# Oxygen Ingress in Titanium and Its Alloys After High-Temperature Oxidation: A Competition Between Strengthening and Embrittlement

Quentin Sirvin, J. Genée, Benjamin Dod, Daniel Monceau, Damien Texier

► To cite this version:

Quentin Sirvin, J. Genée, Benjamin Dod, Daniel Monceau, Damien Texier. Oxygen Ingress in Titanium and Its Alloys After High-Temperature Oxidation: A Competition Between Strengthening and Embrittlement. Metallurgical and Materials Transactions A, 2025, 56, pp.1858-1874. 10.1007/s11661-025-07737-0 . hal-05004906

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# Oxygen ingress in titanium and its alloys after high temperature oxidation: a competition between strengthening and embrittlement

Quentin Sirvinª, Julien Genéeª, Benjamin Dodᵇ, Daniel Monceauᶜ, Damien Texierª

ªInstitut Clement Ader (ICA) - UMR CNRS 5312, Université de Toulouse, CNRS, INSA, UPS, Mines Albi, ISAE-SUPAERO, Campus Jarlard, 81013 Albi Cedex 09, France

ᵇAIRBUS Operation SAS, 316 route de bayonne, 31200 toulouse, France

ᶜCIRIMAT-UMR CNRS 5085, Université de Toulouse, ENSIACET, 4 Allée Emile Monso, BP 44362, 31030, Toulouse Cedex 4, France

## Abstract

High-temperature oxidation of titanium leads to the formation of an external oxide scale and oxygen ingress into the metallic titanium material. Oxygen ingress can be significant due to the high solubility of O within Ti. An oxygen-rich layer (ORL) thus forms beneath the external oxide scale, exhibiting a brittle behavior. Microtensile specimens were used in order to exacerbate surface effects, i.e., surface reactivity in the case of the oxidation of titanium. Playing with the specimen thickness and pre-oxidation durations, it was possible to evaluate the evolution of tensile strength as well as the reduction in ductility for deep extensions of ORL relative to the specimen thickness (high fraction of ORL). In addition, ultrathin specimen extraction at different locations within the ORL depth aimed at better identifying the gradient of properties within the ORL. This micromechanical approach was applied to a commercially pure titanium (CP-Ti grade 2) and to a structural titanium alloy (Ti6242s). Both strengthening and loss of mechanical properties (yield strength and ductility) were observed depending on the material and oxygen ingress. While CP-Ti demonstrated an increase in mechanical strength up to ORL representing 80 pct of the gage section, Ti6242s experienced a loss of mechanical resistance even for the shortest exposure times (the ORL representing 10 pct of the gage section).

**Keywords:** Titanium alloy, Micromechanical testing, Oxygen dissolution, Oxygen embrittlement, Mechanical modelling

## 1. Introduction

Oxygen embrittlement of titanium leads to a strong limitation in use of titanium and its alloys at high and intermediate temperatures for structural applications. Titanium can dissolve a significant amount of oxygen in the metal without necessarily changing crystallographic structure due to high solubility of O within Ti [1], the α phase having a higher solubility limit (33 pct at. of O in Ti) compared to the β phase (8 pct at. of O in Ti). When exposed to oxidative atmosphere at high temperature, titanium components form an external oxide scale and an oxygen-rich layer (ORL)

<footnote>*Corresponding author
Email address: damien.texier@mines-albi.fr (Damien Texier)</footnote>
in the metallic Ti beneath the external oxide scale, in the form of a gradient of oxygen concentration with maximal oxygen content at the metal/oxide interface and minimal content in the bulk. It is worth mentioning that the ORL extension is much greater than the external oxide scale [2–4]. Therefore, the mechanical response and brittleness of the ORL is a central question to better understand the mechanical integrity of oxidation-affected titanium structural components.

The effect of oxygen on different mechanical and physical properties has been extensively reported in the literature. Oxygen insertion within Ti can improve the hardness of the material [4–18] and increase the Young's modulus [7, 11, 16–22]. Such properties evolutions are beneficial for tribological aspects due to superficial strengthening. However, oxidation-affected titanium components subjected to tensile stresses are prone to early cracking, leading to a significant debit in ductility [5–9, 12–14, 18, 20, 23–34]. As far as the beneficial role of oxygen is concerned, the strengthening effect of Ti and its alloys due to O insertion is controversial, depending either on the oxygen content and alloying elements. Barkia *et al.* [35] and Ren *et al.* [28] evidenced change in critical resolved shear stress for the different slip systems, leading to a significant increase in tensile strength of low-content oxygen titanium materials. Such strengthening effects were also reported in the literature [5, 7, 11, 19, 20, 23–25, 28, 31, 33, 35–39]. However, brittle features on fracture surfaces are observed for oxygen contents greater than 0.5–0.6 pct at. [18, 30, 34] for structural Ti alloys. At an oxygen concentration greater than 3.0 pct at., fracture surfaces are clearly brittle with crack paths not necessarily depending on the local crystallographic features (crack path perpendicular to the macroscopic loading direction) [18]. For oxygen concentrations ranging from 0.5-1.0 pct at. to 3.0 pct at., fracture surfaces remain brittle but with a much more intricate crack path following microstructural features depending on the Ti alloy [18, 30]. In addition, Amann *et al.* [33] showed that titanium with an oxygen content up to 2.4 pct at. remains ductile, but significantly strengthened. Therefore, the effect of oxygen insertion on the brittleness of Ti and Ti alloys is assumed to be different due to chemical aspect by atom substitution and interaction. For instance, aluminum [37] and zirconium [33, 36, 38] were found to significantly improve both mechanical strength and ductility of high-content oxygen titanium materials. On the other side, notch sensitivity of Ti-based materials and the presence of a mechanical-property gradient (oxygen-graded titanium due to oxidation), leading to stress/strain concentration under mechanical loading, could also influence the early failure of oxidation-affected Ti-based materials. This point is of high importance to better understand the integrity of components in service. Mechanical characterizations on conventional size specimens are well suited to catch debit in ductility due to oxygen ingress. Subsequently, the fatigue endurance and lifetime of oxidation-affected titanium components subjected to thermomechanical stresses strongly drop-off in oxidizing atmospheres [12–14, 23, 26, 32]. However, conventional size specimens find limitation in the determination of tensile strength evolution due to low fraction of ORL (low extension of the ORL compared to the specimen thickness) in the specimen gage section of pre-oxidized specimens. Modification of the mechanical strength as a function of the fraction of ORL or within the ORL is trivial to assess for further mechanical predictive model (behavior and/or lifetime). More recently, investigation of pre-oxidized ultrathin specimens, i.e., with a thickness comparable to the
<table><caption>Table 1: Nominal Composition of the Commercially Pure Ti (CP-Ti) (in Pct (Wt)).</caption><thead><tr><th>Element % (wt)</th><th>Ti</th><th>Al</th><th>Mo</th><th>Sn</th><th>Si</th><th>Zr</th><th>O</th><th>Fe</th><th>N</th><th>C</th></tr></thead><tbody><tr><td>CP-Ti</td><td>Bal.</td><td>-</td><td>-</td><td>-</td><td>-</td><td>-</td><td>0.14</td><td>0.07</td><td>0.00</td><td>0.00</td></tr><tr><td>Ti6242s</td><td>Bal.</td><td>6.03</td><td>2.08</td><td>1.91</td><td>0.09</td><td>4.03</td><td>0.10</td><td>-</td><td>0.03</td><td>-</td></tr></tbody></table>

extension of the ORL, evidenced change in both mechanical strength and ductility [18]. A similar approach has been conducted on a commercially pure titanium (CP-Ti) and on high temperature structural Ti-based alloy (Ti6242s) to assess the change in mechanical properties in a low yield strength/high work-hardening material (CP-Ti) compared to a high yield strength/low work-hardening material (Ti6242s).

In the current work, the use of ultrathin specimens aims to evidence the effect of oxygen ingress on the oxygen embrittlement response of the CP-Ti material and Ti6242s alloy, both having very different mechanical response. Two tensile testing approaches were used to probe the mechanical response of the ORL: (i) tensile of pre-oxidized microtensile specimens of varying thicknesses and exposure times at elevated temperatures, (ii) extraction of ultrathin specimen within the ORL then tensile testing. The tensile behavior of oxidation-affected CP-Ti material and Ti6242s Ti-based alloy was compared. EPMA analyses were conducted on cross-sections in order to evaluate the oxygen ingress for the different pre-oxidation conditions.

