# Alpha-case promotes fatigue cracks initiation from the surface in heat treated Ti-6Al-4V fabricated by Laser Powder Bed Fusion

Quentin Gaillard, Florian Steinhilber, Amélie Larguier, X Boulnat, Jean-Yves

Buffiere, Guilhem Martin, Sylvain Dancette, Sophie Cazottes, Rémy

Dendievel, Christophe Desrayaud

► To cite this version:

Quentin Gaillard, Florian Steinhilber, Amélie Larguier, X Boulnat, Jean-Yves Buffiere, et al. Alpha-case promotes fatigue cracks initiation from the surface in heat treated Ti-6Al-4V fabricated by Laser Powder Bed Fusion. International Journal of Fatigue, 2025, 190, pp.108621. 10.1016/j.ijfatigue.2024.108621. emse-04717972

HAL Id: emse-04717972

https://hal-emse.ccsd.cnrs.fr/emse-04717972v1

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# Alpha-case promotes fatigue cracks initiation from the surface in heat treated Ti-6Al-4V fabricated by Laser Powder Bed Fusion

Quentin Gaillard<sup>a,b,*</sup>, Florian Steinhilber<sup>b,c</sup>, Amélie Larguier<sup>b</sup>, Xavier Boulnat<sup>b</sup>, Jean-Yves Buffiere<sup>b</sup>, Guilhem Martin<sup>c</sup>, Sylvain Dancette<sup>b,d</sup>, Sophie Cazottes<sup>b</sup>, Rémy Dendievel<sup>c</sup>, Christophe Desrayaud<sup>a</sup>

## ARTICLE INFO

### Keywords:

Ti-6Al-4V
Laser Powder Bed Fusion
Heat treatments
Alpha-case
Fatigue
Embrittlement

## ABSTRACT

This research investigates the effect of the formation of an oxygen-stabilised titanium alpha layer – called alpha-case at the surface – on the fatigue properties of Ti-6Al-4V (Ti64) alloy components produced by Laser Powder Bed Fusion (L-PBF). Three post processing heat treatments with different controlled atmospheres were carried out on samples with as-built surfaces to evaluate how differences in alpha-case layer thickness and hardness affect the material's susceptibility to surface embrittlement and its overall fatigue performance. The investigation includes bulk and subsurface microstructural analysis, surface characterisation by X-ray computed tomography (XCT), and fatigue testing. Key findings show that alpha-case layers can reduce the fatigue resistance of L-PBF fabricated Ti64. The presence of a 70 ± 3 µm thick alpha-case layer was found to promote crack initiation. This is emphasised by a higher density of initiated cracks, thus leading to a reduction in fatigue life. Conversely, thinner alpha-case layers were found to have a reduced impact on the fatigue performance, highlighting the critical role of post processing heat treatments in modulating the fatigue resistance of the material. The use of XIII to characterise the surfaces of the specimens in 3D confirms that fatigue cracks primarily initiate at surface notches, highlighting the predominance of as-built surfaces over microstructure in determining the fatigue resistance of L-PBF Ti64 components.

## 1. Introduction

Among additive manufacturing (AM) processes, Laser Powder Bed Fusion (L-PBF) stands as a prevalent method for producing metallic components [1,2], particularly in the aerospace sector, a major driver for this technology [3]. The fabrication of geometrically complex components in Ti-6Al-4V (Ti64) alloy using L-PBF has proven to be robust [4,5]. Nevertheless, the widespread adoption of Ti64 produced via L-PBF is hampered by several challenges. In particular, as-manufactured parts exhibit significant residual tensile stresses near their surfaces [4, 6], a predominantly fine martensitic $\alpha'$ microstructure [7,8], together with manufacturing defects such as pores [9], and considerable surface roughness including various surface defects [10,11]. The martensitic microstructure confers high strength but results in a poor ductility [10, 12,13], and in conjunction with residual stresses, adversely affects the fatigue performance of the material [14,15].

To improve the mechanical properties, post processing heat treatments (HT) are used to relieve residual stresses [16,17] and to decompose the martensitic microstructure into a balanced $\alpha + \beta$ mixture. Extensive research has been undertaken to develop HT that retain a fine equilibrium microstructure to optimise the strength/ductility trade-off of the alloy. This research has primarily focused on machined and/or polished samples [12,15,18,19], but excellent tensile properties were also found with near-net-shape samples with as-built surfaces [13,20, 21].

However, fatigue properties are significantly affected by a variety of surface defects inherent to the L-PBF process [22,23]. The detrimental effect of as-built surfaces on the fatigue properties of Ti64 alloy produced by L-PBF is well documented, affecting both as-fabricated (untreated) and stress-relieved samples [24,25], sub-transus heat treatments [26], or hot isostatic pressing [27,28], each resulting in different

<footnote><sup>a</sup> Mines Saint-Etienne, Univ. Lyon, CNRS, UMR 5307 LGF, F-42023 Saint-Etienne, France</footnote>

<footnote><sup>b</sup> Univ. Lyon, INSA Lyon, CNRS, UMR 5510 MATEIS, F-69621 Villeurbanne, France</footnote>

<footnote><sup>c</sup> Univ. Grenoble Alpes, Grenoble INP, CNRS, UMR 5266 SIMaP, F-38000 Grenoble, France</footnote>

<footnote><sup>d</sup> ElyTMAx IRL3757, CNRS, Université de Lyon, INSA Lyon, Centrale Lyon, Université Claude Bernard Lyon 1, Tohoku University, Sendai, 980-8577, Japan</footnote>

<footnote>* Corresponding author at: Mines Saint-Etienne, Univ. Lyon, CNRS, UMR 5307 LGF, F-42023 Saint-Etienne, France.
E-mail address: quentin.gaillard@emse.fr (Q. Gaillard).</footnote>
microstructural characteristics. These defects act as stress concentrators, facilitating local plastic deformation and potentially leading to crack initiation early in the fatigue life. Such surface defects significantly reduce the fatigue strength by accelerating the crack initiation stage. In particular, Ti64 alloys produced by additive manufacturing are particularly sensitive to notch-like surface defects [29–31], which can originate from various sources. In up-skin and vertical regions, these defects may be caused by the layering of molten powder, often resulting in a plate-pile or staircase-like surface appearance [29,32,33]. Crack initiation on Ti64 as-built surfaces inherited from L-PBF most commonly occurs at surface notches [23–26]. Various studies have shown that notches are the primary origin of failure in most cases for both L-PBF and E-PBF Ti64 samples with as-built surfaces [24,25,27].

When exposed to high temperatures in an oxygen-containing environment, such as during HT performed in air, titanium and its alloys are susceptible to the formation of oxide layers on the surface. This goes hand in hand with oxygen diffusion from the surface into the bulk of the material [34,35]. The oxygen-enriched layer, commonly referred to as “alpha-case”, is characterised by its high hardness and brittleness and is often associated with a decrease in mechanical properties, primarily due to the initiation of surface cracks under load [36,37]. To date, few studies have specifically investigated the effect of the brittle alpha-case on the fatigue properties of titanium alloys. On a mill-annealed $\alpha + \beta$ alloy with an equiaxed microstructure, Fukai et al. revealed that the bending fatigue strength ($R = -1, f = 25$ Hz) deteriorates in the presence of an alpha-case layer approximately 50 µm thick [38]. In addition, Pilchak et al. found that the fatigue life ($R = 0.05, f = 20$ Hz) of a heat-treated near-$\alpha$ alloy with an equiaxed microstructure is severely reduced by a factor of 10 with an alpha-case layer of approximately $35 \pm 5$ µm [39]. Gaddam et al. demonstrated that the low cycle fatigue life ($R = 0, f = 0.5$ Hz) of a heat treated near-$\alpha$ alloy with a bimodal microstructure decreases by 50% and 90% at a strain amplitude of 0.4 %, with alpha-case layers of approximately 2 µm and 10 µm, respectively [34]. Traditional manufacturing routes for titanium products, such as casting and forging, usually involve a machining step on the heat treated ingot or semi-finished product. The alpha-case layer formed during fabrication and/or HT is therefore removed at this stage. Conversely, in the case of L-PBF products used in near-net-shape (no surface treatment), an alpha-case layer can potentially form during the post-fabrication HT. Recently, it has been shown that the ductility of L-PBF produced Ti64 is affected by alpha-case layers with depths $\ge 50$ µm [13]. However, to the best of our knowledge, the effect of the alpha case on the fatigue properties of Ti64 components produced by L-PBF has not been investigated. It is worth noting that L-PBF Ti64 has specific characteristics compared to the aforementioned studies, such as printing defects, high surface roughness and fine microstructure.

