Influence of Gaseous Hydrogen on Fatigue Behavior of Ferritic Stainless Steel

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Procedia Engineering 133 ( 2015 ) 362 – 378

6th Fatigue Design conference, Fatigue Design 2015 Influence of gaseous hydrogen on fatigue behavior of ferritic stainless steel – a fatigue-life estimation

Georg Schauera,*, Jens Roettinga, Malte Hahnb, Simone Schreijaega, Manfred Bacher- Höchsta, Stefan Weihec

aRobert Bosch GmbH, Robert-Bosch-Campus 1, 71272 Renningen, Germany bL’Orange GmbH, Porschestraße 30, 70435 Stuttgart, Germany cMaterials Testing Institute University of Stuttgart, Pfaffenwaldring 32, 70569 Stuttgart (Vaihingen), Germany

Abstract

The Bosch Group, a leading global supplier of technology and services develops reliable and efficient fuel cell components for future mobility solutions. This work studies the fatigue behavior of ferritic stainless steel 1.4005IA optimized for electromagnetic properties under varying environmental conditions. Fatigue-life of air-tested specimens are compared to specimens which are electrochemically hydrogen precharged and loaded in a pressurized gaseous hydrogen atmosphere. The fracture surfaces of broken specimens are investigated using optical microscope and SEM in order to analyze the failure mechanisms. Furthermore, the hydrogen contents of specimens before and after static and cyclic testing in gaseous hydrogen are measured. Based on strain controlled fatigue tests using unnotched specimens, the influence of hydrogen on cyclic material parameters is determined. Using these parameters, a fatigue-life estimation for notched specimens based on the local strain approach is carried out. Different estimation approaches are evaluated, by comparing experimental data and predicted fatigue-lifes. A modified material model, based on the four-parameter equation according to Manson, Coffin and Morrow is proposed to prescribe the influence of hydrogen on cyclic behavior of ferritic stainless steel in low gas pressure applications (up to 1 MPa). The proposed approach including the modified material model is in a good agreement to the experimental data. This opens a pathway for a reliable component design process for industrial applications.

© 2015 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (©http://creativecommons.org/licenses/by-nc-nd/4.0/ 2015 The Authors. Published by Elsevier Ltd).. PeerPeer-review-review under under responsibility responsibility of CETIM of CETIM.

Keywords: fuel cell; gaseous hydrogen; fatigue design; fatigue-life estimation; ferritic stainless steel; local strain approach; electrochemical hydrogen charging; TDS analysis; hydrogen content; hydrogen embrittlement; hydrogen assisted cracking; Coffin-Manson;

* Corresponding author. Tel.: +49-711-811-10955 ; fax: +49-711-811-51810955 . E-mail address: [email protected]

1877-7058 © 2015 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/). Peer-review under responsibility of CETIM doi: 10.1016/j.proeng.2015.12.669 Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 363

1. Introduction

Fuel cell vehicles (FCV) offer a zero-emission mobility and are one of the future mobility concepts. In fuel cell (FC) systems, components are subjected to pressurized hydrogen gas and cyclic loading, which is the most common load type in automotive applications. For a safe design of FC components exposed to gaseous hydrogen atmosphere the influence of hydrogen on the fatigue behavior of structural steel components has to be quantified. A fatigue design approach for components exposed to gaseous hydrogen, as well as a database of material properties determined in hydrogen atmosphere is essential. The effect of hydrogen on mechanical properties of steels has been investigated for decades. The hydrogen- induced decrease of steel properties such as reduced ductility and tensile strength, subcritical crack growth under monotonic loading, and increased fatigue crack growth (FCG) rates, formally known as hydrogen embrittlement (HE) can cause detrimental effects on structural components [1]. Since then several HE failure mechanisms have been proposed: hydrogen-enhanced decohesion (HEDE) [2,3], hydrogen-enhanced localised plasticity (HELP) [4,5], adsorption-induced dislocation-emission (AIDE) [6], hydride embrittlement [7], pressure theory [8] and surface adsorption theory [9]. The relevance of each HE failure mechanism is not yet clarified and many authors suggest that HE mechanisms occur simultaneously [1]. The influence of hydrogen on mechanical properties of steel components is related to x hydrogen uptake, diffusivity and solubility x microstructure: especially austenitic, martensitic and ferritic, as well as the appearance of inclusions and lattice defects in general x material properties: especially material strength and failure mechanism x loading parameters: type of mechanical loading, testing time, hydrogen gas pressure and hydrogen content However, the influence of hydrogen on fatigue-life and limit of steels is less investigated, which is relevant for the design of automotive components. In low cycle fatigue (LCF) regime usually a hydrogen-induced decrease of fatigue-life has been reported [8]. In high cycle fatigue (HCF) regime a decrease [10] but also a slight increase of fatigue limit have been observed [11, 12]. Moreover, the influence of hydrogen on metal fatigue behavior and its interaction of testing conditions is not yet clearly understood. Murakami et al. [13] reported a hydrogen-induced localized slip band movement and an increase of micro-crack propagation of SUS304 and SUS316 steel due to hydrogen charging, whereas only little effect was observed at SUS405 and SUS316L steel. Lee et al. [14] reported that micro-crack propagation increases, threshold value decreases and slip bands develop faster in hydrogen environment of S10C steel. Crack growth behavior in hydrogen environment was marked by a coalescence of new micro-crack generated ahead of a crack tip, whereas in nitrogen the crack grew continuously in the same total strain range. Uyama et al. [15] observed a hydrogen-induced change in cyclic hardening/softening behavior as a decrease in strain amplitude in the hysteresis loop of hydrogen charged S45C steel specimens with a hydrogen content of 0.5 wppm (weight parts per million). Specimens were cycled with constant stress amplitude and a hydrogen-induced localization of slip bands was found. However no hydrogen effect was observed at S10C steel with a hydrogen content of 0.2 wppm. Han and coworkers [16] observed that 2.25Cr-1Mo steel specimens, charged with a hydrogen content of 4.5 wppm fail from inclusions, whereas the crack of uncharged specimens initiates at the surface. At fatigue test on CT specimens Suresh and Ritchie [17] observed that threshold stress intensity factor range (ΔKth) is reduced in hydrogen environment at load ratio R=0.05. No hydrogen effect was observed at higher load ratio of R=0.75 due to a modified crack closure behavior. Other authors made comparable observations [18, 19]. Nanninga et al. [20] reviewed the influence of gaseous hydrogen on FCG of pipeline steels and concluded that FCG of ferritic steels increases with increasing hydrogen gas pressure. Furthermore, the influence of notches on HE susceptibility [21] and the interaction between three dimensional stress fields and the solute hydrogen atoms is still under discussion [22]. In this paper the fatigue behavior of ferritic stainless steel 1.4005IA is investigated. The influence of hydrogen on the four-parameter equation according to Manson, Coffin and Morrow is demonstrated and a fatigue design 364 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

