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(TiN + TiB) hybrid composite layers produced using liquid phase process
Materials Chemistry and Physics xxx (2014) 1e11
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Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process R. Yazdi*, S.F. Kashani-Bozorg School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, P.O. Box 11155-4563, Tehran, Iran
h i g h l i g h t s In-situ Ti/(TiN þ TiB) hybrid composite layers were synthesized by TIG processing on commercially pure titanium. The microstructure features were characterized by several methods. Microhardness enhanced three to five times higher than that of the CP-Ti substrate after surface modification. The fabricated composite layers improved wear resistance of CP-titanium. Severe adhesive wear mechanism of CP-titanium surface changed to mild abrasive one as a result of the surface modification.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 September 2014 Received in revised form 23 November 2014 Accepted 20 December 2014 Available online xxx
Tungsten inert gas (TIG) technique was conducted on commercially pure (CP)-Ti substrate, which was coated with h-BN-based powder mixture prior to the treatment. The treated surfaces were evaluated and characterized by means of scanning electron microscope (SEM), X-ray diffraction analysis, and electron dispersive spectrometry (EDS). The microhardness and wear experiment were also performed by using a microhardness machine and pin-on-disk tribometer. As h-BN reacted with titanium, an in-situ hybrid composite layer was formed showing near stoichiometric dendrites of TiN, platelets of TiB and interdendritic regions of a0 -Ti martensite crystal structures. The population level of TiN and TiB regions were found to increase using a pre-placed powder mixture with greater h-BN content. However, the fabricated layers exhibited cracking and porosity; these were minimized by adjusting arc energy density and h-BN content of powder mixture. The microhardness value of the fabricated hybrid composite layers was found to be in the range of ~650 HV0.2e1000 HV0.2; this is three to five times higher than that of the untreated CP-Ti substrate. In addition, the in-situ hybrid composite layers exhibited superior wear behavior over CP-Ti substrate; this is attributed to the formation of newly formed ceramic phases in the solidified surface layers and good coherent interface between the composite layer and CP-substrate. Meanwhile, severe adhesive wear mechanism of CP-titanium surface changed to mild abrasive one as a result of surface treatment. © 2014 Elsevier B.V. All rights reserved.
Keywords: Welding Coating Composite materials Microstructure Hardness Wear
1. Introduction Titanium and its alloys have some excellent properties such as a good corrosion resistance and high strength-to-weight ratio. Its light weight and ability to resist high temperature make it suitable for aerospace applications and their good biocompatibility make them suitable as implants for dentistry and orthopedic applications [1,2]. A disadvantage of titanium-based alloys is their high friction
* Corresponding author. E-mail address: [email protected] (R. Yazdi).
and poor wear resistance [3,4]. To overcome this problem, various surface composite layers were fabricated on Ti-based alloys using hard ceramic particles such as SiC [5,6], B4C [7], TiN [8], TiC [9,10], TiB2 [11,12], and Ti3Al [13]. These hard particles were introduced into the surface melt pool generated by different localized heat sources e.g., electron beam [9,11], plasma torch [7,14], laser beam [5,7]. However, such composite layers are usually associated with weak interfacial particulate/metal matrix bonding strength when the material is under cyclic loading. Reactions between boron (B) and nitrogen (N)-bearing ceramic particles with titanium are of special interest due to high affinity of Ti for boron and nitrogen and high thermodynamic stability of TieB
http://dx.doi.org/10.1016/j.matchemphys.2014.12.026 0254-0584/© 2014 Elsevier B.V. All rights reserved.
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and TieN compounds [15]. Euh et al. [16] showed that upon electron beam irradiation of Tie6Ale4V alloyed with pre-placed TiB2based powder layer, ceramic powder was melted and an in-situ surface composite layer was fabricated, which is reinforced with the newly formed phases. Further investigation by Kooi et al. [17] revealed orientation relationships between the titanium matrix and newly formed intermetallic compounds; the coherent interface can be beneficial to the precipitates/matrix bonding strength. According to Mitun Das research [18], the surface of Tie6Ale4V alloy could be modified by laser processing using mixture of BN and Ti powders. Although there are several features associated with techniques such as plasma, laser and electron beam surface modification, but processing time, flexibility in operation, economy in time, energy, and material consumption, and processing precision are the important advantages of tungsten inert gas (TIG) surface alloying process. Limited works were reported on the applications of this process for surface modification of titanium alloys [13,19,20]. In the present investigation, TIG surface alloying of commercial purity titanium (CP-Ti) substrate pre-placed by h-BN-based powder mixtures was conducted under an atmosphere of pure argon gas. Microstructural characterization of the fabricated surface composite layers was carried out. In addition, changes in surface wear characteristics were evaluated.
using TIG process on CP-titanium with h-BN-based layer. Surface melting of the pre-coated substrates was carried out using a TIG torch. A tungsten electrode with a diameter and height of 2.4 and 2 mm was used to create an arc between the tip of the electrode and the substrate surface, respectively. The voltage (V) on tungsten electrode was maintained at 17 V. TIG surface alloying was carried out in a chamber using argon shielding gas with a purity of 99.999% at flow rate of 10 l min1. Alloyed tracks were produced using different substrate traverse speeds (100e200 mm min1) and TIG
2. Experimental procedure Ti metal plates (purity 99%) of 7 mm thickness were supplied. They were cut into 100 mm 25 mm pieces. Contaminations such as oxides, dirt, grease, etc. were removed from the surface of the substrate by grit blasting. All of the grit-blasted test specimens were cleaned in an acetone ultrasonic bath. Fine boron nitride and Ti powder (with a mean particle size of ~10 and ~150 mm, respectively) were blended to make up three powder mixtures containing 5, 10 and 15 mass% h-BN. Before TIG processing, powders were mixed homogeneously using ball mill for half an hour. Organic binder (polyvinyl alcohol) was used to pre-place the powder mixture onto CP-Ti substrates; this prevents powder scattering during fabrication. However, the amount of binder was restricted to only 1 mass% in order to limit pore formation. Finally, the powder mixture was hand-brushed on the substrate surface to a thickness of ~0.8 mm. The pre-coated substrates were then dried in an oven at 150 C for two hours. Fig. 1 shows schematic processing of an in-situ composite layer
Fig. 1. Schematic presentation of the fabrication of an in-situ composite layer using TIG process on the pre-placed titanium with h-BN-based layer.
Fig. 2. Transverse sections of the fabricated composite layers using pre-placed layer with 5 mass% h-BN (a and b, employing TIG arc energy densities of 112 J mm2 and 349 J mm2, respectively) and 10 mass% h-BN (c, employing TIG power density of 349 J mm2) contents.
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currents (60e80 A). TIG heat inputs were calculated using Equation (1) [19]:
Arc Energy Density J mm2 ¼
3
Current ðIÞVoltage ðVÞ Arc Radius mm Velocity mm s (1)
Specimens were cut transversely to the alloyed tracks and polished using standard metallographic procedure and then etched in a solution containing 5 ml HF, 15 ml HNO3 and 80 ml H2O for 10 s. An X-ray diffractometer (Philips X'Pert Pro) was used for the determination of the formed phases on the fabricated surface layers. In addition, microstructural examination was carried out using an optical and a Cam Scan (MV2300) scanning electron
Fig. 3. (a) SEM micrograph of the cross section of an in-situ composite layer produced using pre-placed layer with 5 mass% h-BN content and a TIG arc energy density of 349 J mm2 exhibiting (b) dendritic and (c) platelet structures.
Fig. 4. Energy dispersive spectrum of (a) a fabricated composite layer, (b) dendrite, and (c) platelet.
