Hydrogen Impact on the Mechanical Properties of Three Titanium Alloys Q

Hydrogen impact on the mechanical properties of three titanium alloys Q. Bignon, F. Martin, Q. Auzoux, Y. Wouters

To cite this version:

Q. Bignon, F. Martin, Q. Auzoux, Y. Wouters. Hydrogen impact on the mechanical properties of three titanium alloys. 3rd International conference on metals and hydrogen, May 2018, Gand, Belgium. hal-02416233

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HAL is a multi-disciplinary open access L’archive ouverte pluridisciplinaire HAL, est archive for the deposit and dissemination of sci- destinée au dépôt et à la diffusion de documents entific research documents, whether they are pub- scientifiques de niveau recherche, publiés ou non, lished or not. The documents may come from émanant des établissements d’enseignement et de teaching and research institutions in France or recherche français ou étrangers, des laboratoires abroad, or from public or private research centers. publics ou privés. HYDROGEN IMPACT ON THE MECHANICAL PROPERTIES OF THREE TITANIUM ALLOYS

Quentin Bignon1,2, Frantz Martin1, Quentin Auzoux1, Yves Wouters2

1Den-Service de la Corrosion et du Comportement des Matériaux dans leur Environnement (SCCME), CEA, Université Paris-Saclay, F-91191, Gif-sur-Yvette, France

2Université Grenoble Alpes, SIMaP, F-38402 Saint Martin d’Hères, France

ABSTRACT

During long-term exposure to pressurised hot water (350 °C and 15 MPa), titanium alloys may absorb hydrogen. This study focuses on the determination of the impact of hydrogen on their mechanical properties. Specimens of three different titanium alloys – T40 (a phase), TA6V (4 % of b phase) and Ti10-2-3 (38 % of b phase) – were hydrogen-charged up to 1240, 1630 and 1570 wt. ppm, respectively. Tensile tests at 20 °C and 300 °C under air environment at 10-4 s-1 were then conducted on both uncharged and charged specimens. Fracture surfaces and longitudinal cross-sections of the specimens were eventually observed by Scanning Electron Microscopy (SEM). Results obtained at 20 °C show that concentration of hydrogen in the range of 1200 - 1600 wt. ppm increases the ultimate tensile strength of the three materials and decreases their ductility. These effects are attributed to brittle hydrides owing to the low solubility of hydrogen in a phase at low temperature. Hydrogen also slightly decreased the yield stress of TA6V and Ti10-2-3. The same hydrogen concentration range did not reduce ductility of three materials at 300 °C. It could be due to the increase of hydrogen solubility with temperature or both increase of the hydrides plasticity and decrease of yield stress of the matrix with temperature.

INTRODUCTION

Titanium alloys could be good candidates for nuclear Pressurised Water Reactor (PWR) primary circuit structure components because of their low neutron activation, their good corrosion behavior and their good mechanical properties. However, titanium alloys tend to absorb hydrogen, which may affect their mechanical properties in service.

Hydrogen impact on the mechanical properties of titanium alloys at room temperature have been extensively studied for both a [1-10] and a/b alloys [11-18]. In a titanium alloys, it is well established that hydrogen concentration above the solubility limit (around 20 wt. ppm at ambient temperature) leads to the precipitation of brittle titanium hydride phases [2-3]. In a/b titanium alloys, the role of hydrogen depends on the phases ratio and the microstructure. The b phase increases the solubility limit of hydrogen in the alloy. In the meanwhile, a continuous b phase enhances the diffusion of hydrogen to the a/b interface where hydride precipitation may occur as soon as the solubility limit of the a phase is exceeded [14, 15, 19]. Such a phenomenon may promote the embrittlement of the alloys owing to a weakening point at the a/b interface.

This study aims at measuring the impact of hydrogen on the mechanical properties of three different titanium alloys that were homogeneously hydrogen-charged in the 1200 – 1600 wt. ppm range. Tensile tests at 20 °C and 300 °C under air environment at 10-4 s-1 were conducted on both uncharged and charged specimens. The hydrides distribution according to the temperature and the microstructure as well as the impact of hydrides on the fracture mode of the studied titanium alloys were then investigated by Scanning Electron Microscopy (SEM).

