Unravelling the effect of hydrogen on microstructure evolution under low-cycle fatigue in a high-manganese austenitic TWIP steel

Journal Pre-proof Unravelling the effect of hydrogen on microstructure evolution under low-cycle fatigue in a high-manganese austenitic TWIP steel Dayong An, Waldemar Krieger, Stefan Zaefferer PII:

S0749-6419(19)30674-6

DOI:

https://doi.org/10.1016/j.ijplas.2019.11.004

Reference:

INTPLA 2625

To appear in:

International Journal of Plasticity

Received Date: 13 September 2019 Revised Date:

20 November 2019

Accepted Date: 20 November 2019

Please cite this article as: An, D., Krieger, W., Zaefferer, S., Unravelling the effect of hydrogen on microstructure evolution under low-cycle fatigue in a high-manganese austenitic TWIP steel, International Journal of Plasticity (2019), doi: https://doi.org/10.1016/j.ijplas.2019.11.004. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier Ltd.

Unravelling the effect of hydrogen on microstructure evolution under low-cycle fatigue in a highmanganese austenitic TWIP steel Dayong An*,Waldemar Krieger, Stefan Zaefferer* Max-Planck-Institut für Eisenforschung, Max-Planck-Str. 1, 40237 Düsseldorf, Germany * Corresponding authors: [email protected], [email protected]

Abstract We systematically investigated the effect of hydrogen on the low-cycle fatigue (LCF) behaviour of a Fe-28Mn-0.3C (wt.%) twinning-induced plasticity (TWIP) steel by comparing the fatigue microstructure of samples with and without hydrogen pre-charging using electron channelling contrast imaging (ECCI) and cross-correlation electron backscatter diffraction (CC-EBSD). The results reveal that a complex interplay of several hydrogen-microstructure interaction mechanisms is involved in the observed LCF behaviour of the studied TWIP steel with hydrogen pre-charging; first, hydrogen assists the nucleation of stacking faults (SFs) and deformationinduced ε-martensite in the studied TWIP steel by reduction of local stacking fault energy (Suzuki effect) and stabilization of the hexagonal close-packed (hcp) ε-phase. The evolution of fatigue dislocation pattern is strongly retarded in the presence of hydrogen. The rapid formation of ε-martensite leads to a stronger cyclic hardening. Meanwhile, the impingement of ε-martensite plates at GBs causes high local stress concentration, which plays a dominant role in the observed fatigue cracking. The crack opening angle indicates that the exact associated hydrogen embrittlement mechanism is the hydrogen-enhanced decohesion (HEDE) mechanism instead of

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the hydrogen-enhanced localized plasticity (HELP) mechanism. Among different GB types, annealing twin boundary shows the highest immunity to hydrogen-related fatigue cracking. Key words: TWIP; Low cycle fatigue; Deformation martensite; Hydrogen embrittlement; HEDE

1. Introduction High manganese steels (HMnSs), e.g. steels with twinning-induced plasticity (TWIP) (Bouaziz et al., 2011) and with transformation-induced plasticity (TRIP) (Fischer et al., 2000), exhibit an excellent combination of strength and ductility, which makes these steels candidate weightreduced materials for applications in automotive industry. Nevertheless, these advanced high strength materials suffer premature or unexpected catastrophic failure when exposed to hydrogen, known as hydrogen embrittlement (HE) (Koyama et al., 2012; Liu et al., 2017; Ryu et al., 2013; So et al., 2009). As many of the applications involve cyclic deformation (Hariharan et al., 2011), fatigue failure is also an important damage mechanism for engineering structures. The situations become even worse when cyclic loading acts in association with hydrogen, which was reported to enhance the fatigue crack growth rate (Wan et al., 2019a). Hence, understanding the effect of hydrogen on fatigue behaviour of HMnSs is of great importance to accurately predict the component lifetime and to design reliable advanced materials. A number of mechanisms of HE have been proposed, most prominently hydrogen-enhanced decohesion (HEDE) (Oriani, 1970; Oriani and Josephic, 1974; Troiano, 1960) and hydrogenenhanced localized plasticity (HELP) (Beachem, 1972; Birnbaum and Sofronis, 1994; Robertson and Birnbaum, 1986). The HEDE mechanism claims that the accumulated hydrogen at crack fronts reduces the interatomic cohesive forces and thus leads to crack propagation at a lower stress level than in a low-hydrogen situation. The HELP mechanism proposes that the hydrogen

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atmospheres surrounding the dislocation can shield their elastic stress fields and reduce the interactions between dislocations and obstacles, e.g. precipitates and solute atoms. Therefore, the hydrogen atmospheres can enhance the dislocation mobility and promote localized plastic deformation. Apart from the HEDE and HELP mechanisms, hydrogen was also reported to reduce the stacking fault energy (SFE) of materials (Hermida and Roviglione, 1998; Rozenak and Ellerzer, 1987) and promote the formation of deformation-induced twinning in TWIP steels (Astafurova et al., 2010; Kireeva et al., 2011) and of ε-martensite in austenitic stainless steel (Narita et al., 1982; Rozenak and Ellerzer, 1987). Hydrogen-related fracture in HMnSs has been reported in many studies (Chin et al., 2011; Koyama et al., 2017; Koyama et al., 2012; Koyama et al., 2013; Laureys et al., 2015; Ryu et al., 2013). Results show that the HE susceptibility of HMnSs can be affected by various factors, e.g. texture (Chun et al., 2012b), hydrogen concentration (Wang et al., 2007), residual stress (Chin et al., 2011; Chun et al., 2012b) and phase stability (Koyama et al., 2016; Li et al., 2017). However, most of the studies done so far have been focused on HE of HMnS under monotonic deformation; in contrast, our knowledge on the HE susceptibility of HMnSs under cyclic loading is incomplete. Note that the fatigue damage derives from the accumulation of cyclic slip irreversibilities (Mughrabi, 2009), e.g. formation of complex dislocation patterns, which is determined by dislocation interaction (Li et al., 2011). As proposed by the HELP mechanism (Birnbaum and Sofronis, 1994), hydrogen may strongly influence the interaction between dislocations. Consequently, the evolution of dislocation patterns under cyclic loading may be affected greatly by the presence of hydrogen. Furthermore, the effects of residual stresses and phase stability on HE susceptibility under cyclic loading were also examined in this study. It was reported that the interaction of dislocations with grain boundaries (GBs) plays an important role in the HE of the

