# Altering microstructure and enhancing mechanical properties during direct energy deposition of Ti-6Al-4V via in-process laser heat treatments

Kouider Abdellah Abdesselam, Steve Gaudez, Steven van Petegem, Veijo Honkimäki, Simon Hallais, Louis Cornet, Maxime Vallet, M V Upadhyay

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Kouider Abdellah Abdesselam, Steve Gaudez, Steven van Petegem, Veijo Honkimäki, Simon Hallais, et al. Altering microstructure and enhancing mechanical properties during direct energy deposition of Ti-6Al-4V via in-process laser heat treatments. Materials & Design, 2025, 254, pp.113997. 10.1016/j.matdes.2025.113997. hal-04746656v2

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# Altering microstructure and enhancing mechanical properties during direct energy deposition of Ti-6Al-4V via in-process laser heat treatments

K.A. Abdesselam<sup>a, ID</sup>, S. Gaudez<sup>a,1</sup>, S. Van Petegem<sup>b, ID</sup>, V. Honkimäki<sup>c</sup>, S. Hallais<sup>a</sup>, L. Cornet<sup>d</sup>, M. Vallet<sup>d,e, ID</sup>, M.V. Upadhyay<sup>a, ID, *</sup>

## ARTICLE INFO

Dataset link: 10.15151/ESRF-ES-847864546

**Keywords:**

Additive manufacturing

Ti-6Al-4V

Phase transformation

Synchrotron X-ray diffraction

Rietveld refinement

Scanning electron microscopy

## ABSTRACT

In-process laser-based heat treatments (LHTs) during additive manufacturing are explored as an alternative to post-process heat treatments such as annealing, to alter the microstructure and mechanical properties of Ti-6Al-4V. Three samples were manufactured via laser-based direct energy deposition with the laser operating at 300 W. One as-built sample without any LHTs served as the reference sample. For the remaining two, each deposited layer was subjected to an additional LHT just after its deposition; LHTs were performed at 100 W for one (LHT100) and 300 W for the other (LHT300) without powder addition. LHT100 showed an overall improvement in the strength/ductility trade-off. Whereas LHT300 resulted in a lower strength but a much higher ductility and toughness than the other two samples. Synchrotron X-ray diffraction analysis of as-built samples revealed an insignificant difference in the bcc (β) and combined hcp (martensite α', diffusively formed α<sub>d</sub> and massively transformed α<sub>m</sub>) phases between different samples. However, scanning electron microscopy uncovered a significant effect of LHTs in varying the fractions of different hcp phases between samples, which is the main cause behind mechanical property differences. These results demonstrate the significant potential of in-process LHTs to tailor alloy microstructures for achieving desired mechanical properties.

## 1. Introduction

Ti-6Al-4V is widely used in different industries such as aerospace, automotive, or medical, due to its high strength-to-weight ratio, high corrosion resistance, and good bio-compatibility [1–5]. In recent years, additive manufacturing (AM) of Ti-6Al-4V has gained importance due to the customization capabilities afforded by the layer-by-layer building approach used in AM as opposed to subtractive or mass-conserving conventional manufacturing techniques. AM of Ti-6Al-4V has been explored using processes such as laser powder bed fusion (LPBF), electron-beam powder bed fusion (EPBF), and laser directed energy deposition (LDED) [6–10] to produce geometrically complex parts that are tailored for a desired application with reduced material waste [6,11,12].

For all AM techniques, as-built Ti-6Al-4V parts generally exhibit higher tensile and ultimate strengths but poorer ductility in comparison

to their conventionally manufactured counterparts [7,13]. At the origin of these differences are the variations in the microstructural phenomena that include the constituent phases, grain sizes and morphology, texture, solute trapping, intergranular and intragranular strains, and dislocation density and arrangements.

Ti-6Al-4V exhibits two equilibrium phases: body-centered cubic (bcc) β that is stable above the β transus temperature (~ 980-1000 °C [14,15]) and hexagonal close-packed (hcp) α<sub>d</sub> below the β transus; the subscript d stands for diffusive. When Ti-6Al-4V is slowly cooled from above the β transus, β transforms diffusively into α<sub>d</sub> according to the so-called Burgers orientation relationship (OR): {0001}<sub>α</sub> || {110}<sub>β</sub> (⟨1120⟩<sub>α</sub> || ⟨111⟩<sub>β</sub>). This transformation involves long range atomic diffusion and redistribution of α<sub>d</sub> (aluminum, oxygen, carbon and nitrogen) and β (vanadium and iron) stabilizing elements [1,2,16]. The re-

<footnote><sup>a</sup> Laboratoire de Mécanique des Solides (LMS), École Polytechnique, Institut Polytechnique de Paris, CNRS UMR 7649, 91120 Palaiseau, France</footnote>

<footnote><sup>b</sup> Structure and Mechanics of Advanced Materials, PSI, Center for Photon Science, Forschungsstrasse 111, 5232 Villigen PSI, Switzerland</footnote>

<footnote><sup>c</sup> European Synchrotron Research Facility, Grenoble, France</footnote>

<footnote><sup>d</sup> Université Paris-Saclay, CentraleSupélec, ENS Paris-Saclay, CNRS, LMPS – Laboratoire de Mécanique Paris-Saclay, 91190 Gif-sur-Yvette, France</footnote>

<footnote><sup>e</sup> Université Paris-Saclay, CentraleSupélec, CNRS, SPMS – Laboratoire de Structures, Propriétés et Modélisation des Solides, 91190 Gif-sur-Yvette, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: manas.upadhyay@polytechnique.edu (M.V. Upadhyay).</footnote>

<footnote n="1">¹ Current address: Structure and Mechanics of Advanced Materials, PSI, Center for Photon Science, Forschungsstrasse 111, 5232 Villigen PSI, Switzerland.</footnote>
sulting microstructure features lathy and/or globular (more advanced stage of diffusion) $\alpha_d$ separated by thinner laths of $\beta$.

When Ti-6Al-4V is processed via AM, which involves rapid cooling rates ($10^3$ K/s - $10^6$ K/s), there is a high propensity for martensitic transformation, below the martensitic start temperature ($M_s \sim 780-900^\circ C$ [15,17,18]), of the $\beta$ phase into the metastable martensitic $\alpha'$ hcp phase exhibiting needle-like (acicular) morphology [6,19]. The martensitic transformation is a diffusionless (short-range atomic rearrangement) military phase transformation [16] that involves the shearing of the $\beta$ phase into the $\alpha'$ phase via a cooperative movement of atoms following the aforementioned Burger's OR [2,16,20]. At the time of its formation, $\alpha'$ has the same chemical composition as its parent $\beta$. Furthermore, the transformation strains and the incoherency strains at the $\beta-\alpha'$ interface trigger local plastic deformation that results in a high density of dislocations often observed in AM Ti-6Al-4V [9,21]. Thus, due to solid solution strengthening and high dislocation density, $\alpha'$ results in high strength in AM Ti-6Al-4V, but at the cost of low ductility [22–24]. However, solid-state thermal cycling (SSTC), also known as intrinsic heat treatment, due to subsequent addition of layers or post-process annealing can result in the decomposition of $\alpha'$ into $\alpha_d$ and $\beta$. The amount decomposed highly depends on the temperature amplitudes and heating/cooling rates reached during SSTC and a complete breakdown of $\alpha'$ into equilibrium compositions of $\alpha_d$ and $\beta$ may not be achieved during AM.

Another metastable phase that could occur due to a diffusionless transformation from $\beta$ is the so-called massive-$\alpha$ hcp phase, denoted as $\alpha_m$. In contrast to $\alpha'$ martensite, which forms due to a military transformation, $\alpha_m$ forms via a civilian transformation [16,25], which involves thermally activated interface migration and growth leading to the formation of irregularly-shaped $\alpha_m$ grains. Similar to $\alpha'$, at the time of nucleation, $\alpha_m$ grains have the same composition as their parent $\beta$ phase [16,26]. It is only their irregular shape, which differs significantly from $\alpha'$ and $\alpha_d$, that allows their identification and segmentation [25,26]. $\alpha_m$ was first reported to occur at cooling rates greater than 20 K/s and less than 410 K/s from above the $\beta$ transus by Ahmed et al. [22]. More recently, Nursyifaulkhair et al. [27] reported that the formation of $\alpha_m$ in LDED Ti-6Al-4V was the highest for cooling rates in between 4000 - 5000 K/s but lesser outside this range; in particular, higher cooling rates resulted in a preferential formation of $\alpha'$. Since both ranges can occur during all AM processes, the occurrence of $\alpha_m$ can be expected along with $\alpha'$, $\alpha_d$, and $\beta$ phases in AM Ti-6Al-4V alloys as reported in recent studies [28–32]. It has also been shown that the SSTC occurring due to layer addition during AM can result in the decomposition of $\alpha_m$ into $\alpha_d$ and $\beta$ [28–32]. Based on these observations, the current understanding is that $\alpha_m$ exhibits an intermediate strength/ductility response between $\alpha'$ and $\alpha_d$ [32].