MATERIALS AND PROCEDURES

Materials

Three titanium alloys representative of different metallurgical classes have been studied: T40, TA6V and Ti10-2-3. T40 is a single phase a alloy also called commercially pure titanium (ASTM grade 2); TA6V or Ti-6Al-4V (ASTM grade 5) is a two-phase a/b titanium alloy; and Ti10-2-3 or Ti-10V-2Al-3V is a two-phase b metastable alloy. EBSD analysis were carried out to determine the volumetric phase ratios. TA6V and Ti10-2-3 have respectively 4 % and 38 % of b phase. The three alloys were supplied by TIMET as billets. Their chemical compositions are given in the Table 1.

Billets were hot rolled during two hours at 675 °C, 730 °C and 760 °C for T40, TA6V and Ti10-2-3 respectively. The microstructures obtained are given in Figure 1. T40 has an equiaxed grain microstructure. Average grain size is 60 µm. TA6V and Ti10-2-3 have a finer bimodal microstructure.

Length and thickness of tensile specimens were along radial and longitudinal direction of billets, respectively. The gauge section length (L), width (W) and thickness (T) were 8 mm, 2 mm and 1 mm, respectively. EBSD analysis have shown that c axis of the a hexagonal unit cell are mainly along the length and the width of T40 specimens and mainly along the length and the thickness of TA6V specimens. Ti10-2-3 crystallographic texture has not been characterized yet. Specimens were polished successively with SiC paper, diamond paste and colloidal silica suspensions. The head of specimens were then analysed by total melting technique with thermal conductivity measurement (EMGA-821 HORIBA) in order to determine the initial hydrogen concentration before any hydrogen charging. The initial hydrogen concentrations of specimens dedicated to the tensile tests at 20 °C were 20, 80 and 260 wt. ppm for T40, TA6V and Ti10-2-3 respectively. The initial hydrogen concentrations of specimens dedicated to the tensile tests at 300 °C have not been measured yet.

Table 1: Billets chemical composition (Ti = bal.)

wt. % Al V Fe C O T40 0.01 0.01 0.03 0.003 0.16 TA6V 6.38 4.12 0.16 0.008 0.17 Ti10-2-3 2.96 9.73 2.15 0.011 0.09

Figure 1: Microstructure of the studied T40 observed by optical microscopy with polarised light and microstructures of TA6V and Ti10-2-3 observed by SEM in backscattered electrons mode.

Hydrogen charging procedure

Preliminary tests on 1 mm thick specimens have shown that hydrogen uptake after exposure to simulated PWR primary water during 1000 h were at the most 15 wt. ppm. For such thin parts, a linear extrapolation suggests that the hydrogen uptake may be in the range of 1000 – 1500 wt. ppm after 10 years of exposure to primary water. Tensile specimens were then cathodically charged in order to reach -1 such a hydrogen concentration. Charging were done at 80 °C in H2SO4 at 0.05 mol.L at a current density of - 5 mA.cm-² during 10 h 30 min, 3 h 50 min and 1 h 30 min for T40, TA6V and Ti10-2-3 respectively.

Tensile specimens were then heat-treated in vacuum (pressure below 10-6 mbar) at 300 °C during 5 h and cooled down at approximately 2 °C.min-1 in order to obtain a distribution of hydrogen or hydride as close as possible as the one expected after exposure in nominal conditions (i.e. PWR primary water). This treatment was done under high vacuum in a Thermal Desorption Spectroscopy (TDS) setup. No hydrogen desorption was detected by the mass spectrometer during the heat-treatment. A homogeneous distribution of hydrides was obtained in every specimen. Figure 2 shows the microstructure of the three materials after treatment where hydrides appears in black. In T40, both intragranular and intergranular hydrides are observed. They have a platelet shape and an average diameter close to the grain size (i.e. 60 µm). In TA6V and Ti10-2-3, they are located at the a/b interface. After the treatment, the hydrogen concentrations of specimens dedicated to the tensile tests at 20 °C were 1240, 1630 and 1570 wt. ppm for T40, TA6V and Ti10-2-3 respectively. The hydrogen concentrations of hydrogen charged specimens dedicated to the tensile tests at 300 °C have not been measured yet. Nevertheless, their hydrogen concentration should be very close to the other specimens given that the charging procedure conditions were exactly the same for each alloy.

Figure 2: SEM observation (Backscattered electron) of hydride (in black) distribution in (L,T) plane of T40, TA6V and Ti10-2-3 tensile specimens center after cathodic charging and heat treatment at 300 °C.

