

BMALZ

100 0

10

20

30

Distance (m)

40

50

0

10

20

30

40

50

Distance (m)

Fig. 10: (a) load-depth plot of two distinct zonese.g ENMALZ and BMALZ (b) nanohardness and (c) young’s modulus vs. measurement distance along cross-sectionalofTAPN.

Fig. 11: FESEM photomicrograph of nanoindentation impression along cross section of TAPN.

(a)

(d)

100 m

100 m

(b)

(e)

100 m

100 m

(c)

(f)

100 m

100 m

Fig. 12: Optical micrograph of micro-scratch track of (a, b, c) TA and (d, e, f) TAPN under 5N, 10 N, 15 N loads respectively.

100

TA TAPN

Scratch width (m)

90 80 70 60 50 40 30 20 10 0

1

2 10

5

3 15

Scratch load (N)

Fig. 13: Scratch width variation measured by image analysis softwarefor both TA and TAPN under different loading condition.

1.0 0.8

TA 0.6

Coefficient of friction(

5N 10 N 15 N

(a) Coefficient of friction(

1.0

0.4 0.2 0.0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

Scratch length (mm)

0.8

5N 10 N 15 N

(b)

TAPN

0.6 0.4 0.2 0.0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

Scratch length (mm)

Fig. 14: Co-efficient of friction (COF) vs. Scratch length for micro-scratch of (a) TA and (b) TAPN.

(a)

0.4

2.5 1.5 0.5

0.2 mm

0.6

-0.5 -1.5

0.2 0.2 mm

0

0

0.2

3.6

0.4

0.6

0.8

-3.18

(b)

W~ 53 m

3.047 m

m

5N

4 m

TA mm

mm

0.1 mm

TAPN

-0.5 -2.0

0.2 00

0.2 mm

0.2

0.4

0.6

0.8

W~ 47 m

3.048 m

0.4

3.8 (e) 2.5 1.0

0.2 mm

0.6

m mm

4 m

(d)

-3.5 mm

-5.18

0.1 mm

Fig. 15: Representative 2D optical images along with the confirmatory data on both width (W) and depth (D) of the microscratch grooves measured by noncontact profilometry and the corresponding 3D views case of (a, b, c) TA and (d, e, f) TAPN samples at applied normal load (PN) load of 5N.

TA

10 N

(a)

W~ 73 m

(b)

m

mm

4.23

0.5

-0.2

0.1 mm

0.13 0.2 0.3 0.4 0.5 0.6

-4.27 0.7 mm

TAPN (d) mm

6.2 3.0 0

0.4 0.1 mm

0.196

(e)

m

0.479 0.3

0.1 mm

0.1 mm

0.13 0.2 0.3 0.4 0.5 0.6

-5.7 0.7

W~ 59 m

3.202 m

0.172

4 m

0.1 mm

0.3

1.0 -0.5

4 m

0.4

3.832 m

2.5

mm

0.1 mm

Fig. 16: The typical representative 2D optical images along with the confirmatory data on both width (W) and depth (D) of the microscratch grooves measured by noncontact profilometry and the corresponding 3D viewscase of (a, b, c) TA and (d, e, f) TAPN samples at applied normal load (PN) load of 10N. .

TA (a)

15 N m

mm

W~ 75 m

(b)

5.7

0.4

0 0.1 mm

-5.0

TAPN (d)

0.1 mm

m

mm

0.1 mm

0.6 0.5 0.4 0.3 0.2 0.1 0 0

mm

0.2

0.2 mm

0.4

0.6

0.8 mm

7.0 5.0 3.0 1.0 -1.0 -3.0 -5.0 -7.1

(e)

W~ 62 m

3.706 m

0.15 0.2 0.3 0.4 0.5 0.6 0.7

4 m

0.1 mm

0.167

5.064 m

0.3

4 m

3.0

0.1 mm

Fig. 17: The typical representative 2D optical images along with the confirmatory data on both width (W) and depth (D) of the microscratch grooves measured by noncontact profilometry and the corresponding 3D viewscase of (a, b, c) TA and (d, e, f) TAPN samples at applied normal load (PN) load of 15N.

