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Ductility and fatigue properties of low nickel content type 316L austenitic stainless steel after gaseous thermal pre-charging with hydrogen
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Ductility and fatigue properties of low nickel content type 316L austenitic stainless steel after gaseous thermal pre-charging with hydrogen Thanh Tuan Nguyen, Jaeyeong Park, Seung Hoon Nahm, Naehyung Tak, Un Bong Baek* Center for Energy Materials Measurement, Division of Industrial Metrology, Korea Research Institute of Standard and Science (KRISS), 267 Gajeong-Ro, Yuseoung-Gu, Daejeon 34113, Republic of Korea
highlights Effect of hydrogen on the mechanical properties of low Ni content type 316L was evaluated. HE was more severe in gaseous hydrogen environment than pre-charged hydrogen. Loss of fatigue strength was only observed in the plastic component of fatigue life. Fatigue limit was not degraded in the presence of hydrogen environment. Pronounced strain induced-martensite transformation occurred at high strain amplitude regime.
article info
abstract
Article history:
The susceptibility of low nickel content type 316L austenitic stainless steel to hydrogen was
Received 7 June 2019
quantified using low strain rate tensile tests and strain-controlled low-cycle fatigue life
Received in revised form
measurements. Both tests were performed under air condition after charging with
25 August 2019
high-pressure 10-MPa hydrogen gas at 300 C for eight days. No significant influence of
Accepted 28 August 2019
hydrogen was recognized in 0.2% proof stress, but the strain at fracture and reduction area
Available online 28 September 2019
was decreased significantly in both hydrogen pre-charged and in gaseous hydrogen conditions compared to companion tests conducted in air. The decrease of fatigue life in the
Keywords:
high strain amplitude region was related to a significant decrease in the plastic component
316L
while the effect of hydrogen on the elastic component was negligible. Highly localized
Hydrogen embrittlement
deformation and a pronounced martensite transformation occurred near the site of the
Strain-controlled fatigue
fracture surface in the high strain amplitude regime, resulting in the early formation of
Crack initiation
abundant micro-surface cracks in this regime of the hydrogen pre-charged samples.
Striation
© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Hydrogen has become a choice as an alternative energy source for the replacement of traditional fuels, and is expected as a
primary source for future energy generation to address the challenges of global warming. However, the limited number of hydrogen fueling stations and the relatively high cost of storing methods have been restricting the development of
* Corresponding author. Korea Research Institute of Standard and Science, Republic of Korea. E-mail address: [email protected] (U.B. Baek). https://doi.org/10.1016/j.ijhydene.2019.08.233 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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hydrogen-fueled vehicles. In recent years, research and development on the practical industrial use and storage of hydrogen energy have been accelerated to promote the commercialization of fuel cell vehicles. Furthermore, it is generally accepted that hydrogen degrades the mechanical properties of nearly every structural component, a phenomenon known as hydrogen embrittlement (HE) [1e4]. To ensure the long-term safety and the reliability of hydrogen facilities, suitable structural materials are important in the commercialization of such technologies. In recent decades, several criteria and regulations have been developed for determining the hydrogen compatibility of materials [5e9]. In general, mechanical testing provides the fundamental requirements for material selection and for the design of structural components in the presence of hydrogen. Typically, the mechanical properties in the presence of hydrogen are compared with those in air to understand the differences in fracture behaviors. The testing of the susceptibility of materials to HE can be performed in a hydrogen gas environment or using hydrogen pre-charged material. The main difference in the two testing methods is the distribution of hydrogen within the specimen, resulting in a difference in the fracture mechanisms [10e12]. The