As the fatigue strength of Ti64 components is often critical in aerospace applications [40], a major challenge in manufacturing them by L-PBF is to improve the surface quality. Post processing surface treatments have been shown to significantly improve fatigue resistance [22, 26, 41]. However, the applicability of many surface treatments may be limited by cost or feasibility, particularly for parts with complex geometries or internal features. Therefore, in cases where HT induces an alpha-case layer on the surface of L-PBF Ti64 specimens that cannot be mitigated by surface treatment techniques, a negative effect on fatigue properties is expected. The purpose of this study is to evaluate the influence of different heat treatment conditions on the thickness of the alpha-case layer and to assess the effect of such oxygen-enriched layers on the fatigue life of parts with as-built surfaces.

## 2. Materials and experimental methods

### 2.1. Samples manufacturing

The specimens examined in this study were produced in near-net-shape using a first generation AddUp FormUp 350 machine operated

<table><caption>**Table 1**

Chemical composition of the batch of grade 5 Ti64 powder used in this work.</caption><thead><tr><th></th><th>Ti</th><th>Al</th><th>V</th><th>Fe</th><th>O</th><th>C</th><th>N</th><th>H</th></tr></thead><tbody><tr><td>wt%</td><td>Base</td><td>6.22</td><td>3.87</td><td>0.18</td><td>0.14</td><td>0.009</td><td>0.02</td><td>0.004</td></tr><tr><td>at%</td><td>Base</td><td>10.49</td><td>3.46</td><td>0.15</td><td>0.40</td><td>0.03</td><td>0.07</td><td>0.18</td></tr></tbody></table>

by the French aerospace manufacturer Dassault Aviation in Argonay, Savoie, France. This machine is equipped with two YAG lasers, each with a power of 500 W and operating at a wavelength of 1 070 nm. The build plate was divided equally between the two lasers, each covering half of the build area. A laminar flow of argon (Ar) was used to maintain a controlled atmosphere within the build chamber. Powder deposition was achieved using a roller mechanism. The build plate was not preheated prior to the build process. The scanning strategy involved applying a meander scan pattern and alternating the laser vectors by 90° between successive layers. Due to confidentiality agreements, specific details of other parameters remain undisclosed. The material used for the samples was ASTM Ti64 Grade 5 powder, supplied by TEKNA, Canada, and produced by plasma atomisation. The powder, whose chemical composition is given in Table 1, had a particle size distribution ranging from 5 µm to 25 µm.

The samples investigated in this study are cylindrical fatigue specimens with a gauge length of 3 mm, a diameter of 3 mm (see Fig. 1.a for further details) and with their longitudinal direction parallel to the build direction. Whilst the use of a roller facilitates the creation of a dense powder bed, care must be taken when producing slender elements to prevent the displacement of freshly deposited thin layers by the action of the roller. Consequently, as shown in Fig. 1, the production of vertical specimens required the use of “sarcophagus-like” support structures: Fig. 1.b shows the two halves of the support structure with the vertical fatigue specimen still inside and maintained, while Fig. 1.c shows a similar fatigue specimen outside the support structure along with one half of it.

It should be noted that the supports were designed to closely surround the specimens without making direct contact with the cylindrical section of the gauge. These structures helped to stabilise the powder bed and prevented the specimens from bending in the direction of the roller. After fabrication, the build plates underwent a cleaning process in a water-filled ultrasonic bath, a critical step to remove any remaining loose powder particles. Therefore, the support structures are designed with powder extraction holes to help remove any powder remaining between the supports and the specimens at this stage (see Figs. 1.b and 1.c). In order to preserve the original surface characteristics and to avoid hiding defects such as notches, the build plates and attached specimens were not sandblasted.

### 2.2. Heat treatment conditions

Prior to removing the samples from the base plate, a stress relieving HT at 720 °C for 2 h was applied to the entire build plate. This procedure was designed to mitigate the potential deformation that could occur if the samples were removed prior to stress relief. Preliminary studies showed that this HT effectively eliminated the residual stresses introduced during L-PBF [16]. This HT was carried out in an industrial BMI B84T horizontal gas quenching furnace following a thermal cycle under secondary vacuum at $10 \times 10^{-4}$ mbar. The cycle included heating at a rate of 10 °C/min and annealing for 2 h. This was followed by controlled cooling in the furnace at an average rate of 2 °C/min to allow the diffusion of alloying elements. Once 400 °C is reached, atomic diffusion is too slow to cause additional phase transformation, so argon was introduced into the chamber to increase the cooling rate. This stress relief HT is also described in Ref. [13], where exactly the same parameters and equipment specifications were used. Therefore, the corresponding time-temperature/pressure curves can be found in Ref. [13].
<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Schematic dimensions of cylindrical fatigue specimens given in mm (a). Images of a vertically printed fatigue specimen inside (b) and outside (c) its sarcophagus-like support structure after cutting from the build plates.</figcaption></figure>

The specimens were then removed from the build plate and the support structures were manually removed. A final ultrasonic bath was used to remove any residual powder particles between the specimens and the supports. The specimens were then divided into three groups, each subjected to different HT conditions:

* **Stress-Relieved (SR) condition.** After stress-relieving HT at 720 °C for 2 h, the first group of samples was tested without any further HT.

* **Laboratory Furnace Heat Treatment (LF-HT) condition.** Following the initial stress-relieving HT, the second group underwent an additional HT at 860 °C/4 h in a Nabertherm RHTC 80-450/15 tubular vacuum furnace. This cycle, carried out under a primary vacuum of $10 \times 10^{-2}$ mbar, included heating at 10 °C/min, annealing for 4 h, and controlled cooling at an average rate of 4 °C/min. Similar to the other conditions, rapid cooling was initiated at 400 °C by water quenching.

* **Industrial Furnace Heat Treatment (IF-HT) condition.** The third group was subjected to the same HT at 860 °C/4 h after the initial stress relief, using the same industrial furnace. The parameters and thermal cycle were similar to the LF HT.

The laboratory furnace and the industrial furnace used for the additional HT are similar to those described in a previous paper [13]. In this earlier study it was observed that alpha-case formation occurred during IF-HT whereas LF-HT did not result in alpha-case formation. It was concluded that the vacuum level is not the only parameter to consider in preventing alpha-case formation at the surface of the specimens. Several factors such as the global specific surface area of the heat-treated parts and the dimensions of the chamber can also play a significant role. The industrial furnace has a much larger effective capacity (0.25 m³) than the laboratory furnace ($5 \times 10^{-3}$ m³) which can make it more difficult to maintain air-tightness. In addition, the degassing of fibre/wool insulation elements in the oven may influence oxidation during HT, thereby facilitating the formation of alpha-case. Finally, the formation of alpha-case can be influenced by the "contamination" of the industrial furnace, which is not fully dedicated to the heat treatment of titanium parts. This distinction between IF-HT (alpha-case formation) and LF-HT (no alpha-case formation) allowed a comparative analysis to distinguish the effects of microstructural changes from the effects of alpha-case, both outcomes depending on the HT temperature. The results can be extended to the material studied in the present work, since the samples used in the previous study

were produced with the same Ti64 powders, L-PBF machine, process parameters and underwent HT under the same conditions.

