1 Effect of Trace Lanthanum Hexaboride and Boron Additions on Microstructure, Tensile Properties and Anisotropy of Ti-6Al-4V

Effect of trace Lanthanum Hexaboride and Boron additions on microstructure, tensile properties and anisotropy of Ti-6Al-4V produced by Additive Manufacturing

M.J Bermingham1, S D McDonald, M.S Dargusch

Centre for Advanced Materials Processing and Manufacturing, School of Mechanical and Mining Engineering, The University of Queensland, St Lucia, Queensland, 4072, Australia

Abstract

In this work trace lanthanum hexaboride (LaB6) and elemental boron are alloyed with Ti-6Al-4V and their effects on the microstructure, tensile properties (including anisotropy) and melt pool shape during Additive Manufacturing (AM) are investigated. During the melting process, the LaB6 scavenges oxygen and decomposes into La2O3 and TiB. The presence of the rare earth element drastically changes the apparent surface tension and shape of the deposited layers. This is attributed to the Heiple-Roper effect and could have benefits during AM in producing components with unsupported overhangs. The formation of eutectic TiB during the final stages of solidification results in highly directional TiB needles in between columnar grains that are aligned with the build direction. The slow cooling rate during deposition of approximately 90-100°Cs-1 produces very large TiB particles which can exceed 50µm in length. Although improving strength by up to 10%, under tensile stress the high aspect ratio TiB particles are sites for crack opening which leads to a decline in ductility in the longitudinal test direction and a corresponding increase in anisotropy over unmodified Ti-6Al-4V.

Keywords: titanium alloys; characterization; stress/strain measurements; grains and interfaces

1 Corresponding author: [email protected]

1.0 Introduction

Additive Manufacturing offers many advantages over conventional manufacturing processes and research efforts to advance this technology have grown considerably in recent years. Titanium alloys are strong candidates for AM because of their high material cost and the high cost of traditional manufacturing processes. Ti-6Al-4V is by far the most studied of all titanium alloys in AM and now researchers have a comprehensive understanding of the Ti-6Al-4V microstructure and properties inherent to AM. For instance, across multiple AM technologies with different cooling rates it is now understood that the solidification conditions during AM of Ti-6Al-4V generally favour epitaxial growth of textured columnar-β grains which can create material property anisotropy [1-7]. While incomplete fusion (and therefore porosity) between layers can also contribute to anisotropy during AM of Ti-6Al- 4V [8, 9], Kobryn and Semiatin [8] and Qiu et al. [10] have shown that, even after the removal of porosity through Hot Isostatic Pressing, or in the absence of porosity in the case of Carroll et al. [11], anisotropy still remains, on account of the microstructural texture. It has been proposed that grain boundary-α (which delineates the columnar prior β-grains) is susceptible to failure under tensile loads normal to the build direction and this can account for anisotropy [11, 12]. It is therefore important for future research efforts to find ways to prevent anisotropy. The most obvious way of achieving this is to produce homogeneous equiaxed microstructures during AM.

There are two approaches that may improve the homogeneity in the microstructure of α+β titanium alloys produced during solidification. The first is to eliminate the highly textured columnar β-grains that result from the prevailing solidification conditions and alloy constitution, i.e. promote the Columnar to Equiaxed Transition (CET). The second approach is to refine the α-phase and eliminate unfavourable variants such as grain boundary-α or large α-colonies that can be several millimetres in size and adopt texture from the parent β-grain. Both of these approaches may be achievable through selective chemical alloying, however, to date there has been limited progress made in developing new titanium alloys for this purpose. Silicon is an effective growth restricting solute based grain refiner in titanium [13] and recently Mereddy et al. [14] investigated Ti-Si alloys during Wire Arc Additive Manufacturing (WAAM) in an attempt to achieve homogenous equiaxed grains throughout the build. Although substantial refinement of the grain size occurred (including the presence of some fine grained equiaxed crystals), ultimately the addition was unable to completely suppress epitaxial growth and some columnar grains still formed. Boron has also been studied as a grain refiner in trace additions to Ti-6Al-4V during AM processes. Although the columnar grain structure remained, Bermingham et al. [15] demonstrated substantial refinement of the α-grain morphology, elimination of continuous

segments of grain boundary-α and an improvement in compressive strength with trace boron additions during Ti-6Al-4V AM. In light of the work by Carroll et al. [11], the elimination of grain boundary-α is another promising approach that may eliminate anisotropy and warrants further investigation. Additions of as little as 1000ppm boron to cast Ti-6Al-4V are known to prevent the formation of grain boundary-α [16].