susceptibility of materials to HE often manifests in the decrease of the tensile strength, ductility, and fatigue life, or in the acceleration of fatigue crack growth properties. Considering structural materials, group 300 series austenitic stainless steels are typically less susceptible to HE for the application in hydrogen facilities, due to their excellent HE resistance and based on their satisfactory historical service experience. It has been reported that the grades of susceptibility of austenitic stainless steels strongly depend of various factors such as environmental conditions (gaseous conditions and temperature), alloying components, and loading conditions [13e16]. Among these, nickel content and temperature conditions are two critically important parameters that control the HE effects. Experimentally, it was found that materials with lower Ni content are more susceptible to HE than alloys with greater Ni content [8,15,17e19]. This can be attributed to the difference in the crystal structure of materials and to the stability of the austenite phase. Nevertheless, the temperature factor can be related to the interaction between hydrogen and the metal surface, and to the solubility and diffusivity of hydrogen. The maximum susceptibility of austenitic stainless steel has been reported to be in the range of 200e250 K for austenitic stainless steels. During the testing of different material classes, the individual temperature dependence needs to be considered [15,16,19e21]. Pressure dependency of HE should be also established for the testing under the hydrogen atmosphere. It is reported that HE increased with increasing gas pressure and tend to reach a saturation at a specific pressure [22,23]. No further reduction in relative reduction area is occurred when exceeding the saturation value. It is also reported that increased hydrogen pressure enhanced a larger fraction of intergranular fracture and accelerating fatigue crack growth [24]. Several mechanisms are proposed to describe HE degradation; however, there are two generally accepted mechanisms. The first explains the premature failure due to hydrogen by the increased dislocation mobility at a high hydrogen concentration location (HELP) [25,26] while the
second attributes it to a weakening of interatomic bond at the grain boundaries, which results in brittle fracture, known as hydrogen-enhanced decohesion (HEDE) [27]. In several cases, these mechanisms are likely to be combined, where the dominant mechanism depends on the fracture mode. This distinction can be discussed based on the fractographic observation of the fracture surface. So far, important studies have been conducted to analyze hydrogen susceptibility of materials exposed to hydrogen gas and the effect of hydrogen on particular materials under a specific condition have been investigated, especially the degradation of fatigue properties. In this study, the effects of hydrogen on the mechanical properties of a 316L type austenitic stainless steel were quantified based on low strain rate tensile test results and strain-controlled low-cycle fatigue (LCF) life measurements. Both tests were conducted under air condition with specimens charged with high-pressure hydrogen gas at elevated temperature. The strain at the fracture and reduction area decreased significantly under both hydrogen pre-charged and gaseous hydrogen conditions compared with the corresponding tests conducted in air. A considerable effect of hydrogen pre-charging on the fatigue life properties was only observed in the high strain amplitude regime, while there was a negligible difference in fatigue life in the low strain deformation region. Fractographic analyses were performed to determine the effect of hydrogen on the failure mechanism during the fatigue life test (see Fig. 1).
Experimental procedures The material used in this study was a hot-rolled plate of commercial type 316L austenitic stainless steel with thickness of 25 mm. Its nickel content was approximately 10 wt% according to the specification of the American Society for Testing and Materials (ASTM) standard A240 [28], which was
Fig. 1 e Geometry and dimensions of (a) SSRT test specimen and (b) low-cycle strain-controlled fatigue life test specimen.
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Table 1 e Chemical compositions of type 316L austenitic stainless steel using OES analysis. Mater.
