Additive Manufacturing Letters 6 (2023) 100144

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Short Communication

Microstructure of an additively manufactured Ti-Ta-Al alloy using novel
pre-alloyed powder feedstock material
C. Lauhoﬀ a,∗, T. Arold b, A. Bolender b, M.W. Rackel c, F. Pyczak c, M. Weinmann d, W. Xu a,
A. Molotnikov a, T. Niendorf ba

RMIT Centre for Additive Manufacturing, School of Engineering, RMIT University, Melbourne, VIC, 3000, Australia
Institute of Materials Engineering, University of Kassel, Mönchebergstr. 3, 34125, Kassel, GermanycHelmholtz-Zentrum Hereon, Max-Planck-Str. 1, 21502, Geesthacht, Germanyd
TANIOBIS GmbH, Im Schleeke 78 – 91, 38642, Goslar, Germanyb

a r t i c l e
Keywords:
PBF-EB/M
Alloy formation
EIGA
Chemical homogeneity
Refractory metals
Synchrotron diﬀraction

i n f oa b s t r a c t
Binary Ti-Ta and ternary Ti-Ta-Al alloys attracted considerable attention as new potential biomaterials and/or
high-temperature shape memory alloys. However, conventional forming and manufacturing technologies of refractory based titanium alloys are diﬃcult and cost-intensive, especially when complex shapes are required.
Recently, additive manufacturing (AM) emerged as a suitable alternative and several studies exploited elemental
powder mixing approaches to obtain a desired alloy and subsequently use it for complex shape manufacture.
However, this approach has one major limitation associated with material inhomogeneities after fabrication. In
present work, novel pre-alloyed powder material of a Ti-Ta-Al alloy was additively manufactured. Hereto, electron beam powder bed fusion (PBF-EB/M) technique was used for the ﬁrst time to process such Ti-Ta based alloy
system. Detailed microstructural analysis revealed that additively manufactured structures had a near full density
and high chemical homogeneity. Thus, AM of pre-alloyed feedstock material oﬀers great potential to overcome
major roadblocks, even when signiﬁcant diﬀerences in the melting points and densities of the constituents arepresent as proven in the present case study. The homogeneous microstructure allows to apply short-term thermal
post treatments. The highly eﬃcient process chain detailed will open up novel application ﬁelds for Ti-Ta based
alloys.

1. Introduction
Nowadays, additive manufacturing (AM) techniques are widelyadopted in academia and industry to fabricate functional metal partsand components of unprecedented geometrical complexity. For the two
major categories of metal AM technologies, i.e. powder bed fusion (PBF)and directed energy deposition (DED), structures are fabricated directlyfrom a computer-aided design (CAD) ﬁle through layer wise meltingof either powder or wire feedstock material. The possibility for a toolfree design allows to overcome limitations of conventional manufacturing processes, especially when challenging materials have to be processed causing high production costs [1]. Furthermore, AM techniques
were reported to be employed for direct microstructure design. By controlling the thermal gradient and the solidiﬁcation velocity via an adequate choice of the processing parameters, local microstructural features
and, thus, mechanical properties can be tailored [2–5]. Given these AM
characteristics, topology optimized geometries and locally tailored microstructures can lead to a new generation of lightweight designs be-∗

ing highly attractive for applications in the automotive, aerospace andbiomedical sectors.
Over the last two decades, a variety of materials were qualiﬁed for
AM technologies to open up new industrial markets [6,7]. Among thesematerials, titanium and its alloys are particularly prominent due to theirsuperior properties, i.e. high strength to weight ratio, good biocompatibility, and excellent corrosion resistance [8]. Ti-Ta alloys have attracted
signiﬁcant attention for multipurpose use as bone implants as well as
high-temperature actuators due to an enhanced osseointegration [9–11] and shape memory properties [12], respectively.
Binary Ti-Ta alloys compete with Ti-6Al-4V, the latter being the mostcommon alloy used for orthopedic implants. Tremendous eﬀorts havebeen devoted on processing of Ti-6Al-4V via AM and the relationships
between process, microstructure and mechanical properties are well understood [13–15]. However, the suitability of Ti-6Al-4V as bone implantalloy has recently been questioned due to potential toxicity concerns associated with the alloying elements [16,17] and its inherently high elastic modulus of 113 GPa being much higher than that of human bone (5 –Corresponding author.
E-mail address: christian.lauhoﬀ@rmit.edu.au (C. Lauhoﬀ).

