Promoting the columnar to equiaxed transition and grain refinement of titanium alloys during additive manufacturing

Acta Materialia 168 (2019) 261e274

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Promoting the columnar to equiaxed transition and grain refinement of titanium alloys during additive manufacturing M.J. Bermingham a, *, D.H. StJohn a, b, J. Krynen a, S. Tedman-Jones a, M.S. Dargusch a, b a

Centre for Advanced Materials Processing and Manufacturing, School of Mechanical and Mining Engineering, The University of Queensland, St Lucia, Queensland, 4072, Australia b Defence Materials Technology Centre, VIC, 3122, Australia

a r t i c l e i n f o

a b s t r a c t

Article history: Received 23 August 2018 Received in revised form 10 February 2019 Accepted 12 February 2019 Available online 16 February 2019

Preventing columnar grain formation during additive manufacturing has become a significant challenge. Columnar grains are generally regarded as unfavourable as their presence can impart solidification defects and mechanical property anisotropy, however, the thermal conditions experienced during additive manufacturing make columnar grains difficult to avoid. In this work the thermal conditions during solidification (cooling rate, temperature gradients) are characterised during wire based additive manufacturing. For the selection of deposition conditions that favour equiaxed grain formation, the role of alloy constitution is explored in three classical alloy design regimes: an alloy containing no grain refiners (Tie6Ale4V); an alloy only containing grain refining solutes (Tie3Ale8Ve6Cre4Moe4Zr); and an alloy containing both grain refining solute and nucleant particles (Tie3Ale8Ve6Cre4Moe4Zr þ La2O3). Substantial refinement and equiaxed grain formation is achieved in the latter case which is attributed to b-Ti nucleation on La2O3. However, the thermal environment is dynamic during additive manufacturing and equiaxed grain formation is only achievable when temperature gradients decrease sufficiently to permit constitutional supercooling. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Deposition Grain refinement Solidification microstructure Titanium alloys Columnar-to-equiaxed transition

1. Introduction The great deal of research into metal Additive Manufacturing (AM) technologies over recent years has identified a clear need to prevent the formation of columnar grains that are widely reported when producing titanium and other metal alloy components. The presence of textured grain structures causes mechanical property anisotropy, which is often perceived to be undesirable and which presents challenges in component design, build strategy and certification of products. Producing homogenous microstructures with fine grained equiaxed structures during AM is often highly desirable. The ability to achieve equiaxed grains during AM is dependent on the factors driving nucleation and growth such as the thermal conditions and alloy constitution. Columnar grains form during AM on account of the prevailing solidification conditions that favour epitaxial growth (from prior deposited layers) and a lack of nucleation events ahead of the solid/ liquid (S/L) interface. One approach to promote the columnar to

* Corresponding author. E-mail address: [email protected] (M.J. Bermingham). https://doi.org/10.1016/j.actamat.2019.02.020 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

equiaxed transition (CET) is to manipulate the processing parameters during AM to affect the temperature gradient (G), growth rate of the S/L interface (R) and cooling rate (product of G and R). In general, it is understood that the CET is favoured when R increases and G decreases [1]. Kobryn and Semiatin [2] studied the grain morphology of Tie6Ale4V under a range of solidification conditions from directionally solidified castings through to laser clad deposition and established a G/R solidification map for the alloy. Using this map, Bontha and co-workers [3] modelled the solidification conditions for laser AM of Tie6Ale4V and predicted that increasing the laser power (with other parameters constant) could theoretically produce equiaxed grain morphologies during deposition by reducing G, albeit acknowledging that R is also likely to reduce on account of lower cooling rates. Although there has been limited experimental evidence of forming equiaxed grains during Tie6Ale4V laser AM methods, Xu et al. [4] reported the presence of equiaxed grains in Selective Laser Melting (SLM) of Tie6Ale4V after carefully manipulating the build parameters to affect a reduction in the G. Wu et al. [5] also reported some large equiaxed grains could be formed under high laser power during deposition of Tie6Ale4V which was attributed to a reduction in G. Despite isolated cases, the

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majority of literature relating to AM of Tie6Ale4V does not report the CET. While G and R are undeniably important parameters influencing the CET, other factors such as alloy constitution (which favours nucleation) are also important in achieving equiaxed grain morphologies [1]. During laser metal deposition of Tie6Ale2Sne2Zre3Mo-1.5Cre2Nb, Zhang et al. [6] reported columnar grains at higher laser energy densities and equiaxed grains at lower energy densities. On first appearance this is at odds with the prediction by Bontha et al. [3] (which expects the CET at higher laser energy density), however, an important distinction is that the effect of powder deposition rate or changing nucleation factors (such as nucleation undercooling DTN, nuclei population etc.) were not considered in the model presented in Ref. [3]. As a result Zhang and co-workers [6,7] demonstrated that increasing the powder deposition rate increased the tendency for the CET by providing heterogeneous nucleation on partially unmelted powder particles. Conversely, it was found that increasing the laser energy density reduced the survival rate of unmelted powders and therefore reduced the nucleation rate. Wang et al. [8] also reported this same observation during laser metal deposition of Ti6.5Ale3.5Moe1.5Zre0.3Si as did Wu et al. [9] who found columnar grains when the energy density increased during laser deposition of Tie25Ve15Cre2Ale0.2C (but otherwise reported equiaxed grains at lower laser energy densities). It is worth noting that of the handful of cases where the CET has been achieved during AM, it mainly occurs in titanium alloys other than Tie6Ale4V. For the CET to occur it is essential for supercooled liquid to exist ahead of the columnar front in order for a large proportion of equiaxed grains to nucleate or for detached solid fragments or unmelted powders to survive and grow. It is well known that alloy solute plays an important role in generating constitutional supercooling (DTCS). The rate at which a solute generates DTCS is determined by the growth restriction factor, Q, whereby solutes with large Q values rapidly develop DTCS and are considered as growth restricting solutes that can provide effective grain refinement [10]. The Al and V solute in Tie6Ale4V provides no DTCS (both Al and V solute have negligible Q values in Ti [11]) whereas the alloys mentioned above contain solutes that do (e.g. Cr [12] and C [13]). This means that the CET is far harder to achieve in Tie6Ale4V during AM. This point was clearly demonstrated by Wang et al. [14] when fabricating a compositionally graded Tie6Ale4V to Tie25Ve15Cre2Ale0.2C component by laser metal deposition using a combined Tie6Ale4V wire feed source and Tie25Ve15Cre2Ale0.2C powder. Initially, the 100% Tie6Ale4V end of the component produced columnar grains but equiaxed grains resulted as soon as composition grading began (i.e. resulting in the Tie6Ale4V alloying with Cr and C). The purpose of this study is to investigate the role of alloy constitution on the CET during additive manufacturing of titanium alloys. Existing literature provides evidence that it is possible to achieve equiaxed grains and the purpose here is to understand how alloy chemistry can be manipulated to achieve the CET during AM. First the thermal conditions during AM that favour equiaxed grain formation will be identified (through melt pool characterisation) and then alloy design incorporating growth restricting solute and nucleant grain refiners, is used to achieve substantial refinement. 2. Alloy selection The base alloy selected for this study is Tie6Ale4V which contains a total of 10 wt% Al and V solute that does not provide growth restriction. Previous studies have shown that cast Tie6Ale4V has similar grain size to cast ASTM Grade 2 commercially pure titanium [11]. Oxygen (QO z 10.8C0 [15]) and iron (QFe z 3.8C0 [15]), present