C
Si
Mn
P
S
Cr
Ni
Mo
N
Present ASTM
0.017 0.03
0.48 0.75
0.85 2.0
0.03 0.045
0.001 0.03
16.17 16.0e18.0
10.08 10.0e14.0
2.07 2.0e3.0
0.071 0.1
lower than the 12 mass% lower limit of the Ni content specified by the Japanese Industrial Standards (JIS) for the type SUS316L steel. The chemical compositions were specified by optical emission spectroscopy (OES) analysis, as shown in Table 1 [29]. Slow strain rate tensile (SSRT) tests were performed using smooth specimens under air condition for both hydrogen pre-charged and uncharged specimens. Fig. 1a shows the shape and dimensions of the test specimens. The SSRT tests were performed using a servo-hydraulic testing equipment. A constant crosshead speed of 0.0004 mm/s was applied in the test under hydrogen gas while a constant speed of 0.002 mm/s was used for the pre-charged and uncharged test specimens. Since electrochemical charging can generate a high hydrogen concentration near the surface that can cause some damages to the specimen’s surface, thermal pre-charging in hydrogen gas was used in this study. The precharged specimens were exposed to high-pressure hydrogen gas of 10 MPa and temperature of 300 C for eight days. The charging time was determined by applying the instruction in CSA/CSCM 1e2014 [30]. Before the beginning of the tests, the pre-charged specimens were stored at 80 C to retain the hydrogen concentration in the specimen. In addition, an SSRT test in gaseous hydrogen of 10 MPa was also performed, to compare the results with those under the hydrogen precharging condition. The tests were performed in a special pressure vessel that can withstand hydrogen pressures of up to 100 MPa. Strain was measured using an extensometer inside the pressure vessel while load was measured using a load cell located outside the pressure vessel. The LCF tests were performed at room temperature for both the uncharged and hydrogen pre-charged specimens. The tests were carried out in a strain-controlled mode. Round specimens with a gauge portion of diameter of 6.35 mm and gauge length of 18 mm were used, as shown in Fig. 1b. In SSRT tests the specimen was charged under the conditions mentioned above. An extensometer with a gauge length of 12.5 mm, strain ratio of 1, and strain rate of 0.005 s1 was used in the LCF tests. All tests were performed in ambient air at room temperature. After the tests, the reduction of area (RA) was measured by a travel micrometer across the fracture location. The fracture surfaces of the rupture specimens were analyzed using scanning electron microscopy (SEM). For the phase analysis, an electron backscatter diffraction (EBSD) system was employed with a scanning step size of 1.6 mm at a voltage of 25 kV. The samples for the EBSD measurements were prepared by mechanical polishing using #2400 emery paper and subsequence by electrochemical polishing. To measure the hydrogen content after the SSRT and fatigue life test, the specimens were immediately stored in liquid nitrogen after removing from the testing equipment. Thermal desorption spectroscopy (TDS) was performed at heating rates in the range of 100e600 C/h. A TDS sample was placed in a TDS tube
furnace and the gas was analyzed at 3 min intervals using HE as carrier gas. A standard gas mixture of H2 and He (balance) was used for the calibration.
Results and discussion SSRT test results Fig. 2 shows the stressestrain curves for the uncharged and hydrogen pre-charged specimens and for specimens in 10 MPa hydrogen gas. In each condition, the test was repeated three times to obtain an indication of the reproducibility of the results. The first observation showed that hydrogen had negligible effect on the 0.2% proof strength, and only slightly increased in the hydrogen pre-charged specimen. This indicates that the material retains its stability under the elastic regime regardless of hydrogen conditions, both in the case of hydrogen pre-charged specimen and in the case of treatment in 10 MPa hydrogen gas. The tensile strength in the hydrogen pre-charged specimens were comparable to that in the uncharged specimen; however, it dropped dramatically in specimens in 10 MPa hydrogen gas. Both samples showed premature failure compared to the uncharged specimen results. The average values of strain at fracture were at least 79.4% in the test of the uncharged specimens while hydrogen pre-charged specimens and those in gaseous hydrogen were only 49.6% and 32.5%, respectively. A similar tendency of loss of the strain at fracture could be seen in the results of the reduction area. The average RA was approximately 80.9% when it was tested for the uncharged specimen; however, it significantly decreased to 69.1 for the test with the precharged specimen and decreased below 37.7% in 10-MPa H2 gas. It can also be seen that the loss of ductility (strain at
Fig. 2 e Comparison of nominal stressestrain curves for the uncharged and under hydrogen-charged and gaseous hydrogen treated conditions.