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α-iron (Wan et al., 2019b). The HE-resistance of GBs with different misorientation interacted with dislocation/ε-martensite was also investigated in the studied austenitic steel. Note that the role of twin boundaries (TBs) on fatigue crack resistance is still under debate. For example, TBs are considered to be strongly resistant against cracking in high-SFE materials, while they become preferred crack initiation sites in low-SFE materials (Qu et al., 2008; Zhang et al., 2012). Furthermore, the fatigue cracking susceptibility of TBs is also dependent on the loading direction (Li et al., 2014). Since annealing TBs are particularly common in face-centred cubic (fcc) alloys with low SFE (Jin et al., 2013), it is thus important to study the fatigue behaviour of TBs in the presence of hydrogen. The aim of this study is to shed light on the micro-mechanisms responsible for the observed LCF behaviour of a particular high-Mn TWIP steel under the presence of hydrogen. As shown in a previous study(An and Zaefferer, 2019) , dislocation patterns of a high-manganese TRIP steel have already developed significantly after only 50 cycles at a strain amplitude of around 0.5%. Therefore, low-cycle fatigue (LCF) was performed in this study. The fatigued microstructures of samples with and without hydrogen were compared by electron channelling contrast imaging (ECCI). Compared to transmission electron microscopy (TEM), less artefacts and much larger observation area can be obtained by ECCI (Zaefferer and Elhami, 2014). ECCI has been proved to be successful in observation of fatigue microstructures(An and Zaefferer, 2019). Thus, it is a powerful and well-suited technique to study the effect of hydrogen on fatigue microstructure in a detailed and statistically meaningful way. Electron backscatter diffraction (EBSD) was applied to measure the crystal orientation and phase composition of the deformed microstructures. Furthermore, cross-correlation based EBSD (CC-EBSD) with high angular resolution (Wilkinson et al., 2006) was applied to measure the residual strain/stress distribution after cyclic loading.

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2. Experimental procedures The chemical composition of the TWIP steel investigated in this study is Fe- 27.84 wt.% Mn0.28 wt.% C. The material was strip cast and subsequently homogenized at 1150°C for 2 h. Next, the cast strip was cold rolled from 1.6 mm to 0.8 mm thickness. Thereafter, a recrystallization annealing was performed at 900°C for 20 min, followed by air cooling to room temperature. After heat treatment, the samples showed a fully austenitic (fcc) structure with an average grain size of around 17.8 µm. The SFE of this material, evaluated by a thermodynamic approach (Saeed-Akbari et al., 2009), is around 27 mJ/m2. Deformation twins are clearly observed in a grain under 5% tensile strain, as shown in Fig. 1 (a). To avoid the bulge effect of flat bone-shaped specimen under tensile-compression cyclic loading due to the small thickness, shear specimens with a geometry as shown in Fig. 1 (b1) were prepared by electric discharge machining (EDM). They can be used directly for cyclic shear testing in uniaxial testing machines (ASTM, 2005; Yin et al., 2014). The specimens were polished with SiC grinding papers from 220 grit to 1000 grit, followed by polishing with 1µm diamond suspension. High surface quality was then obtained by polishing with 50 nm colloidal suspension of SiO2 (OPS) for 20 minutes. After sample preparation the thickness of the specimen was around 0.7 mm. Hydrogen charging was conducted electrochemically at ambient temperature (25°C) before cyclic loading in a 0.05 M H2SO4 aqueous solution containing 1.4g/L CH4N2S as hydrogen recombination poison at a cathodic density of 1 mA/cm2 for 120 hours. A platinum plate with dimensions of 25×25×0.1 mm3 was used as the counter electrode. After hydrogen-charging, the specimens were slightly polished by OPS (less than 1 min) to maintain the high surface quality. Note that the charged sample still showed a fully austenitic structure and no surface defects were introduced by the charging process. The hydrogen contents of the

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samples at different conditions, i.e. without hydrogen charging, immediately after hydrogen charging, and after 3 hours under ambient conditions after hydrogen charging, were measured by thermal desorption spectroscopy (TDS). TDS was performed from room temperature to 800°C at a constant heating rate of 26 K min-1. The amount of hydrogen that can diffuse at room temperature is defined as diffusible hydrogen. This includes as well the hydrogen trapped at dislocations (So et al., 2009) and GBs (Du et al., 2011), which plays an important role in HE (Koyama et al., 2013). The diffusible hydrogen content was determined in this study by the cumulatively desorbed hydrogen from room temperature to 300°C. Cyclic shear fatigue tests were conducted with a tensile-compression instrument (Kammrath & Weiss GmbH, Dortmund, Germany) under displacement control (schematically shown in Fig. 1 (b2)) at a speed of 2 µm/s. Cyclic shear fatigue tests of pre-charged samples were performed immediately after charging (within 30 min) for 5 and 50 cycles, respectively. The shear fatigue tests were finished within 3 hours after hydrogen charging for both samples and the deformed microstructures were observed immediately after fatigue tests. For comparison, similar cyclic tests (5 cycles and 50 cycles) were repeated on samples without pre-charged hydrogen. Digital image correlation (DIC) was utilized during the fatigue tests to determine the local strain amplitudes. Fig. 1 (b3) shows one DIC snapshot of the major strain distribution after one quarter cycle. EBSD (Fig. 1 (b4)) and ECCI (Fig. 1 (b5)) were performed on the region of interest (highlighted by the dashed rectangle in Fig. 1(b3)) to measure the crystallographic orientation and dislocation structures, respectively. This integrated experimental approach enables analysing the effect of hydrogen on the fatigue behaviour of the TWIP steel with detailed loading information at a larger field of view and with less artefacts compared to TEM.

