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Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel
Accepted Manuscript Title: Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel Authors: Xiaogang Li, Baoming Gong, Caiyan Deng, Yizhe Li PII: DOI: Reference:
S0010-938X(18)31605-6 https://doi.org/10.1016/j.corsci.2018.12.018 CS 7812
To appear in: Received date: Revised date: Accepted date:
1 September 2018 8 November 2018 12 December 2018
Please cite this article as: Li X, Gong B, Deng C, Li Y, Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel, Corrosion Science (2018), https://doi.org/10.1016/j.corsci.2018.12.018 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel
Xiaogang Li1, Baoming Gong1,*, Caiyan Deng1,*, and Yizhe Li1,*
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Tianjin Key Laboratory of Advanced Joining Technology, Department of Materials Science and Engineering,
The Correspondence E-mail: [email protected] (C. Deng); [email protected] (B. Gong); [email protected] (Y. Li)
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*
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Tianjin University, Tianjin 300350, China.
Highlights
Failure location in hydrogen embrittled stainless steel joint depends on pre-strain
Severe strain-induced α′ martensite in the base metal contributes to failure mechanism transition
Tendency of strain-induced α′ martensitic transformation for the weld metal and base metal is
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Abstract
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strongly determined by local chemical homogeneity
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The effects of pre-strain on hydrogen embrittlement of 304L K-TIG welded joint were investigated. As
pre-strain of joint increased, fracture location of hydrogen-charged joint changed from weld metal to base metal. In low prestrained joints, fracture may be initiated at the sites of clustered dislocations due to
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hydrogen in weld metal with a high density of pre-existing dislocations. In high prestrained joints, some regions in base metal suffered severe strain-induced α′ martensitic transformation due to Ni depletion bands and cracking occurred at phase boundaries where may provide more sites for the accumulation of hydrogen atoms and dislocations.
Keywords: A. Stainless steel; B. EPMA; B. TEM; C. Effects of strain; C. Hydrogen embrittlement; C. 1
Welding.
1. Introduction
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Hydrogen may be introduced into metals over time through environmental exposure and corrosion processes. Hydrogen embrittlement (HE), a process associated with the presence and diffusion of hydrogen atoms, often has the deleterious consequence on mechanical properties, ultimately resulting in fracture. Due to the
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excellent HE resistance, mechanical properties, and weldability, 300 series austenitic stainless steels (ASS) have been widely used in gas and oil production and storage structures that are exposed to extreme hydrogen environments, i.e. hydrogen gas storage system for fuel-cell vehicles, boiling water reactor in nuclear power
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plants, and piping systems for petrochemical industry [1–11]. Cold work hardening through pre-straining is
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usually used to improve the mechanical behaviors of ASS prior to use [12]. It was found that pre-straining of
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ASS can introduce both strain hardening [13–15] and transformations in the phase and/or microstructure, i.e.
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α′ martensite formation [13,14], the diffusion and multiplication of dislocations [15], and the formation of deformation twins [16,17]; these changes alter the HE susceptibility of the ASS. It has been reported that the
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tensile elongation rate in a hydrogen environment may increase with increasing pre-strain [18–20]. However, under the same hydrogen charging condition, the hydrogen content increases with the pre-strain, resulting in
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more severe ductility loss in medium-carbon transformation-induced plasticity (TRIP) steel [21], and a similar phenomenon was reported for 304L steel [22]. Since these aforementioned studies did not thoroughly
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explain the HE mechanisms of pre-straining ASS materials, some contradictory results of the effect of hydrogen on plasticity can occasionally be found in the literature [18–21]. In general, it is widely accepted that during the pre-strain process, metastable ASS undergoes
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strain-induced martensite transformation during plastic deformation [13–15], and α′ martensite can work as a ‘highway’ for hydrogen motion [23–25] due to the higher hydrogen diffusion rate of α′ martensite [26–28]. Therefore, it is expected that pre-strained metastable ASS is more sensitive to HE [22]. In addition, compared with the base metal, the ASS weld metal usually suffers more severe HE [29–31] due to its distinct microstructure [32,33] which obviously affected hydrogen transport [34]. Therefore, the study of the HE mechanism in pre-strained ASS and its weld joint is essential and not exhausted to date. Keyhole tungsten 2
inert gas (K-TIG) welding is an efficient technique widely applied to ASS, and it can allow for deeper penetration than does the classical TIG method [35]. However, there is no literature about the effect of pre-strain on microstructure and hydrogen embrittlement of ASS K-TIG welded joints. For pre-strained ASS widely employed in hydrogen-containing environments, the HE mechanism of pre-strained K-TIG welded joints should be critically investigated. Therefore, the primary aim of the current study is to experimentally investigate the influence of
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pre-strain on microstructure and hydrogen embrittlement in AISI 304L K-TIG welded joints. Pre-strain, electrochemical hydrogen charging, and subsequent tensile tests were conducted to investigate the relationship between the microstructure and failure mechanism using scanning electron microscopy (SEM),
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electron probe microanalysis (EPMA), electron backscatter diffraction (ESBD), and transmission electron microscopy (TEM).
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2. Material and methods
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2.1 Materials and welding parameters
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The base metal used in this study was industrially hot-rolled AISI 304L stainless steel plate solution-annealed at 1050 ℃ for 0.5 h and water-quenched with dimensions 300 mm × 150 mm × 10 mm.
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Butt joints were prepared using a K-TIG torch with DC electrode negative polarity and no metal was filled. The welding parameters used in this study are listed in Table 1. The average chemical compositions (in mass
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percentages) of the weld metal and base metal were measured by optical spark emission spectrometry (OES) and are presented in Table 2. An EBSD system (FE-SEM, JEOL, JSM-7100F) was operated at 20 kV to
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observe the microstructures of the K-TIG weld joint. The specimen for EBSD observations was ground to 3000 grit and electropolished using a solution of ethyl alcohol (95 vol%) and perchloric acid (5 vol%) at 30 V for 30 s. The work distance and step size for EBSD operation were 20 mm and 0.1 μm, respectively. The
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raw data of EBSD observations were processed using Orientation Imaging Microscopy software (TSL-OIM). 2.2 Pre-strain After welding, the specimens for pre-strain were cut from the welded plates, as shown in Fig. 1a, and then an overall plastic pre-strain of either 3%, 6%, 10%, 15%, 20%, or 25% was applied to these specimens by using a 100 kN servo-hydraulic universal testing machine (MTS E64) with a strain rate of 2.7×10-4 s-1. 3
The strain was measured by a clip gauge. After the pre-strain process, the microstructures in the weld and the base metal of the pre-strained specimens were detected by using the EBSD system mentioned above. The local chemical compositions in the weld and the base metal of the welded joint after pre-strain were measured using an EPMA instrument (FE-EPMA, JXA-8530 F). The quantitative calibrations of alloying elements (including Ni, Cr, Mn, Mo, Si, and C) were carried out with the corresponding standard samples. Mapping was carried out for the Ni, Cr, Mn, Mo, Si, and C. During pre-strain process, digital image
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correlation (DIC) method was used to measure the full-field strain distribution of welded joint. The DIC system in this study composes of two high-resolution (4872 × 3248 pixels) GE4900 cameras, two lamps as
Correlated Solutions Inc. 2.3 Cathodic hydrogen charging and measurement of hydrogen content
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the light source, calibration grids, a tripod with supporting hardware, and the software package VIC-3D from
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After pre-strain, the tensile specimens were cut from the pre-strained specimens with dimensions as
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shown in Fig. 1b. The longitudinal directions of the tensile specimens were perpendicular to the weld bead.
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The welded bead was placed in the middle of the specimen. One group of the tensile specimens with different pre-strains was cathodically charged at 25 ℃ to introduce hydrogen in 0.5 mol/L H2SO4 + 3 g/L
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NH4SCN solution with a current density of 50 mA/cm2 for 24 h, and the specimens were completely
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immersed during the hydrogen charging, as shown in Fig. 2. The hydrogen content in the specimen was immediately determined by thermal desorption spectroscopy (TDS) analysis using a quadrupole mass
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spectrometer (R-DEC Company, EX0014) with the accuracy of 0.01 wppm. The hydrogen charging conditions was the same for the tensile specimens and the TDS specimens. The test was conducted at a
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heating rate of 100 °C/h in the temperature range from 25 °C to 800 °C. Under each condition, four samples were repeated to ensure the data reliability. The dimensions of the specimen for hydrogen measurement are shown in Fig. 1b.
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2.3 Tensile tests
After hydrogen charging, tensile tests were immediately conducted for hydrogen-charged specimens
with different pre-strains by using a micro-mechanical testing machine (Instron 5848) at room temperature with a strain rate of 2.7×10-4 s-1. The strain was measured by a clip gauge. The hydrogen-free specimens with different pre-strains were also tested as a reference. Specimens in each condition were tested in quadruplicate to ensure data reliability. Two indexes, Iδ and Iψ, are adopted to quantitatively describe HE sensitivity: 4
I δ (%)
0 H 100% 0
(1)
I ψ (%)
0 H 100% 0
(2)
where δ and Ψ are the elongation to failure and the reduction of area of the specimens, respectively. The subscripts 0 and H denote the uncharged and charged specimens, respectively. After the tensile tests, the fracture morphologies were observed by scanning electron microscopy (SEM,
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Hitachi, S4800). The deformation microstructures around the necking zone of the broken samples were observed by using the EBSD system mentioned above and a field-emission transmission electron microscope
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(FE-TEM, JEOL, JEM-2100F) operated at 200 kV. The samples for TEM observation were prepared by twin-jet electro-chemical polishing using an electrolyte consisting of ethyl alcohol (90 vol%) and perchloric acid (10 vol%). The step size of accelerating voltage for TEM operation was 0.05 kV. Energy dispersive
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X-ray spectrometry (EDS) in the SEM system was performed for local chemical composition analysis.
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3.1 Effect of pre-strain on microstructure evolution
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3. Results and discussion
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The representative microstructure of a welded joint was analysed by EBSD mapping, as shown in Fig. 3. The matrix area and heat-affected zone consisted of equiaxial grains and a fusion zone composed of columnar crystals. The heat-affected zone was very narrow (about 200 ± 20 μm), with grain sizes between 10
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and 80 μm. The average grain size of the base metal was slightly larger than that in the heat-affected zone. After pre-strain, deformed microstructures in the welded joint were examined (Figs. 4, 5, and 6). In this
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paper, α' martensite obtained from the pre-strain process is denoted as “prior α' martensite”, to distinguish it from α' martensite formed during the subsequent tensile testing, which is called “dynamic α' martensite” [36]. Regardless of the pre-strain, there was almost no prior α′ martensite in the weld metal, as shown in Fig. 4a
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and 5b. However, the prior α′ martensite content in the base metal dramatically increased with the pre-strain, especially when the pre-strain from 15%, as indicated in Fig. 7. Strain-induced α′ martensitic transformation is affected by grain size, deformation, and local chemical inhomogeneity. Larger grains are known to facilitate strain-induced α′ martensite formation [37] and the volume fraction of α′ martensite increases with increasing deformation [38]. While the average grain size of the weld metal was larger than that in the base metal (Fig. 3) and the weld metal suffered the highest local strains (Fig. 8), almost no prior α′ martensite was 5
observed in the weld metal. Therefore, we ruled out the influences of grain size and deformation on the strain-induced α′ martensite in this study. The nickel equivalent Nieq influences strain-induced α′ martensitic transformation [39]: the lower Nieq was more beneficial for strain-induced α′ martensitic transformation [40]. Nieq can be calculated as follows:
Nieq =Ni+0.65Cr+0.98Mo+1.05Mn+0.35Si+12.6C
(3)
where all elements are represented in mass fraction. The results of EPMA mapping measurements (Fig. 5d
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and 6d) suggested that the base metal exhibited pronounced Ni segregation bands parallel to the rolling direction, resulting in the Ni depletion bands, which corresponded to the band like distribution of α′
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martensite shown in Fig. 6b. In addition, the results of the local chemistry measurements with EPMA (Figs. 5c and 6c) also demonstrated that mass fractions of main elements (Ni, Cr, Mn, Si, Mo, and C), especially Ni, in the prior α′ martensite region of the base metal were lower than that in the weld metal, resulting in the
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lower Nieq of the prior α′ martensite region of the base metal according to Eq. (3), which indicated that there were some regions in the base metal to prone to the transformation of strain-induced α′ martensite.