https://doi.org/10.1016/j.addlet.2023.100144
Received 23 March 2023; Received in revised form 21 April 2023; Accepted 1 May 2023
2772-3690/© 2023 The Author(s). Published by Elsevier B.V. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/)

C. Lauhoﬀ, T. Arold, A. Bolender et al.

Additive Manufacturing Letters 6 (2023) 100144

Table 1PBF-EB/M processing parameters used for fabrication of
the Ti-Ta-Al samples investigated in the present study.

30 GPa). The latter promotes the stress-shielding eﬀect, bone resorption
and eventually implant failure [18]. Alloying titanium with tantalum, inturn, seems to be a highly promising approach for designing novel biomaterials. Tantalum is known for its nontoxic nature and Ti-Ta alloys
were reported to show reduced elastic moduli [19], leading to betterbiocompatibility and superior implant integration with the human body(compared to Ti-6Al-4V [9–11]), respectively.
Furthermore, Ti-Ta alloys can feature shape memory behavior well
above 100°C and, thus, came into focus of research as high-temperatureshape memory alloy (HT-SMA) candidates [12,20]. Their unique functional properties are based on a thermoelastic, reversible phase transformation between the high-temperature austenitic 𝛽-phase (body-centeredcubic, bcc) and the low-temperature martensitic 𝛼’’-phase (orthorhombic) [12]. Transformation strains of up to 3.6% were reported [21]. Unfortunately, binary Ti-Ta is prone to rapid functional degradation during thermal or thermo-mechanical cycling. Formation of the hexagonal 𝜔-phase leads to the stabilization of the high-temperature 𝛽-phase,resulting in a deterioration or even loss of the functional properties[12,22,23]. However, alloying with ternary elements such as tin andaluminum can improve the functional stability by delaying the formation of the 𝜔-phase [21,24,25]. In addition, compared with Ni-Ti-Hf,being currently the most promising HT-SMA system [26,27], Ti-Ta-X(e.g. X = Al, Sn) HT-SMAs contain more reasonably priced constituentsand provide good workability [12,25].
Despite these advantages, Ti-Ta(-Al) alloys are still not widelyadopted in industrial applications. The challenging alloy formation
caused by the substantial diﬀerences in the alloying elements’ melting
points and densities is a major roadblock towards their widespread use.
The diﬀerences can cause vaporization of the lower-melting elements
and chemical inhomogeneities by segregation of the constituents during
the solidiﬁcation process, respectively [28,29]. In case of conventional
processing, alloy ingots must be remelted and annealed many times toobtain adequate homogeneity [28]. Recently, AM came into focus asa new processing route being capable to produce refractory titanium
alloys (e.g. Ti-Nb, Ti-Ta) via in situ alloying. Since in most cases prealloyed feedstock material was not available so far, it was shown thatmechanical mixing of elemental powders and subsequent processing canlead to the desired alloy fabrication [30–35]. However, it was notedthat in situ alloying results in local inhomogeneity of the microstructure
due to presence of unmolten niobium and tantalum particles, at least
when the energy input is too low. On the other hand, an increase in energy leads to substantial keyhole formation [30–32]. Some success wasreported by Brodie et al. [32–34] for Ti-Ta and Huang et al. [35] for
Ti-Nb. To promote homogeneity, the authors utilized a remelting scanning strategy. However, this approach reduces the productivity of the
AM process and limits its applicability of printing structures with ﬁne
features.
Pre-alloyed feedstock material is an optimal candidate to overcome
aforementioned limitations. One example was reported by Schulze et al.
[36] focusing on a Ti-Nb alloy. Ternary Ti-Nb-Ta pre-alloyed powders
especially designed for application in AM processes have been recently
patented [37]. However, to the best of the authors’ knowledge, no work
has been published on AM of pre-alloyed Ti-Ta based material. Furthermore, all of the previous studies about additively manufactured Ti-Ta
based alloy systems focused on laser beam powder bed fusion (PBFLB/M, abbreviation according to ISO/ASTM 52900 standard terminology). Thus, in order to close this prevailing gap, the present study reports on a pre-alloyed Al-modiﬁed Ti-Ta alloy and its processing byelectron beam powder bed fusion (PBF-EB/M). Aim of the investigations conducted was to shed light on the microstructural evolution along
the whole process chain from feedstock material to additively manufactured bulk material. Hereto, detailed microstructure analysis has been
performed using optical and scanning electron microscopy as well as
high-energy synchrotron diﬀraction. In particular, alloy formation being characterized by a homogenous element distribution is proven forthis highly challenging processing conditions, i.e. processing under vac-