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Besides EBSD, ECCI and DIC, microstructural observations were furthermore performed by secondary electron (SE) imaging, CC-EBSD, and X-ray diffraction (XRD). EBSD measurements were done using an EDAX/TSL Hikari camera with the TSL OIM data collection 7 software in Zeiss XB1540, with a scan step size of 0.1-0.5 µm at 15kV acceleration voltage. SE imaging and ECCI were carried out using a Zeiss Merlin SEM with an electron beam accelerating voltage of 30 kV and a working distance of 17 and 7 mm, respectively. XRD measurements were performed on samples with and without hydrogen pre-charging after 50 cycles at 40 kV/30 mA, an angular step size of 0.03° and count time of 20 s per step with a Co target. A point focus beam with a diameter of 2 mm, which is equal to the gauge length of the shear sample, was applied on the deformed region of the shear sample. CC-EBSD analysis was performed using the crosscourt 4.0 software with EBSD patterns collected in a JEOL JSM 6500F, with a scan step size of 0.1 µm at an electron beam accelerating voltage of 15kV. Detailed information on this technique can be found elsewhere (Wilkinson et al., 2006).

3. Experimental results 3.1 Hydrogen distribution after cathodic charging Fig. 2 shows the hydrogen desorption rate against the temperature. Compared to the uncharged sample (0.3 wt. ppm), about 11.7 wt. ppm of diffusible hydrogen had been charged into the sample after 120 hours of charging. After an exposition for 3 hours to ambient environment conditions, the diffusible hydrogen content does not change much. The fatigue tests were finished within 3 hours after charging. Note that the samples used for TDS measurements were un-deformed. In contrast, cyclic loading introduces dislocations and ε-martensite plates, which are important hydrogen trapping sites (Chun et al., 2012a; Krieger et al., 2018). It is thus

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reasonable to assume that during the fatigue tests the total hydrogen content within the samples decreases even less than in the here performed TDS tests. Note that cathodic charging was reported to generate a hydrogen concentration gradient from surface to centre (Luo et al., 2018; Ronevich et al., 2012): the specimen surface contains a markedly higher amount of hydrogen than that of the specimen interior. To estimate the depth of the hydrogen-affected zone, the fracture surface of a similar cathodic-charged tensile specimen was observed, as shown in Fig. 3. The fracture surface in the interior of the specimen shows typical ductile fracture features, i.e. it is dominated by dimples, e.g. Fig. 3 (c1). In contrast, the fracture surface near the charging surface presents brittle fracture features, i.e. intergranular cracks, as shown in Fig. 3 (c3). In addition, the region bound by the blue and red dashed lines in Fig. 3 (b) displays a transition state, i.e. both intergranular cracking and ductile features coexist, as displayed in the Fig. 3 (c2). The depth of the hydrogen-affected regions, including the brittle cracking and the transition regions, is estimated to be around 100 µm. Alternatively, the diffusion depth, X, can also be roughly estimated based on the equation

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(Gottstein, 2013),

where D is the diffusion rate and t is the charging time. Taken D as 10-11 cm2/s (in austenitic steels (Kanezaki et al., 2008)) and t as 120 h, diffusion depth of 51 µm is obtained, which is in the same order of magnitude as the experimentally obtained one. Note that the thickness of the tensile sample after fracture is around 0.5 mm. Assuming that the hydrogen diffusion depth is similar from both sides, the hydrogen-effected region is thus around 40% (0.1 mm×2/0.5mm) of the whole sample. Based on TDS results shown in Fig. 2 (a) (11.7 wt. ppm), the average diffusible hydrogen content of the hydrogen-affected zone can thus be calibrated to 29.3 wt. ppm. As the observation depth of ECCI and EBSD are less than 100 nm (Zaefferer and Elhami, 2014),

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the fatigue microstructures investigated here are mainly within the sample surface. The hydrogen content of the H-charged sample surface (estimated by Fick’s law) is around 60 wt. ppm.

3.2 Mechanical response with and without hydrogen pre-charging The mechanical responses of the samples with and without hydrogen pre-charging under displacement-controlled shear fatigue are shown in Fig. 4 (a) in terms of maximum stresses versus cycle numbers. Both samples show rapid cyclic hardening during the first 10 cycles, followed by an almost stable response up to 50 cycles. However, the hydrogen-charged sample shows a higher hardening rate and a stronger hardening response (145 MPa) than the noncharged sample (130MPa). The phase composition of the samples after 50 loading cycles was identified by XRD and EBSD, as shown in Fig. 4 (b) and (c). A fully austenitic structure was maintained up to 50 cycles in the sample without hydrogen; in contrast, ε-martensite was detected in the hydrogen pre-charged sample both by XRD and EBSD. Note that due to the resolution limitation of EBSD, some of the fine martensite plates (<~100 nm) are not probed. Furthermore, the effect of hydrogen on the mechanical properties of the studied TWIP material under monotonic deformation was also analysed in this study, as shown in supplementary Fig. 1. Detailed information can be found in the appendix.

3.3 Dislocation patterns with and without the presence of hydrogen Fig. 5 compares the fatigue microstructures of the hydrogen-charged and non-charged samples after 5 loading cycles. Grains with similar loading conditions (local strain amplitude and crystallographic orientation) were selected from the two fatigued samples. In the non-charged sample (Fig. 5 (a1) and (a2)), perfect dislocations are clearly observed. Furthermore, interactions of dislocations from different slip systems lead to the formation of dislocation tangles. The fatigued microstructures of the hydrogen-charged samples are quite different from that of the

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non-charged one. Surprisingly, ε-martensite was already found after 5 cycles even at a relatively low local strain amplitude (0.4-0.6%), as shown in Fig. 5 (b1) and (b2). Furthermore, individual, piled-up (on the same slip plane) and stacked (on neighbouring slip planes) stacking faults (SFs) with dissociation widths varying from hundreds of nanometres to several micrometres were frequently observed, as highlighted by white arrows in Fig. 5 (b1) and (b2). It should be mentioned that, in principle, SF and ε-martensite can be distinguished by contrast difference. As reported by Zaefferer et al. (Zaefferer and Elhami, 2014), the contrast of SFs consists an intense bright line on one side and fades with intensity oscillations, one ideal example is displayed in the white dashed rectangle. In contrast, the contrast of ε-martensite is more intense and homogeneous; it appears as a homogeneous bright band as shown in the red dashed rectangle. Fig. 6 compares the fatigue microstructure of the hydrogen-charged and non-charged samples after 50 loading cycles. The most distinct difference between the two samples is the formation of deformation ε-martensite plates in the hydrogen-charged samples while such structures were never observed in the non-charged sample under the same cyclic loading conditions. For the noncharged samples, after 50 loading cycles dislocations develop into complex dislocation patterns (Fig. 6 (a1)-(a3)), e.g. dislocation veins and dislocation walls. In contrast, in the hydrogencharged sample low-density dislocation structures, e.g. planar arrays and dislocation tangles, develop and remain confined inside the blocks divided by ε-martensite plates. To quantitatively estimate the effect of hydrogen on fatigue dislocation structure evolution, the dislocation densities of charged and non-charged samples were measured from ECC images. For example, the dislocation density of the walls is around 1×1015 m-2 in the non-charged sample (inlet in Fig. 6 (a1)), which is 5 to 10 times higher than that in grains of the hydrogen-charged sample (inlet in

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Fig. 6 (b1)) under the same loading condition. Note that more than 300 grains were observed for each sample. Similar differences also hold for the other grains. Fig. 6 (b) also reveals important details of the interactions between dislocations, SFs and martensite plates in the hydrogen-charged sample. First, it is visible in Fig. 6 (b2) and (b3) that not only parallel ε-martensite variants but also different intersecting variants are activated. Local contrast differences near variant intersections indicate residual strain/stress concentration, e.g. those marked by red arrows in Fig. 6 (b3). A zoom-in into this area, marked by a red rectangle, reveals a very high density of dislocations with distances in the order of < 10 nm running on the same planes as the ε-martensite habit plane and piling up at a variant intersection. Similar morphology is also visible inside the orange-marked rectangle in Fig. 6 (b2). Here dislocations pile up to a distance of about 15 nm in front of a martensite plate. Finally, from the enlarged section of the deformation martensite plate (marked by a green rectangle in Fig. 6 (b3)) it is clearly visible that the wide ε-plate consists of many thin ε-plates separated by residual γ phase. Fig.7 (a)-(c) show an ECC image, a phase map and a kernel average misorientation (KAM) map of the same grain in the hydrogen-charged sample loaded for 50 cycles at a local strain amplitude of 0.7%. Deformation ε-martensite plates in the grain are found to be arrested by an annealing twin boundary (top) and a different ε-martensite variant (bottom), as shown in Fig. 7 (a). The KAM map (calculated from first nearest neighbours) reveals large residual strain concentrations at these impingement sites. In contrast, the strain concentration at the visible triple junction, marked by a red arrow, shows much lower values. CC-EBSD was performed to obtain more detailed information on the stress and strain fields; an example is given in Fig. 7 (d)-(f). The cross-correlation reference points (selected manually in the defect-free regions according to the ECC image) are marked by red dots. The normal stresses σ11 and σ22, displayed in Fig.7 (e) and

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(f), respectively, show the highest values. Assuming that the stress states of the reference points are zero, the positive (red) and the negative (blue) values thus correspond to tensile and compressive stresses, respectively. At the GB with impinging ε-plates an oscillation of stress values between positive and negative values is observed for both normal stresses, σ11 and σ22, as marked by red and blue arrows in Fig. 7 (e) and (f). Note that the epsilon martensite plates here are not properly probed owing to limitation of spatial resolution, which causes severe artefacts in the residual stress measurement.

3.4 Hydrogen-induced intergranular cracks Neither ε-martensite plates nor intergranular cracks were ever observed, even in regions with a relatively high local strain amplitude (1.2-1.8%), in the non-charged sample after 50 cycles. In contrast, nucleation of cracks is already detected after 5 cycles in the sample with hydrogen precharging, as marked by red arrows in Fig. 8 (a) and (b). Typical crack initiation sites are those areas where deformation ε-plates imping on grain boundaries. Surprisingly, no crack nucleation was observed at triple junctions except of those where an impingement of ε-martensite plates was involved. Furthermore, cracks were found to be always initiated at one corner of the ε-GBintersection line, as visible, e.g. with the crack nucleating at the upper part of the intersection in Fig. 8 (a). After 50 cycles, crack nucleation sites were also observed in regions with relatively low local strain amplitude, e.g. Fig.8 (c) (0.7 %) and (d) (0.9%). Fig. 9 (a) shows an example of intergranular crack propagation. The red and green arrows show the nucleation sites and the growth directions, respectively. As mentioned before, the crack always nucleates on one corner of the intersection line and propagates from the nucleation corner along the GB-ε intersection line, as shown in the enlarged view. In this view it is also visible that the crack opening forms a particular angle between the ε-plate and the impinged GB. Fig.9 (b)

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shows crack propagation in a region with a larger local strain amplitude (1.4%). ECCI and phase map of the selected grain are overlaid. The GBs are highlighted by white dashed lines. Generally, the crack propagates along the GB between grain A and B, while short deviations following other crystallographic features, like martensite boundaries or slip bands, are also observed, as shown in details in Fig. 9 (bi) and (bii). This creates a ‘zig-zag’ morphology. As displayed in Fig. 9 (bi), these martensites/slip bands cause surface relief on the hydrogen-charged sample, which are marked as extrusions and intrusions. Based on misorientation relationship, GBs are divided into four groups, i.e. low angle GBs (LAGBs) (misorientation angle below 15°), low-Σ (3< Σ≤ 29) GBs based on the coincidence site lattice theory (Grimmer et al., 1974), TBs (coherent Σ3) and all others as random high angle GBs (HAGBs). From an EBSD overview scan with a size of 400×600 µm2, the fractions of all the four groups were calculated using the OIM software with a critical deviation angle based on Brandon’s criterion (Brandon, 1966). The results are shown by the white columns in Fig. 10 (a). According to the EBSD and ECCI results, no signal of deformation twins was found. Therefore, all the coherent Σ3 GBs are considered to be annealing TBs. To study the hydrogen-assisted intergranular crack susceptibility of different GB types, the misorientation of all the GBs with fatigue cracks (27 GBs in the hydrogen-charged sample after 50 cycles) were determined by correlating ECC/SE images with EBSD data. The proportions of GBs with crack(s) are present by grey columns in Fig. 10 (a). Most of the failed GBs belong to the random HAGBs. Despite the large fraction of annealing TBs (34% of all GBs), no intergranular cracks were found along these GBs. Among the other low-Σ GBs, crack initiations were observed on one Σ17b±3.1° GB and one Σ25b±1.5°. The fraction of GBs with cracks and their corresponding misorientation angles is displayed in Fig. 10 (b). Although no cracks were found to be initiated from LAGB, a

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certain fraction of the misorientation angles of the cracked GBs fall within 15° to 20°, e.g. the upper GB in Fig. 8 (d) (15.4°). Finally, the distribution of loading angles (illustrated in Fig. 10(c)) of cracked GBs is presented in Fig. 10 (d). It is found to be relatively random.

4. Discussion 4.1 Hydrogen-assisted stacking fault and ε-martensite formation The studied material clearly shows twinning-induced plasticity (TWIP, SFE: ~27 mJ/m2) in the absence of hydrogen, as shown in Fig. 1(a). In contrast, after hydrogen charging the material frequently forms deformation-induced ε-martensite plates (no sign of deformation twin was detected) and a high density of SFs with large dissociation widths under cyclic loading, as shown in Figs. 4-6; thus the material changes from a typical TWIP behaviour to a TRIP (transformation-induced plasticity) behaviour. This change of deformation mechanisms suggests a reduction of the SFE in the presence of hydrogen (Luo et al., 2017; Pierce et al., 2014; Rozenak and Ellerzer, 1987). Note that hydrogen is not equally distributed in the material but may localize significantly, particularly at defects, e.g., GBs (especially those with more open structures, e.g. random HAGB (Di Stefano et al., 2015; Du et al., 2011)), dislocations(Zhu et al., 2017), SFs (Suzuki effect(Suzuki, 1962a)) and vacancies(Li et al., 2015), which may change the solubility or/and mobility of hydrogen (Du et al., 2011; Krieger et al., 2018; Lee and Jang, 2007). Especially, the role of the Suzuki effect should be taken into consideration (Yamada et al., 2015), which may explain the formation of widely extended SFs, as schematically illustrated in Fig. 11 (a)-(c); segregation of hydrogen to SFs reduces the local SFE and therefore leads to an increase of the dissociation width of the bounding partials. Provided that the incorporation of hydrogen into SFs reduces the energy of the hydrogen solutes enough to overcompensate the stacking fault energy, the SFs continuously extend and attract more hydrogen as illustrated in Fig.11 (c). 14

A hydrogen-induced martensitic γ-ε transformation has been frequently reported to proceed in metastable austenitic stainless steels (Narita et al., 1982; Rozenak and Ellerzer, 1987; Ulmer and Altstetter, 1993), which suggests that with increasing hydrogen content the hexagonal closepacked (hcp) phase becomes thermodynamically more stable. Furthermore, Hermida el at. (Hermida and Roviglione, 1998) pointed out that the H-H pairs in the SFs stabilize the hcp phase. Nishino el at. (Nishino et al., 1990) examined the hydrogen-induced phase transformation in FeMn-Ni alloys and found that a high manganese composition promotes the formation of εmartensite instead of hydride. The formation of hydrogen-assisted SFs is well explained above and fits well with the experimental results. The formation of ε-martensite in austenite is closely related to the formation of SFs for HMnS (Nishiyama, 1978) and may be achieved by generation of SFs on every second layer of {111}γ lattice planes. As reported by Fujita el at. (Fujita and Ueda, 1972) and further supported by recent publications (Lai et al., 2018; Lu et al., 2018; Wang et al., 2018), SFs first accumulate irregularly; subsequent SFs tend to arrange on every second slip plane to minimize both the bulk free energy and the total SF energy, leading to the formation of ε-martensite nuclei. The here observed morphology of the formed ε-martensite plates, i.e. thin ε-plates separated by the residual γ phase as displayed in green rectangle in Fig. 6 (b3), supports this irregular accumulation process for the present case. Note that in a previously studied TRIP steel (An and Zaefferer, 2019) submitted to cyclic loading , ε-martensite was rarely observed even after 100 cycles with a strain amplitude of around 0.8%. In contrast, ε-martensite plates were already frequently observed in the hydrogen-charged TWIP sample after only 5 cycles with a strain amplitude of around 0.2%. This indicates that the critical resolved shear stress (CRSS) for ε-martensite nucleation is much lower in the hydrogen-charged sample. One possible explanation could be that the presence of hydrogen surrounding the gliding partial dislocations

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reduces the passing stresses when they pass each other on co-planar {111} planes. Therefore, the friction stress resisting the formation of a high density of stacked SFs in local regions as well as the CRSS for ε-martensite nucleation is significantly reduced.

4.2 Effect of hydrogen on fatigue microstructure evolution The high density of piled-up dislocations, e.g. in the enlarged areas in Fig. 6 (b2) and (b3), can be explained by the HELP mechanism(Birnbaum and Sofronis, 1994): the well-known screening of the dislocation stress field by the presence of hydrogen leads to the fact that dislocations in pile-ups move closer together (Gottstein, 2013). Furthermore, the surrounding hydrogen also reduces the attractive force between opposite dislocations, which may thus hinder the formation of dipole structures. Therefore, it can be deduced that cyclic slip reversibility is enhanced and the evolution of fatigue dislocation patterns proceeds more slowly in the hydrogen-charged samples. As a result, low-density dislocation structures are formed in the hydrogen-charged samples compared to densely-arranged dislocation wall structures found in grains with similar cyclic loading conditions but in absence of hydrogen, as seen in Figs. 5 and 6. Note that the formation of deformation martensite accommodates a certain amount of the imposed plastic strain. At the same time the residual stresses caused by the impinging ε-martensite plates act as strong back stresses for the imposed cyclic load, i.e. they show a stress shielding effect. This shielding effect would, consequently, retard the evolution of the dislocation structure. With the present data it is, however, impossible to decide whether the reduced dislocation density in the hydrogen-charged sample is due to the stress and strain shielding by ε-martensite or due to the stress shielding effect of hydrogen in the dislocation cores (HELP) which would lead to lower attractive forces for dipole formation. Most probably both effects are involved.

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Rapid cyclic hardening is observed for the first 10 loading cycles for both samples, with and without pre-charged hydrogen as shown in Fig. 4, however, the maximum cyclic stress of the plateau stage increases by around 10% in the presence of hydrogen (from ~130 MPa to ~145MPa). This increase of the cyclic hardening rate can be attributed to the formation of hydrogen-assisted ε-martensite, which causes microstructure refinement and decrease of the dislocation mean free path, i.e. it causes a dynamic Hall-Petch effect (Bouaziz et al., 2008; Dancette et al., 2012).

4.3 Hydrogen-induced fatigue cracking Severe HE is observed for the studied hydrogen pre-charged material; nucleation of fatigue cracking is already observed after 5 cycles. This process is due to the interaction of several individual effects of hydrogen in the material. Based on our ECCI and CC-EBSD results, a potential collection and order of mechanisms are schematically shown in Fig. 11. As described in section 4.1 and illustrated in Fig. 11 (a)-(d), hydrogen promotes the formation of SFs and εmartensite plates by means of reduction of SFE (most probably in terms of the Suzuki effect (Suzuki, 1962b; Yamada et al., 2015)) and the stabilization of the hcp ε-phase. Furthermore, the impingement of deformation martensite plates on GBs imposes a large local shear strain onto the impinged GB, which further renders the boundary locally more susceptible for hydrogen accumulation so that hydrogen attains concentrations at the impinged GBs that are several orders of magnitude larger than the normal concentration in equilibrium (Oriani, 1970; Oriani and Josephic, 1974; Wilcox and Smith, 1965). Hydrogen accumulation may be further promoted by the cyclic movement of dislocations and gliding partials, through which amounts of hydrogen are collected from the interior matrix and are transported to the boundary as it has been proposed by Novak et al (Novak et al., 2010). Consequently, the cohesive forces at these locations are

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significantly weakened, as proposed by the HEDE mechanism (Oriani, 1970; Oriani and Josephic, 1974; Troiano, 1960). Note that triple junctions, which have been reported to be preferred hydrogen-related crack initiation sites under tensile loading (Koyama et al., 2013), are not involved in crack formation in this study. Instead, crack initiation was always associated with ε-martensite plates. This phenomenon can be explained by the difference in residual strain/stress concentration between triple junctions and martensite-GB impingement locations, as clearly visible in Fig. 7. For deformation ε-martensite, the participating partials always have the same Burgers vector, as illustrated in the enlarged orange rectangle in Fig. 9 (a). As a result, the corresponding shear strain γε is around 0.35 (Lai et al., 2018), which is much larger than the local strain amplitude applied in this study (<0.02). Therefore, the residual stress generated by the impingement of deformation martensite on GBs plays a dominant role in the observed fatigue cracking. The presence of high tensile and compressive residual stresses is probed clearly by the CC-EBSD measurements in Fig.7 (e) and (f). According to the energy theory of Griffith (Griffith and Eng, 1921), the relevant stress for fraction initiation corresponds to the applied normal stress on the crack plane (Rice, 1968; Sih, 1991). Crack nucleation and propagation are promoted by tensile stresses at GBs, while compressive stresses suppress crack formation, as illustrated in Fig. 11(e). Another important feature of the hydrogen-related fatigue cracking is the crack opening angle between the original boundary and the impinged martensite plate, which corresponds, as displayed in the enlarged part of Fig. 9 (a), almost exactly to the shear strain of deformation martensite. This reveals the fact that the impinged GBs where cracks nucleate are completely brittle, without any plastic deformation. Note that since many HE mechanisms, e.g. HELP and HEDE, always lead to similar fracture morphologies (Robertson et al., 2015), it is challenging to

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clarify the exact associated mechanisms. However, in this case we can confidently prove the effect of the HEDE mechanism and exclude that of the HELP mechanism on the nucleation of hydrogen-related fatigue cracking. From Fig. 10 it becomes obvious, that random HAGBs are the most susceptible to crack nucleation while coherent TBs (and probably LAGBs, but the statistical evidence is low) are completely inert. This is in contrast to literature where TBs were reported to be susceptible to crack initiation in hydrogen-charged nickel-based alloys under monotonic deformation (Seita et al., 2015). Although these authors claimed that this observation is not fully understood, they ascribed it mainly to dislocation-mediated GB sliding on the coherent TB planes, which would modify the TB structure by incorporation of dislocations and, consequently, would enhance the solubility for hydrogen in the TB. Our present results indicate that the main reason for fatigue cracking, as aforementioned, is the martensite-induced stress concentration and the associated HEDE mechanism. Due to their close-packed structure (An et al., 2018), the coherent TBs hardly provide hydrogen trapping sites and may even act as diffusion barrier for hydrogen (Di Stefano et al., 2015; Du et al., 2011). In contrast, the open structure units in random HAGBs provide hydrogen trapping sites as well as hydrogen diffusion pathways (Du et al., 2011). Furthermore, in (Seita et al., 2015) the susceptibility to GB cracking is strongly dependent on the loading direction, while in this study a relatively randomly distribution of the loading angles of the cracked GBs is observed (Fig. 10 (d)). This can also be understood by the fact that the stress concentration generated by cyclic loading is much lower than that generated by deformation martensite plates. As indicated by Figs. 8 and 9 (a) cracks nucleate particularly at the largest ε-martensite impingements. At a higher strain amplitude or/and with further cyclic loading, separated

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microcracks may link up to form main cracks, which run generally along complete GBs, as schematically illustrated in Fig. 11 (f). However, short deviations from GBs, e.g. Fig. 9 (bi) and (bii), are also observed. The reason for this has not been fully clarified by our observations yet. Note that extrusions and intrusions are reported to be preferred fatigue crack initiation sites (Mughrabi, 2009). But most probably the observed zig-zag growth is due to the fact that the hydrogen content is locally higher along these features (as discussed in section 4.1) and therefore local embrittlement occurs.

5. Conclusion The effect of hydrogen on low-cycle fatigue behaviour has been investigated in a Fe-28Mn-0.3C (wt. %) steel with twinning-induced plasticity (TWIP). The fatigue behaviour of the studied material are strongly affected by hydrogen, which can be explained by the complex cooperation of several mechanisms known for interaction of hydrogen with metals, including HELP, HEDE, and reduction of SFE. Based on the experimental results, the following detailed conclusions can be drawn: 1) Hydrogen assists the formation of stacking fault (SF) and ε-martensite by means of segregation at SFs (the Suzuki effect) and the stabilization of the hcp ε-phase (reduction of SFE) in the fatigued TWIP samples. 2) Rapid cyclic hardening occurs during the initial cycles in both charged and non-charged samples. The dynamic Hall-Petch effect, caused by formation of ε-martensite, leads to stronger cyclic hardening in the hydrogen-charged sample. 3) Grains in the hydrogen-charged sample are divided into isolated blocks by deformationinduced ε-martensite plates. Fatigue dislocation patterns develop and remain confined in

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these blocks. Compared to the non-charged sample, the evolution of the fatigue dislocation patterns is strongly diminished in the hydrogen-charged sample. 4) The impingement of ε-martensite plates on GBs leads to high local stress concentrations, which change across the plate from positive to negative values. These stresses play a dominant role on the observed fatigue cracking. The fact that the crack opening angle corresponds almost exactly to the shear strain of the deformation martensite suggests that the cracking occurs completely brittle due to a strong HEDE mechanism. 5) Among different GB types, annealing twin boundaries (TBs) (and probably low-angle GBs) show highest immunity to fatigue cracking caused by the impingement of martensite at GBs in the hydrogen-charged sample. The reason for this is probably the high atomic packing density of the coherent TBs which does not allow higher hydrogen concentration to form. The current work, thus, indicates that grain boundary engineering, which increases the fraction of annealing TBs, could be a potential mitigation method for HE in the fatigued TWIP steel.

Acknowledgement The authors thank Benjamin Breitbach for performing XRD measurements. D. An and S. Zaefferer acknowledge funding of this research by the German Research Foundation [Deutsche Forschungsgemeinschaft (DFG)] through SFB 761 “Steel ab Initio”.

Appendix: Deformation behaviour of the hydrogen-charged TWIP steel under monotonic deformation An overview of the HE of the studied TWIP steel under monotonic deformation is present in supplementary Fig. 1. Surprisingly, both the ductility (from 78% to 64%) and the ultimate

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tensile strength (UTS) (from 750 MPa to 675 MPa) decrease in the presence of hydrogen. As highlighted by red arrows in supplementary Fig. 1(b2), faultage zones of DIC speckle pattern were frequently observed during tensile test in the hydrogen-charged sample, which are caused by the formation of macro-cracks (supplementary Fig. 1(c2)), i.e. surface failure. In contrast, no such phenomenon was observed in the non-charged tensile sample (supplementary Fig. 1(b1) and (c1)). Therefore, compared to the non-charged sample, the surface layer of the hydrogencharged sample is much more fragile, which leads to the reduction of the cross section. The observed ‘softening’ behaviour (10 % UTS reduction) is most probably due to the hydrogenrelated surface failure. By comparing the mechanical response of the hydrogen-charged TWIP steel under monotonic and cyclic deformation, it can be seen that hydrogen could show opposite effects on the mechanical response with different loading conditions. This phenomenon can be explained by the fact that several hydrogen-microstructure interaction mechanisms can exist simultaneously while the dominant one varies with different loading conditions (also varies with different hydrogen concentration and distribution). In detail, under cyclic loading the damaging effect of HEDE mechanism is weak due to the low applied strain amplitude, while the dynamic HallPetch effect caused by formation of ε-martensite makes an important contribution on the mechanical response. In contrast, under monotonic deformation severe surface failure (the HEDE effect) occurs with increasing loading strain, which leads to a strong softening/damaging effect that overweighs the hardening effect discussed above.

Declarations of interest: none.