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Table 2 e SSRT test results of type 316L austenitic stainless steel under the three tested conditions. Testing Condition Uncharged Pre-charged 10-MPa H2
Yield Strength (MPa)
Tensile Strength (MPa)
Strain at Fracture (%)
Reduction of Area (%)
233.2 241.7 220.8
591.4 575.1 517.9
79.4 49.6 32.5
80.9 69.1 37.7
fracture and RA) is more severe than the loss of tensile strength. The SSRT results are summarized in Table 2. Figs. 3 and 4 show low- and high-magnification images for the comparison of the fracture surfaces and the longitudinal sections after the SSRT test. The low-magnification SEM images show macroscopic observation of the fracture surface and side-longitudinal view near the fracture location while the high-magnification images were taken from the central area of each rupture surface. As shown in Fig. 3a, a typical cup-cone fracture was observed for the uncharged specimens, which includes the purely normal and shear stress fracture region. The examined state of the hydrogen pre-charged specimen did not reveal noticeable signs of cup and cone fracture and the fracture surface was mixed brittle-ductile, with a predominantly transgranular fracture zone and fine dimples, as shown in Fig. 3b. Several internal and secondary surface cracks, forming from the outer surface, were also found at the rupture location of the hydrogen pre-charged specimen. The test result of the specimen in gaseous hydrogen clearly shows brittle fracture, covered by quasicleavage transgranular fracture surface. The minor necking shown Fig. 3c, indicates a minimal plastic deformation and that the propagation of the crack was nearly perpendicular to the applied tensile force. Compared to the hydrogen precharged specimens, a high density of rough surface cracks was also observed on the outer surface of the specimen treated with gaseous hydrogen. Microscopic observation of the at the center of the fracture surface are presented in Fig. 4. The differences in the ductility loss mentioned above indicate that the grade of susceptibility by hydrogen precharged condition was lower compared with the degraded level by gaseous hydrogen condition. These can be attributed to the significant difference in the hydrogen distributions in the specimen. As the thermal hydrogen charging was carried out by exposing the specimen to high-pressure hydrogen gas at an elevated temperature, an equilibrium hydrogen distribution was established over the entire specimen. Nevertheless, as the hydrogen diffusivity of austenitic stainless steels is in the order of only 1016 m2/s [17], hydrogen is not able to distribute uniformly in the whole specimen during the SSRT test. Thus, the distribution of hydrogen adsorbed near the surface layer was significantly higher than that in the interior. Takakuwa et al. reported that, for the group 300 series austenitic stainless steel, hydrogen diffusivity only occurs in a very thin surface layer that is less than 4 mm thick [12]. Thus, the micro-cracks formed on the outer surface of the gaseous hydrogen specimens can be attributed to the stress/strain incompatibility between the outer surface layer containing hydrogen and the inner material without hydrogen content. The hydrogen-enhanced surface crack occurred during the uniform plastic deformation before reaching ultimate tensile strength, thereby leading to a considerable decrease in the
fracture elongation and the reduction area. It is argued that gaseous hydrogen treatment has a stronger effect on the tensile strength, fracture elongation, and reduction area compared with hydrogen pre-charging. The observed surface cracks in the specimen due to the presence of hydrogen are also consistent with the features of the HELP mechanism reported by other researchers as well [26,27]. It should be noted that the difference in fracture mechanisms due to external and internal hydrogen can only be seen in low Nieq content austenitic stainless steels. For high Nieq content materials, hydrogen cannot enter the inner section of the specimen due to its extremely low diffusivity and the high stability of the austenite crystalline structure [10,11,20] (see Table 3).