## 2. Experimental Procedures

### 2.1. Material

A commercially pure titanium (CP-Ti) and a Ti6242s alloy were used in the present study. The chemical composition of both the materials is presented in Table 1. The oxygen and nitrogen contents were measured using a Oxygen/Nitrogen/Hydrogen Elemental Analyzer using the inert gas (He) fusion technique.

Flat microtensile specimens were machined out of the rolled plates. CP-Ti material was extracted from thin plates using water-jet machining and Ti6242s from thicker blocks using electrodischarge machining. The longitudinal direction (LD) of the microtensile specimens was perpendicular to the rolling direction (RD). All the microtensile specimens, CP-Ti and Ti6242s specimens, were then gently rectified to reduce the machining-affected region. A progressive manual gritting was done on all the specimen surface using P1200 grit papers. The specimens have a dogbone-shape geometry with a gage width of 2.5 mm and gage length of 5.0 mm, as detailed in Ref. [18]. Specimen edges were ground down to a P1200 SiC grad paper to eliminate the water-jet- or EDM-affected region. A high-tolerance polishing technique, i.e., a custom precision Jig with a polishing machine [40, 41], was used to thin down the microtensile specimens to different thicknesses in the range of 800 to 50 µm with a surface finish of P2400 SiC grit paper. Specimens are carefully glued on a rectified glass support using thermal wax. The custom precision
Jig, also rectified, allows maintaining the glass support with glued specimens using vacuum suction. This allows an improved contact between the glass support and the custom precision Jig and a high-tolerance polishing operation. The thickness variation along the specimen gage is less than ± 1 µm (validation in a previous study using optical observations on a specimen cross-section). More details on the preparation technique can be found in Ref. [42]. This preparation technique aims to obtain a repeatable surface finish necessary for pre-oxidation purpose and a highly controlled geometrical tolerance in terms of sample thickness for further mechanical testing.

## 2.2. Pre-oxidation Treatments for Different Microtensile Characterizations

The ORL extension was considered here as the peripheral region beneath the external oxide scale with an oxygen uptake greater than 0.5 pct at.. The fraction of ORL referred to the ratio of peripheral ORL extension (surface) over the gage section. A specimen fully affected by oxygen, i.e., an oxygen uptake higher than 0.5 pct at. on the whole gage section, has an ORL fraction of 100 pct. Two micromechanical approaches were conducted in the present study, leading to slight differences in pre-oxidation treatments and specimen preparation:

* (i) **Tensile testing of different ORL fractions:** Microtensile testing at room temperature of pre-oxidized specimens with different thicknesses and ORL extensions to investigate the role of different ORL fractions on the mechanical response of oxygen-graded specimens. In this approach, microtensile specimens were thinned down to a given thickness, as detailed in the section 2.1. A careful cleaning in acetone then in ethanol using an ultrasonic bath for 15 minutes each was done on the specimens before the pre-oxidation treatment. Pre-oxidation treatments were performed in a Nabertherm C290 furnace with static atmosphere without air stirring (natural convection) in the same batch for a given pre-oxidation duration. Limiting contact was ensured during the pre-oxidation treatment by hanging the microtensile specimens at the two head radii with platinum wires. The temperature inside the furnace was monitored using a K-type thermocouple in the vicinity of the specimens. This procedure aims to ensure reproducibility in atmosphere condition for better controlling oxidation, and more specifically oxygen ingress [3]. For the CP-Ti material, pre-oxidation exposures of 1, 10, 50, and 100 h at 700 °C were investigated. Due to higher oxidation resistance, the Ti6242s specimens were exposed at 700 °C for 50, 100, 200, and 500 h. For both the materials, specimen thicknesses were of 800, 400, 200, 100, and 50 µm, leading to ORL fraction ranging from 0 to 100 pct.

* (ii) **Ultrathin samples inside the ORL:** Microtensile testing at room temperature of ultrathin specimens extracted at different distances from the metal/oxidation interface to screen the gradient of tensile properties as a function of the oxygen concentration. Microtensile specimens with a thickness of 400 µm were first prepared with the high-tolerance polishing technique, then pre-oxidized at 700 °C for 50 h and 200 h for the CP-Ti material and the Ti6242s alloy, respectively. Cleaning then pre-oxidation treatments followed the same procedure as the one described before for "tensile testing of different ORL fractions". Microtensile specimens were
then polished on both sides in order to extract ultrathin specimens, i.e., 50 µm-thick microtensile specimens, at different depths from the metal/oxidation interface within the ORL. The specimen thinning was based on the high-tolerance polishing technique. Successive polishings and observations aimed to first locate the upper surface of the microtensile specimen at the exact location with regard to the ORL extension depth. The second face was then polished to obtain a 50 µm-thick microtensile specimen.

## 2.3. Microtensile Testing Paired with Digital Image Correlation (DIC) Techniques

Microtensile testing was conducted at room temperature using an Instron 5800R electromechanical tensile testing machine with BlueHill2 software. A load-cell of 500 N was used for microtensile specimens of 100 and 50 µm thicknesses, and of 5 kN for microtensile specimens with greater thicknesses. Based on the difference in specimen thickness and thus apparent stiffness when using such a wide range of specimen thicknesses, a non-linear cross-head displacement rate was applied to obtain a quasi-constant strain rate $\dot{\epsilon}$ in the gage area of the specimen. This non-linear displacement motion is designed to compensate for the stiffness of the tensile line, which has a significant effect in the elastic regime, as titanium is particularly sensitive to strain-rate effects at room temperature. The displacement rate $v_{cross-head}$ was purposely monitored faster in the elastic domain compared to the one in the plastic domain, with a smooth transition from elastic to plastic-displacement control (see Eq. 1):

$$v_{cross-head} = \dot{\epsilon} + \alpha_1 \left( 1 - \frac{1}{1 + e^{\alpha_2 - t}} \right) \quad (1)$$

with t the time in seconds, $\alpha_1$ and $\alpha_2$ coefficient to trigger the elasto-plastic transition. Calibration tests were conducted prior to tensile testing in order to determine $\alpha_1$ and $\alpha_2$ coefficients as a function of the specimen thickness. Tensile tests were thus conducted at a quasi-constant strain rate of $1 \times 10^{-4} s^{-1}$. In order to measure the macroscopic deformation of the specimens, non-contact extensometry was used based on digital image correlation techniques (DIC). A speckle pattern was thus necessary in order to track features at the surface of the sample during the deformation process. White then black aerosol paints were applied onto the microtensile specimens to have a homogeneous white background with randomly distributed black dots.

For the macroscopic optical extensometry, the image correlation bench was constituted of a Pike 5B/C camera (5 MPx) from Allied Vision, a Sigma DG Macro lens, a lighting system composed of a Dedolight spot and a KERN Optics light ring. Digital image acquisition was carried out using the VicSnap software from Correlated solution. Tensile test data, i.e., time, load, and cross-head displacement, were simultaneously recorded for each image acquisition with a National Instrument USB module (USB-6009-779026-01). The acquisition frequency was adapted between 0.5 and 12 Hz depending on the intended brittleness of the microtensile specimen (500 to 1500 images per tests). Kinematic fields were computed using PYXEL Python library developed by Passieux et al. [43]. The average strain within the gage length was then calculated from the displacement along the longitudinal direction and considered as the macroscopic strain.
## 2.4. Microprobe Analyses (EPMA)

Microprobe analyses (EPMA) were conducted to assess oxygen profiles for all the investigated conditions, i.e., the oxygen content at different depths from the sample surface. EPMA samples were mounted in a low shrinkage cold epoxy mounting for cross-sectional observations. EPMA samples were then prepared using polishing from P320 to P4000 SiC grit papers, diamond finishing using 1 µm-diamond paste. To limit native oxidation of the sample and alter the oxygen measurements, a thin but continuous carbon layer was deposited using a Leica EM ACE600 sputter coater. Samples were stored in a vacuum chamber (10⁻⁵ mbar) before EPMA analyses. Despite particular preparation, cleaning, and storage cares, surface contamination cannot be fully avoided, as described in References [3, 18, 30]. Therefore, the oxygen uptake was assumed relative to the oxygen content in the bulk by correcting the EPMA intensity by subtracting the intensity relative to oxygen in the sample bulk. The oxygen content presented in the paper thus not correspond to the absolute oxygen content, but the oxygen uptake due to oxygen dissolution during oxidation. EPMA analyses were performed using a CAMECA® SXFive FE at an acceleration voltage of 10 kV and a probe current of 20 nA. Through-thickness profiles were repeated three times per pre-oxidation conditions on 200 µm-thick specimens with a measurement step of 1 µm near the surface then of 5 µm for depths greater than 20 µm. This one-dimensional diffusion approximation is commonly used in the literature to describe oxygen ingress in titanium and titanium alloys [3, 4, 18, 22, 44–46]. For information, the oxygen concentration in bulk material prior to oxidation was measured using instrumental gas analysis (IGA).

