A ternary hybrid Ti alloy (HYTA) was produced by selective laser melting (SLM), combining all three classes of the titanium system. The hybridisation of the α CP Ti, α-β Ti-64 and β Ti-5553 gave rise to a highly heterogeneous microstructure encompassing all the possible phases and structures which do not co-exist in conventional monolithic Ti a l l o y s . T h e d e t a i l s o f t h e c o m p l e x microstructure were comprehensively characterised, revealing the roles played by such important factors as the distribution of the ingredient alloy particles, the Marangoni convection and diffusion of solute during melting, and microsegregation of solute during solidification. The ternary HYTA exhibited an excellent combination of high tensile yield strength of 1000 MPa and enhanced ductility with uniform elongation of ~15% and total elongation of ~20%, superior to conventional Ti alloys of various types. Such exceptional performance is attributed to the high-degree Submitted to Mater. Sci. Eng. A Zafari et al. 2020 2 heterogeneity and the operation of most major deformation mechanisms which coordinate effectively and transfer smoothly between various parts with a wide range of hardness (4–6 GPa). The outcome has further demonstrated the potential of hybridisation of microstructures as a new strategy for future alloy development. Keywords: Titanium alloys, hybrid alloys, uniform elongation, multiphase, selective laser melting. Submitted to Mater. Sci. Eng. A Zafari et al. 2020 3 1. Introduction N a t u r e h a s p r o v i d e d i n s p i r a t i o n s f o r e n g i n e e r i n g materials for a long time, typically by hybridising organic and ceramic structures in a sophisticated way to deliver unique properties. Such complex organisations, however, are hard to realise i n m e t a l s . T h e a d v a n c e d h i g h s t r e n g t h s t e e l (AHSS) comes close with its versatile phase structures and deformation mechanisms, including twinning induced plasticity (TWIP) and transformation induced plasticity (TRIP) [1]. Most of the other alloy systems are not as blessed. Titanium, with its three main classes of alloys, namely α, α-β, and β with the first two dominated by hcp-structured α and the latter by bcc β [2], can offer a wide variety of structures including pure β, lath and acicular hcp martensitic α', orthorhombic martensitic α'' and lamellar α/β. However, the various microstructures cannot co-exist through conventional processing (e.g. upon quenching single β is only achievable in β Ti while martensitic α' only in α or α-β Ti) [2]. One consequence is that most high strength Ti alloys lack ductility, especially uniform elongation (UE) which is desirable in many manufacturing processes. In particular, two of the most used structural alloys, Ti-6Al-4V (Ti-64) and Ti-5Al-5V-5Mo-3Cr (Ti-5553), showed low work hardening, leading to UE of < 6% with strength increased above 1000 MPa [3]. Although TWIP and TRIP in some β a l l o y s d e s i g n e d by t h e d-electron method can improve ductility, their yield strengths (YS) are normally as low as 500-800 MPa [4, 5]. Similarly, a dual phase α/α' Ti-64 produced using the costly elaborative processing of e l e c t r o n b e a m m e l t i n g ( E B M ) , h o t i s o s t a t i c pressing (HIP) and heat treatment exhibited low UE of 4% when YS was raised to > 1000 MPa [6]. I n o u r r e c e n t s t u d y [ 3 ] , w e h a v e e m p l o y e d s e l e c t i ve laser melting (SLM) to create a novel Ti by hybridising microstructures from existing Ti alloys, referred to as HYbrid Ti Alloy (HYTA), achieving an excellent combination of strength and ductility. This strategy is different from the common practice of mixing elemental powders for in-situ alloying with a view to fabricating a homogeneous structure or producing a simple metal matrix composite [7-10]. Instead, a Submitted to Mater. Sci. Eng. A Zafari et al. 2020 4 heterogeneous microstructure combining that of Ti-5553 (dominated by β) and that of Ti-64 (α'), which do not coexist in conventional alloys, has been obtained in HYTA. This leads to an exceptionally high work hardening rate, UE of 10-13%, total elongation (εt) of ~15% and ultimate tensile strength (UTS) of ~1200-1300 MPa. Further, Ti-5553 and Ti-64 can be selectively manipulated using heat treatments to improve εt to ~20% or YS to 1100 MPa [3]. H o w e v e r , c o m p o s i t i o n a l i n t e r a c t i o n s b e t w e e n t h e i ngredient alloys, which has not been investigated before, deserve special attention to understand how HYTA forms and behaves. Several important factors must be considered in the formation of the hybrid microstructure. First, convection may have played an important role since it would create layers of the ingredient alloys in the melt pool. Such layered structures also formed during AM of elemental powders [11], but were ignored since the aim there was to eliminate them and produce homogeneous structures. Second, different sizes and masses of solute atoms would result in different diffusion rates [12], affecting their distribution. Third, different partition coefficients (k) upon solidification would change the composition of the solidifying liquid, causing