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Consequently, the different tendency of strain-induced α′ martensitic transformation for the weld metal and
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base metal was strongly determined by local chemical inhomogeneity due to Ni segregation in the base metal
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with the pre-strain, as shown in Fig. 9.
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in this study. Moreover, it was found that the dislocation density in the weld metal increased significantly
3.2 Effect of pre-strain on hydrogen content
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Figure 10 shows the typical TDS curves of the hydrogen-charged specimens with various pre-strain values. The hydrogen desorption peak was clearly observed at around 120 °C; and the peak generally
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increased with the pre-strain, indicating a higher hydrogen content. Figure 11 illustrates the relationship between the hydrogen content and the pre-strains. The temperature range used for calculation of hydrogen content was from 25 °C to 800 °C. At low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the hydrogen
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content slightly increased. However, a large amount of hydrogen started to be introduced into the specimen at 20% and 25% pre-strain. Meanwhile, the amount of hydrogen was a minimum at 0% pre-strain. In this study, the relationship between hydrogen content (C0) in the whole welded joint and pre-strain (εpre) can be fitted exponentially as follows:
C0 17.0405 7.0398exp(0.1127 pre )
(4)
where a similar tendency was observed in the literature in both experimental [41,42] and theoretical [43] 6
studies. Since hydrogen can be trapped by dislocations, increasing the pre-strain can thus result in a much higher dislocation density (Fig. 9) as there is an increased probability of trapping hydrogen atoms. Furthermore, α′ martensite is considered as a ‘highway’ for hydrogen motion in austenitic stainless steels [23–25] because of its higher hydrogen diffusion coefficient [26–28], and thus the prior α’ martensite increases the hydrogen content in 304L steel [22]. The correlation between strain-induced α′ martensite content and strain increment can be expressed as [38]:
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n f α 1 exp 0 1 exp
(5)
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where fα′ is the strain-induced α′ martensite content and ε is the strain. The parameter α describes the formation of a shear band which depends mainly on the stacking-fault energy, β0 is linearly dependent on the probability of martensite embryo formation, and the exponent n is equal to 4.5. Accordingly, if the pre-strain
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is increased further, the prior α′ martensite content in the base metal becomes higher, as shown in Fig. 7 and
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from dislocation and prior α′ martensite simultaneously.
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predicted by Eq. (5). In summary, the hydrogen content increases with increasing pre-strain, which results
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3.3 Effect of deformed microstructures induced by pre-strain and H atoms on tensile properties
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3.3.1 Tensile properties
All hydrogen-free specimens subjected to different pre-strain fractured in the weld metal. The
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hydrogen-charged specimens with 0%, 3%, 6%, 10%, and 15% pre-strain fractured in the weld metal, while those pre-strained to 20% and 25% fractured in the base metal. The engineering stress–true strain curves for
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hydrogen-free and hydrogen-charged specimens are presented in Fig. 12. For all hydrogen-free specimens, the flow stress increased with pre-strain, while the elongation at failure decreased with increasing pre-strain value. More specifically, at the low pre-strain level (0%, 3%, 6%, 10%, and 15% pre-strain), the hydrogen-charged specimens could withstand higher flow stress compared with the hydrogen-free specimens;
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at high pre-strain deformation (20% and 25% pre-strain), lower flow stresses were measured for hydrogen-charged specimens. In addition, as the pre-strain increased, the ultimate tensile strength of hydrogen-free specimens (UTS0) increased monotonically; however, the ultimate tensile strength of hydrogen-charged specimens (UTSH) initially increased until 20% pre-strain, followed by a sharp decline (Fig. 13a). It was found that at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the UTS was almost 7
identical, regardless of the hydrogen content. Once the pre-strain reached 20%, UTSH deviated from UTS0. In addition, similar effects of the pre-strain on the macroscopic plasticity are demonstrated in Figs. 13b and 13c. The HE sensitivity index Iδ reached 14.3 ± 0.35% and 29.0 + 1.34% at pre-strain values of 10% and 25%, respectively. With increasing pre-strain, the dislocation density in the weld metal increases further (Fig. 9) and tangled dislocations may appear. The interaction between dislocations and H atoms trapped in the dislocation
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could hinder dislocation motion. Therefore, the movement of dislocations requires a higher flow stress, leading to strain hardening. In addition, the formation of strain-induced α′ martensite in the base metal may also result in strain hardening. Therefore, for all hydrogen-free and hydrogen-charged specimens with low
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pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the strain hardening effect led to a noticeable increase in the flow stress and UTS, and a significant reduction in the elongation at failure. When the pre-strain was greater than 15%, the hydrogen content further increased (Fig. 11). A high hydrogen concentration can
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reduce dislocation formation energies [44] and the cohesive energy of twins and grain boundaries [45,46]. As
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a result, the flow stress and the UTSH values of the hydrogen-charged specimens with 15% and 20%
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pre-strain were lower. When the specimen is plastically deformed, two competitive effects may exist, namely,
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an increment in the flow stress and UTSH due to strain hardening and a decrease in the flow stress and UTSH due to hydrogen. Strain hardening was prevalent at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain)
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while the strength reduction due to the high hydrogen concentration was dominant at high pre-strain (20% and 25% pre-strain).
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3.3.2 Fractographic observations
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Figure 14 illustrates the representative fracture morphologies of the fractured samples determined using
low magnification SEM. It is demonstrated in Fig. 14a that the entire fracture surface of hydrogen-free samples was characterized by ductile dimples with severe plastic localization. In contrast, two distinct
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morphologies can be identified on the fracture surface of the hydrogen-charged specimens: (i) brittle fracture features around the edge zone, resulting from the higher hydrogen concentration at the surface of electrolytically-charged samples and the higher stress intensity factor of the surface defect [47], and (ii) a central portion with and dimples. In addition, Fig. 15 shows that the proportion of the brittle zone increased markedly with increasing pre-strain, which is consistent with the loss of macroscopic ductility indicated by Figs. 13b and 13c. It is worth noting that at a pre-strain of 20% and 25%, the crack initiates and propagates 8
in the base metal, and the brittle zone area increased significantly. Figure 16 shows high magnification SEM images of the fracture surface corresponding to the brittle zones shown in Fig. 14. The entire fracture surface of the hydrogen-free samples was dominated by dimples, indicating only ductile fracture. At the pre-strain of 0% and 3%, the brittle fracture surface exhibited quasi-cleavage fractures. As the pre-strain increased further, the cleavage fracture in the brittle zone became the dominant failure mechanism. Moreover, the brittle zone of the hydrogen-charged specimens with 20% and 25% pre-strain was dominated by the typical cleavage
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fracture mode characterized by flat facet features. 3.4 Mechanisms of failure location transition associated with hydrogen and α′ martensite
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Figures 17 and 18 show the deformed microstructures of the hydrogen-charged specimens with different pre-strains. A large amount of α′ martensite was observed in the hydrogen-charged specimen with 20% pre-strain (broken in the base metal) (Fig. 18b) and almost no α′ martensite was observed in the
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hydrogen-charged specimen with 10% pre-strain (broken in the weld metal) (Fig. 17b). Local chemistry
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measurements with EDS in the SEM system were performed and linked to microstructural analysis, as
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shown in Figs. 17c and 18c. Mass fractions of main elements (Ni, Cr, Mn, Si, and Mo) in the base metal
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(especially Ni in some regions, i.e. region 1 and region 3 marked by black boxes in Fig. 18c) were lower than that in the weld metal, resulting in the lower Nieq of the base metal according to Eq. (3) (consistent with the
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results in section 3.1). Moreover, cracks occurred at the phase boundary in the hydrogen-charged specimen with 20% pre-strain (broken in the base metal) close to the fracture surface, as shown in Fig. 19 Therefore, under the condition that the welded joint experienced high pre-strain (20% and 25% pre-strain), large
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amounts of strain-induced α′ martensitic transformation and phase boundaries in the base metal may lead to
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the failure location transition from weld metal to base metal in hydrogen-charged welded joint with different pre-strains. The influence of the interactions between hydrogen and the microstructures induced by pre-strain on the failure mechanisms will be discussed in more detail. After the low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain) of the joints, more dislocations were
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induced in the weld metal than in the base metal due to localized strain in the weld metal (Fig. 8), while the prior α′ martensite content was low in both the weld and base metal. Several studies have reported that the hydrogen diffusivity was not significantly affected by pre-strain in the ASS (fcc) without strain-induced martensitic transformation [48,49]. It is considered that the binding energy of hydrogen to a dislocation is not much higher than the activation energy of hydrogen diffusion in the lattice in fcc metals [48]. Increasing the number of dislocations may increase the number of trap sites for hydrogen [50]. After hydrogen charging, the 9
distributed dislocations in the weld metal trapped hydrogen and the movement of dislocations in subsequent tensile deformation would accompany with transportation of hydrogen. Moreover, grain boundaries could function as hydrogen traps and the hydrogen atoms could aggregate at the grain boundaries [51]. In the subsequent tensile tests, some crystalline defects in ASS, i.e. twin boundaries and grain boundaries, may hinder the movement of dislocations and lead to the pile-up of dislocations. Meanwhile, hydrogen atoms are very likely to migrate to the twin boundaries [52] during plastic deformation and may accumulate at the grain
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boundaries [51]. Furthermore, owing to the increased transport of hydrogen mediated by dislocations, hydrogen may also move and enriches at the sites of these clustered dislocations. Hydrogen atoms can also induce a new dislocation slip plane [46] and may lower the cohesive energy [51,53]. Therefore, the
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concentrated hydrogen promotes dislocation mobility, leading to macroscopic HE, and fracture may initiate and propagate along the interface where dislocations and hydrogen are accumulated [36,54]. As expected, the dislocation distribution was inhomogeneous (Fig. 20a), and a local high density of dislocations could be
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observed at the deformation twin boundaries, grain boundaries, and other sites (Figs. 20b, 20c, and 20d).