approach based on the local strain approach for notched specimens is evaluated for its applicability. In addition the influence of hydrogen on fatigue properties of ferritic steels and fatigue mechanisms are discussed.

2. Experimental methods

2.1. Material

Commercial cold rolled ferritic stainless steel 1.4005IA optimized for electromagnetic properties is investigated. The chemical composition is shown in Table 1, the microstructure in Fig. 1. Its microstructure consists of ferritic grains with an average grain size of 63 to 88 μ m in longitudinal direction and an aspect ratio of 1.5:1, according to DIN EN ISO 643 [23]. Ferritic grains are interspersed with carbides, mostly on grain boundaries and long manganese sulphides (MnS) stretched in rolling direction.

(a) (b) Fig. 1. Longitudinal cross section of ferritic stainless steel 1.4005IA round bars with low (a) and high magnification (b). Long manganese sulphides (MnS) and carbides are visible inside and also on the grain boundaries.

Table 1. Chemical Composition of ferritic steel 1.4005IA in mass%. C Mn Si P S Cr Mo 1.4005IA 0.012 0.81 0.64 0.021 0.31 11.9 0.3

2.2. Testing Procedure

Round specimens are machined from the center of the cold rolled round bars with a diameter of 20 mm. Last manufacturing step was precision turning, resulting in a maximum average surface roughness of Rz=2.5 μ m. Specimen geometries of unnotched and notched specimens (stress concentration factor Kt=2.06 and stress gradient χ*=2 mm-1) are shown in Fig. 2. The small cylinder at the end of the notched specimens was used to measure hydrogen content using thermal desorption spectroscopy (TDS).

Specimens designated for testing in hydrogen atmosphere were electrochemically hydrogen precharged in a 0.5 Mol/l aqueous solution of NaOH. Current density was set to 1 mA/cm² and the charging time to 24 hours. After precharging specimens were stored in liquid nitrogen to prevent hydrogen effusion. Before cyclic loading in hydrogen atmosphere specimens were normalized to room temperature in an isopropanol bath for 6 minutes.

Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 365

Ø6 Ø12 Ø10 15 Ø6 90°

(a) (b) Fig. 2. (a) Unnotched specimens for tensile and strain controlled fatigue testing and (b) notched specimens for stress controlled fatigue tests. The small cylinder at the end of the notched specimen was used to measure the hydrogen content using TDS.

Tensile and fatigue tests in hydrogen atmosphere and strain controlled fatigue tests were carried out on a universal servo-hydraulic testing machine with a hydrogen autoclave chamber, as illustrated in Fig 3(a). After flushing the autoclave for 180 s with nitrogen (1500 l/min) the testing chamber was flooded with hydrogen (3000 l/min) for 100 s. Finally, a hydrogen pressure of pH2=1 MPa was applied. The purity of hydrogen gas was 99.9996 %. To ensure hydrogen diffusion into the bulk material during testing a low strain rate and low test frequency were used. To determine tensile properties in air and hydrogen atmosphere, constant extension rate tensile tests (CERT) with a strain rate of ˙ε=1.23∙10-4 s-1 were used. Test frequencies for fatigue tests in hydrogen atmosphere were 1 Hz up to 200.000 cycles and 20 Hz afterwards to achieve a maximum number of cycles of 5∙106 cycles (see Fig. 3(b)). For tensile tests and strain controlled fatigue tests an extensometer with a gauge length of 12.5 mm was applied.

load cell 10 9 1 Hz Air 1 Hz, 20 Hz Hydrogen (1 MPa) 8 160 Hz hydrogen inlet 7 6 5

upper specimen 4 6 10

holder ∙ 3

Testing time /days Testing time 2 autoclave chamber 1 N = 5 sealing joint and 0 cylinder rod extension 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+7 1.0E+8 Number of cycles, N (a) (b)

Fig. 3. Testing concept for mechanical tests in gaseous hydrogen atmosphere of 1 MPa gas pressure. (a) Servo-hydraulic testing machine with lowered hydrogen autoclave chamber. No specimen is inserted. (b) Testing time depending on the number of cycles for different test frequencies.

Fatigue tests in ambient air of notched specimens were carried out on a resonance testing machine at a frequency of 160 Hz due to time saving reasons, as illustrated in Fig. 3(b). Maximum number of cycles in air was set to 107 cycles. All tests were carried out at temperature of T=20 °C. The notched specimens were loaded with constant stress amplitudes until fracture. A sinusoidal waveform with a stress ratio Rσ=σmin/σmax=0.1 was used. The unnotched specimens were cycled with constant total strain amplitude. A sinusoidal waveform with a strain ratio Rε=εmin/εmax=-1 was applied. The number of cycles to failure (Nf) was defined by a reduction of maximum stress per cycle by 10 % from stabilized condition. Stabilized condition was obtained in midlife cycle (Nf/2). To measure the hydrogen content after testing in hydrogen atmosphere, unnotched specimens loaded until fracture were used. Broken specimens were immediately stored in liquid nitrogen. To avoid heating the cooled 366 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

specimens were cut with a bolt cutter in the reduced section up to a length of 5 to 8 mm. The separated piece was directly measured with TDS. The heating rate for TDS was set to 0.1 K/s. After careful calibration with TiH2, the hydrogen content can be quantitatively determined from the TDS spectrum [24].