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Fig. 5. X-ray diffraction patterns of the fabricated composite layers produced using 5 mass% h-BN (a) and 10 mass% h-BN (b) powder in the pre-placed layer at various TIG arc energy densities: (a) 112 J mm2 and (b) 349 J mm2.
microscope (SEM). The former was linked with an Oxford Inka200 energy dispersive spectrometer (EDS) for the determination of chemical composition. The EDS device has an ultra thin atmospheric beryllium window capable of chemical analysis of boron and nitrogen. Hardness measurements were carried out using a microhardness tester (Micro met) under a 1.962 N load on cross sections of the fabricated surface layers. Dry wear evaluation of the fabricated surface layers and asreceived CP-Ti substrate was carried out using pin-on-disk testing which performed under a constant load of 15 N and at a speed 0.6 m s1. Before measurements, the specimens were ultrasonically cleaned in methanol and dried with blown air. The pin-on-disc tests were performed against quenched and tempered (AISI 52100) steel discs with a hardness of ~61 HRC. Weight losses were measured after certain sliding distances, by weighing the samples to an accuracy of 104 g using an electronic microbalance. Surfaces of the worn samples were investigated using SEM and EDS. 3. Results and discussion 3.1. Effect of process parameters Typical transverse sections of the fabricated tracks are shown in Fig. 2. All sections exhibited a conduction-limited profile. Surface composite layers were achieved with a depth of up to ~2 mm. As it can be seen in Fig. 2a, spherical-type pores were present in the treated track fabricated on the pre-placed CP-Ti with h-BN amount of 5 mass% using arc energy densities of 112 J mm2. Gas entrapment within the fusion zone due to relatively higher rate of solidification process at the lowest arc energy density (i.e., 112 J mm2) forms spherical pores. No pores were present in the tracks using higher arc energy density (e.g. Fig. 2b). Increasing power density decreases substantially the rate of solidification, which provides more time for releasing the unreacted gases from the surface molten layer. Contrary to tracks produced on the pre-placed CP-Ti with h-BN amounts of 5 and 10 mass% (Fig. 2), no arc stability was reached on pre-placed layer containing 15 mass% h-BN. Phase h-BN was
reported to be an excellent dielectric [15]. Therefore, arc discontinuity can be attributed to insulating property of h-BN. It means that there was inadequate titanium powder density in the pre-placed layer providing sufficient arc stability. Fig. 2c also exhibits a processed layer with cracking produced using pre-placed layer with 10 mass% h-BN contents and employing TIG power density of 349 J mm2. In the case of pre-placed CP-Ti with h-BN amount of 10 mass%, tendency to cracking was observed using various energy densities. TieN and TieB phase diagrams [21,22] show that, nitrogen has high solubility in liquid and solid titanium (<8 mass%) due to small atomic radius (0.75 A), and boron with larger atomic radius (1.17 A) has very low solubility (<0.05 mass%) in solid titanium, even in the high-temperature beta phase [23]. The rapid cooling associated with the TIG process may cause metastable retention of boron in the titanium interstices, impeding motion of the geometrically necessary dislocations (GND) needed to accommodate the thermal contraction and the great differences between the coefficient of thermal expansion (CTE) of the surface composite layers and CP-Ti substrate is believed to the punch-out dislocations around the second phase particles and cracking happened consequently. Another factor that may contribute to cracking is the possibility that some amount of nitrogen and boron from dissociated h-BN is dissolving in the molten metal zone. Due to high cooling rate, excessive nitrogen and boron content in the quenched titanium-based solid solution could generated a brittle crystal structure; which is susceptible to cracking [24]. 3.2. Microstructures SEM micrographs and EDS analysis of a fabricated surface layer have been presented in Figs. 3 and 4, respectively. Microstructural evaluation of the cross sections showed dendritic-type in the interdendritic structure (Fig. 3a and b). The secondary arms of the longer dendrites exhibited orthogonal angle with the main stem, revealing a cubic crystal structure. Also, relatively long pillar shaped (platelets) phase were observed in the microstructure (Fig. 3c). In addition, EDS area analysis of the fabricated layers displayed substantial concentration of B and N (Fig. 4a). However, no h-BN
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Fig. 6. SEM micrographs of cross sections of the in-situ composite layers produced using 5 mass% h-BN powder at various TIG arc energy densities: (a) 112 J mm2, and (b) 234 J mm2, showing dendritic and platelet structures in a0 -Ti structure.
particle was detected in the microstructure. Initially, melting occurred in the substrate and titanium powder when applying the liquid processing. The temperature of the molten zone is much higher than that of Ti substrate (1660 C), as a result, all of the BN powder could be uniformly surrounded by liquid titanium into the molten pool. Then, h-BN dissipated in the molten zone during the TIG processing and due to high affinity between titanium and boron or nitrogen [15], more stable phases were formed in the structure. The data are in agreement with the results of Yeh et al. [25] who detected no BN residual in the product of self-propagating hightemperature synthesis of compacted Ti and BN powders. Also, consistency was found with the results of laser surface alloying [18,26]/cladding [27,28] of Ti-based alloys pre-placed with BN layers since no BN was observed in the laser surface alloyed/clad
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layers. Although the concentration of nitrogen and boron, which are very light elements, could not be measured with the sufficient degree of accuracy by EDS analysis, as illustrated in Fig. 4b and c, EDS point analysis on the dendrites exhibited nitrogen and titanium enrichment, and high concentration of Ti and B on pillar shaped phase were shown by EDS point analysis. In order to identification of the phases with different morphologies, XRD is used to analyze. Diffraction patterns of the in-situ composite layers produced using pre-placed CP-Ti with h-BN amounts of 5 and 10 mass% were illustrated in Fig. 5. Reflections with higher intensities were found to be in agreement with TiN (having NaCl type crystal structure), TiB (close-packed hexagonal) and a0 -Ti (close-packed hexagonal) crystal structures. The XRD patterns of the treated surfaces were found to be consistent with the quantitative EDS results (Fig. 4b and c). TiN can form over a wide range of nonstoichiometric compositions with various N:Ti ratios [29]. The near stoichiometric result of the present investigation, which has a dendritic structure (Fig. 3b) is similar to solidstate reactions but unlike substoichiometric TiN produced in vaporphase reactions [30]. TiB phase with platelet morphology (Fig. 3c) was previously observed by many investigators [31]. The titanium borides, especially TiB, are reported to form in a hexagonal pillar shape [26]. Morikawa et al. [32] found clusters of TiB particles with rod-like shape along the grain boundary in the cast Tie1Be1N alloy. In the present work, relatively smaller features were observed; this is attributed to higher cooling rate of TIG compared to cast processing. But, the TiN and TiB phases in the microstructure are coarser compared with those produced with higher energy laser processing. Because lower melt pool heat input will have a comparatively shorter time for growth of precipitates in laser processing, which will result in finer particles [18]. While overlapping with indications of other phases, few indications may correspond to XRD pattern of Ti3B4 crystal structure. But, no microstructural feature related to this structure was detected using SEM and EDS. The existence of Ti3B4 was also reported by Fenish [33] and Neronov et al. [34]. It can be seen that the intensities of XRD peaks of TiB in the pattern do not match the standard intensities (Fig. 5), and they suggest a preferred orientation. The XRD patterns reveal that prominent peaks are present in the surface composite layer synthesized at 349 J mm2 arc energy density and weaker intensities are in the fabricated tracks treated with lower heat inputs (e.g., Fig. 5). The arc energy plays a significant role in the alloying process. Increasing the arc energy decreases the cooling rate and an adequate time will be available for alloying. So, it is reasonable to say that decreasing the arc energy density lowers the alloying process. This idea approved by comparing SEM micrographs of the treated tracks using (a) 112 J mm2, (b) 234 J mm2, and (c) 349 J mm2 TIG arc energy densities which are shown in Figs. 6a,b and 3a, respectively. Therefore, the population level of the randomly grown dendrites and platelet phases increase by increasing the power density from 112 to 349 J mm2. However the strength and stiffness of titanium-base alloys may be increased by the precipitation of TiB intermetallic phase [18]. The presence of high amount of brittle TiN phase in the microstructure is another contributing factor to the crack occurrence of the track fabricated with 10 mass% h-BN presented in the pre-placed layer [35]. The random growth of dendrites and platelet phases in the surface composite layers can be attributed to the stirring of the melt pool by convection and Marangoni flow due to temperature gradients and gradient of the surface tension in the melt pool [36]. The concentration of gradient of diffused nitrogen and boron by this stirring action can result in a homogeneous microstructure. The mechanism of the in-situ composite layer formation during
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Fig. 7. The schematic diagrams illustrating the mechanism of the in-situ composite layer formation with pre-placed h-BN-based powder layer using TIG process.