The HT at 860 °C/4 h was selected based on its ability to achieve an equilibrium $\alpha + \beta$ microstructure, which is known to enhance the strength-ductility trade-off of this alloy [12,13,42]. Furthermore, in our previous study, it was found that alpha-case formation reduced the overall ductility and the elongation to failure of the alloy when subjected to HT for 2 h at 860 °C or higher [13]. Consequently, a similar temperature range was adopted in the current study to assess the effect of alpha-case formation on the fatigue performance of the material.

## 2.3. Microstructure characterisation

Microstructural analysis was performed using Optical Microscopy (OM) and Scanning Electron Microscopy (SEM). After the fatigue tests, samples were prepared from the gauge length of the specimens, including both horizontal (XY) and vertical (Z) cross sections. Preparation for microstructural examination involved grinding the specimens with 320 grit SiC paper followed by polishing with a 9 µm diamond solution. This was followed by chemical and mechanical polishing using a mixture of 50 nm colloidal silica suspension (90%), hydrogen peroxide (6%) and ammonia (4%). For microstructure imaging by Backscattered Electron (BSE) or Electron Backscattered Diffraction (EBSD), the samples were further polished by vibratory polishing with colloidal silica.

Area phase fractions were quantified from BSE images acquired using a ZEISS Supra 55 VP Field Emission Gun SEM. Image analysis, performed using ImageJ software, allowed phase discrimination based on the contrast observed in the images. This analysis provided the average phase fraction (± standard deviations) from 20 images taken in the XY plane for each HT condition. EBSD analysis was performed using the same SEM operating at 20 kV with an Oxford Instruments Symmetry camera. For each condition an EBSD map was collected with a step size of 0.25 µm in the XY plane. The EBSD data were processed using the MTEX toolbox in MATLAB, identifying $\alpha$ lath boundaries with a 5° misorientation threshold between grains. Grains smaller than 10 pixels were excluded from the analysis and the thickness of the $\alpha$ laths was determined by fitting ellipses and measuring the minor axis.

The alpha-case layer on samples polished in the XY plane was revealed using Weck's reagent (ammonium bifluoride dissolved in ethanol and water) for approximately 30 s. Micrographs of the etched samples were taken using an Olympus GX51 light microscope. Quantitative measurements of the thickness of the alpha-case layer were made using ImageJ software, based on 20 images taken around the perimeter of the specimen.
<figure><img src="image_4.png" /><figcaption>Fig. 2. Schematic representation of the locations of the Vickers hardness measurements made at 50 g and 1 kg on horizontal (XY) polished cross-sections of fatigue specimens.</figcaption></figure>

<table><caption>**Table 2**

Microstructural characteristics and bulk hardness measurements for the three HT.</caption><thead><tr><th>Condition</th><th>α laths thickness (μm)</th><th>β phase fraction (%)</th><th>Bulk hardness (HV1)</th><th>Alpha-case thickness (μm)</th></tr></thead><tbody><tr><td>SR (720 °C/2 h)</td><td>0.61 ± 0.05</td><td>2.2 ± 0.3</td><td>395 ± 2</td><td>10 ± 1</td></tr><tr><td>LF-HT (SR + 860 °C/4 h)</td><td>1.02 ± 0.07</td><td>4.2 ± 0.4</td><td>368 ± 2</td><td>—</td></tr><tr><td>IF-HT (SR + 860 °C/4 h)</td><td>1.06 ± 0.08</td><td>4.1 ± 0.3</td><td>366 ± 3</td><td>70 ± 3</td></tr></tbody></table>

## 2.4. Mechanical tests

Fatigue tests were conducted in air at room temperature using an Instron 8516 servo-hydraulic machine equipped with a 100 kN load cell. These tests were performed at a constant stress amplitude, using a sinusoidal waveform at a frequency of 10 Hz, with the stress ratio set at $R = 0.1$. A range of maximum applied tensile stress between 300 MPa and 650 MPa was selected from the fatigue properties found in the literature for L-PBF Ti64 with as-built surfaces. A total of 23, 9, and 14 specimens were tested for the SR, LF-HT and IF-HT conditions respectively. Tests that did not result in failure after $10^7$ cycles for SR specimens were stopped and classified as “run-outs”. Due to time constraints, run-out LF-HT and IF-HT specimens were only stopped after $1.5 \times 10^6$ cycles. For each condition, one specimen subjected to a maximum applied stress of 450 MPa was cycled for approximately 95% of its estimated life and then removed from the fatigue machine unbroken. These specimens are hereafter referred to as “interrupted”. The stresses were determined based on the theoretical initial cross-sectional area, $S = \pi (3 \text{ mm})^2 \approx 7.07 \text{ mm}^2$.

Vickers microhardness tests were performed on horizontal (XY) mirror-polished cross-sections of fatigue specimens using a Buehler microhardness tester. To determine the bulk hardness for the three conditions studied, 20 measurements were taken at random locations within the bulk of the samples, as shown schematically in Fig. 2. These bulk measurements were taken using a load of 1 kg and an indentation time of 15 s, resulting in an average indentation diagonal length of 70 µm. The results are presented in Table 2 as means ± standard deviations.

In addition, to assess the increase in hardness due to the oxygen enriched alpha-case layer, microhardness profiles were obtained at a load of 50 g, resulting in an average indentation diagonal of 15 µm. This load was chosen to minimise the effects of microstructural variability, thereby reducing data scatter whilst maintaining adequate spatial resolution. The initial indentation was placed at 10 µm from the sample surface and measurements were taken in a staggered pattern to achieve a spatial resolution of 10 µm between successive points, as shown in the enlarged view of Fig. 2. Three hardness profiles were obtained for each condition and the results are shown in Fig. 5.b as a curve representing the data together with the scatter plots of each measurement.

## 2.5. 3D characterisation of surfaces using X-ray computed tomography

X-ray computed tomography (XCT) was used to scan the gauge length of all the fatigue specimens, both before the fatigue test and after fatigue failure. The scans were performed using a cone beam phoenix | X-ray V | tome | x laboratory tomograph, set to a voltage of 140 kV and a current of 80 µA, capturing 1200 projections. A 0.1 mm Cu filter was implemented to minimise beam hardening artefacts during acquisition. After acquisition, the images were processed with a $3 \times 3$ median denoising filter. Image reconstruction was performed using a standard filtered back-projection algorithm supported by the proprietary phoenix datos|x software. In addition, a numerical filter was applied during reconstruction to further reduce beam hardening artefacts.

Each specimen was scanned three times along its longitudinal axis prior to fatigue testing with a small overlap of 100 µm. The number of scans was chosen primarily to ensure that failure always occurred in an area previously characterised by SCT. The voxel size for each sample was 3 µm, with a total characterised length from the three scans of 11 mm. After reconstruction, the scans were converted to 8-bit unsigned volumes and merged using a custom Python script. To reduce noise, a $2 \times 2 \times 2$ median filter was applied to the aggregated volume. This processed volume was then used to compute local surface roughness and curvature, as inherited from the L-PBF method. The methodology for calculating these parameters to achieve 3D characterisation of complex surfaces is described in detail in Ref. [43].