microsegregation. Hence, a more in-depth study is essential to reveal the roles played by these factors in the final microstructure of HYTA. On the other hand, the good mechanical properties in the binary HYTA has been attributed to the effective deformation transfer and coordination between the α-β Ti-64 and β Ti-5553 [3]. It appears desirable and of significant interest to add an α type alloy to produce a ternary HYTA with a view to generating a more sophisticated hybridisation of the microstructures covering all the possible phases/structures in the Ti system. It is expected that further property enhancement can be achieved. H e r e w e r e p o r t t h e s u c c e s s f u l p r o d u c t i o n b y S L M o f a n e w t e r n a r y H Y T A w i t h commercially pure (CP) Ti of α type added as the third ingredient to the binary Ti-64 (α-β) and Ti- Submitted to Mater. Sci. Eng. A Zafari et al. 2020 5 5553 (β), focusing on the interactions between the three alloys in the melt pool as well as the resulting tensile properties and deformation mechanisms. A strong influence on CP Ti by the other two alloys was observed while the interaction between Ti-64 and Ti-5553 was limited. Most common phases and structures from all three types of Ti alloys were seen to co-exist in the ternary HYTA, leading to an improved combination of strength and ductility compared to the binary one, with YS of ~1000 MPa, UTS of ~1100 MPa, UE of 15%, and εt of ~20%. These excellent results serve as further proof that the new strategy of hybridising microstructures is promising for future alloy development to achieve unique and exceptional performances in metals. 2. Experimental Materials and Procedures A ternary HYTA was produced by mixing equal portions (~33 wt.%) of plasma atomised CP Ti (Grade 2), Ti-64 and Ti-5553 powders (15–45 µm and distribution of D10 = 22 µm, D50 = 35 µm and D90 = 46 µm) using a mechanical stirrer at 300 rpm. The Renishaw AM250 SLM system was employed for printing rods of 8 mm in diameter and 52 mm in height using stripe scanning strategy on a titanium substrate at room temperature in an atmosphere set at 100 ppm oxygen. The time between two successive depositions (interlayer time) was ~30 s. Other parameters are listed in Table 1. For comparison, a binary HYTA consisting of 50 wt.% Ti-64 and 50 wt.% Ti-5553 was also produced under the same conditions. Nearly full density of > 99.4% was achieved. Table 1 Main printing parameters used. Stripe size (mm) Layer thickness (µm) Point distance (µm) Exposure time (µs) Power (W) Hatch spacing (mm) Laser spot size (µm) Rotation angle between layers (°) 5 30 45 50 200 0.105 66 67 Submitted to Mater. Sci. Eng. A Zafari et al. 2020 6 The distribution of the ingredient alloys before and after SLM was constructed using WDS/EDS elemental maps. The intensity data for Mo, Al and Ti were analysed using MATLAB with the high intensity areas distinguished by setting an empirical threshold of ~70%, i.e. regions with high Mo intensity (i.e. greater than the threshold) identified as Ti-5553 (only Ti-5553 contained Mo), those with high Al intensity but no Mo as Ti-64, and those with Ti only as CP Ti. M i c r o s t r u c t u r e s w e r e c h a r a c t e r i s e d b y s c a n n i n g e l ectron microscopy (SEM, FEI Teneo), and transmission electron microscopy (TEM, FEI Tecnai F20). Backscattered electron (BSE) imaging was performed using analytical mode at an accelerating voltage of 10 keV and current of 13 nA. TEM samples were cut using focused ion beam (SEM/FIB, FEI Nova 200 Nanolab DualBeam). The in-situ lift out FIB technique produced foils of ~2 μm thick which were thinned by milling at 30 keV, resulting in foils of < ~100 nm in thickness after removing contaminations deposited on the surfaces by milling at a lower voltage of 10 keV. Electron probe microanalyser (EPMA, JEOL JXA-8530F) was employed for WDS/EDS analyses using a voltage of 15 keV, dwell time of 50 ms and step size of 1 µm. Further compositional analysis was conducted using atom probe tomography (APT) in a local electrode atom probe (LEAP 4000X Si) at a UV laser pulse repetition rate of 200 kHz, energy of 70 pJ, target evaporation rate of 0.5% and specimen temperature of 50 K. Tensile specimens with a gauge size of 4 mm in diameter and 16 mm in length were cut from the printed rods. The gauge was in the middle section of the rod and the top and bottom layers were not included. The tensile tests were conducted with loading parallel to the build direction at an initial strain rate of 2.5 × 10–4 s–1 according to ASTM E8/E8M-13. A l aser extensometer (MTS LX500) was used to accurately measure strain. Nanoindentation was performed on mirror polished surface using a Hysitron TI950 Triboindenter equipped with a Berkovich pyramidal-shaped indenter tip. Hardness mapping across Submitted to Mater. Sci. Eng. A Zafari et al. 2020 7 an area of 70 × 300 μm2 was performed with spacings of 4 μm between indents, peak load of 10 mN and loading/unloading rate of 1 mN/s. The nanoindentation results were analysed using TriboScan and visualised using MATLAB. 