## 2.5. Scanning Electron Microscopy

Electron backscattered diffraction (EBSD) analyses were carried out to better inform the microstructure and crystallographic texture of both the materials. EBSD characterizations were conducted using a JEOL JSM 7100F scanning electron microscope equipped with a CMOS Symmetry S2 EBSD detector (Oxford Instruments) operating at an acceleration voltage of 20 keV. The step size of the EBSD maps varied between 0.5 and 1.5 µm depending on the coarseness of the microstructure. EBSD post-treatments were done using the MTex library [47].

After tensile tests, fracture surfaces were observed for all the pre-oxidation and specimen thickness conditions using a TM3030Plus tabletop scanning electron microscope from HITACHI® in a secondary electron (SE) mode at 15 kV. Image analyses were performed on the fracture surface to quantify the extension of the brittle region using Fiji software [48].

# 3. Experimental Results

## 3.1. Microstructural Characterization

EBSD characterizations were carried out to document microstructural evolutions occurring during pre-oxidation treatments.
<figure><img src="image_2.png" /><figcaption>Figure 1: Microstructure evolution of the CP-Ti after different heat-treatments: Grain microstructure represented as an inverse pole figures (IFP) EBSD map in reference to the loading direction (LD) for (a) as-received microstructure, (b) the pre-oxidized sample at 700 °C for 10 h, and (c) the pre-oxidized sample at 700 °C for 100 h.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>Figure 2: Microstructure evolution of the Ti6242s material after different heat-treatments: Grain microstructure represented as an inverse pole figure (IFP) map in reference to the loading direction (LD) for (a) as-received microstructure, (b) the pre-oxidized sample at 700 °C for 50 h, and (c) the pre-oxidized sample at 700 °C for 500 h.</figcaption></figure>

At reception, the CP-Ti material had a 15-µm equiaxed-grain microstructure, as shown in Fig. 1(a). Some Fe-rich β phases are sparsely distributed at grain boundaries. The microstructure evolution of the CP-Ti material after pre-oxidation was characterized in the specimen bulk far from the ORL. After 10 and 100 h at 700 °C, the average grain size increased to 35 µm and 75 µm, respectively (Fig. 1(b) and (c)). Despite grain growth, the crystallographic texture was found similar for all the different thermal exposure treatments.

For the Ti6242s material, the microstructure is bimodal with equiaxed α nodules surrounded by α-β metastable regions. It is worth noting that neither grain/nodule size nor crystallographic texture evolves regardless of the thermal treatment duration (see Fig. 2). The Ti6242s material is significantly textured due to rolling with microtextured regions (MTR) morphologically elongated with the rolling direction, corresponding to the further loading direction (LD).

## 3.2. Oxygen Profile Characterization

Microprobe analyses (EPMA) were conducted on thick samples in order to measure the oxygen uptake and ORL extension due to oxygen ingress. Due to greater oxidation resistance of the Ti6242s alloy, longer exposure times were
<figure><img src="image_4.png" /><figcaption>Figure 3: Oxygen profiles measured using microprobe analyses (EPMA) for all the pre-oxidation conditions at 700 °C for (a) the CP-Ti material and (b) the Ti6242s alloy.</figcaption></figure>

investigated compared to the CP-Ti material. The oxygen concentration at the metal/oxide interface $C_s$ was found similar for a given material, regardless of the pre-oxidation time. $C_s$ was found greater for the CP-Ti material ($C_s = 30$ pct at.) compared to the Ti6242s ($C_s = 17.0-18.5$ pct at.). In addition, ORL extension increased with the pre-oxidation time and can reach up to 50 and 95 µm for the longer exposure times for the CP-Ti material (100 h) and Ti6242s alloy (500 h), respectively. For quantitative analyses, the coefficient of diffusion of O within Ti was calculated for both the materials using a unidirectional solid-state diffusion scheme, as described in the Fick's law [49]. Based on the large specimen thickness (400 µm) compared to the ORL extension, the 1D diffusion in a semi-infinite medium approach can be used. The oxygen profile within titanium can thus be described using the following equation:

$$ \frac{C_x - C_0}{C_s - C_0} = 1 - erf(\frac{x}{\sqrt{4Dt}}) \qquad (2) $$

with $C_x$, the concentration of the species that diffuses at position $x$ after a time $t$, $C_0$, the concentration in the specimen bulk, and $D$, the diffusion coefficient of the species in the medium. The coefficient of diffusion of O within Ti was calculated from each individual profile individually, for each profile per pre-oxidation condition. The average coefficient of diffusion and scattering of O within Ti for the different pre-oxidation conditions was found of $D = 5.6 \pm 0.3 \cdot 10^{-16} \text{ m}^2 \cdot \text{s}^{-1}$ for the CP-Ti material and $3.3 \pm 0.5 \cdot 10^{-16} \text{ m}^2 \cdot \text{s}^{-1}$ for the Ti6242s alloy. These coefficients of diffusion are in good agreement with the one reported in the literature [50]. Despite difference in $C_s$ (30 pct at. and 18 pct at. for the CP-Ti material and the Ti6242s alloy, respectively), the coefficient of diffusion of O within Ti was found very similar for both the materials, slightly lower for the Ti6242s alloy.

### 3.3. Tensile Testing of the Non-oxidized CP-Ti and Ti6242s Specimens

Prior to investigate the effect of the pre-oxidation on the mechanical properties of specimens with different specimens, the effect of the specimen size of non-oxidized specimens was studied. Tensile tests were first performed on
<figure><img src="image_5.png" /><figcaption>Figure 4: Stress-strain tensile curves for non-oxidized CP-Ti specimens: (a) Specimen thickness effect on the tensile response of the as-received material, (b) Aging at 700 °C and specimen thickness effects on the tensile response.</figcaption></figure>

the as-received material for specimen thicknesses ranging from 800 to 50 µm. Engineering stress-strain curves for the CP-Ti material and the Ti6242s alloy are shown in Figs. 4(a) and 6, respectively. For the CP-Ti material, microtensile specimens experienced similar tensile behavior, except the 50 µm-thick specimen. The latter specimens have a lower strain-to-failure due to mesoscopic strain localization, thus diluting the necking region in the whole gage length. The 0.2 pct yield strength was calculated for all the specimens and are illustrated in Fig. 5. It is worth noting that the yield strength was constant in the studied thickness domain. As-received CP-Ti specimens had at least three grains in the thickness for the thinnest specimens. Based on the important change in grain size for the CP-Ti material after high temperature exposure, size effect was also studied for two thermal exposures at 700 °C: 10 and 100 h (see Fig. 4(b)). The yield strength was found to be affected both by grain size and by thickness reduction due to Hall-Petch effects and a polycrystalline-to-multicrystalline behavior, respectively. The yield strength decreased with grain size increase as well as with reduction of the specimen size (Fig. 5).

For the Ti6242s material, an effect of the specimen thickness was found, the 100- and 50-µm thick specimens demonstrating lower tensile strength compared to thicker specimens. It is worth noting that the CP-Ti material and the Ti6242s alloy do not exhibit the same tensile behavior. The CP-Ti has a low yield strength and high strain-hardening behavior, while the Ti6242s alloy has a high yield strength and a low strain-hardening behavior. The characterization of these tensile properties is important to further compare the evolution or properties due to pre-oxidation.

### 3.4. Tensile Testing of Different ORL Fractions Using Pre-oxidized CP-Ti and Ti6242s Specimens

Tensile tests were performed on the different pre-oxidized specimens with different thicknesses. Based on the testing conditions, the fraction of ORL in the gage section varied from 0 to 100 pct from specimen to another one. Engineering stress-strain curves are illustrated in Figs. 7 and 9 for all the pre-oxidation conditions. Experiments were
<figure><img src="image_6.png" /><figcaption>Figure 5: Evolution of the yield strength of the as-received and aged CP-Ti specimens as a function of the specimen thickness (Aging was done at 700 °C).</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Figure 6: Stress-strain tensile curves for non-oxidized Ti6242s specimens: Specimen thickness effect on the tensile response of the as-received material</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>Figure 7: Stress-strain tensile curves for different specimen thickness - pre-oxidation time conditions for the CP-Ti material: (a) 1 h, (b) 10 h, (c) 50 h, and (d) 100 h. It is worth mentioning that 50 µm-thick specimens are not displayed for the 50 h pre-oxidation condition due to specimen failure before tensile testing (similar for 50 µm- and 100 µm-thick specimens for the 50 h pre-oxidation condition)</figcaption></figure>

gathered per pre-oxidation duration to better discriminate the role of size effects, i.e., the role of the ORL fraction, on the tensile behavior for both the CP-Ti material and the Ti6242s alloy.