In order to achieve a good strength/ductility trade-off of AM Ti-6Al-4V parts, heat treatments such as annealing, solution treatment, aging, tempering, among others [7,33–35], are often carried out to alter the microstructure and element distribution primarily through the decomposition of $\alpha'$ into $\alpha_d + \beta$. For example, Sabban et al. [36] were able to achieve a bimodal $\alpha_d + \beta$ globular microstructure on AM Ti-6Al-4V parts improving the ductility by 80% and the toughness by 66% by doing heat treatments over 24 h. Li et al. [37] adopted 2-hour multi-step heat treatments at different temperatures below the $\beta$ transus leading to a nearly equiaxed microstructure with a good ductility/ultimate tensile strength balance. De Formanoir et al. [38] investigated heat treatments at sub-transus temperature during 2 hours followed by water quenching to generate an $\alpha' + \alpha_d$ dual-phase microstructure, exhibiting an improvement in strength and ductility. Brown et al. [39] used in situ high energy X-ray diffraction (XRD) to study phase transformations, dislocation density, and internal stress evolutions during post-build heat treatments of LPBF and conventionally manufactured Ti-6Al-4V. They observed that recovery of dislocation density and stress relaxation began at 770 K, with $\beta$ phase transformation initiating around 913 K after recovery was complete.

Although effective, such homogeneous heat treatments require steps that are carried out outside the AM machine. Such steps can be time and energy effective when large number of samples are simultaneously heat treated. However, they can incur significant time and energy costs in the case of smaller batches with samples requiring dissimilar heat treatments; e.g., biomedical applications involving patient specific implants or specialized components in aerospace applications. Furthermore, such heat treatments are impractical when AM is used to repair partially damaged parts. In such cases, homogeneous heat treatments would undesirably alter not only the microstructure and mechanical properties of the portion repaired via AM but also that of the rest of undamaged part.

As an alternative to such external homogeneous heat treatments, localized heat treatments can be performed directly within the AM machine using the primary heat source (for building) or a secondary heat source to locally alter the microstructure. Such treatments can be done during the AM process (in-process) as well as after building (post-process). For instance, Esmaeilzadeh et al. [40] performed post-process laser scanning using their LPBF machine's laser in defocus mode to carry out multiple heat treatments on the same surface in order to initiate the decomposition of the $\alpha'$ phase into $\alpha_d + \beta$. Rietema et al. [41] explored the use of post-process laser scanning to perform layer-by-layer heat treatments on large surfaces of their LPBF samples. Meanwhile, the only instance of in-process laser scanning of Ti-6Al-4V using the primary (printing) laser has been reported for LPBF [42], but not for LDED.

In this work, the aim is to understand the impact of in-process laser heat treatments (LHTs), performed using the primary building laser during LDED of Ti-6Al-4V, on altering the microstructure and improving the mechanical properties (strength/ductility trade-off and toughness) without any post-processing steps. The contribution of each of the microstructural features and internal strains to the overall mechanical response is studied in detail.

## 2. Materials and methods

### 2.1. Material

The material investigated in this study is Ti-6Al-4V grade 23 ELI (extra low interstitial) obtained through inert gas atomization by Oerlikon Metco (USA). Powder composition (provided by the supplier) in weight percent (wt%) is the following: Ti - Bal, Al - 6.13, V - 4.13, Fe - 0.15, O - 0.06, C - 0.03, N - 0.01, H - 0.002. Granulometry indicates that the particles' $d_{10}$, $d_{50}$, and $d_{90}$ are 53 µm, 70 µm, and 110 µm, respectively. Chemical analysis using inductively coupled plasma spectroscopy and total organic carbon analysis (FiLAB, France) of the as-built reference sample revealed the following composition (wt%): Ti - Bal, Al - 5.73, V - 4.21, Fe - 0.16, O - 0.094, C - 0.013, N - 0.02, H - 0.0015.

### 2.2. Additive manufacturing and laser heat treatments

A custom-made miniature laser metal deposition (mini-LMD) machine was used for AM of the studied Ti-6Al-4V; details on the machine design can be found in [43]. It was used to print multiple 60-layered thin wall-shaped samples under an argon atmosphere using a bi-directional single-pass-per-layer printing strategy with the following parameters: laser power = 300 W, printing speed = 10 mm/s, powder flow rate = 2.7 g/min, powder transport gas flow = 0.05 L/min, meltpool developing gas flow = 5.0 L/min, and printing head layer increment = 0.24 mm. The mini-LMD's mobile printing head was used to print the samples while keeping the build plate stationary. The continuous-wave fiber laser (wavelength $\lambda$ = 1080 nm) is inclined at a $\pm 20^\circ$ angle with respect to the z direction in the y-z plane (coordinate system shown in Fig. 1). The powder flow is coaxial with the printing head and the z axis.

The samples were printed on the wide top surface of hot-rolled, annealed and sandblasted grade 5 Ti-6Al-4V substrates (Comptoir General
<figure><img src="image_3.png" /><figcaption>Fig. 1. Representation of the sample printing and additional LHT procedure as described in the text.</figcaption></figure>

des Métaux, France) of dimensions 40 (x) × 140 (y) × 10 (z) mm³. The bottom surface of the substrate was continuously cooled and maintained at ~ 25°C during printing with a closed circuit water and ethylene glycol-based cooling system to facilitate heat dissipation.

Three different samples were printed using the aforementioned strat-
egy. One sample that served as reference was manufactured without
any in-process heat treatment, and it is designated as the no heat treat-
ment (NHT) sample. The other two samples were subjected to in-process
LHTs without powder addition. For one of the two LHT samples, LHT
was performed at 100 W and it is henceforth designated as LHT100. For
the other sample, LHT was performed at 300 W and it is designated as
LHT300.

The following printing and LHT procedure was followed for the LHT100 and LHT300 samples: suppose that *n* layers have already been deposited and the deposition of the (*n* + 1)th and (*n* + 2)th layers is being undertaken as shown in Fig. 1:

(i) The (n + 1)th layer is deposited by moving the printing head along y (or -y) (Fig. 1a).

(ii) Next, for the LHTs, the printing head is then moved back to its position at the start of the printing of the (n + 1)th layer (Fig. 1b). Then, the LHTs are performed along the same direction y (or -y) and using the same printing head speed as the one used during the printing of the (n + 1)th layer (Fig. 1c). Following the LHT, the laser is rotated by 180° (Fig. 1d). This entire step takes a total of 52 s. Meanwhile, for the NHT sample, this step is skipped but a laser rotation of 180° together with a dwell time is imposed for a total duration of 52 s (Fig. 1d).

(iii) Then, for all samples, the (n + 2)th layer is deposited along the di-
rection opposite to that during the printing of the (n + 1)th layer,
i.e., -y (or y) (Fig. 1e).

(iv) Step (ii) is then repeated but lasering is performed in the direction opposite to the one in that step (Fig. 1f-h).

Steps (i) - (iv) ((a) - (h) in Fig. 1) are repeated for the next set of layers until the end of building. From here we can define the building direction (BD) along *z* and the printing direction (PD) along *y*. The final dimensions of NHT, LHT100 and LHT300 (Fig. 2a-c) samples are 1.81 (x) × 94.2 (y) × 14.8 (z) mm³, 1.77 (x) × 95.3 (y) × 16.1 (z) mm³, 1.73 (x) × 93.9 (y) × 15.4 (z) mm³ respectively. Note that the LHT100 and LHT300 samples exhibit a height higher than the NHT sample, which is counter-

intuitive because the LHTs should result in remelting and decrease in the
height of the sample. However, remelting of the few topmost deposited
layers caused by the LHTs results in the working distance between the
printing head nozzle and the top of the sample to vary between sam-
ples, despite the vertical increments of the focusing head remaining the
same for all samples (i.e. 0.24 mm). This change in the working distance
causes variations in the layer heights between samples.

2.3. Ex situ synchrotron X-ray diffraction

Synchrotron XRD was performed on the three as-built samples (Fig. 2) at room temperature at the ID31 beamline of the European Synchrotron Radiation Facility (Grenoble, France) to study their residual (as-built) state. A 75 keV monochromatic X-ray beam with a square cross-section and size (1.7 times the full width at half maximum) 230 µm in transmission mode was used to capture diffraction patterns on ±21 mm (175 data points along the y axis) from the center of the manufactured parts, as highlighted by the white rectangles in Fig. 2a-c. For each sample, the acquisitions were performed starting from above the top of the sample, i.e., in air, and scanning down until the substrate of the sample is reached. A Pilatus 3X CdTe 2M detector (1475 × 1679 pixels with a pixel size of 172 × 172 µm²) located at 0.908 m downstream from the samples was used to record the two-dimensional (2D) Debye-Scherrer patterns. Calibration of the detector position, angles, and tilts was performed using cerium dioxide (CeO₂) standard.

Representative 2D Debye-Scherrer patterns (averaged over the 175 data points per layer along the y axis and 10 layers) for each sample are shown in Fig. 3a-c; these patterns are representative of the microstructure in the gauge area of the tensile testing specimens. These patterns were analyzed in multiple different ways to obtain complementary information.

2.3.1. XRD data treatment for crystallographic texture analysis

The 2D Debye-Scherrer patterns have only been obtained at one rocking angle $\omega = 0°$ (about the *z* axis), which limits the information obtained from the XRD patterns to the *y* – *z* (PD-BD) plane; in particular, texture information in the *x* direction is not obtained from XRD analysis due to lack of patterns with the out-of-plane diffraction vector.