in the alloy as trace elements, are known to be effective grain refiners in Ti castings [12,16] but their small concentrations in Tie6Ale4V make negligible contributions to grain refinement. For this reason a second commercial alloy containing sufficient growth restricting solute was sourced. Tie3Ale8Ve6Cre4Moe4Zr (commonly known as Beta C and ASTM Grade 19) is a metastable btitanium alloy used in aerospace applications and it contains growth restricting solute that should provide the necessary DTCS to activate available nucleant particles. Although containing a total of 25 wt% solute, it is the 6 wt% Cr (QCr z 1.5C0 [15]) in particular that should provide DTCS (similar to Al and V, Zr has a negligible Q value in titanium [11]). The present authors have confirmed that substantial grain refinement occurs in titanium castings with addition of 6 wt% chromium solute [12]. Mo has a partition coefficient greater than unity in the Ti system (k ¼ 1.41 [17]) and Samsonov et al. [18] demonstrated that it offers negligible grain refinement in titanium alloys. Even with 10 wt% Mo, Vrancken et al. [17] reported only minor columnar grain narrowing in a b-Ti alloy produced by SLM. To date no stable and potent1 nucleant particles have been conclusively identified for the b-phase in low concentrations. In the 1970s Crossley [19] achieved grain refinement with nitrogen, carbon and oxygen through the peritectic reaction. Qiu et al. [20] recently confirmed that TiN particles can heterogeneously nucleate b-Ti but require a nitrogen concentration of at least 0.4 wt% to achieve in situ particles. Such a high interstitial content is expected to cause embrittlement. Although stable nucleant particles for b-Ti at low concentrations are yet to be identified, it is understood that titanium alloys contain a naturally occurring population of ‘native’ particles that under favourable conditions are able to nucleate equiaxed grains during solidification (demonstrable by the fact that equiaxed grains are frequently observed in castings in a range of alloys). While unknown in origin, it is reasonable to expect that the native particles in titanium are not particularly potent because there are numerous examples in the literature where equiaxed grains fail to form in casting, welding and AM. Only when growth restricting solutes are introduced does the nucleation rate increase [21]. Although potent foreign nucleant particles for b-Ti are yet to be conclusively identified, there are encouraging reports that several rare earth elements refine the grain size of titanium alloys and could be suitable nucleant catalysts. Yttrium, lanthanum and other rare earth elements are reported to refine the central equiaxed grains in the fusion zone of titanium welds [22e24] and castings [18,25e27]. Although some rare earths theoretically should be effective grain refining solutes (for example, QY ¼ 7.93 C0 and QLa ¼ 3.3 C0 [15]), in reality if small levels of metallic rare earth are alloyed with liquid titanium they should scavenge oxygen from the melt and form oxide particles just as they do in steel melts [28], on account that many rare earth oxides have lower free energy than titanium oxides. Thus metallic rare earth solute should not be available to generate DTCS unless there is a surplus of solute compared to the available oxygen present in the system, in which case not all solute is consumed by oxidation. It was first proposed by Simpson [22] that yttrium scavenges oxygen to form Y2O3 which heterogeneously nucleates titanium, although it has not been supported by definitive evidence and alternative mechanisms for refinement have also been proposed [24]. Nevertheless, the presence of a thermodynamically stable particle in liquid Ti is of interest given the fact that very few compounds are stable in liquid titanium at very small concentrations. In this work we investigate La2O3 as one such potential particle.

1

‘potent’ particles are described here as those having small DTN.

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Several researchers have reported that La2O3 forms by scavenging oxygen if La is added to titanium alloys [29e31]. Recent work by the present authors [32] suggests that La2O3 exists in the liquid Tie6Ale4V prior to solidification in WAAM and small particles several micrometres in diameter were dispersed throughout the solidified microstructure. According to the Free-Growth model for grain refinement [33], the presence of such particles may enhance nucleation under the appropriate thermal environment and in alloys that produce DTCS. To summarise, the grain size and morphology during WAAM is explored in three alloy systems: 1) Tie6Ale4V e a commercial low Q alloy containing no growth restricting solutes or added nucleant particles. 2) Tie3Ale8Ve6Cre4Moe4Zr e a commercial high Q alloy containing growth restricting solute and no added nucleant particles. 3) Tie3Ale8Ve6Cre4Moe4Zr þ La2O3 e a commercial high Q alloy containing growth restricting solute plus added nucleant particles (La2O3). To support the analysis of the macrostructures obtained quantification of the thermal environment, in particular the temperature gradient, is undertaken. 3. Experimental methods The cooling rates during AM are high and collecting experimental data is challenging, particularly in very small melt pools. In this work, the wire þ arc additive manufacturing process (WAAM) is explored because the melt pool is large enough for the cooling rate to be reliably measured using pyrometers, from which the solidification parameters G and R can be measured or calculated. The WAAM equipment used for this research consists of a EWM Tetrix 350 Gas Tungsten Arc Welding torch (GTAW) equipped with a wire feed unit installed into a customised build chamber with a 3axis CNC controlled table. The build chamber is exposed to the open atmosphere and localised shielding (Argon 99.999%) is employed to protect the melt from oxidation. Shielding is provided through the water cooled GTAW nozzle, and where necessary, a trailing shield is also installed behind the GTAW nozzle which provides shielding to the hot metal as it cools. Details of the trailing shield and other equipment are available elsewhere [34]. This research first establishes the thermal environment during WAAM, then moves to explore the role of alloy constitution on the columnar to equiaxed transition. The experimental method for each activity is described in the following sections. 3.1. Establishing the thermal environment in WAAM The seminal empirical work by Kobryn and Semiatin [2] in developing a solidification map for Tie6Ale4V has been universally referenced to predict the solidification mode (i.e. columnar or equiaxed) of Tie6Ale4V. This map will be used in this work to determine if the inherent solidification conditions encountered during WAAM promote dendritic equiaxed solidification. For a given alloy the most important solidification parameters determining the extent of constitutional supercooling, and therefore the growth mode (planar, cellular, dendritic) is the temperature gradient G and the growth rate R. The temperature gradient is particularly important since it determines whether or not a dendritic protuberance develops and can advance into a supercooled liquid. Furthermore, the steep temperature gradients characterised by AM reduce or eliminate constitutional supercooling which can prevent nucleation events ahead of the main interface and favour

