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the thickness of the nitride layer to mitigate wear, and the erties and novel wear-induced deformation structures.Mechanistic studies of the unworn and worn surfaces andsubsurfaces were conducted to fundamentally understandthe structural evolution during wear by using thetechniques of cross-sectional focused ion beam (FIB)

Corresponding author; e-mail: Scharf@unt.edu1These authors contributed equally to this work.

Available online at www.sciencedirect.com

ScienceDirectActa Materialia 83 (2015) 61Keywords: Titanium alloys; Nitriding; Wear; Transmission electron microscopy (TEM); Precession electron diraction (PED)

1. Introduction

Titanium (Ti) alloys have long been used in manydiverse applications ranging from orthopedic biomedicalimplants (femoral stems) to aerospace gas turbine enginecomponents (compressor blade roots) [14]. While Ti alloyssuch as the a + b alloy Ti6Al4V (Ti64) and the b alloyTi35Nb7Zr5Ta (TNZT) are well known for their corro-sion resistance, adequate fracture toughness and fatiguestrength, they exhibit high wear rates, resulting in galling(adhesive) failure and fretting fatigue [48]. Thus, theyrequire surface modication treatments to improve theirwear behavior [412]. One such approach is surface nitrid-ing, which produces a high hardness and a chemically inertlayer that overcomes the poor wear resistance of Ti alloys.The process involves heat treating at elevated temperatures(7001100 C) in a nitrogen-enriched atmosphere orimplantation of nitrogen ions into the surface of Ti alloys[8,10,11,1318]. Surface nitriding is a fast, inexpensiveand simple treatment to form a hardened, wear-resistantcase layer on the surface. However, the penetration depthof nitrogen into the bulk alloy is limited, thus restricting

high processing temperatures can result in alloy softening.This paper focuses on an alternative approach that employslaser engineered net shaping (LENSe) with in situ forma-tion of a uniformly distributed wear-resistant nitrided layer.

Previously, the authors have directly introducedtitanium boride particles during LENSe processing of Tialloy powders to deposit in situ TiBTi composites [7,19].However, the quality of such an interface between exter-nally introduced reinforcements and the alloy matrix canbe rather dicult to control and results in TiB precipitatepull-out during interfacial sliding that leads to acceleratedthird body abrasive wear [7]. Thus, a reinforcement ideallycreated as a product of an in situ gas reaction oers theadvantage of a thermodynamically and mechanically stableinterface with the matrix. Furthermore, by employing anin situ nitride phase formation reaction, such as a nitro-gen-enriched hard a-Ti phase, a more uniform and homo-geneous distribution of the nitride phase can bepotentially achieved throughout the alloy matrix. To thisend, we have fabricated in situ nitride-reinforced TNZTand Ti64 alloys using LENSe in the presence of nitrogengas to study the solidication structure, sliding wear prop-In situ nitrided titanium alloys:solidicati

H. Mohseni,1 P. Nandwana,1 A.

Department of Materials Science and Engineering

Received 20 March 2014; revised 9 SeAvailable onli

AbstractSurface and subsurface structural evolution during sliding wa/b Ti6Al4V (Ti64), was studied by cross-sectional transmission electduring the laser engineered net shaping (LENSe) process resulted inenriched a phase in TNZT and Ti64, respectively. Subsurface structurand to greater extent b grains, to undergo severe plastic deformation, foring precessionorientation imaging phase maps were used to determine tmaps revealed that the nanocrystalline grains of soft/compliant b are mudue to the combination of the above phases and the increase in the aabsence of texture in the highly rened b grains. Conversely, nitridedchange in sliding-induced deformation mechanism where shear bands foparticle generation. 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights resehttp://dx.doi.org/10.1016/j.actamat.2014.09.0261359-6462/ 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rightsicrostructural evolution duringand wear

i, R. Banerjee and T.W. Scharf

versity of North Texas, Denton, TX 76207, USA

er 2014; accepted 12 September 2014October 2014

two in situ nitrided titanium alloys, b Ti35Nb7Zr5Ta (TNZT) andicroscopy coupled with precession electron diraction. In situ nitridingrmation of hard and wear-resistant Ti2N + TiN phases and nitrogen-alyses of the worn nitrided TNZT revealed the tendency of a grains,a heavily grain-rened nanocrystalline a and b tribolayer. Correspond-ientation and percentage of a and b-Ti in the worn nitrided TNZT. Thealler (10100 nm) than hard/sti a grains (>100 nm). Wear reduction isent of {0002}-textured coarser a grains along the sliding direction inexhibited slightly increasing wear, despite higher hardness, due to theand networked leading to brittle fracture and third body abrasive wear

74

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pre-calculated (simulated) diraction patterns using a

sliding distance (m).

Matecross-correlation technique to nd the best match for allpossible material phases and orientations. A major advan-tage of PED over traditional EBSD is the much ner spatialresolution of the TEM probe size that allows previouslyinaccessible phase orientation mapping of nanostructuralmaterials.