In order to study the influence of texture on the mechanical re-
sponse, the following procedure is followed. For all samples, portions
of the Debye-Scherrer rings were partially azimuthally integrated in the
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Images of the as-built (a) NHT, (b) LHT100, and (c) LHT300 samples. The translucent white rectangles represent the XRD investigated zones. The dashed lines delimit the specimens extracted for microstructural investigations. The black squares show the zones for scanning and transmission electron microscopy and electron back-scattered diffraction. Dogbone specimens extracted along *z* (BD) and *y* (PD) for uniaxial tensile testing are also represented.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** Representative 2D Debye-Scherrer patterns of as-built (a) NHT, (b) LHT100 and (c) LHT300 Ti-6Al-4V samples. Each of these 2D Debye-Scherrer patterns have been obtained by integrating data from 175 points in the ±21 mm zone along *y* axis (Fig. 2) per layer and 10 layers, which is representative of the microstructure in gauge area of the tensile testing samples. (d) An XRD line profile (*I* vs. 2θ) obtained from a single data point (i.e., from a zone of 230 µm × 230 µm) after integrating over the entire azimuthal angle and used for Rietveld refinement; the line profile is obtained from NHT and it is also representative of the LHT100 and LHT300 profiles. The experimental intensity, the one calculated from the Rietveld fit and their difference are also shown in (d). The black arrows highlight the faint bcc diffraction peaks. The ±10° portions of the 2D Debye-Scherrer rings used to obtain the lattice parameters along BD and PD, about the vertical and horizontal directions, are also represented by white dashed lines in (a).</figcaption></figure>
range ±10° around PD and BD to obtain the partially integrated I vs. 2θ line profiles along these respective directions; this azimuthal integration accounted for the intermodular distortions of the detector. For both directions, these line profiles were averaged over a set of 10 adjacent layers in the region corresponding to the gauge section of the dogbone tensile testing samples. Equation (1) was then used to calculate and study the ratios R of the maximum intensity of each diffraction peak from the integrated and normalized BD and PD line profiles with respect to the corresponding ones from the simulated random textured Ti-6Al-4V. The results are discussed in section 3.2.

$$R = \left( \max I_{(hkl)_{hcp}}^{\exp} / \max I_{(10\bar{1}1)_{hcp}}^{\exp} \right) / \left( \max I_{(hkl)_{hcp}}^{\text{sim}} / \max I_{(10\bar{1}1)_{hcp}}^{\text{sim}} \right) \quad (1)$$

where $\max I_{(hkl)_{hcp}}^{\exp}$ and $\max I_{(10\bar{1}1)_{hcp}}^{\exp}$ represent the maximum intensity of the experimental $(hkl)_{hcp}$ and $(10\bar{1}1)_{hcp}$ peaks, respectively, and $\max I_{(hkl)_{hcp}}^{\text{sim}}$ and $\max I_{(10\bar{1}1)_{hcp}}^{\text{sim}}$ represent the maximum intensity of the $(hkl)_{hcp}$ and $(10\bar{1}1)_{hcp}$ peaks, respectively, from the simulated random textured Ti-6Al-4V.

### 2.3.2. XRD data treatment for Rietveld refinement to study hcp and bcc phase fractions

In order to obtain the hcp and bcc phase content via Rietveld refinement, the 2D Debye-Scherrer patterns from the as-built samples were azimuthally integrated over the entire 360° angular range using PyFAI (Python library) [44] to obtain the fully azimuthally integrated diffraction (intensity I versus 2θ) line profiles; this azimuthal integration also considered the intermodular distortions of the detector. A fully azimuthally integrated line profile from a single measurement (230 µm × 230 µm) from NHT is shown in Fig. 3d. Rietveld refinement was performed on each such azimuthally integrated I vs. 2θ line profiles [45] from all samples via the FullProf suite [46] with a pseudo-Voigt function in sequential mode. The crystal structures of hcp (space group = P63/mmc) and bcc (space group = Im-3 m) titanium phases considered were obtained from the Crystallographic Open Database (CIF files 9008517, 9012924 [47]). The Rietveld fit of the line profile represented in Fig. 3d is also shown in the same figure. The “goodness of fit” and the residual differences between observed and calculated intensities [48,49] were used to evaluate the quality of the fits.

Note that Rietveld refinement is typically performed on fully azimuthally integrated line profiles in order to minimize any effect of texture. However, in cases of strong texture, even the full azimuthal integration of 2D Debye-Scherrer patterns at just one rocking angle may induce errors in phase fraction predictions.

To test the sensitivity of the predicted phase fractions to any in-plane texture, the representative (summed over 10 layers and 175 points per z position) 2D Debye-Scherrer patterns for all the full rings (10 distinguishable ones from 1010 to 2021 in Fig. 3a-c in each sample including both hcp and bcc rings, except the overlapping 1013hcp and 211bcc ones) are divided into 1° cakes and integrated (over the 1° azimuthal range and 37 Δθ intervals, determined from visual inspection to avoid overlap) and normalized following equation (2) for each ring to obtain the I(η) vs. η profiles (Supplementary material Figure S2); note that in the plot, η is assigned the lowest value of the 1° interval; note also that the division into 1° cakes encompasses significant fluctuations due to the small bin size as well as detector dead zones.

$$I_{hkl}^{\text{norm}}(\eta) = \frac{\sum_{i=1}^{37} I_{hkl}(\eta, \Delta\theta_i)}{\frac{1}{360} \sum_{\eta=0}^{359} \sum_{i=1}^{37} I_{hkl}(\eta, \Delta\theta_i)} \quad (2)$$

where $I_{hkl}^{\text{norm}}(\eta)$ represents the normalized intensity for a given (hkl) and $\eta$, $I_{hkl}(\eta, \Delta\theta_i)$ the measured intensity at a given $\eta$ and on an interval Δθ.

For those reflections showing significantly large amplitudes in I(η), Rietveld refinement at one arbitrarily selected measurement point is performed assuming a preferential orientation corresponding to these

reflections and compared to the one assuming random texture. The results of this comparison are presented in section 3.3.1.

### 2.3.3. XRD data treatment for lattice parameters

In order to obtain the lattice parameters along PD and BD, the partially azimuthally integrated (±10° portions around PD and BD) of the 2D Debye-Scherrer rings were used. Lattice parameters c and a of the hcp phase along PD and BD were obtained from the 0002 and 1120 peaks, respectively, based on [50].

### 2.3.4. XRD data treatment for Williamson-Hall analysis

The aforementioned partially azimuthally integrated diffraction line profiles along PD and BD were also used to obtain the Williamson-Hall (Δθ cos θ vs. 2sin θ/λ) plots in order to study the intragranular strain and crystallite size distributions in samples; Δθ is the full width at half maximum. The contribution of instrumental broadening was obtained from the CeO₂ standard. It was the same across all samples and it was removed from the Δθ of the sample (used in the plots) using equation (3) for a Gaussian profile.

$$\Delta\theta_{\text{sample}} = \sqrt{(\Delta\theta_{\text{total}})^2 - (\Delta\theta_{\text{instrumental}})^2} \quad (3)$$

## 2.4. Scanning and transmission electron microscopy

To investigate the microstructural features, samples were cut across the cross-section (x-z plane), and specimens of size (represented by the dashed black lines in Fig. 2): wall thickness (x) × 5 (y) × wall height (z) mm³ extracted from the middle of the as-built parts using a cutting machine (Struers). They were cold-mounted in an epoxy resin, mechanically polished with silicon carbide (SiC) abrasive sandpaper to P2400, and diamond paste polished down to 1 µm. Finishing step involved a 27-hour vibrometer polishing using a MasterMet solution (Bühler) combined with 10 mL of H₂O₂. A 10 nm layer of carbon was deposited to ensure a good conductivity during scanning electron microscopy (SEM) observations.

Electron back-scattered diffraction (EBSD) imaging of the cross-sections was performed in a FEI quanta 600 FEG SEM equipped with a symmetry detector (Oxford Instruments). EBSD maps of the cross-section were acquired at 15 keV with a spot size of 6.5 and a step size of 0.2 µm with an acquisition time of 5.32 ms and a 10% overlap between each 300 × 400 µm² field. In order to map the entire cross-section of each sample, 7 acquisitions were needed. Each cross-sectional map has an area of 1920 × 400 µm². For each sample, the EBSD acquisition time was 60 hours. EBSD analysis was performed using AZtec 6.0 SP1 software considering the hcp and bcc phases of titanium. The indexation rate prior to image processing was greater than 70% for all the samples. Post-processing was carried out using AZtec Crystal 3.0 software.

Next, energy-dispersive X-ray spectroscopy (EDS) investigations were conducted at Synergie 4 (France) using a flatquad detector (Bruker) in a JEOL IT 500 HR with an acquisition time of 5 hours at 6 keV with ~ 130,000 counts per second.

Following this, a mixture of 10 mL KOH, 5 mL H₂O₂, and 20 mL H₂O was used during 30 s to etch the polished specimens. Back-scattered electron (BSE) imaging was performed using the FEI quanta 600 FEG SEM equipped with a BSE detector (Thermo Fisher). For each cross-section of an etched specimen, 12 regions each of 227 × 331 µm² size were mapped at 0.2 µm resolution. Then, using the Fiji (ImageJ) [51] software, all images were converted to 8-bit (to perform thresholding) and the gray histogram of each acquisition was centered to 128. The 12 treated images from each specimen were then stitched together using Fiji with the procedure described in [52].

Finally, focused ion beam was used to extract thin-film lamellae of approximately 100 nm thick inside a FEI Helios Nanolab 660 dual-beam SEM, which were investigated using a Titan G2 TEM microscope featuring a Cs probe corrector, STEM detectors, and a SuperX detector with EDS capabilities to investigate the chemical heterogeneities and identify the β phase.
<figure><img src="image_6.png" /><figcaption>Fig. 4. (a) Engineering stress vs. engineering strain curves from representative tensile specimen extracted along BD (z) and PD (y). (b) Mean and standard deviation of the yield strength (YS) and ultimate tensile strength (UTS) as a function of maximum elongation prior to necking for NHT, LHT100, and LHT300 along BD and PD.</figcaption></figure>

## 2.5. Uniaxial tensile testing

For the tensile tests, dogbone specimens were machined from the manufactured parts vertically along *z* (BD) and horizontally along *y* (PD), and polished using grit SiC sandpaper up to P2400 to remove surface roughness. All vertical and horizontal dogbones have been extracted at the same height (*z* axis) with respect to the substrate. For NHT and LHT300 samples, one side of the three dogbones was extracted in the Ti-6Al-4V substrate along *z*. The gauge region of each sample had dimensions of 0.98 (*x*) × 2.25 (*y*) × 4.5 (*z*) mm³ to respect a length-to-width ratio of 2.