Recently the ASTM released a standard specification for additive manufacturing Ti-6Al-4V by powder bed fusion (ASTM F2924). The standard provides guidance on the acceptable chemistry tolerance of the base alloy, and interestingly, it indicates a provision for up to 0.1wt% of any ‘other’ elements2, up to a maximum of 0.4wt% in total. Potentially, new alloys could be carefully developed based on Ti- 6Al-4V that contain trace additions capable of eliminating the microstructural directionality and the associated anisotropy of AM while still complying with this new standard. However, there are only a handful of known elements that are powerful enough to impart substantial refinement in such small concentrations. Boron [17, 18], Beryllium [19], Carbon [20] and several Rare Earth elements including La and Y [21-24] have all been reported to refine the microstructure of titanium alloys in small concentrations during solidification. Rare Earth elements are also known to reduce the dissolved oxygen concentration in titanium through a ‘scavenging’ phenomena [25] and this trait could be very beneficial for AM. Not only could this permit the use of lower quality, cheaper feedstock material (i.e. with higher oxygen contents), but it also provides a margin of safety for AM processes that are vulnerable to atmospheric contamination. For example, there is growing interest in developing large scale ‘out of chamber’ AM technologies capable of producing components several metres in size [26] outside of an inert gas or vacuum chamber, and therefore, any trace alloy addition that can reduce the dissolved oxygen content has clear potential in AM processes susceptible to atmospheric contamination.

The purpose of this study is to investigate the effect of trace Lanthanum hexaboride (LaB6) additions on the microstructure and tensile properties, including the effects of anisotropy, of Ti-6Al-4V components produced by Wire Arc Additive Manufacturing (WAAM). The level of addition is to be within the chemical tolerance specified by ASTM F2924, and thus considered to be present in ‘trace levels’. Although ASTM F2924 specifically relates to powder bed fusion AM, it is reasonable to expect that other AM processes including Directed Energy Deposition (such as WAAM, LENS etc.) would

2 ASTM F2924 specifies maximum levels for Al, V, Fe, O, C, N, H and Y but all other elements individually have a maximum allowance of 0.1wt%, up to a combined limit of 0.4wt%.

follow similar guidelines for acceptable alloy chemistry. Only a few researchers have considered using

LaB6 as a trace alloy element in titanium and these have been exclusively in casting and powder metallurgy processes [27-29]. When mixed with titanium the LaB6 is reported to decompose into TiB and La2O3 (via the reaction 2LaB6 + 12Ti + 3O → 12TiB + La2O3), and the subsequent TiB is reported to refine the microstructure while the lanthanum lowers the dissolved oxygen content. Therefore, this is a promising compound to explore as a trace alloy addition to Ti-6Al-4V for AM. For comparison, Ti- 6Al-4V modified with trace boron is also investigated. Although the effect of trace boron additions on the microstructure and compressive strength of Ti-6Al-4V has been reported [15], to date no research has investigated the potential improvement in tensile anisotropy associated with the elimination of grain boundary-α with the introduction of TiB during AM of Ti-6Al-4V.

2.0 Experimental Methods

The wire arc additive manufacturing technique was used for this study (details about the equipment available in [15]). Ti-6Al-4V wire was used as the feedstock and a wrought Ti-6Al-4V base plate was used as a substrate for the deposits. Three alloys were investigated in this study, namely Ti-6Al-4V, Ti-

6Al-4V + boron and Ti-6Al-4V + LaB6. As previously mentioned, the level of addition of B and La was targeted to be within the limits specified by ASTM F2924, i.e. <0.10wt% of each element. The boron and LaB6 containing alloys were prepared by coating the build surface with specially prepared alcohol based paints (each approximately 15wt% of LaB6 or B) prior to each deposition using a similar method as reported elsewhere [14, 15]. The average LaB6 powder size used in the paint was 10µm (Sigma Aldrich, 99%) and the average boron powder size was 1µm (Strem Chemicals, 92-95%). The average alloy chemistry sampled from multiple locations within the built components is shown in Table 1. Each deposit was created by moving the welding torch in a linear direction and feeding wire into the molten pool, which solidified to make a layer. A subsequent layer was then deposited over the first by increasing the height of the torch, however, between layers the direction of torch travel was reversed so that the component was built up in a zig-zag fashion. The temperature at the end of each layer deposition3 was measured using a non-contact IR pyrometer, calibrated for emissivity against ultra high purity titanium (99.995% purity) with the melting temperature known to be 1668°C and the β- transus temperature 882°C. A fitting factor was used to account for the dynamic emissivity during cooling. After each layer the build was allowed to cool to room temperature before the next layer

3 The temperature was measured at the end of the layer after the arc was terminated. It was not possible to measure during deposition as radiation from the arc interfered with the pyrometer and the trailing shield prevented direct measurement.

was deposited. Approximately halfway through the build (when the height was approximately 20mm), deposition was ceased and the components were subjected to a stress-relief at 480°C for 2 hours. After this, deposition continued until the approximate dimensions of the final build were 180mmx12mmx45mm (not measuring the base plate). Details of the deposition parameters are given in Table 2.