the TIG processing has been schematically illustrated in Fig. 7. In the first stage, the h-BN powder dissociated into nitrogen and boron atoms; which then solved into and also reacted with liquid titanium (Fig. 7b). The solidification sequence of the in-situ composite layers is deduced to be formation of TiN dendrites with melting point of 3600 C from the liquid. The remaining liquid is then depleted in nitrogen and enriched with boron (Fig. 7c). Boron element with 2.5e7.5 mass% in the melting pool is a favorite condition for the platelet-like TiB phase formation (Fig. 7d); having 2200 C melting point [18]. After solidification completion, high cooling rate leads to transformation of the solidified b-Ti interdendritic regions to a0 -Ti crystal structure upon cooling to room temperature via a solid state martensitic reaction. The a0 -Ti is believed to be within the interdendritic regions. Although the Ti-rich corner of the TieBeN ternary system includes the various phases [37] and there are several reports on the ternary TieBeN coatings but no chemical composition was often specified [38] and in some investigations [27,37], no ternary
compound was reported in the TieBeN ternary system. Considering the phases characterized in the microstructure, the following reactions may occur in the TIG alloyed layers: BN (s) þ Ti (L) ¼ TiN(s) þ B(s)
(2)
Ti (l) þ B(s) ¼ TiB(s)
(3)
Also, other reactions were proposed by researchers [39]: 3Ti (l) þ 4B(s) ¼ Ti3B4(s)
(4)
Ti (l) þ 2B(s) ¼ TiB2(s)
(5)
The molar Gibbs reaction energy for the four equations as a function of temperature was reported by Lu and Zhang [40,41]. Calculations of Ding et al. [42] showed negative values for the molar Gibbs reaction energy for all of the above equations from 527 C to
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Fig. 8. SEM micrograph exhibiting a martensitic morphology in the heat affected zone (HAZ) of an in-situ composite layer produced on CP-Ti substrate with 5 mass% h-BN content using a TIG power density of 234 J mm2.
Fig. 9. Microhardness profiles of the fabricated composite layers produced with 5 mass% h-BN content using various TIG energy densities.
Fig. 10. Weight loss data of the worn samples of CP-Ti substrate and the composite layers produced with 5 mass% h-BN content using two different TIG energy densities (234 and 349 J mm2).
927 C. However, TiB2 possesses the lowest value. Few reports were found TiB2 on melting reactions at BNeTi interfaces in the literature [26,27,42]:
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(a) Molian and Hualun [27] employed laser surface cladding of Tie6Ale4V with BN powder and reported of the presence of TiN, TiB2 compounds in addition to a0 martensite. Their results were mentioned to be based on the XRD. However, no X-ray diffraction pattern was presented [27]. Contrary to their result, TiB2 was not detected in the present investigation. The greater substrate dilution of the present work compared to laser surface cladding provides excessive Ti for the formation of TiB instead of TiB2. In addition, due to the arc stability limitation, the presence of titanium powder (~95 mass%) in the h-BN based pre-placed layers enhanced the melting zone dilution and this condition also encouraged less thermodynamically stable phase formation, i.e. TiB and Ti3B4. (b) Senthil et al. used laser surface alloying of Tie6Ale4V with pre-placed BN layer and found TiB, Ti3B4 and Ti2N phases in the alloyed layer [26]. Their results were supported by energy dispersive XRD and X-ray photo spectroscopy analysis. Again it can be stated that the higher substrate dilution in laser surface alloying than laser surface cladding can be responsible for the formation of TiB instead of TiB2. But, Ding et al. [42] did not observe Ti3B4 and TiB in the interfacial reaction between BN and Ti during active brazing. The relative low Ti content of the filler can be responsible for the formation of TiB2. This scenario is in agreement with Yeh and Teng [25] who detected TiN and TiB2 products from the reactant of 2Ti þ BN in self-propagating high-temperature synthesis chamber, but in addition to the phases of TiN and TiB2 intermediate compounds including Ti2N and TiB were detected in the final products from the reactants of 3Ti þ BN and 4Ti þ BN. The presence of Ti2N and TiB is mainly attributed to the incompleteness of nitridation resulting from the excessive melting of Ti. Another crystal structure, which was not detected in the present work, is Ti2N. Faran et al. [15,43] investigated the interaction between BN and Ti diffusion couple at 1000 C. The phase sequence in the (Ti powder)/(BN plates) diffusion couple was shown to be Tie(a0 -Ti)eTi2NeTiNeTiBeTiB2eBN. Fine dispersive Ti2N precipitates were detected using high resolution SEM. They could only be formed during annealing in the reaction zones of BN/Ti interfaces below ~1050 C. SEM image of heat affected zone of the treated track has been illustrated in Fig. 8. The treated track microstructure indicates good bonding with the substrate and smoothly metallurgical changes from the interface up to the surface. In addition, no interfacial cracks, and other defects such as lack of fusion, and pores were observed throughout the layers' cross-section. This is due to the complete melting and reaction of pre-placed layer and the substrate at the interface. At the bottom of the composite layers, a relatively fine transformed product can be observed within the substrate heat affected zone grains. The transformation products exhibited almost typical of a0 -Ti martensitic baskets [44]. Such features were reported by various investigators in the heat affected zone of CP-Ti processed by techniques like TIG [20] and etc. [27,45]. The temperature of the heat affected zone was reached to near melting point and due to heat sink effect of the underlining substrate, a0 -Ti martensite formed upon rapid cooling by the reaction of b-Ti to a0 -Ti. Meanwhile, the volume of the heat affected zone was found to increase with increasing TIG heat input. 3.3. Hardness The measured microhardness values of the in-situ composite layers from the near top surface down to the unaffected substrate are shown in Fig. 9. Results show that the microhardness value
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Fig. 11. SEM micrographs of the worn samples: (a) the untreated CP-Ti substrate and (b) the composite layer produced with 5 mass% h-BN content using a TIG power density of 349 J mm2 after a sliding distance of 1000 m.
increases with increasing TIG heat input (i.e., the population level of these hard intermetallic phases). Maximum microhardness values of the in-situ composite layers fabricated at the energy densities of 112, 234 and 349 J mm2 were 650, 874 and 993 HV0.2, respectively. These values are around three to five times higher than that of the CP-Ti substrate (~200 HV0.2). TiB and TiN Phases have the hardness of about 1800 HV and 2100 HV, respectively [46]. So, the formation and the amounts of new phases (titanium boride and nitrides) are believed to be responsible to the increase in the microhardness value. The TiB platelet precipitate can be considered as a fiber
reinforcing element [17]. Also, the hardness increase may benefits from the presence of hard TiN phase and effect of probable lattice deformation or coherent hardening in the surface composite layers. The small variations in hardness values detected in the composite layer derived from hardness measured locations where different amounts of the fabricated phases were present. The maximum obtained microhardness value is higher than that (650e850 HV) achieved on laser boronising of Tie6Ale4V with pre-placed BN using different process parameters [26]; this may be due to greater population level of hard intermetallic phases
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Fig. 12. EDS analysis of wear debris on worn surface of CP-Ti substrate reflecting a relatively significant O, Fe and Cr content from the disk.