To ensure representativeness of the results, three specimens per direction were extracted from each sample as shown in Fig. 2 and tested using an in-house built uniaxial tensile testing machine with a strain rate of 2 µm/s and a 5 kN force cell. A Sill Optics correctal camera and marker tracking were used to measure the deformation during tensile testing. Following failure, secondary electron (SE) imaging was performed with the FEI Quanta 600 FEG SEM to image the fracture surface of the specimens.

## 3. Results

Three samples were built using the printing and LHT strategies described in section 2.2: NHT (without any heat treatment), LHT100 and LHT300 (with in-process laser heat treatment at 100 W and 300 W, respectively, after each layer deposited). *Ex situ* synchrotron XRD was performed on each as-built sample prior to extracting them from the substrate. Then, from the center of the samples, slivers were machined out for microstructure analysis (see section 2.4 for preparation) in the SEM. Finally, dogbone shaped samples were extracted for uniaxial tensile testing (see section 2.5 for preparation). The results of these tensile tests are first presented.

### 3.1. Effect of LHT on the mechanical response

Two sets of three dogbone shaped specimens were extracted from each NHT, LHT100, and LHT300; one set was along BD (*z*) and the second set was along PD (*y*). Fig. 4 illustrates the representative engineering stress-strain curves and the yield and ultimate tensile strengths as a function of maximum elongation prior to necking from the two sets for these samples. As shown in Fig. 5a and Supplementary material Figure S1, all specimens tested satisfied the Considère criterion and failed in the gauge area. Furthermore, fractography analysis presented in Fig. 5b reveals typical ductile fracture features characterized by a homogeneous distribution of dimples on the fracture surface; no abnormal features such as pores or cracks were detected on the fracture surface of any specimen. The results shown in Fig. 5 indicate that the full deformation potential of the material was reached prior to failure. Based on this understanding, a deeper analysis of the results can be undertaken. Table 1 shows the statistics of the Young's modulus and modulus of toughness for all the specimens.

For each sample, the mean values of the Young's moduli are not significantly different between PD and BD (the standard deviation of the Young's moduli is such that there is practically no difference between and within each sample) indicating that the samples are at least transversely isotropic in the *y–z* plane from a macroscale elasticity standpoint. Therefore, the additional LHTs do not impact the Young's modulus.

The representative tensile test curves, presented in Fig. 4a, exhibit a significant difference in the yield strength (at 0.2% plastic strain), ultimate tensile strength, and maximum elongation prior to necking between the three samples along both directions. Along BD, NHT, LHT100, and LHT300 exhibit the highest, intermediate and lowest, respectively, yield and ultimate tensile strengths. This trend is reversed for the maximum elongation prior to necking. Meanwhile, along PD, LHT100 and
<figure><img src="image_7.png" /><figcaption>Fig. 5. (a) Tangent modulus of the true stress-strain curves as a function of true stress for the NHT, LHT100, and LHT300 samples. The solid black line represents the Considère criterion. (b) Fracture surfaces of NHT, LHT100, and LHT300 along *z* (BD) and *y* (PD).</figcaption></figure>

<caption>**Table 1**</caption>

Mean and standard deviation of Young's modulus (E) and modulus of toughness of NHT, LHT100, and LHT300. Percentage difference of the mean values of LHT100 and LHT300 with respect to NHT.

<table><thead><tr><th></th><th></th><th>NHT</th><th>LHT100</th><th>% diff. with NHT</th><th>LHT300</th><th>% diff. with NHT</th></tr></thead><tbody><tr><td rowspan="2">E (GPa)</td><td>BD</td><td>130 ± 4</td><td>121 ± 4</td><td>-6.9</td><td>125 ± 7</td><td>-3.8</td></tr><tr><td>PD</td><td>126 ± 6</td><td>128 ± 2</td><td>1.6</td><td>126 ± 0.6</td><td>0</td></tr><tr><td rowspan="2">Toughness (MPa)</td><td>BD</td><td>107 ± 12</td><td>116 ± 22</td><td>8.4</td><td>150 ± 13</td><td>40.2</td></tr><tr><td>PD</td><td>54 ± 22</td><td>72 ± 9.4</td><td>33.3</td><td>96 ± 60</td><td>77.8</td></tr></tbody></table>

NHT exhibit similar yield and ultimate tensile strengths and higher than those of LHT300. However, interestingly, the trend in the maximum elongation prior to necking is the same as that along BD. Therefore, along PD, LHT100 demonstrates a better strength/ductility trade-off than NHT. For both BD and PD, NHT, LHT100, and LHT300 demonstrate the lowest, intermediate and highest toughness. However, along PD, the standard deviation in toughness of LHT300 is very high.

Collecting these results, it is clear that LHT100 results in better strength than LHT300 but the latter results in better ductility and toughness. In comparison to NHT, LHT100 also results in an overall improvement in the strength/ductility trade-off (at least along PD) and toughness. Meanwhile, LHT300 results in a reduction in strength but a significant increase in ductility and toughness in comparison to NHT.

These differences in the mechanical response must arise from the combined effect of microstructural features and internal (intergranular and intragranular) strains. To understand these contributions, a combination of synchrotron XRD, SEM and TEM analyses was conducted on the as-built samples.

## 3.2. Crystallographic texture

For metals that exhibit strong elastic and plastic anisotropy at the single crystal level, medium or strong crystallographic textures can significantly affect the macroscopic mechanical response.

With respect to single crystal elastic anisotropy, the Zener anisotropic factor for Ti-6Al-4V has approximately a 10% difference with respect to the isotropic one [53]. Thus, even if there were a medium or strong crystallographic texture, it would have a minor effect on the directional macroscopic elastic response. Indeed, there are negligible differences between the elastic moduli along BD and PD for all samples, which indicates the presence of a mild or medium texture, however, it is not possible to conclude this based solely on the macroscopic elastic response.

The single crystal plastic response is anisotropic due to a finite number of slip modes and systems available to accommodate plasticity. Hcp phases in both conventionally manufactured and AM Ti-6Al-4V have prismatic, basal and pyramidal slip systems available to accommodate
<caption>**Table 2**</caption>

R values for the 10 diffraction peaks for NHT, LHT100, and LHT300. Peak 8 has contribution from (1013)$_{hcp}$ and from (211)$_{bcc}$.

<table><thead><tr><th></th><th></th><th>(101̄0)<sub>hcp</sub></th><th>(0002)<sub>hcp</sub></th><th>(110)<sub>bcc</sub></th><th>(101̄1)<sub>hcp</sub></th><th>(101̄2)<sub>hcp</sub></th></tr></thead><tbody><tr><td rowspan="3">NHT</td><td>BD</td><td>1.34</td><td>1.10</td><td>1.09</td><td>1.00</td><td>1.67</td></tr><tr><td>PD</td><td>1.73</td><td>1.03</td><td>1.06</td><td>1.00</td><td>1.44</td></tr><tr><td>BD</td><td>1.81</td><td>1.51</td><td>1.09</td><td>1.00</td><td>1.62</td></tr><tr><td rowspan="3">LHT100</td><td>PD</td><td>2.76</td><td>1.02</td><td>1.18</td><td>1.00</td><td>1.70</td></tr><tr><td>BD</td><td>1.66</td><td>5.95</td><td>1.82</td><td>1.00</td><td>1.48</td></tr><tr><td>PD</td><td>1.22</td><td>1.48</td><td>1.10</td><td>1.00</td><td>1.09</td></tr><tr><td></td><td></td><td>(200)<sub>bcc</sub></td><td>(112̄0)<sub>hcp</sub></td><td>(1013)<sub>hcp</sub> + (211)<sub>bcc</sub></td><td>(112̄2)<sub>hcp</sub></td><td>(202̄0)<sub>hcp</sub></td></tr><tr><td rowspan="3">NHT</td><td>BD</td><td>1.36</td><td>0.27</td><td>0.73</td><td>0.61</td><td>1.13</td></tr><tr><td>PD</td><td>1.17</td><td>2.25</td><td>1.57</td><td>1.47</td><td>1.09</td></tr><tr><td>BD</td><td>1.85</td><td>0.40</td><td>0.65</td><td>0.65</td><td>0.58</td></tr><tr><td rowspan="3">LHT100</td><td>PD</td><td>1.05</td><td>1.95</td><td>1.99</td><td>1.91</td><td>1.50</td></tr><tr><td>BD</td><td>1.24</td><td>0.92</td><td>1.24</td><td>0.57</td><td>0.57</td></tr><tr><td>PD</td><td>1.02</td><td>0.93</td><td>1.09</td><td>1.09</td><td>1.09</td></tr></tbody></table>

plastic deformation. Typically, if the critical resolved shear stress of prismatic slip systems is taken to be $\tau_c$, then that of basal and pyramidal slip systems is ~1.13$\tau_c$ and $3\tau_c$, respectively [53]; these values may differ slightly depending on solid solution strength differences [54] between conventionally manufactured and AM Ti-6Al-4V. Furthermore, the c/a ratio of the equilibrium hcp phase for Ti-6Al-4V is slightly less than 1.6, which in general is less than the ideal one and this facilitates prismatic slip. Thus, if there were to be a strong texture in NHT, LHT100 and LHT300, then it would manifest as a significant difference in the yield and ultimate tensile strengths between their respective BD and PD responses, which is clearly not the case as seen in Fig. 4; even for the cases of LHT100 and LHT300 where yield strength differences between their respective PD and BD are the largest, there is an overlap in their standard deviations for both cases. Note that for all the samples the strains at the onset of necking are less than 0.06, clearly indicating that negligible texture evolution occurs during plastic deformation prior to necking. Therefore, based on the overall mechanical response, one can reasonably argue that none of the samples exhibit a strong texture in the PD-BD plane.