2. Experimental methods

A LENSe 750 system (Optomec, Inc.) with an Nd:YAGlaser at a wavelength of 1.064 lm and a power of 760 Wwas used to deposit blends of elemental powders under100% Ar and 100% N2 atmospheres for deposition of baseand nitrided (TNZT-N and Ti64-N) alloys, respectively.The TNZT and Ti64 alloy powders (manufactured byTIMET, Inc.) were mixed with a twin-roller ball mill inthe ratio of Ti35Nb7Zr11Ta and Ti6Al4V (wt.%).More details of the LENSe processing parameters arereported elsewhere [7,28]. As-deposited base and nitridedTNZT and Ti64 alloys were subsequently surface polishedusing 1200-grade SiC grinding paper followed by ne pol-ishing down to 200 nm arithmetic surface roughness with0.05 lm colloidal silica. X-ray diraction (Rigaku UltimaIII diractometer) with a Cu Ka source was used for phaseidentication. Surface microstructural characterization wascarried out in a eld emission gun scanning electronmicroscope (FEI Nova Nano SEM230).

Site selective, cross-sectional FIB-SEM studies wereperformed on the alloys with an FEI Nova 200 Nanolabdual-beam microscope. Cross-sectional lift-outs acquiredoutside and inside worn areas were then analyzed withbright-eld TEM (BFTEM) and energy-dispersive X-rayspectroscopy (EDS) line proles and maps. These analysesmicroscopy coupled with transmission electron microscopyand precession electron diraction (TEM-PED). Thesecombined techniques will resolve, for the rst time, grain(re)orientation and texture of wear-induced near-surfaceultrane grain (

martensitic in nature. In contrast, the BSE images of thenitrided Ti64 sample shown in Fig. 2c and d exhibit a verydierent microstructure. Comparing the higher-magnica-tion images for both alloys (Fig. 2b and d), there are twoprincipal dierences. First, Fig. 2d shows there is a morepronounced contrast between the a precipitates (darker)

H. Mohseni et al. / Acta Materialia 83 (2015) 6174 63referred to as a basketweave microstructure in a/b Ti alloys,containing multiple crystallographic variants of a intersect-ing within the same prior b grain. Such a basketweave a/bmicrostructure can be attributed to the high cooling ratesinvolved in LENSe deposition due to the substrate actingas an ecient heat sink. These ne scale a laths have beenattributed to martensitic decomposition of the b matrix[15]. However, in the present case, the BSE image shownin Fig. 2b clearly shows a signicant dierence in contrastbetween the darker a laths and the brighter b matrix, indic-ative of the partitioning of alloying elements between thesetwo phases. Hence it is unlikely that these ne scale a laths,observed in the LENSe deposited Ti64 sample, are purely

Fig. 1. XRD patterns obtained from LENS deposited (a) Ti64 andnitrided Ti64 (Ti64-N), and (b) TNZT and nitrided TNZT (TNZT-N)alloys.

Table 1. Lattice parameters and d spacings of signicant planes of a and b

Alloy d(1010)a (A) d(0002)a (A) aa (A)

Ti64 2.526 2.335 2.917Ti64-N 2.535 2.345 2.927TNZT TNZT-N 2.587 2.408 2.987and the b matrix (lighter) in nitrided Ti64, indicating moresubstantial atomic mass dierence between these twophases, presumably arising from the partitioning of V intothe b matrix, while the lighter Al and N partition into thea precipitates. Second, the morphology of a precipitates innitrided Ti64 is substantially dierent compared to baseTi64. The a precipitates in nitrided Ti64 are coarser, exhibita smaller aspect ratio, and in some cases appear to be moreequiaxed in shape. There also appear to be two distinctlydierent morphologies of a precipitates in nitrided Ti64(Fig. 2c and d): a smaller volume fraction of larger elon-gated needle-like or acicular precipitates exhibiting complexinternal contrast variations (referred to as Type 1 inFig. 2d); and smaller-scale precipitates exhibiting a loweraspect ratio (referred to as Type 2 in Fig. 2d).

Further investigation of these two types of precipitateswas carried out by mapping the elemental distribution withEDS. Fig. 3a is a BSE image of the region that was mappedand Fig. 3bd show the EDS maps of nitrogen, aluminumand vanadium, respectively. Interestingly, the nitrogen con-tent of the Type 1 needle-like precipitates is higher com-pared to the Type 2 precipitates. In contrast, theseneedle-like precipitates have a lower Al content comparedto the Type 2 precipitates. The two Type 1 precipitatesdenoted in Fig. 3a appear to have similar composition, thusindicating they are the same phase, while the dierence intheir morphologies may arise from the sectioning eectwhen viewed in the limited 2-D eld of view. It should benoted that both types of precipitates are coarser comparedto the ne scale a laths observed in the LENSe depositedbase Ti64 alloy (Fig. 2b). Therefore, it can be concludedthat in the case of the nitrided Ti64 sample, the microstruc-ture is not martensitic but rather involves diusional parti-tioning of the alloying elements and the introduction ofnitrogen into the a precipitates clearly changes their aspectratio and morphology, in agreement with previous reportsin the literature [28]. These nitrogen-enriched a precipitatesexhibiting a dierent morphology will subsequently bereferred to as a(Ti,N) precipitates in this paper. Therefore,LENSe processing in a nitrogen atmosphere has a pro-nounced eect on the stabilization of the a(Ti,N) solid-solution phase, in agreement with the XRD results. Anenhancement in the precipitation of a phase has also beenobserved on the surface of gas nitrided Ti64 [18]. The vana-dium EDS map in Fig. 3d shows the surrounding b-Tiphase since vanadium is a b stabilizer.