3. Results 3.1. As-mixed powders The binary mixed powder of Ti-64 and Ti-5553 was characterised by the Mo EDS elemental map shown in Fig. 1a. Since Mo only existed in Ti-5553, its intensity was used to distinguish Ti-5553 particles from the Ti-64 ones, resulting in the binary distribution of red (Ti-5553) and blue (Ti-64) in Fig. 1b. Rather than individual particles, clusters of ~100-200 µm of each ingredient alloy (selectively circled) formed. For the ternary mixed powder of CP Ti, Ti-64 and Ti-5553, the Ti, Al and Mo EDS elemental maps are shown in Figs. 1c-e, respectively. The CP Ti particles are those with higher intensity of Ti but free of Al and Mo, Ti-5553 particles are the ones containing both Mo and Al, and Ti-64 particles contain Al without Mo. Quantitative intensity analyses led to the identification of CP Ti particles as yellow, Ti-64 blue and Ti-5553 red in Fig. 1f (this colour scheme is used hereinafter). Compared to Fig. 1b, all three types of clusters are much smaller than 100 µm although a small number of clusters of the order of 100 µm (selectively circled in Fig. 1f) are present. 3.2. As-SLM microstructures Fig. 2a shows BS-SEM of the as-SLM ternary HYTA, revealing the columnar prior β grains (selectively shown by white dotted lines) from epitaxial growth and the dark and bright layers in the melt pool (selectively delineated by yellow dashed lines) with the former containing acicular Submitted to Mater. Sci. Eng. A Zafari et al. 2020 8 martensitic α' plates and the latter β grains confirmed later by high magnification SEM and TEM. The contrast between the layers was resulted from different distributions of heavier and lighter solute atoms, as revealed by WDS elemental mapping in Figs. 2b-g (not very obvious for Cr and Fe because of their much smaller amounts – maximum 3 wt.% Cr and 0.5 wt.% Fe). The three ingredient alloys were identified using similar intensity analyses to those for the as-mixed powders and colour rendered in Fig. 2h. Indeed, WDS/EDS point analyses (not shown here) confirmed their compositions to be close to the nominal ones of the ingredient alloys. It is, however, noted that a significant area in Fig. 2h cannot be associated with any of the ingredient alloys (marked as black). In other words, the black region possesses compositions different from the ingredient alloys. Consequently, the amount of each ingredient alloy was reduced, compared to the as-mixed powder (Fig. 1f), especially for the CP Ti. Closeups of the ternary HYTA are shown in Figs. 3a-d with chemical compositions from WDS analysis at points 1-9 listed in Table 2. In sufficiently thick layers of > ~5-10 µm (Figs. 3a-b) the original compositions of the ingredient alloys are preserved, i.e. points 1 and 4 as Ti-5553, 2 and 5 as Ti-64, and 3 as CP Ti. In contrast, thin layers of < ~1-2 µm were turned into new alloys. For example, points 6 and 7 contained a high amount of Mo and V, indicating their origination from Ti-5553. However, the contents of the β s t a b i l i s i n g e l e m e n t s w e r e m u c h l o w e r t h a n T i - 5 5 5 3 , suggesting that these thin layers had been diluted into a new α-β T i a l l o y n o t s t a b l e e n o u g h t o prevent the formation of needle-shaped martensitic plates (bottom, Fig. 3c). In fact, such a deviation from the ingredient alloy composition occurred even at the boundary of the thick layers (~1 µm in thickness as yellow arrowed in Fig. 3b), reducing β s t a b i l i t y a n d a l l o w i n g m a r t e n s i t i c transformation (inset in Fig. 3b). In addition to the layered structures, there were some isolated regions of single phase acicular α', as shown in Fig. 3d, with compositions different from those of the ingredient alloys (e.g. points 8 and 9). For comparison, the as-SLM binary HYTA shown in Fig. 3e contained mostly thick layers of > 10-20 µm with largely unchanged compositions, resulting in Submitted to Mater. Sci. Eng. A Zafari et al. 2020 9 single β in the Ti-5553 region and α' in the Ti-64 region. In some cases, as selectively delineated by dotted lines in Fig. 3e, the melt pools were entirely occupied by either of the alloys, creating blocks of Ti-5553 (β) and Ti-64 (α'). Table 2 Contents of solute elements (wt.%) at the points numbered 1 to 9 in Fig. 3. Element 1 2 3 4 5 6 7 8 9 Ti-64 Ti-5553 Al 4.72 6.05 - 4.4 5.6 3.89 3.53 4.94 4.58 5.5-6.5 4.4-5.7 V 4.94 3.8 - 4.8 3.5 3.45 3.36 3.87 2.56 3.5-4.5 4-5.5 Mo 5.27 - - 5.21 - 1.97 2.67 - 1.72 - 4-5.5 Cr 2.61 - - 2.88 - 1.07 1.34 - 0.75 - 2.5-3.5 Fe 0.32 - - 0.27 - 0.17 0.27 - 0.2 <0.25 0.3-0.5 The martensitic α' formed in the newly created alloy regions (e.g. Fig. 3d) was further studied using TEM. Acicular α' typical in α-β Ti alloy s was observed (Fig . 