This argument is further supported by the analysis of the synchrotron XRD data collected from the as-built samples prior to the extraction of the dogbone shaped samples. For 10 discernible full Debye-Scherrer patterns, equation (1) is used to calculate the ratio *R* along PD and BD for the three samples. This ratio compares the azimuthally integrated (in ±10° range around PD and BD) and normalized intensities with respect to the corresponding ones obtained from random texture profiles. It is tabulated in Table 2.

Except for the hcp reflection 0002 along BD for LHT300, all other hcp and bcc reflections show an *R* value that is close (most within a factor of 2 and only two between a factor of 2 and 3) to that of the random texture along both BD and PD. Below are the arguments as to why *R* = 5.95 for the 0002 reflection does not imply a strong texture in this direction. Loading a grain normal to its basal plane (corresponding to loading along the 0002 direction) will activate prismatic slip, which has the lowest critical resolved shear stress among all slip modes. This would imply that LHT300 shows a lower yield strength along BD than the other two samples, which it indeed does, and also with respect to its own PD direction that shows an *R* value of 1.48, which is the case but only slightly; the mean yield strength of LHT300 along PD is lower by ~7% than along BD, however, the standard deviations along both directions overlap.

Collecting all this information, and taking into account that the Debye-Scherrer patterns are obtained only at one rocking angle $\omega$ = 0° (see section 2.3.1), which may result in an overestimation of the *R* value for a particular direction, the following two conclusions can be drawn: (i) texture has a mild effect on the mechanical response between BD and PD for all samples, but (ii) the difference in the mechanical response be-

tween NHT, LHT100 and LHT300 certainly cannot be attributed to these differences in crystallographic texture.

### 3.3. Quantifying constituent phase fractions

#### 3.3.1. Rietveld refinement to separate hcp and bcc fractions

In this section, the synchrotron XRD measurements are used to quantify phase fractions of hcp and bcc phases in the as-built samples via Rietveld refinement. The hcp phase could include a combination of $\alpha'$, $\alpha_d$, and $\alpha_m$. However, it is not possible to distinguish these phases from our XRD measurements due to the resolution at which these measurements have been performed; a problem that was also encountered in [55].

Meanwhile, the only known bcc phase in Ti-6Al-4V at room temperature is the $\beta$ phase. In the context of AM, it can either be the retained and/or the transformed $\beta$ phase. Therefore, synchrotron XRD measurements are used only to separate the $\beta$ phase and the collective set of hcp phases.

In the previous section, the texture was found to have a mild effect on the mechanical response along BD and PD. However, it may affect the phase fraction predictions of Rietveld refinement. To assess this, prior to performing Rietveld refinement, radially integrated and normalized intensity $I(\eta)$ vs. $\eta$ profiles are plotted using equation (2) and studied (see Supplementary material Figure S2) for the 10 distinguishable full rings in Fig. 3.

The $I(\eta)$ vs. $\eta$ evolutions of the hcp reflections 101̄0, 101̄1, 112̄0, 112̄2 and 202̄1 are found to be nearly the same for all samples with the maximum intensity ratio with respect to the average being 5 for 112̄0. Meanwhile, for the hcp reflections 0002 and 101̄2, NHT and LHT100 show nearly the same evolutions, however, LHT300 shows significant differences with respect to NHT and LHT100, with peaks of amplitudes ~6 occurring at multiple $\eta$. For the bcc reflections 110 and 200, NHT and LHT100 show lower amplitudes (maxima always less than 3.5) and nearly the same evolutions, but LHT300 once again shows significant differences in the response and slightly higher global maxima (up to 5).

Since the maxima in $I(\eta)$ vs. $\eta$ are generally low for NHT and LHT100, it can be concluded that the texture is mild in the entire PD-BD plane for the hcp phase in these two samples. In order to test whether the larger amplitudes for LHT300 have any significant effect on the predictions of Rietveld refinement, three types of refinements were performed using the procedure outlined in section 2.3.2: one assuming random texture, one assuming that 0002 was the predominant phase and the final one assuming 101̄2 was the predominant phase. In all three cases, the predicted phase fractions were found to be within a difference of 1%, which lies well within the experimental and Rietveld refinement errors.

Meanwhile, for the bcc phase, since there are only two distinguishable rings, it is harder to identify from XRD patterns whether there is a strong or mild texture in this phase. There are differences between
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** Evolution of the β phase fraction (which is equal to 100 - hcp phase fraction in %) as a function of distance from the top of the as-built samples along –z for NHT, LHT100, and LHT300. At each z position, the statistics are computed from 175 measurements along PD in the ±21 mm central region in Fig. 2.</figcaption></figure>

LHT300 on one hand and NHT and LHT100 on the other hand, with *I*(*η*) amplitudes for bcc reflections in the former reaching similar values as those for some of its hcp reflections. Yet, a similar Rietveld analysis as before but for the bcc phase reveals the same negligible difference in the predicted phase fractions.

These results clearly indicate that for all samples, the texture in the PD-BD plane does not affect the phase fractions obtained from Rietveld refinement. In what follows, only the phase fractions from Rietveld refinement performed without any preferential orientation are reported and discussed.

The distribution of the predicted bcc phase fraction (which is equal to 100 - hcp phase fraction in %) as a function of the location in the sample along –z is plotted in Fig. 6. All three samples show nearly the same distribution of the phases along their height. The mean (and standard deviation) values over the entire NHT, LHT100, and LHT300 samples for the hcp phase are 97% (± 1.0%), 96.9% (± 1.1%) and 96.3% (± 1.5%) respectively, and for the bcc phase are 3.0% (± 1.0%), 3.1% (± 1.0%) and 3.7% (± 1.5%), respectively. Phase fraction fluctuations within each plot are influenced by the surface roughness of the samples and the low intensity of the most intense β reflection, i.e. 110, which appears between two very close 0002 and 1011 hcp reflections. Note that the standard deviation for each sample falls into the minimum error of the Rietveld refinement. Nevertheless, the differences in the mean values and standard deviations of each phase between the samples are small. Therefore, it can be concluded that additional LHT does not result in any significant differences in the hcp and bcc phase fractions between the samples.

In all samples, the amount of β is between 3.0% and 3.7%. These values are closer to the equilibrium β content reported, for example, in as-cast Ti-6Al-4V [1,20,39], in comparison to LPBF Ti-6Al-4V, where some works have reported no detectable β [40,56]. Slower cooling rates are typically encountered in LDED in comparison to LPBF. It is therefore unsurprising that the former allows for more β phase formation in comparison to the latter. Nevertheless, such small fractions cannot solely explain the significant differences in the mechanical response. Furthermore, a variation in the phase fraction of β (even a doubling or tripling at such low amounts), which is a soft phase, is not expected to significantly affect the mechanical response; this has been shown in the case of retained austenite and martensite in a steel [57]. More significant contributions are likely coming from other microstructural features.

### 3.3.2. Identifying hcp phase constituents: microstructural investigations

SEM analysis was performed on the cross-sections (x-z plane) of the three as-built samples to identify the hcp phase constituents. As mentioned at the beginning of section 3, the samples have different heights and to facilitate comparison, SEM analysis is performed in the height range between 2.35 - 2.75 mm above the substrate (represented by the

black rectangles in Fig. 2) and across the thickness of all the samples; note that the analyses were restricted to this zone because performing EBSD at a resolution of 0.2 µm, which was deemed to be the minimum necessary to resolve α'/α<sub>d</sub> phases, took a significant amount of time (60 hours) per sample. This height range also corresponds to a zone sufficiently away from the substrate such that the chemical composition of the material is not affected by proximity to the substrate. In addition, this region is at a height much lower than the sample height, which allows for studying the impact of multiple (SSTCs) on the sample.

Fig. 7 shows the EBSD (IPF z coloring) and higher magnification views of EBSD along with BSE maps of the studied portion of the cross-section for the three samples. The β phase, whose volume fraction in each sample is less than 3.7% (section 3.3), could not be identified in any map (see Supplementary material Figure S3), indicating that the grain size of β is smaller than the EBSD resolution (0.2 µm); this result is confirmed by TEM analysis presented in the Supplementary material Figure S4. The EBSD maps allow identifying grains and analyzing them individually. Meanwhile, the zoomed-in BSE images in Fig. 7 allow a clear identification of the different hcp phases through the contrast within each image of a specimen. These differences arise due to the combined contribution of the segregated heavy (vanadium and iron) and light (aluminum and oxygen) elements, which, respectively, create a bright and dark contrast.

To test whether the different phases can be isolated based on their chemical composition, an EDS analysis was carried out in the SEM on a region in LHT300 that is representative of the three hcp phases in each sample. However, the results were inconclusive due to the difficulty of separating the hcp phases using elemental segregation of α<sub>d</sub> and β stabilizers. The images showed a homogeneous distribution of Al and V in the region tested (see Supplementary material Figure S5); therefore, the EDS imaging in SEM was not used further.