Fatigue life test results Strainelife curve, ε-N Fig. 5 shows the relationship between the total strain amplitude and the number of cycles to failure (strainelife curves, ε-N) of the uncharged and hydrogen pre-charged specimens. As the fatigue life tests in this study are performed under the strain-controlled mode, the total strain amplitude can be divided into two distinct components: elastic and plastic srain. A well-known relationship between the strain amplitude and the fatigue life can be approximated by the equation proposed by Coffin and Manson [31,32], as εa ¼ εea þ εpa ¼
b c s0f 2Nf þ ε0f 2Nf E
(1)
where εa is the total strain amplitude, E is Young’s modulus, and Nf is the number of cycles. The relationship between the elastic strain component, εea and the stable stress amplitude can be written as εea ¼ sa/E while the plastic strain amplitude εpa is a measurement of the width of the stressestrain hysteresis loop near the half of the fatigue life. The parameters of fatigue strength coefficient (s0 f), fatigue strength exponent (b), ductility coefficient (εf), and fatigue ductility exponent (c), are considered as material properties. The obtained data of the uncharged and hydrogen-charged specimens shown in Fig. 5 are indicated by blue and red symbols, respectively. Although, there is a certain amount of statistical scatter in the fatigue life test result shown in Fig. 5, a clear tendency can be seen that hydrogen decreases dramatically the fatigue-life in
Table 3 e Cyclic parameters in the CoffineMason equation (Eq. (1)) obtained from the strain-controlled fatigue life test. Material Uncharged Hydrogen Pre-charged
s0 f (MPa)
b
εf
c
1198.2 717.8
0.145 0.0864
0.132 0.044
0.392 0.285
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Fig. 3 e Macroscopic fracture surfaces of type 316L austenitic stainless steel: (a) for uncharged specimen, (b) for hydrogen pre-charged specimen, and (c) for specimen in 10-MPa H2.
the high total strain amplitude regime, (εa 0.50%), while there is no noticeable difference in the fatigue lives in the low strain amplitude regime (εa 0.35%). In addition, it is wellknown that elastic strain is dominant in the low-amplitude regime while in the high strain amplitude level the plastic strains are large compared to the elastic strains. Thus, the
decrease of fatigue life in the LCF regime is related to a significant decrease in the plastic component while the effect of hydrogen on the elastic component was not significant. Furthermore, the strain as a function of the number of cycles to failure tends to decrease slowly at low strain amplitudes, which corresponds to the longest live regime. It can be
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Fig. 5 e Strain amplitude as a function of the number of cycles to failure of type 316L for the uncharged and hydrogen-charged conditions.
Fig. 4 e Comparison of the microscopic images at the center location of type 316L austenitic stainless steel: (a) for uncharged specimen, (b) for hydrogen pre-charged specimen, and (c) for specimen in 10-MPa H2.
concluded that the fatigue limit sw is apparently not degraded in the presence of hydrogen environment. Similar results for the fatigue life properties have also been reported for other alloys of group 300 series austenitic stainless steel, either in the case of hydrogen pre-charged specimens or in specimens treated in gaseous hydrogen [33e36]. Table 2 summarizes the cyclic parameters expressed in Eq. (1) obtained from the fatigue life results. The fatigue strength exponent s0 f and the fatigue ductility exponent εf are theoretically expected to be similar to the true fracture stress and fracture strain from the SSRT test, respectively. The inferior s0 f and εf values of the hydrogen pre-charged sample are in good agreement with the degradation of ductility and strength values in the SSRT results. Fig. 6 shows a comparison of the flow curves of the peak stress between the uncharged and hydrogen pre-charged conditions of the hysteresis loops at the specific strain amplitudes of 1.2% and 0.35%. Both materials exhibit a pronounced cyclic work hardening behavior in the high strain amplitude regime, while a cyclic work softening behavior can be observed in the low strain amplitude regime. In addition, as shown in Fig. 6, compared with the uncharged specimen, the hydrogen pre-charged specimen exhibits a slight increase in the stabilized stress amplitude, which is consistent with the hydrogen-induced yield strength in the SSRT test results. The hydrogen content of the selected specimens at each strain amplitude level was measured by TDS analysis following the fatigue life test for each sample, as shown in Fig. 7. The maximum desorption rate was found to be at approximately 560 C for all specimens, and the average value of the measured hydrogen content was 35.1 wppm and the standard deviation is only ±2.1 wppm. Since the outgassing of hydrogen from the specimen due to the low diffusivity under ambient air conditions during the fatigue tests was reported to be negligible in type 316L steels [12,17], the measured content could be considered the hydrogen content at the initial stage following the fatigue life test. However, the desorption rate of the low-strain amplitude specimen is a slightly higher than that of the high-strain amplitude specimens. It is attributed to
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Fig. 6 e Flow curves of peak stress of hysteresis loops at total strain amplitudes of 1.2% and 0.35% for the as-received and hydrogen-charged conditions.