4a). The α' plates, in particular large ones, were twinned (arrowed) with the twin planes identified to be of {0111} type based on the SAEDP (Fig. 4b) from the circled spot in Fig. 4a. This twinning is often observed in acicular martensite formed upon rapid cooling from the β field [13, 14]. Additionally, dislocations (selectively arrowed in Figs. 4c and d) and stacking domains (Fig. 4e) were observed. Although substructural dislocations and twins have been widely observed, the stacking domains have been reported only once before [15], which are attributed to atom shuffling as an essential mechanism of the bcc to hcp martensitic transformation. To shed light on the interaction between the ingredient alloys, the interfacial area between the ingredient alloys was investigated. One interface between CP Ti and Ti-5553 (dashed line) is shown in Fig. 5a together with the corresponding Al, Mo, V and Ti elemental maps. The higher concentrations of Al and Mo (brighter contrast) below the dashed line and the higher concentration Submitted to Mater. Sci. Eng. A Zafari et al. 2020 10 of Ti above it identify the area below as initially T i - 5 5 5 3 a n d a b o v e a s o r i g i n a l l y C P T i . T h e boundary is, however, no longer sharp, with significant amount of alloying elements in Ti-5553 penetrating into the original Ti region (a composition of ~2 wt.% V, ~1.5 wt.% Al, 1 wt.% Mo and < 1 wt.% Cr was revealed by point EDS analysis above the dashed line). It is apparent that V was almost uniformly distributed across the original interface but there were still higher amounts of Al and Mo on the Ti-5553 side despite their penetration into Ti. The changes in composition across the interface led to the formation of acicular martensitic plates in both regions. More detailed characterisation was performed using TEM at such an interface, as shown in Fig. 5b-h. Four main zones were identified, as delineated by dashed lines and labelled A-D in Fig. 5b. SAEDP from A is shown in Fig. 5c and DF for (110)β in Fig. 5d, both indicating single β phase in A, i.e. Ti-5553. By slight tilting, substructural dislocations in D was observed (Fig. 5e), as typically seen in martensitic α' in CP Ti or diluted Ti alloys. SAEDP from Zone B is shown in Fig. 5f, identifying the martensitic plates observed to be orthorhombic α'' twinned on (111). However, the structure in Zone C looked more like acicular α' seen in Fig. 4a, as confirmed by SAEDP in Fig. 5g taken from multiple α' plates. The STEM-EDS elemental maps for the red-dot framed region in Fig. 5b are shown in Fig. 5h. The Al and V maps revealed higher intensity of both elements in B than A, although V showed a more significant difference. On the other hand, Cr and particularly Mo were higher towards A. Hence, Ti-5553 was slightly depleted of Al and more significantly of V, but enriched mainly in Mo, consistent with the elemental maps in Fig. 5a and point WDS/EDS analyses. The interface between CP Ti and Ti-64 was hard to find since Ti would obtain the same acicular α' microstructure as Ti-64 after SLM (Fig. 4). The interaction of Ti with Ti-64 is, however, expected to be like that with Ti-5553 since they contain similar amounts of Al and V. On the other hand, the interaction between the Ti-64 and Ti-5553 was much more limited, as point WDS/EDS analyses on adjacent Ti-64/Ti-5553 layers revealed almost no change in their compositions except Submitted to Mater. Sci. Eng. A Zafari et al. 2020 11 for regions very close to the interfaces or very thin Ti-64 and Ti-5553 layers (Fig. 3, Table 2 and Ref. [3]). In general, a large number of WDS/EDS point analyses in those regions altered from the original ingredient alloys revealed compositions of ~2-3 wt.% V, ~2-3 wt.% Al and <1-2 wt.% Mo and Cr, indicating that V and Al were the main solute elements responsible for the composition change during SLM. 3.3. Internal cells in Ti-5553 Fig. 6a shows internal cells in the Ti-5553 regions of HYTA revealed after chemical etching. The formation of such cells in Ti-5553 is attributed to the effect of constitutional supercooling on the geometry of growth during solidification [16]. However, it is not known how solute elements are distributed in front of the solid/liquid interface to create such supercooling. This is particularly important in HYTA, since microsegregation of solute e l e m e n t s w o u l d a f f e c t t h e s u r r o u n d i n g solidifying liquids from other ingredient alloys, leading to various microstructures. APT was performed on an as-SLM single Ti-5553 alloy to avoid complications in HYTA. This was justified since the cells were only found in thick Ti-5553 layers in HYTA which were not affected by the other ingredient alloys. Pre-APT SEM of the as-SLM Ti-5553 is shown in Fig. 6b, revealing the locations of the melt pool boundaries and fine cells. Single atom maps for various elements from an area inside a melt pool and in its boundary are shown in Figs. 6c and d, respectively. Visual inspection appeared to suggest that all the elements were uniformly distributed in both regions. Quantitative distribution of the elements other than Mo as a function of the distance from the boundary confirmed this observation (Fig. 6e). Mo, however, tended to enrich in the melt pool boundary. In particular, the concentration of Mo at the boundary (distance 0 in Fig. 6e) was ~6 at.% which gradually reduced to ~4 at.% inside the melt pool. The average chemical compositions inside the melt pool and in the boundary are listed in Table 3, confirming the most significant difference in Mo of ~1 at.%. Submitted to Mater. Sci. Eng. A Zafari et al. 2020 12 Table 3 Average chemical compositions inside the melt pool (MP) and in the melt pool boundary (MPB) analysed by APT. T i A l V M o C r O F e N S i C MP at.% 80.8 8.15 4.60 3.21 2.35 0.43 0.30 0.09 0.05 0.005 wt.