Therefore, the focus in the microstructural analysis is kept on the EBSD and BSE images. Four discernible populations of hcp grains can be identified: acicular grains (in two dimensions, and plates in three dimensions) with the highest intensity in BSE images, lath-shaped grains with medium intensity, globular grains with medium intensity and irregularly-shaped coarser grains with the lowest intensity in the bulk.

The acicular-shaped grains with the highest intensity belong to the martensitic α' phase. The BSE image reveals different sizes of these grains that are not visible in the EBSD image. According to Yang et al. [58], different sized needles of α' form during the initial stages of SSTC, with the largest forming during the first cooling down after deposition.

The lathy and globular shaped grains are typically from the α<sub>d</sub> phase that arises partly during the cooling down from above the β transus (some portion of the material may have lower cooling rates than those that yield the α' phase), and from the breakdown of α' phase into α<sub>d</sub> and β due to the subsequent SSTC. Finally, the irregularly-shaped coarser grains are identified as the α<sub>m</sub> phase. The α' and α<sub>d</sub> phases are rather homogeneously distributed across the microstructure, whereas clusters of α<sub>m</sub>, which are lower in proportion, are heterogeneously distributed.

Based on the morphology of α<sub>m</sub> grains from EBSD maps and the brightest intensity of α' in the BSE images, these phases are manually segmented to obtain comparable measures of their area fractions. These measures are then used to deduce the fraction of the combined globular and lathy α<sub>d</sub> phase. In what follows, the procedure to isolate each hcp phase is described:

* Due to the irregular shape of α<sub>m</sub> grains, there is no established procedure to segment them from images. Furthermore, due to a relatively large area studied, manual segmentation was deemed unfeasible. Instead, the following two-step selection approach that could be implemented in the EBSD software (Aztec) was devised: (i) Select grains that have an area greater than or equal to 40 µm² in the entire microstructure. (ii) Deselect grains that have length-to-width ratio (of an adjusted ellipse) greater than 5. The first criterion avoids selecting globular α<sub>d</sub>, which are smaller in size than the α<sub>m</sub> grains,
<figure><img src="image_9.png" /><figcaption>**Fig. 7.** (a, f, and k) EBSD maps with 0.2 µm resolution from the cross-section of the NHT, LHT100 and LHT300 samples. (b-e), (g-j) and (l-o) EBSD and BSE high magnification views (plain and dashed squares) of different zones in the investigated cross-section. The plain white square in (f) represents thin lamellae extraction zones from an α<sub>m</sub> grain. The white circle in (k) represents the investigated zone for EDS analysis shown in Supplementary material Figure S5. Color code in EBSD images is according to IPF *z* whose key is shown in (a). The coordinate system is also presented in (a). Both apply to all samples. (p) Grain size distribution for NHT, LHT100, and LHT300.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Percentage phase fractions of bcc and adjusted hcp phases obtained from Rietveld refinement in section 3.3.1 and SEM analysis in section 3.3.2.</figcaption></figure>

and the two criteria together avoid selecting acicular $\alpha'$ and lathy $\alpha_d$. The EBSD maps for the distributions of these phases are shown in Supplementary Figure S6. Due to a low amount of these phases in the studied EBSD cross-section, another cross-section (at height 1.95 - 2.35 mm along BD) was also studied to improve statistics (see Supplementary material Figure S7 for the EBSD cross-section and Supplementary material Figure S8 for the phase distribution). Note that the numbers used in the selection criteria are specific to our samples and may differ for samples fabricated using different printing parameters. Furthermore, while the accuracy of the extracted data may not be perfect, nevertheless they highlight the correct trend between samples; computing the error involved would require manual segmentation of each $\alpha_m$ grain, which is also erroneous (based on the investigator) and impractical as mentioned earlier.

* Next, the fraction of the $\alpha'$ phase was obtained from BSE images (Fig. 7). For each cross-sectional image, a manual grayscale thresholding was performed to select only the acicular phases, which appear as the brightest (highest intensity) features.

* Combining the contributions of $\alpha_m$, $\alpha'$, and $\beta$ (from section 2.3), the fraction of combined lathy and globular $\alpha_d$ can be deduced.

* Assuming that the area fractions, measured over a large area, are similar to the volume fractions, the adjusted volume fractions of all the identified hcp phases and the $\beta$ phase for each sample are computed.

Fig. 8 shows the proportions of each phase computed using the aforementioned criteria. Between the hcp phases of each sample, $\alpha_d$ has the highest phase fraction followed by $\alpha'$ and $\alpha_m$. Across samples, the $\alpha'$ fraction is the highest in NHT and lowest in LHT300, $\alpha_m$ is the highest in LHT100 and lowest in LHT300, and $\alpha_d$ is the highest in LHT300 and lowest in NHT.

The dominant hcp phase is $\alpha_d$ followed by $\alpha'$. The ratio of $\alpha'$ to $\alpha_m$ is 8.8, 4, and 10.1 in NHT, LHT100, and LHT300, respectively. Meanwhile, the ratio of $\alpha'$ to $\alpha_m$ is 32, 9.6, and 40.7 in the same order.

For an LDED process with the highest cooling rates that can reach up to $10^5$ K/s [59,60], the proportion of $\alpha'$ martensite appears to be uncharacteristically low. However, note that the microstructure investigations have been conducted on a layer that has been subjected to multiple SSTCs with decreasing temperature amplitudes and heating/cooling rates that should result in the decomposition of $\alpha'$ into $\alpha_d$ and $\beta$. The ratio of $\alpha'/\alpha_d$ is expected to increase with increasing height (along the BD) as has been consistently shown in the literature [17,61].

Finally, it is possible that some or all of the irregular shaped grains identified may not be $\alpha_m$ but could be $\alpha_d$ plates whose normal to the

wide surface is (nearly) perpendicular to the surface of the cross-sections studied. However, even in the extreme case where all these $\alpha_m$ grains turn out to be differently oriented $\alpha_d$, their added contribution to the $\alpha_d$ fractions would alter neither the trend nor the conclusions of this study.

### 3.4. Lattice parameter evolution and Williamson-Hall plots

The synchrotron XRD data used to obtain the hcp and bcc phase fractions in section 3.3.1 is also used to study the lattice parameter distribution and peak widths through the Williamson-Hall plot [62] in the as-built samples as described in section 2.3.

For the lattice parameters, instead of integrating the Debye-Scherrer rings over the entire azimuthal range, they are divided into angular portions of 10° each and the portions $\pm 10^\circ$ around BD and around PD are analyzed (Fig. 3a). Line plots of lattice parameters averaged per layer as a function of the distance from the top of the sample are shown in Fig. 9; a Butterworth low-pass filter was applied to the data to improve visualization. The plots reveal very small differences in magnitudes and evolution between samples in either direction. This similarity indicates that LHTs do not result in any significant differences in lattice parameters between samples.

The mean and standard deviation of a and c for each sample along both BD and PD are shown in Table 3. Their near perfect match implies that there are no major chemical and/or mechanical differences between the different as-built samples.

However, at any height, there is a significant difference in the residual lattice strains between PD and BD but this difference is practically equal across samples. This relative difference in PD and BD lattice strains is caused by the macroscopic residual stress state. Note that the XRD measurements were performed while the as-built walls were still attached to the substrate and the mechanical constraints of the substrate play a significant role in determining the residual stress state. It is well established that AM of thin-walls on large substrates result in the formation of tensile macroscopic residual stresses along PD, and nearly zero macroscopic residual stress along BD (due to traction-free conditions on the sample top). Assuming that the macroscopic residual stress along the incoming X-ray direction is zero (reasonable assumption since thin walls were printed), the residual tensile stress along PD results in a positive shift of lattice parameters along PD with respect to their elastic strain-free counterparts. Meanwhile, due to the Poisson's effect, a negative shift along BD is obtained with respect to the elastic strain-free lattice parameters; note that while it was not possible to measure strain-free lattice parameters during the synchrotron experiments, nevertheless the reasoning remains sound as has been shown in a recent work for the case of 316L stainless steel [63]. Extraction of dogbone samples will alter the macroscopic residual stresses, and their contribution to the shift caused in the lattice parameters. However, this alteration should be the same between samples because of two reasons: the lattice parameters along PD or BD are practically the same between samples, and the dimensions of the dogbones and their locations of extraction are also practically the same between samples.

It can therefore be deduced that residual intergranular strain distributions do not play any significant role in the variation of the mechanical response between NHT, LHT100 and LHT300.

Fig. 10 shows the Williamson-Hall [62] plot obtained from *I* vs. $2\theta$ plots along PD and BD using $\pm 10^\circ$ cakes for all three samples without instrumental broadening. It provides an estimate for the crystallite size averaged over the domain and intragranular lattice strain distributions arising from dislocation populations (both geometrically necessary and statistical i.e., Nye's dislocation density tensor is non-zero and zero, respectively). Apart from one point for the LHT300 sample with a larger spread along both PD and BD, there is a near perfect match between the samples in both directions; the large spread in LHT300 is occurring due to very noisy data that results in a poor fit. Nevertheless, further analysis of this plot is not necessary because this result indicates that the probed volume-averaged intragranular strain distri-
<figure><img src="image_11.png" /><figcaption>Fig. 9. Line plots of a and c lattice parameter of the (combined) hcp phase, averaged per layer as a function of distance from top of the sample for (a) BD (z axis) and (b) PD (y axis) for NHT, LHT100, and LHT300 samples. LHT300 exhibits the widest spread in standard deviation.</figcaption></figure>

<figure><img src="image_12.png" /><figcaption>Fig. 10. (a) and (b) Williamson-Hall plots along with (c) and (d) zooms with the standard deviation, from the diffraction line profiles extracted along (a, c) BD and (b, d) PD using ±10° cakes for all three samples without instrumental broadening. The key is presented in (a).</figcaption></figure>

butions, which arise from dislocations and crystallite size distributions are nearly the same for all samples. These results are also supported by the geometrically necessary dislocation density distributions discussed in section 3.5.

