Journal of Alloys and Compoundsnuvant.com/wp-content/uploads/2017/03/Mechanical... · Mechanical properties and corrosion behavior of the nitriding surface ... (PTFE, (C2F4)n). A

lable at ScienceDirect

Journal of Alloys and Compounds 679 (2016) 138e148

Contents lists avai

Journal of Alloys and Compounds

journal homepage: http: / /www.elsevier .com/locate/ ja lcom

Mechanical properties and corrosion behavior of the nitriding surfacelayer of Tie6Ale7Nb using large pulsed electron beam (LPEB)

Jisoo Kim, Woo Jin Lee, Hyung Wook Park*

Department of Mechanical Engineering, Ulsan National Institute of Science and Technology, UNIST-gil 50, Eonyang-eup, Ulju-gun, Ulsan, 689-798, Republicof Korea

a r t i c l e i n f o

Article history:Received 19 February 2016Received in revised form5 April 2016Accepted 6 April 2016Available online 8 April 2016

Keywords:TitaniumNitridingLarge pulsed electron beamCorrosionHardening

* Corresponding author.E-mail address: [email protected] (H.W. Park).

http://dx.doi.org/10.1016/j.jallcom.2016.04.0600925-8388/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

Large pulsed electron beam (LPEB) irradiation was used as a single surface finishing process for TieA-le7Nb. Nitrogen plasma gas and cathodic apparatus have been adopted to induce nitriding effect ofTie6Ale7Nb during the electron beam irradiation. The atomic concentration of nitrogen atoms at the re-solidified layer could be achieved up to ~18% by LPEB nitriding. Nano-hardness in the re-solidified layerwas improved by ~75% following the irradiation process, as a result of a phase transformation and theformation of TiN. The re-solidified layer induced by the LPEB nitriding, consisted of TiN, TiO2, and TiOxNy,indicated significantly modified corrosion resistance showing a nobler corrosion potential, decreasedcorrosion current density, and improved charge transfer resistance. The increasing fraction of TiN at there-solidified layer, induced by LPEB nitriding, was suggested as being responsible for remarkableimprovement of mechanical properties and corrosion resistance, embedding uniformly noble and stablecharacteristics at the top surface. The corrosion-resistant surface layer with superior mechanical prop-erties on Tie6Ale7Nb has been successfully demonstrated by LPEB nitriding technique.

© 2016 Elsevier B.V. All rights reserved.

1. Introduction

In recent decades, titanium and titanium alloys have beenstudied extensively because of their high weight-to-strength ratio,good corrosion resistance, and biocompatible behavior [1,2]. Ac-cording to previous research, a naturally formed thin oxide layer(TiO2) protects the basic material from corrosion by acting as apassive film [3]. Especially, Tie6Ale4V (Grade 5) alloy has beenused most commonly for orthopedic implants because the combi-nation of a and b phase structures makes an appropriate surfacehardness and elastic modulus [4]. Due to the galvanic couple be-tween the a and b phases, however, it has been demonstrated thatthe corrosion resistance is decreased versus single-phase titaniumalloys [5,6]. Additionally, the b-stabilizer, vanadium (V), of Tie6A-le4V, can form a cytotoxic agent (V2O5) in bodily fluids [7,8]. Forthese reasons, since the 1980s, other aþb phase and b-phase tita-nium alloys have been actively developed. One of the best possiblesubstitutes for Grade 5 is Tie6Ale7Nb alloy (Protasul-100), forwhich the corrosion resistance has been reported to be greatlysuperior and the b-stabilizer, niobium (Nb), has been suggested to

be non-toxic in bodily fluids, unlike V [9,10].Nevertheless, Tie6Ale7Nb has some fundamental problems due

to its readily corrupted thin oxide layer (1e4 nm) with a poorgalvanic couple. As a result, additional corrosion-resistant coatingsare necessary, such as anodization and electrochemical deposition[11e14]. The anodization method is the most frequently used sur-face treatment of titanium and its alloys, and it has been reportedthat anodization of titanium alloys led improved corrosion resis-tance with a nobler TiO2 [11,14,15]. The deposition of CaeP im-proves the corrosion resistance and biocompatibility [16]. However,as Song [17] and Gu et al. [18] referred, mechanical properties ofcorrosion-resistant films have been one of the most importantconcerns on electrochemically stable surface layer. As corrosionresistance of metallic materials is related to chemical stability ofoxide layers and materials themselves, mechanical properties ofsurface films are generally regarded independently [18]. In thiscontext, there is a clear and strong need to develop surface treat-ment that can induce corrosion-resistant surface layer with supe-rior mechanical properties. Recently, nitriding has received wideattention because a nitrogen-implanted surface has been found tobe chemically inert and thermally stable [19,20]. Remarkably, manystudies have supported that titanium nitride (TiN), used as a hardceramic coating material, has nobler corrosion resistance with

Fig. 1. Schematic diagram of the experimental set-up for large pulsed electron beam(LPEB) nitriding.

Table 1Parameters of large pulsed electron beam irradiation process.

Parameters Value

Acceleration voltage 30 kVEnergy density 10 J/cm2

Plasma gas pressure (Ar and N2) 0.05 PaPulse duration 2 msIrradiation frequency 0.1 HzNegative DC bias 0 to �1000 VNumber of pulses 0 to 200 pulses

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148 139

superior surface hardness [21e23].Nitriding methods to date can be classified into three different

coating techniques: gas nitriding, plasma nitriding, and laser gasnitriding (LGN). Gas nitriding, a chemical vapor deposition (CVD)method, implants nitrogen ions using the catalytic decompositionof ammonia gas (NH4) [24,25]. Although it is a cost effectivemethod, it has been reported that it should be avoided for bio-materials because of the toxicity of NH4 [26]. Plasma nitriding, suchas plasma-assisted ion implantation, direct current (DC)magnetronsputtering, and focused ion beam implantation, has been usedextensively industrially [27,28]. However, some disadvantages havebeen reported that the coating film can readily become delami-nated from the substrate due to its brittleness, and surface defects,such as micro-cracks, pores, and cavities, are hardly removed [29].Finally, laser gas nitriding is a direct energymethod that uses a laserbeam for the melted substrate to react with nitrogen (N2) gas.Although, this method enables relatively rapid processing at at-mospheric pressure, LGN also causes coating cracks and it is inef-ficient for complex shapes and large areas [30,31]. To summarizethe nitriding techniques, fabricating a uniformly stable TiN layerwithout defects and inhibition of delamination are important toincrease corrosion resistance.

