Tensile and fatigue properties of 17-4PH martensitic stainless steels in presence of hydrogen

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Procedia Structural Integrity 19 (2019) 249–258

Fatigue Design 2019 Fatigue Design 2019

Tensile and fatigue properties of 17-4PH martensitic stainless steels Tensile and fatigue properties of 17-4PH martensitic stainless steels in presence of hydrogen in presence of hydrogen a a

Jean-Gabriel SEZGINaa*, Junichiro YAMABEa,b Jean-Gabriel SEZGIN *, Junichiro YAMABEa,b

AIST-Kyushu University Hydrogen Materials Laboratory (HydroMate), National Institute of Advanced Industrial Science and Technology AIST-Kyushu University Hydrogen Materials (HydroMate), National819-0395, Institute ofJapan Advanced Industrial Science and Technology (AIST),Laboratory 744 Motooka, Nishi-ku, Fukuoka b (AIST), 744 Fukuoka Motooka,University, Nishi-ku, Fukuoka 819-0395, Jonan-ku, Japan Departement of Mechanical Engineering, 8-19-1 Nanakuma, Fukuoka, 814-0180, Japan b Departement of Mechanical Engineering, Fukuoka University, 8-19-1 Nanakuma, Jonan-ku, Fukuoka, 814-0180, Japan

Abstract Abstract Effects of hydrogen on slow-strain-rate tensile (SSRT) and fatigue-life properties of 17-4PH H1150 martensitic stainless steel Effects of ultimate hydrogentensile on slow-strain-rate tensile (SSRT) and fatigue-life of 17-4PH H1150 martensitic stainless steel having an strength of ~1GPa were investigated. Smooth properties and circumferentially-notched axisymmetric specimens having an ultimate tensile strength of ~1GPa investigated. Smooth andtests circumferentially-notched axisymmetric were used for the SSRT and fatigue-life tests, were respectively. The fatigue-life were done to investigate the hydrogenspecimens effect on were used for growth the SSRT and fatigue-life respectively. The fatigue-life tests were done to investigate the hydrogen effect on fatigue crack (FCG) properties. tests, The specimens, tested in air at ambient temperature, were precharged by exposure to fatigue crack growth (FCG) properties. The at specimens, airSSRT at ambient temperature, were precharged exposure to hydrogen gas at pressures of 35 and 100 MPa 270°C for tested 200 h. in The properties of the H-charged specimensby were degraded hydrogen gas at pressures of 35 and 100 MPa at 270°C for 200 h. The SSRT properties of the H-charged specimens were degraded by hydrogen, showing a relative reduction in area (RRA) of 0.31, accompanied by mixed fracture surfaces composed of quasi-3 Hz by hydrogen, a relativecracking reduction(IG). in area of 0.31, by mixed surfacesranging composed cleavage (QC)showing and intergranular The(RRA) fatigue-life tests,accompanied conducted under wide fracture test frequencies fromof10quasiHz cleavage and three intergranular (IG). The fatigue-life tests, conducted under widefatigue test frequencies from was 10-3 not to 10 Hz,(QC) revealed distinct cracking characteristics in lowand high-cycle regimes and at the limit. The ranging fatigue limit to 10 Hz, revealed three distinct characteristics in lowhigh-cycle regimes at the fatigue The fatigue degraded by hydrogen. In the high-cycle regime, theand hydrogen caused FCGand acceleration withlimit. an upper boundlimit ratiowas of not 30, degraded by by hydrogen. In theInhigh-cycle regime, thethe hydrogen caused upper boundaccompanied ratio of 30, accompanied QC surfaces. the low-cycle regime, hydrogen causedFCG FCGacceleration accelerationwith with an a ratio of ~100, accompanied QCordinary surfaces.models In the low-cycle regime, the hydrogen FCG acceleration with predicted a ratio of the ~100, accompanied by QC and IG.byThe such as process competition andcaused superposition models hardly H-assisted FCG by QC and IG. The ordinary models such as successfully process competition and superposition models predicted H-assisted FCG acceleration; therefore, an interaction model, reproducing the experimental FCGhardly acceleration, wasthe newly introduced. acceleration; therefore, an interaction model, successfully reproducing the experimental FCG acceleration, was newly introduced. © 2019 The Authors. Published by Elsevier B.V. © 2019 The Authors. Published by Elsevier B.V. © 2019 The Authors. Published by B.V. Peer-review underresponsibility responsibility of Elsevier the Fatigue Design 2019 Organizers. Peer-review under of the Fatigue Design 2019 Organizers. Peer-review under responsibility of the Fatigue Design 2019 Organizers. Keywords: hydrogen embrittlement, precipitation hardened stainless steels, fatigue life test, fatigue crack growth modelling Keywords: hydrogen embrittlement, precipitation hardened stainless steels, fatigue life test, fatigue crack growth modelling