For the CP-Ti material, pre-oxidation times as short as 1 or 10 h at 700 °C lead not only to a strengthening of the oxygen-graded materials but also to a decrease of the strain-to-failure with the specimen thickness reduction (Fig. 7(a) and (b)). Furthermore, the uniform deformation decreased also with the specimen thickness reduction. For longer exposure times, the tensile strength increased then decreased with the specimen thickness reduction (Fig. 7(c) and (d)). After 50 hours of pre-oxidation at 700 °C, specimens thicker than 100 µm showed some plasticity, though their ductility decreased as the specimen thickness decreased. Specimens with a thickness of 50 µm exhibited fully brittle behavior under tensile loading. A few specimens had even entirely brittle fracture surfaces: specifically, those with thicknesses of 50 µm and 100 µm after 100 hours of oxidation, and the 50 µm specimens after 50 hours.

Fracture surfaces of all the specimens were observed using a SEM in a secondary electron mode. For representa-
tion purposes, fracture surfaces of the 100 µm-thick specimens are depicted in Fig. 8 for pre-oxidation times of 1, 10, 50, and 100 h. For the 1 hour pre-oxidation condition, the ORL is small enough not to show brittle features (Fig. 8(a)). For the 10 h and 50 h pre-oxidation conditions, a brittle region is observed in the peripheral region of the fracture surfaces, while the specimen core remains ductile with the occurrence of typical dimples. The peripheral brittle region is divided into two regions, i.e., a smooth cleaved region in the extreme surface (region highlighted in red in Fig. 8) and a brittle but lamellar region (region highlighted in yellow in Fig. 8) showing an intermediate response between the very brittle and the ductile region (region highlighted in blue in Fig. 8). The extension of the brittle region increased with the pre-oxidation time. For the 100 µm-thick specimen exposed at 700 °C for 100 h, the fracture surface is nearly fully brittle. This result is in good agreement with the macroscopic stress-strain curve illustrated in Fig. 7(d). For this latter pre-oxidation condition, both intragranular and intergranular features, i.e., grain cleavage and grain boundary cracking, are found on the fracture surface.

Similar investigations were conducted for pre-oxidized Ti6242s specimens. Engineering stress-strain curves are depicted in Fig. 9. For all the pre-oxidation conditions, both increase in pre-oxidation time and decrease in specimen thickness lead to a decrease of both the mechanical strength and the strain-to-failure. The dashed black curve, corresponding to the stress-strain curve of the 800 µm-thick Ti6242s alloy in its as-received form, was added to all the graphs in Fig. 9 for direct comparison. Specimens with a thickness of 800 µm show reduced yield strength and lower strain-to-failure compared to the as-received material. Because of the oxygen ingress, several specimens broke prior to tensile testing, either by pre-oxidation (thinnest specimens) or by handling. The decrease in strain-to-failure and tensile strength due to pre-oxidation and thickness reduction is significant compared to the CP-Ti material.

Fractured specimens were observed in order to characterize the extension of the brittle region of the pre-oxidized Ti6242s specimens (Fig .10). As for the CP-Ti material, a brittle region in the peripheral region was found for the Ti6242s material. This brittle region also increased with the pre-oxidation time. The very brittle region in the extreme surface (region highlighted in red in Fig .10) was not as smooth as for the CP-Ti material, probably due to the more complex and two-phase microstructure of the Ti6242s alloy. Due to the pre-oxidation conditions and studied thicknesses, several specimens had a fully brittle fracture surface, i.e., 50 µm-thick/50 h, 100 µm-thick/100 h and thinner, 100 µm-thick/200 h and thinner, and 200 µm-thick/500 h and thinner.

The yield strength for all the pre-oxidation conditions and materials was calculated as the 0.2.pct offset yield strength or the maximal strength for specimens having a fully elastic behavior. The evolution of the yield strength as a function of both the specimen thickness and pre-oxidation times is depicted in Fig. 11(a) and (b) for the CP-Ti material and the Ti6242s alloy. For the CP-Ti material, the yield strength increased with the oxidation time and specimen thickness reduction for specimens thinner than 400 µm due to oxygen ingress and strengthening (Fig. 11(a)). For the 800 µm-thick specimens, an increase then decrease of the mechanical strength is observed due to competition effects between O strengthening and grain size increase leading to a decrease in yield strength according to the so-called Hall-Petch effect [51, 52]. For longer oxidation times and thin specimens, embrittlement due to oxygen ingress is significant enough to lead to a decrease in tensile strength with specimen thickness reduction. For the Ti6242s
<figure><img src="image_9.png" /></figure>

<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /><figcaption>Figure 8: Fracture surfaces of a 100-µm-thick specimen for different pre-oxidation conditions at 700 °C: (a) 1 h, (b) 10 h, (c) 50 h, and (d) 100 h.</figcaption></figure>
<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /></figure>

<figure><img src="image_15.png" /></figure>

<figure><img src="image_16.png" /><figcaption>Figure 9: Stress-strain tensile curves for different specimen thickness - pre-oxidation time conditions for the Ti6242s material: (a) 50 h, (b) 100 h, (c) 200 h, and (d) 500 h. As for the CP-Ti material, specimens that are not displayed for all the pre-oxidation conditions are due to specimen failure before tensile testing.</figcaption></figure>
<figure><img src="image_17.png" /><figcaption>Figure 10: Fracture surfaces of a 200-µm-thick Ti6242s specimen for different pre-oxidation conditions at 700 °C: (a) 50 h, (b) 100 h, (c) 200 h, and (d) 500 h.</figcaption></figure>
alloy, pre-oxidation leads to a decrease in mechanical strength even for the thicker specimen (Fig. 11(b)). This effect, evidenced for 800 µm-thick specimens, is even more pronounced for longer exposure times and thinner specimens.

Image analyses were performed on the fracture surfaces of all the tensile-tested specimens. The extension of the peripheral brittle region was thus measured for all the pre-oxidation conditions. The extension of the brittle region as a function of the pre-oxidation time for the CP-Ti material and the Ti6242s alloy is shown in Fig. 11(c) and (d), respectively. The black dashed curve corresponds to the fit of a power-law equation for the material of interest, while the gray dashed curve corresponds to the fit of the second material for direct comparison. The extension of the brittle region was not affected by the specimen thickness. It increases with the pre-oxidation time for both the materials. For the Ti6242s, the extension of the brittle region follows a parabolic behavior with an extension coefficient $kp_{Brittle Ti6262s}$ of $7.5 \cdot 10^{-15} \text{ m}^2 \cdot \text{s}^{-1}$ (black dashed curve in Fig. 11(d)). For the CP-Ti material, the extension of the brittle region showed a deviation from the parabolic behavior. Using a power law to fit the data, an exponent close to 3/4 is needed to better fit the measured regions. The extension coefficient $kn_{Brittle CP-Ti}$ of $8.3 \cdot 10^{-12} \text{ m}^{4/3} \cdot \text{s}^{-1}$ (black dashed curve in Fig. 11(c)).

After identification of the extension of the brittle region for the different pre-oxidation conditions, it is possible to evaluate the fraction of gage section affected by oxygen embrittlement. This ratio is considered as two times the extension of the brittle region divided by the specimen thickness. The yield strength previously evaluated for all the pre-oxidation conditions was plotted as a function of the fraction of oxygen embrittlement. Interestingly, the CP-Ti material and the Ti6242s alloy have a different response to oxygen ingress. For the Ti6242s alloy, the yield strength decreased in a nearly linear manner with the increase of the fraction of the brittle region extension (Fig. 11(f)). Yield strength was roughly gathered for a given fraction of the brittle region extension, except when the pre-oxidation is too important. For the CP-Ti material, an increase of the yield strength is observed when increasing the fraction of the brittle region extension (Fig. 11(e)). It is worth noting that an offset in yield strength is observed for longer pre-oxidation time, probably due to the concomitant contribution of the coarsening of the grain microstructure. For fully brittle specimens, the yield strength shows a decrease in mechanical strength.