Large pulsed electron beam (LPEB) treatment has been intro-duced and modified as an innovative finishing process for surfacetreatment and polishing in a single step [32]. Surface finishingusing electron beam irradiation has been widely adopted formagnesium- and iron-based alloys [33e35]. According to previousresearch with LPEB irradiation on Tie6Ale7Nb, it was found thatcorrosion resistance and nano-hardness were improved by LPEB.Additionally, small TiN segments were observed on the LPEB-treated surface [36]. However, nitrogen implantation was obvi-ously small on the re-solidified. Thus, the possibility of titaniumnitriding using LPEB has not yet been reported.

In the present study, we attempted to develop a crack-lessnitriding process by combining a nitrogen ion implantation sys-tem using LPEB and a cathodic process to improve the corrosioncharacteristics and nano-hardness of Tie6Ale7Nb. LPEB irradiationwas performed using argon (Ar) and N2 plasma gas and the resultsof nitriding effects were evaluated in terms of surface morphology,microstructure, nano-hardness, wear resistance, and corrosionresistance.

2. Experimental

2.1. Preparation of samples

A 25-mm diameter and 5-mm-thick Tie6Ale7Nb (supplied byChangsheng Titanium Co., Ltd.) samples were prepared by cuttingextruded bar of the alloy. The samples were ground and polishedfrom 180- to a 1200-grit finish (Allied High Tech Products). AfterLPEB irradiations, the samples were cut in two cross-sectional partsusing a wire electric discharge machine (SL400G, Sodick). Thenthey were mounted and polished up to 4000-grit SiC papers and 1-mm diameter diamond suspension (#90-33015, Applied High TechProducts). Finally, they were etched with Kroll's reagent(H2O:HF:HNO3 ¼ 10:3:6 v/v%) to reveal grain structures at LPEB-irradiated regions.

2.2. Large pulsed electron beam and cathodic apparatus

The LPEB (PF32B, Sodick), shown in Fig. 1, consisted of four mainparts: solenoid, electron acceleration system, plasma source, and atwo-dimensional (2D) translation stage in a vacuum chamber [32].A standard set-up for LPEB irradiation generally uses argon (Ar) gasas a plasma source. For this study, the plasma source was changed

from Ar to nitrogen (N2) for LPEB nitriding process, which repre-sents the LPEB irradiation using N2 plasma gas. Furthermore, anegative DC bias cathodic apparatus was attached to the generalexperimental set-up to attract N2 plasma from environment to thesubstrates. The Tie6Ale7Nb sample was connected to a copperelectrode to apply a negative DC bias as summarized in Fig. 1. Theirradiation frequency of LPEB was 0.1 Hz with 2 ms of pulse durationand 10 s of dwell time. Although the total processing time is relatedwith the number of irradiation pulses, it is relatively rapid processthat the processing time is mostly less than 5 min per sampleincluding vacuuming time due to the short pulse duration and largebeam size with 60 mm in diameter. The detailed parameters ofLPEB irradiation are summarized in Table 1.

2.3. Microstructures characterizations

Before and after LPEB irradiations, the microstructural changesof Tie6Ale7Nb samples were evaluated in terms of each combi-nation of parameters via field emission scanning electron micro-scopy (SEM). Themicrostructural analyses on the surface and cross-section after LPEB irradiations were performed using a FEI Nano-230 Nova NanoSEM equipped with energy dispersive X-ray spec-trometer (EDS).

Surface roughness (Ra) variations were observed following LPEBirradiations using three-dimensional white-interferometer (NV-3000, Nanosystem). The measurements were carried out five times

Fig. 2. Scanning electron microscopy (SEM) images on Tie6Ale7Nb surfaces (a) before and after 10 pulses of LPEB irradiations with (b) 0 V-biased in Ar plasma gas, (c) 0 V-biased inN2 plasma gas, and (d) �100 V-biased in N2 plasma gas.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148140

at different regions on the same sample and Ra was averaged.

2.4. Phase analyses and chemical characterizations

The EDS analyses of elements, including C, N, O, and Ti, on thetop surface were carried out with a 20-kV acceleration voltage.Chemical compositions and states of LPEB-irradiated layer wereanalyzed by X-ray photoelectron spectroscopy (XPS) with a 90�

takeoff angle and 0.45 eV energy step size. A spectroscopy of theelemental compositions were recorded by ESCALAB 250XI (ThermoFisher Scientific) using Al-Ka radiation. Because the top surface ofTie6Ale7Nb was possibly oxidized by exposure to atmosphericconditions after LPEB irradiations, nitriding effects of LPEB irradi-ations could not be exactly evaluated at the top surface. Therefore,not only the binding energies of the top surface but also those of 10-nm deep regions were detected by ion etching.

2.5. Mechanical properties characterization

Nanoindentation tests (NHT2, Anton Paar) were performed by acontinuous stiffness measurement technique. A three-sided pyra-midal Berkovich tip was used as an indentation tool. Nano-indentation was assessed at the top surface after LPEB irradiationsto specify the nano-hardness along depth by increasing indentationload from 10 mN to 500 mN. A step increasing load was set to10 mN and the maximum load at each step was sustained during10 s to obtain steady-state nano-hardness.

2.6. Electrochemical measurements

A standard three-electrode electrochemical cell was designedand manufactured using polytetrafluoroethylene (PTFE, (C2F4)n). A1 cm2 circular hole at the PTFE cell facilitated to expose only LPEB-irradiated surface to an electrolyte and normalized evaluation ofelectrochemical performances. A coiled platinum counter electrodeand a saturated calomel reference electrode (SCE, E0 ¼ 0.24 V vs.

saturated hydrogen electrode) were used for potentiodynamicpolarization tests, electrochemical impedance measurements, andchronoamperometry (CA) analyses. For all types of electrochemicalmeasurements, 1 wt% sodium chloride (NaCl) aqueous solutionwasadopted as the electrolyte. The electrochemical tests were per-formed with multi-channel electrochemical instrument (IVIUMn-STAT, IVIUM Technologies) controlled by IVIUMsoft software.

