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The influences of melting degree of TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites
Materials and Design 136 (2017) 185–195
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The influences of melting degree of TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites Shunyu Liu, Yung C. Shin ⁎ School of Mechanical Engineering, Purdue University, Center for Laser-based Manufacturing, West Lafayette, IN 47906, USA
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The melting degree of reinforced TiC was controlled by the laser energy density and the TiC content. • The high melting degree of reinforced TiC caused the formation of dendritic carbides, inducing premature failure. • The detrimental dendritic carbides were eliminated by adjusting the laser deposition conditions and reinforced TiC size. • The improved mechanical properties and strengthening mechanisms for different melting degrees of TiC were investigated.
a r t i c l e
i n f o
Article history: Received 29 August 2017 Received in revised form 19 September 2017 Accepted 30 September 2017 Available online 02 October 2017 Keywords: Laser direct deposition Ti6Al4V-TiC composite Functionally graded material Microstructure Mechanical property
a b s t r a c t This study is concerned with the influences of melting degree of embedded TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites and a functionally graded material. The melting degree of embedded TiC was controlled by the input laser energy density and the added TiC content. The formation of detrimental primary dendritic TiC grains was successfully avoided by properly adjusting the deposition conditions and the particle size range of TiC reinforcements. The resultant compression test revealed the ultimate strength increasing from 1381 ± 19 MPa to 1636 ± 23 MPa as the premixed TiC content increased from 0 to 15 vol% while a true strain of 0.141 ± 0.002 was still retained for 15 vol% TiC. The primary strengthening mechanism for composites with the most melting control of TiC is the solid solution strengthening induced by carbon, while that for the least melting control is dominated by the unmelted TiC particulates and the refined microstructure resulting from the resolidified carbides. The defect-free functionally graded Ti6Al4V-TiC with 0 to 40 vol% TiC achieved an increased hardness from HRC ~39 to HRC ~65. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Ti6Al4V alloy (Ti64) occupies about 50% of usages among all the titanium alloys due to its high strength, low weight ratio, outstanding ⁎ Corresponding author. E-mail address: [email protected] (Y.C. Shin).
https://doi.org/10.1016/j.matdes.2017.09.063 0264-1275/© 2017 Elsevier Ltd. All rights reserved.
corrosion resistance and superior biocompatibility [1]. Ti64 can also be reinforced with hard particles to build metal matrix composites (MMCs) so as to combine the high stiffness and hardness of the reinforcements with the toughness and damage tolerance of the Ti64 matrix, which leads to increased wear resistance and mechanical properties [2–5]. Due to the high reactivity of titanium and its alloys, the reinforcements have to be carefully chosen. The commonly embedded
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Table 1 Summarized laser deposition conditions for Ti6Al4V and Ti6Al4V-TiC composites with different TiC contents. Sample Premixed TiC size Melting Laser no. TiC (μm) degree power (vol%) (W) 1 2 3 4 5 6 7 8
0 1 1 5 5 10 15 PF1:0 PF2:40
– 90–150 b45 90–150 b45 90–150 90–150 – 45–150
– Least Most Least Most Least Least
300 290 375 265 350 230 200 300 → 200
Powder feed rate (g/min)
Laser scan speed (mm/s)
Layer thickness (μm)
1.5 1.8 1.2 1.8 1.3 1.8 2.0 1.5 → 0 0 → 2.8
15 15 15 15 15 15 15 15
254 254 254 254 254 254 228.6 228.6
Note: No. 8 condition was used for depositing Ti64-TiC functionally graded material. PF1 and PF2 mean powder feeder hopper 1 and 2.
reinforcements for titanium alloys are ceramics such as WC [6], TiN [2], TiB [7], SiC [8], TiC [5,7,9–11] or in situ reacted TiC through adding graphite carbon or carbon black into titanium alloys [12,13]. Among these reinforcements, TiC is highly compatible with titanium as it possesses a similar density to titanium, a high hardness, and chemical and thermal stability. The high biocompatibility also facilitates the potential application of Ti64-TiC MMCs as load-bearing implants. Laser direct deposition (LDD) is an effective and efficient method for preparing MMCs since it can simultaneously feed different powders into the molten pool to provide requisite compositions and properties as well as customized exterior shapes. The process has been successfully used to generate fully dense components to near-net shape tolerances with improved hardness [14,15], strength [5,9], wear resistance and corrosion resistance [15–17] over some other fabrication methods [18–21]. Functionally graded materials (FGMMCs) can also be produced easily by properly adjusting the powder feed rates [16,22,23]. Although TiC is thought to be thermodynamically stable in titanium and titanium alloys, partial melting always happens to the embedded TiC particulates during the LDD process [9,11,14,24]. Some studies showed the diffusion of carbon from the melted TiC into the Ti64 matrix, causing solid solution strengthening [22,25]. The dissolved carbon atoms can in situ react with titanium to form different types of resolidified TiC precipitates [9,12]. With an increase in embedded TiC content, the strength of the composites improves, but accompanies a substantial sacrifice of ductility [9,12,22]. In Wang et al.'s work [9], the Ti64-TiC MMCs almost completely lost their ductility with 8 vol% TiC added. Several factors are responsible for the low ductility of Ti64-TiC MMCs such as the existence of defects like initial voids and lack-of-fusion [26–28], and the high oxygen level and impurity contents involved
during the fabrication process [29,30]. To the authors' knowledge, the microstructure of resolidified TiC and the unmelted TiC (UMC) particles also have a big impact on the mechanical properties of MMCs, but no relative research has been conducted. In the current study, the melting degree of the TiC reinforcements and the morphologies of the resolidifed carbides are controlled by adjusting the LDD parameters and the TiC particle size range. By avoiding the formation of dendritic TiC precipitates, the mechanical properties of the LDD Ti64-TiC MMCs are dramatically improved. A Ti64-TiC FGMMC with an increased TiC content ranging from 0 to 40 vol% within 5 mm was also built. The FGMMC realized a defect-free microstructure and an elevated hardness. The results are presented and discussed in detail. 2. Experimental procedures Gas atomized spherical Ti64 powders (grade 5, b150 μm) obtained from Puris®, along with irregularly shaped TiC particles in the same size range obtained from Atlantic Equipment Engineers®, were used to premix 1 vol%, 5 vol%, 10 vol% and 15 vol% Ti64-TiC precursor powders. LDD Ti64-TiC MMCs were carried out with an Optomec LENS® 750 system, which was equipped with a 500 W continuous wave fiber laser with a wavelength of 1064 nm and a dual-hopper powder feeder system. During the LDD process, the input laser energy density can be roughly calculated as J= αP/VH, where J is the input laser energy density, P is the laser power, V is the laser scan speed, H is the layer thickness, and α is the laser absorptivity coefficient. The input energy density required to deposit a material is mainly determined by the laser absorptivity coefficient and the melting temperature. The laser scan speed was kept constant as 15 mm/s in the current study. Although the melting of TiC reinforcements cannot be avoided, the complete melting of TiC is also difficult to obtain, particularly when the TiC size is large, because the laser processing time is extremely short and the heating and cooling cycles are super-fast [31]. Therefore, in the present work, most melting and least melting conditions were used to indicate two extreme melting degrees. The LENS® 750 system is designed to feed powders with a size range of 45–150 μm. Therefore, 90–150 μm TiC particulates were used for the least melting control of embedded TiC, while b 45 μm TiC particles were used for the most melting control of TiC. The small particles have a larger surface area, and are easier to melt than larger particles. The irregularly shaped TiC powder did not flow well through the powder feed system, particularly when b 45 μm particles were used. Therefore, premixed Ti64-TiC powder mixtures were used to deposit Ti64-TiC MMCs. The mixtures were blended by stirring them for 20 min followed by tumbling them on a roller mill at 90 rpm for 12 h. For building the Ti64-TiC FGMMC, pure Ti64 powder and premixed Ti64-40vol%TiC were fed simultaneously using two powder hoppers.
Fig. 1. (a) Laser direct deposited Ti6Al4V or Ti6Al4V-TiC composite cylinder, (b) the dimension of the machined compression specimen.
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Fig. 2. Morphologies of TiC in Ti6Al4V-15vol%TiC composites deposited with 250 W + 254 μm layer height.
The FGMMC composition varied from 100% Ti64 for the first two layers to an increasing volume of TiC for the subsequent 20 layers. The powder feed rates of the two hoppers were adjusted to obtain a 2 vol% TiC increment for each layer to reach a final composition of 40 vol% TiC. A delay time of 10 s was set between adjacent layers to allow for adequate powder feed stabilization. The substrate used was a 3.5-mm thick hot rolled Ti64 plate whose surface was cleaned and degreased with ethanol prior to usage. Several 10 mm diameter cylinders were built in the argon-controlled working chamber, where the oxygen level was controlled below 10 ppm. After depositing one layer, the orientation of the laser scan was rotated by 60 degrees which can lead to a good cylindrical geometry, and the laser focal point was elevated in Z direction by a pre-set height, which was equal to the measured actual layer thickness. The laser deposition parameters for melting control of TiC particles to obtain a dense and desirable microstructure are summarized in Table 1. A room temperature compression test was carried out on the Ti64 parts and Ti64-TiC MMCs with 1, 5, 10 and 15 vol% TiC at a strain rate
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of 0.005/min on an MTS Sintech 30/D hydraulic test platform with a 100 kN load cell. A CNC lathe was used to turn the surface of the cylinders. The images of the as-built and machined specimens are illustrated in Fig. 1. The true stress and strain data were measured from these tests. To characterize the microstructure, the specimens were sectioned parallel to the building direction, mounted in bakelite, and polished with #240–#1200 SiC papers followed with vibratory polishing for 24 h. To observe the morphologies of the resolidified carbides, the samples were not etched, while the Kroll's etchant was used to reveal the microstructure of the Ti64 matrix. The microstructure details were observed by an optical microscope (Zeiss), and the fracture morphologies were observed using a scanning electron microscope (SEM, JEOL JSMT330) under the secondary electron mode. A LECO RT-370 Rockwell hardness tester was utilized to examine the macrohardness of the Ti64-TiC FGMMC specimens, which indicated the average hardness. A load of 100 kg and 15 s were used. The positions with different TiC contents were marked on the specimen by calculating and measuring the distance to the side with 40 vol% TiC, and the hardness test was preformed along the marked lines. For each TiC content, five indentations were measured on one specimen, and four specimens were measured. The values were finally averaged. 3. Results and discussion 3.1. Laser direct deposited Ti64-TiC composites 3.1.1. Microstructures before etching Fig. 2 shows five morphologies of TiC precipitates in the LDD Ti64TiC MMCs. The big particle is UMC. The melting of the embedded TiC particles can be visually observed according to the boundary shape of the TiC particles. If the melting degree is high, many tiny columnar carbide grains attach to the unmelted TiC (UMC) particles, and the boundary of the UMC particles displays a serrated shape as shown in Fig. 2. The big spherical carbide particles, which are uniformly distributed in the microstructure, are the primary spherical TiC (PSC). According to the Ti-C binary phase diagram [32], the PSC resulting from the in situ reaction of C and Ti first precipitates from the liquid at liquidus temperature. As the PSC stays in the liquid for a relatively long time, the particles grow large. If the melting degree of the embedded TiC particles is high, the primary dendritic TiC (PDC) will form (discussed later). When
Fig. 3. Tensile test fracture morphology of Ti6Al4V-10vol%TiC built with 350 W + 254 μm reveals dendritic TiC grains indicated by arrows.