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After the high pre-strain (20% and 25% pre-strain) of the joints, the prior α′ martensite content in the
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base metal sharply increased (Fig. 7), and thus the hydrogen content in the specimens dramatically increased
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due to prior α′ martensite acting as a path for hydrogen transport. However, the prior α′ martensite content was still low in the weld metal (Fig. 5b) due to the high stability of austenite. As a result of the higher
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specific volume of martensite compared to austenite, the transformation of austenite into martensite may result in strain localization around the phase boundaries. After hydrogen-charging, due to the lower solubility
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and higher diffusion coefficient of hydrogen in martensite than austenite, hydrogen atoms collected at the phase boundary between austenite and prior α′ martensite. The deformation in subsequent tensile testing
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mainly occurred in residual austenite due to the higher strength of α′ martensite than austenite. Therefore, incompatible deformation between α′ martensite and the residual austenite resulted in a high local stress concentration at the phase boundary and the residual austenite near phase boundary is preferentially
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transformed into dynamic α′ martensite. Meanwhile, the transformation of the residual austenite into martensite resulted in the release of excessive hydrogen [55], resulting in more dislocations and a higher concentration of hydrogen atoms at the phase boundary. The extreme local enrichment of hydrogen atoms could enhance the dislocation mobility around the phase boundary and potentially resulted in a lower grain boundary energy, cohesive energy [51,53], and even critical material fracture strain [56]. In addition, the additional tensile stress introduced by hydrogen increased linearly with the increasing hydrogen 10
concentration [57,58]. With the application of an external tensile stress, the additional stress induced by hydrogen will be higher than the bonding force (weakened by the presence of hydrogen) and the fracture was initiated at the phase boundaries (Fig. 19b) where the hydrogen atoms and dislocations are accumulated. As expected, a local high density of dislocations was observed around the phase boundary (Fig. 21). Therefore, at low pre-strain levels (0%, 3%, 6%, 10%, and 15% pre-strain) the tensile fracture in the weld metal of the welded joint was a result of trapped hydrogen atoms at high density dislocations; whereas under high
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pre-strain deformation (20% and 25% pre-strain), failure occurred in the base metal of the welded joint and was caused by cracking along the α′ martensite phase boundaries. Fig. 22 shows failure processes and
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mechanisms under different pre-strain.
5. Conclusion
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The effects of pre-strain on failure mechanism of hydrogen embrittlement in 304L K-TIG welded joint
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were investigated in the study. The correlation between fracture location of hydrogen-charged welded joint
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and the microstructural transformations induced by pre-strain was discussed. The results are concluded as
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follows:
(1) While there was almost no prior α′ martensite in the weld metal regardless of the pre-strain of the welded
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joint, the prior α′ martensite content in the base metal dramatically increased with the pre-strain of the welded joint, resulting from the local chemical inhomogeneity. The weld metal was segregation-free,
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whilst there were some regions of low austenite stability in Ni depletion bands of the base metal. (2) As the pre-strain of joint increased, the fracture location of hydrogen-charged joint moved from the weld
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metal to the base metal; it was attributed to the combined influence of the hydrogen concentration and microstructures induced by pre-straining, i.e. dislocations and prior α′ martensite. (3) In the low pre-strained joints (0%, 3%, 6%, 10%, and 15% pre-strain), the weld metal with a high
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density of pre-existing dislocations was prone to local plastic deformation due to the high mobility of dislocations enhanced by the hydrogen. Hydrogen-induced fracture may be initiated at the sites of clustered dislocations and hydrogen in the weld metal. (4) In the high pre-strained joints (20% and 25% pre-strain), the base metal underwent severe strain-induced α′ martensitic transformation, and brittle fracture occurred at α′ martensite/austenite phase boundaries which may suffer more severely local enrichment of hydrogen atoms and dislocations in the base metal. 11
(5) In the low pre-strained joints (0%, 3%, 6%, 10%, and 15% pre-strain), the flow stress and ultimate tensile strength of the hydrogen-charged joints were higher than those of the hydrogen-free joints at the same pre-strain, whereas in the high pre-strained joints (20% and 25% pre-strain), the flow stress and ultimate tensile strength of the hydrogen-charged joints were lower than those of the hydrogen-free joints at the same pre-strain. A competitive relationship between strain hardening and strength reduction due to hydrogen exists in the pre-strained joints: at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain),
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strain hardening was prevalent while at high pre-strain (20% and 25% pre-strain) the strength reduction was dominant due to the high hydrogen concentration.
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Acknowledgements
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The research is financially supported by National Key R&D Program of China (Grant No.
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2018YFC0310306) and National Natural Science Foundation of China (NSFC) (Grant No. 51305295).
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References
[1] D. Peckner, I.M. Bernstein, Handbook of Stainless Steels, McGraw Hill, 1977.
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[2] GR. Caskey, Hydrogen Compatibility Handbook for Stainless, Aiken, South Carolina: E I du Pont Nemours, Savannah River Laboratory, 1983.
[3] R. Nishimura, O.M. Alyousif, A new aspect on intergranular hydrogen embrittlement mechanism of solution
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annealed types 304, 316 and 310 austenitic stainless steels, Corros. Sci. 51 (2009) 1894-1900. [4] L.W. Tsay, S.C. Yu, R.T. Huang, Effect of austenite instability on the hydrogen-enhanced crack growth of austenitic stainless steels, Corros. Sci. 49 (2007) 2973-2984.
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[5] Y. Mine, K. Hirashita, M. Matsuda, M. Otsu, K. Takashima, Effect of hydrogen on tensile behaviour of micrometre-sized specimen fabricated from a metastable austenitic stainless steel, Corros. Sci. 53 (2011) 529-533. [6] A.M. Brass, J. Chêne, Hydrogen uptake in 316L stainless steel: consequenceson the tensile propertis, Corros. Sci. 48 (2006) 3222-3242.
[7] L.W. Tsay, J.J. Chen, J.C. Huang, Hydrogen-assisted fatigue crack growth of AISI 316L stainless steel weld, Corros.
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Sci. 50 (2008) 2973-2980.
[8] M. Koyama, E. Akiyama, T. Sawaguchi, K. Ogawa, I. Kireeva, Y.I. Chumlyakov, K. Tsuzaki, Hydrogen-assisted quasi-cleavage fracture in a single cystalline type 316 austenitic stainless steel, Corros. Sci. 75 (2013) 345-353.
[9] Y. Jiao, W. Zheng, J.R. Kish, Stress corrosion cracking susceptibility of thermally-aged Type 310S stainless steels in supercritical water, Corros. Sci. 135 (2018) 1-11. [10] C.L. Lai, L.W. Tsay, C. Chen, Effect of microstructure on hydrogen embrittlement of various stainless steels, Mater. Sci. Eng. A 584 (2013) 14-20. [11] X. Zhong, S.C. Bali, T. Shoji, Effects of dissolved hydrogen and surface condition on the intergranular stress 12
corrosion cracking initiation and short crack growth behavior of non-sensitized 316 stainless steel in simulated PWR primary water, Corros. Sci. 118 (2017) 143-157. [12] N. Narita, H.K. Brinbaum, On the role of phase transitions in the hydrogen embrittlement of stainless steels, Scr. Metall. 14 (1980) 1355-1358. [13] A.K. De, J.G. Speer, D.K. Matlock, D.C. Murdock, M.C. Mataya, R.J. Comstock Jr, Deformation-induced phase transformation and strain hardening in type 304 austenitic stainless steel, Metall. Mater. Trans. A 37 (2006) 1875-1886. [14] X.F. Fang, W. Dahl, Strain hardening and transformation mechanism of deformation-induced martensite transformation in metastable austenitic stainless steels, Mater. Sci. Eng. A 141 (1991) 189-198. [15] T. Narutani, Effect of deformation-induced martensite transformation on the plastic behaviour of metastable
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austenitic stainless steel, Mater. Trans. Jim 30 (1989) 33-45.
[16] T.H. Lee, C.S. Oh, S.J. Kim, S. Takaki, Deformation twinning in high-nitrogen austenitic stainless steel, Acta Mater. 55 (2007) 3649-3662.
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[17] Y.F. Shen, X.X. Li, X. Sun, Y.D. Wang, L. Zuo, Twinning and martensite in a 304 austenitic stainless steel, Mater. Sci. Eng. A 552 (2012) 514-522.
[18] K. Keiji, I. Shozo, U. Hitoshi, K. Yasushi, T. Takafumi, Hydrogen embrittlement of sensitized SUS 316 steel, J. Soc. Mater. Sci. Jpn. 45 (1996) 1078-1082. transport by dislocations,Mater. Sci. Eng. A 262 (1999) 173-183.
U
[19] F. Lecoester, J. Chene, D. Noel, Hydrogen embrittlement of the Ni-base alloy 600 correlated with hydrogen [20] T. Takasuji, S.Hanada, The effect of pre-deformation on moisture-induced embrittlement of Ni3Al alloys,
N
Intermetallics 5 (1997) 127-135.
[21] W.J. Hui, Y. Li, Y.J. Zhang, C. Zhou, H. Dong, Effect of prestraining on hydrogen absorption and delayed fracture
A
behaviour of a medium-carbon TRIP steel, T. Mater. Heat Treat. 33 (2012) 42-46.
M
[22] Y. Wang, X. Wang, J. Gong, L. Shen, W. Dong, Hydrogen embrittlement of catholically hydrogen-precharged 304L austenitic stainless steel: Effect of plastic pre-strain, Int. J. Hydrogen Energy 39 (2014) 13909-13918. [23] Y. Murakami, Effects of hydrogen on metal fatigue, in: In: Proceedings of International Hydrogen Energy
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Development Forum, Fukuoka, 2007. [24] M. Martin, S. Weber, C. Izawa, S. Wagner, A. Pundt, W. Theisen, Influence of machining-induced martensite on hydrogen-assisted fracture of AISI type 304 austenitic stainless steel, Int. J. Hydrogen Energy 36 (2011)
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11195-11206.
[25] T. Kanezaki, C. Narazaki, Y. Mine, S. Matsuoka, Y. Murakami, Effects of hydrogen on fatigue crack growth behavior of austenitic stainless steels, Int. J. Hydrogen Energy 33 (2008) 2604-2619.
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[26] G. Han, J. He, S. Fukuyama, K. Yokogawa, Effect of strain-induced martensite on hydrogen environment embrittlement of sensitized austenitic stainless steels at low temperatures, Acta Mater. 46 (1998) 4559-4570. [27] TP. Perng, CJ. Altstetter, Comparison of hydrogen gas embrittlement of austenitic and ferritic stainless steels, Metall. Mater. Trans. A 18 (1987) 123-134. [28] L. Zhang, M. Imade, B. An, M. Wen, T. Iijima, S. Fukuyama, K. Yokogawa, Internal reversible hydrogen
A
embrittlement of austenitic stainless steels based on type 316 at low temperatures, ISIJ Inter. 99 (2013) 294-301.
[29] T. Michler, Y. Lee, R.P. Gangloff, J. Naumann, Influence of macro segregation on hydrogen environment embrittlement of SUS 316L stainless steel, Int. J. Hydrogen Energy 34 (2009) 3201-3209.