XCT scanning of the fatigue specimens enabled the characterisation of the as-built surfaces using the aforementioned ex situ methodology, both before and after fatigue testing. The derived data supported a comprehensive killer defect identification workflow. This workflow involves analysis of the fracture surface of fatigue specimens using SEM in Secondary Electron (SE) or Backscattered Electron (BSE) mode to locate the crack initiation site. It then involves the study of the 3D fracture surface from SCT scans taken after fatigue failure and the identification of the killer defect on the 3D surface from the SCT scan of the pre-fatigue specimen [43,44].
<figure><img src="image_5.png" /><figcaption>Fig. 3. EBSD inverse pole figure α orientation maps (a and b) and SEM-BSE images (c and d) showing the microstructure of an SR (720 °C/2 h) sample (a and c) and of an IF-HT sample (720 °C/2 h + 860 °C/4 h) (b and d). In the EBSD maps, the build direction Z is used as the reference axis. In the SEM-BSE images, the α phase appears in dark contrast, while the β phase appears in white.</figcaption></figure>

# 3. Results and discussion

## 3.1. Bulk microstructure observations

Fig. 3 shows EBSD inverse pole figure maps of the α phase (Figs. 3.a and 3.b), alongside SEM-BSE images (Figs. 3.c and 3.d). Each image was taken in a plane perpendicular to the building direction Z.

Each HT condition is characterised by a two-phase α+β microstructure (α, hexagonal close-packed and β, body-centred cubic). This microstructure results from the decomposition of the as-built α' martensitic phase during post processing heat treatment at annealing temperatures above 500 °C [45,46]. The α' martensite forms on cooling from the parent β grains. In planes normal to Z, the legacy structures from the parent β grains create a “chessboard pattern”, as seen in Figs. 3.a and 3.b. Such patterns are typically observed in Ti64 produced by L-PBF using a 0°/90° scanning strategy between layers [7,47,48]. As up to 12 different α variants can form from a single parent β grain orientation [16,49,50], this contributes to the randomisation of the texture of the α phase as shown in Figs. 3.a and 3.b. The α laths thicknesses were measured from EBSD orientation maps similar to those shown in Fig. 3.a for the SR condition and Fig. 3.b for the IF-HT condition. The average α laths thickness was estimated and the results are summarised in Table 2. In the SR condition, the α laths thickness is 0.61 ± 0.05 µm. These values are slightly higher than those reported in Ref. [13] for the same material using the same method. This is due to the use of a larger EBSD step size which reduces the detection of ultra fine α laths. After LF-HT and IF-HT the α laths are thicker, averaging 1.02 ± 0.07 µm and 1.06 ± 0.08 µm respectively. These values are close to the average α lath thickness reported in the literature for similar HT conditions [42,51].

SEM-BSE images, similar to those shown in Fig. 3.c for the SR condition and Fig. 3.d for the IF-HT and LF-HT conditions, allowed the estimation of the α/β phase fractions, with the results summarised in Table 2. In the SR condition, the β phase fraction barely reaches 2.2 ± 0.3 %, but almost doubles after LF-HT or IF-HT to 4.2 ± 0.4 % and 4.1 ± 0.3 %, respectively.

In summary, the bulk microstructure of the Ti64 alloy in the SR condition consists of finer α lamellae and a lower β phase fraction compared to the bulk microstructure after additional HT at 860 °C/4 hours. The bulk microstructures resulting from LF-HT and IF-HT are almost identical. This correlates with the bulk hardness measurements for each condition as shown in Table 2. The material shows a higher hardness in the SR condition (395 ± 2 HV1) compared to the hardness measured after 860 °C/4 hours, regardless of the furnace used: LF-HT (368 ± 2 HV1) or IF-HT (366 ± 3 HV1).

## 3.2. Alpha-case characterisation

The alpha-case layers identified in the three HT conditions are shown in Fig. 4. The use of Weck's reagent selectively highlights the oxygen rich layer in white, making it easier to measure the thickness of the alpha-case layer. The measurements are given in Table 2.

For the SR condition (Fig. 4.a), a 10 ± 1 µm thick alpha-case layer was revealed after etching, in agreement with the results reported in Ref. [13] for the same condition in the industrial furnace.

Despite the similar initial SR treatment, the LF-HT condition showed no discernible alpha-case layer (Fig. 4.b). This can be attributed to the diffusion of oxygen from the initial alpha-case layer formed during SR deeper into the material during the subsequent treatment at 860 °C/4 h,
<figure><img src="image_6.png" /><figcaption>Fig. 4. Characterisation of the alpha-case in SR (a), LF-HT (b) and IF-HT (c) conditions. These optical micrographs, taken from vertical (Z) cross-sections of fatigue specimens, reveal the alpha-case in white after chemical etching using Weck's reagent. SEM-BSE image (d) showing the surface microstructure of an IF-HT specimen.</figcaption></figure>

resulting in its diffusion within the core material. This implies that no additional oxygen was introduced into the alloy in the laboratory furnace atmosphere, a conclusion also drawn in the aforementioned study [13].

In the IF-HT scenario, the initial $10 \pm 1$ µm-thick alpha-case layer expanded to $70 \pm 3$ µm due to enhanced oxygen diffusion from the specimen surface in the oxidising environment of the industrial furnace (Fig. 4.c). Compared to the alpha-case layers documented in the previous study, this $70 \pm 3$ µm measurement at 860 °C/4 h falls between the values obtained for IF-HT at 860 °C/2 h ($48 \pm 2$ µm) and 920 °C/2 h ($90 \pm 5$ µm) [13]. This observation is consistent with the understanding that alpha-case formation in Ti64 alloy produced by L-PBF is primarily dependent on the HT temperature and is driven by the diffusion of oxygen in α titanium [13,35].

Finally, Fig. 4.d shows an SEM-BSE image taken at the surface of an IF-HT specimen. In the region of the estimated alpha-case thickness ($\approx 70 \pm 3$ µm from Fig. 4.c), a gradient of β phase area fraction can be seen: almost no β is found near the surface (only isolated islands of β phase appearing in bright contrast), then the β phase becomes more visible with a more elongated morphology with increasing depth. Beyond the alpha-case thickness, the bulk microstructure already observed in Fig. 3.d is found. This gradient in the β phase area fraction can be attributed to the presence of an oxygen content gradient from the surface. Indeed, oxygen is a strong alpha-forming element in titanium alloys, thus reducing the stability of the β phase. Oxygen gradients at the surface will be further discussed in the next paragraphs.

Fig. 5.a shows microhardness profiles taken from the surfaces of horizontal (XY) polished cross-sections of fatigue specimens under the

three different conditions (see indent locations in Fig. 2). The scatter plots for the three measurements taken at different depths are shown, together with the fitted curve (dashed line) for each condition.

For the SR condition, a marginal increase in hardness is observed at a distance of 10 µm from the surface, with subsequent measurements closely matching the bulk hardness detailed in Table 2. In the LF-HT condition, the initial hardness measured at 10 µm from the surface exceeds the bulk hardness and shows a gradual decrease as one moves towards the bulk. Conversely, the IF-HT condition shows a significant increase in hardness, with mean values remaining above 400 HV0.05 up to a depth of 70 µm from the surface.