Heating before process start
Start temperature
Hold time before process start

850 °C
1200 sPreheating
Beam speed
Beam current
Focus oﬀset
Number of repetitions

12000 mm/s
18 mA
75 mA
9

Melting
Acceleration voltage
Beam speed
Beam current
Focus oﬀset
Hatch distance
Layer thickness

60 kV
3500 mm/s
10 mA
3 mA
0.1 mm0.05 mmuum of an alloy system featuring melting point and density diﬀerencesin extreme.
2. Material and methods
2.1. Material and processing
In the present study, a Ti-Ta-Al alloy with a nominal chemical composition of Ti-25Ta-5Al (wt.-%) was additively manufactured. While atantalum content of 25 wt.-% was chosen in light of orthopedic implant
applications [10,19], small amounts of Al have been added (regardlessof the potential toxicity concerns for biomaterials) due to its well-known
stabilizing eﬀect on shape memory properties [21,24]. Thereby, an alloy
system with constituents featuring highly diﬀerent physical properties
was evaluated. Pre-alloyed powder feedstock material was produced via
electrode induction melting inert gas atomization (EIGA). For powder
manufacturing, pure elemental powders were blended and subsequently
cold isothermally pressed (CIP) to form rod-shaped electrode material
with a diameter of 45 mm and a length of about 330 mm. Employinga Leco TC-436 analysing unit, the oxygen and nitrogen contents of the
CIP rod were determined to be 2834 μg/g and 292 μg/g, respectively.
The carbon content was found to be 120 μg/g using a Leco TC-444 analyzing system. The atomization process was performed using inert argongas (99.999%) at a gas pressure of 25 bar. From the obtained sphericalpowder, a powder fraction featuring nominal particle sizes between 63
and 125 μm was extracted by sieving.
The PBF-EB/M process was conducted on a GE Additive Arcam A2X
machine. In order to limit the volume of required powder material, abuild plate reduction with dimensions of 50 × 50 mm2 was utilized.
Cuboidal blocks with dimensions of 10 × 10 mm2 base area and 50 mm
in height were fabricated on a steel (AISI 304) build plate employing
beam currents and beam speeds in a range of 7 – 13 mA and 2500
– 3500 mm/s, respectively. A bidirectional meander scanning strategy with 90° rotation between successive layers was employed. In eachlayer, the exposure was accomplished block by block. For the sake ofbrevity, only the material processed using the set of parameters leading
to the highest density is detailed in the present paper. A summary ofthese processing parameters is given in Table 1. Using those values forthe acceleration voltage, beam current, beam speed, and hatch distance,
the resulting energy per unit area is calculated to be 1.71 J/mm2 .
2.2. Sample preparation and characterization
Plates of 1.5 mm thickness were machined along the build direction(BD) from the PBF-EB/M manufactured cuboids by electro-discharge
2C. Lauhoﬀ, T. Arold, A. Bolender et al.

Additive Manufacturing Letters 6 (2023) 100144

Fig. 1. SEM analysis of pre-alloyed Ti-Ta-Al powder material processed by EIGA: (a) SE image with superimposed color coding indicating the size of each powder
particle and (b) cross-section BSE image.

Fig. 2. EDS mappings of pre-alloyed Ti-Ta-Al powder particles showing the element distribution of (a) Ti, (b) Ta, and (c) Al.