3. Experimental results

3.1. Hydrogen precharging

In service FC components are subjected to a pressurized hydrogen atmosphere for several years. In order to increase the hydrogen content of the specimens an electrochemical hydrogen precharging process before mechanical testing in hydrogen atmosphere is performed. Fig. 4(a) shows a deconvolved TDS spectrum of a cylindrical specimen charged for 24 hours. Several peaks in different temperature ranges can be observed. TDS spectra were deconvolved with four Gaussian distributions. The peak temperature correlates with hydrogen binding energies and the area of each signal with the amount of hydrogen. Every peak represents a different hydrogen binding type which can be attributed to atomic hydrogen in lattice or microstructural trapping sites [25]. Peak 4 was also observed at uncharged specimens, also referred to as residual hydrogen. In Fig. 4(b) the development of the hydrogen content after different charging times is shown. Total hydrogen content and hydrogen content of peak 2 signal increases with increasing charging time. Hydrogen content of peak 1 signal saturates after several minutes. The first peak was not observed in specimens that were degassed in a desiccator for one week prior to testing.

2.0E+14 1.2 1.4005IA measured 1.4005IA Total Charging time: 24 h deconvolved 1.0 Heating rate: 0.1 K/s Peak 2 Heating rate: 0.1 K/s Specimen:Specimen Peak 1

Specimen: Ø12 0.8 Ø12

1.0E+14 0.6 2 0.4 1 3 0.2

4 /wppm content Hydrogen

Desorption rate /(molecules/g/s) Desorption 0.0E+00 0.0 100 200 300 400 500 600 700 0110 100 Temperature /°C Charging time /h

(a) (b)

Fig. 4. Measured hydrogen content after electrochemical hydrogen precharging. (a) TDS spectra of a 24 h electrochemically hydrogen precharged cylindrical specimen with a diameter of 12 mm. The spectrum is deconvolved into several peaks. (b) Total hydrogen content and hydrogen content of peak 1 and 2 after different charging times.

3.2. Tensile tests

Tensile tests are a fast and cost effective method to evaluate the material susceptibility to HE, usually by comparing the decrease of reduction of area (RA) in hydrogen and the reference atmosphere (e.g. air or helium). Hydrogen does not significantly affect either Young’s modulus (E), 0.2% yield strength (0.2% YS), ultimate tensile strength (UTS) and RRA, of ferritic steel 1.4005IA, as shown in Table 2. A slight reduction of total elongation can be observed.

Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 367

Table 2. Results of tensile tests conducted in air and hydrogen atmosphere. Atmosphere E /GPa 0.2%YS /MPa UTS /MPA Total Elongation /% RA /% air 176 432 438 37.3 40.6 hydrogen 177 433 440 32.4 39.6

3.3. Fatigue tests

In order to investigate the fatigue behavior of ferritic steel 1.4005IA stress and strain-controlled fatigue tests were conducted in both, air and hydrogen atmosphere. Stress controlled S-N curve obtained were evaluated with PROBIT-method [26, 27]. The failure probability (PF) of 10%, 50% and 90% is determined in LCF regime. The very high cycle fatigue (VHCF) regime was not investigated in this work. Although, only a slight reduction of ductility in tensile tests was observed, the fatigue-life in LCF regime of notched specimens decreases up to one magnitude in hydrogen atmosphere (see Fig. 5(a)). In general the detrimental effect of gaseous hydrogen increases with increasing nominal stress amplitude. At a nominal stress amplitude of Sa=300 MPa only 15% of the fatigue-life in hydrogen remained compared to the fatigue-life in air. No significant hydrogen influence can be observed on fatigue limit. Additionally, strain controlled fatigue tests were conducted in air and hydrogen atmosphere and were approximated with the relationship proposed by Manson [28], Coffin [29] and Morrow [30]. The equation is composed of an elastic and a plastic component of the endurance strain amplitude and can be written as:

' V f HHH N b H ' N)2()2( c , (1) ata el a,,, pl E f

with εa,t the total strain amplitude, εa,el the elastic strain amplitude, εa,pl the plastic strain amplitude, E the elasticity modulus, N the number of cycles, Vf̍ the fatigue strength coefficient, b the fatigue strength exponent, ε̍f the ductility coefficient and c the fatigue ductility exponent. The S-N curve is illustrated in Fig. 5(b). A significant hydrogen- induced reduction of fatigue-life was found at a total strain amplitude of εa,t=0.3 %. No significant influence was observed at εa,t=0.2 %. The reduction of fatigue-life in LCF regime involves a significant decrease of plastic component. The influence of hydrogen on the elastic component was not significant.