A cross-sectional BFTEM image of nitrided Ti64 isshown in Fig. 3e. It is evident from this cross-section thatthere is a combination of an a-lath and b-rib that conrmsthe b phase is present despite its weak XRD signal in

phases in Ti64, Ti64-N, TNZT and TNZT-N alloys.

ca (A) c/a d(110)b (A) ab (A)

4.651 1.59 4.686 1.60

2.342 3.3124.816 1.61 2.333 3.30

LENfor e

MateFig. 2. (a) Low- and (b) high-magnication backscatter SEM images ofSEM images of LENS deposited nitrided Ti64 (Ti64-N) alloy. See text

64 H. Mohseni et al. / ActaFig. 1a. An EDS line prole (not shown) conrmed anincrease in the vanadium X-ray counts with a correspond-ing decrease in titanium counts across the b-rib. The dier-ence in the solubility of vanadium, a b stabilizer, in a-Tiand b-Ti clearly indicates the border between the a-lathand b-rib. The b-Ti has a higher solubility for vanadium anda lower solubility for Al, which is an a stabilizer. From boththe TEM image and the EDS line prole, this b-rib is5070 nm thick.

3.1.2. Nitrided TNZTThe structure of LENSe deposited TNZT is signi-

cantly dierent from that of Ti64 as conrmed by theXRD pattern in Fig. 1b which reveals the presence of peakscorresponding to the b phase and negligible peaks of the aphase. Such an observation was anticipated due to the largeamounts of b stabilizing alloying elements (Nb, Zr and Ta)in TNZT. Conversely, the LENSe deposited nitridedTNZT sample exhibits the precipitation of multiple phasesthat were indexed as tetragonal e-Ti2N (a = 4.822 A andc = 3.067 A), cubic rocksalt d-TiN (a = 4.226 A), a-Tiand b-Ti. Table 1 summarizes the calculated d-spacingsand lattice parameters for b-Ti in TNZT and nitridedTNZT. The incorporation of nitrogen in TNZT shows thatnitrogen has an insignicant eect on the b phase sincethere is no change in the lattice parameter (a = 3.30 A).This observation is expected since nitrogen is known tobe a strong a stabilizer and has limited solubility in the bphase at ambient temperature. However, the XRD patternin Fig. 1b shows that signicant precipitation of the a phaseoccurred due to nitrogen incorporation into LENSedeposited TNZT. A similar eect has been observed onS deposited Ti64 alloy. (c) Low- and (d) high-magnication backscatterxplanation of Type 1 and 2 precipitates.

rialia 83 (2015) 6174the surface of nitrided TNZT [18]. The lattice parametersof the a phase in nitrided TNZT were calculated to bea = 2.987 A and c = 4.816 A, which results in a c/a ratioof 1.61 compared to the ideal ratio of 1.59 for pure a-Ti.The slight increase in the c/a ratio due to nitriding is con-sistent with the values reported in previous studies[28,29]. This increase is an indication of saturation of thea phase by nitrogen and corresponds to the solubility limitof nitrogen (i.e. 2425 at.%) in a-Ti, shown by the TiNbinary phase diagram in Fig. 4 [31].

The inuence of nitrogen on the structure of the LENSedeposited TNZT is also evident by comparing the BSEimages of base (Fig. 5a) and nitrided (Fig. 5b) TNZT.The dendritic microstructure typical of single-phase TNZTcan be seen in Fig. 5a where b dendrites are separated bydarker interdendritic regions that show microsegregationof alloying additions. Such microsegregation is a commonfeature in alloys with a high concentration of b stabilizers[3]. Fig. 5b shows a completely dierent microstructurefor nitrided TNZT where the phase with brighter contrastrepresents the bmatrix and the phases exhibiting the darkercontrast include TiN/Ti2N and a (grain boundary andintragranular a precipitates). The higher-magnicationinset shown in Fig. 5b clearly shows these two dierent pre-cipitate morphologies: ner scale intragranular a precipi-tates with platelet morphology, and coarser TiN/Ti2Nexhibiting a dendritic morphology. Fig. 5c shows a cross-sectional STEM-HAADF image along with EDS elementalmaps of a site-specic sample prepared from the nitridedTNZT sample. It is evident from the compositional mapsthat the intragranular a precipitates (200 nm thick) aredecient in Nb and Ta (b stabilizers), whereas the b matrix

MateH. Mohseni et al. / Actais enriched in Nb and Ta. The Zr map was noisy due to thelower counts and hence has not been included.

3.1.3. Evolution of microstructure in nitrided Ti64 andTNZT

From the XRD, SEM and TEM/EDS analyses, theeect of nitrogen addition on the microstructural evolutionis considerably dierent for LENSe deposited Ti64 andTNZT. The TiN binary phase diagram in Fig. 4 can

Fig. 3. (a) Backscatter SEM image of nitrided Ti64 alloy with correspondingsectional BFTEM image of nitrided Ti64.rialia 83 (2015) 6174 65provide some insights towards the probable sequence ofphase evolution during in situ nitriding of Ti64 and TNZT.Since Ti64 is an a/b alloy, addition of nitrogen to this sys-tem can possibly expand the a phase stability region on thephase diagram. The increased volume fraction of a phase iscapable of absorbing more nitrogen thereby preventing theprecipitation of nitride phases. Conversely, the eect ofnitrogen can be somewhat complicated in nitrided TNZTsince the high concentration of b stabilizing elements

EDS maps of (b) nitrogen, (c) aluminum and (d) vanadium. (e) Cross-

Mate66 H. Mohseni et al. / Acta(Nb, Zr and Ta) can counteract the a stabilizing eects ofnitrogen. Furthermore, the complex interplay of the multi-ple components in this alloy can possibly lead to changes inthe shapes and sizes of the various phase elds on the

Fig. 5. Low-magnication backscatter SEM images of LENS deposited (a)HAADF-STEM image of nitrided TNZT alloy with corresponding titanium

Fig. 4. Binary TiN phrialia 83 (2015) 6174multicomponent phase diagram in such a manner as tonucleate primary TiN precipitates in the liquid phase fol-lowed by nucleation of b grains, as evident by the resultingmicrostructure in Fig. 5b. Consequently, the primary TiN

TNZT and (b) nitrided TNZT (TNZT-N) alloys. (c) Cross-sectional, niobium and tantalum EDS maps.

ase diagram [31].

precipitates forming in the liquid exhibit a dendritic mor-phology, and can act as heterogeneous nucleation sitesfor the b grains during solidication, resulting in moreequiaxed b grains (Fig. 5b) as compared to the dendriticb grains observed in LENSe deposited TNZT in theabsence of nitrogen (Fig. 5a). Despite the prior formationof the d-TiN dendrites, at the elevated temperatures wherethe b grains nucleate, they have a certain solubility of nitro-gen, as evident from the TiN binary phase diagram inFig. 4. Thus on subsequent cooling this retained nitrogenwill facilitate the precipitation of a phase at both heteroge-neous grain boundary sites as well as homogeneouslywithin the b grains. Furthermore, as soon as a precipitateswithin the b matrix, the retained nitrogen within the b par-titions to the a phase, resulting in an increase in the c/aratio of this phase.

To better understand the possible phase formationsequence, isothermal sections of the TiNbN ternary sys-tem were simulated using PANDATe developed by Com-putherm LLC and are shown in Fig. 6ac, correspondingto temperatures of 2000, 1000 and 100 C, respectively.Nb was chosen as the third component since it is the mostabundant element in the alloy after Ti. This ternary phasediagram can provide some valuable insights on the compet-

severity in wear. It is apparent that the surface morphology

H. Mohseni et al. / Acta Mateing nature of the b stabilizing elements and a stabilizingeects of nitrogen on the phase stability in this complexalloy. From Fig. 6a, it evident that fcc rocksalt d-TiN is ahigh-temperature phase and exists at 2000oC in conjunctionwith bcc b phase and liquid for a nominal composition of35 wt.% Nb and 10 wt.% nitrogen. Fig. 6b shows fora similar average composition that bcc b, d-TiN and Ti2Nare the stable phases at 1000oC, indicating the possible for-mation of the second nitride, Ti2N phase. Further loweringthe temperature to 100 C results in the precipitation of thehcp a phase in addition to the bcc b, d-TiN and Ti2Nphases, as shown in Fig. 6c. The retention of d-TiN den-drites, in the LENSe deposited TNZT at room tempera-ture, is explained by the non-equilibrium cooling involvedFig. 6. Isothermal sections of ternary TiNbN phase diagramshowing Ti-rich compositions at (a) 2000 C, (b) 1000 C and (c)100 C. Closed circle denotes the approximate ternary composition.has undergone a plastic shear deformation mechanismtypically associated with less wear volume removal. TheSEM images of the TNZT surface exhibit ductile layeringand smearing of coarse wear platelets with varying sizes(515 lm) aligned along the sliding direction. Thesedierences in wear surface morphology are consistent withprevious sliding wear studies on Ti64 and TNZT [7,32].

Fig. 8 shows SEM images of the worn nitrided Ti64 andTNZT surfaces. From these and other images, the averagedwear track widths calculated in Table 2 were considerablyand the sluggish decomposition kinetics of this nitridephase. Furthermore, it should be noted that the PAN-DATe predictions hold only for the equilibrium phasesand are dependent on the accuracy of the thermodynamicmodels used and their calibration. It can be summarizedthat the basic variations in the starting microstructuresand nature of alloying additions in Ti64 and TNZT areresponsible for the unambiguous dierences that nitrogenaddition has in modifying their respective microstructures.

3.2. Microhardness and wear behavior

The result of microhardness analysis of the LENSedeposited base and nitrided Ti64 and TNZT are summa-rized in Table 2. Ti64 has a higher hardness (419 VHN)compared to TNZT (288 VHN), since the b phase is softerthan a, which makes the Ti64, with a/b structure, harderthan TNZT. The averaged microhardness of both alloysapproximately doubled with nitriding: 1047 and 518VHN for nitrided Ti64 and TNZT, respectively. Theincreasing hardness with in situ nitriding can be attributedto the formation of TiN/Ti2N phases in TNZT, and thenitrogen-enriched a phase, i.e. increasing volume ofa(Ti,N) precipitates in Ti64. This increase in microhardnessis consistent with prior work on nitriding a/b and b Tialloys [16,28,29].

Sliding wear behavior comparisons of the four alloys aresummarized in Table 2. The values of the wear track depthand width along with the wear factor were substantiallyhigher for Ti64 compared to TNZT despite the higher hard-ness of Ti64. These observations do not agree with Arch-ards law for adhesive and abrasive sliding wear,according to which lower wear volume loss correspondsto a harder material. This suggests that hardness alone can-not account for the dierence in wear behavior. Afterexamining the Ti64 and TNZT wear surfaces by SEM, itis evident that the Ti64 surface shown in Fig. 7a exhibitslarge amounts of debris particulates on the surface. EDSconrmed these are SiOx particles that were tribochemicallytransferred from the sliding Si3N4 counterface to the wearsurface. A harder surface such as Ti64 would more easilyabrade the Si3N4 counterface, thus resulting in a larger areaof adhesively transferred SiOx wear debris. In comparison,a much lower concentration of transferred SiOx particles ispresent on the TNZT shown in Fig. 7b, since TNZT issofter and less abrasive. There are also wear surface mor-phology dierences between the alloys. The SEM imagesof the Ti64 surface in Fig. 7a exhibits extensive plasticdeformation, microplowing and cutting, all typical pro-cesses associated with an abrasive wear mechanism. Con-versely, the TNZT wear surface in Fig. 7b shows less

rialia 83 (2015) 6174 67smaller for nitrided Ti64 and TNZT alloys compared tothe base alloys, and the nitrided TNZT exhibited slightly

depce tra

m)

MateTable 2. Summary of averaged Vickers microhardness values, wear trackFor nitrided Ti64 and TNZT alloys, adhesive wear of Si3N4 counterfacalculations.

Alloy Microhardness (HV) Wear track depth (l

Ti64 419 32 20.3 0.7Ti64-N 1047 29 TNZT 288 21 3.9 0.4TNZT-N 518 35

68 H. Mohseni et al. / Actasmaller width than nitrided Ti64. However, the wear trackdepth for both nitrided alloys was not measurable since itwas within the noise of the surface prolometer (50 nmdepth resolution). Hence wear factors could not be calcu-lated, although wear track depths for both nitrided alloyswere orders of magnitude less than the base alloys. Theseobservations indicate a reduction in the real area of contactbetween the nitrided alloys and the Si3N4 counterface dueto increased hardness from nitriding that results in decreas-ing wear. Measuring the wear depth values was also di-cult due to the transfer of SiOx debris from the Si3N4counterface to the wear track thereby preventing accuratewear calculations. An additional comparison between theFig. 8 SEM images of nitrided Ti64 and TNZT wear sur-faces reveal that the nature of surface fracture appears dif-ferent. Surface cracks in nitrided Ti64 occur at more of anoblique angle to the sliding direction, denoted by arrows inthe higher-magnication image in Fig. 8a, while the cracksin nitrided TNZT appear to be conned within thenear-surface platelet wear debris. These dierences inmicrostructural features imply dierent wear mechanism(s)are occurring between nitrided Ti64 and TNZT and will be

Fig. 7. Low- and high-magnication SEM imagth, width and wear factor for Ti64, Ti64-N, TNZT and TNZT-N alloys.nsferred to the wear track prevented accurate wear depth and factor

Wear track width (lm) Wear factor (mm3 (Nm)1)715 39 6.4 0.2 104330 15 420 20 5.3 0.3 105312 11

rialia 83 (2015) 6174discussed in the next section. While the above observationsprovide some insight into sliding-induced deformation,they are conned to the surface and are of lower resolution.Therefore, a more detailed mechanistic study of the surfaceand subsurface structural evolution during sliding wear isnecessary and will be discussed below.

3.3. Wear-induced deformation mechanisms

3.3.1. Cross-sectional TEM analysisTo further investigate the sliding wear mechanisms of

nitrided Ti64 and TNZT, cross-sectional FIB/TEM analy-ses were performed inside the wear tracks. Such wear defor-mation analyses for nitrided Ti alloys are unknown, and theuse of PED in this context is particularly novel. Fig. 9 rep-resents the cross-sectional TEM and PED orientation anal-ysis directions for ball-on-disc sliding, namely SD = slidingdirection, TD = transverse direction and ND = normaldirection. During LENSe processing, the alloy depositionis parallel to the ND. Therefore, the top wear surface cor-responds to the ND of the frame, while for the wear subsur-face viewed in cross-section, the alloy deposition is

es of (a) Ti64 and (b) TNZT wear tracks.

and (b

H. Mohseni et al. / Acta MateFig. 8. Low- and high-magnication SEM images of nitrided (a) Ti64denote cracks.perpendicular to the ND of the frame, i.e. along the TD.Fig. 10 presents the wear-induced surface and subsurfacestructural evolution in nitrided Ti64 acquired along theTD. Both the low- (Fig. 10a) and higher- (Fig. 10b) magni-cation BFTEM images show that on the wear surface,below the protective Pt, is a 430 nm thick layer of amor-phous silica (a-SiOx), which was conrmed by an EDS linescan. This layer transferred from the Si3N4 counterfaceduring the wear process as previously discussed in termsof the SEM surface analysis in Fig. 8a, showing that silicawear debris is present in some locations on the surface. Thesubsurface BFTEM images in Fig. 10 further show that thisporous SiOx transfer layer exhibits cracking along the SD.Below the SiOx transfer layer is a near-surface deformedregion that contains an extensive network of shear bandswith dislocation substructures (tangles/networks) inbetween. The presence of shear bands, occurring to a

Fig. 9. Schematic representation of the cross-sectional TEM and PEDorientation analyses directions for ball-on-disc sliding wear. SD = slid-ing direction, TD = transverse direction, ND = normal direction.) TNZT wear tracks. Arrows in the higher magnication image in (a)

rialia 83 (2015) 6174 69subsurface depth of 2 lm based on this FIB liftout,implies that wear-induced fracture processes are occurringin nitrided Ti64. This process involves continuous forma-tion of sliding-induced shear bands (some up to 1.5 lmlong) that branch, network and evolve into cracks, likethe ones shown on the surface in Fig. 8a. Subsequently,Ti64 third-body abrasive wear particles of varying sizesare generated from these networked cracks. These particles,either by themselves or when combined with SiOx wear deb-ris, result in microplowing and cutting, processes associatedwith an abrasive wear mechanism, which is consistent withthe SEM surface imaging in Fig. 8a. In addition to the brit-tle behavior, there is evidence of wear-induced plasticity innitrided Ti64. The dislocated substructure shown in Fig. 10surrounding the shear bands suggests a work-hardenedzone that would lead to further hardening in the worn area.In comparison, the base Ti64 wear surface, shown inFig. 7a, did not exhibit signs of surface cracking despiteits high wear rate. Therefore, in situ nitriding resulted inincreasing hardness that improved the wear resistance ofTi64; however, the enhancement in the precipitation of aphase with nitriding is likely responsible for surface/subsurface cracking.

Completely dierent wear-induced surface and subsur-face structures were observed in nitrided TNZT. Fig. 11aand b represent two regions of the nitrided TNZT weartrack, acquired along the TD, where the SiOx wear debrisis and is not present on the surface, respectively. This con-rms that there is a heterogeneous distribution of the SiOxdebris on the wear surface. For the case where SiOx is pres-ent, below it is a non-uniform (400900 nm thickness)heavily grain rened zone. This indicates that the depthof plastic strain due to sliding is within this thickness range.

Mate70 H. Mohseni et al. / ActaThe severely deformed zone is composed of equiaxed grainsvarying from 10 to 200 nm size that are presumably a andb grains and will be veried by PED in the next section.Below this wear-induced plastically deformed zone,denoted by the dashed line, is a transition to the unde-formed bulk structure. Fig. 11b shows the same severelydeformed zone with nanocrystalline grain renement,although with a more uniform thickness of 700 nm. Thepresence of this zone without transferred SiOx on top sug-gests counterface wear is not necessary to induce grainrenement. Furthermore, there is no evidence of shearbands that network into cracks, as previously shown fornitrided Ti64, within this heavily grain rened zone orbetween the interfaces. While the nitrided TNZT grainrened zone is less prone to fracture processes, both alloysurfaces are unstable due to localized shear instability dur-ing sliding. As a result, the near-surface material can gener-ate ner loose wear debris particles that can alsomechanically mix with the counterface material, e.g. SiOx.

Fig. 10. (a) Low- and (b) high-magnication cross-sectional BFTEMimages of worn nitrided Ti64 acquired along the TD.rialia 83 (2015) 6174While the hardness of both alloys increases with nitriding,it alone cannot account for the improvements in wearresistance, e.g. the hardness of base Ti64 and nitrided

Fig. 11. Cross-sectional BFTEM images of worn nitrided TNZT (a)with and (b) without SiOx counterface wear debris. The white dashedlines represent the borders between the wear-induced zones. Imagesacquired along the TD.

TNZT are approximately the same despite their large dier-ences in wear values (Table 2). Thus, the dierences in sur-face and subsurface structural evolution during slidingwear are also determining the wear behavior of these alloys.

Recrystallization due to sliding is unlikely since the com-puted interfacial ash temperature due to frictional heatingis low (40 C), due in part to the low contact stress andsliding speed used in this study. Therefore, the wear-induced structural changes are a result of the high deforma-tion strain during sliding, e.g. large plastic shear strainsexceeding 10 are common with sliding contact of metallicmaterials [33]. In the case of nitrided TNZT, formation ofmore stable nanocrystals of equiaxed a and b-Ti withoutfracture, e.g. shear bands, indicates the potential to accom-modate the interfacial shear stresses during sliding contactthat reduces wear. Detailed TEM-based orientation map-ping of the severely deformed zone in nitrided TNZT willbe presented in the following section to explain what rolethese rened a and b grains have on the wear behaviorand mechanism. The presence of nanocrystalline grainrenement has been determined to improve sliding wearresistance in titanium when processed by a severe plasticdeformation (SPD) technique. For example, Stolyarovet al. reported an improvement in the wear resistance ofcommercially pure titanium after equal-channel angularpressing (ECAP) that resulted in an ultrane-grained(UFG) structure that was responsible for a decrease inthe adhesive component of friction and wear [34]. Similarly,

UFG (300 nm size) Ti prepared by ECAP showed thatgrain renement was responsible for decreasing wear ratesin comparison to coarse-grained Ti [35]. In addition, crea-tion of a nanocrystalline surface layer (160 lm thick) onpure Ti using a surface mechanical attrition treatmentwas determined to be an eective procedure to reduce wear[36]. A recent review article on the tribological properties ofUFG Ti and other metallic materials processed by SPDprovides more details on the above ECAP studies [37].

Lastly, there was no evidence of titanium oxide, e.g.TiO2, formation by tribo-oxidation on the wear surfacesfor both nitrided alloys, due in part to the aforementioned,relatively low interfacial ash temperatures during sliding.Under more severe testing conditions, e.g. higher contactstresses and sliding speeds, frictional heating has beenshown to facilitate tribochemical formation of TiO2 surfacelayers when sliding on ECAP UFG Ti that resulted in lowerfriction and wear [35]. However, these sliding conditions,especially the high contact loads/stresses, are more severethan would be exhibited in certain applications, e.g. bio-medical load-bearing applications such as hip implants.Furthermore, such surface layers can mask the role andimportance of the dierent underlying wear-inducedstructures.

3.3.2. Cross-sectional TEM-PED analysisFig. 12a and b show cross-sectional inverse pole gure

(IPF) orientation maps of b-Ti and a-Ti, respectively,

ntatione 2 u

H. Mohseni et al. / Acta Materialia 83 (2015) 6174 71Fig. 12. Cross-sectional IPF maps of worn nitrided TNZT showing orietop portion of the zone 1 worn region, and (d) bottom portion of the zo

dashed lines represent the border between zone 1 and zone 2. In (c) and (d) pfrom zone 1 (yellow dashed line box) and zone 2 (brown dashed line box) regithe reader is referred to the web version of this article.)n map of (a) b-Ti, and (b) a-Ti and phase concentration maps of the (c)nworn region. The insets show grain orientation color codes. The white

hase maps, red represents b grains and green represents a grains takenons. (For interpretation of the references to colour in this gure legend,

acquired from PED analysis of worn nitrided TNZT. Theseorientation IPF b-Ti and a-Ti grain maps, with their colorcode in the inset, were acquired along the TD from thesame FIB liftout as the cross-sectional TEM image shownin Fig. 11a. Based on the b-Ti and a-Ti grain size and dis-tribution, there are two distinct zones in the worn cross-sec-tion, denoted zones 1 and 2 by the white dashed lines inFig. 12a and b. Zone 1 is the top portion with wear-inducedgrain renement as previously observed in the cross-sec-tional BFTEM images in Fig. 11. Zone 1 is composed ofboth ultrane b-Ti grains (10100 nm size) shown inFig. 12a and mostly larger grains of a-Ti shown inFig. 12b. Zone 2 is the bottom portion of the unwornnitrided TNZT bulk structure that was not exposed towear-induced subsurface contact stresses.

The IPF maps in Fig. 12 were also used to observe grainorientation changes and texture evolution resulting fromsliding-induced deformation. Fig. 12a shows that the b-Tigrains in zones 1 and 2 are randomly orientated while themajority of the a-Ti grains (unworn and worn) exhibit(2110) planes that are parallel to the ND, represented on

the IPF map in green in Fig. 12b. Not shown are IPF mapsrotated at 90 to the TD of the sample (viewed along theND) where the (0001) planes, red on the IPF map, are per-pendicular to the ND. Therefore, the a-Ti (0002) planes areparallel to the wear surface, i.e. lie in the sliding plane (SDand TD). From these IPF maps, it is evident, at least basedon this particular TEM liftout, that a pronounced a-Ti tex-ture of the basal type remains after sliding. More details ofthis texture will be presented below with pole gure analy-sis. Furthermore, Fig. 12 shows there are clear dierences inthe b-Ti and a-Ti grain sizes in zone 1 compared to theunworn grain sizes in zone 2, with the former exhibitingmore grain renement. This is due to the softer and morecompliant nature of the b grains in comparison to the sur-rounding hard/sti a grains [11], which results in signicantb grain renement during the wear process. This grain sizecomparison is more evident with viewing the combined b-Tiand a-Ti phase maps of zones 1 (worn) and 2 (unworn) inFig. 12c and d, respectively. Clearly the b-Ti grains,denoted in red, are more rened in zone 1, while the a-Tigrains in green are less rened. Phase concentration

72 H. Mohseni et al. / Acta Materialia 83 (2015) 6174Fig. 13. Comparison of b-Ti and a-Ti pole gures acquired from worn (zonetexture scale bars are in multiples of random.1) and unworn (zone 2) nitrided TNZT viewed along ND. Units for the

sliding thereby reducing wear. Basal slip of hcp metals alongthe sliding direction has been determined to lower the resis-

Med. Devices, Vol. 23, ASM International, Materials Park,OH, 2012.