the higher strain-induced martensite fraction when tested at the high-strain amplitude level, resulting in an increasing amount of hydrogen egress from the specimen during the fatigue life test. Additional errors in thermal desorption analysis results should be included since only one measurement was conducted for each strain amplitude level. Further study is needed to clarify this difference.
Fractography analysis Fig. 8 shows a comparison of typical macroscopic images of the fracture surfaces after the fatigue life test for both hydrogen pre-charged and uncharged specimens. All fracture surfaces can be clearly divided into two separate regions: fatigue and overload fracture areas. On a macroscopic scale, the focused observations were concentrated on the location near the crack initiation sites as shown in Fig. 9. Although certain characteristic of the fracture surface could disappear during the compression mode, formation of striations marks was visually observed from the crack initiation site on the uncharged specimen. Initiation sites of the hydrogen pre-charged specimen surfaces were completely
Fig. 7 e Hydrogen thermal desorption profile of the hydrogen pre-charged samples after the fatigue life test.
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flat and decorated with facets. In general, the fracture process during the fatigue life test can be divided into three stages that can be distinguished based on the analysis of the appearance of the fracture surface. Crack initiation was typically formed from the specimen surface at the first stage and propagated stably in the direction of thickness perpendicular to the loading direction during the second stage. The fracture process in the first and second stages is generally related to the normal fatigue crack growth mechanism while the third stage of the fracture process involves shearing and ductile tearing after the crack has reached a sufficient length. Fig. 10 shows a comparison of SEM images of the longitudinal cross sections near the edge of the fracture surface. No micro-surface cracks can be observed in the uncharged specimen and the gauge length body is stretched. Nevertheless, a substantial surface crack formation can be seen on the outer surface near the fracture edge of the hydrogen-charged specimens and a transgranular internal crack can also be observed in inward direction on the fracture surface. A depth of the surface crack is in the range of approximately 8e10 mm. A detailed microscopic observation was performed at the central location of the fatigue fracture surface shown in Fig. 11. Striation can clearly be identified all over the fracture surface of the uncharged samples, which is a typical feature of fatigue fracture surfaces of ductile materials. This indicates the repeated blunting and resharpening of the crack tip during fatigue cycle loading. Each striation mark can represent one load cycle while every load cycle can result in the formation of a striation mark. On the other hand, the fatigue fracture surface of the hydrogencharged specimens are macroscopically flat, and the fatigue fracture surface is completely covered with quasi-cleavage transgranular fractures. A number of striation marks can be observed at the fatigue fracture region of the low strain amplitude regime of the specimen (εa ¼ 0.35%); however, the fraction area where striations appear is less distinct than those in the uncharged specimen. Fig. 12 shows images of the EBSD measurement performed at the crack initiation site of the hydrogen precharged samples after the fatigue life test. Measurement was also performed on the ruptured SSRT specimen of the hydrogen pre-charged specimen. As shown in Fig. 12a, a highly localized deformation clearly appeared along the fracture surface site in the high strain amplitude regime of the specimen. The fraction of stress-induced martensite (SIM) transformation along the boundaries of the deformed grains is also visually noticeable. By contrast, less evidence of SIM was observed in the low strain amplitude regime of the sample. In the case of SSRT samples, the phase transformation from austenite to martensite occurred extensively due to the severe damage from the deformation at the necking location, as shown in Fig. 12c. Although the EBSD measurement has not been performed at crack initiation site on the uncharged specimen, it is expected that the SIM fraction is higher in the uncharged sample than that in the hydrogen pre-charged specimens due to the abrupt failure. However, the magnitude of the plastic strain within the slip band deformation is higher, which results in an earlier crack formation and an accelerated fatigue crack growth in the hydrogen pre-charged specimens [16].