### 3.5. Tensile Testing of Ultrathin Samples Inside the ORL

Ultrathin specimens, i.e., 50 µm-thick specimens, were extracted at different locations from the metal/oxide interface for the pre-oxidized CP-Ti material and the Ti6242s alloy. The pre-oxidation conditions were different depending on the studied material: 50 h - 700 °C for the CP-Ti material and 200 h - 700 °C for the Ti6242s alloy. The extension of the ORL (resp. brittle region) was approximately of 45 (resp. 42 µm) and 70 µm (resp. 72 µm) based on the EPMA analyses for the CP-Ti material and the Ti6242s alloy, respectively. For the CP-Ti material, ultrathin specimens located at 60 or 90 µm exhibit the same mechanical behavior (Fig. 12(a)). Going a little closer near the metal/oxide interface leads to an increase of the yield strength without significantly changing the ductility. For the specimens extracted right close to the metal/oxide interface or 20 µm beneath, the tensile strength increases at the expense of the strain-to-failure, specimens having a nearly brittle behavior. For the Ti6242s alloy, a similar behavior is observed,
<figure><img src="image_18.png" /></figure>

<figure><img src="image_19.png" /></figure>

<figure><img src="image_20.png" /></figure>

<figure><img src="image_21.png" /></figure>

<figure><img src="image_22.png" /></figure>

<figure><img src="image_23.png" /></figure>

Figure 11: Evolution of the yield strength of the as-received and pre-oxidized specimens as a function of the specimen thickness: (a) Pre-oxidized CP-Ti material, (b) Pre-oxidized Ti6242s material. Evolution of the extension of the brittle region of the pre-oxidized specimens as a function of the pre-oxidation time: (c) Pre-oxidized CP-Ti material, (d) Pre-oxidized Ti6242s material. Evolution of the yield strength of the as-received and pre-oxidized specimens as a function of the fraction of the brittle region extension: (e) Pre-oxidized CP-Ti material, (f) Pre-oxidized Ti6242s material.
<figure><img src="image_24.png" /><figcaption>Figure 12: Stress-strain tensile curves for different specimen thickness - pre-oxidation time conditions: (a) Pre-oxidized CP-Ti material, (b) Pre-oxidized Ti6242s material. Evolution of the yield strength of the as-received and aged specimens as a function of the specimen thickness: (c) Pre-oxidized CP-Ti material, (d) Pre-oxidized Ti6242s material.</figcaption></figure>

except that the samples in the vicinity of the metal/oxide interface were not capable to be tested due to breaking when handling (Fig. 12(b)). The yield strength of the different ultrathin tensile tests were measured and plotted in Figs. 12(c) and (d) for the CP-Ti material and the Ti6242s alloy, respectively. The yield strength was maximal in the vicinity of the metal/oxide interface and progressively decreased down to the specimen core. The yield strength stabilized when the ultrathin specimen is extracted out of the ORL (resp. brittle region).

## 4. Discussion

The effect of the high temperature pre-oxidation on the tensile behavior of a commercial high purity titanium (CP-Ti) material and of a Ti6242s alloy was investigated under tensile testing conditions to emphasize on the mutual contribution of the oxygen strengthening and oxygen embrittlement. Ultrathin specimens and different pre-oxidation
times were used to have high fractions of ORL within the specimen gage, conversely to conventional size tensile specimens. Consequently, a common decrease in strain-to-failure was observed with oxygen ingress. More importantly, changes in mechanical strength were successfully measured. In the present investigation, the CP-Ti material and the Ti6242s alloy demonstrate a different mechanical response to oxygen ingress.

As far as the oxidation behavior is concerned, chemical profiles, i.e., a one-dimensional diffusion approach, were used to describe the oxygen ingress into the CP-Ti material and the Ti6242s alloy. The one-dimensional diffusion approximation is commonly used in the literature to describe oxygen ingress in titanium and titanium alloys [3, 4, 18, 22, 44–46]. In addition, Magazzeni et al. [15] revealed no significant difference in oxygen concentration from grains to grains into the ORL of a pre-oxidized CP-Ti material using 2D EPMA mapping. This confirms the present one-dimensional diffusion approach; the heterogeneous diffusion at the microstructure scale is considered very little at grain boundaries [53] or based on the grain orientation due to anisotropic diffusion [54]. When comparing the oxygen ingress into CP-Ti material and the Ti6242s material, the CP-Ti material experienced a slightly faster diffusion of oxygen within titanium at 700 °C compared to the Ti6242s alloy, as demonstrated by chemical profile analyses using EPMA (Fig. 3). The diffusion coefficient of O within Ti at 700 °C is of $5.6 \pm 0.3 \cdot 10^{-16} \text{ m}^2 \cdot \text{s}^{-1}$ and $3.3 \pm 0.5 \cdot 10^{-16} \text{ m}^2 \cdot \text{s}^{-1}$ for the CP-Ti material and the Ti6242s alloy, respectively. Despite difference in oxygen concentration at the metal/oxide interface, $C_s = 30$ pct at. for the CP-Ti material and $C_s = 18$ pct at. for the Ti6242s alloy, these coefficients of diffusion are in the same order of magnitude and can be considered as equivalent due to discrepancy in such coefficient measurements based on potential experimental errors or misevaluation of the oxygen concentration at the metal/oxide interface. The present values are also relatively close to the ones reported by Vache et al. [50] ($D_O = 4.1 \cdot 10^{-16} \text{ m}^2 \cdot \text{s}^{-1}$). The difference in $C_s$ concentration at the metal/oxide interface can result from nitrogen dissolution being potentially greater for Ti6242s materials, as reported by Gautier et al. [46]. Solutes, both interstitial and substitution elements, can affect the diffusion properties of the α-Ti and β-Ti phases, as reported in the literature based on ab-initio calculations [55, 56]. For instance, molybdenum solutes lead to lowered diffusivity due to negative interaction between Mo and O, thus attracting and trapping diffusing oxygen atoms. In a similar way, alloying elements should affect the solubility of O within the alloy and participate to the difference in diffusion in a given phase. In addition, Ti6252s alloy also contains β phase, in which the solubility of oxygen is significantly lower than in the α phase [57, 58].

The CP-Ti material experienced an overall strengthening behavior (increase in the stress-strain response) when increasing the fraction of ORL on pre-oxidized specimens (see Fig. 11). Indeed, an increase up to 63 pct is observed for a fraction of ORL of 80 pct. For ORL fraction above 80 pct, the embrittling effect due to oxygen ingress is predominant, and the apparent mechanical strength significantly dropped, as shown by the 100 µm-thick/100 h CP-Ti specimen. 50 µm-thick specimens pre-oxidized more than 50 h were even found too brittle to be mechanically tested. They broke during either due to handling or mounting on the tensile test rig. This confirms the severe competition between the concomitant strengthening and embrittling effects of oxygen on titanium. CP-Ti materials have been intensively investigated in the literature and the beneficial contribution of oxygen on the tensile strength increase was
already demonstrated for relatively low oxygen contents (0.7 pct atomic maximum and even nearly 1.5 pct atomic [16], and 2.4 pct atomic [33]). Oxygen addition in a small amount improves the mechanical strength of titanium and helps titanium to prevent from cold creep at low temperatures. Amann et al. [33] showed that oxygen increase up to 2.4 pct at. leads to a significant increase in yield strength from 400 up to 1200 MPa at the expense of the ductility. Increase in yield strength was also found here but to a lesser extent, both for specimens with different ORL fractions (Fig. 11) or ultrathin specimens extracted within the ORL (Fig. 12). Further investigations to better characterize the evolution of the elementary deformation mechanisms with the oxygen addition would be of interest. It is worth mentioning that the pre-oxidation treatment leads to microstructural evolutions for the CP-Ti material. Indeed, a large grain growth was observed at 700 °C in the exposure time window. Grain growth evolution from 5 µm to 100 µm shows a greater spatial evolution than the ORL extension. Despite interest of oxygen mobility with grain boundary migration, this aspect is not covered in the present investigation and will be part of a further study. As far as mechanical aspects are concerned, this change in microstructure is generally accompanied by a decrease of the mechanical strength of the material due to the so-called Hall-Petch effect [51, 52]. This change in yield strength was observed for aged 400 µm-thick specimens, as shown in Figs. 4(b) and 5. In addition to this, size effects can be encountered for thin specimens, when the number of grains in the specimen thickness is below a critical value. This transition is called the polycrystalline to multicrystalline transition and has been reported in the literature to occur for specimens having 3 to 32 grains comprised in the specimen thickness, depending on the tested material [40, 41, 59–66]. For specimens with few grains in the specimen thickness, a softening behavior is observed in the literature due to plastic accommodation at the free surface. This effect was also found in the present investigation for aged CP-Ti specimens (Figs. 4(b) and 5). Both the Hall-Petch and polycrystalline-to-multicrystalline transition effects should lead to a decrease in apparent mechanical strength. However, pre-oxidation of the specimens leads to an opposite effect on the mechanical strength (Fig. 11). Therefore, the expected softening effect can diminish the effective strengthening of the oxygen ingress within Ti. In addition, pre-oxidation does not lead to a free surface condition due to the presence of the oxide. However, oxide cracking and spallation rapidly occur for CP-Ti materials under tensile loading. In addition, the effect of grain orientation on mechanical properties cannot be ignored for specimens having only few grains in the specimen thickness. In the present investigation, long and large specimens were used compared to the literature on such thin specimens, making the present specimen preparation much more complicated. At least 25 grains were found in the specimen width and at least 50 grains along the specimen gage. The gage area thus allows sampling several grain orientations and neighboring. In addition, all the specimens were tested in a similar way, with the same specimen orientation perpendicular to the loading direction. The texture was not significantly modified due to long exposure time at high temperature despite grain growth. Therefore, tensile tests were found particularly reproducible, even for thin specimens, and were used to testify the robustness of the present methodology.

For the Ti6242s alloy, the grain microstructure was sufficiently large enough to consider specimen thickness effects due to the presence of the free surface for the thinnest specimens in the as-received state (Fig. 6). Cave et al. [41] studied this transition behavior at high temperature on the Ti6242s alloy with a Widmanstätten microstructure
and found that the change in mechanical behavior was a function of the colony size and not the grain size for such
complex microstructures. As far as pre-oxidation is concerned, oxygen ingress only evidenced a drop in tensile
strength due to oxygen embrittlement even for ORL fractions as low as 10 pct (Fig. 11). This mechanical response is
very different from the one of the CP-Ti material. Early ORL cracking has been reported for structural titanium alloys
in many studies [6, 8, 12, 29, 30, 34]. The decrease in apparent tensile strength was not necessarily reported due to the
conventional size of the tensile specimens, but the oxygen embrittlement generally leads to a significant decrease in
ductility, as observed in the present study. Metallurgical evolution of the Ti6242s material could occur in the volume
due to long-term exposure at 700 °C but was not investigated in the present study (α2 phase precipitation for instance).
Using ultrathin specimen with ORL fractions greater than those of the literature aims to document both the loss in
ductility but also the drop in tensile strength. This drop in tensile strength for pre-oxidized Ti6242s alloy is consistent
with studies performed on the Ti64 alloy [8, 12, 29]. Ti6242s alloy exposed at high temperature did not show change
in microstructure even for long durations using EBSD characterization. Therefore, change in microstructure is not
responsible for the decrease in apparent tensile strength of the material with ORL fraction increase. This drop in
tensile strength was found to linearly evolve with the ORL fraction (Fig. 11). This mechanical evolution is consistent
with recent observations made for the Ti64 alloy [18, 34]. In other words, Ti64 and Ti6242s similarly behave to the
oxygen ingress in terms of mechanical integrity.

As aforementioned, the CP-Ti material and the structural Ti alloys have a different tensile resistance to the oxygen
ingress. This difference can arise from the alloying capacity of the CP-Ti material compared to structural titanium
alloys having several elements in substitution or insertion. Aluminum was found, in a given content range, to inhibit
oxygen embrittlement [37]. However, oxygen ingress due to high temperature oxidation leads to a much greater
oxygen content within the ORL compared to Chong et al. study [37]. Therefore, substitution or insertion capability
might not be the main reason for such difference in effective tensile strength evolution between the CP-Ti material and
Ti alloys. As far as the tensile behavior of the bulk material is concerned, the CP-Ti material exhibits a much greater
work-hardening and ductility capability compared to either Ti6242s or Ti64 alloys. Therefore, notch sensitivity after
ORL cracking might be different for these materials. For the CP-Ti material, early ORL cracking can continue to
develop based on the work-hardening capacity of the material. For the Ti6242s material, early ORL cracking at the
high end of the elastic domain leads to the reduction of the effective load-bearing section of the specimen and stress
concentration due to occurrence of large and sharp transverse cracks within the ORL. The higher notch sensitivity of
the Ti6242s makes it more prone to early failure due to stress concentration at the crack tip and thus to a decrease in
apparent tensile strength.

**5. Conclusions**

The tensile behavior of non-oxidized and pre-oxidized CP-Ti and Ti6242s specimens was investigated using very
thin specimens. This latter technical capability aimed to test the specimens with high fraction of oxygen-rich layer
(ORL) and thus to study the evolution of the tensile strength as a function of the oxygen uptakes. The ORL fraction
domain investigated in the present study ranged from 0 to 100 pct. Despite a slightly higher/nearly similar diffusion
coefficient of O within Ti for the CP-Ti material compared to the Ti6242s alloy, the CP-Ti material was found less
affected in terms of oxygen embrittlement than the Ti6242s alloy for a given duration. CP-Ti showed a significant
increase of the yield strength (up to 63 pct increase) with the increase in ORL fraction up to a given enrichment (80 pct
of ORL in the section), leading to a decrease of the tensile strength due to oxygen embrittlement counterbalancing the
benefit of the oxygen strengthening effect. For the Ti6242s alloy, a decrease in mechanical strength is observed for
all the pre-oxidized specimens, regardless of the pre-oxidation duration and specimen thickness. Ultrathin specimens
were extracted then tested at different locations of the ORL. For oxygen uptake lower than 1 pct at., the ultrathin
specimens exhibit an increase in tensile strength without significant strain-to-failure reduction for both the CP-Ti
material and the Ti6242s alloy. However, greater oxygen contents lead to a brittle behavior of the specimen for both
the CP-Ti material and the Ti6242s alloy. Despite a decrease in ductility for very rich regions of the ORL in oxygen,
the tensile strength increases with the oxygen uptake for the CP-Ti material.

6. License

This research was funded, in whole or in part, by The European Research Council, project HT-S₄DefOx - Grant number 948007. A [CC-BY public copyright license] has been applied by the authors to the present document and will be applied to all subsequent versions up to the Author Accepted Manuscript arising from this submission, in accordance with the grant's open access conditions.

7. Acknowledgements

The authors are grateful to AIRBUS Operation for financial support and materials. This work was supported by
the European Research Council [project HT-S₄DefOx - Grant number 948007]. The authors particularly acknowledge
Sophie GOUY from the Raimond Castaing Microanalysis Centre (UAR 3623) for performing EPMA analyses and
Arnaud PROIETTI for EBSD characterizations. The authors thank J.-C. Passieux (ICA) for new developments in
https://github.com/jcpassieux/pyxel.

8. Data Availability

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms
part of an ongoing study.

9. Declaration of competing interests

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.
## 10. Author contributions

QS: Data curation, Investigation, Methodology, Visualization, Writing - original draft

JG: Data curation, Formal analysis, Writing - review & editing

DM: Conceptualization, Methodology, Supervision, Validation, Writing - review & editing

BD: Conceptualization, Funding acquisition, Supervision, Writing - review & editing

DT: Funding acquisition, Investigation, Methodology, Project administration, Resources, Supervision, Visualization, Writing - original draft

## References

[1] J. L. Murray, H. A. Wriedt, The O-Ti (Oxygen-Titanium) system, Journal of Phase Equilibria 8 (1987) 148-165. doi:10.1007/BF02873201.

[2] P. Kofstad, P. B. Anderson, O. J. Krudtaa, Oxidation of titanium in the temperature range 800-1200°C, Journal of The Less-Common Metals 3 (1961) 89-97. doi:10.1016/0022-5088(61)90001-7.

[3] C. Dupressoire, A. Rouaix-Vande Put, P. Emile, C. Archambeau-Mirguet, R. Peraldi, D. Monceau, Effect of nitrogen on the kinetics of oxide scale growth and of oxygen dissolution in the Ti6242S titanium-based alloy, Oxidation of Metals 87 (2017) 343-353. doi:10.1007/s11085-017-9729-1.

[4] N. Vaché, D. Monceau, Oxygen Diffusion Modeling in Titanium Alloys: New Elements on the Analysis of Microhardness Profiles, Oxidation of Metals 93 (2020) 215-227. doi:10.1007/s11085-020-09956-9.

[5] W. L. Finlay, J. A. Snyder, Effects of three interstitial solutes (nitrogen, oxygen, and carbon) on the mechanical properties of high-purity, alpha titanium, JOM 2 (1950) 277-286. doi:10.1007/BF03399001.

[6] C. E. Shamblen, T. Redden, Air contamination and embrittlement of titanium alloys, in: The Science, Technology and Application of Titanium, London (UK), 1970, pp. 199-208. doi:10.1016/B978-0-08-006564-9.50027-0.