### 3.5. Grain size and geometrically necessary dislocation density distributions

Fig. 7p shows the cumulative distribution function plot for the grain sizes obtained from the EBSD maps of Fig. 7a, f and k for the three
<figure><img src="image_13.png" /><figcaption>**Fig. 11.** (a, d, and g) Geometrically necessary dislocation density maps, (b,c), (e,f) and (h,i) higher magnification views along with (j) distribution functions of NHT, LHT100, and LHT300.</figcaption></figure>

<caption>**Table 3**</caption>

a and c lattice parameters for both BD and PD for NHT, LHT100 and LHT300.

<table><caption>**Table 4**</caption><thead><tr><th></th><th>NHT (stdev)</th><th>LHT100 (stdev)</th><th>LHT300 (stdev)</th></tr></thead><tbody><tr><td>a<sub>PD</sub> (Å)</td><td>3.004 (±0.002)</td><td>3.004 (±0.0023)</td><td>3.005 (±0.0103)</td></tr><tr><td>a<sub>BD</sub> (Å)</td><td>2.992 (±0.0018)</td><td>2.992 (±0.0016)</td><td>2.995 (±0.019)</td></tr><tr><td>c<sub>PD</sub> (Å)</td><td>4.791 (±0.0052)</td><td>4.790 (±0.0048)</td><td>4.791 (±0.0080)</td></tr><tr><td>c<sub>BD</sub> (Å)</td><td>4.777 (±0.0044)</td><td>4.777 (±0.0038)</td><td>4.778 (±0.0082)</td></tr></tbody></table>

Grain sizes and GND density for NHT, LHT100 and LHT300.

<table><thead><tr><th></th><th>NHT (stdev)</th><th>LHT100 (stdev)</th><th>LHT300 (stdev)</th></tr></thead><tbody><tr><td>Grain size (µm)</td><td>1.75 (±1.27)</td><td>1.58 (±1.16)</td><td>1.48 (±1.04)</td></tr><tr><td>GND ×10<sup>14</sup> (m<sup>-2</sup>)</td><td>3.1 (±1.5)</td><td>3.4 (±1.7)</td><td>3.5 (±1.8)</td></tr></tbody></table>

samples. Table 4 shows the mean and standard deviations of the grain sizes. Given the overlap in the standard deviation between samples, the grain sizes are considered statistically the same. In order to confirm that the strength of the material is not very sensitive to such small differences in grain sizes, the Hall-Petch relationship proposed in [64] for primary α grains, whose tendency is most representative of the response observed in this work, is used: $\sigma_{Y,S}(\text{MPa}) = 230d^{-0.5} + 737$, where d is the grain size in (µm). Of particular importance is the term $230d^{-0.5}$. Using the mean grain sizes reported in Table 3, this term is 174 MPa, 183 MPa, and 189 MPa for NHT, LHT100 and LHT300, respectively. This result implies that, solely on the basis of grain size, LHT300 should result in the highest yield strength followed by LHT100 and NHT. However, the trend observed in Fig. 4 is the opposite. Furthermore, the difference between the highest and lowest magnitudes of the aforementioned term is only

16 MPa, which is an order of magnitude lower than the difference in yield strength between NHT and LHT300. Therefore, small differences in grain sizes are expected to play a minor (if not negligible) role in determining the mechanical response.

Another difference in mechanical response could arise from variations in the strengthening effect caused by different dislocation densities. In section 3.4, the Williamson-Hall plot from the three samples matched nearly perfectly, indicating that the contribution of intragranular strain distributions arising from all dislocation populations to the mechanical response is nearly the same. However, sub-populations of dislocations, specifically geometrically necessary and statistical dislocations, could vary and result in different contributions.

In order to study the role of geometrically necessary dislocations, their density is evaluated from the EBSD results shown in Fig. 7; note that this density magnitude is subject to the EBSD resolution used and yet it will give an indication on the contribution of this population. Fig. 11 shows the geometrically necessary dislocation density maps of the three samples along with their distribution functions.

Overall, there are minor (and possibly negligible) differences in the magnitudes and distributions of these densities between the samples. The mean and standard deviations of the densities over the studied domains are shown in Table 4. They demonstrate a near-perfect match between the samples. Since the strength of the material is proportional to the square root of the dislocation density, the contribution of geometrically necessary dislocation density to the difference in mechanical response between samples is negligible. Combined with the negligible difference in the Williamson-Hall plots, it can be concluded that the geometrically necessary and statistical populations of dislocations are similar between the samples.

However, there are significant differences in the geometrically necessary dislocation densities between $\alpha_m$ and the remaining two hcp phases
($\alpha' + \alpha_d$) in all samples; $\alpha_m$ consistently shows at least an order of magnitude lower densities. However, due to the relatively low fraction of $\alpha_m$ in all samples, and the understanding that this phase has an intermediate strength between $\alpha'$ and $\alpha_d$ [32], this difference in geometrically necessary dislocation densities should not play a significant role in determining the mechanical response; however, a deeper analysis (possibly involving nanoindentation studies) is required to corroborate this statement, which is beyond the scope of this work.

## 4. Discussion

Collecting all the evidence from synchrotron XRD and microstructure analyses performed in sections 3.2 - 3.5, the low phase fractions of $\alpha_m$ and the similarity in $\beta$ fractions, intergranular lattice strains, crystallite/grain sizes, intragranular lattice strains, geometrically necessary dislocation densities and texture between samples should have a minor, if not negligible, contribution to the variation in the mechanical response between them. Thus, by elimination, the possibilities are narrowed down to the fractions of $\alpha'$ and $\alpha_d$ (see Fig. 8), and possible differences in solid trapping, which has not yet been addressed.

The differences in phase fractions of $\alpha'$ and $\alpha_d$, highlighted in Fig. 8, are directly related to the wide range of heating and cooling rates encountered during the process. The microstructure and internal strains just after deposition/solidification are not the same as those obtained at the end of the process. In other words, the SSTC experienced by a material due to subsequent layer addition or LHTs alters the fraction of phases that formed just after solidification. This alteration can occur in two ways: (i) The temperature of the material increases above the $\beta$ transus resulting in complete re-transformation of all hcp phases into $\beta$; this change can occur in the initial stages of thermal cycling. In this case, a new set of hcp phases is formed during the subsequent cool down. (ii) The temperature does not increase above the $\beta$ transus, which happens after a few thermal cycles and until the end of the process. In this case, thermal cycling triggers the diffusive decomposition of the metastable $\alpha'$ into $\alpha_d$ and $\beta$ [17,61,65].

This diffusion-driven phenomenon and the degree of decomposition of $\alpha'$ are highly dependent on the local chemical composition (solid solution) and the time and temperature at which the heat treatment is performed. Cao et al. [66] investigated the degree of martensitic decomposition in annealed LPBF Ti-6Al-4V and the impact of annealing on the tensile properties. They reported a full decomposition of $\alpha'$ into fine $\alpha_d$ and $\beta$ after annealing for at least six hours at 800 °C. With this treatment, they were able to reach a YS and UTS of 937 MPa and 1041 MPa, respectively, with an elongation close to 18.5%. Note that in other works [67–69] annealing of LPBF Ti-6Al-4V is typically performed for 2 hours at 800 °C to obtain a full decomposition of $\alpha'$. Meanwhile, de Formanoir et al. [38] reportedly achieved a complete decomposition of $\alpha'$ in their EPBF Ti-6Al-4V after only 5 minutes of annealing at 700 °C. This clearly demonstrates the influence of the process parameters on the microstructure formation, in particular solute trapping and diffusion during fabrication. In the case of LPBF, due to its higher cooling rates, solute trapping is higher than in EPBF or LDED Ti-6Al-4V, and this could explain why $\alpha'$ in LPBF requires more time at high temperatures to decompose. Note that while it is possible to estimate the temperature occurring during the building and LHTs via X-ray diffraction measurements [55,70], obtaining the local chemical composition is a very difficult, if not impossible, task. Therefore, it is difficult to know a priori how much time would be needed to fully decompose $\alpha'$.

During LDED and LHTs of our Ti-6Al-4V, even though temperatures in the range 700–800 °C are not maintained for a very long time at any material point, evidence of significant (but not complete) decomposition of $\alpha'$ is clear in Fig. 7 where $\alpha'$ needs coexist with $\alpha_m$ and lathy and globular $\alpha_d$, as well as in Fig. 8 where the proportions of $\alpha_d$ are higher than those of $\alpha'$ for all samples. Samples that underwent more thermal cycling (i.e., LHT100 and LHT300) with respect to NHT exhibited a higher degree of martensitic decomposition, with the highest

amount of $\alpha_d$. Interestingly, the amount of $\beta$ remains the same between samples. As the number of thermal cycles decrease along BD, it is expected that the top of each sample present the least (if any) martensite decomposition [17,24,61].