achieved in the in-situ composite layer fabricated by TIG than that of laser surface alloying. Although higher localized heat can be provided by laser but greater heat input is believed to be associated with TIG surface processing of the present investigation. However, the attained microhardness values are lower than those (1600 HV) achieved on laser cladding of Tie6Ale4V with BN [27]; this is believed to be due to lower substrate dilution in cladding than that of alloying process, which provides a greater population level of hard phases in the latter. 3.4. Wear assessment Fig. 10 shows the wear-loss data of CP-Ti substrate and the insitu composite layers fabricated by TIG as a function of sliding distance. The results reveal that weight loss increases as sliding distance increases, irrespective of microstructure features (A nearly linear relationship was found between mass loss and sliding distance). CP-Ti sample exhibited a relatively significant mass loss during sliding process; this is an indication of a relatively high rate of wear. It is reported that titanium alloys have severe wear behavior over a wide range of sliding speed [47]. The results in this study also show that the untreated titanium surface suffers high wear. Contrary to CP-titanium, the in-situ composite layer showed nearly four times lower material removal. Thus, the surface alloying process improved the wear resistance of CP-Ti substrate. According to Archard [48], the wear resistance of a material is proportional to its hardness value. The decrease of wear-loss data was found to be consistent with the increase in microhardness value of the specimens (Fig. 9). Also, the increase in the population level of hard phases in the structure (Figs. 3a and 6) is associated with more improved surface wear properties of the in-situ composite layers. This behavior could be attributed to the presence of hard intermetallic phases in the microstructure of the composite layer. SEM images of wear surfaces and their EDS area analysis are depicted in Figs. 11 and 12, respectively. It could be observed that
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Fig. 13. EDS analysis of wear debris on worn surface of the composite layer reflecting a relatively lower content of Fe and Cr.
surface of untreated CP-titanium sample was severely scored and plastically deformed, which caused a significant roughed worn surface. These features of the surface is indicative of adhesive wear mechanism, which is approved by other researchers [10,47]. The worn surface of CP-Ti also revealed relatively deep grooves and scratches aligned in the wear direction (Fig. 11a) and also with respect to EDS analysis, the wear debris of CP-Ti worn sample reveals significant contamination content of oxygen (Fig. 12); this could be due to the formation of titanium oxide during the sliding due to high temperature at the interface and also, significant Fe and Cr contamination from disk to the pin surface was detected. These elements generated transferred layer during the sliding. The transferred layer associated with the adhesion became workhardened and oxidized after multiple contacts in the wear couple, which in turn resulted in micro-cutting wear damage on the Ti surface. In comparison with wear track of CP-titanium, a distinct change in surface wear mechanism occurred, as exhibited in Fig. 11b and no significant adhesive wear features were detected on the worn surface of the fabricated composite layers by TIG. Likewise, the sample fabricated at 349 J mm2 had a better wear resistance than the specimen processed at lower arc energy (234 J mm2); this is due to the higher hardness of coating and to the coherent interfaces of the hard intermetallic phases and Ti matrix as reported by Kooi et al. [17]. Initially, comparatively softer a0 -Ti worn significantly and, then TiN and TiB phases protruded during sliding contact and prevented substantially high material removal and improved wear resistance of the composite surface. The other cause of the superior wear resistance of hybrid composite layer is less adhesion between the sample surface and steel counterpart in sliding contacts compared with CP-titanium surface. Generally, ceramics have less tendency to adhere firmly to metal surface. So, protruded ceramic phases in the structure prevented from adhesive metal-to-metal contact of the wear couple and subsequently less weight loss arose from sliding surfaces [49]. In the further sliding, the protuberant particles may
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be fractured due to the effects of less support by the matrix and stress concentration exerted on each individual particle and cause that mild scratching, shallower grooves and little material removal have been observed on the wear treated sample after a sliding distance of 1000 m (Fig. 11b). In addition, EDS analysis of the wear debris of the worn sample of the composite layer has been shown in Fig. 13, which reveals minor contamination content of Fe and Cr from disk to the pin surface. From these wear characteristics of hybrid composite layer, it could be concluded that mild abrasion is the prevalent wear mechanism. 4. Conclusion The in-situ composite layers were fabricated using TIG process on CP-Ti pre-placed with h-BN at different arc energy densities, and their microstructures, hardness and wear resistance were investigated with following conclusions. 1. Microstructural characterization of the fabricated composite layers showed that the melt zones consisted of TiN dendrites, TiB platelets and a0 -Ti martensite matrix. The population level of these hard phases increased with increasing the TIG arc energy densities. 2. The presence of the hard phases improved the microhardness value of the surface composite layers up to five times greater than that of the CP-Ti substrate. 3. The wear resistance of the fabricated composite layers was found to be about four times better that of the untreated CP-Ti substrate. This due to higher hardness and lower tendency of Ti/(TiN, TiB) composite for adhesion to steel counterpart as well as strong coherency between composite layer and CP-titanium substrate. 4. Severe wear mechanism of CP-titanium surface changed to mild abrasive one as a result of the surface modification. Acknowledgment The authors wish to thank University of Tehran for partial financial support. References [1] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Ti based biomaterials, the ultimate choice for orthopaedic implants e a review, Prog. Mater. Sci. 54( (2009) 397e425. [2] Sebastian Bauer, Patrik Schmuki, Klaus von der Mark, Jung Park, Engineering biocompatible implant surfaces: part I: materials and surfaces, Prog. Mater. Sci. 58 (2013) 261e326. [3] Xiu-Yan Li, Pei-Qiang Wu, Bin Tang, J.P. Celis, Fatigue behavior of plasma surface modified Tie6Ale4 V alloy, Vacuum 79 (2005) 52e57. [4] Jun Cheol Oh, Dong-Kyun Choo, Sunghak Lee, Microstructural modification and hardness improvement of titanium-base surface-alloyed materials fabricated by high-energy electron beam irradiation, Surf. Coat. Technol. 127 (2000) 76e85. [5] Mitun Das, Sandip Bysakh, Debabrata Basu, T.S. Sampath Kumar, Vamsi Krishna Balla, Susmita Bose, Amit Bandyopadhyay, Microstructure, mechanical and wear properties of laser processed SiC particle reinforced coatings on titanium, Surf. Coat. Technol. 205 (19) (2011) 4366e4373. [6] M. Heydarzadeh Sohi, S.F. Kashani Bozorg, H. Jafari Varzaneh, Formation of metal matrix composite layers with SiCp on titanium using TIG process, in: ICAMT2004, Proc. Int., Manufacturing Conf., K.L., Malaysia, 11e13 May 2004, pp. 1022e1026. [7] Mina Moradi, Maryam Moazeni, Hamid Reza Salimijazi, Microstructural characterization and failure mechanism of vacuum plasma sprayed Ti-6Al-4V/ B4C composite, Vacuum 107 (2014) 34e40. [8] S.S. Akkaya, V.V. Vasyliev, E.N. Reshetnyak, K. Kazmanlı, N. Solak, V.E. Strel'nitskij, M. Ürgen, Structure and properties of TiN coatings produced with PIII&D technique using high efficiency rectilinear filter cathodic arc plasma, Surf. Coat. Technol. 236 (2013) 332e340. [9] E. Yun, K. Lee, S. Lee, Correlation of microstructure with high-temperature hardness of (TiC,TiN)/Tie6Ale4V surface composites fabricated by highenergy electron-beam irradiation, Surf. Coat. Technol. 191 (2005) 83e89.