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Fig. 4. The formation and morphology evolution of dendritic TiC in Ti6Al4V-10vol%TiC composites built with a 15 mm/s laser scan speed, a 254 μm layer height and the same TiC particle size but different laser powers: (a) 230 W, (b) 250 W, and (c) 300 W. Dendritic TiC grains are indicated by arrows.
the temperature decreases to the eutectic point, some eutectic TiC grains form. One is the eutectic grain TiC (EGC) that precipitates around the boundary of the prior β grains and exhibits a loop shape. Another is the eutectic spherical TiC (ESC), which displays a segregated small spherical morphology. In Ref. [12], the ESC grains are found mainly gathering in the triple junction. The LDD process will lead to a non-equilibrium solidification where the liquidus and eutectic temperatures will be a little different from those in the equilibrium phase diagram. The formation of PDC is harmful to the mechanical properties of Ti64-TiC MMCs since the dendritic grains will block the development of the strain, and it is easy for the dendrite tip to initiate cracks. Fig. 3 shows the tensile test fracture morphology of Ti64-10vol%TiC deposited with a high energy input. Some bright dendritic TiC grains, indicated by the arrows, are clearly observed. These dendritic grains display a brittle cleavage fracture behavior, and will change the fracture from ductile
behavior to brittle behavior, rendering the composite loss of the ductility [9]. In the current work, one goal is to avoid the formation of PDC grains by designing proper deposition conditions. Compression tests were carried out since less work was performed on the compression test than the tensile test for Ti64-based composites. However, the tensile fracture can show a clear view of the dendrites. Fig. 4 illustrates the formation mechanism and morphology evolution of PDC in Ti64-10vol%TiC composites. When the laser power was 230 W, only PSC, EGC and ESC were observed, while several short PDC grains were produced for 250 W. As the laser power further increased to 300 W, more and larger PDC grains were produced. Actually, the increased laser power led to an increased laser energy as other parameters were kept constant, and the melting degree of the embedded TiC hence increased. This result reveals the direct relationship between the formation of PDC and the melting degree of embedded TiC particles.
Fig. 5. Potential defects in the least melting control of laser direct deposited Ti6Al4V-TiC composites: (a) lack-of-fusion, and (b) inhomogeneous distribution of unmelted TiC particulates.
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Fig. 6. Microstructures of unmelted TiC particles and resolidified carbides in laser direct deposited Ti6Al4V-TiC composites with the least melting control of TiC reinforcements: (a, b) 1 vol% TiC; (c, d) 5 vol% TiC; (e, f) 10 vol% TiC; (g, h) 15 vol% TiC.
It is believed that as the melting degree of the embedded TiC increased, more carbon dissolved in the liquid. When the carbon content reached to a certain amount, the PDC began to form and attached to PSC as nuclei. The primary carbides consumed the carbon, and hence the eutectic carbide content decreased (Fig. 4(c)).
In addition to the laser energy density, the embedded TiC content also affects the melting degree. According to the microstructural analysis, 5 vol% TiC or less will not produce PDC even when the embedded TiCs were almost completely melted, as will be shown later. To avoid the formation of PDC, the LDD parameters must be carefully selected
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Fig. 7. Microstructures of Ti6Al4V-TiC with the most melting control of TiC: (a) 1 vol% TiC; (b) 5 vol% TiC.
for N5 vol% TiC. Therefore, the most melting control condition was only performed on embedded 1 vol% and 5 vol%TiC. For brevity, hereafter we define the most melting conditions for 1 and 5 vol% TiC as 1M and 5M, and the least melting conditions for 1, 5, 10 and 15 vol% TiC as 1L, 5L, 10L and 15L. For the least melting control, the lowest energy density condition was tested by increasing the powder feed rate so as to increase the layer height for a set laser power and laser scan speed. However, two kinds of defects are easily formed in this process and should be avoided (Fig. 5). If the laser energy density is not high enough for the preset layer thickness, a lack-of-fusion defect will be produced, resulting in a premature failure of the composites [8,26,27]. This lack-of-fusion defect exhibits an irregular-shaped void, and is primarily formed in the interpass overlap (Fig. 5(a)). On the contrary, a high energy density will produce a deep molten pool, leading to a high layer thickness. The vigorous Marangoni flow and buoyancy force in the molten pool will push the UMC particles to the edge zones of the molten pool [33,34]. The nonuniform distribution of the UMCs (Fig. 5(b)) will cause anisotropy in the mechanical properties. Finally, the defect-free condition with the lowest energy input was selected as the optimized deposition condition for the least melting control. For different embedded TiC contents, the optimized laser deposition conditions are different (Table 1). The microstructures of the UMC and resolidified TiC are illustrated in Fig. 6. As can be seen, the deposition conditions produce a dense and defect-free structure where no PDC grains are formed and the UMC particles are uniformly distributed in the Ti64 matrix for all compositions. Fig. 7 shows the microstructures of the resolidified TiC with the most melting control of reinforced TiC in Ti64-TiC MMCs. No PDC was detected even though the embedded TiC particles were almost completely melted. The resolidified TiC in 5M displays a needle-shaped or spherical-shaped eutectic morphology (Fig. 7(b)), but for 1M, no obvious
resolidified carbides are observed (Fig. 7(a)). This indicates that the melted TiC completely dissolves into the Ti64 matrix. 3.1.2. Microstructures after etching Fig. 8 shows the microstructure of the LDD monolithic Ti64 after etching. During LDD fabrication, the high cooling rate always results in the formation of acicular α′ martensite instead of α + β lamellae within the prior β grain in titanium alloys [19,31]. In Fig. 8, the fine needleshaped phase within the prior β grains is the α′ martensite. The prior β grains epitaxially grow and are almost parallel to the deposition direction with slight inclinations due to the scan rotation (Fig. 8(a)) [35]. The length of the β grain could be approximately 20 mm [9,36]. Growing from the boundaries of the prior β grains, a layer of α phase (grain boundary α) is formed, which separates the prior β grains (Fig. 8(b)). Fig. 9 shows the etched microstructures of the LDD Ti64-TiC MMCs. Compared with LDD Ti64 (Fig. 8), the presence of the TiC dramatically changed the Ti64 matrix microstructure. Instead of the acicular α′ martensite, a thin and short lamellar α + β structure was formed. In the optical micrographs, the bright lamella is α-Ti while the dark phase situating around α-Ti is β-Ti. The formed lamellar α + β structure is beneficial for raising the ductility [18]. For 1L, a lamellar α + β structure with 1–1.5 μm thick α-lath was obtained (Fig. 9(a)). This indicates a slower cooling rate as compared to that for Ti64. The slower cooling rate is caused by the more energy input. Since the laser power was decreased by 10 W, the more energy absorbed must come from the 1 vol% TiC. Hence, it is believed that TiC has a higher absorptivity than Ti64 for the wavelength of 1064 nm. As the laser power was increased to 375 W for 1M, the thickness of α-lath increased to 2–4 μm, which was caused by the further decreased cooling rate induced by the high energy density. However, with more TiC added, the Ti64 matrix changed back to a fine structure, where the α-lath is b 1 μm thick and the length is also
Fig. 8. Optical microstructure of laser direct deposited monolithic Ti6Al4V.
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Fig. 9. Optical micrographs of laser direct deposited Ti6Al4V-TiC composites reveal matrix structures: (a) least melted 1 vol% TiC; (b) most melted 1 vol% TiC; (c) least melted 5 vol% TiC; (d) most melted 5 vol% TiC; (e) least melted 10 vol% TiC and (f) least melted 15 vol% TiC.
decreased compared to that in Ti64-1vol%TiC composites. The refining effect is due to the precipitation of the resolidified TiC grains, which serve as the nuclei for the α and β phases and also block the growth of α-lath [21]. As 5M (Fig. 9(d)) shows, the high energy density cannot lead to a coarser structure. In addition, the long and columnar prior β grain in LDD Ti64 was reduced to small equiaxial grains, which can be visually evaluated according to the EGC grain size. The grain boundary α was also eliminated. All these microstructure changes are beneficial for the mechanical strength [29,37,38]. 3.1.3. Compression test Fig. 10 shows the room temperature compression test results of the LDD Ti64 and Ti64-TiC MMCs deformed at the strain rate of 0.005/min. Several duplicate tests were performed to make sure the results were repeatable. As the embedded TiC increased from 0 to 15 vol%, the ultimate strength (US) and yield strength (YS) were elevated from 1381 ± 19 MPa to 1636 ± 23 MPa, and from 997 ± 14 MPa to 1310 ± 22 MPa, respectively. Although the ductility decreased, a true fracture strain value of 0.141 ± 0.002 was still obtained for 15 vol% TiC.
The improved US and YS can be attributed to several factors, such as the lack of PDC grains, the refined α + β lamellar microstructure, the reduced prior β grain size [29,37,38], the thermal stress formation during the LDD process [24], and the increased dislocation density in the matrix [22]. However, the more important strengthening effects come from the embedded TiC particulates and their products, such as the free carbon atoms, the UMC and the resolidified carbides. Ogden et al. [39] and Johnson et al. [22,25] experimentally confirmed the positive role of the solid solution strengthening effect induced by carbon. The authors concluded that carbon in the solid solution provided a greater strengthening contribution than carbides, and the strengthening effect was obvious even if the carbon content was low. According to the compression test results in conjunction with the microstructure analyses as well as the LDD process feature, other strengthening mechanisms besides the solid solution strengthening induced by carbon have been explored in the current study. Actually, the primary strengthening mechanisms are different for different melting degrees and different reinforced TiC contents. For 1 vol% TiC, 1M has a coarser structure (Fig. 9(b)) than 1L (Fig. 9(a)).
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Fig. 10. Compression test results of Ti6Al4V-TiC composites built with different conditions: (a) true stress-strain curves; (b) Ultimate strength; (c) Yield strength; (d) Elongation. M and L indicate the most and least melting control condition of the reinforced TiC particles.
The thick α-lath is known to lower US, YS, and hardness [37,38]. However, 1M had a higher US and YS than 1L, but the ductility is only a little smaller. The primary strengthening for 1M must result from the solid solution strengthening induced by carbon since no resolidified TiC and neglected UMC were detected. However, for 1L, the solution strengthening is limited since the melting of TiC particulates is controlled to a very low level (Fig. 6(b)). Hence, the improved strength compared to LDD Ti64 should be attributed to the fine α + β lamellar structure, the secondary phase strengthening resulting from the UMC particles, and maybe a little bit of solid solution strengthening. As for 5 vol% TiC, 5M and 5L displayed similar US values. Since solid solution strengthening plays a limited role in 5L as compared to 5M, the UMC particles must contribute to the strengthening effect in 5L; otherwise, the mechanical properties for the two melting conditions would show a big difference. The elevated strength from 5L to 15L also confirmed the positive role of UMC in strengthening since the microstructure of the matrix and resolidified TiC are quite similar (Fig. 9(c, e and f)). The large volumes of resolidified carbides formed in the matrix should also cause precipitation strengthening. However, it is difficult to quantitatively evaluate the precipitation strengthening effect.
As α-Ti and β-Ti both have a limited solubility of carbon, and the dissolution of carbon beyond a certain amount will impair the ductility, the Ti64-TiC MMCs have to rely on UMC and resolidified carbides to improve the strength for a high TiC content. However, the UMC and the resolidified TiC increase the dislocation piles up for plastic deformation [40], and TiC particles are hard and brittle. The continuous adding of TiC will finally lead to a deteriorated property. 3.1.4. Compression fracture morphologies and mechanism After the compression tests, SEM was used to characterize the fracture morphologies of LDD Ti64 and Ti64-TiC MMCs to reveal the fracture mechanism. The fracture morphology of 15L is similar to that of 10L except for more UMC particles, so the former is not presented in Fig. 11. With a low TiC content, the fracture displays a ductile morphology (Fig. 11(a–d)), where the deep dimple structure indicates a ductile fracture. The flat zones reveal the slip bands. The highest amount of dimple structure is observed in 1L, but not in LDD Ti64, which may be due to the acicular α′ martensitic structure in LDD Ti64. In 1M, some white spherical carbide spots are detected. This indicates the formation of PSC grains.
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Fig. 11. Compression fracture morphologies of laser direct deposited Ti6Al4V and Ti6Al4V-TiC compositions: (a) Ti6Al4V; (b) most melted 1 vol% TiC; (c) most melted 5 vol% TiC; (d) least melted 1 vol% TiC; (e) least melted 5 vol% TiC; (f) least melted 10 vol% TiC.
As the TiC content increases from 1 vol% to 10 vol%, the deep dimple morphology becomes shallow and reveals a reduction in ductility. The disappearance of the deep dimple is due to the formation of resolidified TiC grains. The UMC particles reveal a smooth and brittle feature. Therefore, the fracture mechanism is a mixture of a subdued ductile mechanism in the matrix and a cleavage mechanism in the UMC particles. The high volume percentage of UMC in 10L and15L speed up the failure of the components. Johnson et al. [14,22] investigated the cracking and failure mechanisms in the Ti64 and Ti64-TiC MMCs during compression tests. In summary, for the Ti64 samples tested at quasi-static small strain rates, the voids forming in the most severely deformed region caused the failure. As for the composites, the particle-matrix interface provided sites for additional void formation through interface separation, and thus the authors concluded that the increased density of these voids accounted for the decreased ductility of the composites with a high TiC content. In the
current work, voids were not observed. However, the debonding of TiC particles from the matrix and the formation of microcracks around the TiC-matrix interface were detected (Fig. 11(f)). The interfacial region between the reinforcements and the matrix governs the mechanical properties of the composites. The debonding or weak bonding and microcrack formed at the interface will lead to the premature failure of the Ti64-TiC MMCs. Therefore, it is necessary to adjust the laser deposition condition to avoid the formation of PDC phase, while producing a firm bonding between the embedded TiC particulates and the Ti64 matrix. 3.2. Laser direct deposited Ti64-TiC functionally graded material Instead of serving as structural materials, the Ti64-TiC MMCs can be built as FGMMC by LDD, and used as surface coatings. The content of the TiC reinforcements and the microstructure of FGMMC change gradually
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Fig. 12. Laser direct deposited Ti64-TiC functionally graded material with (0–40) vol%TiC.
from one side to the other, resulting in gradual changes in the properties. Fig. 12(a) shows the optical micrographs of LDD Ti64-TiC FGMMC. A defect-free structure has been built. A large volume percentage of embedded TiC has been retained as UMC particles. The UMC and resolidified carbides are critical for improving the wear resistance, and the effective fraction of the reinforced TiC was experimentally proved not to be b 15 vol% [9]. As the premixed TiC content exceeds 15 vol%, it is difficult to eliminate the formation of the PDC grains (Fig. 12(b–d)). However, if the Ti64-TiC MMCs are used to improve the erosion and wear resistance of a material, the MMCs can be deposited as a FGMMC coating on the base material. Within a thin coating, a high volume fraction of TiC reinforcements can be embedded. Besides, the formed carbides will also be able to improve the hardness of the matrix. As can be seen from Fig. 13 with the volume percentage of TiC particles varying
from 0 to 40 vol%, the hardness is increased from HRC 39 ± 0.12 to HRC 65 ± 0.28. Considering the high volume of UMC and resolidifed carbides, the improvement of wear resistance of the base material can be expected [9]. 4. Conclusion The influences of the melting degree of embedded TiC on the microstructure evolution and mechanical properties of the Ti64-TiC MMCs and FGMMC were studied. 1. The melting degree of the embedded TiC was determined by the added TiC content and the input laser energy density. The laser energy density cannot cause Ti6Al4V-TiC composites with no N5 vol% TiC to form dendritic carbides. It was difficult to avoid the formation of dendritic carbides when the premixed TiC is larger than 15 vol%. By proper controlling the input laser energy density, the Ti6Al4V-TiC composites with 5–15 vol% TiC were free from dendritic carbides. 2. The dendrites-free compression specimens revealed increased ultimate and yield strengths but correspondingly decreased ductility as the premixed TiC content increased from 0 to 15 vol% while a true strain of 0.141 ± 0.002 was still retained for 15 vol% TiC. 3. The solid solution strengthening induced by carbon and the unmelted TiC particles played the dominant role in strengthening the composites with the most and least melting control of TiC, respectively. The resolidified carbides also caused the precipitation strengthening effect and refined microstructure. The unmelted TiC and resolidified carbides changed the fracture mechanism from the ductile behavior to cleavage behavior. 4. With the TiC increased from 0 to 40 vol%, the defect-free Ti64-TiC FGMMC achieved an increased hardness from HRC ~39 to HRC ~65, and increased unmelted TiC and resolidified TiC contents. Acknowledgement
Fig. 13. Rockwell hardness for laser direct deposited Ti64-TiC functionally graded material with TiC increasing from 0 to 40 vol%TiC.
The authors wish to acknowledge the help of Xianfeng Lu in the compression sample machining and testing.
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References [1] M.J. Donachie, Titanium: a Technical Guide, second ed. ASM International, Metals Park, OH, 1988. [2] Y. Lin, Y. Lin, Y. Chen, Evolution of the microstructure and tribological performance of Ti-6Al-4V cladding with TiN powder, Mater. Des. 36 (2012) 584–589, https://doi. org/10.1016/j.matdes.2011.12.007. [3] F. Weng, C. Chen, H. Yu, Research status of laser cladding on titanium and its alloys: a review, Mater. Des. 58 (2014) 412–425, https://doi.org/10.1016/j.matdes.2014.01. 077. [4] Xiu-Bo Liu, Xiang-Jun Meng, Hai-Qing Liu, Gao-Lian Shi, Shao-Hua Wu, Cheng-Feng Sun, Ming-Di Wang, Long-Hao Qi, Development and characterization of laser clad high temperature self-lubricating wear resistant composite coatings on Ti-6Al-4V alloy, Mater. Des. 55 (2014) 404–409, https://doi.org/10.1016/j.matdes.2013.09. 038. [5] M.F. Erinosho, E.T. Akinlabi, A review article: the mechanical properties and the microstructural behavior of laser metal deposited Ti-6Al-4V and TiC composite, Proc of SPIE 9741, 2016https://doi.org/10.1117/12.2214988 (97410E-1–10). [6] P.K. Farayibi, J.W. Murray, L. Huang, F. Boud, P.K. Kinnell, A.T. Clare, Erosion resistance of laser clad Ti-6Al-4V/WC composite for waterjet tooling, J. Mater. Process. Technol. 214 (2014) 710–721, https://doi.org/10.1016/j.jmatprotec.2013.08.014. [7] X. Wang, X. Ma, X. Li, L. Dong, M. Wang, Effect of boron addition on microstructure and mechanical properties of TiC/Ti6Al4V composites, Mater. Charact. 36 (2012) 41–46, https://doi.org/10.1016/j.matdes.2011.10.040. [8] K.S.S. Aradhya, M.R. Doddamani, Characterization of mechanical properties of SiC/ Ti-6Al-4V metal matrix composite (MMC) using finite element method, Am. J. Mater. Sci. 5 (3C) (2015) 7–11, https://doi.org/10.5923/c.materials.201502.02. [9] F. Wang, J. Mei, H. Jiang, X. Wu, Laser fabrication of Ti6Al4V/TiC composites using simultaneous powder and wire feed, Mater. Sci. Eng. A 445-446 (2007) 461–466, https://doi.org/10.1016/j.msea.2006.09.093. [10] R.M. Mahamood, E.T. Akinlabi, M. Shukla, S. Pityana, Scanning velocity influence on microstructure, microhardness and wear resistance performance of laser deposited Ti6Al4V/TiC composite, Mater. Des. 50 (2013) 656–666, https://doi.org/10.1016/j. matdes.2013.03.049. [11] R.M. Mahamood, E.T. Akinlabi, Laser metal deposition of functionally graded Ti6Al4V/TiC, Mater. Des. 84 (2015) 402–410, https://doi.org/10.1016/j.matdes. 2015.06.135. [12] H.A. Rastegari, S. Asgari, S.M. Abbasi, Producing Ti–6Al–4V/TiC composite with good ductility by vacuum induction melting furnace and hot rolling process, Mater. Des. 32 (2011) 2010–5014, https://doi.org/10.1016/j.matdes.2011.06.009. [13] S. Liu, Z. Liu, Y. Wang, P. Yue, Ti-based composite coatings with gradient TiCx reinforcements on TC4 titanium alloy prepared by laser cladding, Sci. China Technol. Sci. 57 (2014) 1454–1461, https://doi.org/10.1007/s11431-014-5566-5. [14] J.M. Wilson, Y.C. Shin, Microstructure and wear properties of laser-deposited functionally graded Inconel 690 reinforced with TiC, Surf. Coat. Technol. 207 (2012) 517–522, https://doi.org/10.1016/j.surfcoat.2012.07.058. [15] F. Weng, H. Yu, C. Chen, J. Dai, Microstructures and wear properties of laser cladding Co-based composite coatings on Ti-6Al-4V, Mater. Des. 80 (2015) 174–181, https:// doi.org/10.1016/j.matdes.2015.05.005. [16] Y. Zhang, Z. Wei, L. Shi, M. Xi, Characterization of laser powder deposited Ti–TiC composites and functional gradient materials, J. Mater. Process. Technol. 206 (2008) 438–444, https://doi.org/10.1016/j.jmatprotec.2007.12.055. [17] S. Liu, Z. Liu, Y. Wang, J. Tang, A comparative study on the high temperature corrosion of TP347H stainless steel, C22 alloy and laser-cladding C22 coating in molten chloride salts, Corros. Sci. 83 (2014) 396–408, https://doi.org/10.1016/j.corsci. 2014.03.012. [18] B. Vrancken, L. Thijs, J.-P. Kruth, J.V. Humbeeck, Heat treatment of Ti6Al4V produced by selective laser melting: microstructure and mechanical properties, J. Alloys Compd. 451 (2012) 177–185, https://doi.org/10.1016/j.jallcom.2012.07.022. [19] J. Yu, M. Rombouts, G. Maes, F. Motmans, Material properties of Ti6Al4V parts produced by laser metal deposition, Phys. Procedia 39 (2012) 416–424, https://doi. org/10.1016/j.phpro.2012.10.056. [20] M. Shunmugavel, A. Polishetty, G. Littlefair, Microstructure and mechanical properties of wrought and additive manufactured Ti-6Al-4V cylindrical bars, Procedia Technol. 20 (2015) 231–236, https://doi.org/10.1016/j.protcy.2015.07.037.
195
[21] R. Pelletier, L.P. Lefebvre, É. Baril, A MIM route for producing Ti6Al4V-TiC composites, Key Eng. Mater. 704 (2016) 139–147, https://doi.org/10.4028/www.scientific. net/KEM.704.139. [22] A.J.W. Johnson, K.S. Kumar, C.L. Briant, Deformation mechanisms in Ti-6Al-4V/TiC composites, Metall. Mater. Trans. A 34 (2003) 1869–1877, https://doi.org/10.1007/ s11661-003-0152-7. [23] K. Shah, L. Haq, A. Khan, S. Shah, M. Khan, A. Pinkerton, Parametric study of development of Inconel-steel functionally graded materials by laser direct metal deposition, Mater. Des. 54 (2014) 531–538, https://doi.org/10.1016/j.matdes.2013.08.079. [24] B. Zheng, T. Topping, J.E. Smugeresky, Y. Zhou, A.B. As, D. Baker, E.J. Lavernia, The influence of Ni-Coated TiC on laser-deposited IN625 metal matrix composites, Metall. Mater. Trans. A 41 (2010) 568–673, https://doi.org/10.1007/s11661-009-0126-5. [25] A.J.W. Johnson, C.W. Bull, K.S. Kumar, C.L. Briant, The influence of microstructure and strain rate on the compressive deformation behavior of Ti-6Al-4V, Metall. Mater. Trans. A 34 (2003) 295–306, https://doi.org/10.1007/s11661-003-0331-6. [26] N. Biswas, J.L. Ding, V.K. Balla, D.P. Field, A. Bandyopadhyay, Deformation and fracture behavior of laser processed dense and porous Ti6Al4V alloy under static and dynamic loading, Mater. Sci. Eng. A 549 (2012) 213–221, https://doi.org/10.1016/j. msea.2012.04.036. [27] E. Wycisk, A. Solbach, S. Siddique, D. Herzog, F. Walther, C. Emmelmann, Effects of defects in laser additive manufactured Ti-6Al-4V on fatigue properties, Phys. Procedia 56 (2014) 371–378, https://doi.org/10.1016/j.phpro.2014.08.120. [28] P.-H. Li, W.-G. Guo, W.-D. Huang, Y. Su, X. Lin, K.-B. Yuan, Thermomechanical response of 3D laser-deposited Ti-6Al-4V alloy over a wide range of strain rates and temperatures, Mater. Sci. Eng. A 647 (2015) 34–42, https://doi.org/10.1016/j.msea. 2015.08.043. [29] B.E. Carroll, T.A. Palmera, A.M. Beese, Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing, Acta Mater. 87 (2015) 309–320, https://doi.org/10.1016/j.actamat.2014.12.054. [30] X. Li, J. Xie, Y. Zhou, Effects of oxygen contamination in the argon shielding gas in laser welding of commercially pure titanium thin sheet, J. Mater. Sci. 40 (2005) 3437–3443, https://doi.org/10.1007/s10853-005-0447-8. [31] L.E. Murra, S.A. Quinones, S.M. Gaytan, M.I. Lopez, A. Rodela, E.Y. Martinez, D.H. Hernandez, E. Martinez, F. Medina, R.B. Wicker, Microstructure and mechanical behavior of Ti-6Al-4V produced by rapid-layer manufacturing, for biomedical applications, J. Mech. Behav. Biomed. Mater. 2 (2009) 20–32, https://doi.org/10.1016/j. jmbbm.2008.05.004. [32] J. Murray, ASM Metals Handbook—Alloy Phase Diagrams, Springer, New York, 1992. [33] R.T.C. Choo, J. Szekely, The possible role of turbulence in GTA weld pool behavior, Weld. J. 73 (1994) 25–31. [34] G. Yang, J. Ma, H. Wang, B. Carlson, R. Kovacevic, Studying the effect of lubricant on laser joining of AA 6111 panels with the addition of AA 4047 filler wire, Mater. Des. 116 (2017) 176–187, https://doi.org/10.1016/j.matdes.2016.12.014. [35] A.A. Antonysamy, J. Meyer, P.B. Prangnell, Effect of build geometry on the β-grain structure and texture in additive manufacture of Ti-6Al-4V by selective electron beam melting, Mater. Charact. 84 (2013) 153–168, https://doi.org/10.1016/j. matchar.2013.07.012. [36] X. Wu, J. Liang, J. Mei, C. Mitchell, P.S. Goodwin, W. Voice, Microstructures of laserdeposited Ti-6Al-4V, Mater. Des. 26 (2004) 137–144, https://doi.org/10.1016/j. matdes.2003.09.009. [37] Y. Zhai, H. Galarraga, D.A. Lados, Microstructure evolution, tensile properties, and fatigue damage mechanisms in Ti-6Al-4V alloys fabricated by two additive manufacturing techniques, Process. Eng. 114 (2015) 658–666, https://doi.org/10. 1016/j.proeng.2015.08.007. [38] N. Hrabe, T. Quinn, Effects of processing on microstructure and mechanical properties of a titanium alloy (Ti–6Al–4V) fabricated using electron beam melting (EBM), part 1: Distance from build plate and part size, Mater. Sci. Eng. A 573 (2013) 264–270, https://doi.org/10.1016/j.msea.2013.02.064. [39] H. Ogden, R. Jaffee, F. Holden, Structure and properties of Ti-C alloys, J. Mater. 7 (1) (1955) 73–80, https://doi.org/10.1007/BF03377457. [40] G. Welsch, G. Lütjering, K. Gazioglu, W. Bunk, Deformation characteristics of age hardened Ti-6AI-4V, Metall. Mater. Trans. A 8 (1977) 169–177, https://doi.org/10. 1007/BF02677278.