Table 1. Average chemical analysis of the components determined by ICP AES and Leco Combustion (in wt%). The composition of all alloys is within the tolerance specified in ASTM F2924.

Alloy O N Al V B La Ti-6Al-4V 0.07 0.02 6.09 3.86 - - Ti-6Al-4V + Boron 0.07 0.02 5.98 3.76 0.05 -

Ti-6Al-4V + LaB6 0.06 0.02 6.07 3.91 0.03 0.08

Table 2. Deposition parameters used during fabrication of the components.

Deposition Parameters Peak Current: 150Amp Base Current: 75Amp Arc Pulse: 5kHz Wire feed: 1.5m/min Wire: Ti-6Al-4V, ø=1.0mm Deposition speed: 50mm/min electrode-substrate gap: ≈5mm Vertical build interval: ≈3mm Substrate: Ti-6Al-4V Electrode: ø=2.4mm tungsten-rare earth Argon: 99.999% purity

After deposition the components were subjected to Hot Isostatic Pressing (HIPing) in a 150MPa argon atmosphere at 920°C for 4 hours. Tensile bars complying with ASTM E8 (ø4mmx20mm gauge) were machined from the builds in both the longitudinal (horizontal) and transverse (vertical) build orientations and tested at 0.5mm/min using an Instron machine. Specimens were extracted from the builds for metallographic analysis and prepared using conventional techniques. Photographs and cross-sections of the specimens and the orientations of the tensile bars are shown in Figure 1.

Figure 1. Cross sections of the built components. The addition of LaB6 (C) caused instability during deposition resulting in a highly undulated surface. The orientation of the tensile bars machined from the builds are shown: T (Transverse/Vertical) and L (Longitudinal/Horizontal). The instability in the melt pool with the addition of LaB6 may be exploitable if it is desired to produce unsupported overhangs (demonstrated by arrow in C).

3.0 Results and Discussion

Effect of LaB6 and boron on liquid metal bead shape

The addition of LaB6 to Ti-6Al-4V had a marked effect on the molten metal’s apparent surface tension and the associated geometry of the deposited bead produced during wire arc additive manufacturing. The aspect ratio of the bead (ratio of width to height) substantially reduced from approximately 4 (in

Ti-6Al-4V and Ti-6Al-4V+B) to ~1.7 in the alloy containing LaB6. The narrowing of the bead width and the extension of its height promoted unstable deposition and the resulting build structure exhibited extreme undulations in topography, as shown in Figure 1. This can be attributed to the presence of lanthanum as there was no discernible difference between surface tension and bead shape when comparing the Ti-6Al-4V and the Ti-6Al-4V + boron build.

The decrease in the deposited bead’s aspect ratio may suggest that the surface tension (γ) of the melt has increased when LaB6 is introduced. However, this is at odds with conventional wisdom given that rare earths are well known for their surface active properties and have been reported to reduce the liquid surface tension in many alloys of aluminium [30, 31], Co-Fe [32], silver [33] and tin based solders [34-36]. Lanthanum has a much lower surface energy than titanium [37], and on this basis alone would be expected to wet the surface and lower the bead aspect ratio, however, the chemical and physical interactions that occur in the melt are far more complex and these interactions may act to reverse the effect. For instance, in liquid steels the addition of cerium is known to initially cause a rapid decrease in surface tension as it is being dissolved, however, surface tension quickly increases as the rare earth begins to react with dissolved oxygen to form oxides [38]. Furthermore, oxygen is also highly surface active and is known to strongly decrease the surface tension in many metals [39], so the removal of dissolved oxygen and the formation of La2O3 in the present case could also act to increase the surface tension in Ti-6Al-4V.