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Morii, Microstructure and mechanical properties of Ti-B-N cast alloys prepared by reactive arc-melting, Mater. Trans. 43 (2002) 2193e2196. [33] X. Ma, C. Li, W. Zhang, Study on the phase diagram of the TieBeN system and the interfacial reaction of the Ti/BN joints, Mater. Sci. Engine A 392 (2005) 394e402. [34] V.A. Neronov, M.A. Korchagin, V.V. Aleksandrov, S.N. Gusenko, Investigation of the interaction between boron and titanium, J. Less Common Met. 82 (1981) 125e129. [35] S. Mridha, T.N. Baker, Crack-free hard surfaces produced by laser nitriding of commercial purity titanium, Mater. Sci. Eng. A188 (1994) 229e239. [36] T.R. Anthony, H.E. Cline, Surface rippling induced by surface-tension gradients during laser surface melting and alloying, J. Appl. Phys. 48 (1977) 3888e3894. [37] R.G. Fenish, A new intermediate compound in the titanium-boron system, Ti3B4, Trans. AIME 236 (1966) 804. [38] E. Broszeit, B. Matthes, W. Herr, K.H. 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Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process R. Yazdi*, S.F. Kashani-Bozorg School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, P.O. Box 11155-4563, Tehran, Iran
h i g h l i g h t s In-situ Ti/(TiN þ TiB) hybrid composite layers were synthesized by TIG processing on commercially pure titanium. The microstructure features were characterized by several methods. Microhardness enhanced three to five times higher than that of the CP-Ti substrate after surface modification. The fabricated composite layers improved wear resistance of CP-titanium. Severe adhesive wear mechanism of CP-titanium surface changed to mild abrasive one as a result of the surface modification.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 September 2014 Received in revised form 23 November 2014 Accepted 20 December 2014 Available online xxx
Tungsten inert gas (TIG) technique was conducted on commercially pure (CP)-Ti substrate, which was coated with h-BN-based powder mixture prior to the treatment. The treated surfaces were evaluated and characterized by means of scanning electron microscope (SEM), X-ray diffraction analysis, and electron dispersive spectrometry (EDS). The microhardness and wear experiment were also performed by using a microhardness machine and pin-on-disk tribometer. As h-BN reacted with titanium, an in-situ hybrid composite layer was formed showing near stoichiometric dendrites of TiN, platelets of TiB and interdendritic regions of a0 -Ti martensite crystal structures. The population level of TiN and TiB regions were found to increase using a pre-placed powder mixture with greater h-BN content. However, the fabricated layers exhibited cracking and porosity; these were minimized by adjusting arc energy density and h-BN content of powder mixture. The microhardness value of the fabricated hybrid composite layers was found to be in the range of ~650 HV0.2e1000 HV0.2; this is three to five times higher than that of the untreated CP-Ti substrate. In addition, the in-situ hybrid composite layers exhibited superior wear behavior over CP-Ti substrate; this is attributed to the formation of newly formed ceramic phases in the solidified surface layers and good coherent interface between the composite layer and CP-substrate. Meanwhile, severe adhesive wear mechanism of CP-titanium surface changed to mild abrasive one as a result of surface treatment. © 2014 Elsevier B.V. All rights reserved.
Keywords: Welding Coating Composite materials Microstructure Hardness Wear
1. Introduction Titanium and its alloys have some excellent properties such as a good corrosion resistance and high strength-to-weight ratio. Its light weight and ability to resist high temperature make it suitable for aerospace applications and their good biocompatibility make them suitable as implants for dentistry and orthopedic applications [1,2]. A disadvantage of titanium-based alloys is their high friction
* Corresponding author. E-mail address: [email protected] (R. Yazdi).
and poor wear resistance [3,4]. To overcome this problem, various surface composite layers were fabricated on Ti-based alloys using hard ceramic particles such as SiC [5,6], B4C [7], TiN [8], TiC [9,10], TiB2 [11,12], and Ti3Al [13]. These hard particles were introduced into the surface melt pool generated by different localized heat sources e.g., electron beam [9,11], plasma torch [7,14], laser beam [5,7]. However, such composite layers are usually associated with weak interfacial particulate/metal matrix bonding strength when the material is under cyclic loading. Reactions between boron (B) and nitrogen (N)-bearing ceramic particles with titanium are of special interest due to high affinity of Ti for boron and nitrogen and high thermodynamic stability of TieB
http://dx.doi.org/10.1016/j.matchemphys.2014.12.026 0254-0584/© 2014 Elsevier B.V. All rights reserved.
Please cite this article in press as: R. Yazdi, S.F. Kashani-Bozorg, Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process, Materials Chemistry and Physics (2014), http://dx.doi.org/10.1016/j.matchemphys.2014.12.026
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and TieN compounds [15]. Euh et al. [16] showed that upon electron beam irradiation of Tie6Ale4V alloyed with pre-placed TiB2based powder layer, ceramic powder was melted and an in-situ surface composite layer was fabricated, which is reinforced with the newly formed phases. Further investigation by Kooi et al. [17] revealed orientation relationships between the titanium matrix and newly formed intermetallic compounds; the coherent interface can be beneficial to the precipitates/matrix bonding strength. According to Mitun Das research [18], the surface of Tie6Ale4V alloy could be modified by laser processing using mixture of BN and Ti powders. Although there are several features associated with techniques such as plasma, laser and electron beam surface modification, but processing time, flexibility in operation, economy in time, energy, and material consumption, and processing precision are the important advantages of tungsten inert gas (TIG) surface alloying process. Limited works were reported on the applications of this process for surface modification of titanium alloys [13,19,20]. In the present investigation, TIG surface alloying of commercial purity titanium (CP-Ti) substrate pre-placed by h-BN-based powder mixtures was conducted under an atmosphere of pure argon gas. Microstructural characterization of the fabricated surface composite layers was carried out. In addition, changes in surface wear characteristics were evaluated.
using TIG process on CP-titanium with h-BN-based layer. Surface melting of the pre-coated substrates was carried out using a TIG torch. A tungsten electrode with a diameter and height of 2.4 and 2 mm was used to create an arc between the tip of the electrode and the substrate surface, respectively. The voltage (V) on tungsten electrode was maintained at 17 V. TIG surface alloying was carried out in a chamber using argon shielding gas with a purity of 99.999% at flow rate of 10 l min1. Alloyed tracks were produced using different substrate traverse speeds (100e200 mm min1) and TIG
2. Experimental procedure Ti metal plates (purity 99%) of 7 mm thickness were supplied. They were cut into 100 mm 25 mm pieces. Contaminations such as oxides, dirt, grease, etc. were removed from the surface of the substrate by grit blasting. All of the grit-blasted test specimens were cleaned in an acetone ultrasonic bath. Fine boron nitride and Ti powder (with a mean particle size of ~10 and ~150 mm, respectively) were blended to make up three powder mixtures containing 5, 10 and 15 mass% h-BN. Before TIG processing, powders were mixed homogeneously using ball mill for half an hour. Organic binder (polyvinyl alcohol) was used to pre-place the powder mixture onto CP-Ti substrates; this prevents powder scattering during fabrication. However, the amount of binder was restricted to only 1 mass% in order to limit pore formation. Finally, the powder mixture was hand-brushed on the substrate surface to a thickness of ~0.8 mm. The pre-coated substrates were then dried in an oven at 150 C for two hours. Fig. 1 shows schematic processing of an in-situ composite layer
Fig. 1. Schematic presentation of the fabrication of an in-situ composite layer using TIG process on the pre-placed titanium with h-BN-based layer.
Fig. 2. Transverse sections of the fabricated composite layers using pre-placed layer with 5 mass% h-BN (a and b, employing TIG arc energy densities of 112 J mm2 and 349 J mm2, respectively) and 10 mass% h-BN (c, employing TIG power density of 349 J mm2) contents.
Please cite this article in press as: R. Yazdi, S.F. Kashani-Bozorg, Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process, Materials Chemistry and Physics (2014), http://dx.doi.org/10.1016/j.matchemphys.2014.12.026
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currents (60e80 A). TIG heat inputs were calculated using Equation (1) [19]:
Arc Energy Density J mm2 ¼
3
Current ðIÞVoltage ðVÞ Arc Radius mm Velocity mm s (1)
Specimens were cut transversely to the alloyed tracks and polished using standard metallographic procedure and then etched in a solution containing 5 ml HF, 15 ml HNO3 and 80 ml H2O for 10 s. An X-ray diffractometer (Philips X'Pert Pro) was used for the determination of the formed phases on the fabricated surface layers. In addition, microstructural examination was carried out using an optical and a Cam Scan (MV2300) scanning electron
Fig. 3. (a) SEM micrograph of the cross section of an in-situ composite layer produced using pre-placed layer with 5 mass% h-BN content and a TIG arc energy density of 349 J mm2 exhibiting (b) dendritic and (c) platelet structures.
Fig. 4. Energy dispersive spectrum of (a) a fabricated composite layer, (b) dendrite, and (c) platelet.