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cellular or dendritic columnar grain morphologies. Accurately measuring G and R during AM is challenging. While the growth rate R can be approximated from the travel/scan speed or a simple trigonometrical relationship therein [24], measuring the temperature gradient can only be achieved experimentally. Traditional measurement via thermocouples is challenging because the Tie6Ale4V build layer must be deposited over the thermocouples and the heat source and liquid metal must not melt or dissolve the thermocouples. Non-contact infrared pyrometers overcome these problems but only the surface temperature can be accurately measured, thus are limited in that they may not be representative of changing thermal conditions below the surface. To overcome these challenges a hybrid process was used to approximate the average thermal gradient across the melt pool using surface temperature data and post-mortem microstructural analysis. This approach firstly involves measuring the temperature of the liquid at the top surface during deposition using a pyrometer. In this work a pyrometer with effective range of 800e2100  C was used (Micro-Epsilon, spectral range 1 mm, ±7  C accuracy at 1680  C, 1 ms response time), with emissivity of 0.288 calibrated against high purity titanium using the same technique outlined elsewhere [32]. After deposition, the layer is cross-sectioned and polished to reveal the extent of melting in order to measure the depth of the layer. The cross-section naturally consists of a melted region and an unmelted region from the previous layer or the base plate if it is the first layer. The average thermal gradient across the bulk liquid pool can then be determined given that the temperature of this interface during deposition equals the freezing temperature of the alloy and the distance between it and the top surface (where the temperature is known) is measurable. However, the effectiveness of this method requires distinguishing between the S/L interface corresponding to the melted region and the unmelted prior layer or baseplate, which becomes almost impossible if the same alloy is used. To overcome this problem, single Tie3Ale8Ve6Cre4Moe4Zr layers were deposited onto Tie6Ale4V substrates which allowed a very distinctive boundary to be observed when examining the metallurgical cross-section, an example is given in Fig. 1 (B). This technique was used on both flat Tie6Ale4V substrates (representing layer 1) as well as previously deposited Tie6Ale4V layers (at least 3 layers high) which are curved. Similar results were obtained irrespective of the shape of the substrate (curved vs flat). The deposition parameters investigated are shown in Table 1. The purpose of this activity was not to map all possible deposition parameter combinations, but only focus on a select range of deposition speeds that are frequently reported in the literature for wire based AM processes. The other parameters such as wire feed rate, arc power etc. have been used in other research to produce quality components with minimum oxidation [34,35]. It was not possible to use a trailing shield in this initial stage of experiments because it obstructed temperature measurements. While the absence of the trailing shield may slightly affect the solid state cooling rate (through convective cooling associated with the argon gas), it is not expected to significantly influence the cooling rate of the melt pool (or the measured G) because this is protected by argon from the GTAW nozzle and not by the trailing shield. Our previous work in Tie6Ale4V WAAM with similar build parameters to those studied here, found that complete solidification occurs while still protected by the cover gas flowing from the GTAW nozzle [34]. 3.2. Experimental methods investigating the CET As previously mentioned, the three alloys studied in this work are: Tie6Ale4V, Tie3Ale8Ve6Cre4Moe4Zr, and Tie3Ale8Ve

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M.J. Bermingham et al. / Acta Materialia 168 (2019) 261e274 Table 1 Deposition parameters used to build the components. Deposition Parameters Base Current: Wire feed: Wire: Deposition speeds: Electrode-substrate gap: Substrate: Electrode: Argon:

113Amp 1.5 m/min Tie6Ale4V, Tie3Ale8Ve6Cre4Moe4Zr; both ø ¼ 1.0 mm 50 mm/min, 100 mm/min, 200 mm/min, 400 mm/min 5 mm Tie6Ale4V ø ¼ 2.4 mm tungsten-rare earth 99.999% purity, 20 L/min

paint containing La2O3 powder (with approximately 50% by weight La2O3) was painted onto the surface before deposition, in the same way as described in earlier work [32]. The La2O3 powder selected for study contained a distribution of large particles (D50 4.362 mm; D90 26.0 mm, Alfa Aesar, 99.9%). Four layer high deposits were fabricated using the parameters identified in Table 1 with the most appropriate travel speed identified in Section 3.1 (50 mm/min). The chemical composition of the alloys after deposition is given in Table 2. Each deposit was approximately 130 mm long. A trailing shield was used to minimise contamination during deposition (details of this shield are available in Ref. [34]). An unavoidable consequence of using the trailing shield is that it makes thermal measurements during deposition impossible (since it encloses the hot metal), however, it was still possible to measure the stationary melt pool at the end of deposition since this was not obstructed by the trailing shield. After each layer the component was allowed to cool to room temperature before the next layer was deposited. A Tie6Ale4V baseplate was used for deposition. After deposition, several cross sectional samples were machined from the components and prepared using conventional metallographic methods (grinding on SiC paper followed by polishing with 0.05 mm colloidal Silica) followed by etching with Kroll's reagent (5% HF, 30% HNO3, 65% H2O). The samples were characterised using optical and electron microscopy techniques and the grain size was measured using the linear intercept method [36].

4. Results 4.1. Thermal conditions during WAAM

Fig. 1. (A) Diagram and photograph showing the experimental set up. A pyrometer (fixed in relative position to the work piece) measures the real time cooling rate as the torch passes by (depositing along the Y-axis). A video camera (synchronised with the pyrometer) observes the process in the Y-Z plane. (B) Diagram with example optical and Backscatter SEM images at deposition speed 400 mm/min showing how the average temperature gradient G was determined. Similar results were obtained when depositing on curved surfaces (i.e. on a pre-existing layer) and flat substrates (i.e. first layer). T1 is measured by the pyrometer and T2 equals the melting point of Tie6Ale4V. EDS maps show minimal mixing between the Tie3Ale8Ve6Cre4Moe4Zr and Tie6Ale4V layers making the interface easily discernible.