Prior to testing, the substrate surfaces were exposed to theelectrolyte for approximately 30 min to attain a steady-state open-circuit potential (OCP). Following stabilization, potentiodynamicpolarization tests were carried out with a scan rate of 1 mV/s over arange of �500 to þ1000 mV/SCE with respect to the OCP. The re-sults of three independent polarization tests were obtained forsamples treated by the same LPEB parameters, and the represen-tative polarization curves were plotted. Polarization electro-chemical parameters, including corrosion potential (Ecorr),corrosion current density (icorr), and cathodic and anodic Tafelslopes (bc and ba) were averaged from the three independent tests.

The electrochemical impedance spectroscopy (EIS) was ob-tained in the frequency range from 2500 Hz to 0.03 Hzwith a 10mVvs. OCP level of alternating current (AC) amplitude. The EIS datawere investigated through Nyquist plots, Bode plots, and hypo-thetical impedance parameters specified from two-time-constantequivalent-circuit model.

The CA analyses were conducted to subsidiary compare the rateof corrosion progressed in a 1 wt% NaCl solution before and afterLPEB irradiations with various irradiating conditions. A constantpotential, þ100 mV/SCE, was applied to adopt the potential atsurely active region of polarization. The tests were conducted usinghigh-resolution electrochemical instrument (EZstat-Pro, NuVantSystems Inc.) with current density resolution of 3 nA/cm2. The datawere transiently recorded every 100 ms during 30 min.

Fig. 3. (a, b, c, d) White-interference micrographs on the surface of Tie6Ale7Nb after LPEB irradiations, and (e) a SEM image and EDS result near the crater-generated regioncorresponding to (d).

Fig. 4. Cross-sectional SEM images of LPEB-treated Tie6Ale7Nb samples with (a) Ar plasma gas and a single pulse, (b) N2 plasma gas and a single pulse, (c) Ar plasma gas and 10pulses, and (d) N2 plasma gas and 10 pulses.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148 141

Table 2Energy dispersive X-ray spectroscopy results of Tie6Ale7Nb following LPEB irra-diation with 10 pulses in terms of negative DC bias.

Type of plasma gas Negative DC bias (V) Atomic concentration (at.%)

C N O Ti

Untreated 0 9.86 3.23 13.92 55.77Ar 0 5.42 5.87 8.69 63.48N2 0 3.81 17.08 3.84 64.75N2 �100 3.51 18.27 3.26 63.68N2 �500 5.35 18.45 4.17 60.70N2 �1000 5.00 18.88 3.09 62.90

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148142

3. Results and discussions

3.1. Surface morphology and phase analyses

Fig. 2 shows the SEM images and corresponding roughness atthe surface following LPEB irradiation using Ar and N2 plasma gas.The surface roughness (Ra) was reduced significantly followingLPEB irradiation with both Ar and N2 plasma gas. The lowest valueof Ra was obtained from 10 pulses of LPEB irradiations without anegative DC bias; Ra was reduced from ~540 nm to ~90 nm on theLPEB irradiated surface both with Ar and N2 plasma gas comparingto the bare surface. This may be attributable to surface dissolutionby LPEB, which could melt the polishing mark and generate a phasetransformation on the surface [37]. However, when applying anegative DC bias for the nitriding process, Ra on the surface wasslightly larger than that on the surface with no DC bias becausenumerous crater-like defects were generated following LPEB irra-diation as specified by white-interference micrographs in Fig. 3a-3d. Especially, with a negative DC bias larger than �100 V, thedefects were formed frequently and their sizes were larger. Thepartial evaporation of non-metallic inclusions, such as manganesesulfide (MnS) and metal carbide, is well known to be the majorreason for crater generation during the LPEB irradiation process[38]. However, in the case of Tie6Ale7Nb, there are almost no non-metallic inclusions, suggesting that the mechanism of crater gen-eration differ from iron-based alloys. Thus, SEM images and cor-responding EDS analyses were performed to reveal the cratergeneration mechanism during the LPEB nitriding process (Fig. 3e).The EDS results indicated that small particles consisted of iron atthe center of the craters. Thus, it can be concluded that theincreasing density of craters with increasing negative DC biasresulted mainly from sputtering effects on the surface by gener-ating a glow discharge of plasma ions remaining in the vacuumchamber. Thus, it is preferred that the DC bias level should notexceed �100 V, because this is relatively low compared with thethreshold voltage.

Fig. 4 shows cross-sectional SEM images of Tie6Ale7Nb alloysfollowing LPEB irradiationwith Ar and N2 plasma gas. With a singlepulse of LPEB irradiation, the depth of the re-solidified layer was4.66 mm with Ar plasma gas; this decreased to 2.61 mm using N2

Table 3Energy dispersive X-ray spectroscopy results of Tie6Ale7Nb after LPEB irradiation as a f

Type of plasma gas Negative DC bias (V) Pulses (shots)

N2 �100 10N2 �100 20N2 �100 30N2 �100 50N2 �100 100N2 �100 200

plasma gas during LPEB irradiation. The depth of the re-solidifiedlayer increased with more pulses; however, it did not vary mark-edly above 10 pulses. The depth of the re-solidified layer induced byLPEB irradiationwas deeper with Ar (Fig. 4a and 4c) than N2 plasmagas (Fig. 4b and 4d). This could be a result of a change in beamenergy density transferred to the surface of the substrates. Thebeam energy induced from the electron gun is partially absorbed byplasma gas ionizing atoms. The ionization energy of each plasmasource used in LPEB irradiation is expressed as:

Ar þ ð1;520:6kJ=molÞ/Arþ þ e� (1)

12N2 þ ð1;874:5kJ=molÞ/Nþ þ e� (2)

As shown in Eqs. (1) and (2), Ar plasma is induced simply, onlyconsuming the ionized enthalpy of 1520.6 kJ/mol in a single ioni-zation step. In contrast, N2 plasma is generated in multiple steps tomake the same number of electrons by consuming a total enthalpyof 1874.5 kJ/mol to break the powerful triple bond of the nitrogenmolecule (472.5 kJ/mol) and to ionize (1402 kJ/mol). Thus, energyabsorbed on the surface of substrates is relatively smaller with N2plasma gas, leading to a thinner depth of the re-solidified layer thanwith Ar.