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Fig. 8 e Comparison of typical macroscopic images of fracture surfaces between as-received and hydrogen pre-charged specimen after the fatigue life test for (a) at strain amplitude of 0.35%, (b) at strain amplitude of 0.75%, and (c) at strain amplitude of 1.2%.
Fatigue failure mechanism The degradation of mechanical properties in austenitic stainless steel exposed to hydrogen is typically attributed to the phase transformation from austenite to martensite. In this aspect, Ni has been reported to be important in the stability of the austenite phase and in suppressing strain-induced
martensitic transformation. The higher Ni content stabilizes the austenitic matrix, suppress martensitic transformation, and increases the stacking fault energy [14,15]. The relative reduction area of the present material is nearly identical to the value of type 316 with nickel content of 10.12 wt%, but they are relatively lower than the value of type 316L, which contains
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Fig. 9 e Origin of fatigue crack, in which crack initiated at a strain amplitude of 0.35%: (a) uncharged specimen and (b) hydrogen pre-charged specimen.
12.22 wt% [37]. Michler et al. reported that the loss of RA decreased with decreasing Ni content from 9.07 wt% to 12.75 wt% and RA is inversely proportional to the straininduced martensite transformation fraction [38]. In that way, HE could be negligible when the Ni content is above 11.5 wt% [21]. Other studies in the literature have also presented an empirical parameter to relate the HE resistance to the Ni content and other alloying elements, known as Nickelequivalent, Nieq [8,9,15e19,39e41]. Also, based on the quantitative variation of Nieq from 23.1 to 35.7, the fracture morphologies of material with the presence of hydrogen can be
comprehensively classified into five mechanisms. The Nieq value of the material in this study, calculated using Takakuwa’s equation [11], is 25.11, which is lower than the saturation value of 27, at which the effect of hydrogen environment is eliminated for austenitic stainless steel. As shown by the comparison in Fig. 11a, b, in the high strain amplitude regime, the fraction of the deformed structures is higher and the SIM transformation appearing along the deformed grain on the initiation site is more pronounced than those in the low strain amplitude regime specimen. As the SIM transformation zone can form a hydrogen diffusion pathway, hydrogen atoms are
Fig. 10 e Images of the outer surface of the rupture specimen after the fatigue life test at the strain amplitude of 1.2%: (a) uncharged specimen and (b) hydrogen pre-charged specimen.
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Fig. 11 e Comparison of microscopic fatigue fracture surfaces between uncharged specimen and hydrogen pre-charged specimen for (a), (a) at strain amplitude of 0.35%, (b) at strain amplitude of 0.75%, and (c) at strain amplitude of 1.2%.