There are also complex physical fluid flow processes that are likely to have a strong effect on the shape of the molten bead. Several fluid flow mechanisms are known to operate during Gas Tungsten Arc Welding (and therefore WAAM) including thermocapillary (Marangoni), electromagnetic (Lorentz), aerodynamic plasma drag and buoyancy forces. However, under normal welding conditions it has been determined from multiple mathematical models that Marangoni convection is the dominant fluid flow force influencing the molten pool shape [40]. Marangoni forces operate in the molten pool due to steep temperature gradients that exist across the pool which, firstly create a surface tension gradient and then a surface flow from the low to high surface tension regions [41]. Since the surface tension of most pure metals decreases with increasing temperature (i.e. the surface tension temperature coefficient (dγ/dT) is negative) [42], the fluid flow during WAAM is expected to move from the centre of the weld pool outwards, as depicted in Figure 2. However, extensive experimental testing in steel welds has revealed a powerful effect generated by the addition of surface active elements to the melt. Heiple & Roper [43] found that the addition of trace surface active elements such as sulphur in steel reverse the surface tension temperature coefficient (dγ/dT) so that it becomes a positive value. In other words, the colder liquid now has a lower surface tension than the hotter liquid. This has a dramatic effect on the weld penetration and pool shape because the direction of fluid flow reverses and now moves towards the centre of the weld [42]. In the present case, if lanthanum was to act in a similar way to how sulphur acts in steel, then it is conceivable that marangoni forces could be responsible for the taller and thinner beads observed in the Ti-6Al-4V+LaB6 deposits. Yin et al. [44] recently reported the narrowing of build walls produced by WAAM with the

addition of CaF2 (believed to be a surface active addition) to Ti-6Al-4V components built by WAAM and also attributed this change to a reversal in the direction of fluid flow. With such a powerful effect on the shape of the component, further study is needed in this important area in order to understand the role of lanthanum and other surface active elements on the shape of titanium components produced by AM. This phenomenon could be exploited, for example, it may be possible to fabricate components with unsupported overhangs during AM (see arrow in Figure 1C). Wu et al. [45] propose low heat input arc pulsing during WAAM enables unsupported Ti6Al-4V overhangs and after many progressive layers was able to produce a 90° unsupported overhang. With the addition of LaB6 is appears that similar overhangs could be possible with only 1 layer, however, more work is needed in this space.

Figure 2. Schematic diagram showing how the Heiple-Roper effect may influence the deposited bead shape in AM. Marangoni forces create fluid flow from regions of small tension (γ) to regions of large surface tension. In most metals (A), hotter liquid will have smaller γ than colder liquid. Consequently, fluid flows out from underneath the arc (hottest region) to the cooler sides, producing wide melt pools with lower bead height. (B) - The Heiple-Roper effect associated with the introduction of surface active elements (La in this case) causes hotter liquid to have a larger γ than colder liquid, and this causes a reverse in fluid flow directions [42]. The diameter of the melt bead decreases and its height increases.

3.2 Effect of LaB6 and boron on solidification process

Representative cooling curves from the liquid state for each alloy are given in Figure 3. It is important to note that these measurements were collected at the end of a layer after the arc was terminated and may not exactly resemble the cooling rates during steady-state deposition. Nevertheless, they provide useful insights into the effects of boron and LaB6 additions on the solidification process of Ti- 6Al-4V. From these curves it is clear that both additions have a noticeable effect. Initially, all alloys cool rapidly during the liquid state where the temperatures exceeds 1700°C (average cooling rate ≈1100°Cs-1). As the temperature continues to fall the cooling rate decreases and a clear thermal arrest with duration of around 1.5s is apparent on the curves of each alloy, representing the solidification of primary Ti. During this plateau, the latent heat of fusion that is evolved, more or less balances with the rate heat extraction. During the final stages of solidification the rate of heat extraction exceeds the rate of latent heat release, which again leads to an increase in cooling rate. The average cooling rate immediately preceding the thermal arrest was 93°Cs-1. This cooling rate was averaged over a ΔT=20°C period before recalescence and provides an approximation for the cooling rate during initial solidification. It is also worth noting that this cooling rate is far lower than typical cooling rates reported during laser based additive manufacturing processes (which are estimated to range between 103-108 °Cs-1 [46]).

The presence of boron or LaB6 decreases the solidification temperature of Ti-6Al-4V from around 1679°C to approximately 1660-1662°C and this is not unexpected as boron decreases the equilibrium liquidus temperature. The cooling curves of both Ti-6Al-4V and Ti-6Al-4V+B display a degree of recalescence prior to the primary solidification plateau of ≈17°C and 9°C respectively but this is not present when LaB6 is alloyed. It is noted that the absence of recalescence has been associated with grain refinement in the literature (see Figure 9 in Shabestari et al. [47]) and this is the focus of ongoing research.

The derivatives of the cooling curves (cooling rate) for Ti-6Al-4V+B and Ti-6Al-4V+LaB6 also show clear evidence of a reaction occurring around 1465°C in these alloys. This is likely to be the final stages of eutectic TiB solidification given that TiB has a large exothermic heat of formation [48]. The reaction is not observed in Ti-6Al-4V. These events provide an indication of the total solidification time which was measured to be 1.6s, 1.8s and 1.9s for Ti-6Al-4V, Ti-6Al-4V + B and Ti-6Al-4V + LaB6, respectively. No

other reactions that possibly correspond to La2O3 or La-containing intermetallics can be discerned during the cooling of Ti-6Al-4V + LaB6.

Figure 3. Typical cooling curves and derivatives (over a 100ms period) for each alloy. Left images show the cooling curve from liquid to approximately 800°C, and the Right images show a magnified view of the primary solidification event.

3.3 Tensile Properties and Microstructure

Figure 4 shows the results of tensile testing. The addition of trace boron and LaB6 had a strong influence on the mechanical properties and generally increased strength at the expense of ductility.

The addition of LaB6 had the strongest effect: tensile strength increased by about 10% but the ductility, expressed as elongation at failure (εf) and reduction in area, both decreased significantly. The typical

εf of Ti-6Al-4V produced by wire arc additive manufacturing has been reported by others to be between 7-10% in the longitudinal direction and 12-16% in the transverse build direction [2-5]. Thus the transverse orientation is generally about 60% more ductile than the longitudinal orientation in these works. However, in the present study the HIPed Ti-6Al-4V components were much more isotropic in strength and ductility (εf ≈20% in the longitudinal orientation and εf ≈23.5% in the transverse orientation). The initial intention of adding the boron was that it may refine the α-phase and eliminate the α-phase that forms along β-grain boundaries (known as grain boundary-α) as was previously found [15]. However, instead of improving isotropy, the additions of boron and LaB6 actually had the reverse effect and increased anisotropy. The longitudinal orientation was about 50% less ductile than the transverse orientation with either boron or Lab6 additions. Ductility in the transverse orientation was satisfactory with both boron and LaB6 additions (εf ≈19.5% and 13.5% respectively, exceeding the 10% minimum required for wrought Ti-6Al-4V according to AMS4928) but the ductility in the transverse orientation was considerably lower (εf <8.5% in both cases).

Figure 4. Tensile properties of each alloy in the Transverse and Longitudinal build orientations. Error bars represent ±1 standard deviation.

Examination of the microstructure in Figure 5 reveals the presence of highly orientated TiB ‘needles’ aligned with the build direction in both Ti-6Al-4V+B and Ti-6Al-4V+LaB6. Some of the borides were very large and almost perfectly straight (>50µm long and ≈5mm thick) and were orientated in near perfect vertical directions (aligned with the transverse direction) while other regions contained clusters of much smaller borides that, while still orientated in a general vertical direction, took on a more dendritic appearance and outlined columnar grain boundaries.

The generalised preference for vertical TiB alignment parallel with the build direction is the likely microstructural feature accounting for the strong anisotropy measured. The alignment occurs because the eutectic TiB forms during the final stages of solidification which occurs at the edges of the columnar grains and/or within interdendritic pockets. As the primary β-Ti dendrite (or cell) grows it rejects boron solute which accumulates between growing grains. The steep thermal gradients in AM ensure that the columnar β-Ti grain continues to advance with the moving heat source but the presence of boron solute depresses the freezing range (by over 100°C) and results in an accumulation of solute rich liquid between the columnar grains. This effectively restricts growth in the lateral (horizontal) direction resulting in thinner columnar grains [15]. The concentration of the boron rich liquid (with low melting point) between the columnar grains continues to increase to the eutectic composition where borides are formed and solidification is complete. The solidification of the Ti-TiB eutectic is exothermic [48] and is detected on the cooling curves presented in Figure 3, and is consistent with the temperatures corresponding to formation of TiB according to the Ti-B phase diagram. The very small concentration of boron solute in the alloys (≈0.05wt%) essentially creates a divorced TiB eutectic at the final stage of solidification, predominately in between the columnar grains. The nature of this solidification process during AM is the subject of ongoing work.