3.2. Chemical component

In the case of nitriding processes, the fraction of nitrogen plays amajor role in evaluating performance [31]. The EDS results whichindicate variations in atomic concentrations following LPEB irra-diation are summarized in Tables 2 and 3, in terms of the negativeDC bias and number of pulses, respectively. Corresponding varia-tions in nitrogen concentration are shown in Fig. 5. Compared withbare Tie6Ale7Nb, the atomic concentration of nitrogen wasincreased slightly after LPEB irradiation with Ar plasma gas(Fig. 5a). This could be a result of the flushing gas, which was usedto clean up the vacuum chamber between pulses [36]. Becausenitrogenwas used as the flushing gas during LPEB irradiation, someof the gas remained in the chamber, forming a small fraction of TiNat the re-solidified layer. By changing the plasma source for LPEBfrom Ar to N2, it was possible to achieve a much higher fraction ofnitrogen in the re-solidified layer. LPEB nitriding without a negativeDC bias increased the nitrogen fraction, to 17.08 at.%, at the topsurface. Moreover, the nitrogen fraction was increased further byapplying a negative DC bias to the substrate. It was increased toover 18 at.% with a negative DC bias ranging from 0 to�1000 V. Theoptimized number of pulses in terms of nitrogen fraction was wellmatched to that of the re-solidified depth. The atomic concentra-tion at the re-solidified layer was almost unaffected by increasingthe number of pulses above 10 (Fig. 5b). The effect of negative DCbias was clear; the nitrogen fraction was increased with �100 V ofbias compared with 0 V. However, no significant change in nitrogenfraction was observed with various DC biases above �100 V; all

unction of the number of pulses.

Atomic concentration (at.%)

C N O Ti

3.51 18.27 3.26 63.685.31 17.80 3.84 61.535.95 18.12 3.85 61.844.87 17.96 3.21 63.465.63 18.36 3.27 64.037.07 17.43 3.93 63.02

Fig. 5. Atomic concentration of nitrogen after LPEB irradiation in terms of the (a)negative DC bias and (b) number of pulses.

Fig. 6. X-ray photoelectron spectroscopy (XPS) spectra of Tie6Ale7Nb after LPEBnitriding with the negative DC bias in the range of binding energy corresponding to (a)Ti-2p, (b) N-1s, and (c) O-1s.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148 143

substrates were similar, ~18.5 at.%. Thair et al. [39] showed that TiNwas not formed if the nitrogen concentration was lower than~6 at.%; otherwise, TiN was formed uniformly at the surface if ni-trogen concentration was ~20 at.%. Thus, we may conclude that aneffective nitriding process was achieved using LPEB irradiationwithN2 plasma gas and a negative DC bias. Also, �100 V of DC bias wasthe optimal condition because larger biases did not modify thenitrogen concentration significantly and induced the small crater-like defects shown in Fig. 3c and 3d.

In addition to EDS analyses, XPS results specified the formationof TiN in the re-solidified layer after LPEB nitriding. The XPS resultsare presented in Fig. 6. The spectral data were obtained at the topsurface and 10 nm-deep regions with 0, �100, �500, and �1000 Vbias in N2 gas. It is clear that TiN and TiOxNy were formed suc-cessfully on Tie6Ale7Nb after LPEB nitriding. In Fig. 6a, the bindingenergy of the Ti-2p3/2 peak or shoulder was shifted from the458e460 eV range to 453e456 eV. Also, the Ti-2p1/2 peak wasmoved from the 464e465 eV range to 459e461 eV, with a higheroverall intensity of the spectrum. These XPS results showed simi-larities with other nitriding researches [40e42]. It has been docu-mented that TiN and TiOxNy corresponding to the Ti-2p3/2 peakappeared at the ~454.5 [40] and the Ti-2p1/2 peak indicating TiN

was measured at ~460.85 eV [42]. Thus, the shift of the Ti-2pspectrum suggested that TiN was formed inside the re-solidified

Fig. 7. Variations in nano-hardness at re-solidified layer as a function of depth.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148144

layer, and a compound layer, composed mainly of TiOx, was formed

Fig. 8. White-interference micrographs of wear tracks after a ball-on-disc wear test on Tie60 V-biased in N2 gas, (d) �100 V-biased in N2 gas, and corresponding (e) friction coefficien

at the top surface. Additionally, the spectrum of N-1s in Fig. 6bshowed the extent of TiOxNy at the top surface and TiN at 10 nmbelow the top surface [43]. These results werewell agreed alsowiththe spectrum of Ti-2p and O-1s, as shown in Fig. 6a and 6c. Toconclude, LPEB nitriding could facilitate nitrogen implantation inthe re-solidified layers and the formation of TiOxNy layers on thenitrided layers.

3.3. Mechanical properties

TiN is known to have much higher hardness than conventionaltitanium alloys, such as Tie6Ale4V and Tie6Ale7Nb. Moreover,surface hardening effect of LPEB irradiations has been firmlyestablished in previous researches [44]. Thus, comparisons of nano-hardness profiles among bare and LPEB-treated Tie6Ale7Nb withAr and N2 plasma gas are appropriate to specify the surface hard-ening and extent of the nitrogen diffusion into the substrates [23].Fig. 7 shows variations in nano-hardness as a function of the depthprofile after LPEB irradiation with Ar and N2 plasma gas. The nano-hardness near the top surface was increased slightly following LPEBirradiation with Ar plasma gas. At the top surface, it was modifiedby ~15%, and the depth of modification was observed to be

Ale7Nb surfaces (a) before and after LPEB irradiation with (b) 0 V-biased in Ar gas, (c)ts, (f) widths and depths of wear tracks measured at (a)e(d).

Fig. 9. (a) Potentiodynamic polarization curves of LPEB-irradiated Tie6Ale7Nb and (b)corresponding corrosion potentials and corrosion current densities.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148 145

~1500 nm. Using N2 plasma gas during LPEB irradiationwithout DCbias, a further increase in nano-hardness was observed at the topsurface. LPEB irradiation with N2 plasma gas with DC bias couldsignificantly modify nano-hardness, as shown in Fig. 7. Indeed, itwas modified ~75% at the top surface (~5.7 GPa) comparing to thebare surface (~3.25 GPa). Although there was no major difference innano-hardness depending on the DC bias, themost improved nano-hardness was obtained with �100 V.