Tensile tests

Tensile tests were done under air environment at 20 °C and 300 °C on an certified INSTRON electromechanical tensile testing machine equipped with a 10 kN load cell and a ventilated heating chamber. The strain rate was 10-4 s-1. For tensile tests at 300 °C, temperature was measured by a thermocouple in contact with the gauge length of the specimen. Area reductions at failure were measured by optical binocular. Fracture surfaces and longitudinal cross sections were observed by SEM in secondary electrons (SE) and backscattered electrons (BSE) mode, respectively.

RESULTS

Figure 3 shows the tensile curves obtained for uncharged and charged tensile specimens at 20 °C and 300 °C.

TA6V and Ti10-2-3 yield stress at 0.2% (YS0.2) and ultimate tensile strength (UTS) are higher than T40 ones. Indeed, YS0.2 are equal to 360 MPa, 860 MPa and 866 MPa and UTS are equal to 454 MPa, 922 MPa and 946 MPa for T40, TA6V and Ti10-2-3 uncharged specimens respectively.

Marked impacts of hydrogen charging on the mechanical properties at 20 °C were found. T40 hydrogen charging caused the increase of YS0.2 and UTS and a drastic decrease of plastic elongation at fracture (EF). TA6V and Ti10-2-3 hydrogen charging also caused an increase of UTS and a decrease of EF, which was more important for TA6V than for Ti10-2-3. Contrary to T40, TA6V and Ti10-2-3 hydrogen charging induced a slight softening at the very beginning of the plasticity. Indeed, YS0.2 decreased from 860 MPa to 790 MPa for TA6V and from 880 MPa to 855 MPa for Ti10-2-3.

At 300 °C, the mechanical behavior of hydrogen charged and uncharged specimens were globally similar.

However, some differences are noticeable. First, YS0.2 of T40 slightly increased with hydrogen charging whereas UTS of T40 slightly decreased. Concerning TA6V, hydrogen charging induced an increase of EF and the occurrence of serrations with an amplitude of 5 MPa during plastic flow.

Tensile tests at 20 °C Tensile tests at 300 °C 600 200 500 Charged T40 Uncharged T40 400 150 Uncharged T40 Charged T40 300 100 (MPa) (MPa) 200 T40 50 100 Conventional Stress Conventional Stress 0 0 0,00 0,05 0,10 0,15 0,20 0,25 0 0,1 0,2 0,3 0,4 Conventional Strain (-) Conventional Strain (-) 1200 800 Charged 1000 Uncharged TA6V 800 Charged TA6V 600 670 600 TA6V 400 660 Uncharged 400 (MPa) 650 (MPa) TA6V TA6V 200 640 200 0,05 0,115 0,18

0 Conventional Stress

Conventional Stress 0 0,00 0,05 0,10 0,15 0,20 0,25 0 0,1 0,2 0,3 Conventional Strain (-) Conventional Strain (-) 1200 1000 Charged Ti10-2-3 1000 800 Charged Ti10-2-3

800 Uncharged Uncharged

3 600 - 600

2 Ti10-2-3 - 400 Ti10-2-3 (MPa) (MPa) 400

Ti10 200 200 Conventional Stress Conventional Stress 0 0 0 0,05 0,1 0,15 0,2 0 0,05 0,1 0,15 0,2 Conventional Strain (-) Conventional Strain (-) Figure 3: Conventional stress - strain curves of tensile tests obtained at 10-4 s-1 on both uncharged and charged specimens of the studied materials, at 20 °C and at 300 °C.

Optical binocular observations indicated a significant difference between the macroscopic fracture surfaces of the hydrogen charged and uncharged specimens tested at 20 °C. The uncharged specimens exhibited slanted fracture surfaces with normal direction oriented 45° from the loading direction in the LT plan for the TA6V and in the LW plan for the T40 and Ti10-2-3. Area reduction at failure of the uncharged specimens was high: 41 %, 43 % and 32 % for T40, TA6V and Ti10-2-3, respectively. On the contrary, all hydrogen charged specimens exhibited a flat fracture surface perpendicular to the tensile loading direction. Due to hydrogen charging, area reduction at failure decreased to 2 %, 3 % and 7 % for T40, TA6V and Ti10-2-3 respectively.

At 300 °C, all uncharged and charged specimens exhibited macroscopic ductile fracture. Area reductions at failure slightly decreased from 78 %, 63 % and 66 % to 55 %, 56 % and 47 % for T40, TA6V and Ti10-2- 3, respectively.