Figure 5. SEM images of the alloys – (A) Ti-6Al-4V, (B) Ti-6Al-4V + B, (C) Ti-6Al-4V + LaB6. These images were taken in the lower third section of the components (i.e. at build height equivalent to approximately 15-17mm) so have likely undergone remelting (and grain growth through reheating) as new layers are deposited on top. Due to atomic contrast the dark phase indicates the presence of light atomic elements (e.g. boron) and the white phase indicates heavy elements (e.g. lanthanum).

The microstructure of the HIPed Ti-6Al-4V is predominantly widmanstatten-α structure in a β-matix. The relatively low tensile strength (around 800MPa) is attributed to the coarse α-phase that results

from the slow cooling from the β-phase field into the α+β phase field (cooling rates during deposition of 10-20°Cs-1 in the present work). It is worth noting that this is a key difference compared to other AM technologies such as SLM which can experience cooling rates exceeding 400°Cs-1 during the β→α transformation. The already relatively coarse α-phase then further coarsens through grain growth during the high temperature HIPing that was performed post build. The α-colony size (and α-lath size in highly Widmanstatten structures) is an important microstructural feature that determines the slip length in titanium alloys and has been shown to follow a Hall-Petch relationship with tensile strength [49]. A positive benefit of the coarse α-morphology is that very high ductility is achievable which exceeded 20% in the present study. This is high for Ti-6Al-4V and is actually twice the ductility requirement for wrought Ti-6Al-4V components according to AMS4928. The effects of trace boron on the morphology of the α-phase of Ti-6Al-4V has been reported in detail elsewhere [15, 16] and will not be explored further here, however, the same refinement effect (decrease in α-lath aspect ratio) on the α-phase was observed in the present study for both alloys containing boron. Similarly, the strengthening mechanisms of TiB reinforcement in Ti-6Al-4V is covered in detail elsewhere [50-52].

The addition of trace LaB6 into Ti-6Al-4V during WAAM decomposed into TiB particles and La rich particles. The low volume fraction of these particles (0.08wt% La was added) made detection using lab XRD techniques difficult (results not shown) but it has been shown with XRD that in higher concentrations (Ti-6Al-4V-1LaB6) the addition of LaB6 decomposes into La2O3 [28]. It was noted that the TiB plates adopting the same morphology as that observed in the Ti-6Al-4V+B alloy. However, unlike the TiB, the La2O3 forms as very small round particles up to ≈3µm in diameter (but usually much smaller) which are dispersed evenly throughout the microstructure. This could indicate that they are not formed during solidification, but rather, are present prior to the first β-titanium crystal nucleating.

Yang et al. [53] proposed that La2O3 does not form immediately during decomposition of LaB6 and suggested it forms during the β→α transformation (on account of La having ≈8wt% theoretical solubility in β-Ti but almost no solubility in α-Ti). However, this scenario is less likely than La2O3 spontaneously forming during decomposition of LaB6 on account of the low Gibbs energy of the rare earth oxide. When metallic rare earths are introduced to liquid steels the rare earths are known dissolve and immediately scavenge oxygen (forming oxides) [38]. Furthermore, the cooling curves presented in Figure 3 do not reveal any evidence of other reactions that could indicate the formation of La based compounds during cooling, although the upper limit on the pyrometer is 1800°C so any reactions above this temperature are not detected. Lastly, the uniform distribution of La2O3 throughout the alloy would not be expected if La2O3 precipitated during the β→α transformation because the partion coefficient (determined from binary Ti-La phase diagram) is 0.43 [19] which would

result in segregation effects creating localised regions richer in La solute and these regions would match the location of the borides, which clearly does not occur. In any case, the homogenous distribution of small oxides is unlikely to contribute to the tensile property anisotropy but they have been attributed to dispersion strengthening mechanisms [53] which could account for the higher strength (and loss of ductility) in alloys containing LaB6 compared to boron only additions.

3.4 Sources of anisotropy

Examination of the fracture surfaces and the microstructure within the broken tensile bars confirms microstructural factors that contributed to fracture. Firstly, in standard Ti-6Al-4V free from boron or

LaB6 additions, it was identified that prior-β grain boundaries and the associated grain boundary-α phase was prone to fracture. Grain boundary-α is known to nucleate cracks under fatigue loading [54] and evidence of void coalescence and crack propagation along grain boundary-α under tension is shown in Figure 6. Carrol et al. [11] found that the grain boundary-α is particularly important during AM processes because it provides a pathway for damage accumulation and correlated the columnar grain texture to anisotropic ductility. Therefore, highly elongated columnar grains are susceptible to forming continuous grain boundary-α and ductility is lower in the transverse direction because these segments are unfavourably orientated with the tensile loading axis and are prone to Mode I fracture. Wilson-Heid and co-workers expanded upon this work and proposed that anisotropy in the ductility of Ti-6Al-4V can be predicted based on the grain aspect ratio (grain height/grain width) that develops during deposition and correlated lower anisotropy with lower β-grain aspect ratio [55].