Please cite this article in press as: R. Yazdi, S.F. Kashani-Bozorg, Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process, Materials Chemistry and Physics (2014), http://dx.doi.org/10.1016/j.matchemphys.2014.12.026
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Fig. 5. X-ray diffraction patterns of the fabricated composite layers produced using 5 mass% h-BN (a) and 10 mass% h-BN (b) powder in the pre-placed layer at various TIG arc energy densities: (a) 112 J mm2 and (b) 349 J mm2.
microscope (SEM). The former was linked with an Oxford Inka200 energy dispersive spectrometer (EDS) for the determination of chemical composition. The EDS device has an ultra thin atmospheric beryllium window capable of chemical analysis of boron and nitrogen. Hardness measurements were carried out using a microhardness tester (Micro met) under a 1.962 N load on cross sections of the fabricated surface layers. Dry wear evaluation of the fabricated surface layers and asreceived CP-Ti substrate was carried out using pin-on-disk testing which performed under a constant load of 15 N and at a speed 0.6 m s1. Before measurements, the specimens were ultrasonically cleaned in methanol and dried with blown air. The pin-on-disc tests were performed against quenched and tempered (AISI 52100) steel discs with a hardness of ~61 HRC. Weight losses were measured after certain sliding distances, by weighing the samples to an accuracy of 104 g using an electronic microbalance. Surfaces of the worn samples were investigated using SEM and EDS. 3. Results and discussion 3.1. Effect of process parameters Typical transverse sections of the fabricated tracks are shown in Fig. 2. All sections exhibited a conduction-limited profile. Surface composite layers were achieved with a depth of up to ~2 mm. As it can be seen in Fig. 2a, spherical-type pores were present in the treated track fabricated on the pre-placed CP-Ti with h-BN amount of 5 mass% using arc energy densities of 112 J mm2. Gas entrapment within the fusion zone due to relatively higher rate of solidification process at the lowest arc energy density (i.e., 112 J mm2) forms spherical pores. No pores were present in the tracks using higher arc energy density (e.g. Fig. 2b). Increasing power density decreases substantially the rate of solidification, which provides more time for releasing the unreacted gases from the surface molten layer. Contrary to tracks produced on the pre-placed CP-Ti with h-BN amounts of 5 and 10 mass% (Fig. 2), no arc stability was reached on pre-placed layer containing 15 mass% h-BN. Phase h-BN was
reported to be an excellent dielectric [15]. Therefore, arc discontinuity can be attributed to insulating property of h-BN. It means that there was inadequate titanium powder density in the pre-placed layer providing sufficient arc stability. Fig. 2c also exhibits a processed layer with cracking produced using pre-placed layer with 10 mass% h-BN contents and employing TIG power density of 349 J mm2. In the case of pre-placed CP-Ti with h-BN amount of 10 mass%, tendency to cracking was observed using various energy densities. TieN and TieB phase diagrams [21,22] show that, nitrogen has high solubility in liquid and solid titanium (<8 mass%) due to small atomic radius (0.75 A), and boron with larger atomic radius (1.17 A) has very low solubility (<0.05 mass%) in solid titanium, even in the high-temperature beta phase [23]. The rapid cooling associated with the TIG process may cause metastable retention of boron in the titanium interstices, impeding motion of the geometrically necessary dislocations (GND) needed to accommodate the thermal contraction and the great differences between the coefficient of thermal expansion (CTE) of the surface composite layers and CP-Ti substrate is believed to the punch-out dislocations around the second phase particles and cracking happened consequently. Another factor that may contribute to cracking is the possibility that some amount of nitrogen and boron from dissociated h-BN is dissolving in the molten metal zone. Due to high cooling rate, excessive nitrogen and boron content in the quenched titanium-based solid solution could generated a brittle crystal structure; which is susceptible to cracking [24]. 3.2. Microstructures SEM micrographs and EDS analysis of a fabricated surface layer have been presented in Figs. 3 and 4, respectively. Microstructural evaluation of the cross sections showed dendritic-type in the interdendritic structure (Fig. 3a and b). The secondary arms of the longer dendrites exhibited orthogonal angle with the main stem, revealing a cubic crystal structure. Also, relatively long pillar shaped (platelets) phase were observed in the microstructure (Fig. 3c). In addition, EDS area analysis of the fabricated layers displayed substantial concentration of B and N (Fig. 4a). However, no h-BN
Please cite this article in press as: R. Yazdi, S.F. Kashani-Bozorg, Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process, Materials Chemistry and Physics (2014), http://dx.doi.org/10.1016/j.matchemphys.2014.12.026
R. Yazdi, S.F. Kashani-Bozorg / Materials Chemistry and Physics xxx (2014) 1e11
Fig. 6. SEM micrographs of cross sections of the in-situ composite layers produced using 5 mass% h-BN powder at various TIG arc energy densities: (a) 112 J mm2, and (b) 234 J mm2, showing dendritic and platelet structures in a0 -Ti structure.
particle was detected in the microstructure. Initially, melting occurred in the substrate and titanium powder when applying the liquid processing. The temperature of the molten zone is much higher than that of Ti substrate (1660 C), as a result, all of the BN powder could be uniformly surrounded by liquid titanium into the molten pool. Then, h-BN dissipated in the molten zone during the TIG processing and due to high affinity between titanium and boron or nitrogen [15], more stable phases were formed in the structure. The data are in agreement with the results of Yeh et al. [25] who detected no BN residual in the product of self-propagating hightemperature synthesis of compacted Ti and BN powders. Also, consistency was found with the results of laser surface alloying [18,26]/cladding [27,28] of Ti-based alloys pre-placed with BN layers since no BN was observed in the laser surface alloyed/clad
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layers. Although the concentration of nitrogen and boron, which are very light elements, could not be measured with the sufficient degree of accuracy by EDS analysis, as illustrated in Fig. 4b and c, EDS point analysis on the dendrites exhibited nitrogen and titanium enrichment, and high concentration of Ti and B on pillar shaped phase were shown by EDS point analysis. In order to identification of the phases with different morphologies, XRD is used to analyze. Diffraction patterns of the in-situ composite layers produced using pre-placed CP-Ti with h-BN amounts of 5 and 10 mass% were illustrated in Fig. 5. Reflections with higher intensities were found to be in agreement with TiN (having NaCl type crystal structure), TiB (close-packed hexagonal) and a0 -Ti (close-packed hexagonal) crystal structures. The XRD patterns of the treated surfaces were found to be consistent with the quantitative EDS results (Fig. 4b and c). TiN can form over a wide range of nonstoichiometric compositions with various N:Ti ratios [29]. The near stoichiometric result of the present investigation, which has a dendritic structure (Fig. 3b) is similar to solidstate reactions but unlike substoichiometric TiN produced in vaporphase reactions [30]. TiB phase with platelet morphology (Fig. 3c) was previously observed by many investigators [31]. The titanium borides, especially TiB, are reported to form in a hexagonal pillar shape [26]. Morikawa et al. [32] found clusters of TiB particles with rod-like shape along the grain boundary in the cast Tie1Be1N alloy. In the present work, relatively smaller features were observed; this is attributed to higher cooling rate of TIG compared to cast processing. But, the TiN and TiB phases in the microstructure are coarser compared with those produced with higher energy laser processing. Because lower melt pool heat input will have a comparatively shorter time for growth of precipitates in laser processing, which will result in finer particles [18]. While overlapping with indications of other phases, few indications may correspond to XRD pattern of Ti3B4 crystal structure. But, no microstructural feature related to this structure was detected using SEM and EDS. The existence of Ti3B4 was also reported by Fenish [33] and Neronov et al. [34]. It can be seen that the intensities of XRD peaks of TiB in the pattern do not match the standard intensities (Fig. 5), and they suggest a preferred orientation. The XRD patterns reveal that prominent peaks are present in the surface composite layer synthesized at 349 J mm2 arc energy density and weaker intensities are in the fabricated tracks treated with lower heat inputs (e.g., Fig. 5). The arc energy plays a significant role in the alloying process. Increasing the arc energy decreases the cooling rate and an adequate time will be available for alloying. So, it is reasonable to say that decreasing the arc energy density lowers the alloying process. This idea approved by comparing SEM micrographs of the treated tracks using (a) 112 J mm2, (b) 234 J mm2, and (c) 349 J mm2 TIG arc energy densities which are shown in Figs. 6a,b and 3a, respectively. Therefore, the population level of the randomly grown dendrites and platelet phases increase by increasing the power density from 112 to 349 J mm2. However the strength and stiffness of titanium-base alloys may be increased by the precipitation of TiB intermetallic phase [18]. The presence of high amount of brittle TiN phase in the microstructure is another contributing factor to the crack occurrence of the track fabricated with 10 mass% h-BN presented in the pre-placed layer [35]. The random growth of dendrites and platelet phases in the surface composite layers can be attributed to the stirring of the melt pool by convection and Marangoni flow due to temperature gradients and gradient of the surface tension in the melt pool [36]. The concentration of gradient of diffused nitrogen and boron by this stirring action can result in a homogeneous microstructure. The mechanism of the in-situ composite layer formation during
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Fig. 7. The schematic diagrams illustrating the mechanism of the in-situ composite layer formation with pre-placed h-BN-based powder layer using TIG process.