Table 2
Chemical composition determined by EDS of the pre-alloyed Ti-Ta-Al powder
and PBF-EB/M processed material in both as-built and annealed condition. For
direct comparison, the nominal composition is listed to the top.

machining (EDM). In order to remove the EDM-aﬀected surface layer,
the plates were ground down to 5 μm grit size. The additively manufactured Ti-Ta-Al was investigated in two diﬀerent material states: asbuilt and after a heat treatment at 1200 °C for 21 h followed by waterquenching. All plates were encapsulated in quartz glass tubes under argon atmosphere to avoid oxidation.
For microstructure characterization, plates in both conditions werevibration-polished for 1 h with a 0.04 μm colloidal silica suspension (OPS NonDry, Struers). Etching was performed for 30 s using Kroll’s reagent.
Analysis of the process-induced defect distribution was conducted using optical microscopy (OM). For analysis of the chemical composition,
local segregations and crystallographic texture, scanning electron microscopy (SEM) including energy-dispersive X-ray spectroscopy (EDS)
and electron backscatter diﬀraction (EBSD) using a Zeiss Ultra Plus Gemini microscope was used. The SEM measurements were performed withan acceleration voltage of 20 – 30 kV. All microsections shown in thefollowing were recorded in a plane parallel to the lateral surfaces, i.e.
parallel to the build direction (BD), and depict representative areas fromthe center of the PBF-EB/M processed cuboids being not aﬀected by surface phenomena.
Phase analysis was carried out at the Deutsches ElektronenSynchrotron (DESY) at beamline P61A in Hamburg, Germany. Highenergy synchrotron diﬀraction allows to probe sample volumes of several mm3 and, thus, provides for a detailed high-resolution microstructure analysis. Polychromatic synchrotron radiation, covering an energy
range of 20 – 200 keV, and a Mirion high purity Germanium point detector with collimator-slit system were used. For further details on thesynchrotron beamline P61A, the reader is referred to Reference [38].

Composition (wt.-%)NominalPowderAs-built
Annealed

Ti

Ta

Al

70
68.3
69.4
68.0

25
25.9
25.5
27.6

5
5.8
5.1
4.4

tion has no detrimental eﬀect on the processability (powder ﬂowability
and printability) of the powder material at all. Structures with high relative density could be fabricated in this study (cf. Fig. 3 and see details
below).
A cross-section backscattered electron (BSE) micrograph of the powder particles is shown in Fig. 1b, indicating a dendritic microstructure.
EDS mappings of a polished particle cross-section in Fig. 2 reveal thechemistry of the powder material. Within the limits of EDS accuracy,
the quantitative results of the EDS analysis (Table 2) show that theoverall chemical composition of the particles is in good agreement with
the nominal composition of the Ti-25Ta-5Al alloy. As can be deduced
from the mappings, however, the powder particles feature dendritetype structures, i.e. chemical inhomogeneity associated to microsegregation during solidiﬁcation. The inter-dendritic phase is enriched intitanium (Fig. 2a). The inherent cooling rates during powder synthesis by EIGA are not suﬃciently high to fully prevent segregation processes and, thus, ﬁndings are fully in line with results recently reported
on gas-atomized Ti-42Nb [36] and tungsten containing Ti-Al powders[39]. Nonetheless, it will be shown below that the PBF-EB/M processed
Ti-Ta-Al does not comprise such dendritic microstructural features. Additively manufactured structures with very homogenous element distribution could be obtained using the pre-alloyed powder material for
fabrication.
For initial microstructure characterization after PBF-EB/M processing, OM was conducted. A representative micrograph is shown in

3. Results and discussion
Fig. 1 shows SEM images of the pre-alloyed Ti-Ta-Al powder in asatomized and sieved condition. Following the EIGA process, the particles
are highly spherical and feature smooth surfaces being free of satellites
and any defects like cracks. The powder material is fully deagglomerated as seen from the secondary electron (SE) image in Fig. 1a. It should
be noted that a slight residual amount of particles with diameters below
25 μm has remained in the sieved powder fraction. However, this frac3