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The influences of melting degree of TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites Shunyu Liu, Yung C. Shin ⁎ School of Mechanical Engineering, Purdue University, Center for Laser-based Manufacturing, West Lafayette, IN 47906, USA
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The melting degree of reinforced TiC was controlled by the laser energy density and the TiC content. • The high melting degree of reinforced TiC caused the formation of dendritic carbides, inducing premature failure. • The detrimental dendritic carbides were eliminated by adjusting the laser deposition conditions and reinforced TiC size. • The improved mechanical properties and strengthening mechanisms for different melting degrees of TiC were investigated.
a r t i c l e
i n f o
Article history: Received 29 August 2017 Received in revised form 19 September 2017 Accepted 30 September 2017 Available online 02 October 2017 Keywords: Laser direct deposition Ti6Al4V-TiC composite Functionally graded material Microstructure Mechanical property
a b s t r a c t This study is concerned with the influences of melting degree of embedded TiC reinforcements on microstructure and mechanical properties of laser direct deposited Ti6Al4V-TiC composites and a functionally graded material. The melting degree of embedded TiC was controlled by the input laser energy density and the added TiC content. The formation of detrimental primary dendritic TiC grains was successfully avoided by properly adjusting the deposition conditions and the particle size range of TiC reinforcements. The resultant compression test revealed the ultimate strength increasing from 1381 ± 19 MPa to 1636 ± 23 MPa as the premixed TiC content increased from 0 to 15 vol% while a true strain of 0.141 ± 0.002 was still retained for 15 vol% TiC. The primary strengthening mechanism for composites with the most melting control of TiC is the solid solution strengthening induced by carbon, while that for the least melting control is dominated by the unmelted TiC particulates and the refined microstructure resulting from the resolidified carbides. The defect-free functionally graded Ti6Al4V-TiC with 0 to 40 vol% TiC achieved an increased hardness from HRC ~39 to HRC ~65. © 2017 Elsevier Ltd. All rights reserved.
1. Introduction Ti6Al4V alloy (Ti64) occupies about 50% of usages among all the titanium alloys due to its high strength, low weight ratio, outstanding ⁎ Corresponding author. E-mail address: [email protected] (Y.C. Shin).
https://doi.org/10.1016/j.matdes.2017.09.063 0264-1275/© 2017 Elsevier Ltd. All rights reserved.
corrosion resistance and superior biocompatibility [1]. Ti64 can also be reinforced with hard particles to build metal matrix composites (MMCs) so as to combine the high stiffness and hardness of the reinforcements with the toughness and damage tolerance of the Ti64 matrix, which leads to increased wear resistance and mechanical properties [2–5]. Due to the high reactivity of titanium and its alloys, the reinforcements have to be carefully chosen. The commonly embedded
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Table 1 Summarized laser deposition conditions for Ti6Al4V and Ti6Al4V-TiC composites with different TiC contents. Sample Premixed TiC size Melting Laser no. TiC (μm) degree power (vol%) (W) 1 2 3 4 5 6 7 8
0 1 1 5 5 10 15 PF1:0 PF2:40
– 90–150 b45 90–150 b45 90–150 90–150 – 45–150
– Least Most Least Most Least Least
300 290 375 265 350 230 200 300 → 200
Powder feed rate (g/min)
Laser scan speed (mm/s)
Layer thickness (μm)
1.5 1.8 1.2 1.8 1.3 1.8 2.0 1.5 → 0 0 → 2.8
15 15 15 15 15 15 15 15
254 254 254 254 254 254 228.6 228.6
Note: No. 8 condition was used for depositing Ti64-TiC functionally graded material. PF1 and PF2 mean powder feeder hopper 1 and 2.
reinforcements for titanium alloys are ceramics such as WC [6], TiN [2], TiB [7], SiC [8], TiC [5,7,9–11] or in situ reacted TiC through adding graphite carbon or carbon black into titanium alloys [12,13]. Among these reinforcements, TiC is highly compatible with titanium as it possesses a similar density to titanium, a high hardness, and chemical and thermal stability. The high biocompatibility also facilitates the potential application of Ti64-TiC MMCs as load-bearing implants. Laser direct deposition (LDD) is an effective and efficient method for preparing MMCs since it can simultaneously feed different powders into the molten pool to provide requisite compositions and properties as well as customized exterior shapes. The process has been successfully used to generate fully dense components to near-net shape tolerances with improved hardness [14,15], strength [5,9], wear resistance and corrosion resistance [15–17] over some other fabrication methods [18–21]. Functionally graded materials (FGMMCs) can also be produced easily by properly adjusting the powder feed rates [16,22,23]. Although TiC is thought to be thermodynamically stable in titanium and titanium alloys, partial melting always happens to the embedded TiC particulates during the LDD process [9,11,14,24]. Some studies showed the diffusion of carbon from the melted TiC into the Ti64 matrix, causing solid solution strengthening [22,25]. The dissolved carbon atoms can in situ react with titanium to form different types of resolidified TiC precipitates [9,12]. With an increase in embedded TiC content, the strength of the composites improves, but accompanies a substantial sacrifice of ductility [9,12,22]. In Wang et al.'s work [9], the Ti64-TiC MMCs almost completely lost their ductility with 8 vol% TiC added. Several factors are responsible for the low ductility of Ti64-TiC MMCs such as the existence of defects like initial voids and lack-of-fusion [26–28], and the high oxygen level and impurity contents involved
during the fabrication process [29,30]. To the authors' knowledge, the microstructure of resolidified TiC and the unmelted TiC (UMC) particles also have a big impact on the mechanical properties of MMCs, but no relative research has been conducted. In the current study, the melting degree of the TiC reinforcements and the morphologies of the resolidifed carbides are controlled by adjusting the LDD parameters and the TiC particle size range. By avoiding the formation of dendritic TiC precipitates, the mechanical properties of the LDD Ti64-TiC MMCs are dramatically improved. A Ti64-TiC FGMMC with an increased TiC content ranging from 0 to 40 vol% within 5 mm was also built. The FGMMC realized a defect-free microstructure and an elevated hardness. The results are presented and discussed in detail. 2. Experimental procedures Gas atomized spherical Ti64 powders (grade 5, b150 μm) obtained from Puris®, along with irregularly shaped TiC particles in the same size range obtained from Atlantic Equipment Engineers®, were used to premix 1 vol%, 5 vol%, 10 vol% and 15 vol% Ti64-TiC precursor powders. LDD Ti64-TiC MMCs were carried out with an Optomec LENS® 750 system, which was equipped with a 500 W continuous wave fiber laser with a wavelength of 1064 nm and a dual-hopper powder feeder system. During the LDD process, the input laser energy density can be roughly calculated as J= αP/VH, where J is the input laser energy density, P is the laser power, V is the laser scan speed, H is the layer thickness, and α is the laser absorptivity coefficient. The input energy density required to deposit a material is mainly determined by the laser absorptivity coefficient and the melting temperature. The laser scan speed was kept constant as 15 mm/s in the current study. Although the melting of TiC reinforcements cannot be avoided, the complete melting of TiC is also difficult to obtain, particularly when the TiC size is large, because the laser processing time is extremely short and the heating and cooling cycles are super-fast [31]. Therefore, in the present work, most melting and least melting conditions were used to indicate two extreme melting degrees. The LENS® 750 system is designed to feed powders with a size range of 45–150 μm. Therefore, 90–150 μm TiC particulates were used for the least melting control of embedded TiC, while b 45 μm TiC particles were used for the most melting control of TiC. The small particles have a larger surface area, and are easier to melt than larger particles. The irregularly shaped TiC powder did not flow well through the powder feed system, particularly when b 45 μm particles were used. Therefore, premixed Ti64-TiC powder mixtures were used to deposit Ti64-TiC MMCs. The mixtures were blended by stirring them for 20 min followed by tumbling them on a roller mill at 90 rpm for 12 h. For building the Ti64-TiC FGMMC, pure Ti64 powder and premixed Ti64-40vol%TiC were fed simultaneously using two powder hoppers.
Fig. 1. (a) Laser direct deposited Ti6Al4V or Ti6Al4V-TiC composite cylinder, (b) the dimension of the machined compression specimen.
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Fig. 2. Morphologies of TiC in Ti6Al4V-15vol%TiC composites deposited with 250 W + 254 μm layer height.
The FGMMC composition varied from 100% Ti64 for the first two layers to an increasing volume of TiC for the subsequent 20 layers. The powder feed rates of the two hoppers were adjusted to obtain a 2 vol% TiC increment for each layer to reach a final composition of 40 vol% TiC. A delay time of 10 s was set between adjacent layers to allow for adequate powder feed stabilization. The substrate used was a 3.5-mm thick hot rolled Ti64 plate whose surface was cleaned and degreased with ethanol prior to usage. Several 10 mm diameter cylinders were built in the argon-controlled working chamber, where the oxygen level was controlled below 10 ppm. After depositing one layer, the orientation of the laser scan was rotated by 60 degrees which can lead to a good cylindrical geometry, and the laser focal point was elevated in Z direction by a pre-set height, which was equal to the measured actual layer thickness. The laser deposition parameters for melting control of TiC particles to obtain a dense and desirable microstructure are summarized in Table 1. A room temperature compression test was carried out on the Ti64 parts and Ti64-TiC MMCs with 1, 5, 10 and 15 vol% TiC at a strain rate
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of 0.005/min on an MTS Sintech 30/D hydraulic test platform with a 100 kN load cell. A CNC lathe was used to turn the surface of the cylinders. The images of the as-built and machined specimens are illustrated in Fig. 1. The true stress and strain data were measured from these tests. To characterize the microstructure, the specimens were sectioned parallel to the building direction, mounted in bakelite, and polished with #240–#1200 SiC papers followed with vibratory polishing for 24 h. To observe the morphologies of the resolidified carbides, the samples were not etched, while the Kroll's etchant was used to reveal the microstructure of the Ti64 matrix. The microstructure details were observed by an optical microscope (Zeiss), and the fracture morphologies were observed using a scanning electron microscope (SEM, JEOL JSMT330) under the secondary electron mode. A LECO RT-370 Rockwell hardness tester was utilized to examine the macrohardness of the Ti64-TiC FGMMC specimens, which indicated the average hardness. A load of 100 kg and 15 s were used. The positions with different TiC contents were marked on the specimen by calculating and measuring the distance to the side with 40 vol% TiC, and the hardness test was preformed along the marked lines. For each TiC content, five indentations were measured on one specimen, and four specimens were measured. The values were finally averaged. 3. Results and discussion 3.1. Laser direct deposited Ti64-TiC composites 3.1.1. Microstructures before etching Fig. 2 shows five morphologies of TiC precipitates in the LDD Ti64TiC MMCs. The big particle is UMC. The melting of the embedded TiC particles can be visually observed according to the boundary shape of the TiC particles. If the melting degree is high, many tiny columnar carbide grains attach to the unmelted TiC (UMC) particles, and the boundary of the UMC particles displays a serrated shape as shown in Fig. 2. The big spherical carbide particles, which are uniformly distributed in the microstructure, are the primary spherical TiC (PSC). According to the Ti-C binary phase diagram [32], the PSC resulting from the in situ reaction of C and Ti first precipitates from the liquid at liquidus temperature. As the PSC stays in the liquid for a relatively long time, the particles grow large. If the melting degree of the embedded TiC particles is high, the primary dendritic TiC (PDC) will form (discussed later). When
Fig. 3. Tensile test fracture morphology of Ti6Al4V-10vol%TiC built with 350 W + 254 μm reveals dendritic TiC grains indicated by arrows.