% 80.94 4.60 4.90 6.45 2.56 0.14 0.35 0.03 0.03 ~0.00 MPB at.% 80.4 8.29 4.35 4.21 1.99 0.45 0.17 0.07 0.02 0.006 wt.% 79.85 4.64 4.60 8.38 2.15 0.15 0.2 0.02 0.01 ~0.00 3.4. Hardness distribution and tensile properties F i g . 7 a s h o w s t h e h a r d n e s s m a p o b t a i n e d f r o m n a n o indentation along the build direction, and the WDS elemental maps of Ti, Al and Mo are shown in Figs. 7b-d, respectively. The same method used for the production of Fig. 2h was employed to construct the distribution of the ingredient alloys as well as the altered regions from the elemental maps, as shown in Fig. 7e. The white dots with white contour in Fig. 7e represent hardness data points of > 5.2 GPa in Fig. 7a, associated with α' with high contents of solute elements (e.g. Ti-64). On the other end of the hardness spectrum, the green dots with green contour indicate regions of < 4.5 GPa, related to areas largely free of acicular α' (e.g. Ti-5553 and CP Ti). The vast majority of the area exhibited hardness values between these two limits. In general, the distribution of hardness is very heterogeneous thanks to hybridisation. Fig. 7f shows the tensile engineering stress-strain curves for the binary and ternary HYTAs as well as ingredient alloys together with derived properties. The ternary HYTA exhibited much higher YS of ~1000 MPa, εt of ~20%, and UE of ~15%, compared to YS of ~750 MPa, εt of ~14%, and UE of ~11% for the binary HYTA. Although the monolithic Ti-64 showed higher strength as a result of the higher amount of Ti-64 (the hardest ingredient alloy), its UE of ~5% was significantly Submitted to Mater. Sci. Eng. A Zafari et al. 2020 13 lower compared to both HYTAs, and εt of ~14% was much inferior to 20% for the ternary HYTA. While εt for Ti-5553 and CP Ti were comparable to that for the ternary HYTA, their strengths were considerably lower by > 200 MPa. The work hardening rate for the binary and ternary HYTAs, calculated as dσ/dϵ with σ and ϵ as true stress and true strain, respectively, is shown in Fig. 7g, displaying much higher values in the binary HYTA at ϵ > 5%. The calculated UE using the criterion dσ/dϵ = σ gave rise to 11-12% for the binary and 14-15% for the ternary, consistent with the values obtained from the engineering stress and strain curves (Fig. 7f). 3.5. Microstructures after tensile deformation T h e d e f o r m a t i o n s t r u c t u r e s a f t e r t e n s i l e s t r a i n i n g to ~1.5% are shown for the binary (Fig. 8a) and ternary (Figs. 8b-j) HYTA. Slip bands were present in both although their number was higher and spacing smaller in the ternary HYTA (Figs. 8a and b). In addition to dislocation slip, lens-shaped plates were formed in the β r e g i o n o f t h e t e r n a r y H Y T A ( S T E M - H A A D F i n F i g . 8 c ), identified as stress induced martensitic α'' using the inset SAEDP which also revealed diffraction from athermal ω (ωath). Although most of the strain was accommodated in β, dislocations were also generated in α' (yellow triangles) near the α'/β interface (dashed-line), as shown in Fig. 8d. At some α'/β interface, rotation of β crystal can be observed in a thin band (separated by dashed lines) in Fig. 8e, as proven by SAEDPs in Figs. 8f-h corresponding to areas f, g and h in Fig. 8e. In particular, Fig. 8f confirms area f as β. When diffraction was taken from g which crossed the dashed line, extra spots in addition to those identified in Fig. 8f (pointed by triangles) were present, forming the pattern shown by solid lines (Fig. 8g). This new pattern turned out to be the same as that shown in Fig. 8h taken from location h in Fig. 8e (i.e. entirely within the thin band) and was indexed as β. In other words, both areas f and h are β but with different orientations. DF from the circled spot in Fig. Submitted to Mater. Sci. Eng. A Zafari et al. 2020 14 8g confirmed the presence of crystals with different orientations in area h (Fig. 8i). The lens-shaped martensitic plates found in area h (Fig. 8j) was identified to be α'' using the corresponding SAEDP (inset) as well as the DF. Further straining to 6.5% led to the formation of more slip bands and dislocation pileups at the α'/β interface in the ternary HYTA (Fig. 9a). At another location, Fig. 9b shows a thin stripe delineated by dashed lines sandwiched between two Ti-5553 regions. Chemical analysis showed that the contents of β s t a b i l i s i n g e l e m e n t s i n t h i s a r e a w e r e s l i g h t l y l ower than in Ti-5553. The plates observed in this region were identified as martensitic α'' using electron diffraction (not shown here). The interaction between slip bands and α'' created dislocation pileups (Fig. 9b). DF TEM for α'' in Fig. 9c shows that some of the slip bands (selectively shown by dashed lines) were able to cut through the martensitic plates to create square-shaped α'', as seen in a commercial β Ti-5553 [17]. At tensile fracture, steps at the α'/β interfaces (shown by dashed lines) were formed (Fig. 9d), aligned with slip bands (selectively shown by yellow dotted lines). Tilting the sample revealed that the steps were able to penetrate into α' by creating slip bands, as arrowed in Fig. 9e. Dislocation