350 10.00 a

S Air Air Hydrogen (1MPa) /% Hydrogen (1MPa) 300 a,t ε 1.4005IA 1.4005IA R =0.1 1.00 R =0.1 250 V ε notched (Kt=2.06) unnotched T=20 °C T=20 °C

/MPa 200 fracture x1 0.10 runout x1 150 PF=10% total PF=50% VHCF not elastic investigated Nominal stress amplitude, amplitude, stress Nominal plastic PF=90% Total srain amplitude, 100 0.01 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+7 1.0E+8 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+7 Number of cycles, N Number of cycles, N

(a) (b)

Fig. 5. S-N curves of ferritic stainless steel 1.4005IA conducted in air and hydrogen atmosphere at pH2=1 MPa and T=20 °C. (a) Stress controlled

S-N curves of notched specimen (Kt=2.06) tested at RV=0.1. The S-N curves were evaluated with PROBIT-method with maximum number of cycles 1∙107 in air and 5∙106 in hydrogen atmosphere. VHCF regime was not investigated (b) Strain controlled S-N curves of unnotched specimens tested at Rε=-1. 368 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

The cyclic stress strain curves, are evaluated using Ramberg-Osgood equation [31]:

nc c ' V f §VV aa · , b K ata el HHH a,,, pl ¨ ¸ with nc and nc (2) © KE c ¹ c H cf

where Va is the stress amplitude, K ̍ the cyclic hardening coefficient and n ̍ the cyclic strain hardening exponent. Cyclic stress strain curve in hydrogen corresponds to that in air, as illustrated in Fig. 6(a). There is a slight increase of stabilized stress amplitude at εa,t=0.2 % and εa,t=0.3 % in hydrogen. This is a result of evaluating the stabilized stress amplitude at Nf/2. The cyclic hardening/softening behavior of material investigated is not influenced by hydrogen, as illustrated in Fig. 6(b). In Fig. 6(b) the cyclic flow curves of fatigue tests conducted at a strain amplitude of εa,t=0.3 % in air and hydrogen atmosphere are compared. The material shows pronounced softening behavior. The stress amplitude in hydrogen falls abruptly after approximately 6800 cycles. In air the stress amplitude continues to decrease steadily and decreases significantly after approximately 21.000 cycles. Cyclic parameters obtained of strain controlled fatigue tests are shown in Table 3.

600

a 600

σ Air Air Hydrogen (1MPa) 500 400 Hydrogen (1MPa)

400 200 /MPa

300 σ 0 /MPa 200 -200 1.4005IA

1.4005IA stress , unnotched unnotched 100 Rε=-1 R =-1 -400 ε εa,t=0.3 % ° Stabilized stress amplitude, Stabilized stress amplitude, T=20 C T=20 °C 0 -600 0.0 0.5 1.0 1.5 1.0E+0 1.0E+1 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 Total strain amplitude, ε /% Number of cycles, N a,t (a) (b)

Fig. 6. (a) Stabilized cyclic stress strain curve and (b) flow curves conducted at a strain amplitude of εa=0.3 %. Unnotched specimens of ferritic

stainless steel 1.4005IA were tested in air and hydrogen atmosphere at Rε=-1, pH2=1 MPa and T=20 °C.

Table 3. Cyclic parameters of ferritic stainless steel 1.4005IA. Unnotched specimens of ferritic stainless steel 1.4005IA were tested in air and

hydrogen atmosphere at Rε=-1, pH2=1 MPa and T=20 °C.

Atmosphere Vf̍ /MPa b ε̍f /- c K̍ /MPa n̍ air 468.5 -0.0426 0.108 -0.4167 588.0 0.102 hydrogen 379.1 -0.0236 0.0194 -0.2948 519.8 0.080

3.4. Fracture surface analyses

Fracture surfaces of broken notched specimens were investigated by optical and scanning electron microscope (SEM). In Fig. 7 fracture surfaces of broken notched specimens tested in air and hydrogen atmosphere in LCF regime (Sa=300 MPa) are compared. In air multiple crack origins form a ring. In hydrogen only a few crack origins can be observed. Fracture surfaces of air tested specimens reveal rough fatigue fracture surfaces (see Fig. 7(c)), whereas hydrogen tested specimens show cleavage like fracture surfaces (see Fig. 7(d)). Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 369

Air Hydrogen Air Hydrogen (c) (d)(d)

(a) (b) (c) (d)

Fig. 7. Fracture surfaces of notched specimens tested in air (a, c) and hydrogen (b,d) atmosphere at Sa=300 MPa, RV=0.1, pH2=1 MPa and T=20 °C. Overview of fracture surfaces (a,b) and crack initiation area (c,d). Number of cycles until fracture were N=14,000 in air and N=3,359 in hydrogen atmosphere. Arrows indicate the area of crack initiation.

Fracture surfaces of notched specimens that were tested at in HCF regime (Sa=170 MPa) are illustrated in Fig. 8. Mostly, one or two crack origins from specimen surface in air and hydrogen were observed. Fatigue fracture mode near crack origin is intergranular independently of testing atmosphere (see Fig. 8(a,b)). The ratio of intergranular T fracture mode increases significantly of hydrogen tested specimens after a crack length of a max=230 to 270 μm (percentage intergranular fracture surface, air: 5-15 %, hydrogen: 60-70 %). At all specimens crack initiated at the surface in both, air and hydrogen atmosphere tested specimens. The critical crack length (acrit) was not significantly influenced by hydrogen (see also Table 5). Air Hydrogen Air Hydrogen

(a) (b) (c) (d)

Fig. 8. Fracture surfaces of notched specimens tested in air (a, c) and hydrogen (b,d) atmosphere at Sa=170 MPa, RV=0.1, pH2=1 MPa and T=20 °C. Area of fracture surfaces of stable fatigue crack growth (a,b) and crack initiation area (c,d). Number of cycles until fracture were N=1,531,000 in air and N=1,023,000 in hydrogen atmosphere. Arrows indicate the area of crack initiation.