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Fig. 4. The formation and morphology evolution of dendritic TiC in Ti6Al4V-10vol%TiC composites built with a 15 mm/s laser scan speed, a 254 μm layer height and the same TiC particle size but different laser powers: (a) 230 W, (b) 250 W, and (c) 300 W. Dendritic TiC grains are indicated by arrows.
the temperature decreases to the eutectic point, some eutectic TiC grains form. One is the eutectic grain TiC (EGC) that precipitates around the boundary of the prior β grains and exhibits a loop shape. Another is the eutectic spherical TiC (ESC), which displays a segregated small spherical morphology. In Ref. [12], the ESC grains are found mainly gathering in the triple junction. The LDD process will lead to a non-equilibrium solidification where the liquidus and eutectic temperatures will be a little different from those in the equilibrium phase diagram. The formation of PDC is harmful to the mechanical properties of Ti64-TiC MMCs since the dendritic grains will block the development of the strain, and it is easy for the dendrite tip to initiate cracks. Fig. 3 shows the tensile test fracture morphology of Ti64-10vol%TiC deposited with a high energy input. Some bright dendritic TiC grains, indicated by the arrows, are clearly observed. These dendritic grains display a brittle cleavage fracture behavior, and will change the fracture from ductile
behavior to brittle behavior, rendering the composite loss of the ductility [9]. In the current work, one goal is to avoid the formation of PDC grains by designing proper deposition conditions. Compression tests were carried out since less work was performed on the compression test than the tensile test for Ti64-based composites. However, the tensile fracture can show a clear view of the dendrites. Fig. 4 illustrates the formation mechanism and morphology evolution of PDC in Ti64-10vol%TiC composites. When the laser power was 230 W, only PSC, EGC and ESC were observed, while several short PDC grains were produced for 250 W. As the laser power further increased to 300 W, more and larger PDC grains were produced. Actually, the increased laser power led to an increased laser energy as other parameters were kept constant, and the melting degree of the embedded TiC hence increased. This result reveals the direct relationship between the formation of PDC and the melting degree of embedded TiC particles.
Fig. 5. Potential defects in the least melting control of laser direct deposited Ti6Al4V-TiC composites: (a) lack-of-fusion, and (b) inhomogeneous distribution of unmelted TiC particulates.
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Fig. 6. Microstructures of unmelted TiC particles and resolidified carbides in laser direct deposited Ti6Al4V-TiC composites with the least melting control of TiC reinforcements: (a, b) 1 vol% TiC; (c, d) 5 vol% TiC; (e, f) 10 vol% TiC; (g, h) 15 vol% TiC.
It is believed that as the melting degree of the embedded TiC increased, more carbon dissolved in the liquid. When the carbon content reached to a certain amount, the PDC began to form and attached to PSC as nuclei. The primary carbides consumed the carbon, and hence the eutectic carbide content decreased (Fig. 4(c)).
In addition to the laser energy density, the embedded TiC content also affects the melting degree. According to the microstructural analysis, 5 vol% TiC or less will not produce PDC even when the embedded TiCs were almost completely melted, as will be shown later. To avoid the formation of PDC, the LDD parameters must be carefully selected
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Fig. 7. Microstructures of Ti6Al4V-TiC with the most melting control of TiC: (a) 1 vol% TiC; (b) 5 vol% TiC.
for N5 vol% TiC. Therefore, the most melting control condition was only performed on embedded 1 vol% and 5 vol%TiC. For brevity, hereafter we define the most melting conditions for 1 and 5 vol% TiC as 1M and 5M, and the least melting conditions for 1, 5, 10 and 15 vol% TiC as 1L, 5L, 10L and 15L. For the least melting control, the lowest energy density condition was tested by increasing the powder feed rate so as to increase the layer height for a set laser power and laser scan speed. However, two kinds of defects are easily formed in this process and should be avoided (Fig. 5). If the laser energy density is not high enough for the preset layer thickness, a lack-of-fusion defect will be produced, resulting in a premature failure of the composites [8,26,27]. This lack-of-fusion defect exhibits an irregular-shaped void, and is primarily formed in the interpass overlap (Fig. 5(a)). On the contrary, a high energy density will produce a deep molten pool, leading to a high layer thickness. The vigorous Marangoni flow and buoyancy force in the molten pool will push the UMC particles to the edge zones of the molten pool [33,34]. The nonuniform distribution of the UMCs (Fig. 5(b)) will cause anisotropy in the mechanical properties. Finally, the defect-free condition with the lowest energy input was selected as the optimized deposition condition for the least melting control. For different embedded TiC contents, the optimized laser deposition conditions are different (Table 1). The microstructures of the UMC and resolidified TiC are illustrated in Fig. 6. As can be seen, the deposition conditions produce a dense and defect-free structure where no PDC grains are formed and the UMC particles are uniformly distributed in the Ti64 matrix for all compositions. Fig. 7 shows the microstructures of the resolidified TiC with the most melting control of reinforced TiC in Ti64-TiC MMCs. No PDC was detected even though the embedded TiC particles were almost completely melted. The resolidified TiC in 5M displays a needle-shaped or spherical-shaped eutectic morphology (Fig. 7(b)), but for 1M, no obvious
resolidified carbides are observed (Fig. 7(a)). This indicates that the melted TiC completely dissolves into the Ti64 matrix. 3.1.2. Microstructures after etching Fig. 8 shows the microstructure of the LDD monolithic Ti64 after etching. During LDD fabrication, the high cooling rate always results in the formation of acicular α′ martensite instead of α + β lamellae within the prior β grain in titanium alloys [19,31]. In Fig. 8, the fine needleshaped phase within the prior β grains is the α′ martensite. The prior β grains epitaxially grow and are almost parallel to the deposition direction with slight inclinations due to the scan rotation (Fig. 8(a)) [35]. The length of the β grain could be approximately 20 mm [9,36]. Growing from the boundaries of the prior β grains, a layer of α phase (grain boundary α) is formed, which separates the prior β grains (Fig. 8(b)). Fig. 9 shows the etched microstructures of the LDD Ti64-TiC MMCs. Compared with LDD Ti64 (Fig. 8), the presence of the TiC dramatically changed the Ti64 matrix microstructure. Instead of the acicular α′ martensite, a thin and short lamellar α + β structure was formed. In the optical micrographs, the bright lamella is α-Ti while the dark phase situating around α-Ti is β-Ti. The formed lamellar α + β structure is beneficial for raising the ductility [18]. For 1L, a lamellar α + β structure with 1–1.5 μm thick α-lath was obtained (Fig. 9(a)). This indicates a slower cooling rate as compared to that for Ti64. The slower cooling rate is caused by the more energy input. Since the laser power was decreased by 10 W, the more energy absorbed must come from the 1 vol% TiC. Hence, it is believed that TiC has a higher absorptivity than Ti64 for the wavelength of 1064 nm. As the laser power was increased to 375 W for 1M, the thickness of α-lath increased to 2–4 μm, which was caused by the further decreased cooling rate induced by the high energy density. However, with more TiC added, the Ti64 matrix changed back to a fine structure, where the α-lath is b 1 μm thick and the length is also
Fig. 8. Optical microstructure of laser direct deposited monolithic Ti6Al4V.
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Fig. 9. Optical micrographs of laser direct deposited Ti6Al4V-TiC composites reveal matrix structures: (a) least melted 1 vol% TiC; (b) most melted 1 vol% TiC; (c) least melted 5 vol% TiC; (d) most melted 5 vol% TiC; (e) least melted 10 vol% TiC and (f) least melted 15 vol% TiC.
decreased compared to that in Ti64-1vol%TiC composites. The refining effect is due to the precipitation of the resolidified TiC grains, which serve as the nuclei for the α and β phases and also block the growth of α-lath [21]. As 5M (Fig. 9(d)) shows, the high energy density cannot lead to a coarser structure. In addition, the long and columnar prior β grain in LDD Ti64 was reduced to small equiaxial grains, which can be visually evaluated according to the EGC grain size. The grain boundary α was also eliminated. All these microstructure changes are beneficial for the mechanical strength [29,37,38]. 3.1.3. Compression test Fig. 10 shows the room temperature compression test results of the LDD Ti64 and Ti64-TiC MMCs deformed at the strain rate of 0.005/min. Several duplicate tests were performed to make sure the results were repeatable. As the embedded TiC increased from 0 to 15 vol%, the ultimate strength (US) and yield strength (YS) were elevated from 1381 ± 19 MPa to 1636 ± 23 MPa, and from 997 ± 14 MPa to 1310 ± 22 MPa, respectively. Although the ductility decreased, a true fracture strain value of 0.141 ± 0.002 was still obtained for 15 vol% TiC.