Figure 6. Ti-6Al-4V fractured tensile test pieces showing that fracture has occurred along the grain boundary-α that forms along prior-β grain boundaries. The prior-β grain boundaries have been enhanced (red) to more clearly show these in the top left figure. This particular test specimen failed at εf=12%, which is significantly below the average Ti-6Al-4V( εf≈21%).

Figure 7 shows microstructures of failed tensile specimens in both orientations and confirms that the highly orientated borides are responsible for increased anisotropy. In both Ti-6Al-4V + B and Ti-6Al-

4V + LaB6, TiB needles were found to contain severe cracking normal to the direction of tensile loading. The tensile specimens tested in the transverse (vertical) orientation (which contains the majority of borides aligned in this direction) resulted in the borides cracking multiple times throughout their length but due to the high aspect ratio of the TiB needles these cracks were limited to the width of the boride (about 5µm). However, when the tensile load was applied in the longitudinal orientation the cracking occurs along the length of the boride, which can exceed 50µm in large borides. The La2O3 particles were also found to separate from the titanium matrix but because these particles were very small the crack size was smaller than those found on TiB. The presence of large borides in the microstructure of AM parts is therefore unfavourable due to these particles nucleating large cracks.

Evidence of the cracked borides (and La2O3) was also observed on the fracture surface as seen in Figure 8. Others have also reported cracking of TiB under tension in cast Ti-6Al-4V-0.14B [56] and Ti-5Al-5Zr-

0.5Mo-0.25Si-0.5B [57]. Although fatigue properties were not evaluated here it would be reasonable to expect that the presence of large borides may be detrimental to fatigue life because the internal strain incompatibilities foster crack nucleation. However, Sen et al. [58] demonstrated that an increasing fraction of TiB reinforcement in cast Ti-6Al-4V-0.55B increased the high cycle fatigue strength4 by over 50%. The approximate length of the borides in the cast alloy was 20-50µm long and they were evenly dispersed. This is a surprising discovery and it was suggested that the removal of grain boundary-α with increased TiB volume fraction could have resulted in better fatigue performance. Improved fatigue endurance was also reported with boron additions (up to 0.5wt%) in Ti-6Al-4V produced by metal injection moulding [59, 60]. Nevertheless, the effect of such large and highly textured TiB on the fatigue properties in the present study requires further investigation.

4 Fatigue strength was defined by Sen et al. as the maximum stress that at least 75% of samples survived 106 cycles.

Figure 7. Microstructures of fractured tensile specimens showing cracking of the borides which generally align with the build direction. In specimens containing LaB6, separation of the La2O3 particles from the titanium matrix was also observed.

Figure 8. Example of fracture surface in boron containing alloys tested in the longitudinal (horizontal) orientation. A) Ti-

6Al-4V + B; B) high magnification of box in A, C) Ti-6Al-4V + LaB6 showing TiB and La2O3 particles at fracture surface. Arrows indicate cracked TiB.

Several researchers have produced titanium-boron alloys (or metal-matrix composites) during additive manufacturing but the deleterious effect of large textured borides on anisotropy has not yet

been reported. This in part is related to the variations in cooling rate during deposition and the composition of alloys studied. Banerjee et al. [48] produced Ti-2wt%B composites during blown powder laser deposition and compared this with the same alloy produced by arc-melting and casting. In both cases the alloys contained TiB particles randomly distributed throughout the microstructure but the main difference being that the additive manufactured components had considerably smaller TiB than the cast samples (20-100µm in cast vs 0.5-3µm in AM). Attar et al. [61] also compared Ti- 1.55%B alloys produced via casting and SLM and reported the same finding as Banerjee and co- workers, that is, the microstructure in the AM produced alloy contained randomly distributed TiB that was significantly smaller (typical length 0.1-0.5µm in AM) than the cast alloys. Wang, Mei and Wu [62] and Genc et al. [63] produced Ti-6Al-4V-TiB metal matrix composites (containing up to 3.1wt% boron) through blown powder laser deposition and were able to achieve homogenous microstructures containing fine TiB and reported no elongated TiB in build direction. In all of these studies the alloy compositions were at near eutectic or hypereutectic compositions. This will have important implications for the solidification sequence because in hypereutectic alloys the TiB will form before the β-Ti phase, and therefore, their orientation in the microstructure is not dictated by the solidifying β-Ti. This is in stark contrast to the present study where trace boron (hypoeutectic) precipitates as TiB from the final pools of liquid, which under the prevailing solidification conditions, occurs between solidified columnar β-Ti grains.

