Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement

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Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement S.M. Teus, D.G. Savvakin, O.M. Ivasishin, V.G. Gavriljuk* G.V. Kurdyumov Institute for Metal Physics, Vernadsky Blvd. 36, 03680 Kiev, Ukraine

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abstract

Article history:

CrNi austenitic steels and titanium alloy Ti-10V-2Fe-3Al are studied aiming to clarify a

Received 23 July 2016

reason for difference between two classes of engineering materials in their sensitivity to

Received in revised form

hydrogen brittleness. Using ab initio calculations, it is found that hydrogen increases

29 September 2016

density of electron states at the Fermi level in both materials except for its decrease in the

Accepted 30 September 2016

titanium alloy at extremely high hydrogen contents. Migration of hydrogen atoms and

Available online xxx

their interaction with dislocations are studied using mechanical spectroscopy. The enthalpies of hydrogen atoms migration and their binding to dislocations, as well as tem-

Keywords:

perature for condensation of hydrogen clouds around dislocations, are shown to be

Hydrogen embrittlement

significantly larger in austenitic steels in comparison with the b titanium alloy. This is a

Austenitic steels

reason for lower temperature range of hydrogen embrittlement in the titanium alloys. The

Titanium alloy

different hydrogen effect in the studied materials and usage of hydrogen as temporary

Electron structure

alloying element increasing plasticity of titanium alloys in the course of their processing

Hydrogen migration

are interpreted within the frame of HELP theory.

Dislocations

© 2016 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.

Introduction Despite numerous studies, a physical mechanism for hydrogen brittleness of metals still remains to be debatable. Excluding hydride formation which can be applied only to hydride-forming metals like V, Zr, Nb, Ta and partially Ti, the available hypotheses can be listed in the chronological order as those concerned with hydrogen enhanced decohesion (HEDE), e.g. Refs. [1e4], adsorption-induced dislocation emission (AIDE), e.g. Refs. [5e9], hydrogen-enhanced localized plasticity (HELP), e.g. Refs. [10e15], and hydrogen-enhanced strain-induced vacancies (HESIV), e.g. Refs. [16e18].

The essence of HEDE hypothesis, as formulated by Oriani [3], is that the hydrogen-induced propagation of a crack occurs if the local tensile stress normal to the plane of this crack becomes equal to the cohesive force per unit area reduced by hydrogen. The increase in the hydrogen concentration at the crack tip is achieved due to the effect of elastic stresses on the chemical potential of hydrogen. A detailed scheme for HEDE proposed by Knott [4] includes the following steps: (i) the transfer of hydrogen atoms by moving dislocations into existing cracks with subsequent formation of gaseous hydrogen, (ii) a partial dissociation of hydrogen molecules at the crack tip followed by dissolution of hydrogen atoms in the crystal lattice and their migration into a hydrostatic region

* Corresponding author. E-mail address: [email protected] (V.G. Gavriljuk). http://dx.doi.org/10.1016/j.ijhydene.2016.09.212 0360-3199/© 2016 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article in press as: Teus SM, et al., Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement, International Journal of Hydrogen Energy (2016), http://dx.doi.org/10.1016/ j.ijhydene.2016.09.212