3.5. Hydrogen content after testing

As expected, the total hydrogen content of tensile tested specimens increased by almost 36 % compared to the charged state, as illustrated in Fig. 9. The total hydrogen content of cyclically tested unnotched specimens does not increase significantly compared to the precharged state. Likewise, strain amplitude and average testing time in hydrogen atmosphere does not significantly increase the total hydrogen content of cyclic tested specimens. A small decrease of total hydrogen content of unnotched specimens loaded with εa,t=0.3 % can be observed. It can be concluded that with the chosen precharging parameters an equilibrium of saturation is achieved. This means that the precharged state should be representative of a component which has been in service under hydrogen (1 MPa) for years. 370 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

Average testing time in hydrogen atmosphere /h Total hydrogen content Average testing time 0 15 30 45 60 75 90 in hydrogen atmosphere cylinder specimen cylinder specimen, charged charged cut unnotched specimen Kt1 specimen,charged, charged, tensile tested tensile in hydrogen test atm. cut unnotched specimen Kt1 specimen, charged, cycled, Rε=-1, … charged, fatigue tested in hydrogen atm. (Rε=-1, εa,t=0.3%) Kt1 specimen, charged, cycled,cut unnotched Rε=-1, specimen … charged, fatigue tested in hydrogen atm. (Rε=-1, εa,t=0.2%) 0.0 0.2 0.4 0.6 0.8 1.0 1.2 Total hydrogen content /wppm

Fig. 9. Total hydrogen content of hydrogen precharged cylindrical specimens and cut unnotched specimens. In the presence of hydrogen unntoched specimens were loaded in tensile and fatigue tests until fracture. The unnotched specimens were cut with a bolt cutter to avoid heating of the specimens. The separated pieces were directly measured with TDS.

4. Fatigue-life Estimation

4.1. Local strain approach

Fracture surface investigations reveal fatigue crack initiation from specimen surface regardless of testing atmosphere. In order to estimate the fatigue-life of notched specimens (Kt=2.06) the local strain approach [32] is used. Calculation procedure of the local strain approach is shown in Fig. 10. In the first step local stresses and strains in the notch are calculated using the cyclic stress strain curve obtained of strain controlled fatigue tests of unnotched specimens. Afterwards a damage model (e.g. Smith Watson Topper [33]) using local stress/strains and cyclic parameters obtained from strain S-N curves are used to estimates the fatigue-life until a ‘technical crack’ has initiated (crack length 0.5 to 1mm [34,35]). The hydrogen influence in LCF regime is modeled with cyclic parameters obtained of strain controlled fatigue tests in hydrogen atmosphere with a decreased plastic strain component compared to air.

Fig. 10. Calculation procedure for fatigue-life estimation in hydrogen atmosphere according to local strain approach following Seeger [32].

4.2. Local stresses and strains

In experiments no significant hydrogen-influence on cyclic hardening/softening was found. Therefore the cyclic stabilized stress strain curve of air tested specimens is used due to the fact that more experimental data in air are available. Local stress/strains in the notch root were calculated using finite element method (FEM). Results of FEM Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 371 are compared with classical NEUBER approach [36], as illustrated in Fig. 11. FEM values are interpolated linearly between the calculated breakpoints. In ABAQUS elastic-plastic material behavior with combined (isotropic, kinematic) hardening, cyclic true stress true strain curve, CAX8R elements and von Mises yield criterion were used. Stress/strains obtained from FEM analyses were normal stresses/strains in axial direction.

600 0.7 NEUBER: amplitude

NEUBER: amplitude /% 500 NEUBER: mean stress 0.6 FEM: amplitude

FEM: amplitude local ε K =2.06 FEM: mean stress 0.5 t /MPa 400 RV=0.1 σ Kt=2.06 0.4 300 RV=0.1 0.3 200 0.2 Local stress,Local 100 0.1

0 strain amplitude, Local 0.0 0 100 200 300 0 100 200 300 Nominal stress amplitude, S /MPa Nominal stress amplitude, Sa /MPa a (a) (b)

Fig. 11. Evolution of (a) local stresses and (b) local strains of notched specimens (Kt=2.06) calculated with NEUBER approach and elastic-plastic

FEM analyses. Load ration applied was RV=0.1. Specimens were axially loaded.

In general local stress/strains in the notch root differ significantly from nominal values. Local stress amplitude (σa) and local strain amplitude (εa,t) increases with increasing nominal stress amplitude. The local mean stress (σm) calculated with FEM analysis increases, then decreases and then increases again. Latter is due to high deformation of the notch root at high nominal stress amplitudes, leading to high transverse contraction. Additionally, local stress amplitudes calculated with FEM analyses are increased and local strain amplitudes are lowered compared to NEUBER approach, as a result of modelling the three-dimensional stress state with FEM.

4.3. Damage modeling in air and hydrogen atmosphere

Smith Watson Topper (PSWT) [33], ε-Morrow [37] and Heitmann (PHe) [38] damage models were used for fatigue-life estimations according to local strain approach (see Fig. 10). It should be noted that PHE considers microcrack propagation, modelling the crack closure behavior with an empirical approach. Notch support factor n was calculated according to FKM-guideline [39] (based on the work of Liu [40]) where deformation mechanical part was not included, as suggested by Seeger [41]. The deformation mechanical part is already implemented in the elastic-plastic calculation of local stress/strains in the notch root. The notch support factor involves the statistical size effect (n=1.12). It is proposed to implement n in the elastic component, following Hahn [42]. The damage parameter P and the corresponding P-N curve equations including n are shown in Table 4. The fatigue-life until technical crack initiation is calculated by equating the damage parameter and P-N curve term 5 5 and solving the equation for N. Hydrogen decreases PSWT and PHe significantly in LCF-regime up to N=10 to 2∙10 cycles, as illustrated in Fig. 12. For both, PSWT and PHE, n results in a parallel shift of the curves to higher P values.

372 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

Table 4. Equations of damage parameters (P) and P-N curves for fatigue-life estimation of notched specimens according to local strain approach [32]. Notch support factor n is implemented in the elastic component, following Hahn [42].