1. Introduction 1. Introduction The recent interest attributed to hydrogen in the energy industry raised several problematics related to exposure of The recent interest attributed to hydrogen in the several of problematics related exposure of material to aggressive environments. In the case of energy metallicindustry alloys, araised degradation the mechanical andtometallurgical material to aggressive environments. In the case of metallic alloys, a degradation of the mechanical and metallurgical properties in presence of hydrogen, also called hydrogen embrittlement (HE), was observed (Brass and Chene 1998; properties in presence of hydrogen, also called hydrogen embrittlement (HE), was observed (Brass and Chene 1998; 2452-3216 © 2019 The Authors. Published by Elsevier B.V. 2452-3216 2019responsibility The Authors. of Published by Elsevier B.V. Organizers. Peer-review©under the Fatigue Design 2019 Peer-review under responsibility of the Fatigue Design 2019 Organizers.

2452-3216 © 2019 The Authors. Published by Elsevier B.V. Peer-review under responsibility of the Fatigue Design 2019 Organizers. 10.1016/j.prostr.2019.12.027

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Gangloff and Somerday 2012; Hirth 1980). Hydrogen-related failure could be attributed to several mechanisms depending on factors such as the source of hydrogen, the metallurgy of the alloy, or the thermomechanical load. The principal mechanisms are the hydrogen-enhanced localized plasticity (HELP) (Birnbaum and Sofronis 1994), the hydrogen-enhanced decohesion (HEDE) (Oriani and Josephic 1974), the adsorption-induced dislocation emission (AIDE) (Lynch 2011), the hydrogen induced cracking (HIC) (Zapffe and Sims 1940), the defactant concept (Kirchheim 2012), and the hydrogen-enhanced strain-induced vacancies (HESIV) (Nagumo 2001). To optimize weight and performance of components, high-strength steels are pertinent. However, when an ultimate tensile strength (UTS) exceeds 1 GPa, the susceptibility of such steels to HE becomes significant (Bandyopadhyay, Kameda, and Mcmahon 1983; Chandler and Walter 1974, 1975; Gangloff 2003; Matsuoka et al. 2017; San Marchi and Somerday 2012; Sandoz 1972; Toplosky and Ritchie 1981; Walter and Chandler 1968; Yamabe et al. 2016). For low and medium-strength steels with the UTS of < 900 MPa, there exists an upper bound for the fatigue crack growth (FCG) acceleration ratio of ~30 (Matsuoka et al. 2017). In this case, the steels failed by quasi-cleavage (QC) and the hydrogen-induced successive crack growth (HISCG) model justified such a bounded acceleration. In contrast, SAE52100 (JIS-SUJ2) with the UTS of 1900 MPa presented intergranular (IG) facets without any upper bound for the FCG acceleration ratio (Yamabe et al. 2012). In this case, the IG failure was justified by the hydrogen-enhanced deformation twin model. These facts sustained an interest for the H-assisted FCG acceleration and related fracture surface morphology for steels with the UTS of around 1 GPa. The FCG rate is a parameter of great interest and some models have been proposed in the literature (Wei and Gangloff 1989). However, the effects of hydrogen on the FCG properties are multiple and have not been fully understood yet, as explained in (Wei and Simmons 1973; Nanninga et al. 2010). Some complex hydrogen effects on the FCG properties exist and therefore, the interested reader is invited to refer to comprehensive review such as (Nanninga et al. 2010; Petit et al. 1994). The fundamental mechanisms of the H-assisted FCG acceleration being not elucidated, the present paper adopts a phenomenological approach to model and predict the H-assisted FCG acceleration ratio. Literature exists about modelling aspects of such acceleration based on multiple mechanisms. The common models are the superposition model (Chen and Wei 1998; Landes and Wei 1969; Wei 2002; Wei and Gao 1983) and the process competition model proposed by Austen et al.(Amaro et al. 2014; Austen and Mcintyre 2014). The validity of these models in regard with the present context is discussed hereafter. The objective of this study was to investigate the effects of hydrogen on 17-4PH H1150 steel with the UTS nearly equal to 1 GPa under both monotonic and cyclic loading and characterize the frequency dependence of the H-assisted FCG acceleration. An interaction model considering both cycle- and time-dependent mechanisms was newly introduced and provided results in coherence with the experimental results. Nomenclature f R 𝜎𝜎𝑎𝑎 ( (