[30] C.M. Younes, A.M. Steele, J.A. Nicholson, C.J. Barnett, Influence of hydrogen content on the tensile properties and fracture of austenitic stainless steel welds, Int. J. Hydrogen Energy 38 (2013) 4864-4876. [31] T. Michler, Toughness and hydrogen compatibility of austenitic stainless steel welds at cryogenic temperatures, Int. J. Hydrogen Energy 32 (2007) 4081-4088. [32] H.F. Jackson, K.A. Nibur, C.S. Marchi, J.D. Puskar, B.P. Somerday, Hydrogen-assisted crack propagation in 13
304L/308L and 21Cr-6Ni-9Mn/308L austenitic stainless steel fusion welds, Corros. Sci. 60 (2012) 136-144. [33] J. Naumann, T. Michler, Hydrogen environment embrittlement of orbital welded austenitic stainless steels at -50°C, Int. J. Hydrogen Energy 34 (2009) 6478-6483. [34] V. Olden, A. Alvaro, O.M. Akselsen, Hydrogen diffusion and hydrogen influenced critical stress intensity in an API X70 pipeline steel welded joint-Experiments and FE simulations, Int. J. Hydrogen Energy 37 (2012) 11474-11486. [35] M. Lohse, U. Fuessel, H. Schuster, J. Friedel, M. Schnick, Keyhole welding with CF-TIG (cathode focussed GTA), Weld. World 57 (2013) 735-741. [36] L. Zhang, Z. Li, J. Zheng, Y. Zhao, P. Xu, C. Zhou, X. Li, Effect of strain-induced martensite on hydrogen embrittlement of austenitic stainless steels investigated by combined tension and hydrogen release methods, Int. J.
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Hydrogen Energy 38 (2013) 8208-8214.
[37] C. Yoo, Y. Park, Y. Jung, Y. Lee, Effect of grain size on transformation-induced plasticity in an ultrafine-grained metastable austenitic steel, Scr. Mater. 59 (2008) 71-74.
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[38] G.B. Olson, M. Cohen, Kinetics of strain-induced martensite nucleation, Metall. Mater. Trans. A 6 (1975) 791-795. [39] T. Hirayama, M. Ogirima, Influence of chemical composition on martensitic transformation in Fe-Cr-Ni stainless steel, J. Jpn. Inst. Met. Mater. 34 (1970) 507-510.
[40] L. Zhang, M. Wen, M. Imade, S. Fukuyama, K. Yokogawa, Effect of nickel equivalent on hydrogen gas embrittlement of austenitic stainless steels based on type 316 at low temperatures, Acta Mater. 56 (2008)
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3414-3421.
[41] M. Nagumo, K. Ohta, H. Saitoh, Deformation induced defects in iron revealed by thermal desorption spectroscopy
N
of tritium, Scr. Mater. 40 (1999) 313-319.
[42] M. Nagumo, K. Takai, N. Okuda, Nature of hydrogen trapping sites in steels induced by plastic deformation, J.
A
Alloys Compd. 293 (1999) 310-316.
M
[43] S. Wang, S. Ohnuki, N. Hashimoto, K. Chiba, Hydrogen effects on tensile property of pure iron with deformed surface, Mater. Sci. Eng. A 560 (2013) 332-338.
[44] Caskey, G.R. Jr, Fractography of hydrogen-embrittled iron-chromium-nickel alloys, in: In: For Presentation at the
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13th Annual Technical Meeting of the International Metallographic Society, Brighton England, 1980. [45] V.G. Gavriljuk, V.N. Bugaev, Y.N. Petrov, A.V. Tarasenko, B.Z. Yanchitski, Hydrogen-induced equilibrium vacancies in fcc iron-base alloys, Scr. Mater. 34 (1996) 903-907.
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[46] T.D. Lee, I.M. Bernstein, Effects of hydrogen on dislocation morphology in spheroidized steel, Acta Metall. Mater. 39 (1991) 363-372.
[47] G.E. Dieter, Mechanical Metallurgy, third ed., McGraw Hill, London, 1986.
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[48] Y. Mine, C. Narazaki, K. Murakami, S. Matsuoka, Y. Murakami, Hydrogen transport in solution-treated and pre-strained austenitic stainless steels and its role in hydrogen-enhanced fatigue crack growth, Int. J. Hydrogen Energy 34 (2009) 1097-1107.
[49] T.P. Perng, C.J. Altstetter, Effects of deformation on hydrogen permeation in austenitic stainless steels, Acta Metall. 34 (1986) 1771-1781.
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[50] X. Li, Y. Wang, P. Zhang, B. Li, X. Song, J. Chen, Effect of pre-strain on hydrogen embrittlement of high strength steels, Mater. Sci. Eng. A 616 (2014) 116-122.
[51] M. Yamaguchi, K. Ebihara, M. Itakura, T. Kadoyoshi, T. Suzudo, H. Kaburaki, First-principles study on the grain boundary embrittlement of metals by solute segregation: Part II. Metal (Fe, Al, Cu)-hydrogen (H) systems, Metall. Mater. Trans. A 42 (2011), 330-339. [52] H. Ji, I. Park, S. Lee, Y. Lee, The effect of pre-strain on hydrogen embrittlement in 310S stainless steel, J. Alloy. Compd. 598 (2014) 205-212. [53] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation II. 14
Experimental evidence and consequences, Acta Mater. 55 (2007) 5139-5148. [54] P. Novak, R. Yuan, B.P. Somerday, P. Sofronis, R.O. Ritchie, A statistical, physical-based, micro-mechanical model of hydrogen-induced intergranular fracture in steel, J. Mech. Phys. Solids. 58 (2010) 206-226. [55] Y. Mine, Z. Horita, Y. Murakami, Effect of hydrogen on martensite formation in austenitic stainless steels in high-pressure torsion, Acta Mater. 57 (2009) 2993-3002. [56] Y.A. Du, L. Ismer, J. Rogal,T. Hickel, J. Neugebauer, R. Drautz, First-principles study on the interaction of H interstitials with grain boundaries in alpha- and gamma-Fe, Phys. Rev. B 84 (2011) 667-673. [57] R. Wang, Effects of hydrogen on the fracture toughness of a X70 pipeline steel, Corros. Sci. 51 (2009) 2803-2810. [58] T. Zhang, W.Y. Chu, K.W. Gao, L.J. Qiao, Study of correlation between hydrogen-induced stress and hydrogen
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embrittlement, Mater. Sci. Eng. A 347 (2003) 291-299.
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Figure Captions
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Fig. 1. Dimensions and sampling location of specimens for: (a) pre-strain and (b) tensile tests and hydrogen content measurements (in mm).
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Fig. 2. Schematic of the electrolytic hydrogen charging system.
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Fig. 3. EBSD images of the fusion zone, heat-affected zone, and base metal of the welded joint: (a) inverse pole figure (IPF) map and (b) misorientation distribution.
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Fig. 4. EBSD analyses the welded joint after 10% pre-strain: (a) weld metal and (b) base metal.
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Fig. 5. Microstructure and local chemical composition analysis of the weld metal of the welded joint after 20% pre-strain: (a) detected region, (b) IPF and phase map measured by EBSD, (c) local chemistry measurements with EPMA obtained from the locations marked by black boxes in Fig. 5 (b), and (d) the distributions of different alloying elements (Ni, Cr, Mn, Si, Mo, and C) of the region shown in Fig. 5 (b), obtained by EPMA mapping.
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Fig. 6. Microstructure and local chemical composition analysis of the base metal of the welded joint after 20% pre-strain: (a) detected region, (b) IPF and phase map measured by EBSD, (c) local chemistry measurements with EPMA obtained from the locations marked by black boxes in Fig. 6 (b), and (d) the distributions of different alloying elements (Ni, Cr, Mn, Si, Mo, and C) of the region shown in Fig. 6 (b), obtained by EPMA mapping.
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Fig. 7. Average prior α′ martensite content (from EBSD measurements) of six randomly selected regions in base metals of welded joints with different pre-strains.
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Fig. 8. Map of strain distribution of the hydrogen-free specimen with 15% pre-strain obtained from digital image correlation measurement.
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Fig. 9. TEM analyses of dislocation in weld metals of welded joints after different pre-strain: (a) 6% pre-strain, (b) 10% pre-strain, and (c) 15% pre-strain.
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Fig. 10. TDS curves of hydrogen-charged specimens at various pre-strain values.
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Fig. 11. Hydrogen content determined at various pre-strain values, as measured by TDS analysis using a quadrupole mass spectrometer.
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Fig. 12. Representative engineering stress–true strain curves of hydrogen-free and hydrogen-charged specimens at different pre-strain values: (a) 6% pre-strain, (b) 15% pre-strain, (c) 20% pre-strain, and (d) 25% pre-strain.
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Fig. 13. Tensile properties of hydrogen-free and hydrogen-charged specimens at different pre-strain values: (a) ultimate tensile strength UTS, (b) elongation and hydrogen embrittlement sensitivity index Iδ, and (c) reduction of area and hydrogen embrittlement sensitivity index IΨ. Specimens in each condition were tested in quadruplicate.
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Fig. 14. Low magnification SEM images of morphologies of fractured specimens.
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Fig. 15. Areal proportion of brittle zone vs. the pre-strain, as observed in the fracture morphologies.
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Fig. 16. High magnification SEM images of morphologies of fractured specimens.
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IP T SC R U N A M ED PT CC E A Fig. 18. EBSD analyses of the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal): (a) IPF and (b) phase map. (c) Local chemistry measurements with EDS in the SEM system obtained from the locations marked by black boxes in Fig. 18 (b).
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Fig. 19. EBSD images of cracks along phase boundary on the broad transverse side of the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal), close to the fracture surface: (a) IPF and (b) phase map (The locations marked by white arrows in Fig. 19 (b) are phase boundaries).
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Fig. 20. TEM analyses of an area on the broad transverse side of the broken hydrogen-charged specimen with 10% pre-strain (broken in the weld metal) close to the fracture surface: (a) inhomogeneous dislocation distribution, (b) local high density of dislocations, (c) local high density of dislocations around the deformation twins, and (d) local high density of dislocations around the grain boundary.
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Fig. 21. Local high density of dislocations around the α′ martensite/austenite phase boundary in the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal), located in an area on the broad transverse side close to the fracture surface. The selected-area electron diffraction pattern obtained from the location marked by the red dot shows the microstructure corresponding to α′ martensite.
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Fig. 22. Schematic diagram showing failure processes under different pre-strain.
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Table 1 Welding parameters. Table 1 Welding speed
(A)
(mm/s)
520
6
Arc voltage (V)
Arc length (mm)
Gas flow rate (L/min)
16.5
2.5
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Welding current
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Table 2 Average chemical compositions of the base metal and weld metal (wt.%).
Table 2 Cr
Ni
Mn
Cu
Mo
Si
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S
C
Fe
Base metal
18.11
8.03
0.74
0.54
0.18
0.42
0.088
0.029
0.018
Balance
Weld metal
18.84
9.23
0.98
0.09
0.30
0.68
0.021
0.002
0.012
Balance
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Element
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S0010-938X(18)31605-6 https://doi.org/10.1016/j.corsci.2018.12.018 CS 7812
To appear in: Received date: Revised date: Accepted date:
1 September 2018 8 November 2018 12 December 2018
Please cite this article as: Li X, Gong B, Deng C, Li Y, Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel, Corrosion Science (2018), https://doi.org/10.1016/j.corsci.2018.12.018 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.