6Cre4Moe4Zr þ La2O3. Both Tie6Ale4V and Tie3Ale8Ve6Cre4Moe4Zr were commercially sourced wires having 1 mm diameter. Where it was necessary to add La2O3 particles, an alcohol based

The thermal conditions measured before, during and after solidification are provided in Fig. 2 (A) for four deposition speeds. For each deposition speed the liquid is superheated in excess of 2100  C and the temperature rapidly fluctuates in the liquid under the electric arc most likely due to rapid fluid flow generated by Marangoni forces. Bai et al. [37] modelled the thermal environment within the melt pool during plasma arc additive manufacturing of Tie6Ale4V under comparable deposition conditions to that here. The authors predicted that liquid temperatures under the arc can exceed 2400  C and steep temperature gradients within the liquid produce strong and complex Marangoni currents causing rapid mixing. The cooling rates depend strongly on the travel speed and from multiple tests the average cooling rates during solidification ranged from 101  C/s at 50 mm/min deposition speed to 642  C/s at 400 mm/min. After solidification, a significant exothermic reaction (latent heat release) is observed between 1400 and 1500  C and corresponds to oxide formation in Fig. 2 (B). This occurred because it was not possible to use a trailing shield to provide localised inert gas shielding to the hot solid metal which would otherwise prevent

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Table 2 Chemical analysis of each alloy determined by ICP-AES and Leco Combustion. Note that the accuracy range of the O, N, Fe and La detection is approximately ±0.005e0.016 wt% and for all other elements the accuracy in the range of ±0.06e0.08 wt% (these ranges represent 1 standard deviation). Alloy

Tie6Ale4V Tie3Ale8Ve6Cre4Moe4Zr Tie3Ale8Ve6Cre4Moe4Zr þ La2O3

O

N

Al

V

Fe

Mo

Cr

Zr

La

wt%

wt%

wt%

wt%

wt%

wt%

wt%

wt%

wt%

0.08 0.11 0.15

0.012 0.032 0.035

5.91 3.85 3.99

3.74 7.51 7.6

0.11 0.18 0.12

e 3.74 3.83

e 5.55 5.58

e 3.56 3.6

e e 0.31

thermal measurements. Finally, Fig. 2 (C) shows the solidification map for Tie6Ale4V at the four different WAAM deposition speeds. The growth rate R is approximated by the deposition speed and the average thermal gradients presented are calculated from the bottom of the melt pool to the external surface being measured by the pyrometer. The maximum temperature of the liquid exceeds the pyrometer's effective range (i.e. 2100  C) and as such an upper and lower range of G values are shown based on minimum surface temperature of 2100  C to an upper temperature of 2400  C reported by Bai et al. [37]. For comparison, the typical operating range for LENS (blown powder-laser AM) is provided based on the work of Bontha et al. [3]. The solidification map predicts a mixed equiaxed þ columnar grain morphology during deposition at each travel speed. In light of this, deposition at a speed of 50 mm/min was selected for the next experimental section studying the composition effects on the alloys because this speed had the lowest G of the conditions explored. Fig. 3 shows the cooling curves for each individual alloy at the end of deposition when the arc was terminated. Without interference from the electric arc, this shows sensitive thermal events such as recalescence and changing cooling rates associated with this latent heat release. Initially the liquid in all alloys cools very rapidly from above 1700  C (exceeding 5000  C/s) but dramatically slows and eventually plateaus during solidification. This happens when the rate of latent heat release balances the rate of heat extraction, and occurs over a duration of about 1e2 s. The Tie6Ale4V and Tie3Ale8Ve6Cre4Moe4Zr alloys also recalesce by 11e16  C, however, the addition of La2O3 eliminates this phenomena.

4.2. Promoting the CET Fig. 4 shows the microstructures and measured grain size for Tie6Ale4V, Tie3Ale8Ve6Cre4Moe4Zr and Tie3Ale8Ve6Cre4Moe4Zr þ La2O3 deposited under thermal conditions that should be favourable to a mixed equiaxedcolumnar microstructure. The Tie6Ale4V build contains very large grains which are mostly columnar but with some isolated equiaxed grains which is consistent with the prediction from the solidification map. The introduction of solute (Tie3Ale8Ve6Cre4Moe4Zr) substantially refines the grain size including the width of the columnar grains, however, fails to prevent columnar grain formation. This is a similar trend to that observed by Mereddy et al. with the introduction of Si solute to CPeTi [38] or carbon solute in Tie6Ale4V [13]. Although this is a significant refinement, the greatest change in the microstructure occurred with the introduction of La2O3 particles to the Tie3Ale8Ve6Cre4Moe4Zr which produced substantial refinement while increasing the proportion of equiaxed grains within the microstructure. Close examination of the microstructure reveals spherical La2O3 particles distributed throughout. Fig. 5 shows the size and shape of the La2O3 before and after deposition, as well as the measured size distributions of La2O3 within the Tie3Ale8Ve6Cre4Moe4Zr microstructure.

Fig. 2. (A) Example cooling rates for Tie6Ale4V during WAAM at deposition speeds between 50 and 400 mm/min (B) Shows photographs at three time intervals (I, II, III) during deposition at 100 mm/min as an example. It is clear that temperatures near the arc exceed 2100  C and considerable thermal fluctuations occur within the melt pool. After solidification and once the torch nozzle moves past there is no longer inert gas shielding and oxidation occurs which releases latent heat at time III. (C) Tie6Ale4V solidification map for WAAM (based on the original map by Kobryn and Semiatin [2]). A comparison to the laser þ powder AM process is also provided based on the work and map presented by Bontha et al. [3] in modelling a range of typical laser powers and scanning speeds encountered during LENS that yield energy densities between 40 and 260 J/mm.

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Fig. 3. Example of typical cooling curves collected at the end of deposition during solidification of Tie6Ale4V, Tie3Ale8Ve6Cre4Moe4Zr and Tie3Ale8Ve6Cre4Moe4Zr þ La2O3. Both Tie6Ale4V and Tie3Ale8Ve6Cre4Moe4Zr recalesce by 11e16  C but this is eliminated with the addition of La2O3. The cooling rate (DT/Dt) is averaged over a 20 ms period.

5. Discussion In this work we determined the approximate thermal conditions (G, R and cooling rate) during WAAM of Tie6Ale4V and confirm the accuracy of the solidification map [2] in predicting a mixed columnar þ equiaxed growth morphology for the conditions studied here. Compared to other directed energy deposition processes such as LENS, the formation of some equiaxed grains is possible in Tie6Ale4V WAAM due to the lower thermal gradients. As a general rule, WAAM will have a higher linear energy density (also known as heat input) than LENS and this creates larger melt pools, slower cooling rates and lower thermal gradients at