The observed increase in hardness is associated with an increase in oxygen concentration, a correlation supported by numerous studies [37,52,53]. To estimate the oxygen content corresponding to these hardness measurements, oxygen diffusion profiles were modelled based on the solution of Fick's second law for a semi-infinite solid (Eq. (1)). This model assumes a steady-state concentration $C_S$ at the surface, a uniform initial concentration $C_0$ within the bulk, and a constant diffusion coefficient $D$ which can be calculated using an Arrhenius law under isothermal conditions (Eq. (2)) [35,54,55].

$$C(x,t) = C_0 + (C_S - C_0) \times \left(1 - erf\left(\frac{x}{2\sqrt{Dt}}\right)\right) \quad (1)$$

$$D(T) = D_0 \times \exp\left(-\frac{Q}{RT}\right) \quad (2)$$

The pre-exponential factor $D_0$ and the activation energy $Q$ were set to $1.31 \times 10^{-3}$ m² s⁻¹ and 225.8 kJ mol⁻¹ respectively. These values correspond to the experimentally determined oxygen diffusion coefficients in the same L-PBF Ti64 alloy from the previous study [13]. They
<figure><img src="image_7.png" /><figcaption>**Fig. 5.** Microhardness (HV0.05) profiles (a) and simulated oxygen concentrations (b) for fatigue specimens under SR, LF-HT, and IF-HT conditions. Microhardness data and simulated oxygen concentration gradients are plotted against the depth from the surface. Vertical dashed lines mark the depths at which an oxygen concentration of 1 at% is reached.</figcaption></figure>

are close to those found in the literature for oxygen diffusion in L-PBF Ti64 [35] or in α titanium [56]. The initial concentration parameter $C_0$ was set to 0.48 at%, corresponding to the bulk material concentration as determined by infrared absorption analysis of combustion gas with a LECO TC-436 (O,N) analyser. This value is slightly higher than the 0.40 at% reported in Table 1 for the powder batch used in this work. In fact, it was reported in Ref. [45] that oxygen pick-up occurs during L-PBF. For both SR and IF-HT conditions, the surface concentration $C_S$ was set to 33 at%, reflecting the maximum solubility limit of oxygen in the titanium α phase. Using these parameters in Eqs. (1) and (2), the oxygen concentration profiles for the SR and IF-HT conditions (blue and red solid lines, respectively) were simulated and are shown in Fig. 5.b. For the LF-HT condition, it is hypothesised that the oxygen initially absorbed into the alpha-case layer during the SR treatment further diffused into the material during the LF-HT treatment without any additional oxygen intake. This idea is consistent with the absence of a visible white alpha-case layer on LF-HT samples after etching with Weck's reagent, as shown in Fig. 4.b. Consequently, when modelling

the oxygen concentration profile for the LF-HT condition, the surface concentration parameter $C_S$ was adjusted to ensure that the total oxygen content (represented by the area under the curve) was consistent between the SR and LF-HT conditions. This adjustment resulted in a $C_S$ value of approximately 5.2 at% and allowed the green solid line profile to be plotted on Fig. 5.b. Based on the simulation, the oxygen content at the extreme surface (depth 5 µm) is significantly reduced for the LF-HT condition compared to the SR condition, but higher from 10 µm depth from the surface. However, in the first 70 µm from the surface, the oxygen concentration is much lower for the LF-HT condition compared to the IF-HT condition. In each simulated oxygen concentration profile, vertical dashed lines mark the depths at which an oxygen concentration of 1 at% is reached. This threshold indicates the level of oxygen above which the ductility of Ti64 decreases significantly [57,58]. Finally, when the fitted microhardness data profiles (Fig. 5.a) are compared with the simulated oxygen concentration profiles (Fig. 5.b) between 10 µm and 120 µm depths from the surface, a good correlation is found.
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** Wöhler curves for SR, LF-HT and IF-HT conditions. Fatigue tests are performed at *R* = 0.1. The straight lines correspond to the mean curves that cross each mean value of the number of cycles before failure for the different applied stresses.</figcaption></figure>

### 3.3. Fatigue properties

Fig. 6 shows the Wöhler (S-N) curves for the three conditions evaluated. Each data point, classified by different markers, illustrates the result for a tested specimen. The markers (circles, crosses, stars) correspond to three possible scenarios: failure, test interruption or survival for the duration of the test (run-out). The solid lines represent the average trend lines, each intersecting the average number of cycles to failure at different maximum applied stresses.

The plot shows no noticeable difference in fatigue resistance between the SR and LF-HT conditions. Firstly, this suggests that the microstructural changes resulting from the LF-HT at 860 °C/4 hours have a limited effect on fatigue resistance. Secondly, it questions the real effect of the 10 µm alpha case layer found on the SR samples. However, a reduction in fatigue life is demonstrated for the IF-HT condition at maximum applied stresses greater than 350 MPa. The microhardness profiles shown in Fig. 5.a suggest that the alpha-case layer increases brittleness, a phenomenon more evident in the IF-HT condition which developed a 70 µm thick alpha-case layer during industrial heat treatment at 860 °C/4 hours (Figs. 4 and 5.a). Our previous study reported that alpha-case layers ≥ 50 µm reduce the elongation to failure of near-net-shape Ti64 samples produced by L-PBF [13]. The present study further confirms that such alpha-case layers also negatively affect the fatigue resistance of L-PBF Ti64 specimens.

In this study the fatigue strength at 10⁷ cycles is estimated to be around 250 MPa. As expected, the scatter of fatigue life increases as the fatigue strength is approached, making the fatigue properties more sensitive to small variations in defect size and harmfulness.

Despite the differences observed between the IF-HT condition and the SR and LF-HT conditions, the Wöhler curves for all three conditions are consistent with the fatigue strengths reported in the literature for specimens with as-built surfaces, typically ranging from 200 to 300 MPa [22,23,27,59,60]. In contrast, the average fatigue properties for L-PBF Ti64 specimens with machined and polished surfaces, which are significantly higher, typically range between 400 and 600 MPa [22, 24,27,44,59–61]. It is therefore essential to evaluate the impact of the as-built surface state on the crack initiation mechanisms. This will be discussed in the next sections for all HT conditions.

### 3.4. Identification of killer defects and crack initiation mechanisms

In order to assess the influence of the as-built surface and the presence of alpha-case layers of different thicknesses on the fatigue properties, the fracture surfaces of the specimens were first examined by SEM. The analyses revealed that, for all the HT conditions studied, all failure initiating cracks originated from the surface, see some examples of each condition in Figs. 7a-c. Furthermore, as can be seen from the magnified views in Fig. 7, no pore was found to cause initiation in the crack initiation zone. This conclusion was indeed confirmed for all the conditions studied for all specimens tested until failure based on fracture surface observations using SEM.

XCT and SEM were used together to identify the initiation sites. This involved scanning the fatigue specimens by SCT both before and after failure. The procedure for identifying initiation sites is illustrated in Figs. 8a-d. First, the fracture surface is examined by SEM. For all specimens, the SEM observations did not allow the identification of a single, clearly localised defect that would have caused the crack to initiate. Rather, the cracks appear to initiate from extended zones, as shown in Fig. 7. In some cases, multiple initiation sites are identified on the surface, as shown in the BSE image in Fig. 8a. This scenario is common under all conditions, especially under high stresses. Having identified the main initiation site in the SEM image, the same fracture surface is then observed in 3D using SCT. This step requires the surface to be extracted from the SCT scan as shown in Fig. 8b. Then, by comparing 8a with 8b, the initiation site on the SCT fracture surface is identified. The next step is to compare the original SCT scan (taken before the fatigue test) with the SCT scan of the fracture surface to identify the surface defect responsible for crack initiation. As demonstrated in Refs. [43,44], the use of minimum curvature analysis is an efficient way to capture surface notches. Thus, for killer defect identification, surfaces were visualised using a point cloud where the colour bar represents the local minimum curvature of the surface, called $κ_{min}$, as illustrated in Fig. 8b-d. In these visualisations, surface notches typically appear in dark blue, indicating areas of lowest minimum curvature. These notches are perpendicular to the loading direction
<figure><img src="image_9.png" /><figcaption>**Fig. 7.** SEM-BSE topological mode observation of fracture surfaces for an SR specimen that failed after 73 547 cycles at 300 MPa (a), an LF-HT specimen that failed after 40 901 cycles at 450 MPa (b), and an IF-HT specimen that failed after 198 645 cycles at 300 MPa (c). In each case, crack initiation occurs from a large portion of the surface as shown in the magnified views.</figcaption></figure>

and run along the specimen surface unless interrupted by significant protrusions. Using this procedure, initiation sites were identified for 10 SR, 6 LF-HT and 6 IF-HT specimens. For 20 out of 22 specimens, failure cracks originated from plate-pile horizontal surface notches, as typically shown in Figs. 8.c-d. Remarkably, one specimen tested at $\sigma_{max}$ = 450 MPa from the IF-HT condition exhibited a crack initiation site without any noticeable defect on the initial surface scanned by/XCT. This situation was not observed in the SR or LF-HT specimens. This may indicate a shift in the crack initiation mechanisms due to the thicker alpha case, although it may be due to the resolution limitations of/XCT, which may not capture the killer defect.