Figure 4, 5 and 6 present SEM observations of the fracture surfaces and longitudinal cross sections obtained after testing at 20°C of both uncharged and hydrogen charged specimens of T40, TA6V and Ti10-2-3, respectively. Figure 4 illustrates the ductile fracture of the uncharged T40 specimen and the brittle fracture of the charged T40 specimen. Figure 5 shows the ductile fracture of the uncharged TA6V specimen and the brittle fracture of the charged TA6V specimen (with both a/b interfacial fracture and a transgranular brittle fracture). Figure 6 show the ductile fracture of the uncharged Ti10-2-3 specimen and the mixed fracture of the charged Ti10-2-3 specimen (with both a/b interfacial fracture and b transgranular ductile fracture).

Uncharged T40 Charged T40

surfaces Fracture Fracture

Longitudinal cross section

Figure 4: SEM observations of fracture surfaces in SE mode (on top) and longitudinal cross sections in BSE (at the bottom) of T40 uncharged (on the left) and charged (on the right) specimens after test at 20 °C.

Uncharged TA6V Charged TA6V

surfaces Fracture

Longitudinal cross section

Figure 5: SEM observations of fracture surfaces in SE mode (on top) and longitudinal cross sections in BSE mode (at the bottom) of TA6V uncharged (on the left) and charged (on the right) specimens after test at 20 °C. Arrows show a/b interfacial fracture.

Uncharged Ti10-2-3 Charged Ti10-2-3

surfaces Fracture Fracture

Longitudinal cross section

Figure 6: SEM observations of fracture surfaces in SE mode (on top) and longitudinal cross sections in BSE mode (at the bottom) of Ti10-2-3 uncharged (on the left) and charged (on the right) specimens after test at 20 °C. Arrows show a/b interfacial fracture.

DISCUSSION

Hydride distribution according to temperature and microstructure

The presence of homogeneously distributed hydrides in T40 hydrogen charged sample (figure 2) is consistent with the values of the solubility and the diffusion coefficient of hydrogen in titanium a phase at 300 °C. Indeed, during the heat treatment conditions, most of the hydrogen atoms are in solid a solution due to the high value of hydrogen solubility (SH (300 °C) = 1100 wt. ppm [4]). Following a -7 references [20] and [21], the diffusion coefficient DH (300 °C) is around 10 cm²/s, which gives a theoretical diffusion length of around 1 mm during the heat treatment – i.e. equal to the thickness of tensile specimens. Hydrogen can therefore diffuse to the middle of T40 specimen during the heat treatment temperature plateau. Eventually, hydrogen precipitates into hydrides during cooling due to the decrease of hydrogen solubility with decreasing temperature.

The hydrogen/hydride distribution in two-phase a/b TA6V and Ti10-2-3 specimens is linked to the b phase ratio and continuity. Diffusion coefficient of hydrogen in b phase at 300 °C is two orders of b -5 magnitude higher than in a phase (DH (300 °C) = 3.10 cm²/s [22]). Thereby, a continuous b phase provides a quick diffusion pathway for hydrogen [13]. Nevertheless, hydride precipitation occurs at the a/b interface as soon as the solubility limit of hydrogen in the a phase is exceeded. Hydrogen concentrations of TA6V and Ti10-2-3 specimens (respectively 1630 and 1570 wt. ppm) are higher than hydrogen solubility in a phase at 300 °C. It means that hydride precipitation at a/b interface in the center of specimens should have begun during the heat treatment temperature plateau. During cooling, hydrides growth occurs due to the decrease of hydrogen solubility with decreasing temperature.

Hydrogen and hydride distribution inside specimens during the tensile tests at 300 °C is equivalent to the distribution at the end of the heat treatment temperature plateau. Thus, at 300 °C, T40 specimen should

7 contain 1100 wt. ppm of solid solution hydrogen and around 140 wt. ppm of hydrogen under the form of hydrides. Similarly, TA6V and Ti10-2-3 should both contain 1100 wt. ppm of solid solution hydrogen equally divided between a and b phase and around 530 and 470 wt. ppm of hydrogen under the form of hydrides, respectively, located at the interface a/b.