comparable growth velocities (hence smaller G/R ratio). The linear energy density (LED) for WAAM in this study was 1.49 kJ/mm at a deposition speed of 50 mm/min (compared to 0.06e0.26 kJ/mm for LENS conditions shown in Fig. 2). While this is favourable in terms of lower average thermal gradients that permit the formation of some equiaxed grains, the undesirable consequence is a large grain size. The average grain size in Tie6Ale4V determined using the linear intercept method exceeded 1.5 mm, which is three times larger than the equiaxed grain size in Tie6Ale4V castings produced with similar cooling rates [11]. However, some individual columnar grains can be much larger than the average value as evident in Fig. 4. LENS, with its lower LED and higher cooling rates typically forms much thinner columnar grains. For example, Carroll et al. [39] reported an average columnar width of 0.375 mm during LENS of Tie6Ale4V with a LED of 0.189 kJ/mm. The challenge, therefore, is to not only achieve equiaxed grain formation (through a reduction in G), but also to refine the grain size at the same time. This is achievable with grain refiners. In the present example for WAAM, the growth restricting solute (especially Cr) in Tie3Ale8Ve6Cre4Moe4Zr generates a DTCS zone ahead of the S/ L interface and equiaxed grain initiation will occur when DTCS equals or exceeds the potency of the nucleant particles (DTN) present in this zone. However, for the CET to occur there must be suitable particles present within the supercooled zone. Titanium alloys naturally contain a large population of uncharacterised, semi-potent nucleant particles [21]. From Fig. 4 (B) it is clear that the introduction of growth restricting solute has not promoted a complete CET but instead adopts a mixed mostly columnar with isolated equiaxed grain structure similar to Tie6Ale4V. However, the grain size, including the width of the columnar grains, is less than half of that measured in the Tie6Ale4V build. The mechanism for this columnar refinement has been previously discussed for boron [40] and silicon [38] additions in WAAM but in summary manifests by lateral solute rejection and accumulation between columnar grains and the associated growth restriction. Although some nucleation events occur ahead of the columnar front which generate some equiaxed grains, there are also many cases where this does not occur and instead very long and narrow columnar grains result. Achieving the CET during WAAM requires the addition of extra nucleating particles that are more potent or more numerous than the native population found in titanium alloys. The introduction of La2O3 into Tie3Ale8Ve6Cre4Moe4Zr further refines the grain size while also increasing the fraction of equiaxed grains, however, does not eliminate all columnar grains. The average grain size of Tie3Ale8Ve6Cre4Moe4Zr reduced from 685 ± 222 mm to 252 ± 12 mm with La2O3 addition and was much more consistent in comparison to the grain size of Tie6Ale4V and Tie3Ale8Ve6Cre4Moe4Zr which had a large standard deviation due to the presence of both small and very large columnar grains. SEM examination reveals that the microstructure contains an abundance of La2O3 particles and it is probable that some of these have nucleated new equiaxed grains. 5.1. Refinement by La2O3 The presence of thermodynamically stable particles within the melt is the first key requirement for heterogeneous nucleation. Given the reactivity of liquid titanium, identifying a stable solid particle has traditionally been problematic as most high melting point compounds dissolve in liquid titanium. However, rare earth metal oxides (including La2O3) have lower free energies than titanium dioxide and should be stable in titanium melts. Metallic erbium and yttrium additions to titanium have long been known to

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Fig. 4. Microstructures of as-built Tie6Ale4V (A), Tie3Ale6Cre4Moe4Zr (B) and Tie3Ale6Cre4Moe4Zr þ La2O3 (C). (D) Shows the average grain size for each alloy (error bars ±1 standard deviation). Deposition speed 50 mm/min. Note that the grain boundaries have been enhanced in (A1), (B1) and (C1) for visibility. The large red box overlay marks the approximate locations for all SEM images (all from top layer) except for (C4) which shows the interface between the Tie6Ale4V substrate and the Tie3Ale6Cre4Moe4Zr þ La2O3 at the first layer and is approximately located by the small blue box in image (C1). The white ‘dots’ in images (C2), (C3) and (C4) are La2O3 particles which have strong contrast during the backscatter imaging mode. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

scavenge oxygen and form oxide particles during melting and solidification [41e44]. The same phenomena was recently observed with LaB6 which decomposes into La2O3 in Tie6Ale4V melts produced by WAAM [32]. Complete melting is not required and this oxygen scavenging reaction (to form La2O3) also occurs during solid state sintering [29]. Thus La2O3 belongs to a small group of compounds that are thermodynamically stable at titanium's freezing temperature and may be able to facilitate heterogeneous nucleation under appropriate thermal conditions. The traditional consideration for effective nucleation is that the crystallography of the nucleating phase should be close to that of the growing solid phase [45,46]. The edge-to-edge crystallographic model developed by Zhang et al. [47] builds upon early planematching theory with a focus on atom-row matching at the interface and has been effective in predicting potent nucleating particles in a number of systems including Mg [48] and Al [49]. Applied to the current investigation, the following orientation relationship between La2O3 and b-Ti was identified:




 1011

h i 1121

La2O3

La2O3

.. ð1 0 0ÞbTi

.. ½0 1 0bTi

Using published crystallographic data of La2O3 [50] and b-Ti [51] at elevated temperatures, the misfit of this orientation relationship was evaluated to be ~5%, with lattice parameters of ~3.17 Å and ~3.33 Å along the ½1 1 0 1La2O3 and ½0 1 0bTi directions respectively. Given that the lattice parameter of the b-Ti phase is likely smaller in Tie3Ale8Ve6Cre4Moe4Zr due to the substitutional content of Cr, Mo, V and Al, the actual lattice disregistry between La2O3 and b-Ti is expected to be less than ~5%. It is therefore plausible that La2O3 particles could act as favourable nucleation substrates for the b-Ti phase during solidification. While the orientation relationship between the nucleating particle and the new solid phase is important, perhaps an even more important consideration influencing the potency of a particle

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Fig. 5. Backscatter SEM images showing the size and morphology of the initial La2O3 powder (A) compared to the morphology of the La2O3 particles after deposition (B). The shape of the particles becomes more spherical after the deposition process and the size decreases. (C) La2O3 particle size distribution after deposition and (D) area fraction of La2O3 particles distributed within the Tie3Ale8Ve6Cre4Moe4Zr at different layers (box plots showing 25th,50th & 75th percentiles; whiskers showing 5e95th percentiles). Note that data for (C) and (D) was measured from 60 SEM images containing over 13,500 particles and calculated using ImageJ software assuming an average spherical shape.

is its size. Classical nucleation theory indicates that free growth of a grain can only occur if the nuclei embryo exceeds the critical size (r*) to overcome the interfacial energy barrier of formation. The Free Growth theory proposed by Greer et al. [33,52] predicts that the presence of large particles in the melt enables the liquid to wet a larger surface and form a thin hemispherical cap, effectively increasing its radius of curvature above r* and facilitating further growth. In contrast, free growth on smaller particles can only occur by growing the height of the hemispherical cap which effectively reduces the radius and in some cases this may fall below r*. Therefore, the potency of particles is inversely proportional to their size, with larger particles having smaller DTN. An analysis of particle size and potency of TiB2 nucleants for aluminium alloys found that particles above 3 mm had high potency (DTN z 0.2  C) and that it was the large particles that were responsible for most nucleation events [52]. Particles less than 1 mm are generally ineffective because DTN is large and this amount of supercooling is rarely achieved, especially once latent heat and recalesence reduces available supercooling (generated from the growing solid that may have nucleated on larger particles). The La2O3 powder used in this study had a distribution of large particles, with 50% exceeding 4.362 mm in diameter and 10% exceeded 26.0 mm in size. However, close examination of the as-built microstructures (see Fig. 5) revealed that the typical La2O3 particle size is much smaller.2 In addition, it is clear that the particles have changed shape during AM, which could indicate that the particles themselves melt/ decompose and reform or change shape through solid state