The improved US and YS can be attributed to several factors, such as the lack of PDC grains, the refined α + β lamellar microstructure, the reduced prior β grain size [29,37,38], the thermal stress formation during the LDD process [24], and the increased dislocation density in the matrix [22]. However, the more important strengthening effects come from the embedded TiC particulates and their products, such as the free carbon atoms, the UMC and the resolidified carbides. Ogden et al. [39] and Johnson et al. [22,25] experimentally confirmed the positive role of the solid solution strengthening effect induced by carbon. The authors concluded that carbon in the solid solution provided a greater strengthening contribution than carbides, and the strengthening effect was obvious even if the carbon content was low. According to the compression test results in conjunction with the microstructure analyses as well as the LDD process feature, other strengthening mechanisms besides the solid solution strengthening induced by carbon have been explored in the current study. Actually, the primary strengthening mechanisms are different for different melting degrees and different reinforced TiC contents. For 1 vol% TiC, 1M has a coarser structure (Fig. 9(b)) than 1L (Fig. 9(a)).
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Fig. 10. Compression test results of Ti6Al4V-TiC composites built with different conditions: (a) true stress-strain curves; (b) Ultimate strength; (c) Yield strength; (d) Elongation. M and L indicate the most and least melting control condition of the reinforced TiC particles.
The thick α-lath is known to lower US, YS, and hardness [37,38]. However, 1M had a higher US and YS than 1L, but the ductility is only a little smaller. The primary strengthening for 1M must result from the solid solution strengthening induced by carbon since no resolidified TiC and neglected UMC were detected. However, for 1L, the solution strengthening is limited since the melting of TiC particulates is controlled to a very low level (Fig. 6(b)). Hence, the improved strength compared to LDD Ti64 should be attributed to the fine α + β lamellar structure, the secondary phase strengthening resulting from the UMC particles, and maybe a little bit of solid solution strengthening. As for 5 vol% TiC, 5M and 5L displayed similar US values. Since solid solution strengthening plays a limited role in 5L as compared to 5M, the UMC particles must contribute to the strengthening effect in 5L; otherwise, the mechanical properties for the two melting conditions would show a big difference. The elevated strength from 5L to 15L also confirmed the positive role of UMC in strengthening since the microstructure of the matrix and resolidified TiC are quite similar (Fig. 9(c, e and f)). The large volumes of resolidified carbides formed in the matrix should also cause precipitation strengthening. However, it is difficult to quantitatively evaluate the precipitation strengthening effect.
As α-Ti and β-Ti both have a limited solubility of carbon, and the dissolution of carbon beyond a certain amount will impair the ductility, the Ti64-TiC MMCs have to rely on UMC and resolidified carbides to improve the strength for a high TiC content. However, the UMC and the resolidified TiC increase the dislocation piles up for plastic deformation [40], and TiC particles are hard and brittle. The continuous adding of TiC will finally lead to a deteriorated property. 3.1.4. Compression fracture morphologies and mechanism After the compression tests, SEM was used to characterize the fracture morphologies of LDD Ti64 and Ti64-TiC MMCs to reveal the fracture mechanism. The fracture morphology of 15L is similar to that of 10L except for more UMC particles, so the former is not presented in Fig. 11. With a low TiC content, the fracture displays a ductile morphology (Fig. 11(a–d)), where the deep dimple structure indicates a ductile fracture. The flat zones reveal the slip bands. The highest amount of dimple structure is observed in 1L, but not in LDD Ti64, which may be due to the acicular α′ martensitic structure in LDD Ti64. In 1M, some white spherical carbide spots are detected. This indicates the formation of PSC grains.
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Fig. 11. Compression fracture morphologies of laser direct deposited Ti6Al4V and Ti6Al4V-TiC compositions: (a) Ti6Al4V; (b) most melted 1 vol% TiC; (c) most melted 5 vol% TiC; (d) least melted 1 vol% TiC; (e) least melted 5 vol% TiC; (f) least melted 10 vol% TiC.
As the TiC content increases from 1 vol% to 10 vol%, the deep dimple morphology becomes shallow and reveals a reduction in ductility. The disappearance of the deep dimple is due to the formation of resolidified TiC grains. The UMC particles reveal a smooth and brittle feature. Therefore, the fracture mechanism is a mixture of a subdued ductile mechanism in the matrix and a cleavage mechanism in the UMC particles. The high volume percentage of UMC in 10L and15L speed up the failure of the components. Johnson et al. [14,22] investigated the cracking and failure mechanisms in the Ti64 and Ti64-TiC MMCs during compression tests. In summary, for the Ti64 samples tested at quasi-static small strain rates, the voids forming in the most severely deformed region caused the failure. As for the composites, the particle-matrix interface provided sites for additional void formation through interface separation, and thus the authors concluded that the increased density of these voids accounted for the decreased ductility of the composites with a high TiC content. In the
current work, voids were not observed. However, the debonding of TiC particles from the matrix and the formation of microcracks around the TiC-matrix interface were detected (Fig. 11(f)). The interfacial region between the reinforcements and the matrix governs the mechanical properties of the composites. The debonding or weak bonding and microcrack formed at the interface will lead to the premature failure of the Ti64-TiC MMCs. Therefore, it is necessary to adjust the laser deposition condition to avoid the formation of PDC phase, while producing a firm bonding between the embedded TiC particulates and the Ti64 matrix. 3.2. Laser direct deposited Ti64-TiC functionally graded material Instead of serving as structural materials, the Ti64-TiC MMCs can be built as FGMMC by LDD, and used as surface coatings. The content of the TiC reinforcements and the microstructure of FGMMC change gradually
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Fig. 12. Laser direct deposited Ti64-TiC functionally graded material with (0–40) vol%TiC.
from one side to the other, resulting in gradual changes in the properties. Fig. 12(a) shows the optical micrographs of LDD Ti64-TiC FGMMC. A defect-free structure has been built. A large volume percentage of embedded TiC has been retained as UMC particles. The UMC and resolidified carbides are critical for improving the wear resistance, and the effective fraction of the reinforced TiC was experimentally proved not to be b 15 vol% [9]. As the premixed TiC content exceeds 15 vol%, it is difficult to eliminate the formation of the PDC grains (Fig. 12(b–d)). However, if the Ti64-TiC MMCs are used to improve the erosion and wear resistance of a material, the MMCs can be deposited as a FGMMC coating on the base material. Within a thin coating, a high volume fraction of TiC reinforcements can be embedded. Besides, the formed carbides will also be able to improve the hardness of the matrix. As can be seen from Fig. 13 with the volume percentage of TiC particles varying
from 0 to 40 vol%, the hardness is increased from HRC 39 ± 0.12 to HRC 65 ± 0.28. Considering the high volume of UMC and resolidifed carbides, the improvement of wear resistance of the base material can be expected [9]. 4. Conclusion The influences of the melting degree of embedded TiC on the microstructure evolution and mechanical properties of the Ti64-TiC MMCs and FGMMC were studied. 1. The melting degree of the embedded TiC was determined by the added TiC content and the input laser energy density. The laser energy density cannot cause Ti6Al4V-TiC composites with no N5 vol% TiC to form dendritic carbides. It was difficult to avoid the formation of dendritic carbides when the premixed TiC is larger than 15 vol%. By proper controlling the input laser energy density, the Ti6Al4V-TiC composites with 5–15 vol% TiC were free from dendritic carbides. 2. The dendrites-free compression specimens revealed increased ultimate and yield strengths but correspondingly decreased ductility as the premixed TiC content increased from 0 to 15 vol% while a true strain of 0.141 ± 0.002 was still retained for 15 vol% TiC. 3. The solid solution strengthening induced by carbon and the unmelted TiC particles played the dominant role in strengthening the composites with the most and least melting control of TiC, respectively. The resolidified carbides also caused the precipitation strengthening effect and refined microstructure. The unmelted TiC and resolidified carbides changed the fracture mechanism from the ductile behavior to cleavage behavior. 4. With the TiC increased from 0 to 40 vol%, the defect-free Ti64-TiC FGMMC achieved an increased hardness from HRC ~39 to HRC ~65, and increased unmelted TiC and resolidified TiC contents. Acknowledgement
Fig. 13. Rockwell hardness for laser direct deposited Ti64-TiC functionally graded material with TiC increasing from 0 to 40 vol%TiC.
The authors wish to acknowledge the help of Xianfeng Lu in the compression sample machining and testing.
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References [1] M.J. Donachie, Titanium: a Technical Guide, second ed. ASM International, Metals Park, OH, 1988. [2] Y. Lin, Y. Lin, Y. Chen, Evolution of the microstructure and tribological performance of Ti-6Al-4V cladding with TiN powder, Mater. Des. 36 (2012) 584–589, https://doi. org/10.1016/j.matdes.2011.12.007. [3] F. Weng, C. Chen, H. Yu, Research status of laser cladding on titanium and its alloys: a review, Mater. Des. 58 (2014) 412–425, https://doi.org/10.1016/j.matdes.2014.01. 077. [4] Xiu-Bo Liu, Xiang-Jun Meng, Hai-Qing Liu, Gao-Lian Shi, Shao-Hua Wu, Cheng-Feng Sun, Ming-Di Wang, Long-Hao Qi, Development and characterization of laser clad high temperature self-lubricating wear resistant composite coatings on Ti-6Al-4V alloy, Mater. Des. 55 (2014) 404–409, https://doi.org/10.1016/j.matdes.2013.09. 038. [5] M.F. Erinosho, E.T. Akinlabi, A review article: the mechanical properties and the microstructural behavior of laser metal deposited Ti-6Al-4V and TiC composite, Proc of SPIE 9741, 2016https://doi.org/10.1117/12.2214988 (97410E-1–10). [6] P.K. Farayibi, J.W. Murray, L. Huang, F. Boud, P.K. Kinnell, A.T. Clare, Erosion resistance of laser clad Ti-6Al-4V/WC composite for waterjet tooling, J. Mater. Process. Technol. 214 (2014) 710–721, https://doi.org/10.1016/j.jmatprotec.2013.08.014. [7] X. Wang, X. Ma, X. Li, L. Dong, M. Wang, Effect of boron addition on microstructure and mechanical properties of TiC/Ti6Al4V composites, Mater. Charact. 36 (2012) 41–46, https://doi.org/10.1016/j.matdes.2011.10.040. [8] K.S.S. Aradhya, M.R. Doddamani, Characterization of mechanical properties of SiC/ Ti-6Al-4V metal matrix composite (MMC) using finite element method, Am. J. Mater. Sci. 5 (3C) (2015) 7–11, https://doi.org/10.5923/c.materials.201502.02. [9] F. Wang, J. Mei, H. Jiang, X. Wu, Laser fabrication of Ti6Al4V/TiC composites using simultaneous powder and wire feed, Mater. Sci. Eng. A 445-446 (2007) 461–466, https://doi.org/10.1016/j.msea.2006.09.093. [10] R.M. Mahamood, E.T. Akinlabi, M. Shukla, S. Pityana, Scanning velocity influence on microstructure, microhardness and wear resistance performance of laser deposited Ti6Al4V/TiC composite, Mater. Des. 50 (2013) 656–666, https://doi.org/10.1016/j. matdes.2013.03.049. [11] R.M. Mahamood, E.T. Akinlabi, Laser metal deposition of functionally graded Ti6Al4V/TiC, Mater. Des. 84 (2015) 402–410, https://doi.org/10.1016/j.matdes. 2015.06.135. [12] H.A. Rastegari, S. Asgari, S.M. Abbasi, Producing Ti–6Al–4V/TiC composite with good ductility by vacuum induction melting furnace and hot rolling process, Mater. Des. 32 (2011) 2010–5014, https://doi.org/10.1016/j.matdes.2011.06.009. [13] S. Liu, Z. Liu, Y. Wang, P. Yue, Ti-based composite coatings with gradient TiCx reinforcements on TC4 titanium alloy prepared by laser cladding, Sci. China Technol. Sci. 57 (2014) 1454–1461, https://doi.org/10.1007/s11431-014-5566-5. [14] J.M. Wilson, Y.C. Shin, Microstructure and wear properties of laser-deposited functionally graded Inconel 690 reinforced with TiC, Surf. Coat. Technol. 207 (2012) 517–522, https://doi.org/10.1016/j.surfcoat.2012.07.058. [15] F. Weng, H. Yu, C. Chen, J. Dai, Microstructures and wear properties of laser cladding Co-based composite coatings on Ti-6Al-4V, Mater. Des. 80 (2015) 174–181, https:// doi.org/10.1016/j.matdes.2015.05.005. [16] Y. Zhang, Z. Wei, L. Shi, M. Xi, Characterization of laser powder deposited Ti–TiC composites and functional gradient materials, J. Mater. Process. Technol. 206 (2008) 438–444, https://doi.org/10.1016/j.jmatprotec.2007.12.055. [17] S. Liu, Z. Liu, Y. Wang, J. Tang, A comparative study on the high temperature corrosion of TP347H stainless steel, C22 alloy and laser-cladding C22 coating in molten chloride salts, Corros. Sci. 83 (2014) 396–408, https://doi.org/10.1016/j.corsci. 2014.03.012. [18] B. Vrancken, L. Thijs, J.-P. Kruth, J.V. Humbeeck, Heat treatment of Ti6Al4V produced by selective laser melting: microstructure and mechanical properties, J. Alloys Compd. 451 (2012) 177–185, https://doi.org/10.1016/j.jallcom.2012.07.022. [19] J. Yu, M. Rombouts, G. Maes, F. Motmans, Material properties of Ti6Al4V parts produced by laser metal deposition, Phys. Procedia 39 (2012) 416–424, https://doi. org/10.1016/j.phpro.2012.10.056. [20] M. Shunmugavel, A. Polishetty, G. Littlefair, Microstructure and mechanical properties of wrought and additive manufactured Ti-6Al-4V cylindrical bars, Procedia Technol. 20 (2015) 231–236, https://doi.org/10.1016/j.protcy.2015.07.037.
195
[21] R. Pelletier, L.P. Lefebvre, É. Baril, A MIM route for producing Ti6Al4V-TiC composites, Key Eng. Mater. 704 (2016) 139–147, https://doi.org/10.4028/www.scientific. net/KEM.704.139. [22] A.J.W. Johnson, K.S. Kumar, C.L. Briant, Deformation mechanisms in Ti-6Al-4V/TiC composites, Metall. Mater. Trans. A 34 (2003) 1869–1877, https://doi.org/10.1007/ s11661-003-0152-7. [23] K. Shah, L. Haq, A. Khan, S. Shah, M. Khan, A. Pinkerton, Parametric study of development of Inconel-steel functionally graded materials by laser direct metal deposition, Mater. Des. 54 (2014) 531–538, https://doi.org/10.1016/j.matdes.2013.08.079. [24] B. Zheng, T. Topping, J.E. Smugeresky, Y. Zhou, A.B. As, D. Baker, E.J. Lavernia, The influence of Ni-Coated TiC on laser-deposited IN625 metal matrix composites, Metall. Mater. Trans. A 41 (2010) 568–673, https://doi.org/10.1007/s11661-009-0126-5. [25] A.J.W. Johnson, C.W. Bull, K.S. Kumar, C.L. Briant, The influence of microstructure and strain rate on the compressive deformation behavior of Ti-6Al-4V, Metall. Mater. Trans. A 34 (2003) 295–306, https://doi.org/10.1007/s11661-003-0331-6. [26] N. Biswas, J.L. Ding, V.K. Balla, D.P. Field, A. Bandyopadhyay, Deformation and fracture behavior of laser processed dense and porous Ti6Al4V alloy under static and dynamic loading, Mater. Sci. Eng. A 549 (2012) 213–221, https://doi.org/10.1016/j. msea.2012.04.036. [27] E. Wycisk, A. Solbach, S. Siddique, D. Herzog, F. Walther, C. Emmelmann, Effects of defects in laser additive manufactured Ti-6Al-4V on fatigue properties, Phys. Procedia 56 (2014) 371–378, https://doi.org/10.1016/j.phpro.2014.08.120. [28] P.-H. Li, W.-G. Guo, W.-D. Huang, Y. Su, X. Lin, K.-B. Yuan, Thermomechanical response of 3D laser-deposited Ti-6Al-4V alloy over a wide range of strain rates and temperatures, Mater. Sci. Eng. A 647 (2015) 34–42, https://doi.org/10.1016/j.msea. 2015.08.043. [29] B.E. Carroll, T.A. Palmera, A.M. Beese, Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing, Acta Mater. 87 (2015) 309–320, https://doi.org/10.1016/j.actamat.2014.12.054. [30] X. Li, J. Xie, Y. Zhou, Effects of oxygen contamination in the argon shielding gas in laser welding of commercially pure titanium thin sheet, J. Mater. Sci. 40 (2005) 3437–3443, https://doi.org/10.1007/s10853-005-0447-8. [31] L.E. Murra, S.A. Quinones, S.M. Gaytan, M.I. Lopez, A. Rodela, E.Y. Martinez, D.H. Hernandez, E. Martinez, F. Medina, R.B. Wicker, Microstructure and mechanical behavior of Ti-6Al-4V produced by rapid-layer manufacturing, for biomedical applications, J. Mech. Behav. Biomed. Mater. 2 (2009) 20–32, https://doi.org/10.1016/j. jmbbm.2008.05.004. [32] J. Murray, ASM Metals Handbook—Alloy Phase Diagrams, Springer, New York, 1992. [33] R.T.C. Choo, J. Szekely, The possible role of turbulence in GTA weld pool behavior, Weld. J. 73 (1994) 25–31. [34] G. Yang, J. Ma, H. Wang, B. Carlson, R. Kovacevic, Studying the effect of lubricant on laser joining of AA 6111 panels with the addition of AA 4047 filler wire, Mater. Des. 116 (2017) 176–187, https://doi.org/10.1016/j.matdes.2016.12.014. [35] A.A. Antonysamy, J. Meyer, P.B. Prangnell, Effect of build geometry on the β-grain structure and texture in additive manufacture of Ti-6Al-4V by selective electron beam melting, Mater. Charact. 84 (2013) 153–168, https://doi.org/10.1016/j. matchar.2013.07.012. [36] X. Wu, J. Liang, J. Mei, C. Mitchell, P.S. Goodwin, W. Voice, Microstructures of laserdeposited Ti-6Al-4V, Mater. Des. 26 (2004) 137–144, https://doi.org/10.1016/j. matdes.2003.09.009. [37] Y. Zhai, H. Galarraga, D.A. Lados, Microstructure evolution, tensile properties, and fatigue damage mechanisms in Ti-6Al-4V alloys fabricated by two additive manufacturing techniques, Process. Eng. 114 (2015) 658–666, https://doi.org/10. 1016/j.proeng.2015.08.007. [38] N. Hrabe, T. Quinn, Effects of processing on microstructure and mechanical properties of a titanium alloy (Ti–6Al–4V) fabricated using electron beam melting (EBM), part 1: Distance from build plate and part size, Mater. Sci. Eng. A 573 (2013) 264–270, https://doi.org/10.1016/j.msea.2013.02.064. [39] H. Ogden, R. Jaffee, F. Holden, Structure and properties of Ti-C alloys, J. Mater. 7 (1) (1955) 73–80, https://doi.org/10.1007/BF03377457. [40] G. Welsch, G. Lütjering, K. Gazioglu, W. Bunk, Deformation characteristics of age hardened Ti-6AI-4V, Metall. Mater. Trans. A 8 (1977) 169–177, https://doi.org/10. 1007/BF02677278.