pileups at α'/β interfaces could trigger twinning in the entire layer of a diluted Ti alloy (Fig. 9f), identified as {1011} type using the inset SAEDP. The red dotted line in Fig. 9f distinguishes two β regions with different orientations from crystal rotation, similar to observation in Fig. 8e. Fig. 9f also shows dislocation pileups at both β/β (red dotted line) and α'/β interfaces. The former can be seen more clearly using STEM-HAADF in Fig. 9g. Submitted to Mater. Sci. Eng. A Zafari et al. 2020 15 4. Discussion H y b r i d i s a t i o n o f m i c r o s t r u c t u r e s f r o m d i f f e r e n t a lloys promises to offer new opportunities for materials development by providing exceptional performances not possible in individual ingredient alloys. SLM is shown to be most potent in producing hybrid alloys thanks to ultrahigh cooling rates (> 106 K/s) which largely preserves the identities of the ingredient alloys. The first type of such materials was successfully realised from SLM of pre-alloyed Ti-64 and Ti-5553 powders, leading to a binary HYTA with excellent mechanical properties [3]. In the present work, a third alloy, CP Ti of α t y p e , w a s a d d e d t o t h e α/β a n d β t y p e s t o m a k e a t e r n a r y H Y T A encompassing all possible phases and structures in the Ti system. The resulting microstructure turns out to be much more complex, compared to that of the binary HYTA, leading to improved tensile properties. Several important factors involved in SLM processing such as the distribution and proportion of ingredient powders, convection in the m e l t p o o l , d i f f u s i o n i n l i q u i d a n d m i c r o -segregation need to be studied to understand their roles in the formation of the sophisticated microstructure. It is also critical to understand the fundamentals responsible for the excellent mechanical properties in the ternary HYTA. 4.1. Effects of convection, diffusion and microsegregation Convection T h e s c h e m a t i c s i n F i g . 1 0 s h o w t h e f o r m a t i o n o f t he final microstructures in one melt pool (of the order of 100 µm in size) in the binary (rows a, b and c) and ternary (row d) HYTAs. Column 1 displays the possible distributions of the pre-alloy ed powders of Ti-64 (blue), Ti-5553 (red) and Ti (yellow). Column 2 shows the distributions of the alloys after melting, formed by the Marangoni convection (illustrated at the bottom) resulting from a steep thermal gradient in the melt pool [18]. Column 3 illustrates the microstructures immediately upon solidification, exhibiting prior Submitted to Mater. Sci. Eng. A Zafari et al. 2020 16 β g r a i n b o u n d a r i e s ( w h i t e l i n e s ) a n d a l l o y d i s t r i b u tions unchanged from column 2. Column 4 represents the final microstructures after quenching to low temperatures. In the binary HYTA, many of the clusters of Ti-64 and Ti-5553 are <~100 µm in size (Fig. 1b), filling the melt pool as in Fig. 10a1, although some clusters are as large as ~200 µm, resulting in the entire melt pool being occupied by either alloy (Figs. 10b1 and c1). When there are two alloys in the melt pool, the Marangoni convection would create liquid layers of the alloys (Fig. 10a2) and significant mechanical mixing, which could result in one uniform composition in the whole melt pool, is prevented by the ultrahigh cooling rates during SLM (i.e. > 106 K/s). Following rapid solidification, prior β columns are formed in both but the shape and distribution of the alloy layers are unchanged (Fig. 10a3). The original composition of each ingredient alloy (i.e. Ti-64 or Ti-5553) is largely preserved although some dilution could occur at the Ti-64/Ti-5553 interface [3]. Since the ingredient alloys have similar melting points and both solidify into β below the solidus temperatures, perfect fusion between particles is achieved in HYTA with density of > 99.4%. The internal cells formed in Ti-5553 (Fig. 6a) are represented by the black lines in Figs. 10a3 and c3 (discuss later). Quenching to low temperatures transforms the Ti-64 layers (blue) into acicular martensitic α', but high contents of β stabilising elements in Ti-5553 (red) suppress any m ar t e n s i t i c t r a n s f or m a t i o n s (Fig. 10a4). In the cases of single occupations (Figs. 10b1 and c1), there is only one alloy in the melt pool, resulting in either a Ti-64 (Fig. 10b4) or a Ti-5553 (Fig. 10c4) microstructure, as observed in Fig. 3e. T h e c l u s t e r s i n t h e t e r n a r y H Y T A a r e c o n s i d e r a b l y smaller (Fig. 1f) and most melt pools would contain all three alloys (Fig. 10d1), making the microstructure more complicated since layers created by the convection can have a wide range of thicknesses (Fig. 10d2 and observed in Figs. 3a-d) and in most cases smaller than those in the binary HYTA. The inside of the relatively thick layers (> ~5-10 µm) can maintain the original compositions, but the entire thinner ones (< ~1-2 µm) as Submitted to Mater. Sci. Eng. A Zafari et al. 2020 17 well as a thin layer at the interface of the thick layers would undergo compositional changes (Table 2). These layers with altered compositions are represented by orange in Fig. 10d2, taking into consideration