Effect of hydrides on mechanical properties

Brittle fracture of T40 hydrogen charged specimen tested at 20 °C is linked to the brittle hydride phase. Indeed, Figure 4 indicates that flat facets distinctive of brittle fracture (surrounded by small area of a ductile fracture) correspond to areas where hydride are located. Fracture seems to occur inside hydrides rather than at the hydride/matrix interfaces. This observation is consistent with conclusions of several studies on both hydrided titanium and zirconium alloys [23] [24] [25]. Increase of YS0.2 and UTS by hydrogen charging could be also attributed to the presence of hydrides. Indeed, Xu et al. have measured a titanium hydride yield stress of 895 MPa much higher than the initial annealed commercially pure titanium (∼ 450 MPa) [26].

Fracture at a/b interfaces in TA6V and Ti10-2-3 hydrogen charged specimens tested at 20 °C (figure 5 and figure 6) could also be attributed to hydride located at these interfaces (figure 2). It is consistent with conclusion of other studies concerning hydrided a-b titanium alloys [12] [13]. Nevertheless, SEM observations do not allow to identify the exact fracture location (inside the hydride, at the a/hydride interface or at b/hydride interface).

At 300 °C, high ductility of hydrogen charged specimens may result from three reasons. First, the low hydride concentration due to the increase of hydrogen solubility with temperature. Second, the decrease of yield stress of the matrix with temperature, which lowers the stress imposed to the hydrides and may prevent its fracture. Third, the temperature may also decrease the yield strength of hydrides which may flow with the matrix without cracking. Several studies show similar results concerning zirconium alloys. Indeed, Arsene et al. found that ductile-brittle transitions of recrystallized Zircaloy-2 occurred for hydrogen concentrations of 1000 and 2000 wt. ppm at the temperatures of 20 and 300 °C, respectively [24]. In the same way, Yangnik et al. noted that Zircaloy-4 containing 200 wt. ppm of hydrogen tested at 20 °C and 300 °C exhibits a loss of ductility compared to uncharged material only at 20 °C [27]. Finally, Choubey et al. have observed that number of crack initiation of a-b Zr-2.5Nb alloy containing 100 wt. ppm of hydrogen drop drastically starting from 150 °C [28].

Effect of hydrogen in solid solution on mechanical properties

Hydrogen-induced softening which occurs at the onset of plastic stage of TA6V and Ti10-2-3 specimens tested at 20 °C (figure 3) has already been observed on a-b titanium alloys in other studies [15] [16] [17]. This phenomenon may be attributed to a solid solution hydrogen effect in a phase. Indeed, it is well established than hydrogen enhances dislocations movement in a titanium [9]. Two different hypothesis are proposed in the literature to explain this effect. According to the first one, solid solution hydrogen may decrease dislocation activation volume and then increase their mobility [17]. According to the second one, the softening effect of hydrogen comes from the weakening of the interactions between solute element (O, C…) and dislocations when both temperature and strain rate conditions are suitable

8 for Dynamic Strain Ageing (DSA) to occur [16] [18]. More tests at different temperatures and strain rates are necessary to identify the cause of hydrogen-induced softening in the present conditions.

At 300 °C, serrations which appear during plastic flow of the TA6V hydrogen charged specimen are similar to Portevin-Le Chatelier effect characteristic of DSA phenomenon. Jousset found that DSA of Ti6242 attributed to oxygen occurred in temperature range of 200 °C – 400 °C for 10-4 s-1 strain rate [7]. Then, it could be suggested that serrations observed are due to oxygen DSA phenomenon. Nevertheless, it is not possible to explain why no serration has been observed on non-hydrogen charged TA6V specimen. More tensile tests at different temperatures, hydrogen concentrations and strain rates will be necessary to conclude on this point.

CONCLUSION

Hydrogen effects on the mechanical properties of three titanium alloys at 20 °C and 300 °C have been investigated in concentration range of 1200 – 1600 wt. ppm. The low solubility of hydrogen in a phase at room temperature is responsible for the hydride precipitation in both a phase and a/b titanium alloys. In T40, hydrides obtained in the present study have a platelet shape and are both trangranular and intergranular with an average diameter close to the grain size. In TA6V and Ti10-2-3, hydrides are located at the a/b interface. At 20 °C, hydrogen charged specimens exhibit hardening and embrittlement effects due to presence of hydrides. Brittle rupture was observed to occur at places where hydrides are localised. At 300 °C, no hydrogen embrittlement has been observed. It could be due to the low hydride concentration due to the increase of hydrogen solubility with temperature or both increase of the hydrides plasticity and decrease of yield stress of the matrix with temperature.

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