Effect of pre-strain on microstructure and hydrogen embrittlement of K-TIG welded austenitic stainless steel
Xiaogang Li1, Baoming Gong1,*, Caiyan Deng1,*, and Yizhe Li1,*
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Tianjin Key Laboratory of Advanced Joining Technology, Department of Materials Science and Engineering,
The Correspondence E-mail: [email protected] (C. Deng); [email protected] (B. Gong); [email protected] (Y. Li)
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*
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Tianjin University, Tianjin 300350, China.
Highlights
Failure location in hydrogen embrittled stainless steel joint depends on pre-strain
Severe strain-induced α′ martensite in the base metal contributes to failure mechanism transition
Tendency of strain-induced α′ martensitic transformation for the weld metal and base metal is
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Abstract
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strongly determined by local chemical homogeneity
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The effects of pre-strain on hydrogen embrittlement of 304L K-TIG welded joint were investigated. As
pre-strain of joint increased, fracture location of hydrogen-charged joint changed from weld metal to base metal. In low prestrained joints, fracture may be initiated at the sites of clustered dislocations due to
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hydrogen in weld metal with a high density of pre-existing dislocations. In high prestrained joints, some regions in base metal suffered severe strain-induced α′ martensitic transformation due to Ni depletion bands and cracking occurred at phase boundaries where may provide more sites for the accumulation of hydrogen atoms and dislocations.
Keywords: A. Stainless steel; B. EPMA; B. TEM; C. Effects of strain; C. Hydrogen embrittlement; C. 1
Welding.
1. Introduction
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Hydrogen may be introduced into metals over time through environmental exposure and corrosion processes. Hydrogen embrittlement (HE), a process associated with the presence and diffusion of hydrogen atoms, often has the deleterious consequence on mechanical properties, ultimately resulting in fracture. Due to the
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excellent HE resistance, mechanical properties, and weldability, 300 series austenitic stainless steels (ASS) have been widely used in gas and oil production and storage structures that are exposed to extreme hydrogen environments, i.e. hydrogen gas storage system for fuel-cell vehicles, boiling water reactor in nuclear power
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plants, and piping systems for petrochemical industry [1–11]. Cold work hardening through pre-straining is
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usually used to improve the mechanical behaviors of ASS prior to use [12]. It was found that pre-straining of
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ASS can introduce both strain hardening [13–15] and transformations in the phase and/or microstructure, i.e.
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α′ martensite formation [13,14], the diffusion and multiplication of dislocations [15], and the formation of deformation twins [16,17]; these changes alter the HE susceptibility of the ASS. It has been reported that the
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tensile elongation rate in a hydrogen environment may increase with increasing pre-strain [18–20]. However, under the same hydrogen charging condition, the hydrogen content increases with the pre-strain, resulting in
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more severe ductility loss in medium-carbon transformation-induced plasticity (TRIP) steel [21], and a similar phenomenon was reported for 304L steel [22]. Since these aforementioned studies did not thoroughly
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explain the HE mechanisms of pre-straining ASS materials, some contradictory results of the effect of hydrogen on plasticity can occasionally be found in the literature [18–21]. In general, it is widely accepted that during the pre-strain process, metastable ASS undergoes
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strain-induced martensite transformation during plastic deformation [13–15], and α′ martensite can work as a ‘highway’ for hydrogen motion [23–25] due to the higher hydrogen diffusion rate of α′ martensite [26–28]. Therefore, it is expected that pre-strained metastable ASS is more sensitive to HE [22]. In addition, compared with the base metal, the ASS weld metal usually suffers more severe HE [29–31] due to its distinct microstructure [32,33] which obviously affected hydrogen transport [34]. Therefore, the study of the HE mechanism in pre-strained ASS and its weld joint is essential and not exhausted to date. Keyhole tungsten 2
inert gas (K-TIG) welding is an efficient technique widely applied to ASS, and it can allow for deeper penetration than does the classical TIG method [35]. However, there is no literature about the effect of pre-strain on microstructure and hydrogen embrittlement of ASS K-TIG welded joints. For pre-strained ASS widely employed in hydrogen-containing environments, the HE mechanism of pre-strained K-TIG welded joints should be critically investigated. Therefore, the primary aim of the current study is to experimentally investigate the influence of
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pre-strain on microstructure and hydrogen embrittlement in AISI 304L K-TIG welded joints. Pre-strain, electrochemical hydrogen charging, and subsequent tensile tests were conducted to investigate the relationship between the microstructure and failure mechanism using scanning electron microscopy (SEM),
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electron probe microanalysis (EPMA), electron backscatter diffraction (ESBD), and transmission electron microscopy (TEM).
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2. Material and methods
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2.1 Materials and welding parameters
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The base metal used in this study was industrially hot-rolled AISI 304L stainless steel plate solution-annealed at 1050 ℃ for 0.5 h and water-quenched with dimensions 300 mm × 150 mm × 10 mm.
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Butt joints were prepared using a K-TIG torch with DC electrode negative polarity and no metal was filled. The welding parameters used in this study are listed in Table 1. The average chemical compositions (in mass
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percentages) of the weld metal and base metal were measured by optical spark emission spectrometry (OES) and are presented in Table 2. An EBSD system (FE-SEM, JEOL, JSM-7100F) was operated at 20 kV to
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observe the microstructures of the K-TIG weld joint. The specimen for EBSD observations was ground to 3000 grit and electropolished using a solution of ethyl alcohol (95 vol%) and perchloric acid (5 vol%) at 30 V for 30 s. The work distance and step size for EBSD operation were 20 mm and 0.1 μm, respectively. The
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raw data of EBSD observations were processed using Orientation Imaging Microscopy software (TSL-OIM). 2.2 Pre-strain After welding, the specimens for pre-strain were cut from the welded plates, as shown in Fig. 1a, and then an overall plastic pre-strain of either 3%, 6%, 10%, 15%, 20%, or 25% was applied to these specimens by using a 100 kN servo-hydraulic universal testing machine (MTS E64) with a strain rate of 2.7×10-4 s-1. 3
The strain was measured by a clip gauge. After the pre-strain process, the microstructures in the weld and the base metal of the pre-strained specimens were detected by using the EBSD system mentioned above. The local chemical compositions in the weld and the base metal of the welded joint after pre-strain were measured using an EPMA instrument (FE-EPMA, JXA-8530 F). The quantitative calibrations of alloying elements (including Ni, Cr, Mn, Mo, Si, and C) were carried out with the corresponding standard samples. Mapping was carried out for the Ni, Cr, Mn, Mo, Si, and C. During pre-strain process, digital image
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correlation (DIC) method was used to measure the full-field strain distribution of welded joint. The DIC system in this study composes of two high-resolution (4872 × 3248 pixels) GE4900 cameras, two lamps as
Correlated Solutions Inc. 2.3 Cathodic hydrogen charging and measurement of hydrogen content
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the light source, calibration grids, a tripod with supporting hardware, and the software package VIC-3D from
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After pre-strain, the tensile specimens were cut from the pre-strained specimens with dimensions as
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shown in Fig. 1b. The longitudinal directions of the tensile specimens were perpendicular to the weld bead.
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The welded bead was placed in the middle of the specimen. One group of the tensile specimens with different pre-strains was cathodically charged at 25 ℃ to introduce hydrogen in 0.5 mol/L H2SO4 + 3 g/L
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NH4SCN solution with a current density of 50 mA/cm2 for 24 h, and the specimens were completely
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immersed during the hydrogen charging, as shown in Fig. 2. The hydrogen content in the specimen was immediately determined by thermal desorption spectroscopy (TDS) analysis using a quadrupole mass
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spectrometer (R-DEC Company, EX0014) with the accuracy of 0.01 wppm. The hydrogen charging conditions was the same for the tensile specimens and the TDS specimens. The test was conducted at a
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heating rate of 100 °C/h in the temperature range from 25 °C to 800 °C. Under each condition, four samples were repeated to ensure the data reliability. The dimensions of the specimen for hydrogen measurement are shown in Fig. 1b.
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2.3 Tensile tests
After hydrogen charging, tensile tests were immediately conducted for hydrogen-charged specimens
with different pre-strains by using a micro-mechanical testing machine (Instron 5848) at room temperature with a strain rate of 2.7×10-4 s-1. The strain was measured by a clip gauge. The hydrogen-free specimens with different pre-strains were also tested as a reference. Specimens in each condition were tested in quadruplicate to ensure data reliability. Two indexes, Iδ and Iψ, are adopted to quantitatively describe HE sensitivity: 4
I δ (%)
0 H 100% 0
(1)
I ψ (%)
0 H 100% 0
(2)
where δ and Ψ are the elongation to failure and the reduction of area of the specimens, respectively. The subscripts 0 and H denote the uncharged and charged specimens, respectively. After the tensile tests, the fracture morphologies were observed by scanning electron microscopy (SEM,
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Hitachi, S4800). The deformation microstructures around the necking zone of the broken samples were observed by using the EBSD system mentioned above and a field-emission transmission electron microscope
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(FE-TEM, JEOL, JEM-2100F) operated at 200 kV. The samples for TEM observation were prepared by twin-jet electro-chemical polishing using an electrolyte consisting of ethyl alcohol (90 vol%) and perchloric acid (10 vol%). The step size of accelerating voltage for TEM operation was 0.05 kV. Energy dispersive
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X-ray spectrometry (EDS) in the SEM system was performed for local chemical composition analysis.
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3.1 Effect of pre-strain on microstructure evolution
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3. Results and discussion
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The representative microstructure of a welded joint was analysed by EBSD mapping, as shown in Fig. 3. The matrix area and heat-affected zone consisted of equiaxial grains and a fusion zone composed of columnar crystals. The heat-affected zone was very narrow (about 200 ± 20 μm), with grain sizes between 10
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and 80 μm. The average grain size of the base metal was slightly larger than that in the heat-affected zone. After pre-strain, deformed microstructures in the welded joint were examined (Figs. 4, 5, and 6). In this
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paper, α' martensite obtained from the pre-strain process is denoted as “prior α' martensite”, to distinguish it from α' martensite formed during the subsequent tensile testing, which is called “dynamic α' martensite” [36]. Regardless of the pre-strain, there was almost no prior α′ martensite in the weld metal, as shown in Fig. 4a
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and 5b. However, the prior α′ martensite content in the base metal dramatically increased with the pre-strain, especially when the pre-strain from 15%, as indicated in Fig. 7. Strain-induced α′ martensitic transformation is affected by grain size, deformation, and local chemical inhomogeneity. Larger grains are known to facilitate strain-induced α′ martensite formation [37] and the volume fraction of α′ martensite increases with increasing deformation [38]. While the average grain size of the weld metal was larger than that in the base metal (Fig. 3) and the weld metal suffered the highest local strains (Fig. 8), almost no prior α′ martensite was 5
observed in the weld metal. Therefore, we ruled out the influences of grain size and deformation on the strain-induced α′ martensite in this study. The nickel equivalent Nieq influences strain-induced α′ martensitic transformation [39]: the lower Nieq was more beneficial for strain-induced α′ martensitic transformation [40]. Nieq can be calculated as follows:
Nieq =Ni+0.65Cr+0.98Mo+1.05Mn+0.35Si+12.6C
(3)
where all elements are represented in mass fraction. The results of EPMA mapping measurements (Fig. 5d
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and 6d) suggested that the base metal exhibited pronounced Ni segregation bands parallel to the rolling direction, resulting in the Ni depletion bands, which corresponded to the band like distribution of α′
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martensite shown in Fig. 6b. In addition, the results of the local chemistry measurements with EPMA (Figs. 5c and 6c) also demonstrated that mass fractions of main elements (Ni, Cr, Mn, Si, Mo, and C), especially Ni, in the prior α′ martensite region of the base metal were lower than that in the weld metal, resulting in the
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lower Nieq of the prior α′ martensite region of the base metal according to Eq. (3), which indicated that there were some regions in the base metal to prone to the transformation of strain-induced α′ martensite.