In addition to the increasing fraction of TiN, microstructuralchanges of Tie6Ale7Nb itself following LPEB irradiation possiblyaffect the modification of nano-hardness. The formation of TiN atthe re-solidified layer generates a phase transformation during theprocess because the nitrogen atoms are implanted interstitiallybetween titanium atoms in the re-solidified layer. Thus, latticestructures at the re-solidified layer can become distorted, resultingin surface dislocations, which are well-established hardeningmechanisms of metal alloys [45]. Consequently, the nano-hardness

Table 4Polarization electrochemical parameters of untreated and LPEB-treated Tie6Ale7Nb dep

Type of plasma gas Negative DC bias (V) Ecorr (mV/SCE)

Untreated 0 �404.7 ± 5.58Ar 0 �303.5 ± 5.25N2 0 �243.0 ± 9.45N2 �100 43.2 ± 8.00

of Tie6Ale7Nb could be effectively modified by LPEB irradiationforming TiN at the re-solidified layer. The modification on me-chanical properties at the re-solidified layers was alsowell reflectedin ball-on-disc wear tests. Fig. 8 shows white-interference micro-graphs, friction coefficients, and widths and depths of wear tracksafter the wear tests. Although the friction coefficients were similarfor the bare and LPEB-treated samples regardless of LPEB param-eters (Fig. 8e), configurations of wear tracks were significantlydifferent. The lowest worn depth of ~8 mm and width of ~600 mmwere obtained from the LPEB-nitrided Tie6Ale7Nb biasedwith �100 V among the four samples. The greater wear resistanceof re-solidified layer following LPEB irradiations could be associatedwith its higher surface hardness; the decreasing tendency of weartrack width and depth followed the increasing tendency of nano-hardness depending on LPEB parameters.

3.4. Corrosion performance

From the phase, microstructural, and nano-hardness analyses, itwas clear that 10 irradiation pulses and a �100 V DC bias were theoptimal conditions because the fraction of nitrogen did not varywith increasing number of pulses to more than 10 or a DC biasvoltage larger than�100 V. Thus, the corrosion characteristics wereevaluated with a sample irradiated through 10 LPEB pulses anda �100 V DC bias. Fig. 9 shows representative potentiodynamicpolarization curves of the untreated and LPEB-treated Tie6Ale7Nballoys and corresponding variations of corrosion potentials andcorrosion current densities in a 1 wt% NaCl aqueous solution atroom temperature. The electrochemical parameters estimated fromthe Tafel extrapolation, including corrosion potential, corrosioncurrent density, and Tafel slopes, are summarized in Table 4. Thecorrosion potential became nobler following LPEB irradiation; itincreased from �404.7 mV/SCE on the bare surface to �303.5 mV/SCE on the LPEB-treated surface with Ar plasma gas. The corrosionpotential on the LPEB-treated surface with N2 plasma gas wasnobler than that on the LPEB-treated surface with Ar plasma gas.Moreover, it was possible to achieve a positive level of corrosionpotential following LPEB nitriding on Tie6Ale7Nb using a biasof �100 V. Thus, we can summarize that the corrosion resistancewas increased by LPEB irradiation and further increased by LPEBnitriding from the nobler corrosion potential, as summarized inFig. 9b. In addition to the corrosion potential, the changes incorrosion current densities also indicated an improvement in thecorrosion resistance following LPEB irradiation with Ar and N2plasma gas. Comparing to the bare surface (194.9 nA/cm2), muchlower corrosion current density was obtained on LPEB-irradiatedsurface with Ar plasma gas, N2 plasma gas, and negative DC biasof �100 V (<~50 nA/cm2). The corrosion rate corresponding tocorrosion current density has been studied previously. The rela-tionship between corrosion current density and corrosion rate canbe expressed, as follows, using Faraday's law [46]:

n ¼ MzFr

icorr (3)

where M is the molar mass for Tie6Ale7Nb, icorr is the corrosion

ending on the type of plasma gas and negative DC bias.

icorr (nA/cm2) ba (mV/dec) bc (mV/dec)

194.9 ± 5.04 971 ± 20.6 �372 ± 10.549.9 ± 4.60 1254 ± 22.9 �602 ± 22.945.4 ± 2.41 1450 ± 40.0 �440 ± 22.051.6 ± 5.01 943 ± 11.0 �790 ± 16.0

Fig. 10. (a) Nyquist plots and (b, c) Bode plots of Tie6Ale7Nb before and after LPEB irradiations, and the equivalent circuit model corresponding to (d) bare and (e) LPEB-treatedsurfaces.

Table 5Electrical parameters specified from the fitted lines of EIS spectra in Fig. 10a-10c.

Type of plasma gas Negative DC bias (V) Rs (kU cm2) CPEc (mF/cm2) nc Rp (kU cm2) CPEdl (mF/cm2) ndl Rct (kU cm2)

Untreated 0 0.1611 8.228 0.9296 0.1683 6.963 0.9078 34.40Ar 0 0.1626 2.890 0.8984 0.5455 1.020 0.9837 2013N2 0 0.1610 0.3827 0.8096 0.3792 0.3790 0.9887 5327N2 �100 0.1657 0.8379 0.8849 0.5559 0.2449 0.9515 6194

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148146

current density, z denotes the electron number, F is the Faradayconstant, and r is the density. It is clear that corrosion rate, which isthe corrosion propagation speed once it occurs, is linearly propor-tional to the corrosion current density. Consequently, it can beconcluded that the corrosion rate is decreased significantly by LPEBirradiation on Tie6Ale7Nb both with Ar and N2 plasma gas.

This modification of corrosion resistance following LPEB irra-diation and LPEB nitriding is strongly related to the formation of astable and passive re-solidified layer. As shown in the polarizationcurves (Fig. 9a), a slight active-passive transition region at ~100mV/SCE was observed on the bare surface of Tie6Ale7Nb. How-ever, spontaneously passive characteristics in the anodic environ-ment were obtained on LPEB-treated samples with both Ar and N2

plasma gas. It is also known that oxide layers formed at the surface,consisting mainly of TiO2, are responsible for the high corrosionresistance of titanium-based alloys [47]. A typical titanium oxidelayer that forms at the surface of titanium alloys is generally weakand thin; however, more stable and firmly established oxide layerscan be achieved following LPEB irradiation.