the TIG processing has been schematically illustrated in Fig. 7. In the first stage, the h-BN powder dissociated into nitrogen and boron atoms; which then solved into and also reacted with liquid titanium (Fig. 7b). The solidification sequence of the in-situ composite layers is deduced to be formation of TiN dendrites with melting point of 3600 C from the liquid. The remaining liquid is then depleted in nitrogen and enriched with boron (Fig. 7c). Boron element with 2.5e7.5 mass% in the melting pool is a favorite condition for the platelet-like TiB phase formation (Fig. 7d); having 2200 C melting point [18]. After solidification completion, high cooling rate leads to transformation of the solidified b-Ti interdendritic regions to a0 -Ti crystal structure upon cooling to room temperature via a solid state martensitic reaction. The a0 -Ti is believed to be within the interdendritic regions. Although the Ti-rich corner of the TieBeN ternary system includes the various phases [37] and there are several reports on the ternary TieBeN coatings but no chemical composition was often specified [38] and in some investigations [27,37], no ternary
compound was reported in the TieBeN ternary system. Considering the phases characterized in the microstructure, the following reactions may occur in the TIG alloyed layers: BN (s) þ Ti (L) ¼ TiN(s) þ B(s)
(2)
Ti (l) þ B(s) ¼ TiB(s)
(3)
Also, other reactions were proposed by researchers [39]: 3Ti (l) þ 4B(s) ¼ Ti3B4(s)
(4)
Ti (l) þ 2B(s) ¼ TiB2(s)
(5)
The molar Gibbs reaction energy for the four equations as a function of temperature was reported by Lu and Zhang [40,41]. Calculations of Ding et al. [42] showed negative values for the molar Gibbs reaction energy for all of the above equations from 527 C to
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Fig. 8. SEM micrograph exhibiting a martensitic morphology in the heat affected zone (HAZ) of an in-situ composite layer produced on CP-Ti substrate with 5 mass% h-BN content using a TIG power density of 234 J mm2.
Fig. 9. Microhardness profiles of the fabricated composite layers produced with 5 mass% h-BN content using various TIG energy densities.
Fig. 10. Weight loss data of the worn samples of CP-Ti substrate and the composite layers produced with 5 mass% h-BN content using two different TIG energy densities (234 and 349 J mm2).
927 C. However, TiB2 possesses the lowest value. Few reports were found TiB2 on melting reactions at BNeTi interfaces in the literature [26,27,42]:
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(a) Molian and Hualun [27] employed laser surface cladding of Tie6Ale4V with BN powder and reported of the presence of TiN, TiB2 compounds in addition to a0 martensite. Their results were mentioned to be based on the XRD. However, no X-ray diffraction pattern was presented [27]. Contrary to their result, TiB2 was not detected in the present investigation. The greater substrate dilution of the present work compared to laser surface cladding provides excessive Ti for the formation of TiB instead of TiB2. In addition, due to the arc stability limitation, the presence of titanium powder (~95 mass%) in the h-BN based pre-placed layers enhanced the melting zone dilution and this condition also encouraged less thermodynamically stable phase formation, i.e. TiB and Ti3B4. (b) Senthil et al. used laser surface alloying of Tie6Ale4V with pre-placed BN layer and found TiB, Ti3B4 and Ti2N phases in the alloyed layer [26]. Their results were supported by energy dispersive XRD and X-ray photo spectroscopy analysis. Again it can be stated that the higher substrate dilution in laser surface alloying than laser surface cladding can be responsible for the formation of TiB instead of TiB2. But, Ding et al. [42] did not observe Ti3B4 and TiB in the interfacial reaction between BN and Ti during active brazing. The relative low Ti content of the filler can be responsible for the formation of TiB2. This scenario is in agreement with Yeh and Teng [25] who detected TiN and TiB2 products from the reactant of 2Ti þ BN in self-propagating high-temperature synthesis chamber, but in addition to the phases of TiN and TiB2 intermediate compounds including Ti2N and TiB were detected in the final products from the reactants of 3Ti þ BN and 4Ti þ BN. The presence of Ti2N and TiB is mainly attributed to the incompleteness of nitridation resulting from the excessive melting of Ti. Another crystal structure, which was not detected in the present work, is Ti2N. Faran et al. [15,43] investigated the interaction between BN and Ti diffusion couple at 1000 C. The phase sequence in the (Ti powder)/(BN plates) diffusion couple was shown to be Tie(a0 -Ti)eTi2NeTiNeTiBeTiB2eBN. Fine dispersive Ti2N precipitates were detected using high resolution SEM. They could only be formed during annealing in the reaction zones of BN/Ti interfaces below ~1050 C. SEM image of heat affected zone of the treated track has been illustrated in Fig. 8. The treated track microstructure indicates good bonding with the substrate and smoothly metallurgical changes from the interface up to the surface. In addition, no interfacial cracks, and other defects such as lack of fusion, and pores were observed throughout the layers' cross-section. This is due to the complete melting and reaction of pre-placed layer and the substrate at the interface. At the bottom of the composite layers, a relatively fine transformed product can be observed within the substrate heat affected zone grains. The transformation products exhibited almost typical of a0 -Ti martensitic baskets [44]. Such features were reported by various investigators in the heat affected zone of CP-Ti processed by techniques like TIG [20] and etc. [27,45]. The temperature of the heat affected zone was reached to near melting point and due to heat sink effect of the underlining substrate, a0 -Ti martensite formed upon rapid cooling by the reaction of b-Ti to a0 -Ti. Meanwhile, the volume of the heat affected zone was found to increase with increasing TIG heat input. 3.3. Hardness The measured microhardness values of the in-situ composite layers from the near top surface down to the unaffected substrate are shown in Fig. 9. Results show that the microhardness value
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Fig. 11. SEM micrographs of the worn samples: (a) the untreated CP-Ti substrate and (b) the composite layer produced with 5 mass% h-BN content using a TIG power density of 349 J mm2 after a sliding distance of 1000 m.
increases with increasing TIG heat input (i.e., the population level of these hard intermetallic phases). Maximum microhardness values of the in-situ composite layers fabricated at the energy densities of 112, 234 and 349 J mm2 were 650, 874 and 993 HV0.2, respectively. These values are around three to five times higher than that of the CP-Ti substrate (~200 HV0.2). TiB and TiN Phases have the hardness of about 1800 HV and 2100 HV, respectively [46]. So, the formation and the amounts of new phases (titanium boride and nitrides) are believed to be responsible to the increase in the microhardness value. The TiB platelet precipitate can be considered as a fiber
reinforcing element [17]. Also, the hardness increase may benefits from the presence of hard TiN phase and effect of probable lattice deformation or coherent hardening in the surface composite layers. The small variations in hardness values detected in the composite layer derived from hardness measured locations where different amounts of the fabricated phases were present. The maximum obtained microhardness value is higher than that (650e850 HV) achieved on laser boronising of Tie6Ale4V with pre-placed BN using different process parameters [26]; this may be due to greater population level of hard intermetallic phases
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R. Yazdi, S.F. Kashani-Bozorg / Materials Chemistry and Physics xxx (2014) 1e11
Fig. 12. EDS analysis of wear debris on worn surface of CP-Ti substrate reflecting a relatively significant O, Fe and Cr content from the disk.