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ahead of the crack tip, (iii) the hydrogen-caused weakening of interatomic bonds in this region, opening of a microcrack and its joining to the growing macrocrack. Numerous fractographic studies of fracture surfaces in the hydrogen-charged metals are at variance with HEDE hypothesis. For example, a strongly localized dislocation slip was observed in austenitic steels AISI 310S and 304 as a prerequisite for the opening of microcracks [19,20]. Critical for this hypothesis is also the observation of hydrogen embrittlement in the amorphous materials where both plastic deformation and fracture proceed only by the shear and the role of normal stresses approaches to zero, see e.g. Ref. [21]. According to AIDE hypothesis [5e9], the dissolved hydrogen is not responsible for embrittlement. Instead, be adsorbed at the surface of the crack tips, the hydrogen atoms facilitate the nucleation of dislocations due to weakening the interatomic bonds over several atomic distances. Once nucleated, dislocations can readily move away from the crack tip under the applied stress. Furthermore, the crack growth under sustained or monotonically increasing stresses occurs not only by dislocation emission from the crack tip but also involves the nucleation and growth of microvoids (or nanovoids) ahead of it. In both cases, HEDE and AIDE, the existence of cracks is suggested as a prerequisite for hydrogen brittleness. Such cracks are being formed during electrolytic hydrogen charging which is accompanied by plastic deformation (see, e.g., [22]). However, the gaseous hydrogenation does not cause microcracking. In contrast, the HELP hypothesis [10e15] is based on the effect of the dissolved hydrogen on dislocation properties, namely the increase in mobility of dislocations and decrease in the distance between them in dislocation ensembles. According to calculations within the frame of continuum mechanics [11,15], this effect originates from the hydrogen shielding of elastic stresses caused by dislocations. Another approach to HELP hypothesis is proposed in Refs. [23e27] taking into account the hydrogen-increased concentration of free electrons, which decreases the shear modulus and, correspondingly, eases the start of dislocation sources, decreases the line tension, i.e. the specific energy of dislocations, and the distance between dislocations in the pileups. All these factors facilitate the opening of microcracks and their propagation. The HESIV hypothesis [16e18] manifests “deformationinduced vacancies and their clusters, being enhanced and stabilized by hydrogen, to play the primary role in HE”, whereas “the role of hydrogen in embrittlement is indirect and rather subsidiary”, as it is claimed, e.g., in Ref. [17]. The nucleation of dislocations and their mobility, as well as the nature of hydrogen-caused shear localization are completely ignored in this hypothesis. Its more detailed analysis is beyond the scope of the present paper. The aim is, based on the above mentioned hypotheses, to analyze the role of hydrogen-dislocation interaction in hydrogen embrittlement of two classes of engineering materials, austenitic steels having the fcc crystal lattice and a bphase titanium alloy with the bcc crystal lattice. These objects of studies were chosen for two reasons. First, the both materials are comparatively hydrogen-resistant in contrast to the

a-iron base steels and the hcp Ti alloys. Second, their sensitivity to hydrogen brittleness is strikingly different. Hydrogen can be harmful in austenitic steels even at hydrogen contents of about several wppm, see, e.g. Ref. [28]. In the hcp and two phase hcp þ bcc Ti alloys, hydrogen causes brittleness due to easy formation of hydrides (e.g. Refs. [29,30]). However, in the single phase b Ti alloys with the bcc crystal lattice, hydrogen embrittlement occurs only at extremely high ratio H/M > 0.21 [31]. For this reason, the saturation of b Ti alloys by hydrogen is used to increase their technological plasticity, see e.g. Refs. [32,33]. Such processing is unimaginable for austenitic steels.

Calculations and experimental The calculations of the electron structure were carried for the fcc iron and bcc titanium aiming to estimate the hydrogencaused change in the interatomic bonds. The energy bands, the total potential, the electron density and total energy per cell were calculated using the modern computational programme package Wien2k developed by the European scientific group [34]. These calculations are based on the KohnHohenberg-Sham density functional theory, DFT [35,36]. The initial electron density for solving the KohneSham equations is taken in the local spin density approximation, LSDA [37]. A procedure for solving the KohneSham equations is the full-potential-linearized-augmented plane wave method, FLAPW. The calculation of the exchange-correlation potential is carried out using the generalized gradient approximation in the parameters of Perdew, Burke, and Ernzernhoff [38] where the gradient terms of electron density are added to the exchange-correlation energy and its potential. The integration over the Brillouin zone is carried out using the modified tetrahedron method [39] over the 14  14  14 points in k-mesh. The wave functions, the charge densities and the potential were expanded with L  10 spherical harmonics inside each ‘muffin-tin’ radius, RMT, of 2.17 a.u. and 1.09 a.u. for Fe and H atoms, respectively, and 2.17 a.u. and 1.17 a.u. for Ti and H atoms, respectively. The mentioned values of RMTs and the magnitude of the largest k-vector in the basis set were kept constant for all FeeH and TieH structures for further comparison. All calculations were performed in the full relativistic approach with the spin polarized electron states at the temperature T ¼ 0 K. Self-consistency was achieved when the root-mean square distances between the j- and (j-1) steps of the iteration procedure for the total charges and spin densities were smaller than 1.0  104. The energy values were found with an accuracy of 103 eV. A number of austenitic steels having different compositions were chosen for this study. A reason to vary alloying elements is concerned with their effect on the electron state density at the Fermi level and, correspondingly, on the concentration of free electrons [40]. As the change in the electron structure is considered to control hydrogen embrittlement, it is important to clarify the effect of chemical composition on migration of hydrogen atoms and their binding to dislocations. Steels, of which the compositions are presented in Table 1, were melted in an induction furnace under argon atmosphere as the ingots of 0.175 kg in the weight followed by the

Please cite this article in press as: Teus SM, et al., Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement, International Journal of Hydrogen Energy (2016), http://dx.doi.org/10.1016/ j.ijhydene.2016.09.212

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Table 1 e Chemical compositionsa. Steel Cr15Ni25 Cr15Ni40 Cr25Ni25 Cr15Ni25Cu2 Cr15Ni25Si2 Cr15Ni25Al2 Cr15Ni25Mn15 a

Cr

Ni

Cu

Si

Al

14.66 14.71 24.63 14.69 14.99 15.12 15.31

25.8 40.5 25.9 25.6 25.5 26 25.6

e e e 2.01 e e e

e e e e 1.85 e e

e e e e e 1.98 e

Mn e e e e e 14.77

Fe Balance Balance Balance Balance Balance Balance Balance

Impurity elements are present in usual limits.

homogenization at 1150  C for 24 h. After hot forging, drawing and electrolytic etching, the samples of different diameters, 0.35 mme0.8 mm, have been prepared and subjected to solution treatment at 1050  C for 30 min followed by quenching in water. A single b-phase titanium alloy was obtained by powder metallurgy (PM) approach using the titanium sponge TG-110 produced by Zaporozhye Titanium-Magnium Enterprise. This material was hydrogenated up to hydrogen concentration of 3.5 mass %, which corresponds to the composition of Ti hydride. It was then milled into powder with the size of particles smaller than 100 mm, blended with V-Al-Fe master alloy powder (less than 40 mm) in the ratio needed to reach the Ti10V-2Fe-3Al composition. The obtained powder blend was compacted at room temperature under pressure of 640 MPa. Then, this powder compact has been heated in vacuum up to 1250  C with 5 h holding at this temperature for the synthesis which included dehydrogenation, chemical homogenization and consolidation in one cycle. In more detail, the procedure for obtaining the dense PM titanium alloy using the preliminary hydrogenated titanium and the V-Fe-Al master alloy powders is described in Ref. [41]. Post sintering heat treatment (isothermal exposure at 900  C followed by subsequent water quenching) was used to obtain a single b-phase state of the alloy at room temperature. The samples of 1  1 mm2 in square and 50 mm of length were prepared for measurements using spark cutting followed by the grinding and polishing of their surface. Mechanical spectroscopy was used for studies of hydrogen migration, dislocation mobility and interaction between hydrogen atoms and dislocations. Two kinds of measurements were carried out: temperature dependence of internal friction, TDIF, and its amplitude dependence, ADIF. A frequency shift of the Snoek relaxation peak temperature allows to determine the enthalpy for migration of hydrogen atoms, Hm, by using the Arrhenius equation t ¼ t0exp (Hm/kT), where t is relaxation time, 1/t0 ¼ f0 is the frequency factor. The enthalpy of binding between hydrogen atoms and dislocations, Hb, was obtained due to measurements of ADIF at different temperatures and treatment of the obtained data in the Arrhenius co-ordinates. Measurements were carried out using an automated inverted pendulum operating at frequencies of about 1 Hz at temperatures from 80 to 580 K. The strain amplitude was about 0.5$106 in the measurements of the temperaturedependent internal friction and within 0.5$106 to 5$104 in case of the amplitude-dependent internal friction.

Hydrogen charging was carried out in the aerated 1 N H2SO4 solution containing 100 mg/l NaAsO2 at the current density of 50 mA/cm2 for 72 h. The chosen time of charging could not be sufficient for homogeneous hydrogen distribution, and the obtained hydrogen content is the highest in the surface layer of the samples. Nevertheless, taking into account that torsion deformation by the inverted pendulum in the internal friction relaxometer causes the highest stress in the surface layer, it is reasonable to relate the experimental data obtained using mechanical spectroscopy to hydrogen concentrations estimated by means of X-ray diffraction. As will be shown, hydrogen mobility in the titanium alloy is significantly higher in the titanium alloy in comparison with that in the austenitic steels, which suggests a higher homogeneity of hydrogen distribution for the same charging time. The atomic ratio hydrogen/metal, H/M, in the austenitic steels was estimated using the Baranowski relation [42] between the volume increase and H/M for fcc metals and constituted about 0.2. For estimation of hydrogen content in the bcc b titanium alloy, we used the data obtained by Teter et al. [31] who presented a series of data about the change of the lattice parameter in the b titanium alloy Timetal® 21S obtained by means of X-ray diffraction and the hydrogen content measured using a gas extraction system coupled to a gas chromatograph. According to a correlation derived from their experimental data, the hydrogen content in our titanium samples, H/M, was estimated as 0.036.

Results Electron structure of hydrogen-containing fcc FeeH and TieH solid solutions The density of electron states, DOS, as a function of electron energy is presented in Fig. 1 for different hydrogen contents. The effect of hydrogen on the DOS at the Fermi level is shown in the inserts of Fig. 1a for the g-iron and Fig. 1b for the b-Ti. DOS at the Fermi level characterizes the thermodynamic stability of phases as well as the concentration of free electrons. A monotonous hydrogen-caused increase of the concentration of free electrons in the austenitic steels has been confirmed using the electron spin resonance [43]. In contrast, the effect of hydrogen on the DOS at the Fermi level in the b-Ti is non-monotonous. It decreases in the TiH and TiH2 compositions and this is consistent with formation of the Ti hydride. At its smaller contents, in Ti2H, hydrogen

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Fig. 1 e Effect of hydrogen on the density of electron states, DOS, in the fcc Fe (a) and bcc Ti (b). DOS at the Fermi level, EF, is shown in the inserts. increases DOS and, correspondingly, is expected to increase the concentration of free electrons.

Temperature-dependent internal friction, TDIF The Snoek-like relaxation caused by stress-induced diffusion jumps of hydrogen atoms in their complexes with substitution atoms was used to estimate the enthalpy of hydrogen migration. As example, TDIF for three hydrogen-charged austenitic steels and the titanium alloy are presented in Figs. 2 and 3, respectively. In case of austenitic steels, one can see the Snoek-like relaxation peak located between 200 and 250 K, as well as transient peaks accompanying hydrogen degassing at higher temperatures (Fig. 2). For correct determination of peak temperatures, the experimental data were presented as a function of the inverse temperature, thereafter the background was subtracted and fitting has been carried out. An example of fitting is presented in Fig. 2d. As seen, the Snoek-like peak in austenitic steels consists of two components with the same activation enthalpy Hm and different t0. A reason for that is “the frozen splitting” which appears in case of a rather high content of substitution solutes resulting in the s1-i-s2 neighborhoods for an interstitial atom i with two substitution atoms s. This orthorhombic defect

having <110> orientation in the fcc crystal lattice causes two shear relaxations: in directions <100> with the modulus defect dm<100> ¼ dS44 and <111> with dm<111> ¼ (4/3) d(S11eS12) þ (1/3)dS44. Correspondingly, two different relaxation times t1(S44) ¼ 2n12 þ 4n13 and t1(S11eS12) ¼ 6n13 are expected, where Si are the components of the tensor of elastic compliances and ni are the frequencies. Details are described by Novick and Berry [44], whereas the experimental evidence can be also found in Refs. [45,46] for nitrogen and hydrogen in austenitic steels, respectively. The frequency shift of Snoek-like peak temperature is shown as the inserts in Fig. 2aec. The obtained data of hydrogen migration enthalpy, Hm, are presented in Table 2. The Hm value of about 0.5 eV is typical for hydrogen in austenitic steels, see, e.g., [46,47]. One can see that the change in chemical compositions affects it. Nevertheless, some trend is clearly visible: nickel and copper facilitate hydrogen migration in the fcc iron, whereas chromium, silicon and manganese retard it. Quite different is the hydrogen-caused internal friction in the single b-phase alloy Ti-10V-2Fe-3Al (Fig. 3a). First, the Snoek peak is located at much lower temperature. As consequence, the enthalpy of hydrogen migration, ~0.27 eV, as determined after treatment of spectra obtained at different frequencies, corresponds to a higher mobility of hydrogen atoms in comparison with that in austenitic steels (see Table 2). Second, along with the Snoek peak, the peak of Snoek€ ster relaxation appears, which is expected in metals having Ko the bcc crystal lattice. According to Seeger's theory of SeK relaxation [48], which is generally considered to be the most reliable, its mechanism amounts to the formation of paired kinks accompanied by the migration of hydrogen atoms. The relaxation enthalpy of SeK relaxation, 1.45 ± 0.25 eV, was roughly estimated due to fitting of spectra measured at different frequencies (Fig. 3b). It is equal to the sum of paired kink formation enthalpy and that for hydrogen migration. The obtained value is consistent with the hydrogen migration enthalpy obtained in this study and the enthalpies of SeK relaxation and paired kinks formation in the refractory metals [49,50]. The hydrogen degassing during heating up to 550 K in the course of second and third measurements of the same sample is accompanied by the decrease in the intensity of Snoek and SeK peaks (Fig. 3c). This is the additional confirmation of the hydrogen nature for the both relaxations. Two conclusions can be also derived from Fig. 3c: (i) the PM Ti-10V-2Fe-3Al alloy used in this study contains some residual hydrogen, (ii) hydrogen brought into the solid solution by cathodic charging is not completely released by degassing during heating up to 550 K.

Amplitude-dependent internal friction, ADIF These measurements aimed to study the interaction between hydrogen atoms and dislocations. The effect of hydrogen on the ADIF is shown in Figs. 4 and 5 for steels Cr15Ni25, Cr15Ni40 and in Fig. 6 for the alloy Ti-10V-2Fe-3Al. Beyond the temperature range of relaxation processes, the internal friction background is caused by the vibrations of dislocation segments under applied stress. Its value is the

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Fig. 2 e Hydrogen effect on the internal friction in austenitic steels: (a) Cr25Ni25, (b) Cr15Ni25Si2, (c) Cr15Ni25Al2. Measurements were carried out at different frequencies in order to determine the enthalpy of activation for Snoek-like relaxation (shown in the inserts). As example, fitting of experimental curves is presented in Fig. 2d.

higher the larger the area crossed by the moving dislocations in the course of one cycle of vibrations [51,52]. Four features are distinctive for data presented. First, hydrogen increases the IF background, i.e. enhances mobility of dislocations (Figs. 4a and 5a). The subsequent degassing, as shown in Fig. 5a, restores the initial level of damping with some insignificant excess caused by the increase in dislocation density due to electrolytic hydrogen charging. Second, the internal friction starts to depend on the strain as soon as some definite strain, εcr, is reached. This critical strain is decreased by hydrogen, which is clearly seen in Fig. 5a. Third, the slope DQ1/Dε of the strain-dependent IF part is higher in the hydrogen-charged samples in comparison with that for the hydrogen-free ones. Fourth, with increasing temperature of measurements, the DQ1/Dε slope in the hydrogen-charged samples starts to increase above some critical temperature Tc (see the inserts in Figs. 4b and 5b). The first theory for the amplitude-dependent internal friction was proposed by Granato and Lu¨cke [53]. Their main idea was that the internal friction becomes strain-dependent if the applied stress is sufficient to break dislocations away from their pinning points. Later on [54], the authors have improved their theory taking into account the thermoactivated character of the interaction between dislocations and point defects. However, as shown in the experiments, e.g. Refs. [55,56], and confirmed theoretically, e.g. Ref. [57], the stress needed for

breakaway of dislocations from interstitial atoms in annealed metals is too high and even exceeds the yield strength. Therefore, the emission of new dislocations by dislocation sources occurs before the unpinning of the existing ones. Consequently, the deformation εcr corresponds to the start of microplastic deformation and hydrogen decreases it. In the temperature range where the interstitial atoms are not sufficiently mobile to follow the dislocations, the slope DQ1/Dε in the strain-dependent part of the Q1(ε) curves is controlled by the intersection of the pinned dislocation forest by the moving dislocations and it is only slightly depends on temperature. If the temperature increases, so that interstitial atoms can follow dislocations during their vibrations, the slope DQ1/Dε should be substantially dependent on the migration of interstitial atoms and their binding to dislocations, which results in the temperature dependent DQ1/Dε, as shown in the inserts to Figs. 4b and 5b. The critical temperature Tc corresponds to the condensation of hydrogen atoms around dislocations. The further heating makes these hydrogen clouds diluted depending on the enthalpy of binding between hydrogen atoms and dislocations. Therefore, treating in the Arrhenius co-ordinates the experimental data of Q1(ε) at different temperatures and supposing that DQ1/Dε z c⊥/c0 ¼ exp (Hb/kT), where c⊥ and c0 are the concentrations of interstitial atoms at dislocations and far from them, respectively, one can find the binding enthalpy Hb (see Table 3).

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Table 2 e Activation enthalpy for hydrogen migration, dH ¼ ±0.005. Material Cr15Ni25 Cr15Ni40 Cr15Ni25Cu2 Cr15Ni25Al2 Cr15Ni25Si2 Cr15Ni25Mn15 Cr25Ni25 Ti10V-2Fe-3Al

Fig. 3 e Hydrogen-caused internal friction in the Ti-10V2Fe-3Al alloy as measured at different frequencies (a), the fitting of spectrum (b) and the effect of repeated heating (c).

The measurements of ADIF in the titanium alloy turned out to be more complicated because the DQ1/Dε experimental data at temperatures below the condensation temperature Tc fall within the temperature range of extremely huge Snoek relaxation, which does not allow to determine Tc (compare Figs. 3a and 6). The Snoek peak in the hydrogen-charged titanium alloy is ten times as intensive as that in austenitic steels, see Figs. 2 and 3. Usually, the relaxation strength does

H1, eV

H2, eV

0.544 0.496 0.517 0.555 0.563 0.567 0.589 0.27 ± 0.01

0.543 0.495 0.516 0.557 0.562 0.567 0.590 0.28 ± 0.01

Fig. 4 e Amplitude-dependent internal friction in steel Cr15Ni25: (a) effect of hydrogen, (b) effect of temperature in the hydrogen-charged samples. The insert in Fig. 4b shows how mobility of dislocations depends on the temperature. Tc is the condensation temperature of hydrogen clouds around the dislocations, above which hydrogen atoms accompany moving dislocations and. not essentially depend on the strain amplitude. Nevertheless, the Snoek relaxation can be enhanced in presence of dislocations [58]. At the same time, the enthalpy Hb of binding between hydrogen atoms and dislocations in the titanium alloy can be correctly determined from the slope DQ1/Dε because the corresponding experimental data are obtained at tempera€ ster relaxation tures beyond the both Snoek and Snoek-Ko peaks. Results obtained are presented in Table 3.

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Table 3 e Binding enthalpy for hydrogen atoms and dislocations and condensation temperature Tc. Material Cr15Ni25 Cr15Ni40 Cr15Ni25Cu2 Cr15Ni25Al2 Cr15Ni25Si2 Cr15Ni25Mn15 Cr25Ni25 Ti-10V-2Fe-3Al

HB, ±0.005 eV

Tc, ±3 K

0.107 0.096 0.102 0.110 0.112 0.114 0.120 0.035

173 163 174 173 186 182 189 e

temperature Tc is expected to be located at much lower temperatures.

Discussion

Fig. 5 e Amplitude-dependent internal friction in steel Cr15Ni40: (a) effect of hydrogen charging, (b) effect of temperature of measurements in the hydrogen-charged sample.

Fig. 6 e Effect of temperature on the amplitude-dependent internal friction in the hydrogen-charged Ti-10V-2Fe-3Al alloy. Treatment in the Arrhenius co-ordinates is presented in the right upper corner. It is seen that the enthalpy of binding between hydrogen atoms and dislocations in the titanium b-phase alloy is by one order of magnitude smaller in comparison with that in austenitic steels. Correspondingly, the condensation

The obtained results allow to analyze hydrogen embrittlement of austenitic steels and b titanium alloy in terms of hydrogen effect on the electron structure, migration of hydrogen atoms and their interaction with dislocations. The hydrogen-increased DOS at the Fermi level enhances the metallic character of interatomic bonds, which should result in a decrease of the shear modulus. As mentioned in the Introduction, the electron concept of hydrogen embrittlement [23e27] supports the HELP theory earlier developed in Refs. [10e15] within the elasticity approach with the only distinction that, alternative to the shielding of elastic interaction between the dislocations, a decisive role is attributed to the hydrogen-caused change in atomic interactions. The hydrogen migration in the solid solution and the interaction between hydrogen atoms and dislocations are important in the HELP theory. A significance of these factors is refused in the AIDE theory [5e9], where only the adsorption of hydrogen atoms at the crack tips surface is decisive for initiation of dislocation emission, whereas their further slip is not controlled by hydrogen dissolved in the crystal lattice. Let us discuss the obtained data on hydrogen-dislocation interactions in the studied alloys. The hydrogen-induced decrease of the strain above which the internal friction depends on the strain amplitude is consistent with the hydrogen effect on the electron properties. As shown in Part 3.1, hydrogen increases the density of electron states at the Fermi level. According to [26,27], the hydrogen atoms in the octahedral interstitial sites are surrounded by the clouds of conduction electrons. Along with the increase in the concentration of free electrons throughout the crystal lattice, this should result in the early start of dislocation sources and enhance the dislocation slip, particularly at the condition that hydrogen atoms accompany the moving dislocations. A remarkable feature of hydrogen brittleness is its appearance within a definite range of temperatures and strain rates (e.g. Refs. [59,60]), which strongly supports the idea of joint movement of dislocations and hydrogen clouds as a precondition of embrittlement. Therefore, the critical microplastic strain εcr should be decreased and the slope of amplitude-dependent internal friction is expected to be increased by hydrogen. The slope

Please cite this article in press as: Teus SM, et al., Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement, International Journal of Hydrogen Energy (2016), http://dx.doi.org/10.1016/ j.ijhydene.2016.09.212

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DQ1/Dε is proportional to the area swept by dislocations during their oscillations. In other words, this slope characterizes velocity of dislocations which is increased by hydrogen. As follows from Tables 2 and 3, some correlation exists between the migration enthalpy Hm of hydrogen atoms, the enthalpy Hb of their binding to dislocations and condensation temperature Tc. Nickel enhances the migration of hydrogen atoms, weakens the hydrogen-dislocation binding and decreases the condensation temperature. Silicon, manganese and chromium retard the hydrogen migration, enhance hydrogen-dislocation binding and assist the hydrogen condensation at dislocations. This correlation is consistent with the effect of alloying elements on the electron structure [40] and mechanical properties of hydrogen-charged austenitic steels [26,27]. It should be also noted that a negative effect of nickel at its high contents is at variance with its generally accepted positive influence on the hydrogen resistance of austenitic steels. At the same time, a harmful nickel effect on the hydrogen resistance of high nickel austenitic steels and the inconel alloys is confirmed by the experiments with gaseous [61] and electrolytic [62] hydrogen charging. Nickel stabilizes the gamma phase in steel which is more stable to hydrogen brittleness in comparison with the alpha one. However, as nickel increases the concentration of free electrons in the iron-based alloys [40], one can suppose that the corresponding increase in mobility of dislocations combined with the same effect of hydrogen prevails over the stabilizing Ni effect on the austenitic structure. Therefore, particularly interesting is a possible correlation between the condensation temperature Tc, the enthalpies of hydrogen migration Hm, binding to dislocations Hb and the temperature range of hydrogen brittleness. Above Tc, the hydrogen atoms start to follow the dislocations in the course of dislocation slip. At the same time, the hydrogen clouds around the dislocations start to be diluted. Both these phenomena depend on the hydrogen migration enthalpy Hm and the hydrogen-dislocation binding enthalpy Hb. The larger Hm and Hb the higher is the temperature above which the hydrogen atoms can follow dislocations. Correspondingly and in consistency with HELP theory, the increased Tc is expected and the temperature range of hydrogen brittleness should be shifted to higher temperatures. Thus, at the condition of constant strain rate, the condensation temperature Tc characterizes a low temperature limit for hydrogen embrittlement. At the same time, hydrogen clouds start to be diluted above Tc, the binding between hydrogen atoms and dislocations weakens and the density of hydrogen clouds decreases. Therefore, as mobility of hydrogen atoms and their binding to dislocations control dilution of hydrogen clouds, the upper limit in the temperature range of hydrogen brittleness will be the lower the smaller Hb and Hm are. With increasing hydrogen content, this temperature range should be increased. The data of Teter et al. [31] on the Timetal® 21S alloy can be used for a qualitative comparison of the temperature range for hydrogen brittleness with the data of Hm and Hb obtained in the present study for the Ti-10V-2Fe-3Al alloy of nearly the same composition. In Ref. [31], according to results of tension tests, the hydrogen-induced brittle transition started at 248 K for the

ratio H/M > 0.21. To obtain the brittle fracture at 298 K, the ratio H/M > 0.27 was needed, and the brittleness disappeared if the tests were carried out at 373 K. Substantial is the extremely high hydrogen content needed to initiate the embrittlement of the titanium-based alloy. For comparison, as mentioned above, the hydrogen brittleness of austenitic steels appears at the H contents exceeding 2 to 3 wppm [28], despite the retarding effect of the fcc crystal structure on the hydrogen migration. In accordance with the foregoing analysis, the small values of hydrogen migration enthalpy Hm and hydrogen binding to dislocations Hb in the titanium alloy should be responsible for the low temperature range, within of which the hydrogen brittleness manifests itself. The increase in the hydrogen content shifts its temperature range to higher temperatures. This is why so much hydrogen is needed to cause brittleness of the b titanium alloy at 248 K, and this critical value increases with increasing temperature. For this reason, a softening effect of hydrogen on the mechanical properties of the b titanium alloy can be revealed at ambient temperatures. Really, at hydrogen contents of about 0.3e1 mass %, see e.g. Ref. [32], hydrogen improves plasticity of titanium alloys and is used as a temporary alloying element in the course of their processing (see also [33]). Teter et al. [31] concluded that a mechanism for the hydrogen-caused brittle fracture of the Timetal® 21S alloy corresponds to hydrogen adsorption at the crack surfaces in consistency with the HEDE theory. However, this conclusion was not really supported by the analysis of fracture surface. At variance with the HEDE and AIDE theories, the experimental data obtained in the present study and supported by the ab initio calculated softening of the crystal lattice in presence of hydrogen give the evidence that, in both classes of engineering materials, austenitic steels and titanium alloys, the hydrogen-dislocation interaction controls hydrogen brittleness in accordance with HELP theory.

Conclusions 1. Hydrogen in the g-iron and b-titanium increases the density of electron states, DOS, at the Fermi level, which corresponds to the increase of the concentration of free electrons and enhances metallic character of interatomic bonds. In case of titanium, the decrease of DOS at rather high hydrogen contents occurs because of hydride formation, which does not occur in the iron. 2. There is a critical temperature, Tc, above which the hydrogen clouds around the dislocations can follow dislocations in the course of plastic deformation. The further increase in temperature is accompanied by the dilution of hydrogen clouds. 3. The hydrogen-dislocation interaction controls hydrogen brittleness due to the local increase in the concentration of free electrons within the hydrogen clouds around the moving dislocations. 4. The enthalpy of hydrogen migration in the austenitic steels is by two times larger and the enthalpy of hydrogen binding to dislocations is by one order of magnitude as large as that in the titanium alloy. Correspondingly, the temperatures for condensation of hydrogen clouds around the

Please cite this article in press as: Teus SM, et al., Hydrogen migration and hydrogen-dislocation interaction in austenitic steels and titanium alloy in relation to hydrogen embrittlement, International Journal of Hydrogen Energy (2016), http://dx.doi.org/10.1016/ j.ijhydene.2016.09.212

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dislocations and that for release of dislocations from hydrogen atoms are expected to be significantly lower in the titanium alloy. This decreases the temperature range of hydrogen brittleness down to temperatures below the room one and allows to use the hydrogen-caused softening to improve deformability of titanium alloys in the singlephase b state in the course of their processing.

Acknowledgements This study was supported by National Academy of Sciences of Ukraine, Project No. III-10-13.

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