𝑑𝑑𝑑𝑑

testing frequency stress ratio stress amplitude )

𝑑𝑑𝑑𝑑 𝑖𝑖 𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝑖𝑖

fatigue crack growth (FCG) rate resulting in 𝑖𝑖 type failure

fatigue crack growth (FCG) acceleration ratio resulting in 𝑖𝑖 type failure

2. Material characterization 2.1. Metallurgical aspects The alloy in interest was the 17-4PH H1150 stainless steel, composed of 0.04 C, 0.31 S, 0.87 Mn, 0.034 P, 0.004 S, 3.3 Cu, 4.24 Ni, 15.57 Cr, and 0.34 Nb (in mass %); the remainder was iron. This grade was obtained by application of a solution treatment at 1040°C during 1 h (water quenching) prior a precipitation hardening treatment at 620°C during 2 h and air cooling. Observations by means of scanning electron microscopy (SEM) / electron backscatter diffraction (EBSD) showed a martensitic microstructure with a residual austenite content lower than 4%.



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2.2. Samples and experimental protocol Slow-strain-rate tensile (SSRT) and fatigue-life tests were performed on non-charged and H-charged specimens to quantify the effects of hydrogen on tensile and FCG properties of this alloy. The SSST tests were carried out on smooth axisymmetric specimens with 7.5 mm in diameter at ambient temperature and a displacement rate of 1 mm.min-1. Fatigue-life tests were carried out on circumferentially-notched specimens. The gross-section diameter of the specimen was 12.8 mm. The net-section one was 6.4 mm and the notch radius was 0.083 mm. The corresponding stress concentration factor was then 6.6. This geometry respects with the prescriptions of the ASTM G142-98 standard (ASTM 2004). The tests of the non-charged and H-charged specimens were performed in air at ambient temperature under various stress amplitudes and test frequencies. The stress ratio, R, was equal to -1 and the test frequencies were included between 0.001 Hz and 10 Hz. The H-charged specimens were charged by exposure to gaseous hydrogen at pressures of 35 MPa and 100 MPa at 270°C for 200 h prior testing. According to calculations using the solution of a diffusion equation, these charging conditions result in a uniform distribution of dissolved hydrogen through the specimens. In order to quantify a potential outgassing of hydrogen in solution during testing, the hydrogen content was measured on 12.8-mm-diameter and 5mm-thick cylindrical specimens withdrawn from the fatigue-life specimens after testing (at 5 mm from the failure surface). The hydrogen content was determined by thermal desorption analysis (TDA) at a heating rate of 100°C.h-1. 3. Results and discussion 3.1. Tensile properties The results of three conditions were gathered in Figure 1-a). The nominal stress-strain curves suggested that nosignificant effect on neither the elastic domain nor the UTS was observable. The UTS was indeed equal to 1016 MPa (35 MPa H-charged), 1026 MPa (100 MPa H-charged), and 1013 MPa (non-charged). In contrast, the tensile ductility was significantly degraded by the presence of hydrogen. Namely, the relative reduction of area (RRA), which translates the ductility loss of the material in presence of hydrogen (ANSI 2014), was 0.43 (35 MPa H-charged) and 0.31 (100 MPa H-charged). Figure 1-b) shows the fracture surface morphology of the non-charged specimen, whereas Figure 1-c) presents the one of the H-charged specimens. The failure mode characterized by a cup-and-cone failure in the non-charged specimen was affected by hydrogen, resulting in brittle failure consisting in a mixture of quasicleavage (QC) and intergranular cracking (IG).

Figure 1 – Effect of hydrogen on the SSRT property: a) nominal stress-strain curves related to non- and H-charged specimens, b) fractographic observation of the non-charged specimen, c) fracture surface observed on the H-charged specimen

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3.2. Hydrogen content The hydrogen content was measured after exposing H-charged specimens to air at ambient temperature for durations from 24 h up to 358 h. The resulting hydrogen contents were nearly constant in the whole duration range. In the case of 100 MPa H-charged specimens, the average content was 16.5 mass ppm ± 1.9 mass ppm (based on 27 measures). In the case of 35 MPa H-charged specimens, the average content was 9.6 mass ppm ± 0.5 mass ppm (based on 3 measures). All the fatigue-life tests were completed within 358 h; therefore, the effect of outgassing from the Hcharged specimen during the test was concluded to be negligible. 3.3. Experimental evidence of fatigue-life properties in presence of hydrogen The results of fatigue-life tests were provided by Figure 2 for different test frequencies and stress amplitudes. On the S-N diagram, the data related to the non-charged and H-charged specimens were respectively plotted in blue and red. A notable H-induced degradation of the fatigue life was observed in the low-cycle regime, especially at lower test frequencies. Oppositely, a moderate H-induced degradation of the fatigue life was observable in the high-cycle regime. Additionally, no H-induced degradation of the fatigue limit was observed. The fatigue limit is the maximal stress amplitude resulting to non-failed specimens at 1 × 107 cycles. In the present case, the fatigue limits of non-charged and H-charged specimens were equal to 110 MPa. Figure 3Figure shows fatigue cracks emanating from the notch root observed in the non-charged and H-charged specimens tested at the stress amplitude corresponding to the fatigue limit (σa = 110 MPa). These cracks were not detected in the non-failed specimens tested at stress amplitudes of 105 MPa and 100 MPa. Although we do not have a lot of experimental evidences, these cracks are considered to be nonpropagating cracks. A series of the experimental results suggest that the fatigue limit of the circumferentially-notched specimen was not degraded by hydrogen.

Figure 2 – S-N diagram for the circumferentially-notched specimens tested under a stress amplitude in the [100 MPa; 400 MPa] range, a stress ratio of -1 and frequencies from 10-3 Hz up to 10 Hz. Non-charged and H-charged specimens were tested in air at room temperature.



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Figure 3 – Non-propagating cracks observed at a stress amplitude of σa = 110 MPa on non-charged (a) and H-charged (b) specimens

The H-induced degradation of the fatigue life was then discussed. As shown in Figure 2, the H-induced degradation of the fatigue life was notable in the low-cycle regime and moderate in the high-cycle regime. According to previous research on a low-alloy steel with the UTS of ~950 MP by (Yamabe et al. 2015), the fatigue-life of circumferentiallynotched specimen having the same geometry could be predicted from the FCG properties. In the present case, the fatigue-life was also predictable from the FCG properties. Therefore, the H-induced degradation of the fatigue-life revealed in the present study was attributed to the H-assisted FCG acceleration. Hence, experimental results emphasized the presence of two different effects of hydrogen on the FCG acceleration occurring in the high- and lowcycle regimes. To clarify the effects of hydrogen in the different regimes, the fracture surfaces were analysed. Figure 4 shows fracture surface morphologies of the non- and H-charged specimens in the high- and low-cycle regimes obtained by SEM. The SEM micrographs (a) and (c), classified in the high-cycle regime, were obtained at a stress amplitude of 120 MPa and a test frequency of 10 Hz. The micrographs (b) and (d), classified in the low-cycle regime, were taken at a stress amplitude of 400 MPa and a test frequency of 1 Hz. The micrographs (a) and (b) were related to the non-charged specimens and the micrographs (c) and (d) to the H-charged specimens. The non-charged specimens presented striation or microstructure-dependent surfaces whereas the H-charged ones presented some brittle fracture surfaces. At low-stress amplitude levels in the high-cycle regime (c), the specimen failed by QC, although at high-stress amplitude levels in the low-cycle regime (d), the specimen failed by a mixture of QC and IG. The fractographic observations then suggested that different mechanisms of the H-assisted FCG acceleration took place depending on the testing conditions. The introduction of hydrogen led to QC surfaces in the whole testing range. However, the high-stress amplitude in the low-cycle regime promoted the occurrence of another mechanism leading to the IG cracking.

Figure 4 – Fracture surfaces observed on non-charged (a,b) and H-charged (c,d) specimens at: a low-stress level in the high-cycle regime, 𝝈𝝈𝒂𝒂 =120 MPa, 𝒇𝒇=10 Hz (a,c) and a high-stress level in the low-cycle regime, 𝝈𝝈𝒂𝒂 =400 MPa, 𝒇𝒇=1 Hz (b,d)

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3.4. Test frequency-dependent failure mechanism It has been established that the fatigue-life of the present specimen is closely related to the FCG properties. Thus, the H-assisted FCG acceleration ratio, defined as the ratio of fatigue-life of non-charged to H-charged specimen Nair/NH2, was analysed. Figure 5 shows the test-frequency dependence of the H-assisted FCG acceleration ratio measured for stress amplitudes between 150 MPa and 400 MPa. For all the testing conditions, the FCG acceleration ratio appeared to be bounded. The literature suggested that in case of the HISCG model, an upper bound equal to 30 exists for the BCC structure, as illustrated by the example of low-alloy steels plotted with a dashed line (Matsuoka et al. 2017). In addition, the fracture surface within the framework of the HISCG model was covered with only QC. In the present case, in high-cycle regime, the upper bound of the H-assisted FCG acceleration ratio was around 30 and the fracture surface was covered with QC; thus, the observed FCG acceleration could be justified by the HISCG model. However, in low-cycle regime (e.g. stress amplitude of 300 MPa), the upper bound exceeded the value of 30, up to around 100 and, additionally, the fracture surface was covered with QC and IG. This fact translated that the IG fracture provoked an FCG acceleration produced by a mechanism different from the HISCG model, leading to an increased value of the upper bound. The graph below clearly shows an increase of the FCG acceleration ratio with the stress amplitude and the decrease of the testing frequency. For this reason, it was relevant to quantify the amount of IG facets covering the fracture surface in order to verify the occurrence of this additional failure mechanism.

Figure 5 – Test-frequency dependence of the H-assisted FCG acceleration ratio illustrated for four stress amplitudes. The acceleration ratio appears to be bounded at different levels depending on the stress amplitude. The data related to SCM435 have been taken from (Matsuoka et al. 2017)

Percentage of IG surfaces was measured on the dataset corresponding to a stress amplitude of 300 MPa, the SEM observations being available in a wide range of test frequencies. These measurements were conducted by observing fracture surface located at 1 mm from the notch root by SEM. Figure 6 illustrates the measurement results by highlighting in red the IG facets in the 10 Hz (a) and the 10-3 Hz (b) cases. The measurements were realized in an area judged representative of the fracture mode and extended enough to limit localized measurements. The micrographs showed that the fraction of IG surface was 3.6% at 10 Hz (a) and 15.3% at 10-3 Hz (b). The results of the measurements in the whole frequency range were represented in Figure 7. This figure clearly shows a correlation between the percentage of IG facets covering the fracture surface and the testing frequency. The data were then fitted and the results were given in the frame and represented by the dashed line. 3.5. Prediction of the test frequency-dependent fatigue crack growth rate The experimental facts suggested that in the H1150 steel, the underlying mechanism of H-assisted FCG acceleration presented some similitude with the HISCG mechanism. The HISCG mechanism has been identified in low-alloy steels with the UTS of ≤ 900 MPa (e.g. JIS-SCM435, JIS-SCM439) (Matsuoka et al. 2017). However, the experimental facts have concluded that at the low-cycle regime, the upper bound of the FCG acceleration ratio exceeded 30 and IG facets were observed on the fracture surface. In this case, the fracture surface has somehow



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presented some analogies with JIS-SUJ2, whose failure mechanism is time-dependent, resulting in both an unbounded FCG acceleration ratio and an IG fracture. These points imply that the H-assisted FCG acceleration may be out of scope of the HISCG mechanism. For these reasons, the assumption of a failure mechanism issued from the combination of two elementary mechanisms was investigated. A combination rule for the HISCG model (cycledependent and bounded at 30) and a time-dependent model (unbounded) was established in regard with experimental measurements. Figure 8 illustrates the combination of the mechanisms using the fatigue life (a), the translated FCG rate (b), and the FCG acceleration ratio (c). The data related to the cycle-dependent mechanism were extrapolated from the literature of JIS-SCM435 (Matsuoka et al. 2017) as represented in Figure 5. The value at a frequency of 10 Hz being unavailable, the value of FCG acceleration ratio for H1150 at a stress amplitude of 200 MPa were considered since the IG failure was not significantly observed for this condition. In the case of a purely time-dependent mechanism, the FCG rate and the FCG acceleration ratio can be expressed as (with the nomenclature introduced in Figure 8):

Figure 6 –Proportion of IG surface highlighted in red given for two frequencies at a stress amplitude of 300 MPa: a) f=10 Hz and b) f=10-3 Hz

Figure 7 – Percentage of IG failure measured on the basis of fractographic analysis for five frequencies (fitted by the expression in the frame)

(

(

𝑑𝑑𝑑𝑑

)

𝑑𝑑𝑑𝑑 𝐼𝐼𝐼𝐼 𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

=(

𝐼𝐼𝐼𝐼

𝑑𝑑𝑑𝑑

=(

1 𝑑𝑑𝑑𝑑

(1)

) + ( )

𝑑𝑑𝑑𝑑 𝐹𝐹 𝑑𝑑𝑑𝑑

𝑓𝑓

) /(

𝑑𝑑𝑑𝑑 𝐼𝐼𝐼𝐼

𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑

(2)

)

𝑑𝑑𝑑𝑑 𝐹𝐹

In case of cycle-dependent mechanism, these expressions of the FCG rate and FCG acceleration ratio are expressed by (3) and (4). 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 (3) ( ) = ( ) +( ) (

𝑑𝑑𝑑𝑑 𝑄𝑄𝑄𝑄 𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝑄𝑄𝑄𝑄

𝑑𝑑𝑑𝑑 𝐹𝐹 𝑑𝑑𝑑𝑑

=(

)

𝑑𝑑𝑑𝑑 𝑄𝑄𝑄𝑄

𝑑𝑑𝑑𝑑 𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻 𝑑𝑑𝑑𝑑

/(

(4)

)

𝑑𝑑𝑑𝑑 𝐹𝐹

From the expression (1) and (2), the FCG acceleration ratio related to time-dependent mechanism was expressed in (5) as a function of the testing frequency. ln ((

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝐼𝐼𝐼𝐼

𝑑𝑑𝑑𝑑

𝑑𝑑𝑑𝑑

− 1) = − ln 𝑓𝑓 + ln ( ) − ln ( ) = − ln 𝑓𝑓 + ln 𝑓𝑓0 𝑑𝑑𝑑𝑑

𝑑𝑑𝑑𝑑 𝐹𝐹

(5)

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Considering the values of the FCG acceleration ratio in the [10 -3 Hz; 10 Hz] frequency range, it was assumed that Nair/NH >> 1, which led to the expression (6). The relevance of this assumption was then retrospectively verified. 𝑁𝑁 (6) ln ( 𝑎𝑎𝑎𝑎𝑎𝑎) = − ln 𝑓𝑓 + ln 𝑓𝑓0 𝑁𝑁𝐻𝐻

𝐼𝐼𝐼𝐼

As observed on Figure 7, the IG fraction depended on the test frequency. By assuming that an IG fraction of zero translated the non-occurrence of the time-dependent mechanism, the value of 𝑓𝑓0 was deduced from the results of the fitting in Figure 7 and this value was found to be equal to 134 Hz. In the literature, some models exist to take into account such a combined mechanism: the superposition model (Chen and Wei 1998; Landes and Wei 1969; Wei 2002; Wei and Gao 1983) and the process competition model (Amaro et al. 2014; Austen and Mcintyre 2014). In the present study, the mechanisms were combined in regard with the fraction of IG fraction noted hereafter α. The superposition model relies on a linear combination of expressions (1) and (3) translating the time dependent and cycle dependent mechanisms respectively. As a result, the FCG rate and FCG acceleration ratio are expressed by (7) and (8) (with the nomenclature introduced in Figure 8). 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 𝑑𝑑𝑑𝑑 𝛼𝛼 𝑑𝑑𝑑𝑑 = (1 − 𝛼𝛼) ( ) + 𝛼𝛼 ( ) = ( ) + (1 − 𝛼𝛼) ( ) + ( ) (7) ( ) (

𝑑𝑑𝑑𝑑 𝑄𝑄𝑄𝑄+𝐼𝐼𝐼𝐼 𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝑑𝑑𝑑𝑑 𝑄𝑄𝑄𝑄

𝑄𝑄𝑄𝑄+𝐼𝐼𝐼𝐼

= {(1 − 𝛼𝛼) (

𝑑𝑑𝑑𝑑

)

𝑑𝑑𝑑𝑑 𝑄𝑄𝑄𝑄

𝑑𝑑𝑑𝑑 𝐼𝐼𝐼𝐼

+ 𝛼𝛼 (

𝑑𝑑𝑑𝑑

𝑑𝑑𝑑𝑑 𝐹𝐹

) }/ (

𝑑𝑑𝑑𝑑 𝐼𝐼𝐼𝐼

𝑑𝑑𝑑𝑑

𝑑𝑑𝑑𝑑 𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻𝐻

) = (1 − 𝛼𝛼) (

𝑑𝑑𝑑𝑑 𝐹𝐹

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝑄𝑄𝑄𝑄

𝑓𝑓

𝑑𝑑𝑑𝑑

+ 𝛼𝛼 (

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝐼𝐼𝐼𝐼

(8)

The process competition model consists in combining the time and cycle dependent mechanisms, (2) and (4) respectively, by considering a pondered harmonic combination (9). (

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 −1 𝑁𝑁𝐻𝐻

)

𝑄𝑄𝑄𝑄+𝐼𝐼𝐼𝐼

=

(1−𝛼𝛼)

𝑁𝑁 ( 𝑎𝑎𝑎𝑎𝑎𝑎) 𝑁𝑁𝐻𝐻 𝑄𝑄𝑄𝑄

+

𝛼𝛼

𝑁𝑁 ( 𝑎𝑎𝑎𝑎𝑎𝑎) 𝑁𝑁𝐻𝐻 𝐼𝐼𝐼𝐼

(9)

Figure 8 – Decomposition of the FCG acceleration ratio into two mechanisms in terms showed on different diagrams. The S-N diagram (a) shows the different regimes. The da/dN–ΔK diagram (b), obtained by rotation of (a), shows the mechanisms. Effect of test frequency on the H-enhanced FCG acceleration (c) shows the combination of the mechanisms of cycle- and time-dependence properties.

Figure 9-a) compares the results obtained by the superposition and the process competition models to the experimental values. In the graph, the elementary cycle- and time-dependent mechanisms were respectively represented by a dashed line with diamond shaped markers and dotted line with cross shaped markers. The results of the superposition and process competition models were respectively marked with green triangles and black squares linked by solid lines and the experimental results with blue circles. The graph showed that neither the superposition model nor the process competition model firmly supported the decomposition of the FCG acceleration ratio into two elementary mechanisms. The first one led to an unbounded behaviour, whereas the later one led to the FCG acceleration ratio almost equal to the cycle-dependent mechanism (the HISCG mechanism). These facts suggest that the cycle- and time-dependent crack growths do not occur individually but interact with each other since the present H-assisted FCG acceleration was accompanied by QC and IG. Taking into account these situations, the third model was then proposed to predict more accurately the experimental results. Considering the elementary mechanisms expressed in (2) and (4), the proposed model, hereafter called interaction model, is expressed in (10). (

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝑁𝑁𝐻𝐻

)

𝑄𝑄𝑄𝑄+𝐼𝐼𝐼𝐼

=(

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 1−𝛼𝛼 𝑁𝑁𝐻𝐻

)

𝑄𝑄𝑄𝑄

(

𝑁𝑁𝑎𝑎𝑎𝑎𝑎𝑎 𝛼𝛼 𝑁𝑁𝐻𝐻

)

𝐼𝐼𝐺𝐺

(10)



Jean-Gabriel SEZGIN et al. / Procedia Structural Integrity 19 (2019) 249–258 Jean-Gabriel Sezgin et al./ Structural Integrity Procedia 00 (2019) 000–000

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The FCG acceleration ratio calculated by the interaction model was low compared to the superposition model but larger than the ratio issued from the process competition model. The results issued from the interaction model are given on Figure 9-b) and show that not only the interaction model didn’t mimic any of the elementary mechanisms, but also successfully reproduced the test frequency dependence of the FCG acceleration ratio of the H1150 steel.

Figure 9 – Prediction of test-frequency dependence on the hydrogen-enhanced FCG acceleration ratio from elementary mechanisms by the superposition and competition models (a) and interaction model (b)

4. Conclusions This paper investigated the effects of hydrogen on tensile and fatigue crack growth (FCG) properties of the 174PH H1150 steel. Smooth and circumferentially-notched axisymmetric specimens were used for slow-strain-rate tensile (SSRT) and fatigue-life tests, respectively. The specimens were precharged by exposure to hydrogen gas at pressures of 35 and 100 MPa at 270°C for 200 h. The notched fatigue test was done to determine the FCG properties. Several testing conditions were considered: the stress amplitude was in the [120 MPa; 400 MPa] range at stress ratio of -1 and the test frequency in the [10-3 Hz; 10 Hz] range. These investigations have led to the following conclusions: 1. The SSRT tests showed no degradation of an ultimate tensile strength (UTS). The relative reduction in area (RRA) was 0.31 for 100 MPa H-charged specimens and the fractographic analysis showed a mixture of QC and IG surfaces, suggesting a high susceptibility to hydrogen embrittlement (HE). 2. In the fatigue-life test, no degradation of the fatigue limit was observed. Fatigue cracks emanating from notch roots, considered to be non-propagating cracks, were observed in both the non- and H-charged specimens. 3. At low stress amplitudes in the low-cycle regime, the H-assisted FCG acceleration ratio showed an upper bound of ~30 and the H-charged specimens failed accompanied by quasi-cleavage (QC), in accordance with the hydrogen-induced successive crack growth (HISCG) mechanism. In contrast, at high stress amplitudes in the low-cycle regime, the upper bound of the FCG acceleration ratio was nearly equal to 100 and the Hcharged specimens failed by a mixture of QC and intergranular cracking (IG). Additionally, the proportion of IG facets were quantified and correlated to the test frequency. 4. The H-assisted FCG acceleration ratio was decomposed into cycle-dependent and time-dependent mechanisms. On the hypothesis that the cycle-dependent mechanism follows the HISCG one, the superposition and process competition models available in the literature did not provide satisfactory results. An interaction model, in good agreement with the experimental FCG acceleration ratio, was then proposed. References Amaro, Robert L., Neha Rustagi, Kip O. Findley, Elizabeth S. Drexler, and Andrew J. Slifka. 2014. “Modeling the Fatigue Crack Growth of X100 Pipeline Steel in Gaseous Hydrogen Q.” International Journal of Fatigue 59:262–71. ANSI. 2014. “ANSI/CSA CHMC 1 - 2014: Test Method for Evaluating Material Compatibility in Compressed Hydrogen Applications - Phase I Metals.” CSA Group. ASTM. 2004. “G142-98" Standard Test Method for Determination of Susceptibility of Metals to Embrittlement in Hydrogen Containing Environments at High Pressure, High Temperature, or Both.” ASTM International, West Conshohocken, PA 4.

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