diffusion driven by a surface area reduction. The melting temperature of La2O3 is 2313  C [53] and Bai et al. [37] predict that this temperature is exceeded during WAAM of Tie6Ale4V. Weng et al. [54] also report that Y2O3, which has a melting temperature of over 2400  C, also decomposes into Y and O atoms in laser heated Tie6Ale4V melt pools before reforming as spherical Y2O3 particles. An undesirable consequence of La2O3 decomposition is that the particle size is reduced which may reduce its potency. Like other heterogeneous nucleant particles, it is expected that the potency of La2O3 will be linked to its size and that only the largest particles will be responsible for nucleation events, with many of the smaller particles lacking the potency to activate under the available thermal or constitutional supercooling. The above factors are all important considerations affecting nucleant potency but ultimately such particles can only seed new grains if they are present at the right location in time. The Interdependence model developed by StJohn et al. [55] predicts the grain size (dgs ) by considering these important interrelated factors including alloy chemistry (solute effects including concentration C0, solute diffusion rate D, solute growth restriction factor Q, solute partition coefficient k, and solute concentration at the interface C *l ); nucleant particle potency (DTN) and particle spacing (xSD ). Note that due to the inherent distribution of particles sizes, DTN referred to in this paper and in Equation (1) represents the undercooling required to activate the largest particles that have the lowest value of DTN. The Interdependence model is given by Equation (1) and is separated into its three key terms in Equation (2):

dgs 2

The particle size in the as-built microstructure is still quite large (up to ~3.5 mm). It is worth noting that particles sizes above ~0.5 mm might be considered ‘coarse’ and not helpful for dispersion strengthening and may even negatively influence some properties such as ductility.

C *  C0 D,z,DTn 4:6,D , *l þ ¼ R R,Q C l ð1  kÞ

dgs ¼ xCS þ xDL þ xSD

! þ xSD

(1)

(2)

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In essence, the Interdependence model predicts that the grain size is determined by the summation of three distinct length scales: xCS , xDL , and xSD . The first term, xCS , represents the distance (or size) that a grain must grow before it generates sufficient DTCS to activate nearby nucleant particles (i.e. where DTN  DTCS). zDTN is the incremental amount of undercooling required to achieve DTN  DTCS for subsequent nucleation of equiaxed grains. The second term, xDL , relates to the diffusion dependent distance of the solute profile (and therefore DTCS profile) from the S/L interface to the moment sufficient DTCS is achieved for nucleation. The final term, xSD , relates to the distance from xDL to the nearest nucleant particle where the condition DTN  DTCS is satisfied. Considering that particle potency is inversely related to its size, xSD is therefore influenced by both the number of particles present in the melt, their size distribution and the probability that sufficiently potent (large) particles will be nearby the interface. It is worth noting that to achieve the CET during casting, welding or AM, sufficient particles need to be present along the entire growth front otherwise only a few new equiaxed grains will grow where particles are available and areas devoid of suitable particles will solidify into columnar grains. 5.2. Challenges in achieving fully equiaxed microstructures during AM Recalling the microstructures in Fig. 4, the addition of both segregating solute and nucleant particles refine the columnar grains and promote the formation of equiaxed grains. Each new layer in Fig. 4 (C1) is a mixture of columnar grains and equiaxed grains. It appears that the columnar grains form first with the deposition of each layer followed by a region of equiaxed grains. The question remains, why do columnar grains still form? The Interdependence model provides two possibilities. The first is that suitable nucleant particles are not present at the start (bottom) of each layer (i.e. xSD is very large). However, this scenario is unlikely given the abundance and homogenous distribution of particles seen in Fig. 4 (C4) and Fig. 5 (D), which are much more closely spaced than the columnar length scale (which can be hundreds or even thousands of micrometres). The second and more likely scenario is that thermal conditions at the bottom of the layer are not yet ready for nucleation. It is important to recognise that the first nucleation event occurs epitaxially where DTN is effectively zero. Therefore, solidification of the columnar front begins instantaneously when the freezing temperature of the alloy is reached. Any foreign nucleant particle within the melt (including La2O3 in the present case) will require some amount of supercooling (DTN > 0  C). This can occur once constitutional supercooling develops ahead of the columnar front which in turn can only occur once the thermal gradient G is sufficiently low. Additive manufacturing processes are widely believed to form columnar microstructures on account of steep thermal gradients. Even for high energy density, large melt pool AM processes such as WAAM, the conservative estimates of average bulk thermal gradients across the melt pool is of the order of 100  C/mm (presented in Fig. 2), which is in stark contrast to that experienced in typical castings which are almost flat. Under such steep thermal gradients it is surprising that any equiaxed grains can form. Given the available evidence that equiaxed grains do indeed form, the most likely scenario is that G is not constant across the melt pool, but rather, dynamically reduces during different stages of solidification. Eventually a reducing thermal gradient will allow constitutional supercooling to develop. Multiple factors could cause G to decrease throughout solidification of a layer during AM. One important physical attribute of solidification is the generation of latent heat. The effects of latent heat release are evident on the solidification curves presented in

269

Fig. 3. The magnitude of latent heat release during solidification of each titanium alloy3 is so significant that it can temporarily match or exceed the rate of heat extraction. To put this into perspective, the cooling rate of the liquid prior to solidification exceeded 5000  C/s but during solidification the heat release was so intense that not only did the cooling rate decrease to zero, but latent heat actually caused a net heating effect in Tie6Ale4V and Tie3Ale8Ve6Cre4Moe4Zr whereby the actual temperature increased by 11e16  C (recalesence). Such significant latent heat will cause a lowering of G as solidification proceeds.3 Melt turbulence, and therefore thermal mixing will also cause G to reduce before or during solidification. It is well known that Marangoni forces are active during AM which will cause rapid fluid flow and mixing within the melt pool. This mixing will continue as long as the heat source is present but once the heat source passes the melt may thermally equilibrate before complete solidification. This could be similar to the situation in castings when the application of external ultrasonic fields causes rapid mixing and consequently thermal gradients flatten (approaching zero) prior to solidification [57]. Not much is known about the temperature distribution within molten titanium pools during AM, however, in Fig. 6 we present new data from recently published research simulating the WAAM process as well as new experimental data confirming that temperature gradients are changing throughout the molten pool. Fig. 6 (A) is a reproduction of work presented by Bai et al. [37] in modelling the heat transfer during WAAM of Tie6Ale4V using a plasma heat source. The figure shows the local temperatures within the melt pool which vary significantly and generate significant Marangoni convection. Using Bai's data, we have calculated the temperature gradient within the liquid at different positions along the melt pool and presented this in Fig. 6 (B). The temperature gradients are predicted to be very high initially (over 500  C/mm directly underneath the heat source) but rapidly flatten out due to efficient liquid mixing and decreases to approximately 50e60  C/ mm at the tail of the melt pool. In Fig. 6 (D) we present new experimental data collected using tandem pyrometers to monitor the melt pool at fixed distance apart (see Fig. 6 (C) for experimental schematic). This data confirms that the temperature gradient within the liquid at the rear of the melt pool is much smaller than previously expected (fluctuating between 0 and 50  C/mm). Interestingly, even under very high cooling rates approaching 105  C/s, Zhao et al. [58] showed through in situ synchrotron imaging that the growth rate (and by extension G) decreased during solidification of very small melt pools in SLM of Tie6Ale4V. Thus, although initial thermal conditions directly underneath the heat source may preclude nucleation ahead of the interface, a reduction in G eventually generates DTCS which makes nucleation of equiaxed grains possible. Given that the whole solidification process takes only a few seconds or fractions of a second with a very high cooling rate and steep initial temperature gradient, it is surprising that a model (Equation (2)) that was developed for quiescent, low G conditions and potent particles can still shed light on the mechanisms occurring. The application of the model to date has not specifically considered high G environments although we know that a relationship between grain size and the inverse of Q applies in high pressure die casting environments [59]. The most influential terms affecting both the CET and equiaxed grain size in Equation (2) are

3 Latent heat of fusion for titanium is approximately 14,550 ± 500 J mol1 and molten titanium has a specific heat of 46.29 ± 1.7 J mol1 [56], so as an example a mass of solidifying titanium will release heat equivalent to it cooling from about 2000  C to 1680  C.

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Fig. 6. Temperature distribution within Tie6Ale4V melt pools during WAAM. (A) and (B) are based on the analysis of Bai et al. [37] who developed a model for heat transfer during plasma arc additive manufacturing of Tie6Ale4V. (A) is a reproduction from the original research and shows a temperature profile within the melt pool. Note that the arc position is at approximately Y ¼ 0 mm on this figure. Using this map the present authors have calculated local temperature gradients along the melt pool, presented in (B). (D) Shows new experimental data for the dynamic temperature gradient between tandem pyrometers positioned a fixed distance apart (experimental schematic shown in (C). The result shows that the temperature gradient fluctuates within the liquid at the back of the melt pool (T > 1700  C) and is less than 50  C/mm.

xCS and xSD . xCS is the amount of growth required to achieve sufficient DTCS for nucleation. In the case of epitaxial nucleation (DTN ¼ 0) DTCS equal to DTN of a suitable nucleant particle, needs to be generated which is given by the first term in Eq. (2) when z ¼ 1. This situation corresponds to the equation developed by Qian et al. [60] for the case of nucleation without considering the effect of casting conditions or the availability of nucleant particles. Simultaneously, latent heat is released during growth to xCS which lowers the temperature gradient. Once nucleation on the first available potent nucleant particle occurs, xCS for subsequent nucleation is much smaller (i.e. z is a fraction of 1) promoting equiaxed nucleation. The size of the initial xCS after epitaxial nucleation is important as this represents the amount of unimpeded growth of the columnar grains before CET becomes possible. In alloys containing high Q solutes, (e.g. Tie3Ale8Ve6Cre4Moe4Zr) a concentration gradient may develop quickly ahead of the interface but no DTCS occurs until G sufficiently reduces. However, in alloys absent of high Q solutes (e.g. Tie6Ale4V), no DTCS occurs due to the lack of segregating solute, regardless of whether G is low or even completely flat. Thus for the CET to occur a concurrent requirement is an abundance of potent particles (small DTN) which will determine xSD . The addition of La2O3 particles to Tie3Ale8Ve6Cre4Moe4Zr appears to substantially reduce xSD , although it remains unknown whether this occurs because the particles have a smaller DTN than the native nucleants, or have equal potency and are simply more numerous and therefore closely packed, or even a combination of both. The microstructure reveals that La2O3 particles are abundant and quite closely packed even in areas at the base of the layers where columnar grains form so it is clear that xCS is the dominant term here. This highlights a challenging reality for some AM processes with steep initial thermal

gradients in that xCS may exceed the depth of the melt pool and thus make the CET impossible in alloys not containing potent nucleant particles and sufficiently high Q solute. Fig. 7 illustrates the principles of the Interdependence model applied during AM solidification at different time intervals. As previously mentioned, the first growth occurs epitaxially at DTN ¼ 0  C under a steep thermal gradient (t1) where no DTCS or solute diffusion into the liquid has occurred. After a small amount of epitaxial growth (t2) solute rejection into the melt will begin generating a constitutional profile ahead of the growing cellular/ columnar grains, however, G is still too steep to permit DTCS and hence nearby particles (with DTN > 0  C) cannot activate. With further growth of the columnar front, G continues to decrease and after time t3 a small constitutionally supercooled zone exists, however, this DTCS may be insufficient to activate nearby particles (DTCS < DTN). With further growth eventually G decreases sufficiently where DTCS  DTN and thus nucleation can occur if a suitable particle is nearby, however, at t4 no such particle is nearby and thus columnar growth continues by distance xSD1 until the next particle is reached and an equiaxed grain can form (t5). Once nucleation on the first available potent nucleant particle occurs, the first term of Eq. (1) is then relevant where much less growth (i.e. z is a fraction of 1) is required to trigger subsequent equiaxed nucleation events (t6). From this point on, the temperature gradient is low enough and sufficient DTCS exists that new nucleation events are mostly determined by particle spacing xSD . Fig. 8 shows the relationship between the measured grain size and the inverse of the growth restriction factor (1/Q) for the alloys studied here and compared with the WAAM data presented by Mereddy et al. [38] (TieSi alloys) and previous data reported by Bermingham et al. [12,15,61,62] in casting for a range of TieBe,

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Fig. 7. Diagram illustrating the key principles of the Interdependence model applied to AM at six time intervals, t1-t6 (refer to text for a detailed description of the event sequence).

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Fig. 8. Grain size plotted against the inverse Q for a range of titanium alloys produced by casting and WAAM. The presence of La2O3 to Tie3Ale8Ve6Cre4Mo-4Z reduced the grain size beyond levels achievable with solute effects alone (more nucleation events) and as such was excluded from the line of best fit calculation for the WAAM series.

TieCr, TieFe, TieB and TieSi alloys. Presenting grain size data this way is a useful way to understand the relationship between the nucleant and solute parameters of the Interdependence model. A considerable amount of empirical evidence has been generated over recent decades suggesting a linear relationship occurs in alloys where only solutes are changed (Q term) in titanium, aluminium and magnesium systems [55,63,64]. Assuming sufficiently low G where constitutional supercooling is achieved, the Interdependence model predicts that only changing solute terms will decrease the grain size to a minimum theoretical value where the xCS and xDL terms approach zero (1/Q / zero). This means that the grain size is controlled by the spacing of the particles (xSD ) which corresponds to the y-intercept on the grain size vs 1/Q plot. Further refinement from here is only possible by reducing xSD (increasing the number of activated particles so that they are closer together and smaller grains result). A clear trend emerges in that titanium alloys produced by WAAM (containing Si and Cr as the most powerful growth restricting solute) exhibit much larger grain size than the same alloys cast into copper moulds. The best fit lines through the respective casting and WAAM data sets are similarly sloped, which from the Interdependence theory, suggests that the potency of the native nucleant population is similar between these data sets. The main difference, however, is that the y-intercept (xSD ) is much larger for WAAM than for the cast alloys, although it is recognised that the growth required to reduce G within xCS may also account for this difference. The addition of La2O3 to Tie3Ale8Ve6Cre4Moe4Zr substantially reduced the grain size beyond the xSD value for the WAAM data set. This implies that the nucleation events within this alloy are more closely spaced than in alloys not containing La2O3 and infers that some La2O3 particles are responsible for the enhanced nucleation. The cooling curves in Fig. 3 shows no evidence of recalescence during solidification when La2O3 is alloyed with Tie3Ale8Ve6Cre4Moe4Zr indicating that the added particles have sufficient potency to readily nucleate grains. Due to the nucleation of grains, latent heat is released and as the grains grow the larger surface area increases the rate of latent heat release. This was a repeatable observation from multiple tests. This same observation was apparent when LaB6 was alloyed with Tie6Ale4V and decomposed into La2O3 during WAAM [32]. In castings, the absence of recalesence has long been associated with grain refinement and

particularly the presence of potent nucleant particles. For example, Shabestari and Malekan [65] investigated the effects of potent titanium-boron grain refiners on the solidification of aluminium alloys and found that recalescence disappeared completely in alloys containing potent nucleant particles. It is believed that fewer grains (larger grain size) result in melts that undergo large recalescence because some smaller grains that may have initially nucleated are remelted by the temperature rise [33], so the absence of this reheating effect enhances survivability of newly nucleated grains. The low undercooling also means a slower growth rate and a balance is achieved between heat released and heat extracted. In this condition, G will begin to decrease due to thermal diffusion sooner than the case when recalesence occurs. 5.3. Other considerations for refinement with La2O3 Recently it was observed that the Heiple-Roper effect occurred during WAAM of Tie6Ale4V alloyed with LaB6 [32], as did Yin et al. during WAAM of Tie6Ale4V alloyed with CaF2 [66]. The HeipleRoper effect occurs when surface active solutes reverse the surface tension-temperature coefficient of the liquid metal solvent which, due to Marangoni forces, changes the direction of fluid flow and the melt pool shape [67]. The effect is exploited during steel welding using trace sulphur additions to enhance weld penetration [68]. The effect was also observed here when La2O3 was added to Tie3Ale8Ve6Cre4Moe4Zr as evidenced by the deeper penetration along the centreline into the baseplate (Fig. 4 (C1)) in comparison to Tie3Ale8Ve6Cre4Moe4Zr (Fig. 4 (B1)) and the tendency for taller, narrower deposited layers. One consequence of this is that localised temperature gradients within the melt pool will change which could influence xCS . Another consequence is that this effect could amplify other mechanisms for equiaxed grain formation such those covered by the Dendrite Arm Remelting Theory [69], Free Chill Theory [70] or the Separation Theory [71]. Readers interested in the details of these theories are referred to the original works but the common theme to all is that growing dendrites (which initiate at the mould wall in a casting or the previous layer in AM) become separated from the interface and are swept into the interior of the melt pool where they grow into equiaxed grains. In addition to nucleation on La2O3 particles, a reversal of fluid flow direction with La2O3 addition may enhance the

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survivability of detached dendrites at the edges and top surface of the melt while also accounting for coarser grain sizes along the centreline of the component that is exposed to hotter liquid. Another notable observation that requires further research is the potential grain boundary Zener pinning with La2O3 particles. The La2O3 particles were widely distributed throughout the microstructure but occasionally some particles were observed at the grain boundaries. If such particles pin the boundaries, then grain growth in the existing solid would be reduced during AM as new layers are being deposited. A number of rare earth oxides (including La2O3) are known to pin b-grain boundaries during heat treatment of titanium alloys at levels as low as 0.01 wt% [29,72] and thus their potential exploitation for grain size control during AM requires ongoing investigation. This work has demonstrated the importance of nucleant particles in promoting equiaxed grain formation during AM. The La2O3 particles selected for this study are a good starting point but more work is needed to identify better, more potent nucleant particles for b-Ti. The initially large xCS at the start of each layer is underpinned by the difference between epitaxial nucleation (DTN ¼ 0  C) and the potency of nulceant particles (DTN > 0  C) which will always exist unless perfect nucleant particles (i.e. DTN approximately equal to zero) can be identified. The discovery of very potent particles with smaller DTN than La2O3 would make the CET more achievable during AM of titanium alloys.

6. Conclusions This work explored the thermal and composition sensitive factors affecting the columnar to equiaxed transition (CET) during wire based additive manufacturing of titanium alloys. Titanium alloy melt pools were thermally characterised and average temperature gradients varied from 100 to 500  C/mm across a range of deposition rates. A solidification map for Tie6Ale4V was constructed which predicts a mixed columnar þ equiaxed grain morphology. To promote the CET, classical solidification theory indicates that segregating solute is required to provide constitutional supercooling and potent nucleant particles are required to facilitate heterogeneous nucleation of equiaxed grains. Three representative alloys were selected to study the relevance of this theory to additive manufacturing. The first alloy investigated was Tie6Ale4V and represents the control containing no added grain refiners. The second alloy was a commercial alloy containing only growth restricting solutes (Tie3Ale8Ve6Cre4Moe4Zr). The third alloy contained both growth restricting solute and added nucleant particles (Tie3Ale8Ve6Cre4Moe4Zr þ La2O3 particles). As predicted by the solidification map the control alloy (Tie6Ale4V) produced a mixed columnar-equiaxed grain morphology with a very large grain size. The presence of grain refining solute in Tie3Ale8Ve6Cre4Moe4Zr refined the grain size by over 55% principally through a narrowing in columnar grain width but failed to produce a substantial fraction of equiaxed grains. However, the addition of La2O3 particles refined the grain size by 85% and produced an equiaxed zone at the top of the layer and columnar grains at the base. The La2O3 particles are believed to act as heterogeneous nucleation sites for b-Ti when present at the right location and time. The Interdependence model for equiaxed grain formation accounts for the columnar grains at the base of each layer given that constitutional supercooling is not generated until the thermal gradient locally reduces. In-situ melt temperature monitoring confirmed that temperature gradients decrease during layer solidification which permits constitutional supercooling and equiaxed grain nucleation.

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