From the/XCT data obtained prior to the fatigue tests, it was also possible to extract the 3D surface roughness of the 22 vertical specimens (10 SR, 6 LF-HT and 6 IF-HT) investigated. The calculation of the 3D surface roughness is described in detail in Refs. [43,44]. The roughness calculation region was performed over a length of 9 mm along the specimen axis. This includes part of the specimen fillet radius. Several roughness parameters ($S_a$ is the arithmetic mean height, $S_q$ the root mean square height, $S_z$ the maximum height and $S_v$ the maximum notch depth) were calculated and the results are summarised in Table 3. From these values, it can be concluded that the surface roughness is similar for the three conditions studied. This result suggests that the observed differences in fatigue behaviour (Fig. 6) are due to a difference in alpha case thickness and cannot be attributed to a difference in surface roughness. The mechanisms that could explain the difference in fatigue resistance in the presence of a thick alpha-case layer are discussed in Section 3.5.

The conclusion of these investigations was that the cracks initiated from shallow elongated surface notches perpendicular to the loading

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

Roughness parameters for the specimens of the three HT, calculated from/XCT data using the 3D characterisation workflow described in Refs. [43,44].

<table><thead><tr><th>Condition</th><th>S<sub>a</sub> (µm)</th><th>S<sub>q</sub> (µm)</th><th>S<sub>z</sub> (µm)</th><th>S<sub>v</sub> (µm)</th></tr></thead><tbody><tr><td>SR (720 °C/2 h)</td><td>7.2 ± 0.6</td><td>10.1 ± 0.8</td><td>143 ± 25</td><td>64 ± 20</td></tr><tr><td>LF-HT (SR + 860 °C/4 h)</td><td>6.7 ± 0.3</td><td>9.5 ± 0.5</td><td>137 ± 26</td><td>60 ± 11</td></tr><tr><td>IF-HT (SR + 860 °C/4 h)</td><td>6.9 ± 0.4</td><td>9.8 ± 0.6</td><td>129 ± 13</td><td>61 ± 10</td></tr></tbody></table>

direction, and multiple crack initiations were observed in some specimens. On several occasions the surface crack initiation zone was found to follow the notch root, sometimes forcing it to follow a tortuous path. This illustrates well how the surface defects, particularly notches, dominate the fatigue properties of specimens with as-built surfaces. The influence of these surface notches can be explained by the stress concentration they generate at their root, also known as the notch effect. From Table 3, one can see that the calculated surface roughness parameters are similar regardless of the condition, including the maximum notch depth parameter $S_v$. Thus, the notch effect is expected to be the same for all fatigue specimens tested. Furthermore, the extent to which this effect affects the fatigue resistance depends on the notch sensitivity of the material. Alloys characterised by high strength and low ductility tend to have higher notch sensitivity, making them particularly susceptible to the notch effect [62]. As the mechanical properties of titanium alloys are highly sensitive to oxygen content [63], the formation of an alpha-case layer, which correlates with an increase in hardness near the surface (see Fig. 5.a), could lead to an increase in the notch sensitivity of the material near the surface, resulting in lower fatigue resistance. For example, this embrittlement could accelerate the initial phase of crack initiation.
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Workflow for crack initiation site identification using both SEM and/XCT scans. The example specimen from the SR condition failed after 129 339 cycles at 350 MPa. SEM-BSE fracture surface observation (a). Fracture surface obtained fromXCT scans after failure (b). The colour bar corresponds to the minimum curvature Kmin, estimated by integral invariant estimation as described in Refs. [43,44]. A crack initiation site is observed on the fracture surface from the side of the specimen (c). Same area as in (c), identified fromXCT scans of the initial state before the fatigue test (d). A notch is identified at the origin of the crack initiation by comparison with (c).</figcaption></figure>

However, it is unclear whether there are multiple initiation sites
along the notches or whether a crack initiated at one point propagates
along the notch root. In attempting to answer this question, it is
instructive to draw a comparison with the study by Vie et al. who
investigated the effect of a 10 µm brittle alumina oxide layer on
the fatigue behaviour of 7075-T6 aluminium alloy (R = 0.1) [64].
Their investigation revealed a threefold reduction in fatigue life in
the presence of the oxide layer. Fig. 9 is a reproduction of an SEM
fractography analysis presented in the work of Vie et al. (Figure 13 of
their paper). The magnified view of the fracture surface of an anodised
and polymer-coated aluminium fatigue specimen (Fig. 9.b) shows that
the main crack grows from multiple crack initiation sites associated
with microcracks in the oxide layer. The multiplicity of crack initiation
sites results in a very large crack initiation zone as indicated by the
red arrows in 9.a. In contrast, in a polymer-coated aluminium fatigue

specimen with no oxide layer, the main crack propagates from a sin-
gle narrow initiation site associated with intermetallic particle failure
(Figs. 9.c and 9.d). These observations, coupled with in situ acoustic
emission during fatigue cycling, shed light on several mechanisms for
the anodised samples with the brittle alumina layer. First, multiple
cracking of the oxide layer early in the fatigue life generates multi-
ple stress concentration regions. Due to the strong adhesion between
alumina and the aluminium substrate, these microcracks in the oxide
layer promote multiple crack initiation into the underlying material.
Second, the cracks initiated in the substrate grow and converge to
form a single main crack, resulting in a large crack initiation zone.
Third, cracking assisted by brittle fracture of the oxide layer makes
easier the propagation of the main crack, a phenomenon referred to
as the “continuous initiation process”. To draw a parallel with our
study, the brittle oxygen-enriched metallic layer can be compared to
the oxide layer in Vie et al.’s work. Furthermore, Fig. 8 highlights
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** SEM observations of crack initiation for an anodised and polymer coated (APT) fatigue specimen (a,b) and a polymer coated (PT) fatigue specimen (c,d) made from 7075-T6 aluminium alloy. This figure is reproduced from the paper by Vie et al. [64].</figcaption></figure>

that the main crack initiation stems from the surface notches present in the as-built specimens. This observation raises the possibility that surface notches either contribute to crack initiation or support crack propagation by a mechanism similar to continuous initiation. Finally, from the examination of the fracture surfaces in the present study, which also show a large crack initiation zone (Fig. 7), the hypothesis of multiple cracking of the alpha-case at the notch roots leading to a single main crack distributed over a large crack initiation zone and followed by a continuous initiation process seems plausible.

More sophisticated analyses would be required to further elucidate these mechanisms. For example, *in situ* acoustic emission analysis during fatigue testing, similar to the method used by Vie et al. [64], could provide valuable insights. From the same perspective, fatigue testing coupled with *in situ* X-ray tomography setups using synchrotron X-ray sources should be carried out [65]. Finally, it would also be instructive to investigate the effects of SR, LF-HT and IF-HT conditions on specimens that have been machined and polished. Such an investigation could reveal whether, in the absence of critical surface defects, the microstructure or the presence of an alpha-case layer has a more significant influence on fatigue resistance.