able to accumulate before the crack tip region during cyclic loading. This can result in a localized planar deformation and stress concentration, which results in the rapid formation of micro-cracks after the beginning of the fatigue life test. Once a micro-crack is formed, a stress concentration zone appears at the crack tip, which accumulates the high SIM transformation. As mentioned previously, due to the appearance of several micro-cracks from the beginning of the test process (as shown in Fig. 10), the failure mechanism in the hydrogen precharged specimen can be considered as a competition of crack growth processes under cyclic load among the crack initiations. Once a crack propagation at one surface crack dominates the others, the propagation at this crack proceeds
continuously while the others are nearly terminated. According to the hydrogen-assisted surface crack growth (HASCG) model [6,42], the cracks are generated successively and remain in a sharp shape because the localized slip deformation in the vicinity of the crack due to hydrogen, results in an accelerated crack growth rate and induces premature failure during fatigue life test. By contrast, the failure in the uncharged specimen is completely dominated by the propagation of a crack. As localized slip deformation did not occur, a large crack blunting resulted in clearly visible striation, as shown in Fig. 10. It should be noted, that the cyclic stressestrain behavior for hydrogen pre-charged specimen was similar to those of the uncharged specimens; however, in
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Fig. 12 e EBSD measurement results of the fracture edge: (a) after fatigue life test, εa ¼ 1.2%; (b) after fatigue life test, εa ¼ 0.35; and (c) after SSRT test of the hydrogen pre-charged specimen. In each figure, the inverse pole figure image is on the left side and the phase map is located on the right side.
the high strain amplitude regime a failure can occur sooner in the pre-charged specimens. Nagging a reported that, compared with hydrogen-induced surface crack initiation sites in the hydrogen pre-charged samples, the fatigue crack initiation mechanism for the uncharged specimens could be attributed to the irreversible plastic strain-induced surface roughening [43].
Conclusions An SSRT and strain-controlled fatigue life test of low nickel content type 316L austenitic stainless steels was performed and the results of uncharged and hydrogen pre-charged
specimens were compared, from which the following conclusions can be drawn: 1. The SSRT results clearly reveal that hydrogen changes the fracture mechanism of type 316L steel and results in significant loss in the reduction area and strain at the fracture. Treatment with gaseous hydrogen has a stronger effect than pre-charging specimens with hydrogen. 2. The reduction of fatigue life in the high strain amplitude region is related to a significant decrease of the plastic component while the effect of hydrogen on the elastic component was negligible. The fatigue limit of the precharged specimens was consistently equivalent with the results of the uncharged sample.
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3. The difference in fracture appearances is mainly related to the accelerated crack initiation due to hydrogen during the strain-controlled LCF life test. The highly localized deformation and pronounced martensite transformation occur near the side of the fracture surface in the high strain amplitude regime, resulting in an early formation of abundant surface cracks.
[13]
[14]
Acknowledgements
[15]
This research was supported by Development of Reliability Measurement & Standard Technology for Hydrogen Fueling Station Funded by Korea Research Institute of Standards and Science (KRISS-2019-GP2019-0012).
[16]
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[18] [1] Hirth JP. Effects of hydrogen on the properties of iron and steel. Metall Trans A 1980;11:861e90. [2] Louthan MR, Caskey GR, Donovan JA, Rawl DE. Hydrogen embrittlement of metals. Mater Sci Eng A 1972;10:357e68. [3] Gangloff RP, Somerday BP. Gaseous hydrogen embrittlement of materials in energy technologies: mechanisms, modelling and future developments. Elsevier; 2012. [4] Caskey GR. Hydrogen compatibility handbook for stainless steels No DP-1643. Du Pont de Nemours (EI) and Co. Aiken, SC (USA): Savannah River Lab; 1983. [5] ASTM G142-98. Standard test method for determination of susceptibility of metals to embrittlement in hydrogen containing environments at high pressure, high temperature, or both, 4. West Conshohocken, PA: ASTM International; 2014. [6] Matsuoka S, Yamabe J, Matsunaga H. Criteria for determining hydrogen compatibility and the mechanisms for hydrogenassisted, surface crack growth in austenitic stainless steels. Eng Fract Mech 2016;153:103e27. [7] San Marchi C, Somerday BP, Nibur KA. Development of methods for evaluating hydrogen compatibility and suitability. Int J Hydrogen Energy 2014;39:20434e9. [8] Takaki S, Nanba S, Imakawa K, