Mechanical and wear properties of ultrafine-grained pure Ti produced by multi-pass equal-channel angular extrusion

Materials Science and Engineering A 517 (2009) 97–104

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Mechanical and wear properties of ultrafine-grained pure Ti produced by multi-pass equal-channel angular extrusion G. Purcek a,∗ , O. Saray a , O. Kul a , I. Karaman b , G.G. Yapici b , M. Haouaoui b , H.J. Maier c a b c

Department of Mechanical Engineering, Karadeniz Technical University, 61080 Trabzon, Turkey Department of Mechanical Engineering, Texas A&M University, College Station, TX 77843-3123, USA Lehrstuhl für Werkstoffkunde, Universität Paderborn, Germany

a r t i c l e

i n f o

Article history: Received 23 October 2008 Received in revised form 14 March 2009 Accepted 23 March 2009 Keywords: Equal-channel angular extrusion Ultrafine-grained materials Titanium Mechanical behavior Wear

a b s t r a c t In this study, pure grade 2 Ti was processed via equal-channel angular extrusion (ECAE) for 8 and 12 passes following route-E at 300 ◦ C. After processing, the microstructural evolution, tensile properties and wear behavior were investigated. ECAE-processed Ti exhibited a significant improvement in strength values with a slight decrease in ductility. However, the wear test results surprisingly showed that the strengthening of titanium by ECAE processing does not lead to the improvement of wear resistance at least for the pressures and sliding distances used in this study. This finding was mainly attributed to the tribochemical reaction leading to oxidative wear with the abrasive effect in Ti. Three distinct regions were formed on the subsurface of CG and UFG Ti after sliding wear, which are the tribolayer including titanium oxide with smeared wear material at the top, a deformed region having material structure oriented along the sliding direction in the middle, and the original unaffected bulk material at the bottom. © 2009 Elsevier B.V. All rights reserved.

1. Introduction Pure titanium and its alloys are used in many applications ranging from biomedical to aerospace due to their low density, high specific strength, excellent corrosion resistance and high biocompatibility [1,2]. One important application for commercial purity (CP) Ti and its well-known alloy, Ti–6Al–4V, is for bone and other medical implants in human body [3]. Ti–6Al–4V was first developed for the aerospace industry. The alloying elements, Al and V, are toxic and may potentially cause a series of ailments including cancer [3]. Pure titanium has several advantages for medical use such as better corrosion resistance, higher inertness and biocompatibility as compared with Ti–6Al–4V alloy which is currently the material of choice for most medical implants due to its high strength [3–5]. Therefore, it is desired to replace Ti–6Al–4V with high strength pure Ti in medical applications. However, the application of ordinary coarse-grained (CG) Ti is limited due to its low strength and fatigue limit needed for medical implants [6]. Thus, there have been many processing efforts in order to improve the strength of this material without significant decrease in its ductility. Grain refinement via severe plastic deformation (SPD) techniques has been shown to be an effective route for achieving this goal [4]. The equal-channel angular extrusion (ECAE) is one of the SPD pro-

∗ Corresponding author. Tel.: +90 4623772941; fax: +90 4623255526. E-mail address: [email protected] (G. Purcek). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.03.054

cessing techniques which allow the application of extremely large imposed strains in bulk samples without fracture, realizing grain refinement and significant strengthening through specific structure formation by simple shear [2,7,8]. The advantage of ECAE is its capability of maintaining the initial dimensions of the work piece so that the process can be repeated to obtain the desired strain and grain size. Most studies, which have been focused on structural and mechanical characteristics of ECAE-processed materials, emphasize that a combination of a high strength with a good plasticity can be achieved after SPD [2,9]. Some of them have been focused on the ECAE processing of pure Ti with different purity and grain sizes, and titanium alloys. These studies are of great interest, since the plastic deformability of these materials is normally inferior to that of cubic-structured metals and alloys due to their hexagonal close packed (hcp) structure that has limited number of slip systems available. The work on SPD processed Ti and its alloys started extensively after 2000, most of which is on the ECAE processing of pure Ti and Ti–6Al–4V. Stolyarov et al. [10] studied the microstuctural evolution and corresponding mechanical properties in pure Ti processed using ECAE and ECAE followed by cold extrusion. They found that route-Bc among the standard ECAE routes was the best route for achieving equiaxed grains and effectively refined microstructure. Further cold extrusion and cold rolling resulted in improved strength levels [3,11,12]. Kim et al. [13] studied the effect of ECAE temperature on the microstructural evolution of pure Ti. The reason for abnormal strain hardening in pure Ti after ECAE plus cold

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rolling has been investigated by Wang et al. [14]. They have found that deformation twinning may have played a significant role in plastic deformation of UFG Ti in addition to dislocation slip. Yapici et al. [4] studied the anisotropic features of the material flow after ECAE plus cold rolling, and they observed that the post-ECAE cold rolling produces stronger crystallographic texture. In addition, the cyclic behavior of UFG Ti produced using ECAE was investigated, and a significant enhancement in fatigue life were reported [2,6,15]. Besides pure Ti, some studies have focused on Ti–6Al–4V (Ti64) and its structural and mechanical property improvements after ECAE [1,5,16,17]. There are limited numbers of studies on the friction and wear properties of the ECAE-processed materials although they are important for certain applications, especially those involving contact deformation and damage [18,19]. Increase in wear resistance along with improved strength can potentially make CP Ti attractive for wear-sensitive biomedical applications like dentistry and hip-joints. However, little information exists regarding the friction and wear properties of ECAE-processed pure Ti [20,21]. An earlier work showed that the wear resistance of the annealed CG Ti slightly increased after ECAE processing under dry sliding condition [20]. Also, the friction coefficient of ECAE-processed Ti was found to be slightly lower than that of the CG Ti. Another study was on the friction properties of UFG Ti processed by ECAE and ECAE + cold rolling [21]. This study showed that the friction coefficient of UFG Ti is almost similar to that of CG Ti at room temperature but is markedly lower than that of CG Ti at elevated temperatures. Recently, Garbacz et al. [22] studied the tribological properties of nano-Ti obtained by hydrostatic extrusion and they found no improvement in wear resistance of nanostructured pure Ti after processing although a substantial enhancement was obtained in its strength. In view of the above, the objective of this work is to study the influence of ECAE on the mechanical properties and wear behavior of CP grade 2 Ti. Also, the microstructural evolution due to ECAE has been examined. The results obtained and potential operating wear mechanisms in UFG Ti ECAE processed using route-E for 8 passes and 12 passes are presented and compared with those from CG Ti. 2. Materials and experimental procedures Commercial purity (CP) grade 2 Ti initially hot rolled with an average grain size of about 110 ␮m and with impurities including 0.15 wt.% O, 0.002 wt.% H, 0.006 wt.% N, 0.041 wt.% Fe and 0.008 wt.% C was used as the starting material. The as-received Ti bars were extruded in a sharp 90◦ angle ECAE die which was preheated to 300 ◦ C. The extrusion rate was chosen as 1.27 mm s−1 . Between each pass, the billets were heated in the die for 15–30 min at the deformation temperature. Following each pass, the billets were water quenched to maintain the microstructure achieved during ECAE. The minimum deformation temperature which allowed processing without shear localization and macroscopic cracking was determined to be 300 ◦ C. ECAE at the lowest possible temperature is crucial in preventing recrystallization, and partly achieved at 300 ◦ C in this study using the sliding walls concept in the ECAE die which helps reducing the die frictional effects [23]. The die and the billets are coated with graphite base lubricant to minimize the friction and eliminate sticking of the specimens to the die during the extrusion. In this study, in quest for achieving the best mechanical response, an alternative ECAE route, i.e. route-E was applied. This route is a combination of the standard routes of ECAE, in which the billet is rotated 180◦ and again 180◦ in subsequent passes (2C) followed by a 90◦ rotation and then another 2C extrusion [24]. The billets were processed for a total of 8 and 12 passes using this route.

Optical microscopy (OM) and transmission electron microscopy (TEM) were used to investigate the microstructural evolution of pure Ti. Optical microscopy has only been utilized to investigate the as-received microstructure. Due to the high dislocation density and the refined grain size of the as-processed material, TEM was used for microstructural investigations. The samples were prepared from transverse and longitudinal sections of the ECAE-processed billets. OM samples were mechanically ground down to 1200 SiC grit and then polished with 9 ␮m diamond suspension and 0.05 ␮m colloidal silica mixed with 30% H2 O2 . They were chemically etched with Kroll’s reagent before optical investigation. TEM investigations were carried out in a Philips CM200 microscope operated at a nominal accelerating voltage of 200 kN. Tensile tests were performed at room temperature using an MTS 810 servo-hydraulic test frame. Monotonic tension tests were performed at the quasi-static strain rate regime (10−4 to 10−3 ). The tension samples had a dog-bone shape with a gage section of 1.5 mm × 3 mm × 8 mm where their tensile axis were oriented parallel to the extrusion direction of the ECAE-processed billets. Strain was measured using a miniature 3 mm gage length extensometer directly attached onto the gage sections of the specimens. Experiments were repeated on two to three companion specimens to check the repeatability of the results. Dry sliding wear tests were conducted using a pin-on-disc type test rig (ASTM G-99) at room temperature (19 ± 2 ◦ C) and at a relative humidity of about 47 ± 5%. The wear test rig is schematically illustrated in Fig. 1. The samples for wear tests were machined into a cylindrical pin (5 mm diameter and 20 mm long) which slides against an AISI 1045 steel disc with a diameter of 80 mm, hardness of 56 RC , and surface roughness (Ra ) of about 0.3 ␮m. The wear tests were performed for a total sliding distance of 10.40 km corresponding to a period of 2.60 h under a constant pressure of 1.5 MPa using a sliding speed of 1.1 m s−1 . Prior to the each test, the surface of the pins were grinded with 1500 grit SiC emery paper, and then cleaned in an acetone with an ultrasonic bath for 10 min and dried in hot air. The roughness of the pin wear surfaces was measured in the range of 0.2–0.3 ␮m after this procedure. The wear resistance was measured by a weight loss technique using a microbalance with an accuracy of ±0.01 mg. The wear samples were cleaned first with carbon tetra chloride (CCl4 ) and then acetone before and after the wear test in order to determine the weight loss. To reveal the effect of applied pressure on the wear resistance, the experiments were performed at five fixed values for the mean pressure at the interface: 0.5, 1.0, 1.5, 2.0, 2.5 MPa. The sliding distance was 1.52 km with constant sliding speed of 1.1 m s−1 for each test. Morphologies of worn surfaces and subsurfaces of as-received and as-processed samples were examined using a scanning electron microscope (SEM) in the secondary electron image mode at 15 keV

Fig. 1. A schematic representation of the pin-on-disc wear test configuration.

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3. Results and discussion 3.1. Microstructural evolution

Fig. 2. Optical micrograph showing the as-received microstructure of CP grade 2 Ti.

in order to clarify the operative wear mechanisms. For subsurface studies, the worn samples were sectioned parallel to the sliding direction and then etched using Kroll’s reagent after polishing in order to observe microstructural changes. Spectral analysis of the worn surfaces was made using XRD.

The initial microstructure of CP grade 2 Ti is shown in Fig. 2. The microstructure of pure Ti consists of equiaxed coarse grains (CG) with an average grain size of about 110 ␮m. Figs. 3 and 4 demonstrate the TEM micrographs and selected area diffraction (SAD) patterns of the ECAE-processed billets following route-E for 8 (8E) and 12 (12E) passes in low and high magnifications. Fig. 3 was taken from the transverse plane of the as-processed billet while Fig. 4 was from the longitudinal plane. It is clearly seen from these micrographs that ECAE processing leads to a significant refinement in the microstructure. A homogeneous UFG structure with a mean grain sizes of about 300 nm for 8E and 250 nm for 12E are formed after ECAE. The grains are almost equiaxed with no discernable preferred morphological alignment in the microstructure of the transverse plane. Grain refinement and the homogeneity are more pronounced after 12-pass ECAE (Fig. 3(c) and (d)). In contrast, the longitudinal plane displays both equiaxed and slightly elongated grains (Fig. 4(a) and (b)). The elongated grains have their long axes parallel to the direction of shear. It is also noted from these figures that the microstructure consists of a mixture of grains with low-angle and high-angle boundaries that is typical of ECAEprocessed materials. This is suggested by the SAD patterns, having a high number of clustered spots. The micrographs also showed

Fig. 3. The TEM micrographs and SAD patterns from transverse planes of CP Ti billet processed by ECAE following route-E. Eight passes at (a) low and (b) high magnifications, and 12 passes at (c) low and (d) high magnifications.

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Fig. 4. The TEM micrographs and SAD patterns from longitudinal planes of CP Ti billet processed by ECAE following route-E. Eight passes at (a) low and (b) high magnifications, and 12 passes at (c) low and (d) high magnifications.

that dislocation cells and extremely high density of dislocations are the typical microstructural features in all ECAE-processed samples, indicating a non-equilibrium state with high internal stresses. High internal stresses are also revealed by the bend extinction contours inside many grains. The dislocations are distributed in such a way that the small grains are rather free of dislocations and most of dislocations are attracted to the grain boundaries. The SAD patterns present numerous spots arranged in circles that is typical for a fine granular type microstructure with high-angle grain boundaries and significant spreading of spots confirms the existence of high internal stresses [12].

gation to failure for 8E and for 12E billets, respectively. As compared with the elongation to failure of 33% in CG Ti, there is a decrease in ductility in UFG Ti which may be related to the decrease of strain hardening capability after ECAE [6]. After increasing the number of passes from 8 to 12, tensile ductility is somewhat improved with

3.2. Mechanical properties The true stress–true strain curves of the CG and UFG Ti samples are shown in Fig. 5. From this figure, it can be seen that the ECAE processing significantly affects the tensile properties of CG Ti. The summary of tensile test results taken from the stress–strain curves is given in Table 1. The table shows that the ECAE leads to a considerable increase in strength values with sufficient ductility. The yield strength values for 8E and 12E processed UFG Ti are about 102% and 84% higher than that of CG Ti, respectively, whereas the ultimate tensile strength values increased almost 43% and 45%, respectively. These improvements in strength values were obtained while maintaining a sufficient ductility of about 21% and 23% elon-

Fig. 5. True stress–true strain response of the CG and UFG Ti.

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Table 1 Summary of the tensile test results conducted on the CG and UFG Ti samples at room temperature. Processing states

Grain size (␮m)

 0.2 (MPa)

 u (MPa)

ı (%)

As-received CG Ti ECAE-processed UFG Ti (8E) ECAE-processed UFG Ti (12E)

110 0.30 (300 nm) 0.25 (250 nm)