that CP Ti is affected considerably more than the other two alloys (Fig. 2h). The compositions of the orange layers fall in the range of the α-β Ti alloys (Table 2, WDS/EDS point analyses and elemental maps presented in 3.2), as shaded in the schematic phase diagram in Fig. 11a. Therefore, after quenching to low temperatures, these orange layers with new compositions would also transform into acicular α' (Fig. 10d4 as observed in Figs. 4 and 5). The much smaller amount of remaining CP Ti would undergo β to massive martensitic α' transformation [2], creating large α' plates with substructural dislocations (Figs. 5b and e). Although overlapping melt pools in one track and two consecutive layers deposited could lead to further mixing and diffusion, the effects on the overall microstructure as described here are minimal. Additionally, the low temperature of the substrate and long interlayer time (~30 s) minimised the heat accumulation during SLM, resulting in the precipitation of only a small amount of α i n i s o l a t e d r e g i o n s i n t he bottom layers. Diffusion in the liquid state C o n v e c t i o n i n c l u d e s t w o m a i n s u b - m e c h a n i s m s o f a d vection (directional bulk flow transfer of heat, in which the Marangoni effect falls) and diffusion (non-directional transfer of mass particles stemmed from a concentration gradient). The Marangoni convection has been identified as the cause for forming the alloy layers but not for altering the alloy compositions in the thin layers and at the interface. It is time to ask what mechanisms are responsible for it and why CP Ti is affected more than the other two alloys. It is recalled that V and Al are the main elements responsible for the changing compositions. For example, Figs. 5a and h clearly show rejection of V and Al from Ti-5553, leaving Mo behind. This confirms that the convection (refers to the sub-mechanism advection) is not the reason since it would not selectively reject or enrich any particular Submitted to Mater. Sci. Eng. A Zafari et al. 2020 18 elements. Instead, solute diffusion in the liquid state and microsegregation should be considered. The diffusion coefficient of solute x in liquid Ti (Dx Ti) can be calculated using [12]  =./   (,)∗( ∗) (1) where p denotes pressure in atmospheres, T temperature (K), dx Ti the average atomic diameters of x and Ti (i.e. # =() ), mx a n d mTi t h e a t o m i c m a s s e s o f x a n d T i , r e s p e c t i v e l y , a n d Ω (,)∗(% ∗) the collision integral taken as unity for rigid spheres. At a given temperature, the atomic size and mass of the solute element determine its diffusivity in liquid Ti. Table 4 lists #  and &+&/2&& calculated for the four main solute elements of V, Al, Cr and Mo in the ternary HYTA. Al diffuses faster than all the other elements and Mo is the slowest, while Cr and V diffuse at the same rate between Al and Mo. The diffusion coefficients calculated for Ti-x (x = V, Al, Mo or Cr) at the temperature range of 1600-2000 °C (i.e. in liquid) using CALPHAD are shown in Fig. 11b, confirming that Mo diffuses at a much slower speed. The faster diffusion of Al and V allows them to move into the neighbouring layer in much shorter times and to play more important roles in compositional changes. Although Cr can diffuse as fast, its concentration is lower, making it less visible. Table 4 The atomic diameter (dx) and mass (mx), )* +,= -*+-+,/.-*-+, and dx Ti for solute x (= Al, V, Mo or Cr) in Ti, giving rise to Dx Ti/DAl Ti denoting the diffusion coefficient of x in liquid Ti normalised by that of Al. Element dx (pm) mx (u) Mx Ti /* +,. Dx Ti/DAl Ti Ti 430 47.867 - - - Al 286 26.981 0.0289 128164 1 V 410 50.941 0.02 176400 0.61 Mo 418 95.95 0.0.016 179776 0.52 Cr 400 51.996 0.02142 172225 0.62 Submitted to Mater. Sci. Eng. A Zafari et al. 2020 19 In addition to diffusion coefficients, the amount of diffusion is also dependent on the concentration gradient of the solute element and this explains why the interaction between CP Ti and the other two alloys is more significant than that between Ti-5553 and Ti-64. CP Ti contains almost no solute elements and when coupled with Ti-64 or Ti-5553, which is rich in solute contents, a large concentration gradient is established for each solute element, providing a great driving force for its diffusion into CP Ti. At the same time, the concentration of Ti in CP Ti is higher than in Ti-64 and Ti-5553 by >10 and >18 wt.