### 3.5. Crack initiation site analysis

To investigate whether the presence of the alpha-case layer could promote crack initiation, the crack initiation site density was estimated in specimens under SR, LF-HT and IF-HT conditions. The idea was to study cracks that propagated during fatigue cycling, even if they did not propagate through the entire specimen section to cause failure. These are hereafter referred to as “**Secondary Fatigue Cracks**” (SFC). Fig. 10 shows an example of such a SFC, revealed after chemical etching of the longitudinal cross-section of one half of a fractured specimen tested at 450 MPa. As shown in the overall view of the specimen, a total of 6

SFC were identified and are marked with red arrows in the magnified views. SFC are oriented perpendicular to the loading direction (Mode I propagation) with a relatively small opening within the enlarged views in Fig. 10. Their initiation can often be correlated with the presence of surface notches and their propagation often extends beyond the alpha case layer.

A distinct type of crack, observed predominantly on IF-HT specimens after fracture, deserves attention. Such cracks are specifically associated with the final stage of failure, see an example marked with green arrows in Fig. 10. It is noteworthy that such cracks were only identified on the surfaces of specimens that experienced complete failure, and not on those of interrupted tests. It is therefore likely that these cracks are associated with the final stage of failure, but not with the initial stages of crack initiation or propagation. Accordingly, they are defined as “**Final Failure Cracks** (FFC)”. Fig. 11 shows some of these FFC. Their relatively large opening and limited depth, correlating with the local thickness of the alpha case, are highlighted in Fig. 11.b. These features are similar to the cracks observed in the post-failure tensile specimens in Ref. [13]. In the final stages of fatigue life, the material is subjected to significantly high stresses towards the end of crack propagation, which explains the similarities between the fracture surfaces observed in the final failure region and those in the tensile specimens. Close examination reveals a dense cluster of FFC at the edge of the final failure zone, as shown schematically in Fig. 11.c. The occurrence of such FFC highlights the potential of the alpha-case layer to reduce fracture toughness, thereby accelerating final failure. This observation is in good agreement with the findings from the tensile test results [13]. From this perspective, the FFC are not the objects of primary interest to study the effect of the alpha layer on the initiation stage and overall fatigue life. Therefore, the following paragraphs will focus exclusively on the SFC.
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** Longitudinal cross-section analysis with optical microscopy of secondary fatigue cracks (SFC) in an IF-HT specimen tested at 450 MPa. The cross-section of the specimen was polished and etched with Weck's reagent. Magnified views highlight the fatigue cracks with red arrows, while green arrows near the final fracture zone indicate the location of the final fracture cracks shown in Fig. 11.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 11.** Observation of final fracture cracks (FFC) in an IF-HT specimen tested at 350 MPa. SEM-SE image (a) and optical microscope image of a longitudinal cross-section after chemical etching with Weck's reagent to reveal the alpha-case in white (b). Schematic representation of an IF-HT fracture surface, showing the preferential locations of FFC (c).</figcaption></figure>

From the XCT data obtained after fatigue testing, it was found that FFC with a sufficiently large opening were detectable, whereas all SFC showed a rather small opening and thus could not be detected by XCT. Therefore, in order to investigate the effect of alpha-case layers

of different depths on the prevalence of SFC, a comprehensive light microscopy analysis was carried out on 11 SR, 9 LF-HT and 11 IF-HT specimens subjected to different stresses. After each test, a longitudinal cross-section of the specimen (half for fractured specimens) was
<figure><img src="image_14.png" /><figcaption>**Fig. 12.** Fatigue crack linear density measurements on SR, LF-HT and IF-HT specimens at different stresses. Schematic representation of the height *h* within which cracks are counted for each specimen (a). SFC initiation site densities measured for all specimens as a function of the maximum stress during the fatigue test (b).</figcaption></figure>

prepared. This cross-section was then polished and etched with Weck's reagent. On this cross-section of thickness *e* near the fracture area, the SFC were identified and counted within a distance *h* measured between the middle of the fracture area and the beginning of the fillet radius defined as $1.1 \times e$, as illustrated in Fig. 12.a. For interrupted and run-out specimens, the *h* value was measured over the entire gauge length, between the beginning of the two fillet radii at $1.1 \times e$, and the thickness at the centre of the gauge area. The number of SFC counted was then normalised by *h* to calculate a linear SFC initiation site density in $mm^{-1}$. The results are presented in Fig. 12.b. The analysis shows an increase in SFC initiation density with maximum stress for all conditions, a trend consistent with the understanding that higher stresses lead to an increase in the number of nucleated cracks as reported in previous studies [62]. Closer examination of Fig. 12.b indicates that the IF-HT specimens exhibit significantly higher SFC initiation site densities at stresses greater than 300 MPa. This observation supports the hypothesis that the 70 µm thick alpha case layer in the IF-HT specimens acts as a catalyst for crack initiation during fatigue loading, potentially reducing fatigue life. This is consistent with the findings of Pilchak et al. who reported a tenfold reduction in the fatigue life of a near α Ti alloy with an alpha-case layer approximately 35 ± 5 µm thick [39]. In their work this significant reduction was attributed to both a reduced resistance to crack initiation and an accelerated crack growth rate within the oxygen-enriched region. The SFC initiation site densities for the SR and LF-HT conditions were found to be comparable, with a slight increase observed in the SR specimens. This minor difference could be attributed to the presence of a harder oxygen rich alpha-case layer close to the surface of the SR specimens, despite its relatively small thickness of 10 µm (Figs. 4 and 5).

Interestingly, previous research on the same material also reported cracking of the alpha-case under tensile loading [13]. In situ micro-tensile tests in the SEM were carried out on specimens after IF-HT at 860 °C/2 h, which exhibited an alpha-case layer of $48 \pm 2$ µm at the surface. By observing the surface of the gauge at various applied tensile stresses, it was demonstrated that cracks with a very small opening began to be visible over the entire surface of the sample at 460 MPa, i.e. still in the elastic regime (with an estimated yield stress of about

$1\,000$ MPa). In fatigue testing at $R = 0.1$, multiple tensile loading occurs in the elastic regime, which is likely to facilitate rapid surface crack initiation in the presence of an alpha-case layer. This brittle layer should therefore be considered as a source of multiple potential fatigue initiation defects (i.e. cracks), the length of which is approximately equal to the alpha-case depth. Consequently, for the IF-HT condition of this study, the reduced fatigue life (except at 300 MPa) is likely to be explained by the presence of an alpha case layer of 70 µm.

The study of tensile properties showed that with alpha-case layers of $48 \pm 2$ µm (IF-HT at 860 °C/2 h), $90 \pm 5$ µm (IF-HT at 920 °C/2 h), and $185 \pm 8$ µm (IF-HT at 980 °C/2 h), the elongations to failure of the alloy decreased by an average of 25%, 50% and 80% respectively, compared to those obtained for the same LF-HT without alpha-case formation [13]. In terms of fatigue properties, it would therefore be expected that deeper alpha-case layers could lead to reduced fatigue life. Finally, regarding the limited effect of the alpha-case layer on the fatigue property for the SR condition, a comparison can be made with its negligible impact on ductility. Indeed, the elongations to failure measured for tensile specimens with the as-received surface condition subjected to IF-HT or LF-HT at 720 °C/2 h were similar. The limited effect on the fatigue properties, even at higher stresses, could be explained by the small depth of initiated cracks in the $10 \pm 1$ µm alpha case layer compared to the depth of the surface notches (see Fig. 4 and $S_v$ values in Table 3). Therefore, if an alpha-case layer is likely to be expected after a post-processing HT carried out on L-PBF fabricated parts, applying a heat treatment at 720 °C/2 h enables the complete stress relief and partial martensite decomposition of the material [16] without leading to a severe reduction of the static and dynamic properties. Even if a higher temperature subtransus HT gives better results in terms of strength-ductility trade-off [13,42], the presence of an alpha-case layer will negatively affect both ductility and fatigue life. For example, a post-LPBF HT at 860 °C results, on average, in a 25% reduction in elongation to failure [13] and a 50% reduction in fatigue life (see Fig. 6).