-12

-10

10

10

TA TAPN

-13

8x10

TA TAPN

-11

W V = 2x10 . (PN)

-0.22

TA

6x10

-11

10

-1

Wear Rate, WR (m N m )

3

Wear Volume, WV (m )

-13

W V = 7x10 . (PN)

-13

TA

-11

W V = 4x10 . (PN)

-0.57

TAPN

3

4x10

0.78

-1

-14

TAPN -13

2x10

-13

W V = 10 . (PN)

0.426

-12

10

-13

10

(a)

(b) -14

4

5

6

7

8

9

10

10

11 12 13 14 15 16

4

Applied Normal Load, PN (N)

5

6

7

8

9

10

11 12 13 14 15 16

Applied Normal Load, PN (N)

Fig. 18: Wear Data both experimental and predicted (Following Equations 3, 4, 5 and 6) of TA and TAPN as a function of applied normal load (PN) of 5, 10 and 15 N: (a) wear volume and (b) wear rate.

10

4

(GPa) max

TAPN 6

8

P (N)

10

12

14

0.1 16 18

0.1

4

6

8

10

12

0.1 14 16 18

P (N)

Fig. 19: Maximum tensile stresses active at the wake of the indenter due to static (σstmax) and dynamic (σdtmax) contacts as a function of the applied normal load (PN) of 5, 10 and 15 N: (a) TA and (b) TAPN.

t

1

t

max

1

TA 0.1

10

(b)

d

(GPa) max

t

s

d

(GPa)

1

t

max

1

(GPa)

10

(a)

s

10

TA 5N

2 m

20 m

10 N

(c)

20 m

15 N

20 m

(b)

(a)

(d)

2 m

(f)

(e)

2 m

Fig. 20: The FESEM images of the scratch grooves made by applied normal loads of 5, 10 and 15 N respectively in TA samples at lower (a, c, e) and corresponding higher magnifications (b, d, f). In all the FESEM photomicrographs scratch directions are shown by white arrows.

TAPN 5N

2 m

20 m

10 N

20 m

(d)

(c)

20 m

15 N

(b)

(a)

2 m

(f)

(e)

2 m

Fig. 21: The FESEM images of the scratch grooves made by applied normal loads of 5, 10 and 15 N respectively in TAPN samples at lower (a, c, e) and corresponding higher magnifications (b, d, f). (Please add arrows showing scratch directions in all the Figures a-f). In all the FESEM photomicrographs scratch directions are shown by white arrows.

14

TA TAPN

HS (GPa)

12 10 8 6 4 2 0 1

5

2 10

Scratch load (N)

3 15

Fig. 22: Scratch hardness (Hs) of TA and TAPN as a function of the applied normal loads (PN) of 5, 10 and 15 N.

0.8

8 0.6 6 0.4 4 0.2

2 0 0.0

0.2

0.4

0.6

0.0 1.0

0.8

10

1.0

(b)

TAPN 0.8

8 0.6 6 0.4 4 0.2

2 0 0.0

0.2

0.4

0.6

0.8

Scratch distance (mm)

Scratch distance (mm)

Fig. 23: Nano-scratch analysis of (a) TA and (b) TAPN.

TA

(a)

(c)

TAPN

(b)

(d)

Fig. 24: 2D AFM micrographs of nano-scratch track of (a) TA and (b) TAPN (c, d) corresponding 3D views respectively.

0.0 1.0

Coefficient of friction

TA

(a)

Scratch depth (µm)

10

12

1.0

Coefficient of friction

Scratch depth (µm)

12

(a) TA

(b) TAPN

Fig. 25: Optical micrograph of Wear track of (a) TA (b) TAPN.

TA

(b)

(a)

15 m

50 m (c)

(d)

10 m

Fig. 26: The SEM images of the sliding wear track portions (a, b) for the TA and (c, d) for the TAPN samples, wear directions are given by white arrows.

Percentage of viability

140 120

5 Days

100 80 60 40 20 0

TA

TAPN

Fig. 27: Plot of percentage of cell viability in case of TA and TAPN after 5 days.

Fig. 28: SEM images of mouse embryonic fibroblast cell (NIH3T3)proliferated over after 5 days respectively on (a) TA (b) TAPN.