[7] H. Conrad, Effect of interstitial solutes on the strength and ductility of titanium, Progress in Materials Science 26 (1981) 123-403. doi:https://doi.org/10.1016/0079-6425(81)90001-3.

[8] L. Bendersky, A. Rosen, The effect of exposure on the mechanical properties of the Ti-6Al-4v alloy, Engineering Fracture Mechanics 20 (1984) 303-311. doi:10.1016/0013-7944(84)90135-8.

[9] R. Shenoy, J. Unnam, R. Clark, Oxidation and embrittlement of Ti-6Al-2Sn-4Zr-2Mo alloy, Oxidation of Metals 26 (1986) 105-124. doi:10.1007/BF00664276.

[10] J. Oh, B. Lee, S.-W. Cho, S.-W. Lee, G.-S. Choi, J.-W. Lim, Oxygen effects on the mechanical properties and lattice strain of Ti and Ti-6Al-4V, Metals and Materials International 17 (2011) 733-736. doi:10.1007/s12540-011-1006-2.

[11] F. B. Vicente, D. R. N. Correa, T. A. G. Donato, V. E. Arana-Chavez, M. A. R. Buzalaf, C. R. Grandini, The influence of small quantities of oxygen in the structure, microstructure, hardness, elasticity modulus and cytocompatibility of Ti-Zr alloys for dental applications, Materials 7 (2014) 542-553. doi:10.3390/ma7010542.

[12] H. Hornberger, C. Randow, C. Fleck, Fatigue and surface structure of titanium after oxygen diffusion hardening, Materials Science and Engineering A 630 (2015) 51-57. doi:10.1016/j.msea.2015.02.006.

[13] A. A. Hidalgo, T. Ebel, W. Limberg, F. Pyczak, Influence of oxygen on the fatigue behaviour of Ti-6Al-7Nb alloy, Key Engineering Materials 704 (2016) 44-52. doi:10.4028/www.scientific.net/KEM.704.44.

[14] D. Satko, J. Shaffer, J. Tiley, S. Semiatin, A. Pilchak, S. Kalidindi, Y. Kosaka, M. Glavicic, A. Salem, Effect of microstructure on oxygen rich layer evolution and its impact on fatigue life during high-temperature application of α/β titanium, Acta Materialia 107 (2016) 377-389. doi:10.1016/j.actamat.2016.01.058.
[15] C. M. Magazzeni, H. M. Gardner, I. Howe, P. Gopon, J. C. Waite, D. Rugg, D. E. Armstrong, A. J. Wilkinson, Nanoindentation in multi-modal map combinations: a correlative approach to local mechanical property assessment, Journal of Materials Research 36 (2021) 2235-2250. doi:10.1557/s43578-020-00035-y.

[16] Z. Cai, T. Xiang, W. Bao, J. Chen, T. Gao, G. Xie, Enhancing strength and ductility of pure titanium by interstitial oxygen atoms, Materials Science, SSRN Electronic Journal 14 (2022).

[17] D. Texier, Q. Sirvin, A. Dziri, H. Proudhon, V. Yastrebov, M. Legros, Effect of oxygen content on the sub-grain nanoindentation response in titanium affected by high temperature oxidation, in: Titanium 2023, Edimbourg, 2023.

[18] D. Texier, A. Palchoudhary, J. Genée, Q. Sirvin, Y. Zhang, G. Kermouche, D. Monceau, D. Poquillon, E. Andrieu, Effect of oxygen dissolution on the mechanical behavior of thin ti-6al-4v specimens oxidized at high temperature: Experimental and modeling approach, Corrosion Science 235 (2024) 112177. doi:https://doi.org/10.1016/j.corsci.2024.112177.

[19] Y. Lee, G. Welsch, Young's modulus and damping of Ti6Al4V alloy as a function of heat treatment and oxygen concentration, Materials Science and Engineering A 128 (1990) 77-89. doi:10.1016/0921-5093(90)90097-M.

[20] M. C. Brandes, M. Baughman, M. J. Mills, J. C. Williams, The effect of oxygen and stress state on the yield behavior of commercially pure titanium, Materials Science and Engineering A 551 (2012) 13-18. doi:10.1016/j.msea.2012.04.058.

[21] P. Kwasniak, M. Muzyk, H. Garbacz, K. Kurzydlowski, Influence of oxygen content on the mechanical properties of hexagonal Ti-First principles calculations, Materials Science and Engineering A 590 (2014) 74-79. doi:10.1016/j.msea.2013.10.004.

[22] J. Baillieux, C. Archambeau, P. Emile, D. Poquillon, Effet de la diffusion de l'oxygène sur le comportement mécanique du Ti-6Al-2Sn-4Zr-2Mo-0.1Si, in: CFM2015, Lyon (France), 2015, pp. 22-25.

[23] C. Beevers, J. Robinson, Some observations on the influence of oxygen content on the fatigue behaviour of α-titanium, Journal of The Less-Common Metals 17 (1969) 345-352. doi:10.1016/0022-5088(69)90061-7.

[24] G. Baur, P. Lehr, Influence spécifique de l'oxygène sur quelques propriétés physiques, les caractéristiques mécaniques et la dynamique de déformation plastique du titane de haute pureté, Journal of The Less-Common Metals 69 (1980) 203-218. doi:10.1016/0022-5088(80)90055-7.

[25] M. Wasz, F. Brotzen, R. McLellan, A. Griffin, Effect of oxygen and hydrogen on mechanical properties of commercial purity titanium, International Materials Reviews 41 (1996) 1-12. doi:10.1179/imr.1996.41.1.1.

[26] T. A. Parthasarathy, W. J. Porter, S. Boone, R. John, P. Martin, Life prediction under tension of titanium alloys that develop an oxygenated brittle case during use, Scripta Materialia 65 (2011) 420-423. doi:https://doi.org/10.1016/j.scriptамат.2011.05.025.

[27] M. Yan, W. Xu, M. S. Dargusch, H. P. Tang, M. Brandt, M. Qian, Review of effect of oxygen on room temperature ductility of titanium and titanium alloys, Powder Metallurgy 57 (2014) 251-257. URL: https://doi.org/10.1179/1743290114Y.0000000108. doi:10.1179/1743290114Y.0000000108.

[28] J. Q. Ren, Q. Wang, X. F. Lu, W. F. Liu, P. L. Zhang, X. B. Zhang, Effect of oxygen content on active deformation systems in pure titanium polycrystals, Materials Science and Engineering A 731 (2018) 530-538. doi:10.1016/j.msea.2018.06.083.

[29] D. Texier, Q. Sirvin, V. Velay, M. Salem, D. Monceau, B. Mazères, E. Andrieu, R. Roumiguier, B. Dod, Oxygen/nitrogen-assisted embrittlement of titanium alloys exposed at elevated temperature, MATEC Web Conf. 321 (2020) 06004. doi:10.1051/matecconf/202032106004.

[30] A. Casadebaigt, J. Hugues, D. Monceau, High temperature oxidation and embrittlement at 500-600 °C of Ti-6Al-4V alloy fabricated by Laser and Electron Beam Melting, Corrosion Science 175 (2020) 108875. doi:10.1016/j.corsci.2020.108875.

[31] Q. Sirvin, D. Texier, J. Genée, V. Velay, D. Monceau, B. Dod, Oxygen content effect on mechanical behavior of CP titanium exposed at elevated temperature, in: Titanium 2023, Edimbourg, 2023.

[32] A. Amherd Hidalgo, T. Ebel, R. Frykholm, E. Carreño-Morelli, F. Pyczak, High-oxygen MIM Ti-6Al-7Nb: Microstructure, tensile and fatigue properties, Materials Today Communications 34 (2023). doi:10.1016/j.mtcomm.2022.104982.

[33] F. Amann, R. Poulain, S. Delannoy, J. Couzinié, E. Clouet, I. Guillot, F. Prima, An improved combination of tensile strength and ductility in titanium alloys via oxygen ordering, Materials Science and Engineering: A 867 (2023) 144720. doi:https://doi.org/10.1016/j.msea.2023.144720.
[34] A. Casadebaigt, D. Monceau, J. Hugues, Oxygen embrittlement kinetics at 500 – 600 °C of the Ti-6Al-4V alloy fabricated by Laser and Electron Powder Bed Fusion, high temperature corrosion and protection of materials (????) 107–122. doi: underreview.