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Fig. 1 (a) Deformation microstructure of the studied TWIP steel under tensile loading with global strain of around 5 %. From top to bottom: LD-IPF map, IQ map and phase map of the same area. (b) Overview of the experimental procedures. (b1) Sample geometry; (b2) Schematic of the displacement controlled symmetric cyclic loading; (b3) Digital image correlation (DIC) snapshot showing the local strain distribution of the sample at ¼ cycle. (b4) and (b5) are the EBSD scan displayed in LD-IPF and ECCI overview of the deformation region of interest marked by dashed line in (b3). (TWIP: Twinning induced plasticity, LD: Loading direction, IPF: Inverse Pole Figure, IQ: Image Quality, γ: Austenite, ε: Epsilon martensite).

Fig. 2 Hydrogen desorption rate curve of the samples at different sample conditions: black line: directly after hydrogen charging, red line: after hydrogen charging and keeping at ambient environment for 3 hours, pink line: uncharged. HD indicates the diffusible hydrogen content.

Fig. 3 (a) Fracture surface overview of the tensile specimen with hydrogen pre-charging; (b) Magnified image of the region highlighted by black dashed rectangle in (a). Three different fracture morphologies ((c1)-(c3)) are characterized and separated by the red and blue dashed lines in (b): (c1) ductile fracture surface, (c2) transition stage between ductile and intergranular fracture surface and (c3) intergranular cracking surface.

Fig. 4 (a) Global cyclic deformation response of the studied TWIP sample with (red) and without (black) hydrogen pre-charging. (b) XRD profiles of the fatigued TWIP sample with (red) and (without) hydrogen; (c) Phase map overlaid with IQ of the fatigued TWIP sample: (c1) without hydrogen pre-charging, (c2) with hydrogen pre-charging.

Fig. 5 Comparisons of the fatigued microstructures of the un-charged sample ((a1)-(a2)) and the hydrogen-charged sample ((b1)-(b2)) after 5 cycles at local strain amplitude of ~0.4-0.6%. (SD: Shear direction, SF: Stacking Fault, DT: Dislocation Tangle).

Fig. 6 Comparisons of the fatigued microstructures between the un-charged sample ((a1)-(a3)) and hydrogen-charged sample ((b1)-(b3)) after 50 cycles at local strain amplitude of ~0.6-0.8%. The blue rectangles in (a1) and (b1) illustrate the process of dislocation density determination from ECCI: the intersections of individual dislocations with the surface are marked by red dots. The dislocation density is obtained by counting the total number of the red dots and dividing it by the area of blue rectangle. (PA: Planar Array, DW: Dislocation Wall).

Fig. 7 (a) ECCI, (b) phase map and (c) KAM map based on conventional EBSD of a selected area in the hydrogen-charged sample after 50 cycles at a local strain amplitude of 0.7%; (d) ECCI, (e, f) normal stress distribution based on CC-EBSD of another selected area in the hydrogen-charged sample after 50cycles at a local strain amplitude of 0.7%. Note in the CCEBSD example, the ε-martensite cannot be recognized due to the limitation of spatial resolution (KAM: Kernel Average Misorientation, TJ: Triple Junction).

Fig. 8 ECC images showing intergranular crack initiation in the hydrogen-charged sample after (a, b) 5 cycles and (c, d) 50 cycles. Red arrows highlight the fatigue crack nucleation sites.

Fig. 9 (a) ECC images showing intergranular crack propagation in the hydrogen-charged sample after 50 cycles at a local strain amplitude of 0.9%. The red and green arrows in the magnified region highlight the nucleation and propagation direction of the intergranular cracks, respectively; (b) ECCI overlaid with phase map showing the intergranular crack passing through the whole grain boundary after 50 cycles at a strain amplitude of 1.4%. Crack deviations are displayed in the magnified images (i) and (ii).

Fig. 10 Statistical analysis of GB cracking in the hydrogen-charged sample after 50 cycles (a) GB fractions of four different GB types; white bars: including all GBs, grey bars: including particular GBs with crack(s); (b) Misorientation angle distribution frequency of GBs with cracks; (c) Example illustrating the definition of loading angle; (d) Loading angle distribution frequency of the GBs with cracks. (HAGB: High Angle Grain Boundary, LAGB: Low Angle Grain Boundary).

Fig.11 Schematic describing the fatigue microstructure evolution in the H-charged TWIP sample and the involved hydrogen-microstructure interaction mechanisms. (HEDE: hydrogen-enhanced decohesion).

Supplementary Fig. 1 Overview of hydrogen embrittlement of the studied TWIP sample under tensile deformation (initial strain rate 10-3s-1): (a) Engineering stress-strain curves of the studied TWIP sample with (red) and without (black) hydrogen pre-charging. (b) Snapshots of tensile sample with DIC speckle pattern under 20%, 40% and 50% global engineering strain, respectively; (b1) without hydrogen, (b2) with hydrogen. The red arrows highlight the macro cracks (surface failure in the hydrogen-charged sample) that destroy the speckle patterns. (c) Comparison of fracture surface of tensile sample without hydrogen pre-charging (c1) and that of the sample with hydrogen pre-charging (c2). (TD: tensile direction).

Highlights: •

Fatigue microstructures of high-manganese TWIP samples with and without hydrogen are characterized and compared.

•

Effects of hydrogen on the evolution of dislocation pattern under low-cycle fatigue are discussed.

•

The exact associated mechanisms of hydrogen-related fatigue cracking are clarified.

•

Distribution of residual stress generated by the impingement of ε-martensite at grain boundaries is probed and its role on hydrogen embrittlement susceptibility is examined.

Author Statement Dayong An: Conceptualization, Methodology, Investigation, Writing-Original Draft, Writing-Review&Editing. Waldemar Krieger: Methodology. Stefan Zaefferer: Conceptualization, Methodology, Validation, Investigation, WritingReview&Editing, Supervision, Project administration

Declarations of interest: none