Matetance to shear in sliding contacts [38]. The b-Ti pole guresin Fig. 13 show that the rened and unworn b-Ti grains donot exhibit any specic texture, which is corroborated by thelower spread in texture intensity values compared to a-Tivalues. Along with the higher degree of randomness of thezone 1 b-Ti grains, there is also the possibility that thesesofter/compliant b-Ti rened grains exhibit grain rotationduring the wear process.

4. Summary and conclusions

Detailed studies of processingstructureproperty rela-tions were conducted for in situ nitrided titanium alloysTNZT and Ti64 with comparisons to untreated TNZTand Ti64. Dierences in microstructural evolution duringsolidication and wear have been determined as follows:1. Nitriding during LENSe deposition resulted in a sub-

stantial increase in the microhardness of TNZT andTi64 due to the precipitation of TiN/Ti2N phases and/or nitrogen-enriched a phase, respectively. This phaseevolution with nitrogen addition during solidicationwas relatively well predicted with thermodynamic calcu-lation software.

2. Increased hardness in both nitrided alloys was partiallyresponsible for the reduction in real area of contact dur-ing interfacial sliding that resulted in improved wearresistance. However, hardness alone cannot accountfor the dierences in wear behavior with and withoutnitriding since there were clear dierences in sliding-induced deformation mechanisms. In the case of nitridedTi64, strain-induced brittle fracture in the form of sur-face and subsurface shear bands resulted in crackingand subsequent third-body abrasive wear debris genera-tion. The networked shear bands were surrounded by acalculations determined that zone 1 contains 60% a-Ti and40% b-Ti, whereas this proportion changes to 47% and53%, respectively, in zone 2. While this is not a detailed sta-tistical phase analysis based on several FIB liftouts alongthe wear track, it does show that more a-Ti is present inthe wear-induced zone 1 compared to unworn zone 2. Thisdistribution is likely a result of the aforementioned a-Tiphase exhibiting higher hardness and stiness than b-Ti.Thus, the a-Ti phase has a higher resistance to interfacialshear stresses present during sliding wear, whereas thesofter/compliant b-Ti grains become more rened with anincreasing probability of becoming loose, ejected wear par-ticles that would account for their deceasing concentrationin zone 1.

Corresponding b-Ti and a-Ti pole gures acquired fromworn (zone 1) and unworn (zone 2) nitrided TNZT are pre-sented in Fig. 13. These pole gures are viewed down theND (rotated 90 to the TD) to facilitate interpretation ofthe a-Ti (0002) texture evolution during wear. In comparingthe a-Ti worn and unworn < 0001 > pole gures, there is anincrease in the alignment of the a-Ti basal planes along theSD due to the wear process. In contrast, the unworn a-Ti < 0001 > pole gure shows there are more spots with lessalignment along the SD. This sliding-induced alignment ofa-Ti basal planes would lower interfacial shear stress during

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3. Orientation imaging analysis of nitrided TNZT usingTEM-PED revealed that the a-Ti grains exhibit basaltexture and become more aligned parallel to the slidingdirection during wear, while the softer and compliantnanocrystalline b-Ti grains are randomly orientatedand likely prone to rotation during wear. These slid-ing-induced phenomena coupled with the hard phasesthat form during solidication are responsible forimprovements in wear resistance with in situ nitriding.

Acknowledgements

The authors would like to acknowledge the support of the AirForce Research Laboratory (AFRL)-sponsored UNT Institute forScience and Engineering Simulation (ISES) with Grant FA8650-08-C-5226. We also acknowledge the UNT Center for AdvancedResearch and Technology (CART), UNT College of Engineeringfor sponsoring A.T. as part of its SUPER program, and AnchalSondhi for assistance with XRD acquisition.

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74 H. Mohseni et al. / Acta Materialia 83 (2015) 6174

In situ nitrided titanium alloys: Microstructural evolution during solidification and wear1 Introduction2 Experimental methods3 Results and discussion3.1 Microstructure3.1.1 Nitrided Ti643.1.2 Nitrided TNZT3.1.3 Evolution of microstructure in nitrided Ti64 and TNZT

3.2 Microhardness and wear behavior3.3 Wear-induced deformation mechanisms3.3.1 Cross-sectional TEM analysis3.3.2 Cross-sectional TEM-PED analysis

4 Summary and conclusionsAcknowledgementsReferences