achieved in the in-situ composite layer fabricated by TIG than that of laser surface alloying. Although higher localized heat can be provided by laser but greater heat input is believed to be associated with TIG surface processing of the present investigation. However, the attained microhardness values are lower than those (1600 HV) achieved on laser cladding of Tie6Ale4V with BN [27]; this is believed to be due to lower substrate dilution in cladding than that of alloying process, which provides a greater population level of hard phases in the latter. 3.4. Wear assessment Fig. 10 shows the wear-loss data of CP-Ti substrate and the insitu composite layers fabricated by TIG as a function of sliding distance. The results reveal that weight loss increases as sliding distance increases, irrespective of microstructure features (A nearly linear relationship was found between mass loss and sliding distance). CP-Ti sample exhibited a relatively significant mass loss during sliding process; this is an indication of a relatively high rate of wear. It is reported that titanium alloys have severe wear behavior over a wide range of sliding speed [47]. The results in this study also show that the untreated titanium surface suffers high wear. Contrary to CP-titanium, the in-situ composite layer showed nearly four times lower material removal. Thus, the surface alloying process improved the wear resistance of CP-Ti substrate. According to Archard [48], the wear resistance of a material is proportional to its hardness value. The decrease of wear-loss data was found to be consistent with the increase in microhardness value of the specimens (Fig. 9). Also, the increase in the population level of hard phases in the structure (Figs. 3a and 6) is associated with more improved surface wear properties of the in-situ composite layers. This behavior could be attributed to the presence of hard intermetallic phases in the microstructure of the composite layer. SEM images of wear surfaces and their EDS area analysis are depicted in Figs. 11 and 12, respectively. It could be observed that
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Fig. 13. EDS analysis of wear debris on worn surface of the composite layer reflecting a relatively lower content of Fe and Cr.
surface of untreated CP-titanium sample was severely scored and plastically deformed, which caused a significant roughed worn surface. These features of the surface is indicative of adhesive wear mechanism, which is approved by other researchers [10,47]. The worn surface of CP-Ti also revealed relatively deep grooves and scratches aligned in the wear direction (Fig. 11a) and also with respect to EDS analysis, the wear debris of CP-Ti worn sample reveals significant contamination content of oxygen (Fig. 12); this could be due to the formation of titanium oxide during the sliding due to high temperature at the interface and also, significant Fe and Cr contamination from disk to the pin surface was detected. These elements generated transferred layer during the sliding. The transferred layer associated with the adhesion became workhardened and oxidized after multiple contacts in the wear couple, which in turn resulted in micro-cutting wear damage on the Ti surface. In comparison with wear track of CP-titanium, a distinct change in surface wear mechanism occurred, as exhibited in Fig. 11b and no significant adhesive wear features were detected on the worn surface of the fabricated composite layers by TIG. Likewise, the sample fabricated at 349 J mm2 had a better wear resistance than the specimen processed at lower arc energy (234 J mm2); this is due to the higher hardness of coating and to the coherent interfaces of the hard intermetallic phases and Ti matrix as reported by Kooi et al. [17]. Initially, comparatively softer a0 -Ti worn significantly and, then TiN and TiB phases protruded during sliding contact and prevented substantially high material removal and improved wear resistance of the composite surface. The other cause of the superior wear resistance of hybrid composite layer is less adhesion between the sample surface and steel counterpart in sliding contacts compared with CP-titanium surface. Generally, ceramics have less tendency to adhere firmly to metal surface. So, protruded ceramic phases in the structure prevented from adhesive metal-to-metal contact of the wear couple and subsequently less weight loss arose from sliding surfaces [49]. In the further sliding, the protuberant particles may
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be fractured due to the effects of less support by the matrix and stress concentration exerted on each individual particle and cause that mild scratching, shallower grooves and little material removal have been observed on the wear treated sample after a sliding distance of 1000 m (Fig. 11b). In addition, EDS analysis of the wear debris of the worn sample of the composite layer has been shown in Fig. 13, which reveals minor contamination content of Fe and Cr from disk to the pin surface. From these wear characteristics of hybrid composite layer, it could be concluded that mild abrasion is the prevalent wear mechanism. 4. Conclusion The in-situ composite layers were fabricated using TIG process on CP-Ti pre-placed with h-BN at different arc energy densities, and their microstructures, hardness and wear resistance were investigated with following conclusions. 1. Microstructural characterization of the fabricated composite layers showed that the melt zones consisted of TiN dendrites, TiB platelets and a0 -Ti martensite matrix. The population level of these hard phases increased with increasing the TIG arc energy densities. 2. The presence of the hard phases improved the microhardness value of the surface composite layers up to five times greater than that of the CP-Ti substrate. 3. The wear resistance of the fabricated composite layers was found to be about four times better that of the untreated CP-Ti substrate. This due to higher hardness and lower tendency of Ti/(TiN, TiB) composite for adhesion to steel counterpart as well as strong coherency between composite layer and CP-titanium substrate. 4. Severe wear mechanism of CP-titanium surface changed to mild abrasive one as a result of the surface modification. Acknowledgment The authors wish to thank University of Tehran for partial financial support. References [1] M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, Ti based biomaterials, the ultimate choice for orthopaedic implants e a review, Prog. Mater. Sci. 54( (2009) 397e425. [2] Sebastian Bauer, Patrik Schmuki, Klaus von der Mark, Jung Park, Engineering biocompatible implant surfaces: part I: materials and surfaces, Prog. Mater. Sci. 58 (2013) 261e326. [3] Xiu-Yan Li, Pei-Qiang Wu, Bin Tang, J.P. Celis, Fatigue behavior of plasma surface modified Tie6Ale4 V alloy, Vacuum 79 (2005) 52e57. [4] Jun Cheol Oh, Dong-Kyun Choo, Sunghak Lee, Microstructural modification and hardness improvement of titanium-base surface-alloyed materials fabricated by high-energy electron beam irradiation, Surf. Coat. Technol. 127 (2000) 76e85. [5] Mitun Das, Sandip Bysakh, Debabrata Basu, T.S. Sampath Kumar, Vamsi Krishna Balla, Susmita Bose, Amit Bandyopadhyay, Microstructure, mechanical and wear properties of laser processed SiC particle reinforced coatings on titanium, Surf. Coat. Technol. 205 (19) (2011) 4366e4373. [6] M. Heydarzadeh Sohi, S.F. Kashani Bozorg, H. Jafari Varzaneh, Formation of metal matrix composite layers with SiCp on titanium using TIG process, in: ICAMT2004, Proc. Int., Manufacturing Conf., K.L., Malaysia, 11e13 May 2004, pp. 1022e1026. [7] Mina Moradi, Maryam Moazeni, Hamid Reza Salimijazi, Microstructural characterization and failure mechanism of vacuum plasma sprayed Ti-6Al-4V/ B4C composite, Vacuum 107 (2014) 34e40. [8] S.S. Akkaya, V.V. Vasyliev, E.N. Reshetnyak, K. Kazmanlı, N. Solak, V.E. Strel'nitskij, M. Ürgen, Structure and properties of TiN coatings produced with PIII&D technique using high efficiency rectilinear filter cathodic arc plasma, Surf. Coat. Technol. 236 (2013) 332e340. [9] E. Yun, K. Lee, S. Lee, Correlation of microstructure with high-temperature hardness of (TiC,TiN)/Tie6Ale4V surface composites fabricated by highenergy electron-beam irradiation, Surf. Coat. Technol. 191 (2005) 83e89.
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Please cite this article in press as: R. Yazdi, S.F. Kashani-Bozorg, Microstructure and wear of in-situ Ti/(TiN þ TiB) hybrid composite layers produced using liquid phase process, Materials Chemistry and Physics (2014), http://dx.doi.org/10.1016/j.matchemphys.2014.12.026