[35] B. Barkia, V. Doquet, J. P. Couzinié, I. Guillot, E. Héripré, In situ monitoring of the deformation mechanisms in titanium with different oxygen contents, Materials Science and Engineering A 636 (2015) 91–102. doi:10.1016/j.msea.2015.03.044.

[36] R. Poulain, S. Delannoy, J.-P. Couzinié, I. Guillot, E. Clouet, F. Prima, Extra-high oxygen addition as a new strengthening strategy to overcome strength/ductility trade-off in fully biocompatible hexagonal titanium alloys, MATEC Web of Conferences 321 (2020) 05012. doi:10.1051/matecconf/202032105012.

[37] Y. Chong, R. Zhang, M. S. Hooshmand, S. Zhao, D. C. Chrzan, M. Asta, J. W. Morris, A. M. Minor, Elimination of oxygen sensitivity in α-titanium by substitutional alloying with Al, Nature Communications 12 (2021) 1–9. doi:10.1038/s41467-021-26374-w.

[38] R. Poulain, F. Amann, J. Deya, J. Bourgon, S. Delannoy, F. Prima, Short-time heat treatments in oxygen strengthened ti-zr α titanium alloys: A simple approach for optimized strength-ductility trade-off, Materials Letters 317 (2022) 132114. doi: https://doi.org/10.1016/j.matlet.2022.132114.

[39] Y. Fu, W. Xiao, S. Zhao, L. Ren, J. Wang, J. Rong, J. Li, X. Zhao, C. Ma, Substantially strengthening a dual-phase titanium alloy by moderate oxygen doping, Scripta Materialia 226 (2023) 115236. doi: https://doi.org/10.1016/j.scriptamat.2022.115236.

[40] C. Romain, D. Texier, C. Desgranges, J. Cormier, S. Knittel, D. Monceau, D. Delagnes, Oxidation of thin nickel-based superalloy specimens: Kinetics study and mechanical integrity, Oxidation of Metals 96 (2021) 169–182. doi:10.1007/s11085-021-10075-2.

[41] K. Cave, D. Texier, E. Fessler, D. Monceau, D. Poquillon, Size effect on the tensile mechanical behavior of thin Ti6242S specimens at 723 K and 823 K, Metallurgical and Materials Transactions A: Physical Metallurgy and Materials Science 54 (2023) 549–561. doi:10.1007/s11661-022-06898-6.

[42] D. Texier, D. Monceau, J.-C. Salabura, R. Mainguy, E. Andrieu, Micromechanical testing of ultrathin layered material specimens at elevated temperature, Materials at high temperatures 33 (2016) 325–337. doi:10.1080/09603409.2016.1182250.

[43] J.-c. Passieux, R. Bouclier, J.-n. Périé, A space-time PGD-DIC algorithm, Experimental Mechanics 58 (2018) 1195–1206. doi:10.1007/s11340-018-0387-2.

[44] K. Cavé, D. Texier, S. Vallot, N. Chanfreau, E. Fessler, M. Dehmas, D. Monceau, D. Poquillon, Effect of creep loading on the oxygen diffusion of Ti6242S at 650 °C, Scripta Materialia (2023) submitted.

[45] Relationship between the volume of the unit cell of hexagonal-close-packed Ti, hardness and oxygen content after α-case formation in Ti-6Al-2Sn-4Zr-2Mo-0.1Si alloy, Journal of Applied Crystallography 49 (2016) 175–181. doi:10.1107/S1600576715022906.

[46] K. Gautier, E. Epifano, T. Gheno, D. Connétable, D. Monceau, Effects of al and refractory alloying elements (w, ta and hf) on oxidation kinetics, oxygen dissolution and diffusion in titanium alloys, Corrosion Science 237 (2024) 112330. doi: https://doi.org/10.1016/j.corsci.2024.112330.

[47] R. Hielscher, H. Schaeben, A novel pole figure inversion method: Specification of the MTEX algorithm, Journal of Applied Crystallography 41 (2008) 1024–1037. doi:10.1107/S0021889808030112.

[48] J. Schindelin, I. Arganda-Carreras, E. Frise, V. Kaynig, M. Longair, T. Pietzsch, S. Preibisch, C. Rueden, S. Saalfeld, B. Schmid, J.-Y. Tinevez, D. James White, V. Hartenstein, K. Eliceiri, P. Tomancak, A. Cardona, Fiji: an open-source platform for biological-image analysis, Nature Methods 9 (2012) 676–682. doi:10.1016/B978-0-12-415853-5.00046-7.

[49] A. Fick, Ueber Diffusion, Annalen der Physik 170 (1855) 59–86. doi: https://doi.org/10.1002/andp.18551700105.

[50] N. Vaché, Y. Cadoret, B. Dod, D. Monceau, Modeling the oxidation kinetics of titanium alloys: Review, method and application to Ti-64 and Ti-6242s alloys, Corrosion Science 178 (2021) 109041. doi:10.1016/j.corsci.2020.109041.

[51] E. Hall, The deformation and ageing of mild steel: III Discussion of results, Proceedings of the Physical Society 64 (1951) 747–753. doi:10.1088/0370-1301/64/9/303.

[52] N. Petch, The cleavage strength of polycrystals, Journal of the Iron and Steel Institute 174 (1953) 25–28. doi:10.1007/BF01972547.

[53] M. Hooshmand, C. Niu, D. Trinkle, M. Ghazisaeidi, First-principles prediction of oxygen diffusivity near the (10-12) twin boundary in titanium, Acta Materialia 156 (2018) 11–19. doi: https://doi.org/10.1016/j.actamat.2018.05.076.
[54] X. Chen, L. Huang, Y. Jiao, S. Wang, Q. An, Y. Bao, L. Geng, Mechanisms of oxidation anisotropy between -ti (0001) and (01-10) crystallographic planes in titanium matrix composites, Materials Letters 286 (2021) 129230. doi: https://doi.org/10.1016/j.matlet. 2020.129230.

[55] H. H. Wu, D. R. Trinkle, Direct diffusion through interpenetrating networks: oxygen in titanium, Phys. Rev. Lett. 107 (2011) 045504. doi:10.1103/PhysRevLett.107.045504.

[56] H. H. Wu, D. R. Trinkle, Solute effect on oxygen diffusion in -titanium, Journal of Applied Physics 113 (2013) 223504. doi:10.1063/1. 4808283.

[57] M. Song, S. Han, D. Min, G. Choi, J. Park, Diffusion of oxygen in β-titanium, Scripta Materialia 59 (2008) 623–626. doi: https://doi. org/10.1016/j.scriptamat.2008.05.037.

[58] Z. Liu, G. Welsch, Literature survey on diffusivities of oxygen, aluminum, and vanadium in alpha titanium, beta titanium, and in rutile, Metallurgical Transactions A 19 (1988) 1121–1125. doi: https://doi.org/10.1007/BF02628396.

[59] R. Armstrong, On size effects in polycrystal plasticity, Journal of the Mechanics and Physics of Solids 9 (1961) 196–199. doi:10.1016/ 0022-5096(61)90018-7.

[60] A. Thompson, Use of non-polycrystal specimens in mechanical behavior tests, Scripta metallurgica 8 (1974) 145–148. doi:10.1017/ CB09781107415324.004.

[61] N. Hansen, The effect of grain size and strain on the tensile flow stress of aluminium at room temperature, Acta Metallurgica 25 (1977) 863–869. doi:10.1016/0001-6160(77)90171-7.

[62] S. Miyazaki, K. Shibata, H. Fujita, Effect of specimen thickness on mechanical properties of polycrystalline aggregates with various grain sizes, Acta Metallurgica 27 (1979) 855–862. doi:10.1016/0001-6160(79)90120-2.

[63] C. Thompson, The yield stress of polycrystalline thin films, Journal of Materials Research 8 (1993) 237–238.

[64] C. Keller, E. Hug, X. Feaugas, Microstructural size effects on mechanical properties of high purity nickel, International Journal of Plasticity 27 (2011) 635–654. doi:10.1016/j.ijplas.2010.08.002.

[65] E. Hug, P. Dubos, C. Keller, L. Duchêne, A. Habraken, Size effects and temperature dependence on strain-hardening mechanisms in some face centered cubic materials, Mechanics of Materials 91 (2015) 136–151. doi: https://doi.org/10.1016/j.mechmat.2015.07.001.

[66] D. Texier, J. Genée, V. Velay, A. C. Moreno, D. Monceau, E. Andrieu, Size effects on the plastic behavior of polycrystalline materials: Grain size, precipitation state and free-surface effects, International Journal of Plasticity 188 (2025) 104284. doi: https://doi.org/10.1016/ j.ijplas.2025.104284.