The corrosion-inhibiting effects of these passivation, oxide, andre-solidified layers can be characterized by the charge transferresistance. The charge transfer resistance, obtained from the EISspectra, can be analyzed simply using Nyquist plots, Bode plots, and

corresponding equivalent circuit modeling. Fig. 10aec showsNyquist plots and Bode plots of Tie6Ale7Nb before and after LPEBirradiation with/without DC bias. It is clear that LPEB irradiationwith both Ar and N2 plasma gas enhanced the impedance of thesurface by forming a stable and passive re-solidified layer combinedwith oxygen. All of the LPEB-treated surfaces of Tie6Ale7Nb indi-cated a larger radius on Nyquist plots compared with that on thebare surface. The largest radius of the semi-circle on Nyquist plotswas obtained with the sample treated by LPEB with N2 plasma gasand �100 V of DC bias. Comparing the samples treated by LPEBwith Ar plasma gas and N2 plasma gas, a higher impedance at highfrequencies was observed at the surface of the sample treated withAr plasma gas; however, this was reversed at low frequencies, sothat the impedancemeasured at frequencies lower than ~10 Hzwaslarger on the LPEB-treated sample with N2 plasma gas than Arplasma gas (Fig. 10b). Thus, it was well supported by the Nyquistand Bode plots that corrosion resistance on the LPEB-treated sur-face was hugely increased, and the most improved corrosionresistance was obtained by LPEB irradiationwith N2 plasma gas anda DC bias.

For further discussions on the corrosion characteristics at thesurface of Tie6Ale7Nb before and after LPEB irradiation, equivalentelectrical circuit models corresponding to the corrosion

Fig. 11. Chronoamperometry of Tie6Ale7Nb before and after LPEB irradiations.

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148 147

mechanisms of LPEB-treated Tie6Ale7Nb are proposed, as illus-trated in Fig. 10d and e. The corresponding electrical parametersinterpreted and fitted from EIS data are summarized in Table 5. Theequivalent circuit models consist of the solution resistance (Rs), theconstant phase elements indicating capacitance of passive oxidefilms (CPEc), the reaction resistance considering the semi-conducting properties of defects in the passive film (Rp), the chargetransfer resistance of substrates (Rct), and the double layer capaci-tance (CPEdl). The analysis results using the proposed two-timeconstant equivalent circuit model matched well with the varia-tions in corrosion resistance indicated by the polarization curvesand EIS spectra. The Rct for bare Tie6Ale7Nb was 34.4 kU/cm2,which was much lower than that of the LPEB-treated samples. Incontrast, LPEB irradiation with Ar plasma gas could lead to theformation of a stable re-solidified layer on the surface. Thus, thecorresponding Rct at the re-solidified layer was hugely increased to2013 kU/cm2 with Ar gas, and further increased to 5327 kU/cm2

with N2 gas. The highest value of Rct, 6194 kU/cm2, was obtainedfrom the LPEB-nitrided Tie6Ale7Nb sample with �100 V bias.Although LPEB irradiations with Ar gas modifies Rct by generatingelectrochemically stable re-solidified layer, further modification ofRct through LPEB nitriding could be a result of the increasing frac-tion of TiN formed on the re-solidified layer by changing the plasmagas from argon to nitrogen. A higher TiN fraction was achieved byusing N2 plasma gas versus Ar plasma gas. Rossi et al. [48] revealedthat the chemical stability of TiN itself leads to the highly corrosion-resistant characteristic of TiN, as assessed by an analysis of

corrosion performance in extremely aggressive electrolytes. Thus,the most modified corrosion resistance analyzed from the polari-zation curves and EIS spectra could be supported by the modifiedRct on LPEB-nitrided samples with or without DC bias.

In order to clearly specify the effect of corrosion-resistant re-solidified layer induced by LPEB nitriding on corrosion reactions, CAcurves were analyzed potentiostatically at a certain anodicpotential, þ100 mV/SCE. Fig. 11 shows the CA curves on Tie6A-le7Nb samples before and after LPEB irradiations. The currentdensities measured from samples were increased dramatically in afew seconds immediately after the anodic potential applied. Thisincrease of current density is mainly due to the dissolution of apassive oxide film [49]. After the momentary increase of currentdensity, the curves soon decreased continuously till the oxide filmto be fully dissolved. Finally the current density could be saturatedat a certain level indicating continuous corrosion reaction of thebare Tie6Ale7Nb and re-solidified layer. As shown in Fig. 11, thehighest level of current density was obtained on the bare surfaceindicating the most reactive characteristic to corrosion, and it wasnot saturated till 30 min testing. The current densities obtainedfrom LPEB-treated samples were significantly lower than that fromthe bare surface. The LPEB-treated sample with Ar plasma gasindicated 40 nA/cm2 of current density at the first, and it wassaturated at ~10 nA/cm2 after 30 min. The LPEB-nitrided samplesboth with and without DC bias showed the lowest current density;it was initially increased up to ~30 nA/cm2, and then saturated atnearly zero after ~200 s indicating no corrosion reactions. Thus CAanalyses confirmed that the re-solidified layer induced by LPEBnitriding has superior corrosion resistance.

4. Conclusions

The LPEB nitriding process, describing LPEB irradiation using N2plasma gas and cathodic apparatus to apply a negative DC bias, hasbeen introduced as a unique, effective, and crack-free nitridingmethod for Tie6Ale7Nb alloys. Most importantly, increasing frac-tion of TiN on re-solidified layers following LPEB nitriding facili-tated to fabricate the surface layer with high corrosion resistanceandmechanical properties. Although only small fraction (~6 at.%) ofnitrogen was implanted by conventional LPEB irradiations with Argas, the LPEB nitriding increased the fraction of implanted nitrogenat the re-solidified layer up to ~18 at.% with �100 V of bias. Ahardened surface by distorting the lattice structure due to inter-stitially implanted nitrogen atoms, forming TiN in the re-solidifiedlayer, led improved nano-hardness and wear resistance. Further-more, the corrosion resistance of Tie6Ale7Nb was modified byconstructing a passive and stable re-solidified layer consisted ofTiN, TiO2, and TiOxNy. The chronoamperometry measurementsrevealed that no reactions occur at the LPEB-nitrided surface astheir current densities were saturated at nearly zero after dissolu-tion of oxide layers. Hence, the LPEB nitriding could offer promisingprospects for industrial applications as a single-step finishingprocess which could simultaneously modify mechanical propertiesand corrosion resistance.

Acknowledgment

This work was supported by the Mid-Career Researcher Pro-gram through the National Research Foundation of Korea funded bythe Ministry of Education (No. 2015R1A2A2A01005499) andDevelopment of the High Speed Ecological Finishing Process forprecision and micro pattern products (No. 20100110038656).

J. Kim et al. / Journal of Alloys and Compounds 679 (2016) 138e148148

References

[1] C. Castellani, R.A. Lindtner, P. Hausbrandt, E. Tschegg, S.E. Stanzl-Tschegg,G. Zanoni, S. Beck, A.-M. Weinberg, Boneeimplant interface strength andosseointegration: biodegradable magnesium alloy versus standard titaniumcontrol, Acta Biomater. 7 (2011) 432e440.

[2] M. Long, H.J. Rack, Titanium alloys in total joint replacementda materialsscience perspective, Biomaterials 19 (1998) 1621e1639.

[3] T. Niemeyer, C.R. Grandini, L. Pinto, A. Angelo, S. Schneider, Corrosionbehavior of Tie13Nbe13Zr alloy used as a biomaterial, J. Alloys Compd. 476(2009) 172e175.

[4] I. Cvijovi�c-Alagi�c, Z. Cvijovi�c, S. Mitrovi�c, V. Pani�c, M. Rakin, Wear and corro-sion behaviour of Tie13Nbe13Zr and Tie6Ale4V alloys in simulated physi-ological solution, Corros. Sci. 53 (2011) 796e808.

[5] M. Dimah, F.D. Albeza, V.A. Borr�as, A.I. Mu~noz, Study of the biotribocorrosionbehaviour of titanium biomedical alloys in simulated body fluids by electro-chemical techniques, Wear 294 (2012) 409e418.

[6] E. Blasco-Tamarit, A. Igual-Mu~noz, J.G. Ant�on, D. García-García, Corrosionbehaviour and galvanic coupling of titanium and welded titanium in LiBrsolutions, Corros. Sci. 49 (2007) 1000e1026.

[7] A. Montiel-D�avalos, A. Gonzalez-Villava, V. Rodriguez-Lara, L.F. Monta~no,T.I. Fortoul, R. L�opez-Marure, Vanadium pentoxide induces activation anddeath of endothelial cells, J. Appl. Toxicol. 32 (2012) 26e33.

[8] N.T. Oliveira, E.A. Ferreira, L.T. Duarte, S.R. Biaggio, R.C. Rocha-Filho, N. Bocchi,Corrosion resistance of anodic oxides on the Tie50Zr and Tie13Nbe13Zr al-loys, Electrochim. Acta 51 (2006) 2068e2075.

[9] M. Khan, R. Williams, D. Williams, The corrosion behaviour of Tie6Ale4V,Tie6Ale7Nb and Tie13Nbe13Zr in protein solutions,, Biomaterials 20 (1999)631e637.

[10] V. Challa, S. Mali, R. Misra, Reduced toxicity and superior cellular response ofpreosteoblasts to Ti-6Al-7Nb alloy and comparison with Ti-6Al-4V, J. Biomed.Mater. Res. A 101 (2013) 2083e2089.

[11] J. Liu, A. Alfantazi, E. Asselin, The anodic passivity of titanium in mixed sulfate-chloride solutions,, J. Electrochem. Soc. 162 (2015) E289eE295.

[12] C.-F. Huang, H.-C. Cheng, C.-M. Liu, C.-C. Chen, K.-L. Ou, Microstructure andphase transition of biocompatible titanium oxide film on titanium by plasmadischarging, J. Alloys Compd. 476 (2009) 683e688.

[13] F. Cheng, P. Shi, G. Pang, M. Wong, H. Man, Microstructural characterization ofoxide film formed on NiTi by anodization in acetic acid, J. Alloys Compd. 438(2007) 238e242.

[14] Z. Liu, X. Zhong, J. Walton, G. Thompson, Anodic film growth of titanium oxideusing the 3-electrode electrochemical technique: effects of oxygen evolutionand morphological characterizations, J. Electrochem. Soc. 163 (2016)E75eE82.

[15] L. Mohan, C. Anandan, N. Rajendran, Electrochemical behavior and effect ofheat treatment on morphology, crystalline structure of self-organized TiO 2nanotube arrays on Tie6Ale7Nb for biomedical applications, Mater. Sci. Eng.C 50 (2015) 394e401.

[16] L. Benea, E. Mardare-Danaila, M. Mardare, J.-P. Celis, Preparation of titaniumoxide and hydroxyapatite on Tie6Ale4V alloy surface and electrochemicalbehaviour in bio-simulated fluid solution, Corros. Sci. 80 (2014) 331e338.

[17] G.-L. Song, The grand challenges in electrochemical corrosion research, Front.Mater. 1 (2014) 1e3.

[18] C.-D. Gu, J.-L. Zhang, W.-Q. Bai, Y.-Y. Tong, X.-L. Wang, J.-P. Tu, Electro-brushplating from deep eutectic solvent: a case of nanocrystalline Ni coatings withsuperior mechanical property and corrosion resistance, J. Electrochem. Soc.162 (2015) D159eD165.

[19] L. Mohan, C. Anandan, Effect of gas composition on corrosion behavior andgrowth of apatite on plasma nitrided titanium alloy Beta-21S, Appl. Surf. Sci.268 (2013) 288e296.

[20] V. Oliveira, C. Aguiar, A. Vazquez, A. Robin, M. Barboza, Improving corrosionresistance of Tie6Ale4V alloy through plasma-assisted PVD deposited nitridecoatings, Corros. Sci. 88 (2014) 317e327.

[21] E. Zalnezhad, A.A. Sarhan, M. Hamdi, Investigating the fretting fatigue life ofthin film titanium nitride coated aerospace Al7075-T6 alloy, Mater. Sci. Eng. A559 (2013) 436e446.

[22] T. Zakroczymski, N. Lukomski, J. Flis, Entry and transport of hydrogen in ionnitrided iron, J. Electrochem. Soc. 140 (1993) 3578e3583.

[23] T.M. Manhabosco, S. Tamborim, C. Dos Santos, I. Müller, Tribological, elec-trochemical and tribo-electrochemical characterization of bare and nitridedTi6Al4V in simulated body fluid solution, Corros. Sci. 53 (2011) 1786e1793.

[24] A. Zhecheva, S. Malinov, W. Sha, Titanium alloys after surface gas nitriding,Surf. Coat. Technol. 201 (2006) 2467e2474.