As illustrated in Fig. 6, the data at 300 MPa show significant scatter, with none of the conditions presenting clear advantages in fatigue properties over the others. Similarly, Fig. 12 shows that the SFC
initiation site densities are low at 300 MPa, regardless of the conditions tested. This observation suggests that a maximum applied fatigue stress of 300 MPa is not sufficient to significantly influence SFC propagation, even in the presence of a 70 ± 3 µm thick alpha case layer.

From the *in situ* micro-tensile tests observations of Ref. [13], it is considered that cracking occurs in the alpha-case layer during the early stages of the fatigue tests. Therefore, it is possible to calculate $K_{I,max}$ associated with these defects, which is the maximum value of the mode I (opening mode) stress intensity factor along the crack front. Since the thickness *a* of the alpha-case is small relative to the specimen thickness, $K_{I,max}$ can be calculated using Eq. (3) for a half penny shaped crack whose depth is equal to the alpha-case thickness *a* [66]:

$$K_{I,max} = 0.629 \times \sigma_{max} \times \sqrt{\pi \times \sqrt{\frac{\pi \times a^2}{2}}} \quad (3)$$

Applying Eq. (3) for the alpha-case thickness of 70 ± 3 µm under IF-HT conditions gives $K_{I,max}$ values approximately equal to 3.1 MPa m¹/², for a maximum applied stress of 300 MPa. Interestingly, this value of the same order of magnitude as the fatigue crack growth threshold $\Delta K_{th}$ of Ti64 alloy fabricated by L-PBF, even it was not possible to measure this property within the scope of this study. For instance, Leuders et al. reported a $\Delta K_{th}$ of 3.9 ± 0.4 MPa m¹/² at R = 0.1 for Ti64 fabricated by L-PBF and subjected to a post-processing heat treatment at 800 °C [67]. This suggests that at a maximum applied stress of 300 MPa, even the cracks initiated in the deepest 70 µm alpha case layer do not propagate due to the proximity of the calculated $K_{I,max}$ to the potential fatigue crack growth threshold of the material. Therefore, at 300 MPa, crack propagation only occurs for sufficiently critical surface defects (i.e. higher $K_{I,max}$), resulting in the high scatter in fatigue properties and the low density of SFC initiation sites observed.

The effect of an alpha-case layer of varying depth on near-net-shape L-PBF Ti64 specimens was investigated by various analyses: determination of fatigue properties, identification of killer defects by XCT, measurement of crack initiation site densities. The as-built surface condition, with most cracks initiated at the surface notch roots, was shown to have a predominant effect on the fatigue properties. Therefore, the microstructural changes that occur during sub-transus post-processing HT have a limited effect on fatigue resistance. However, Ti64 parts are subject to the formation of an oxygen-enriched brittle alpha-case layer below the surface if the HT protective atmosphere (vacuum, argon gas, ...) is not properly controlled. If this layer cannot be removed by post-HT surface treatments, this work suggests two possible scenarios:

* The first is that the alpha-case depth is small compared to the surface defect size (i.e. notch depth). In this study, the effect of a 10 ± 1 µm brittle layer was limited and did not exacerbate the detrimental effects of surface notches. However, such a small alpha-case layer could have a large effect on polished samples, according to Refs. [34,38,39].

* The second scenario is that the alpha case depth is large compared to the surface defect size. In this study, the 70 ± 3 µm brittle layer reduced the fatigue life of the near-net-shape samples by approximately half for applied stresses above the assumed fatigue crack growth threshold. Such a large alpha case layer could have a severe effect on polished specimens.

The lesson to be learnt from this study is to be cautious of alpha-case formation which can occur during the post-processing HT of near-net-shape Ti64 L-PBF parts. A large alpha-case layer (≥ 50 µm) is detrimental to both ductility and fatigue life. However, a small alpha case layer (≈ 10 µm) does not appear to be of primary concern for both tensile [13] and fatigue properties.

## 4. Conclusions

In this study, the fatigue properties of Ti-6Al-4V (Ti64) alloy fabricated by Laser Powder Bed Fusion (L-PBF) with as-built surfaces and subjected to three different post-processing heat treatments (HT) were investigated. The use of two annealing temperatures, Stress Relieving (SR) at 720 °C and HT at 860 °C, together with the use of IF-HT (industrial furnace) and LF-HT (laboratory furnace) conditions, allowed the generation of alpha-case layers of different thicknesses on the sample surface. It was therefore possible to investigate the effect of alpha-case layers on the surface embrittlement and fatigue properties of the material. The main results of this work are summarised below:

* The bulk microstructures of all samples after HT consist of a mixture of $\alpha + \beta$ phases in a chessboard pattern of prior $\beta$-grain structures observable in planes perpendicular to the building direction. The size of $\alpha$ laths and the fraction of $\beta$ phase are larger in the IF-HT and LF-HT conditions (approximately 1 µm and 4%) compared to the SR condition (approximately 0.6 µm and 2%).

* The SR and IF-HT conditions show a distinct alpha-case layer at the surface, with thicknesses estimated to be 10 ± 1 µm and 70 ± 3 µm respectively. These layers result from oxygen diffusion from the surface, leading to oxygen concentration profiles associated with increased surface hardness. In contrast, the LF-HT condition shows no discernible alpha-case layer and only a moderate increase in surface hardness.

* The fatigue properties are consistent with those reported in the literature for L-PBF Ti64 specimens with as-built surfaces. The fatigue strength at 10⁷ cycles was estimated to be around 250 MPa. Among the conditions studied, only the IF-HT condition, which has the thickest alpha-case layer, exhibited a decrease in fatigue resistance, while the SR and LF-HT conditions demonstrated similar properties.

* Under all conditions, fatigue crack initiation occurred systematically at the surface of the specimens. 3D surface characterisation using X-ray Computed Tomography (XCT) before and after fatigue testing revealed that, in most cases, cracks initiated at surface notches. This finding emphasises the primary influence of as-built surface conditions over microstructure for Ti64 components fabricated by L-PBF.

* Post-mortem examination of the specimens allowed identification and counting of secondary fatigue cracks (SFC) distributed along the height of the specimen and extending beyond the alpha-case layer. The linear densities of SFC initiation sites in specimens from the IF-HT condition were significantly higher than those in the other two conditions for applied stresses above 300 MPa. This emphasises the role of the 70 ± 3 µm deep alpha case layer in promoting crack initiation and thus reducing fatigue life.

This work shows that end-users should consider more carefully the possibility of forming an alpha-case layer during post processing HT, as this can reduce the fatigue life of the material.

### CRediT authorship contribution statement

**Quentin Gaillard:** Writing – original draft, Visualization, Investigation, Conceptualization. **Florian Steinhilber:** Writing – original draft, Visualization, Investigation, Conceptualization. **Amélie Larguier:** Investigation. **Xavier Boulnat:** Writing – review & editing, Validation, Supervision, Conceptualization. **Jean-Yves Buffiere:** Writing – review & editing, Validation, Supervision, Funding acquisition, Conceptualization. **Guilhem Martin:** Writing – review & editing, Supervision. **Sophie Cazottes:** Writing – review & editing, Supervision. **Rémy Dendievel:** Validation, Supervision, Funding acquisition. **Christophe Desrayaud:** Validation, Supervision, Funding acquisition.
Declaration of competing interest

The authors declare that they have no known competing finan-
cial 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.

Acknowledgements

This work, conducted at LGF, MATEIS and SiMaP laboratories, falls
within the framework of the AEROPRINT project, which received sup-
port from the Auvergne-Rhône-Alpes Region, France and the company
Dassault Aviation, France. The authors are grateful for their financial
support and extend special thanks to the project team at the Argonay
production site for fabricating the samples used in this study.

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