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Additive Manufacturing Letters 6 (2023) 100144

responding diﬀraction patterns obtained at room temperature from theadditively manufactured material in both as-built and annealed condition. Please note that the latter state will be considered later. The
microstructure of the Ti-Ta-Al in as-built condition, i.e. without conducting a post-process heat treatment, mainly consists of the hexagonalclose-packed (hcp) 𝛼-phase. In addition, minor fractions of the bcc 𝛽phase are also present (cf. the low intensity diﬀraction peak at around87.5 keV). The lattice parameters are a𝛼 = 0.2930 nm / c𝛼 = 0.4694 nm
and a𝛽 = 0.3286 nm for the 𝛼- and 𝛽-phase, respectively. All parameters are in accordance with data reported in literature for binary Ti-Ta[45,46]. It should be mentioned that no detailed phase diagram is available for the alloy composition investigated. Tantalum and aluminum
have opposite inﬂuence on the phase formation of titanium alloys; tantalum is a 𝛽-stabilizer whereas aluminum is known to stabilize the 𝛼-phase[8,45]. Based on the empirical molybdenum equivalency ([Mo]eq) andaluminum equivalency ([Al]eq) [47–49], the Ti-25Ta-5Al alloy can beassessed in the context of 𝛽-phase stability, possible constituent phases
and resulting microstructure. With a [Mo]eq of 5.5 and an [Al]eq
of 5.0, the current Ti-25Ta-5Al alloy, to a large extent, is similar toTi-6Al-2Sn-4Zr-6Mo (Ti-6246) [47,48,50] as a heat-treatable 𝛼+𝛽 dualphase alloy. Consequently, the 𝛼+𝛽 dual-phase microstructure observed
in PBF-EB/M Ti-25Ta-5Al is supposed to be close to equilibrium state.
Due to the relative high processing temperatures of around 850 °C (cf.
Section 2) as well as the inherent slow cooling within the process chamber after melting the uppermost layer in PBF-EB/M, decomposition of
the high-temperature 𝛽-phase into the low-temperature 𝛼-phase takes
place. In addition, there is no evidence for the existence of the nonequilibrium martensitic phases 𝛼’ (hcp) and 𝛼’’ (orthorhombic) and thenon-equilibrium 𝜔-phase (hexagonal) in the as-built condition (Fig. 4),
again indicating near-equilibrium state [45].
The prior-𝛽 grain structure as well as the morphology and crystallographic orientation of the volume-dominant 𝛼-phase are clariﬁed by the
SEM EBSD analysis shown in Fig. 5. The 𝛼-phase was indexed with an
hcp crystal structure (P63 /mmc) and the lattice parameters determinedfrom the synchrotron diﬀraction pattern in Fig. 4, i.e. a𝛼 = 0.2930 nm/ c𝛼 = 0.4694 nm. The inverse pole ﬁgure (IPF) mapping and the corresponding IPF (Fig. 5b and c, respectively) reveal a well-known Widmanstätten patterned microstructure without an obvious global texture of the 𝛼 lamellae formed upon PBF-EB/M processing. In bulk
metallic materials, microstructures formed during AM, e.g. PBF-EB/M,
PBF-LB/M as well as directed energy deposition (DED), are often dominated by columnar grains oriented in BD due to epitaxial growth along
the main direction of heat ﬂow [1]. As can be deduced from the SE
image (Fig. 5a), the aspect ratios (length/width) of the prior-𝛽 grains

Fig. 3. Optical micrograph depicting the microstructure of PBF-EB/M processed
Ti-Ta-Al in the as-built condition. BD is from bottom to top.

Fig. 3 revealing a crack-free microstructure for the Ti-Ta-Al blocks(10 × 10 × 50 mm3 ) in the as-built condition. In AM processes, highresidual stresses may result from steep thermal gradients [40,41]. However, elevated base plate and built temperatures eﬀectively reduce
process-induced residual stresses [40,41], eventually hampering substantial crack formation as in present case. At this point, size eﬀects
are known to have a signiﬁcant inﬂuence on cracking in PBF-LB/M processes, however, PBF-EB/M as a hot-bed AM process does not suﬀer fromsuch issues to the same extent. Thus, it is expected that the ﬁndings discussed and presented here will be transferable to real components. Suchassessment, however, is out of the scope of present work and will be thefocus of follow-up studies. Only a small contribution of porosity is visible after processing. Using ImageJ software, a relative density of 99.86%
has been determined from a series of optical micrographs (not shown).
Pore sizes of up to 157 μm (cf. Fig. 3) and an average pore sphericity of
0.96, being very close to a spherical shape, were found. The coincidenceof high sphericity and small diameters of pores points at gas porosity induced from gas entrapment in the initial powder material, while thelarger pores (cf. defect in the middle of Fig. 3) are likely resulting from
keyholing eﬀects [42–44]. Lack of fusion porosity [42], however, hasbeen avoided with the set of processing parameters used in the present
work.
In order to shed light on the phase composition of the PBF-EB/M
processed Ti-Ta-Al, structure identiﬁcation was conducted employing
high-energy synchrotron diﬀraction experiments. Fig. 4 shows the cor-

Fig. 4. Synchrotron diﬀraction patterns of PBF-EB/M processed Ti-Ta-Al in as-built and annealed condition. Additionaldiﬀraction peaks of the elements Ta, W and Pb are visible, arising from conventional ﬂuorescence (Ta) and beamline components (W, Pb). Note, the ﬂuorescence peaks of Ti and Al areout of the displayed energy range. The OM image in the inset shows the martensitic microstructure after annealing andetching.

4

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Additive Manufacturing Letters 6 (2023) 100144

Fig. 5. SEM EBSD analysis of PBF-EB/M processed Ti-Ta-Al in the as-built condition: (a) SE image, (b) IPF mapping, and (c) IPF; (b) and (c) illustrate the crystallographic orientation of the 𝛼-phase with respect to BD. The reference coordinate system showing BD in horizontal direction and the color-coded standard triangle for
the IPF mapping are shown to the lower right.

Fig. 6. SEM EDS analysis of PBF-EB/M processed Ti-Ta-Al in the (a-d) as-built and (e-h) annealed condition: (a,e) SE images and (b-d) and (f-h) corresponding EDS
mappings. BD is from left to right.

are > 1. Thus, in accordance with grain structures reported in otherstudies for additively manufactured Ti-Ta alloys [31,51,52], a clear tendency to columnar prior-𝛽 grain formation is also visible in the present
study. However, it has to be noted that microstructures after PBF-EB/M
processing can signiﬁcantly diﬀer from PBF-LB/M microstructures interms of the phase composition. The inherent high cooling rates of the
PBF-LB/M process eﬀectively hamper diﬀusion-controlled processes
and, thus, in situ quenched-in non-equilibrium martensitic phases are often reported in PBF-LB/M fabricated 𝛼 and 𝛼+𝛽 Ti alloys [13,31,34,53].
In contrast, (near)-equilibrium phases are typically seen upon PBF-EB/M
as rationalized before [13,54].
Upon formation of the 𝛼-phase within the parent 𝛽-phase by a solidstate transformation, the 𝛼-phase enriches in 𝛼-stabilizer and depletesin 𝛽-stabilizer, and vice versa [8]. Beside decomposition, aluminum andtantalum are also known to form intermetallic compounds such as Al3 Ta
and Al69 Ta39 [55,56]. However, the synchrotron results (Fig. 4) provide for clear evidence that these intermetallic phases have not formedin the alloy system under consideration, which is also perfectly in linewith detailed microstructure analysis conducted on conventionally processed Ti-Ta-Al alloys in the past [57]. The alloying element partitioning
observed in present work is conﬁrmed by EDS results for the Ti-Ta-Al
material in the as-built condition (Fig. 6b-d). While in the present study
the 𝛼-phase (dark areas in Fig. 6a) is enriched in aluminum up to 6.5%
(77.7Ti-15.8Ta-6.5Al), a chemical composition of 61.3Ti-34.1Ta-4.6Al
with pronounced enrichment in tantalum up to 34.1% is found for the
𝛽-phase (bright areas in Fig. 6a). To better assess the overall chemical
homogeneity of the PBF-EB/M processed Ti-Ta-Al, a thermal treatment
at 1200 °C for 21 h has been conducted. After annealing in the single
𝛽-phase region and subsequent water quenching (suppressing diﬀusion
activities eﬀectively), a non-equilibrium, fully martensitic microstruc-ture consisting of the orthorhombic 𝛼’’-phase evolved as can be seenfrom the diﬀractogram and the optical micrograph in Fig. 4. This isakin to the presence of 𝛼’’ martensite in Ti-6Al-2Sn-4Zr-6Mo [48] additively manufactured by PBF-LB/M, where a much higher cooling rate isoften achieved compared to PBF-EB/M. The lattice parameters are determined to be a𝛼’’ = 0.3083 nm, b𝛼’’ = 0.4948 nm, and c𝛼’’ = 0.4587 nm.
Following both PBF-EB/M processing (as-built condition) and annealing
treatment, a slight decrease in the aluminum content can be observed(on the average, global scale) due to evaporation when compared with
the initial powder feedstock material (Table 2). However, most importantly the overall element distribution signiﬁcantly changed upon annealing. As is evident from the EDS mappings in Fig. 6f-h, the initialdendritic microstructure of the powder feedstock material (cf. Fig. 1band 2) completely vanishes. Accordingly, the constituents, i.e. titanium,
tantalum as well as aluminum, are dispersed homogeneously on a submicron scale in the annealed and subsequently quenched condition. Insummary, despite obviously tremendous diﬀerences in the alloying element densities and melting points, the use of pre-alloyed Ti-Ta based
feedstock material, including even low-melting and light elements suchas aluminum, allows to overcome current limitations using mixed elemental powders [31,33,34,51,52]. In this regard, the pre-alloyed Ti-TaAl powder allows direct fabrication of bulk metallic components featuring microstructures with excellent chemical homogeneity and concomitantly highest density.
4. Conclusions
In the present study, an Al-modiﬁed Ti-Ta alloy was successfully additively manufactured by electron beam powder bed fusion (PBF-EB/M)
using novel pre-alloyed feedstock material. In order to investigate defect5C. Lauhoﬀ, T. Arold, A. Bolender et al.