%, respectively, driving Ti to diffuse from CP Ti to the alloys. Consequently, Ti-64 and Ti-5553 are diluted into the range of α-β alloys, while CP Ti is enriched in solute elements so it also moves from the α to the α-β region (shaded in Fig. 11a). The formation of new α-β alloys is readily observed in the ternary HYTA, as typified in Figs. 3c-d, 4 and 5, and they exhibit a wide range of compositions. For example, in Zone B of Fig. 5b, which is very close to Ti-5553 (Zone A), the amount of β stabilising elements is less than that of Ti-5553, but still high to prevent any α' from forming upon quenching. Instead, the orthorhombic α'' is formed (Fig. 5f). In other words, the local composition in B would be close to the α-β/β Ti boundary, as indicated by Alloy 1 in Fig. 11a. On the other hand, Zone C in Fig. 5b was closer to CP Ti (Zone D) with lower contents of β stabilising elements, creating an alloy closer to that of Ti-64 (e.g. Alloy 2 is in Fig. 11a), which undergoes β to acicular α' transformation. Finally, it is worth noting that the diffusion length was apparently of the order of 1 µm since the composition changes only occurred in a slice (~1 µm) at the interface in the thick layers of > 5 µm but throughout the entire volume of the thin layers of < 1-2 µm. Microsegregation and formation of internal cells M i c r o s e g r e g a t i o n c a n i n f l u e n c e t h e l i q u i d c o m p o s i tion in front of the solidifying alloy. No microsegregation is expected in Ti-64 since the partition coefficients (k) for V and Al are very close to unity [19]. Microsegregation in Ti-5553, however, is significant and might be responsible for the Submitted to Mater. Sci. Eng. A Zafari et al. 2020 20 formation of the internal cells (Figs. 6a-b). Such internal cells in as-SLM Ti-5553 was attributed to constitutional supercooling [16] although no compositional analysis was conducted to show which elements created the supercooling. The partition coefficients of V, Cr and Fe are smaller than unity (i.e. k < 1), making them depleted at the beginning of solidification and enriched towards the end while those of Al and Mo are larger than unity (i.e. k > 1 ) , m a k i n g t h e m e n r i c h e d a t t h e s t a r t . During SLM, solidification starts at the melt pool boundary and consequently, Fe and Cr (with effective k of 0.6 and 0.8, respectively [20, 21]) are expected to be depleted and Mo and Al (with k of 1.41 and ~1.1-1.2, respectively [22, 23]) enriched at the boundary. Indeed, there are more Mo and Al and less Fe and Cr in the melt pool boundary based on APT results (Table 3). However, the most considerable segregation happens to Mo with its content ~1-2 at.% higher at the melt pool boundary whereas the segregation is less visible for the other elements (Table 3 and Fig. 6e). It is therefore concluded that the supercooling required for cellular growth is mostly provided by Mo microsegregation. Indeed, it was shown that adding Mo to Ti-64 could change its growth into cellular during SLM [23]. This can also be another factor affecting the composition of the solidifying liquid in the ternary HYTA. The enrichment of Mo in Ti-5553 reduces its amount in the surrounding liquid which eventually creates new α-β Ti with less Mo than Al and V. 4.2. Tensile behaviour T a b l e 5 c o m p a r e s t e n s i l e p r o p e r t i e s o f t h e t e r n a r y H Y T A t o t h o s e o f b i n a r y H Y T A , monolithic Ti-64 and Ti-5553 which, besides being the ingredient alloys in HYTA, are two important commercial Ti alloys on their own right, and a range of TWIP/TRIP β Ti alloys whose high work hardenability owing to twinning and stress induced martensitic transformation (SIMT) has made them attractive. The data in Table 5 were all obtained from similar testing conditions (i.e. room temperature and quasi-static deformation), although not all followed ASTM E8. The ternary HYTA clearly exhibits an improved combination of strength and ductility. Compared to Ti-64, it Submitted to Mater. Sci. Eng. A Zafari et al. 2020 21 possesses similar YS of ~1000 MPa and UTS of ~1100 MPa but with substantially higher UE (~15%) and εt (~20%). Although the strength of Ti-5553 can be higher from precipitation hardening and total elongation reach ~18-20% (at a lower YS of 800 MPa though), its UE is in the range of 0-8% as a result of very small work hardening. The ternary HYTA is also superior to its binary counterpart with higher YS by 200-600 MPa, UE by 2-5% and εt by > 5%. As far as TWIP/TRIP β Ti alloys are concerned, they generally exhibit considerably lower strength with YS and UTS of 400-800 MPa and < 990 MPa, respectively, although ductility can be comparable or higher (when strength is at the lower end of the spectrum). Table 5 Engineering yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE) and strain-to-fracture (εt) of ternary and binary HYTA, and Ti-64, Ti-5553, and TWIP/TRIP β Ti produced using different methods. Alloy YS (MPa) UTS (MPa) UE (%) εt R e f e r e n c e Ternary HYTA 1000 1100 ~15 20 Present work Binary HYTA 750 1205 11 14 Present work Binary HYTA 370 1200 11 18-20 [3] Binary HYTA 1100 ~1250 8 13-14 [3] Binary HYTA ~800 ~1250 13 15 [3] Ti-64 790 870 < 6 18 [24] Ti-64 990 1100 < 6 8 [24] Ti-64 945 1070 ~8 10 [25] Ti-64 875-939 971-1192 4-10 4-12 [6] Ti-64 1100 1200-1300 2-5 6-15 [26-28] Ti-5553 800 800 0 18-20 [3] Ti-5553 1245 1310 8 17 [29] Ti-5553 1050-1220 1150-1300 < 8 4-10 [30] TWIP/TRIP β Ti Ti-15Mo-5Zr 760 801 10 19 [31] Ti-10Mo-2Fe 866 944 8 17 [31] Ti-15Mo 439 714 27 49 [32] Ti-15Mo-1Fe 837 837 0.2 20 [32] Ti-10Mo-1Fe 563 825 24 43 [33] Ti-10Mo-1Fe 626 833 24 32 [34] Ti-20V-2Nb-2Zr 550-600 < 700 11-20 25-35 [35] Ti-16V-1Fe 760 980 20 27 [36] Submitted to Mater. Sci. Eng. A Zafari et al. 2020 22 Ti-12V-2Fe-1Al 760 900 30 37 [36] Ti-12V-2Fe-2Al 980 990 7 17 [36] Ti-12Mo* ~500 ~700 30 ~48 [4] Ti-8.5Cr-1.5Sn* ~550 ~920 35 ~40 [37] Ti-18Zr-13Mo* 800 ~900 18 18 [5] * True stresses and strains given in the