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Consequently, the different tendency of strain-induced α′ martensitic transformation for the weld metal and
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base metal was strongly determined by local chemical inhomogeneity due to Ni segregation in the base metal
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with the pre-strain, as shown in Fig. 9.
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in this study. Moreover, it was found that the dislocation density in the weld metal increased significantly
3.2 Effect of pre-strain on hydrogen content
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Figure 10 shows the typical TDS curves of the hydrogen-charged specimens with various pre-strain values. The hydrogen desorption peak was clearly observed at around 120 °C; and the peak generally
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increased with the pre-strain, indicating a higher hydrogen content. Figure 11 illustrates the relationship between the hydrogen content and the pre-strains. The temperature range used for calculation of hydrogen content was from 25 °C to 800 °C. At low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the hydrogen
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content slightly increased. However, a large amount of hydrogen started to be introduced into the specimen at 20% and 25% pre-strain. Meanwhile, the amount of hydrogen was a minimum at 0% pre-strain. In this study, the relationship between hydrogen content (C0) in the whole welded joint and pre-strain (εpre) can be fitted exponentially as follows:
C0 17.0405 7.0398exp(0.1127 pre )
(4)
where a similar tendency was observed in the literature in both experimental [41,42] and theoretical [43] 6
studies. Since hydrogen can be trapped by dislocations, increasing the pre-strain can thus result in a much higher dislocation density (Fig. 9) as there is an increased probability of trapping hydrogen atoms. Furthermore, α′ martensite is considered as a ‘highway’ for hydrogen motion in austenitic stainless steels [23–25] because of its higher hydrogen diffusion coefficient [26–28], and thus the prior α’ martensite increases the hydrogen content in 304L steel [22]. The correlation between strain-induced α′ martensite content and strain increment can be expressed as [38]:
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n f α 1 exp 0 1 exp
(5)
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where fα′ is the strain-induced α′ martensite content and ε is the strain. The parameter α describes the formation of a shear band which depends mainly on the stacking-fault energy, β0 is linearly dependent on the probability of martensite embryo formation, and the exponent n is equal to 4.5. Accordingly, if the pre-strain
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is increased further, the prior α′ martensite content in the base metal becomes higher, as shown in Fig. 7 and
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from dislocation and prior α′ martensite simultaneously.
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predicted by Eq. (5). In summary, the hydrogen content increases with increasing pre-strain, which results
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3.3 Effect of deformed microstructures induced by pre-strain and H atoms on tensile properties
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3.3.1 Tensile properties
All hydrogen-free specimens subjected to different pre-strain fractured in the weld metal. The
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hydrogen-charged specimens with 0%, 3%, 6%, 10%, and 15% pre-strain fractured in the weld metal, while those pre-strained to 20% and 25% fractured in the base metal. The engineering stress–true strain curves for
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hydrogen-free and hydrogen-charged specimens are presented in Fig. 12. For all hydrogen-free specimens, the flow stress increased with pre-strain, while the elongation at failure decreased with increasing pre-strain value. More specifically, at the low pre-strain level (0%, 3%, 6%, 10%, and 15% pre-strain), the hydrogen-charged specimens could withstand higher flow stress compared with the hydrogen-free specimens;
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at high pre-strain deformation (20% and 25% pre-strain), lower flow stresses were measured for hydrogen-charged specimens. In addition, as the pre-strain increased, the ultimate tensile strength of hydrogen-free specimens (UTS0) increased monotonically; however, the ultimate tensile strength of hydrogen-charged specimens (UTSH) initially increased until 20% pre-strain, followed by a sharp decline (Fig. 13a). It was found that at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the UTS was almost 7
identical, regardless of the hydrogen content. Once the pre-strain reached 20%, UTSH deviated from UTS0. In addition, similar effects of the pre-strain on the macroscopic plasticity are demonstrated in Figs. 13b and 13c. The HE sensitivity index Iδ reached 14.3 ± 0.35% and 29.0 + 1.34% at pre-strain values of 10% and 25%, respectively. With increasing pre-strain, the dislocation density in the weld metal increases further (Fig. 9) and tangled dislocations may appear. The interaction between dislocations and H atoms trapped in the dislocation
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could hinder dislocation motion. Therefore, the movement of dislocations requires a higher flow stress, leading to strain hardening. In addition, the formation of strain-induced α′ martensite in the base metal may also result in strain hardening. Therefore, for all hydrogen-free and hydrogen-charged specimens with low
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pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain), the strain hardening effect led to a noticeable increase in the flow stress and UTS, and a significant reduction in the elongation at failure. When the pre-strain was greater than 15%, the hydrogen content further increased (Fig. 11). A high hydrogen concentration can
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reduce dislocation formation energies [44] and the cohesive energy of twins and grain boundaries [45,46]. As
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a result, the flow stress and the UTSH values of the hydrogen-charged specimens with 15% and 20%
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pre-strain were lower. When the specimen is plastically deformed, two competitive effects may exist, namely,
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an increment in the flow stress and UTSH due to strain hardening and a decrease in the flow stress and UTSH due to hydrogen. Strain hardening was prevalent at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain)
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while the strength reduction due to the high hydrogen concentration was dominant at high pre-strain (20% and 25% pre-strain).
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3.3.2 Fractographic observations
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Figure 14 illustrates the representative fracture morphologies of the fractured samples determined using
low magnification SEM. It is demonstrated in Fig. 14a that the entire fracture surface of hydrogen-free samples was characterized by ductile dimples with severe plastic localization. In contrast, two distinct
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morphologies can be identified on the fracture surface of the hydrogen-charged specimens: (i) brittle fracture features around the edge zone, resulting from the higher hydrogen concentration at the surface of electrolytically-charged samples and the higher stress intensity factor of the surface defect [47], and (ii) a central portion with and dimples. In addition, Fig. 15 shows that the proportion of the brittle zone increased markedly with increasing pre-strain, which is consistent with the loss of macroscopic ductility indicated by Figs. 13b and 13c. It is worth noting that at a pre-strain of 20% and 25%, the crack initiates and propagates 8
in the base metal, and the brittle zone area increased significantly. Figure 16 shows high magnification SEM images of the fracture surface corresponding to the brittle zones shown in Fig. 14. The entire fracture surface of the hydrogen-free samples was dominated by dimples, indicating only ductile fracture. At the pre-strain of 0% and 3%, the brittle fracture surface exhibited quasi-cleavage fractures. As the pre-strain increased further, the cleavage fracture in the brittle zone became the dominant failure mechanism. Moreover, the brittle zone of the hydrogen-charged specimens with 20% and 25% pre-strain was dominated by the typical cleavage
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fracture mode characterized by flat facet features. 3.4 Mechanisms of failure location transition associated with hydrogen and α′ martensite
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Figures 17 and 18 show the deformed microstructures of the hydrogen-charged specimens with different pre-strains. A large amount of α′ martensite was observed in the hydrogen-charged specimen with 20% pre-strain (broken in the base metal) (Fig. 18b) and almost no α′ martensite was observed in the
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hydrogen-charged specimen with 10% pre-strain (broken in the weld metal) (Fig. 17b). Local chemistry
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measurements with EDS in the SEM system were performed and linked to microstructural analysis, as
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shown in Figs. 17c and 18c. Mass fractions of main elements (Ni, Cr, Mn, Si, and Mo) in the base metal
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(especially Ni in some regions, i.e. region 1 and region 3 marked by black boxes in Fig. 18c) were lower than that in the weld metal, resulting in the lower Nieq of the base metal according to Eq. (3) (consistent with the
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results in section 3.1). Moreover, cracks occurred at the phase boundary in the hydrogen-charged specimen with 20% pre-strain (broken in the base metal) close to the fracture surface, as shown in Fig. 19 Therefore, under the condition that the welded joint experienced high pre-strain (20% and 25% pre-strain), large
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amounts of strain-induced α′ martensitic transformation and phase boundaries in the base metal may lead to
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the failure location transition from weld metal to base metal in hydrogen-charged welded joint with different pre-strains. The influence of the interactions between hydrogen and the microstructures induced by pre-strain on the failure mechanisms will be discussed in more detail. After the low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain) of the joints, more dislocations were
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induced in the weld metal than in the base metal due to localized strain in the weld metal (Fig. 8), while the prior α′ martensite content was low in both the weld and base metal. Several studies have reported that the hydrogen diffusivity was not significantly affected by pre-strain in the ASS (fcc) without strain-induced martensitic transformation [48,49]. It is considered that the binding energy of hydrogen to a dislocation is not much higher than the activation energy of hydrogen diffusion in the lattice in fcc metals [48]. Increasing the number of dislocations may increase the number of trap sites for hydrogen [50]. After hydrogen charging, the 9
distributed dislocations in the weld metal trapped hydrogen and the movement of dislocations in subsequent tensile deformation would accompany with transportation of hydrogen. Moreover, grain boundaries could function as hydrogen traps and the hydrogen atoms could aggregate at the grain boundaries [51]. In the subsequent tensile tests, some crystalline defects in ASS, i.e. twin boundaries and grain boundaries, may hinder the movement of dislocations and lead to the pile-up of dislocations. Meanwhile, hydrogen atoms are very likely to migrate to the twin boundaries [52] during plastic deformation and may accumulate at the grain
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boundaries [51]. Furthermore, owing to the increased transport of hydrogen mediated by dislocations, hydrogen may also move and enriches at the sites of these clustered dislocations. Hydrogen atoms can also induce a new dislocation slip plane [46] and may lower the cohesive energy [51,53]. Therefore, the
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concentrated hydrogen promotes dislocation mobility, leading to macroscopic HE, and fracture may initiate and propagate along the interface where dislocations and hydrogen are accumulated [36,54]. As expected, the dislocation distribution was inhomogeneous (Fig. 20a), and a local high density of dislocations could be
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observed at the deformation twin boundaries, grain boundaries, and other sites (Figs. 20b, 20c, and 20d).