The other notable difference between the present study and the prior literature is the size of the TiB that forms which is a function of cooling rate. The relatively slow cooling rate in the present study (of the order of 90-100°Cs-1) produced extremely coarse TiB that tended to accumulate in a single region. Other AM processes experience higher cooling rates which produces TiB that are orders of magnitude smaller and more evenly dispersed. In their study Banerjee et al. [48] predicted cooling rates in their AM process range between 200-6000°Cs-1. Under very high cooling rates it was proposed that the boron, which has extremely low solubility in α and β Ti, becomes trapped in a supersaturated titanium phase and can only precipitate during thermal re-exposure as the next layer is deposited (simulating ageing). This leads to the formation of extremely fine nanosized TiB evenly dispersed throughout the component produced by AM [64]. Higher cooling rates also produced finer columnar grains (and smaller dendrite arm spacing) so eutectic TiB precipitation will be spread over a more uniform area, as opposed to a highly concentrated region between large grains. Mahbooba et al. [65] studied the effects of hypoeutectic boron during electron beam deposition and reported very small borides (about 1µm in size) that were uniformly dispersed within columnar grains. The phenomena observed here of large concentrated borides at columnar grain boundaries is a product of solidification conditions (low

cooling rate and growth mode, i.e. cellular vs dendritic) and is the subject of ongoing research. Reduced anisotropy and improved mechanical properties may be achievable in higher cooling rate AM processes that produce finer and more evenly dispersed TiB.

5.0 Conclusions

Trace additions of LaB6 and boron were introduced to Ti-6Al-4V during Wire Arc Additive Manufacturing and the response on the microstructure and tensile properties were examined. The key findings were:

• The alloying of lanthanum drastically changed the shape of the molten pool and subsequently the deposited bead shape. The bead aspect ratio (width:height) decreased from

approximately 4 in Ti6Al-4V and Ti-6Al-4V+B to around 1.7 when LaB6 was added, resulting in a poor surface finish. The presence of the rare earth is proposed to reverse the direction of Marangoni flow resulting in a taller but narrower bead and may be exploitable in the design of components with unsupported overhangs.

• LaB6 decomposes during additive layer deposition to form La2O3 and TiB. The La2O3 is evenly dispersed as small particles (<3µm) and likely forms prior to solidification of β-Ti. The TiB forms during the final stages of eutectic solidification which under the prevailing thermal gradients occurs between columnar grains resulting in highly orientated TiB needles aligned with the build direction. These needles vary in size but TiB >50µm long were observed.

• The tensile strength increased with additions of trace boron and LaB6 by about 10%, but this came at the expense of ductility. • The presence of directional TiB in the components had a marked effect on anisotropy. The

relative anisotropy increased dramatically with the addition of boron and LaB6 which was attributed to the highly aligned TiB needles. The transverse (vertical) build orientation was far

more ductile (εf ≈14-18%) than the longitudinal (horizontal) build orientation (εf ≈5-8%) in the presence of these aligned TiB. In the absence of TiB, the anisotropy in the ductility of Ti-6A-4V

reduced (εf ≈24% in transverse orientation, εf ≈20% in longitudinal orientation). In Ti-6Al-4V cracking was observed along grain boundary-α. • The aligned TiB are susceptible to cracking under tensile stress and the large size of the TiB produced by slow cooling during solidification (measured to be 90-100°Cs-1) creates opportunities for large cracks to open. It is anticipated that an improvement in the properties would occur if the size of the TiB reduced and were more uniformly dispersed. This may be possible with higher cooling rates during AM solidification.

Acknowledgements

The authors would like to acknowledge to support of the School of Mechanical and Mining Engineering and the Queensland Centre for Advanced Materials processing and Manufacturing. MJB acknowledges the support of the Australian Research Council Discovery Program and is in receipt of Discover Early Career Researcher Award (DE160100260). All authors acknowledge the support of the Australian Research Council Research Hub for Advanced Manufacturing of Medical Devices (IH150100024). MSD also acknowledges the Australian Research Council Research Hub to Transform Transforming Australia’s Manufacturing Industry through High Value Additive Manufacturing (IH130100008).

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