List of tables Table 1: Literature survey on quasi-static/dynamic micromechanical properties Plasma nitride/Surface Modified Ti6Al4V Sustrate

Quasi Static micro mechanical Properties

Remarks

Ref.

Microhardness decreased from increase distance from the surface and -TiN and -Ti2N are formed due to nitridation Microhardness improved due to nitridation but decrease from increase distance from the surface and -TiN and -Ti2N are formed due to nitridation Microhardness increased with thick nitrided layer formation of -TiN and -Ti2N -TiN and -Ti2N phases formed and 40% increment in microhardness in case of nitrided sample and microhardness decreased with increase in depth from the surface -TiN and -Ti2N phases formed and surface roughness, microhardness and compound layer thinkness increased with increase in nitriding time Formation of dendritic TiN phase, microhardness, nanohardness and young’s modulus improved by laser surface nitriding technique Plasma ion Nitriding processing parameters controlled phase modifications, surface microstructure and micro-hardness Friction co-efficient increased with the increased number of cycles during multi-pass scratch tests, nitrided Ti6Al4V was more scratch resistant than Ti6Al4V Increment of microhardness happened due to formation of TiN phase with N/Ti ratio of 0.75 to 1.05 as confirmed by XPS

Yilbas et al., 1996

Micro indentation

Micro Scratch

TA



x

TA



x

TA



x

TA



x

TA



x

TA



x

TA



x

TA

x



TA



x

TA

x

x

Ti–6Al–4V with a hard and well-adhered surface compound layer had superior adhesion strength and improved mechanical properties

Cassar et al., 2012

TA

x

x

Plasma nitriding effectively enhanced the wear resistance, and reduced the friction coefficient

She et al., 2015

TA, Ti-6Al-7Nb

x

x

Improvement of wear and corrosion resistance properties occurred after surface nitriding of both the alloys

Kim et al., 2016

TA



x

Mahdipoor et al., 2016

TA

x

x

TA

x

x

Gas nitriding generated a very hard and brittle compounds layer, it resulted in 44% higher erosion resistance along with increment in micro hardness as a function of nitriding temperature and time Mechanical Wear and tribocorrosion performance of the conventional metallic biomaterials with plasma nitriding was improved Water erosion resistance increased significantly in case of low-temperature RF plasma nitriding for Ti6Al4V alloy

*Ti-6Al-4V~TA,

Molinari et al., 1997

Fouquet et al., 2004 Fernandes et al., 2006

Yildiz et al., 2008

Biswas et al., 2008

Yilmazer et al., 2009

Yildiz and Alsaran, 2010

Vasylyev et al., 2012

Zhao et al., 2016

Batory et al., 2016

Table 2: Literature survey on quasi-static nanomechanical properties Plasma nitrided Ti6Al4V Sustrate

Quasi Static Nano mechanical Properties Nano indentation

Nano Scratch

TA



x

TA



x

TA



x

TA



x

TA, Ti-6Al-7Nb



x

TA



x

Remarks

Ref.

Fivefold increment of nanohardness occurred due to the formation of hard -Ti2N and -TiN nitrided phases during plasma jet nitriding Nanohardness and Fretting response of surfacemodified (plasma nitrided or multilayered) alloys was much more improved due to enhancement in surface properties Formation of dendritic TiN phase, nanohardness and young’s modulus improved by laser surface nitriding technique Both nanohardness and fretting wear resistance in dry and bovine serum lubrication condition increased after surface modification by nitrogen Increment of nanomechanical properties occurred after surface nitriding of both the alloys

Barbieri et al., 2002

Nano mechanical properties increased as thick brittle compound layer of Ti2N/TiN formed on the surface of Ti6Al4V

Vadiraj and Kamaraj, 2007

Biswas et al., 2008

Luo and Ge, 2009

Kim et al., 2016

Batory et al., 2016

*Ti-6Al-4V~TA

Table 3: Literature Survey on Test Techniques used for Tribological Characterization of bioimplant materials Material TA TA TA and TO-TA

Test Technique Microindentation, Sliding Wear, Ruby Ball, ET025 oil Microindentation, Dry Sliding Wear, TA Pin

Ref. Yilbas et al., 1996 Molinari et al., 1997

TA

Rolling Sliding Wear under boundary lubrication condition Nanoindentation

Barbieri et al., 2002

TA

Microindentation

Fouquet et al., 2004

TA, Ti–6Al– 2Sn–4Zr–2Mo TA,TAPN

Dry Sliding Wear with SS440C, Al2O3, Si3N4 and PTFE Pins, Sliding Wear

Qu et al., 2005

TA

Microindentation

TA,TAPN

Dry Sliding Wear, UHMWPE pin

TA

Nanoindentation

TA

Nanoindentation

Vadiraj and Kamaraj, 2007 Biswas et al., 2008

TA

Microindentation, Dry Sliding Wear, WC-Co Ball,

Yildiz et al., 2008

TA

Microindentation

Biswas et al., 2008

Dong et al, 2000

Filemonowicz et al., 2005 Fernandes et al., 2006 Xiong et al., 2007

TA

Sliding Wear, Al2O3 ball, Ringer’s solution

Yildiz et al.,2009

TA TA

Yilmazer et al., 2009

TA, TAPN TA, TAPN TA, TAPN TA

Microindentation Nanoindentation, Fretting Wear, Si3N4 ball, Bovine Serum Dry Sliding Wear, WC-Co Ball Fretting Wear Microscratch Microscratch

TA

Sliding Wear, WC ball, SBF

Ti–35Nb– 7.2Zr–5.7Ta TA

Sliding Wear, High Cr Steel ball, SBF Microindentation

Majumdar et al., 2011 Vasylyev et al., 2012

TA

Adhesion Strength

Cassar et al., 2012

Ti6Al7Nb

Fretting Wear, ZrO2 ball, Naturally aerated solution

Cai et al., 2013

Co-Cr-Mo

Sliding Wear, Al2O3 ball, SBF

Çelik et al., 2014

TA

Sliding Wear, AISI 52100 steel ball

She et al., 2015

CP Ti TA, TAPN

Sliding Wear, Si3N4 ball, DI water and 10% fetal bovine serum Sliding Wear, SS316L ball, SBF

Sahasrabudhe et al., 2015 Kao et al., 2015

Co-Cr-Mo

Sliding Wear, Al2O3 ball, SBF

TA

Sliding Wear, Al2O3 ball, PBS

TA

Sliding Wear, Al2O3 ball, SBF

TA

Dry Sliding Wear, UHMWPE Pin

TA

TA

Dry and Lubricated Sliding Wear with UHMWPE Pin Nanoindentation, Sliding Wear and Corrosion Resistance Microindentation and erosion resistance

Julian and Munoz, 2015 Bartolomeu et al., 2016 Yang and Hutchinson, 2016 Saravanan et al., 2016 Guezmil et al., 2016

TA

Sliding Wear and Tribocorrosion

Mahdipoor et al., 2016 Zhao et al., 2016

TA

Nanoindentation, Water Erosion Resistance

Batory et al., 2016

TA, TAPN

Sliding Wear, TAPN ball, Fetal bovine serum

Chan et al., 2017

TA, Ti-6Al-7Nb

Luo and Ge, 2009 Yetim et al., 2009 Ali and Raman, 2010 Yetim et al., 2010 Yildiz and Alsaran, 2010 Vangolu et al, 2011

Kim et al., 2016

*Ti6Al4V~TA, *Plasma Nitrided Ti6Al4V~TAPN, *Thermally oxidized Ti6Al4V~TO-TA, *Commercially available Ti~CPTi

Table 4: Chemical composition of Ti6Al4V Elements

Ti

Al

V

Composition (%)

88.72

6.67

4.26

Table 5:Ion concentrations of a simulated body fluid (SBF) and human blood plasma Ion

Simulated body fluid (mM)

Human blood Plasma (mM)

Na+

142.0

142.0

Ca+2

2.5

2.5

5.0 1.5

5.0 1.5

+

K Mg+2 Cl-

147.8

103.0

-

4.2

27.0

-2

1.0

1.0

0.5

0.5

HCO3 HPO4 SO4

-2

Table 6: Values of cell parameters, cell volume, average crystallite size and average lattice strain as obtained by Rietveld analysis Sample

Wt. % of Phases

Cell parameters (Ǻ)

Cell

Crystallite

Lattice

Volume (Ǻ)3

Size (nm)

Strain (%)

α Ti: 86.10 %

a= b = 2.938578, c= 4.681871

35.01262

21.89

0.074

β Ti: 13.90 %

a= b = 3.228645

33.65586

5.07

0.218

α Ti: 47.60 %

a= b = 2.925826, c= 4.684947

34.7322

27.69

0.052

TiN: 17.70 %

a= b = 4.223302

75.32794

24.00

0.064

Ti2N: 34.70 %

a= b = 4.9410, c= 3.0300

73.97285

72.70

0.185

TA

TAPN

Table 7: Lit survey on COF of TA and TAPN Substrate

Method of Determination of COF (Wear/Scratch)

COF

Ref.

TA and TAPN

Wear

TA~0.4, TAPN~0.13

Filemonowicz et al., 2005

TA and TAPN

Wear

TA~0.25, TAPN~0.05

Xiong et al., 2007

TA and TAPN

Wear

TA~0.42-0.46, TAPN~0.42-0.55

Yildiz et al., 2008

TA and TAPN

Micro Scratch

TA~0.45, TAPN~0.49, TA/DLC~0.21

Yetim et al., 2010

CPTi

Wear

CPTi~1.1-1.2, TiN/CPTi~0.7-0.9

Sahasrabudhe et al.,2015

TA

Wear

TA~0.50

Bartolomeu et al., 2016

TA and TAPN

Wear

TA~0.4, TAPN~0.15

Chan et al., 2017

TA and TAPN

Micro Scratch

TA~0.30-0.45, TAPN~0.13

Present study

*Ti6Al4V~TA, *Plasma Nitrided Ti6Al4V~TAPN, *Commercially available Ti~CPTi

Table 8: AFM data on Surface roughness parameters (Rq, Ra, Rp, Rv, Rp-v) of Ti6Al4V alloy (TA)

Grand Average Standard Deviation

(RMS) Surface Roughness (Rq) in (nm)

Avg Surface Roughness (Ra) in (nm)

Maximum profile peak height of the surface (Rp) in (nm)

Maximum profile depth of valley region (Rv) in (nm)

Maximum Peak to Valley Height Difference (Rp-v) in (nm)

9.67

8.05

14.63

-24.85

39.48

10.38

8.46

15.90

-27.20

43.10

10.01

8.39

14.98

-26.38

41.36

10.56

8.94

15.06

-27.65

42.71

10.76

9.12

16.16

-27.64

43.80

10.28

8.59

15.35

-26.74

42.09

0.44

0.43

0.65

1.18

1.70

Table 9: AFM data on Surface roughness parameters (Rq, Ra, Rp, Rv, Rp-v) of TAPN Rms Roughness (Rq) of surface in (nm)

Ave Roughness of surface (Ra) in (nm)

Peak Asperity Height (Rp) on surface (nm)

Maximum Asperity Valley Depth from surface in (nm)

Asperity Peak to Valley Height Difference (Rp-v) in (nm)

4.45

4.07

7.23

-6.49

13.72

4.89

4.25

6.67

-11.07

17.74

2.23

1.76

3.62

-7.00

10.62

2.89

1.95

4.43

-9.95

14.38

2.57

1.99

4.55

-7.77

12.32

Grand Average

3.40

2.80

5.30

-8.46

13.76

Standard Deviation

1.19

1.24

1.57

1.97

2.66

Table 10: The Sliding Wear Volumes and Wear Rates of the TA and TAPN Samples Sample

Average Wear Volume (mm3)

Wear rate (mm3.N-1m-1)

TA

Average cross sectional area (mm2) 14.7 x 10-3

4.642 x 10-1

5.158 x 10-4

TAPN

0.178 x 10-3

5.432 x 10-3

6.2 x 10-6

*Ti6Al4V~TA, *Plasma Nitrided Ti6Al4V~TAPN

Highlights:    

The embedded nitrided metallic alloy layer zone (ENMALZ) was formed. Nanocrystalline Ti2N, TiN etc. phases are formed due to plasma nitriding. Nano/micro- tribological properties enhanced significantly after plasma nitriding. TAPN material is sufficiently biocompatible.