Damage model Damage parameter P-N curve

' n VV mf ε-Morrow [37] H N b H ' N)2()2( c ,ta E f

Smith Watson P )( HVV E nP V )2()( 22' b HV '' NEnN )2( cb Topper [33] ,tama f ff

2 ª n V ' º 9.2 ª .3 72 2V a º 5.2 ' b f b 10 ' c Heitmann [38] P 22 HV V f NnP « .0)2( 645 N)2( H f N)2( » « .1 74 » ' aa , pl E ' 2E ¬« R)3( ¼» 1 n ¬« 1 n ¼»

10,000 1,000 Air Air Hydrogen (1MPa) Hydrogen (1MPa) 100 /MPa

1,000 /MPa 10 He SWT P P 1 n n

SWT damage SWT damage parameter, =1.12 =1.12

n=1.0 parameter, damage Heitmann n=1.0 100 0 1.0E+1 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+1 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 Number of cycles, N Number of cycles, N

(a) (b)

Fig. 12. P-N curves of cyclic parameters conducted in air and hydrogen atmosphere. (a) Smith-Watson-Topper damage parameter PSWT and (b) Heitmann damage parameter PHe. The Notch support factor n is implemented in the elastic component, following Hahn [42].

4.4. Fatigue-life estimation of notched specimens

In Fig. 13 results of estimated fatigue-lifes until a technical crack has initiated are compared with experimental determined fatigue-lifes until fracture of notched specimens stressed with a load ratio RV=0.1 in air (a) and hydrogen (b) atmosphere. Fatigue-life estimation using SWT and ε-Morrow approach are conservative, whereas the Heitmann approach is not conservative at lower stress amplitudes, for both testing atmospheres. Moreover, PHe reveals good correlation at higher stress amplitudes with experimental results. The deviation of fatigue-life estimated and experimentally determined in HCF regime is attributed to be the result of conducting strain controlled S-N curves mainly in LCF regime. The legitimacy of comparing the fatigue-life until a technical crack has initiated (estimation with local strain approach) and the fatigue-life until fracture (experiment) is discussed in section 5.1. Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 373

350 350 a a S Air S Hydrogen (1MPa)

300 1.4005IA 300 1.4005IA notched (Kt=2.06) notched (Kt=2.06) 250 RV=0.1 250 RV=0.1 T=20 °C T=20 °C /MPa fracture /MPa fracture 200 200 runout runout x1 PF=50% x1 PF=50% 150 SWT 150 SWT ε-Morrow ε-Morrow Heitmann Heitmann Nominal stress amplitude, amplitude, stress Nominal Nominal stress amplitude, amplitude, stress Nominal 100 100 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+7 1.0E+8 1.0E+2 1.0E+3 1.0E+4 1.0E+5 1.0E+6 1.0E+7 1.0E+8 Number of cycles, N Number of cycles, N

(a) (b)

Fig. 13. Fatigue-life estimations of notched specimens (Kt=2.06) tested in air (a) and hydrogen (b) atmosphere at load ratio Rσ=0.1 and pH2=1 MPa and T=20 °C.

5. Discussion

Several works in the past showed that hydrogen can decrease the fatigue-life in LCF and the fatigue limit in HCF regime, influence slip band movement, increase crack initiation and crack propagation. However, only few studies were carried out to transfer these observations into a fatigue material model to quantify the hydrogen influence and develop a fatigue design approach for structural components cycled in hydrogen atmosphere. Therefore, this work suggests a fatigue design approach based on the local strain approach, modelling the hydrogen influence on strain S-N curve with a reduced plastic strain component in LCF regime.

5.1. Evaluation of fatigue design approach

For the evaluation of the fatigue design approach the ratio of fatigue-life of notched specimens determined in hydrogen (Nhydrogen) and air (Nair) of estimated and experimental conducted fatigue-lifes for nominal stress amplitudes Sa=300, 250, 225 and 200 MPa are compared in Fig. 14. SWT and ε-Morrow approach conservatively model the hydrogen-influence, except at Sa=200 MPa. Heitmann approach fits remarkably well for Sa=300 MPa, but for lower nominal stress amplitudes the hydrogen influence is described non-conservative. It is assumed that this is the result of the overestimated fatigue-life by Heitmann approach at lower Sa in both environments (see Fig 13). Consequently, SWT and ε-Morrow approach should be used for a conservative fatigue-life estimation in hydrogen atmosphere. As already mentioned, notches can influence the effect of hydrogen, e.g. leading to a significant decrease of notched specimens’ UTS [21]. The three dimensional stress state causes high hydrostatic stresses near the notch root, resulting in an increase of hydrogen solubility in the lattice [43] and increasing the detrimental effect of hydrogen. Nonetheless the parameters obtained from unnotched specimens seem to be capable of a fatigue-life estimation of notched specimens. Hence the concept should be applicable to other notch geometries and different fatigue loading types (e.g. R-ratio) as well. This highlights the advantages of the fatigue design concept proposed. Finally, the question arises whether it is meaningful to compare the estimated fatigue-life until a ‘technical crack’ has initiated (crack length 0.5 to 1 mm [34,35]) and the fatigue-life until fracture of notched specimens. The difference between total fatigue-life until fracture and fatigue-life until a ‘technical crack’ hast initiated is assumed to be smaller than factor of two for Nf>100 [44]. In fatigue tests of notched specimens critical crack length from 0.4 to 1.8 mm were observed depending on Sa values (see Table 5). Hence, the difference should be not significant. 374 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

1.4 10 000 % Air , 5 V Hydrogen (1MPa) 1.2 Δ N - 1 000

1.0 MPa Sa=300 , 15% V air,estimated Sa=250 MPa Sa=250 Δ N /N 0.8 MPa Sa=225

= N Sa=200 MPa Sa=200 100 Δ 0.6

0.4 10

Hydrogen,estimated 1.4005IA

Morrow crack cycles Number of of

N unnotched 0.2 SWT propagation, propagation, Rε=-1 Heitmann T=20 °C 0.0 1 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 0.1 0.2 0.3 0.4 0.5 N /N Total strain amplitude, ε /% Hydrogen,experiment air,experiment a,t Fig. 14. Comparison of the ratio of fatigue-life of notched specimens Fig 15. Number of cycles of crack propagation (ΔN) in air and determined in hydrogen (Nhydrogen) and air (Nair) of estimated and hydrogen atmosphere of unnotched specimens at different total strain experimental conducted results. amplitudes εa,t.