A mechanism behind $\alpha'$ decomposition into $\alpha_d$ and $\beta$ was recently proposed by Kaschel et al. [71]. Based on their analysis, the $\alpha'$ unit cell is initially in a compressed state. Upon heating the material from room temperature to 400 °C, the a and c lattice parameters associated with $\alpha'$ slightly increase due to the redistribution of Al and V atoms yielding an undeformed martensite phase. Further heating causes a linear increase in a and c and the solute atoms, in particular V, diffuse out of the grain as $\alpha_d$ and $\beta$ form out of $\alpha'$ at ~550 °C. The diffusion of Al and V and substitution by Ti results in an asymmetrical lattice due to varying chemical compositions, and an independent increase in basal and prismatic planes. The rate of diffusion increases with increasing temperature until the lattice is saturated in Ti. This saturation is accompanied by an expansion in the lattice size due to the larger size of Ti. This additional expansion provides sufficient room within the unit cell for elements to rearrange leading to full decomposition of $\alpha'$ into $\alpha_d$ and $\beta$. This full decomposition was observed to occur at 700 °C and increasing the temperature beyond this point but below the $\beta$ transus (995 °C) simply accelerates the decomposition process. Note that these temperature values may not correspond exactly to the ones experienced by the Ti-6Al-4V studied in this work because the heating rate (200 °C/min) used in the work of Kaschel et al. [71] is typically lower than the highest ones experienced during LDED. Also, the higher cooling/heating rates should result in lower decomposition due to less time at high temperatures.

As mentioned earlier, the variation in the number of thermal cycles that each layer experiences as a function of the BD also affects the chemical composition as a function of sample height. Such a gradient in chemical composition has a direct consequence on the material strength through the change in solute trapping. In particular, in the case of Ti-6Al-4V, addition of layers and LHTs result in the evaporation of Al [69,72,73]. Elemental composition analysis carried out on the NHT sample revealed a noticeable decrease in the Al wt% from 6.13 (powder) to 5.73 (see section 2.1). Additional LHTs would result in a further decrease in Al content. In particular, LHT300 should result in more evaporation of Al than LHT100; chemical composition analysis could not be carried out on these samples due to lack of sufficient material. It is well known that the evaporation of Al reduces the solid solution strength [69,72,73]. Obtaining the precise contribution of this aspect to the overall mechanical response requires performing local mechanical tests such as nanoindentation all along the height of the sample and in multiple locations, which was beyond the scope of this work. Despite the additional Al evaporation that LHT100 would experience with respect to NHT, it should be noted that along PD the former shows a better strength/ductility trade-off than the latter along with nearly the same YS and UTS. Therefore, it can be argued that the differences in $\alpha_d$ and $\alpha'$ between samples have a much more important contribution to variation in mechanical response than the differences in Al content.

A reduction in Al content also results in an increase in the weight percent of other elements including the $\beta$ stabilizing element V. The change in composition then also causes a change in the $\beta$ transus and martensite start temperatures that can affect the hcp phase content. According to Zhang et al. [74], for Ti-alloys, a decrease in the $\alpha$ stabilizing elements such as Al results in a decrease in the $\beta$ transus and martensite start temperature. Furthermore, an increase in $\beta$ stabilizers such as V, results in further decrease in these temperatures. These differences between samples further contribute to the differences in the hcp phases that form. A clearer understanding of the change in chemical composition as a function of height would lead to a better understanding of the phase fractions formed and their impact on the mechanical response.

The heating and cooling rates encountered during LDED and LHTs also led to the occurrence of the $\alpha_m$ phase and its presence was higher in LHT100 than in NHT and LHT300; although it should be noted that this observation is made based on a relatively small region in the sam-
ple. This phase also presented a much lower GND density than $\alpha'$ and $\alpha_d$ phases. Furthermore, in contrast to recent studies [28–32], in the present study this phase did not decompose into ultra-fine $\alpha_d$ and $\beta$ lamellae as evidenced from the lack of structuring within this phase in the BSE images (Fig. 7). For instance, Suprobo et al. [29] performed annealing at 850 °C during 30 min to reach such a stage of decomposition of $\alpha_m$. Karimi et al. [32] used a secondary laser during printing which lead to the decomposition of $\alpha_m$. This phase has been reported to have a lower strength and higher ductility than $\alpha'$ and increasing its fraction could improve the strength/ductility trade-off. However, despite the fact that a lower fraction of $\alpha_m$ is present in LHT300, this sample exhibited a much higher ductility/toughness than NHT and LHT100, indicating that $\alpha'$ and $\alpha_d$ phases are playing the most important role in determining the differences in the mechanical response between samples.

## 5. Conclusion

In this work, the impact of in-process laser heat treatments (LHTs) on altering the microstructure and the mechanical response of laser direct energy deposited (LDED) Ti-6Al-4V is studied. The reference as-built sample is manufactured with a laser power of 300 W and two other samples are prepared with LHTs performed after the deposition of each layer at 100 W (LHT100) and 300 W (LHT300). Uniaxial tensile test response is studied along the building and printing directions. All samples exhibit the same Young's moduli indicating that printing and additional LHTs result in at least a transversely isotropic elastic response in the plane formed by the building and printing directions. However, LHT100 exhibits a better strength/ductility trade-off than NHT along the printing direction. While LHT300 exhibits a poorer strength than both NHT and LHT100, a significant improvement in ductility and toughness is obtained. These alterations in the mechanical response are obtained simply from the LHTs and without any *in situ* alloying or post-process heat treatments.

Synchrotron XRD and electron microscopy measurements were performed prior to the extraction of dogbones to understand the contribution of different microstructural features and internal strains to the variations in mechanical response. The analysis reveals a similarity in the $\beta$ phase fraction, intergranular lattice strains, grain sizes, intragranular lattice strains, geometrically necessary dislocation densities and texture between the samples. These features have a minor, if not negligible, contribution to the variation in the mechanical response between samples. Additional LHTs also result in the evaporation of Al, and a change in the solid solution strengthening effect. However, this difference is argued to have a negligible role due to the similarity in strength along the printing direction between NHT and LHT100. The only remaining possibility is then the difference in the hcp phase content between samples.

Among the different possible hcp phases, variations in the stable diffusively formed $\alpha_d$ (soft) and the metastable martensitic $\alpha'$ (hard) between samples are found to have a primary contribution to the difference in mechanical properties. LHT300 and NHT exhibit respectively the highest and lowest amount of $\alpha_d$ and the inverse for $\alpha'$, which explains the observed differences in mechanical properties. Another hcp phase, massive-$\alpha$ ($\alpha_m$), also occurs in all samples but its proportion is an order of magnitude lower than $\alpha'$ and $\alpha_d$. Nevertheless, its presence is the highest in LHT100 and along with its intermediate $\alpha'$ and $\alpha_d$ content, it contributes to a better strength/ductility trade-off than NHT.

This study highlights the tremendous potential of in-process LHTs in engineering the microstructure of LDED Ti-6Al-4V during manufacturing while exhibiting a good strength/ductility trade-off as well as improving ductility and toughness without the need of performing additional heat treatments outside an AM machine. In this sense, performing in-process LHTs can be further explored as an alternative to traditional post-process heat treatments such as annealing, which are particularly impractical when it comes to repairing parts via AM. Immediate perspectives include optimizing LHT strategies and processing parameters such as laser power, scanning speed, or applying multiple LHTs after

each layer deposition. In addition, post-process LHTs inside the AM machine can also be performed to further modify the microstructure close to the surface of the samples as done in [75] for the case of an LDED 316L stainless steel; the treatments performed in that work resulted in a sandwiched microstructure with significant intragranular refinement close to the sample surface causing a simultaneous enhancement in strength, ductility and fatigue limit of LDED 316 L. For Ti-6Al-4V, the strategies should be optimized to trigger martensitic decomposition in order to increase the $\beta$ and globular $\alpha_d$ fractions in order to improve ductility.

## CRediT authorship contribution statement

K.A. Abdesselam: Writing – review & editing, Writing – original draft, Visualization, Validation, Methodology, Investigation, Formal analysis, Conceptualization. S. Gaudez: Writing – review & editing, Visualization, Validation, Investigation, Conceptualization. S. Van Petegem: Writing – review & editing, Visualization, Investigation. V. Honkimäki: Resources, Investigation. S. Hallais: Resources, Investigation. L. Cornet: Writing – review & editing, Resources, Investigation. M. Vallet: Writing – review & editing, Resources, Investigation. M.V. Upadhyay: Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Resources, Project administration, Methodology, Investigation, Funding acquisition, Conceptualization.

## Funding

This work was supported by the Horizon 2020 - EXCELLENT SCIENCE - European Research Council Starting Grant 2020 project GAMMA (Grant agreement ID: 946959).

## 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.

## Acknowledgements

The authors acknowledge the support of the European Synchrotron Radiation Facility (ESRF) for beamtime access via proposal no. MA-5222 and thank Florian Russello for assistance and support in using beamline ID31. The authors also thank Hugo Pinsard in helping conducting the experiments. The authors are grateful to Phillipe Lasson and Ihssan Hendi from Synergie 4 for the EDS maps using their Flatquad detector (Bruker). The authors thank Kewei Chen, Alexandre Tanguy, Nathalie Isaac and Kassigié Dembélé for the support during experimental characterization. The authors would also like to thank Sandrine Tuseau-Nenez, Kim Vameensel and Pushkar Prakash Dhekne for the fruitful discussions on the XRD data. KAA, SG and MVU are grateful to European Research Council (ERC) for their support through the European Union's Horizon 2020 research and innovation program for Starting Grant project GAMMA (Grant agreement ID: 946959).

## Appendix A. Supplementary material

Supplementary material related to this article can be found online at https://doi.org/10.1016/j.matdes.2025.113997.

## Data availability

Data will be made available on request. Open access to synchrotron experimental raw data will be available on July 2025, see https://doi.org/10.15151/ESRF-ES-847864546.
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