307 ± 11 620 ± 3 564 ± 21

532 ± 18 760 ± 8 771 ± 17

33 ± 1 20.7 ± 4.2 22.8 ± 1.1

slight decrease in yield strength but an increase in ultimate tensile strength. This may be because the samples processed by ECAE for 12 passes have higher fraction of high-angle grain boundaries than the samples processed by ECAE for 8 passes [11]. A positive attribute is that the strain hardening rate also increased after 12 passes. From Fig. 5, the CG Ti exhibits significant strain hardening behavior and correspondingly large elongation. After ECAE, the stress–strain curves of UFG Ti shows that negligible strain hardening occurs as compared to that of the CG microstructure, which leads to a relatively early start of stress drop together with a localized neck formation, so-called ‘geometrical softening’ in stress–strain curves [25]. Park and Shin [26] attributed this diminished strain hardening rate to the decreased mean free path in the UFG materials. Also, dynamic recovery actively takes place during the tensile deformation of UFG materials due to the relatively fast spreading kinetics of trapped lattice dislocations, causing a negligible increase in the lattice dislocation density and resulting in the lack of strain hardening. Ko et al. [25] studied the possible reasons of low strain hardening of UFG pure Ti during tensile deformation and attributed it to a dynamic recovery process balancing the dislocation generation rate with the dissociation rate of trapped lattice dislocations in the non-equilibrium grain boundaries, i.e. as a kind of relaxation process in UFG structures. Clearly, almost one fold increase in yield strength of the CG Ti is obtained after 8 and 12 passes following route-E. Such an increase in strength values is related not only to the decreasing grain size (grain boundary strengthening) but also to effect of severe plastic deformation that caused high density of dislocations in the grain interiors (dislocation strengthening). This fact suggests that the UFG structure of pure Ti is very advantageous for improving its strength without alloying, so that it would be suitable for use in medical devices [25]. Also, this could open new frontiers in the applicability of pure Ti in more strength demanding environments and replace the widely used extensive Ti alloys. 3.3. Wear behavior The relation between the sliding distance and weight loss under an applied pressure of 1.5 MPa and at a sliding speed of 1.1 m s−1 for CG and UFG Ti is shown in Fig. 6. The weight loss of the samples in both conditions increases almost linearly with the sliding distance. The variation of weight loss of CG and UFG Ti with applied pressure is shown in Fig. 7. It can be seen that the weight loss of Ti in the CG and UFG states increases with increasing applied pressure and the increase is more pronounced at higher pressures. There is also a critical applied pressure above which the weight loss suddenly increases in both conditions. This pressure level for CG Ti and UFG Ti is about 2 MPa. It is likely that above such a critical pressure level the wear mechanisms changes from mild to severe wear. It is well known that Ti is one of the most reactive materials, readily interacting with O2 to form titanium oxide [20]. The high reactivity of Ti with relatively high temperature reached during rubbing results in a thick oxide layer on the surface which plays a dominant role in wear response of the materials. This layer stays stable up to a critical surface pressure. This results in a lower weight loss during this period of sliding due to the oxidative wear mechanism. Above this critical pressure level, the stability of surface layer deteriorates and it gets smeared off during rubbing the surface, and consequently

Fig. 6. The change in weight loss as a function of sliding distance for CG Ti and UFG Ti processed by ECAE following 8E and 12E, during the dry sliding wear experiments on a pin-on-disc apparatus.

the severe wear regime starts which lead to substantial decrease in wear resistance. The results from Figs. 6 and 7 surprisingly show that the strengthening of pure titanium by multi-pass ECAE processing did not cause any considerable improvement in wear resistance at least for the pressure levels and sliding distances used in this study. Similar results were also obtained in few previous studies [22,27] in both CG and UFG pure Ti. The main factor affecting the wear resistance of CG and UFG Ti is seen to be the dominant wear mechanism operating during sliding. It is well known that the dominant wear mechanism is the oxidative wear in pure Ti and also Ti alloys [27], which controls the weight loss during sliding. Besides, the adhesion, materials transfer between contacting surfaces, and abrasion are the other operative wear mechanisms, but they are not as effective as oxidative wear mechanism [22]. The studies on wear surface

Fig. 7. Variations in weight loss of CG and UFG Ti with applied pressure during the dry sliding wear experiments on a pin-on-disc apparatus.

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showed that a thick mechanically unstable oxide layer appeared on the sliding surface of the samples in both CG and UFG conditions as a result of tribochemical reactions and frictional temperature [28]. High affinity of Ti for oxygen results in the formation of adherent surface oxides by tribochemical reactions [29]. It has been proposed that when normal or shear stresses are high enough to introduce breakdown of the surface passive layer, the oxide will be disrupted [30]. This oxide layer is easily removed by spalling or microfragmentation during sliding [31]. The exposed metal surface may then either reform a passive layer or adhesively bond to the counter surface. The latter situation leads to continuous removal and reformation of the passivating layer and results in gradual consumption of Ti material. Concurrently, the roughness of the material increases which results in yet higher metal loss. Ultimately, the breakdown of the oxide layer creates the potential for abrasive wear, where the hard oxide debris acts as third body abrasive components. In such friction and wear conditions, the effect of titanium strengthening by ECAE processing on the improvement of wear resistance seems to be unimportant. In other words, in the present wear experiments titanium oxide layer played a dominant role in determining the wear resistance of CG and UFG Ti, and it masked the effect of strengthening of UFG structure on the wear resistance. The previous studies on friction and wear behaviors of pure titanium and its alloys in the CG and UFG forms have also shown poor wear characteristics [27,30]. The poor wear behavior was attributed to the properties of the oxide layer formed on the surface during sliding, its poor integrity with the matrix, and the deformation behavior of subsurface regions [32]. CP titanium, a relatively low shear strength hcp material, exhibited higher friction coefficient values, but also greater material transfer, due to its high reactivity, to the counterfaces, than higher strength materials [30]. Hence, pure Ti and Ti alloys was considered to have poor oxidative wear resistance when tribochemical reaction occurs at the contact area. XRD analysis in previous study [30] showed that TiO2 was a dominant oxide, suggesting that the formation of TiO2 during tribo-oxidation destroys the protective oxide layer and therefore increases friction and wear rate. XRD analysis of worn surfaces of the CG and UFG Ti samples was also conducted in this study and the result is shown in Fig. 8. From this figure, the surfaces of both CG and UFG (12E) Ti were mainly covered by anatase and rutile modification of TiO2 during sliding. Also, ␣-Ti peaks appeared on the XRD pattern. This was basically due to the penetration of X-ray beyond the oxide layer through the Ti matrix. In addition, the XRD spectra of titanium oxide on the worn surface of the CG Ti are identical to those of the UFG Ti, suggesting that the UFG structure did not affect the wear products on the sample surface and thus improvement in wear resistance was not observed due to the dominant oxidative wear mechanism. The worn surface of pure Ti was also examined by La et al. [20] using XRD analysis and they found that the XRD spectra of TiO2 on the worn surface of the CG Ti are identical to those of the UFG Ti. The surface and subsurface studies also supported this finding as explained below. Thus, an almost similar wear resistance was observed for both CG and UFG Ti as seen in Figs. 6 and 7, in spite of the strengthening in UFG Ti. The morphologies of the worn surfaces of the CG and UFG Ti after wear test for a sliding distance of 10.40 km under a pressure of 1.5 MPa and at a sliding speed of 1.1 m s−1 are shown in Fig. 9. Generally no qualitative differences in surface topography were observed among the worn surfaces of CG and UFG titanium after wear. Typically, pure Ti in both conditions revealed extensive wear-induced plastic deformation and surface damage. The worn surfaces exhibit mainly a limited number of deeper and defined grooves, some micro-cracking and gouging along with smearing over the worn surface. Presence of deep grooves and scratches on the wear surfaces could be attributed to the abrading action of the hard debris particles in the form of titanium oxide and asperities of mating surfaces.

Fig. 8. XRD spectra of worn surfaces of CG (as-received) and UFG (12E) Ti tested at a sliding speed of 1.1 m s−1 under a pressure of 1.5 MPa for a sliding distance of 10.42 km.

Some of the debris particles get entrapped in between the mating surfaces and abrade the specimen surface during sliding [33]. The smearing effect due to the adhesion is also evident on the surface of the samples of the CG and UFG Ti. The smeared materials had been detached from the sample surface by adhesion to the surface of the steel disc. During subsequent sliding, some of the transferred material was lost and some re-embedded and smeared over the sample surface, with some degree of oxidation [34]. These observations show that the main wear mechanisms predominating during friction of CG and UFG titanium against steel counterface are tribochemical reaction leading to oxidative wear with the abrasive effect, and adhesion by material transfer and back-transfer. The micrographs in Fig. 9 also shows that worn surface of the CG Ti are identical to those of the UFG Ti, suggesting that the UFG structure did not affect the wear products and also dominant wear mechanisms. Fig. 10 shows the cross-sectional view of the wear surfaces of the CG and UFG Ti samples taken parallel to the sliding direction. Microstructural changes were observed in the subsurface regions due to the wear-induced plastic deformation involving thermal and thermomechanical processes. The micrographs in Fig. 10 reveal mainly three distinct regions near the surface. The upper-most region called tribolayer mainly consists of oxide layer with smeared and/or embedded wear material. This layer seams to be weakly bonded to the bulk and shows microcracks along the sliding direction [28]. Additionally, the smeared and embedded wear materials can be seen in the tribolayer. This layer forms as a discontinuous discrete layer of compacted wear fragments suggesting that mechanical instability of this layer was responsible for the high friction leading to high wear rate. The tribolayer is very indicative in all micrographs of CG and UFG Ti. The maximum thickness of the layer is roughly 5 ␮m in the CG Ti, 7 ␮m in the UFG after 8E pro-

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Fig. 9. SEM micrographs showing the appearance of worn surfaces of (a) CG Ti and UFG Ti processed using ECAE through (b) 8E and (c) 12E. The wear tests were performed for a period of 2.63 h corresponding to a sliding distance of 10.40 km under a pressure of 1.5 MPa and at a sliding speed of 1.1 m s−1 .

cessing and 8 ␮m in the UFG Ti after 12E processing, which mainly show an increase in the thickness of the layer after ECAE processing. Moreover, the thickness of this layer increases with increasing number of passes. The region underneath the tribolayer is called

as deformed region and constitutes the transition region between tribolayer and unaffected bulk material. The severity of deformation reduces and material flow oriented along the sliding direction can also be seen in this region. The deformed region is thin in the

Fig. 10. Subsurface regions of (a) CG Ti and UFG Ti processed using ECAE through (b) 8E and (c) 12E. The wear tests were performed for a period of 2.63 h corresponding to a sliding distance of 10.40 km under a pressure of 1.5 MPa and at a sliding distance of 1.1 m s−1 .

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CG Ti compared to the UFG Ti. The bottom region, called unaffected bulk material has the original microstructure practically without any wear-induced deformation. It may be noted that, the extent of wear-induced deformation decreases gradually away from the wear surface ultimately becoming undetectable in the unaffected bulk material. 4. Conclusions The tensile properties and wear characteristics of pure Ti processed by multi-pass equal-channel angular extrusion (ECAE) following route-E were examined. The main findings and conclusions of this study can be summarized as follows. • Equal-channel angular extrusion of coarse-grained (CG) Ti produces homogeneous ultra-fine grained (UFG) structures with a mean grain sizes of about 0.3 ␮m (300 nm) for 8E processed samples and 0.250 ␮m (250 nm) for 12E processed samples as compared to the CG Ti with the grain size of about 110 ␮m. • Significant improvement in the strength values of CG Ti was obtained via ECAE processing because of refinement grain size down to UFG regime and increasing dislocation density without alloying. Furthermore, this process retains adequate ductility (more than 20%). • ECAE processing has no significant effect on the sliding wear behavior of CG Ti. Lack of discernible improvement in wear resistance after ECAE is surprising considering the exceptional increase in strength. This was mainly attributed to the operating wear mechanisms in which the titanium oxide surface layer plays a dominant role in determining the wear resistance of CG and UFG Ti, and it masks the effect of strengthening due to UFG structure on the wear resistance. • Microstructural changes were observed in the subsurface region of the worn CG and UFG Ti samples due to the wear-induced plastic deformation involving thermal and thermomechanical processes. The three distinct regions detected are the tribolayer including titanium oxide with smeared and/or embedded wear material at the top, a deformed region with material flow oriented in the sliding direction in the middle and the original unaffected bulk material at the bottom. The dominant wear mechanism of the CG and UFG Ti is the tribochemical reaction leading to oxidative wear with the abrasive effect. Acknowledgements This research was supported by The Scientific and Technological Research Council of Turkey (TUBITAK), Engineering Research Group (MAG), under Grant No: 104M289, and by Scientific

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