5.2. Hydrogen influence on fatigue properties

The present study reveals that the main effect of hydrogen is a decrease of fatigue-life in LCF regime. Different fracture surface topographies observed of air and hydrogen tested specimens indicate that crack propagation is affected in hydrogen atmosphere. In LCF regime at Sa=300 MPa quasi-cleavage and less ductile fracture was observed (see Fig. 7(d)). The fracture surfaces are characterized by flat areas marked by ‘tear lines’ [45] and look strikingly similar to fast fracture surface (FFS) observed by Gamboa and Atrens [45]. FFS was correlated to stress corrosion cracking (SCC) and leading the crack propagating quickly towards the critical growth area. In HCF regime no significant hydrogen influence on fracture surfaces for short cracks was observed, whereas for longer cracks intergranular fracture mode was predominant only in the presence of hydrogen. Ritchie and Suresh T [17] observed an increase of crack growth rate exceeding a critical maximum stress intensity K max=20 MPa√m, leading to intergranular fracture mode. They investigated the crack growth behavior of low strength steel in hydrogen environment with a gas pressure of pH2=0.138 MPa. They found that crack growth in hydrogen environment is influenced by true corrosion fatigue (TCF) for low cyclic stress intensities (ΔK) and stress corrosion T fatigue (SCF) for exceeding K max in hydrogen atmosphere, according to McEvily [46]. TCF decreases ΔKth at Rσ=0.05 and SCF enhances crack propagation up to 20 times. They suggested that the appearance of SCF is dependent of a critical hydrogen concentration ahead of the crack tip. This critical hydrogen concentration depends on hydrostatic stress applied and the hydrogen concentration in the lattice [22, 43]. The latter is related to hydrogen T uptake and diffusivity. Thus an increased hydrogen gas pressure or lower frequency would decrease K max. In this T T study K max can be calculated using the crack length after transition to intergranular cracking a max. According to elastic fracture mechanics, stress intensity (K) is related to crack length as follows:

2 Sa I ),( max S daYaSdaYK ),( Sa , (3) RV )1(

where Y(a,d) is the geometrical factor as a function of a the crack length and d the specimen diameter. The geometrical factor was calculated using the equation for a tension rod with one surface crack proposed by Richard T [48]. According to equation (4) the average value of K max is approximately 11.8 MPa√m, as shown in Table 5. Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 375

T Table 5. Calculated K max and fracture toughness (KIC) of notched 1.4005IA ferritic steel specimens tested in air and hydrogen atmosphere. Notched specimens of ferritic stainless steel 1.4005IA were tested in air and hydrogen atmosphere at Rσ=0.1, pH2=1 MPa and T=20 °C. Air Hydrogen

T T Sa /MPa acrit /mm KIC /MPa√m acrit /mm KIC /MPa√m a max /mm K max /MPa√m 300 0.6 34.5 0.7 38.1 - - 300 0.7 38.1 0.4 27.0 - - 170 1.8 47.1 1.7 44.3 0.23 11.3 170 1.4 36.7 1.6 41.6 0.27 12.3 Average 39.1 37.8 11.8

T In this study lower K max values were obtained compared to the work of Richtie and Suresh (pH2=0.138 MPa), suggested as the result of the higher hydrogen gas pressure (pH2=1 MPa). However, the influence of an increased crack propagation on total fatigue-life should be negligible in HCF regime, whereas a decrease of ΔKth by TCF would lead to a decrease of fatigue limit. Latter was not observed in fatigue tests (see Fig. 13(a), which indicates that TCF did not occur at low nominal stress amplitudes. Concluding that hydrogen does not influence crack initiation at lower nominal stress amplitudes. Fracture surface investigations suggest that hydrogen accelerates crack propagation. In order to compare crack growth rates in air and hydrogen environment the number of cycles to failure of strain controlled fatigue tests are evaluated with different failure conditions. Assuming a crack has already initiated at a drop of maximum stress per cycle by 5% (Δσ5%) from stabilized condition and propagates until a drop by 15% (Δσ15%) a number of cycles of crack propagation ( ΔN) is obtained. In Fig. 15 Δ N is shown for both, air and hydrogen tested specimens at εa,t=0.3 %. In the presence of hydrogen ΔN decreases up to one order of magnitude. Insofar as the reduction of ΔN is the result of a higher crack propagation, hydrogen accelerates FCG. A hydrogen-induced reduction of fracture toughness (KIC) would also lead to lower fatigue-life in LCF regime. Hence, following equation (4) KIC can be calculated using the critical crack length acrit. There is no significant effect on KIC in the presence of hydrogen, as presented in Table 5. The observations above suggest that crack propagation is significantly increased and KIC stays unaffected in hydrogen atmosphere, leading to a reduction of fatigue-life mainly in LCF regime. Additionally, hydrogen assisted cracking (HAC) occurs if a critical hydrogen concentration is exceeded, influenced by applied stresses. Similar observations were also obtained by Gaddam et al. They observed a hydrogen-induced degradation of fatigue-life in LCF regime, no hydrogen-influence in HCF regime and an increase of crack growth rates at cast titanium alloy loaded in high-pressure gaseous hydrogen (10 MPa) [49]. They concluded that the reduction of fatigue-life in LCF regime is due to the increased crack propagation but did not exclude that hydrogen influences crack initiation. It is likely that the detrimental effect of hydrogen on fatigue-life in LCF regime is due to the fact that hydrogen accelerates crack growth but there is no sure evidence yet that hydrogen does not affect crack initiation in LCF regime.