[25] H. Li, Z. Cui, Z. Li, S. Zhu, X. Yang, Surface modification by gas nitriding forimproving cavitation erosion resistance of CP-Ti, Appl. Surf. Sci. 298 (2014)164e170.

[26] Z.R. Ismagilov, N.V. Shikinaa, E.V. Bessudnovaa, D.V. Korneevc,A.V. Ishchenkoa, Y.A. Chesalova, A. Vladimirovab, E.I. Ryabchikovab, The effectof chemical treatment conditions of titanium dioxide sols on their dispersionand cytotoxic properties,, Chem. Eng. 27 (2012) 1e6.

[27] J.A.-B. Wilson, S. Banfield, J. Housden, C. Olivero, P. Chapon, On the response ofTie6Ale4V and Tie6Ale7Nb alloys to a Nitron-100 treatment, Surf. Coat.Technol. 260 (2014) 335e346.

[28] T. Borkar, S. Nag, Y. Ren, J. Tiley, R. Banerjee, Reactive spark plasma sintering(SPS) of nitride reinforced titanium alloy composites, J. Alloys Compd. 617(2014) 933e945.

[29] I. Pohrelyuk, V. Fedirko, O. Tkachuk, R. Proskurnyak, Corrosion resistance ofTie6Ale4V alloy with nitride coatings in Ringer's solution, Corros. Sci. 66(2013) 392e398.

[30] S. Sathish, M. Geetha, N. Pandey, C. Richard, R. Asokamani, Studies on thecorrosion and wear behavior of the laser nitrided biomedical titanium and itsalloys, Mater. Sci. Eng. C 30 (2010) 376e382.

[31] J. Abboud, Effect of processing parameters on titanium nitrided surface layersproduced by laser gas nitriding, Surf. Coat. Technol. 214 (2013) 19e29.

[32] Y. Uno, A. Okada, K. Uemura, P. Raharjo, S. Sano, Z. Yu, S. Mishima, A newpolishing method of metal mold with large-area electron beam irradiation,,J. Mater. Process. Technol. 187 (2007) 77e80.

[33] J. Murray, P. Kinnell, A. Cannon, B. Bailey, A. Clare, Surface finishing of intricatemetal mould structures by large-area electron beam irradiation, Precis. Eng.37 (2013) 443e450.

[34] Y. Samih, G. Marcos, N. Stein, N. Allain, E. Fleury, C. Dong, T. Grosdidier,Microstructure modifications and associated hardness and corrosion im-provements in the AISI 420 martensitic stainless steel treated by high currentpulsed electron beam (HCPEB), Surf. Coat. Tehcnol. 259 (2014) 737e745.

[35] Y. Su, G. Li, L. Niu, S. Yang, J. Cai, Q. Guan, Microstructure modifications andassociated corrosion improvements in GH4169 superalloy treated by highcurrent pulsed electron beam, J. Nanomater. 2015 (2015) 1e6.

[36] J. Kim, H.W. Park, Influence of a large pulsed electron beam (LPEB) on thecorrosion resistance of Ti� 6Al� 7Nb alloys, Corros. Sci. 90 (2015) 153e160.

[37] X. Zhang, S. Hao, T. Grosdidier, J. Zou, B. Gao, B. Bolle, N. Allain-Bonasso, Y. Qin,X. Li, C. Dong, Surface modification of light alloys by low-energy high-currentpulsed electron beam, J. Metall. 2012 (2012) 1e10.

[38] K. Zhang, J. Zou, T. Grosdidier, C. Dong, Formation and evolution of craters incarbon steels during low-energy high-current pulsed electron-beam treat-ment, J. Vac. Sci. Technol. A 27 (2009) 1217e1226.

[39] L. Thair, U.K. Mudali, S. Rajagopalan, R. Asokamani, B. Raj, Surface character-ization of passive film formed on nitrogen ion implanted Tie6Ale4V andTie6Ale7Nb alloys using SIMS, Corros. Sci. 45 (2003) 1951e1967.

[40] I. Bertoti, M. Mohai, J. Sullivan, S. Saied, Surface characterisation of plasma-nitrided titanium: an XPS study, Appl. Surf. Sci. 84 (1995) 357e371.

[41] M. Krawczyk, W. Lisowski, J.W. Sobczak, A. Kosi�nski, A. Jablonski, Studies ofthe hot-pressed TiN material by electron spectroscopies, J. Alloys Compd. 546(2013) 280e285.

[42] L. Assaud, K. Pitzschel, M. Hanbücken, L. Santinacci, Highly-conformal TiN thinfilms grown by thermal and plasma-enhanced atomic layer deposition, ECS J.Solid State Sci. 3 (2014) P253eP258.

[43] D. Jaeger, J. Patscheider, A complete and self-consistent evaluation of XPSspectra of TiN, J. Electron. Spectrosc. Relat. Phenom. 185 (2012) 523e534.

[44] J. Zou, K. Zhang, S. Hao, C. Dong, T. Grosdidier, Mechanisms of hardening, wearand corrosion improvement of 316L stainless steel by low energy high currentpulsed electron beam surface treatment, Thin Solid Films 519 (2010)1404e1415.

[45] A. Pogrebnjak, V. Ladysev, N. Pogrebnjak, A. Michaliov, V. Shablya, A. Valyaev,A. Valyaev, V. Loboda, A comparison of radiation damage and mechanical andtribological properties of a-Fe exposed to intense pulsed electron and ionbeams, Vacuum 58 (2000) 45e52.

[46] E. Sherif, S.-M. Park, Inhibition of copper corrosion in acidic pickling solutionsby N-phenyl-1, 4-phenylenediamine, Electrochim. Acta 51 (2006) 4665e4673.

[47] W.-q. Yu, J. Qiu, L. Xu, F.-q. Zhang, Corrosion behaviors of TiO2 nanotubelayers on titanium in Hank's solution, Biomed. Mater. 4 (2009) 065012.

[48] S. Rossi, L. Fedrizzi, T. Bacci, G. Pradelli, Corrosion behaviour of glow dischargenitrided titanium alloys, Corros. Sci. 45 (2003) 511e529.

[49] E.-S.M. Sherif, H.S. Abdo, K.A. Khalil, A.M. Nabawy, Corrosion properties insodium chloride solutions of AleTiC composites in situ synthesized by HFIHF,Metals 5 (2015) 1799e1811.