Additive Manufacturing Letters 6 (2023) 100144

populations, the chemistry, crystallographic texture, and phase compositions, detailed microstructure analysis was conducted employing
optical microscopy (OM), scanning electron microscopy (SEM) including energy-dispersive X-ray spectroscopy (EDS) and electron backscatter
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- Following electrode induction melting inert gas atomization (EIGA),
pre-alloyed feedstock material with highly spherical powder particles is obtained. Slight chemical inhomogeneities, i.e. Ta-rich dendritic segregations embedded in a Ti-rich inter-dendritic phase, arefound within the particles.
- Ti-Ta-Al bulk structures with densities > 99.86% are fabricated using
PBF-EB/M technique. Residual pores are likely resulting from gas
entrapments and keyholing eﬀects.
- An 𝛼+𝛽 dual-phase microstructure with columnar prior-𝛽 grain
structure is observed in the as-built condition. The volume-dominant𝛼-phase is enriched in Ti and Al, featuring a lamellar morphology
with no preferred crystallographic orientation.
- After a post-process thermal treatment at 1200 °C for 21 h in argon atmosphere followed by water quenching, a non-equilibrium,
fully martensitic microstructure (𝛼’’-phase) is present. In this postprocessed condition, the Ti-Ta-Al bulk material is characterized bya very homogenous distribution of the alloying elements.

Declaration of Competing Interest
The authors declare that they have no known competing ﬁnancial
interests or personal relationships that could have appeared to inﬂuence
the work reported in this paper.
Data availability
Data will be made available on request.

Acknowledgments
The authors gratefully acknowledge DESY (Hamburg, Germany),a member of the Helmholtz Association HGF and Helmholtz-Zentrum
Hereon, for the provision of experimental facilities. Parts of this researchwere carried out at PETRA III. Dr. Philipp Krooß and Dr. Guilherme
Abreu Faria are thanked for assistance with the experiments and in using P61A – WINE, respectively. Beamtime was allocated for proposal
ID: 11012003. C. Lauhoﬀ acknowledges funding by the Alexander von
Humboldt Foundation.

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