5.3. Hydrogen distribution in the microstructure

The influence of hydrogen on the mechanical properties of steel can be linked to hydrogen diffusion and trapping phenomena. In metals hydrogen is dissolved in the lattice and in reversible and irreversible trap sites, with low binding energy for lattice hydrogen and high binding energy for irreversible trapped hydrogen [50, 51]. TDS spectra of electrochemical hydrogen charged specimens showed several peaks rising with longer charging time. Except peak 1, which saturates after several minutes at a low hydrogen content of approximately 0.07 wppm (see Fig. 4). Additionally, peak 1 diminishes in specimens degassed in a desiccator for one week prior to testing, indicating low binding energy and high mobility. This suggests that peak 1 is linked to the highly mobile hydrogen. Rough estimation using an analytical solution of Fick’s law [52] reveal a hydrogen diffusion coefficient of -7 approximately DH≈10 m²/s. 376 Georg Schauer et al. / Procedia Engineering 133 ( 2015 ) 362 – 378

Higher peak temperatures of peak 2,3 and 4 indicate higher binding energies compared to peak 1. This suggests that these observed peaks can be linked to the trapped hydrogen. Peak 2 represents the greatest amount on total hydrogen content. Microstructural heterogeneities, like MnS [53, 54], carbides [25], grain boundaries [55], dislocations [25], vacancies [56] etc. provide hydrogen trapping sites with their specific binding energy in steels. Unfortunately, results of trapping energies obtained by Kissinger’s method [57, 25] were not plausible and are not shown in this work. It is suggested that the reason are large thermal and diffusion gradients caused by cylindrical specimens’ diameter of 12mm. However, the increased hydrogen content of tensile tested specimens can be attributed to high plastic deformation involving dislocation and void formation and hence providing a higher trap density [58]. The hydrogen content of cyclically loaded specimens stays nearly unaffected regardless of average testing time. On the contrary, a slight decrease of hydrogen content at εa,t=0.3 % can be observed. This suggests that dislocation traps are decreased due to a reduction of dislocation density by distinctive softening behavior of ferritic stainless steel 1.4005IA (see Fig. 6(b)). This leads to the conclusion that a significant amount of hydrogen is trapped on dislocations. Additionally, it is likely that hydrogen is dissolved at grain boundaries, especially since they contain numerous carbide precipitates and considering intergranular fracture at low Sa values. It can be noted that a high amount of hydrogen is trapped in microstructural heterogeneities of ferritic stainless steel 1.4005IA after electrochemical precharging, as well as after testing in hydrogen atmosphere.

5.4. Possible hydrogen mechanisms

Finally, the question arises as to the existence of the predominant hydrogen failure mechanism. For steels most relevant hydrogen failure mechanisms discussed are HEDE and HELP. At HEDE hydrogen reduces the atomic bond strength at sufficiently high hydrogen concentration [2,3]. Intergranular cracking (see Fig. 8(d)) can but need not be an evidence for HEDE [47]. Furthermore, very high hydrogen contents are necessary for HEDE and as far as currently known HEDE is more relevant for high strength steels [47]. HELP implies that solute hydrogen facilitates the movement of dislocations, leading to a localized deformation and reduction of cross-slip [4,5]. But cleavage-like fracture is not explainable by HELP [47]. In contrast to the work of several authors [15, 58-64] a significant hydrogen effect on monotonic and cyclic stress strain curve was not observed, pointing to the fact that interatomic bonds or dislocation motion is not influenced on the macroscopic scale. This suggests, that the presence of hydrogen locally effects material properties, e.g. slip band localization or increased FCG. From observations it appears that the presence of hydrogen gives rise to local decohesion effects rather than dislocation motion-based ones.

6. Conclusion

Stress and strain controlled S-N curves of ferritic strainless steel 1.4005IA are determined in ambient air and gaseous hydrogen (1 MPa) atmosphere and a new fatigue design approach for ferritic stainless steel components loaded in hydrogen atmosphere is proposed following the local strain approach. The conclusions are as follows: (1) The detrimental effect of hydrogen on fatigue-life gets apparent in LCF regime. No significant influence on fatigue limit was observed. In tensile tests only a small decrease of total elongation was found. (2) Fracture surface investigations reveal crack initiation from surface for both, air and hydrogen tested specimens. Fracture surfaces of hydrogen tested specimens show quasi-cleavage at high nominal stress amplitudes. At lower nominal stress amplitudes an increase of intergranular fracture mode in the presence of hydrogen can be observed. (3) The hydrogen-influence on fatigue-life can be modeled with the four-parameter equation according to Manson, Coffin and Morrow by a reduction of plastic strain component in LCF regime. No significant hydrogen-influence on elastic component and cyclic hardening/softening behavior was observed. Georg Schauer et al. / Procedia Engineering 133 ( 2015) 362 – 378 377

(4) Fatigue-life estimations in LCF regime based on the local strain approach show good correlation with experimental results in hydrogen atmosphere. The hydrogen-induced degradation on fatigue-life in LCF regime is modeled mostly conservative and with minor deviation to experiments. (5) It is likely that the effect of hydrogen on fatigue-life in LCF regime is due to hydrogen assisted cracking but there is no evidence that hydrogen does not influence crack initiation at high nominal stress amplitudes.

Acknowledgements

This work was developed in the framework of the German research project “MatFuel” (FKZ: 03ET2051A) founded by the Federal Ministry for Economic Affairs and Energy (BMWi). Furthermore the author thanks DEW GmbH for the provision of material and Dr. Michael Hirscher from Max Planck Institute for Intelligent Systems for the support with TDS analyses. The authors wish to acknowledge the SEM work performed by Andrej Turk.

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