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After the high pre-strain (20% and 25% pre-strain) of the joints, the prior α′ martensite content in the
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base metal sharply increased (Fig. 7), and thus the hydrogen content in the specimens dramatically increased
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due to prior α′ martensite acting as a path for hydrogen transport. However, the prior α′ martensite content was still low in the weld metal (Fig. 5b) due to the high stability of austenite. As a result of the higher
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specific volume of martensite compared to austenite, the transformation of austenite into martensite may result in strain localization around the phase boundaries. After hydrogen-charging, due to the lower solubility
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and higher diffusion coefficient of hydrogen in martensite than austenite, hydrogen atoms collected at the phase boundary between austenite and prior α′ martensite. The deformation in subsequent tensile testing
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mainly occurred in residual austenite due to the higher strength of α′ martensite than austenite. Therefore, incompatible deformation between α′ martensite and the residual austenite resulted in a high local stress concentration at the phase boundary and the residual austenite near phase boundary is preferentially
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transformed into dynamic α′ martensite. Meanwhile, the transformation of the residual austenite into martensite resulted in the release of excessive hydrogen [55], resulting in more dislocations and a higher concentration of hydrogen atoms at the phase boundary. The extreme local enrichment of hydrogen atoms could enhance the dislocation mobility around the phase boundary and potentially resulted in a lower grain boundary energy, cohesive energy [51,53], and even critical material fracture strain [56]. In addition, the additional tensile stress introduced by hydrogen increased linearly with the increasing hydrogen 10
concentration [57,58]. With the application of an external tensile stress, the additional stress induced by hydrogen will be higher than the bonding force (weakened by the presence of hydrogen) and the fracture was initiated at the phase boundaries (Fig. 19b) where the hydrogen atoms and dislocations are accumulated. As expected, a local high density of dislocations was observed around the phase boundary (Fig. 21). Therefore, at low pre-strain levels (0%, 3%, 6%, 10%, and 15% pre-strain) the tensile fracture in the weld metal of the welded joint was a result of trapped hydrogen atoms at high density dislocations; whereas under high
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pre-strain deformation (20% and 25% pre-strain), failure occurred in the base metal of the welded joint and was caused by cracking along the α′ martensite phase boundaries. Fig. 22 shows failure processes and
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mechanisms under different pre-strain.
5. Conclusion
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The effects of pre-strain on failure mechanism of hydrogen embrittlement in 304L K-TIG welded joint
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were investigated in the study. The correlation between fracture location of hydrogen-charged welded joint
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and the microstructural transformations induced by pre-strain was discussed. The results are concluded as
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follows:
(1) While there was almost no prior α′ martensite in the weld metal regardless of the pre-strain of the welded
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joint, the prior α′ martensite content in the base metal dramatically increased with the pre-strain of the welded joint, resulting from the local chemical inhomogeneity. The weld metal was segregation-free,
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whilst there were some regions of low austenite stability in Ni depletion bands of the base metal. (2) As the pre-strain of joint increased, the fracture location of hydrogen-charged joint moved from the weld
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metal to the base metal; it was attributed to the combined influence of the hydrogen concentration and microstructures induced by pre-straining, i.e. dislocations and prior α′ martensite. (3) In the low pre-strained joints (0%, 3%, 6%, 10%, and 15% pre-strain), the weld metal with a high
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density of pre-existing dislocations was prone to local plastic deformation due to the high mobility of dislocations enhanced by the hydrogen. Hydrogen-induced fracture may be initiated at the sites of clustered dislocations and hydrogen in the weld metal. (4) In the high pre-strained joints (20% and 25% pre-strain), the base metal underwent severe strain-induced α′ martensitic transformation, and brittle fracture occurred at α′ martensite/austenite phase boundaries which may suffer more severely local enrichment of hydrogen atoms and dislocations in the base metal. 11
(5) In the low pre-strained joints (0%, 3%, 6%, 10%, and 15% pre-strain), the flow stress and ultimate tensile strength of the hydrogen-charged joints were higher than those of the hydrogen-free joints at the same pre-strain, whereas in the high pre-strained joints (20% and 25% pre-strain), the flow stress and ultimate tensile strength of the hydrogen-charged joints were lower than those of the hydrogen-free joints at the same pre-strain. A competitive relationship between strain hardening and strength reduction due to hydrogen exists in the pre-strained joints: at low pre-strain (0%, 3%, 6%, 10%, and 15% pre-strain),
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strain hardening was prevalent while at high pre-strain (20% and 25% pre-strain) the strength reduction was dominant due to the high hydrogen concentration.
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Acknowledgements
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The research is financially supported by National Key R&D Program of China (Grant No.
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2018YFC0310306) and National Natural Science Foundation of China (NSFC) (Grant No. 51305295).
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References
[1] D. Peckner, I.M. Bernstein, Handbook of Stainless Steels, McGraw Hill, 1977.
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[2] GR. Caskey, Hydrogen Compatibility Handbook for Stainless, Aiken, South Carolina: E I du Pont Nemours, Savannah River Laboratory, 1983.
[3] R. Nishimura, O.M. Alyousif, A new aspect on intergranular hydrogen embrittlement mechanism of solution
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annealed types 304, 316 and 310 austenitic stainless steels, Corros. Sci. 51 (2009) 1894-1900. [4] L.W. Tsay, S.C. Yu, R.T. Huang, Effect of austenite instability on the hydrogen-enhanced crack growth of austenitic stainless steels, Corros. Sci. 49 (2007) 2973-2984.
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[5] Y. Mine, K. Hirashita, M. Matsuda, M. Otsu, K. Takashima, Effect of hydrogen on tensile behaviour of micrometre-sized specimen fabricated from a metastable austenitic stainless steel, Corros. Sci. 53 (2011) 529-533. [6] A.M. Brass, J. Chêne, Hydrogen uptake in 316L stainless steel: consequenceson the tensile propertis, Corros. Sci. 48 (2006) 3222-3242.
[7] L.W. Tsay, J.J. Chen, J.C. Huang, Hydrogen-assisted fatigue crack growth of AISI 316L stainless steel weld, Corros.
A
Sci. 50 (2008) 2973-2980.
[8] M. Koyama, E. Akiyama, T. Sawaguchi, K. Ogawa, I. Kireeva, Y.I. Chumlyakov, K. Tsuzaki, Hydrogen-assisted quasi-cleavage fracture in a single cystalline type 316 austenitic stainless steel, Corros. Sci. 75 (2013) 345-353.
[9] Y. Jiao, W. Zheng, J.R. Kish, Stress corrosion cracking susceptibility of thermally-aged Type 310S stainless steels in supercritical water, Corros. Sci. 135 (2018) 1-11. [10] C.L. Lai, L.W. Tsay, C. Chen, Effect of microstructure on hydrogen embrittlement of various stainless steels, Mater. Sci. Eng. A 584 (2013) 14-20. [11] X. Zhong, S.C. Bali, T. Shoji, Effects of dissolved hydrogen and surface condition on the intergranular stress 12
corrosion cracking initiation and short crack growth behavior of non-sensitized 316 stainless steel in simulated PWR primary water, Corros. Sci. 118 (2017) 143-157. [12] N. Narita, H.K. Brinbaum, On the role of phase transitions in the hydrogen embrittlement of stainless steels, Scr. Metall. 14 (1980) 1355-1358. [13] A.K. De, J.G. Speer, D.K. Matlock, D.C. Murdock, M.C. Mataya, R.J. Comstock Jr, Deformation-induced phase transformation and strain hardening in type 304 austenitic stainless steel, Metall. Mater. Trans. A 37 (2006) 1875-1886. [14] X.F. Fang, W. Dahl, Strain hardening and transformation mechanism of deformation-induced martensite transformation in metastable austenitic stainless steels, Mater. Sci. Eng. A 141 (1991) 189-198. [15] T. Narutani, Effect of deformation-induced martensite transformation on the plastic behaviour of metastable
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austenitic stainless steel, Mater. Trans. Jim 30 (1989) 33-45.
[16] T.H. Lee, C.S. Oh, S.J. Kim, S. Takaki, Deformation twinning in high-nitrogen austenitic stainless steel, Acta Mater. 55 (2007) 3649-3662.
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[17] Y.F. Shen, X.X. Li, X. Sun, Y.D. Wang, L. Zuo, Twinning and martensite in a 304 austenitic stainless steel, Mater. Sci. Eng. A 552 (2012) 514-522.
[18] K. Keiji, I. Shozo, U. Hitoshi, K. Yasushi, T. Takafumi, Hydrogen embrittlement of sensitized SUS 316 steel, J. Soc. Mater. Sci. Jpn. 45 (1996) 1078-1082. transport by dislocations,Mater. Sci. Eng. A 262 (1999) 173-183.
U
[19] F. Lecoester, J. Chene, D. Noel, Hydrogen embrittlement of the Ni-base alloy 600 correlated with hydrogen [20] T. Takasuji, S.Hanada, The effect of pre-deformation on moisture-induced embrittlement of Ni3Al alloys,
N
Intermetallics 5 (1997) 127-135.
[21] W.J. Hui, Y. Li, Y.J. Zhang, C. Zhou, H. Dong, Effect of prestraining on hydrogen absorption and delayed fracture
A
behaviour of a medium-carbon TRIP steel, T. Mater. Heat Treat. 33 (2012) 42-46.
M
[22] Y. Wang, X. Wang, J. Gong, L. Shen, W. Dong, Hydrogen embrittlement of catholically hydrogen-precharged 304L austenitic stainless steel: Effect of plastic pre-strain, Int. J. Hydrogen Energy 39 (2014) 13909-13918. [23] Y. Murakami, Effects of hydrogen on metal fatigue, in: In: Proceedings of International Hydrogen Energy
ED
Development Forum, Fukuoka, 2007. [24] M. Martin, S. Weber, C. Izawa, S. Wagner, A. Pundt, W. Theisen, Influence of machining-induced martensite on hydrogen-assisted fracture of AISI type 304 austenitic stainless steel, Int. J. Hydrogen Energy 36 (2011)
PT
11195-11206.
[25] T. Kanezaki, C. Narazaki, Y. Mine, S. Matsuoka, Y. Murakami, Effects of hydrogen on fatigue crack growth behavior of austenitic stainless steels, Int. J. Hydrogen Energy 33 (2008) 2604-2619.
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[26] G. Han, J. He, S. Fukuyama, K. Yokogawa, Effect of strain-induced martensite on hydrogen environment embrittlement of sensitized austenitic stainless steels at low temperatures, Acta Mater. 46 (1998) 4559-4570. [27] TP. Perng, CJ. Altstetter, Comparison of hydrogen gas embrittlement of austenitic and ferritic stainless steels, Metall. Mater. Trans. A 18 (1987) 123-134. [28] L. Zhang, M. Imade, B. An, M. Wen, T. Iijima, S. Fukuyama, K. Yokogawa, Internal reversible hydrogen
A
embrittlement of austenitic stainless steels based on type 316 at low temperatures, ISIJ Inter. 99 (2013) 294-301.
[29] T. Michler, Y. Lee, R.P. Gangloff, J. Naumann, Influence of macro segregation on hydrogen environment embrittlement of SUS 316L stainless steel, Int. J. Hydrogen Energy 34 (2009) 3201-3209.
[30] C.M. Younes, A.M. Steele, J.A. Nicholson, C.J. Barnett, Influence of hydrogen content on the tensile properties and fracture of austenitic stainless steel welds, Int. J. Hydrogen Energy 38 (2013) 4864-4876. [31] T. Michler, Toughness and hydrogen compatibility of austenitic stainless steel welds at cryogenic temperatures, Int. J. Hydrogen Energy 32 (2007) 4081-4088. [32] H.F. Jackson, K.A. Nibur, C.S. Marchi, J.D. Puskar, B.P. Somerday, Hydrogen-assisted crack propagation in 13
304L/308L and 21Cr-6Ni-9Mn/308L austenitic stainless steel fusion welds, Corros. Sci. 60 (2012) 136-144. [33] J. Naumann, T. Michler, Hydrogen environment embrittlement of orbital welded austenitic stainless steels at -50°C, Int. J. Hydrogen Energy 34 (2009) 6478-6483. [34] V. Olden, A. Alvaro, O.M. Akselsen, Hydrogen diffusion and hydrogen influenced critical stress intensity in an API X70 pipeline steel welded joint-Experiments and FE simulations, Int. J. Hydrogen Energy 37 (2012) 11474-11486. [35] M. Lohse, U. Fuessel, H. Schuster, J. Friedel, M. Schnick, Keyhole welding with CF-TIG (cathode focussed GTA), Weld. World 57 (2013) 735-741. [36] L. Zhang, Z. Li, J. Zheng, Y. Zhao, P. Xu, C. Zhou, X. Li, Effect of strain-induced martensite on hydrogen embrittlement of austenitic stainless steels investigated by combined tension and hydrogen release methods, Int. J.
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Hydrogen Energy 38 (2013) 8208-8214.
[37] C. Yoo, Y. Park, Y. Jung, Y. Lee, Effect of grain size on transformation-induced plasticity in an ultrafine-grained metastable austenitic steel, Scr. Mater. 59 (2008) 71-74.
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[38] G.B. Olson, M. Cohen, Kinetics of strain-induced martensite nucleation, Metall. Mater. Trans. A 6 (1975) 791-795. [39] T. Hirayama, M. Ogirima, Influence of chemical composition on martensitic transformation in Fe-Cr-Ni stainless steel, J. Jpn. Inst. Met. Mater. 34 (1970) 507-510.
[40] L. Zhang, M. Wen, M. Imade, S. Fukuyama, K. Yokogawa, Effect of nickel equivalent on hydrogen gas embrittlement of austenitic stainless steels based on type 316 at low temperatures, Acta Mater. 56 (2008)
U
3414-3421.
[41] M. Nagumo, K. Ohta, H. Saitoh, Deformation induced defects in iron revealed by thermal desorption spectroscopy
N
of tritium, Scr. Mater. 40 (1999) 313-319.
[42] M. Nagumo, K. Takai, N. Okuda, Nature of hydrogen trapping sites in steels induced by plastic deformation, J.
A
Alloys Compd. 293 (1999) 310-316.
M
[43] S. Wang, S. Ohnuki, N. Hashimoto, K. Chiba, Hydrogen effects on tensile property of pure iron with deformed surface, Mater. Sci. Eng. A 560 (2013) 332-338.
[44] Caskey, G.R. Jr, Fractography of hydrogen-embrittled iron-chromium-nickel alloys, in: In: For Presentation at the
ED
13th Annual Technical Meeting of the International Metallographic Society, Brighton England, 1980. [45] V.G. Gavriljuk, V.N. Bugaev, Y.N. Petrov, A.V. Tarasenko, B.Z. Yanchitski, Hydrogen-induced equilibrium vacancies in fcc iron-base alloys, Scr. Mater. 34 (1996) 903-907.
PT
[46] T.D. Lee, I.M. Bernstein, Effects of hydrogen on dislocation morphology in spheroidized steel, Acta Metall. Mater. 39 (1991) 363-372.
[47] G.E. Dieter, Mechanical Metallurgy, third ed., McGraw Hill, London, 1986.
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[48] Y. Mine, C. Narazaki, K. Murakami, S. Matsuoka, Y. Murakami, Hydrogen transport in solution-treated and pre-strained austenitic stainless steels and its role in hydrogen-enhanced fatigue crack growth, Int. J. Hydrogen Energy 34 (2009) 1097-1107.
[49] T.P. Perng, C.J. Altstetter, Effects of deformation on hydrogen permeation in austenitic stainless steels, Acta Metall. 34 (1986) 1771-1781.
A
[50] X. Li, Y. Wang, P. Zhang, B. Li, X. Song, J. Chen, Effect of pre-strain on hydrogen embrittlement of high strength steels, Mater. Sci. Eng. A 616 (2014) 116-122.
[51] M. Yamaguchi, K. Ebihara, M. Itakura, T. Kadoyoshi, T. Suzudo, H. Kaburaki, First-principles study on the grain boundary embrittlement of metals by solute segregation: Part II. Metal (Fe, Al, Cu)-hydrogen (H) systems, Metall. Mater. Trans. A 42 (2011), 330-339. [52] H. Ji, I. Park, S. Lee, Y. Lee, The effect of pre-strain on hydrogen embrittlement in 310S stainless steel, J. Alloy. Compd. 598 (2014) 205-212. [53] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation II. 14
Experimental evidence and consequences, Acta Mater. 55 (2007) 5139-5148. [54] P. Novak, R. Yuan, B.P. Somerday, P. Sofronis, R.O. Ritchie, A statistical, physical-based, micro-mechanical model of hydrogen-induced intergranular fracture in steel, J. Mech. Phys. Solids. 58 (2010) 206-226. [55] Y. Mine, Z. Horita, Y. Murakami, Effect of hydrogen on martensite formation in austenitic stainless steels in high-pressure torsion, Acta Mater. 57 (2009) 2993-3002. [56] Y.A. Du, L. Ismer, J. Rogal,T. Hickel, J. Neugebauer, R. Drautz, First-principles study on the interaction of H interstitials with grain boundaries in alpha- and gamma-Fe, Phys. Rev. B 84 (2011) 667-673. [57] R. Wang, Effects of hydrogen on the fracture toughness of a X70 pipeline steel, Corros. Sci. 51 (2009) 2803-2810. [58] T. Zhang, W.Y. Chu, K.W. Gao, L.J. Qiao, Study of correlation between hydrogen-induced stress and hydrogen
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embrittlement, Mater. Sci. Eng. A 347 (2003) 291-299.
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Figure Captions
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Fig. 1. Dimensions and sampling location of specimens for: (a) pre-strain and (b) tensile tests and hydrogen content measurements (in mm).
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Fig. 2. Schematic of the electrolytic hydrogen charging system.
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Fig. 3. EBSD images of the fusion zone, heat-affected zone, and base metal of the welded joint: (a) inverse pole figure (IPF) map and (b) misorientation distribution.
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Fig. 4. EBSD analyses the welded joint after 10% pre-strain: (a) weld metal and (b) base metal.
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Fig. 5. Microstructure and local chemical composition analysis of the weld metal of the welded joint after 20% pre-strain: (a) detected region, (b) IPF and phase map measured by EBSD, (c) local chemistry measurements with EPMA obtained from the locations marked by black boxes in Fig. 5 (b), and (d) the distributions of different alloying elements (Ni, Cr, Mn, Si, Mo, and C) of the region shown in Fig. 5 (b), obtained by EPMA mapping.
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Fig. 6. Microstructure and local chemical composition analysis of the base metal of the welded joint after 20% pre-strain: (a) detected region, (b) IPF and phase map measured by EBSD, (c) local chemistry measurements with EPMA obtained from the locations marked by black boxes in Fig. 6 (b), and (d) the distributions of different alloying elements (Ni, Cr, Mn, Si, Mo, and C) of the region shown in Fig. 6 (b), obtained by EPMA mapping.
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Fig. 7. Average prior α′ martensite content (from EBSD measurements) of six randomly selected regions in base metals of welded joints with different pre-strains.
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Fig. 8. Map of strain distribution of the hydrogen-free specimen with 15% pre-strain obtained from digital image correlation measurement.
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Fig. 9. TEM analyses of dislocation in weld metals of welded joints after different pre-strain: (a) 6% pre-strain, (b) 10% pre-strain, and (c) 15% pre-strain.
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Fig. 10. TDS curves of hydrogen-charged specimens at various pre-strain values.
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Fig. 11. Hydrogen content determined at various pre-strain values, as measured by TDS analysis using a quadrupole mass spectrometer.
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Fig. 12. Representative engineering stress–true strain curves of hydrogen-free and hydrogen-charged specimens at different pre-strain values: (a) 6% pre-strain, (b) 15% pre-strain, (c) 20% pre-strain, and (d) 25% pre-strain.
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Fig. 13. Tensile properties of hydrogen-free and hydrogen-charged specimens at different pre-strain values: (a) ultimate tensile strength UTS, (b) elongation and hydrogen embrittlement sensitivity index Iδ, and (c) reduction of area and hydrogen embrittlement sensitivity index IΨ. Specimens in each condition were tested in quadruplicate.
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Fig. 14. Low magnification SEM images of morphologies of fractured specimens.
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Fig. 15. Areal proportion of brittle zone vs. the pre-strain, as observed in the fracture morphologies.
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Fig. 16. High magnification SEM images of morphologies of fractured specimens.
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IP T SC R U N A M ED PT CC E A Fig. 17. EBSD analyses of the broken hydrogen-charged specimen with 10% pre-strain (broken in the weld metal): (a) IPF and (b) phase map. (c) Local chemistry measurements with EDS in the SEM system obtained from the locations marked by black boxes in Fig. 17 (b). 34
IP T SC R U N A M ED PT CC E A Fig. 18. EBSD analyses of the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal): (a) IPF and (b) phase map. (c) Local chemistry measurements with EDS in the SEM system obtained from the locations marked by black boxes in Fig. 18 (b).
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Fig. 19. EBSD images of cracks along phase boundary on the broad transverse side of the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal), close to the fracture surface: (a) IPF and (b) phase map (The locations marked by white arrows in Fig. 19 (b) are phase boundaries).
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Fig. 20. TEM analyses of an area on the broad transverse side of the broken hydrogen-charged specimen with 10% pre-strain (broken in the weld metal) close to the fracture surface: (a) inhomogeneous dislocation distribution, (b) local high density of dislocations, (c) local high density of dislocations around the deformation twins, and (d) local high density of dislocations around the grain boundary.
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Fig. 21. Local high density of dislocations around the α′ martensite/austenite phase boundary in the broken hydrogen-charged specimen with 20% pre-strain (broken in the base metal), located in an area on the broad transverse side close to the fracture surface. The selected-area electron diffraction pattern obtained from the location marked by the red dot shows the microstructure corresponding to α′ martensite.
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Fig. 22. Schematic diagram showing failure processes under different pre-strain.
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Table 1 Welding parameters. Table 1 Welding speed
(A)
(mm/s)
520
6
Arc voltage (V)
Arc length (mm)
Gas flow rate (L/min)
16.5
2.5
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Welding current
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Table 2 Average chemical compositions of the base metal and weld metal (wt.%).
Table 2 Cr
Ni
Mn
Cu
Mo
Si
N
S
C
Fe
Base metal
18.11
8.03
0.74
0.54
0.18
0.42
0.088
0.029
0.018
Balance
Weld metal
18.84
9.23
0.98
0.09
0.30
0.68
0.021
0.002
0.012
Balance
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Element
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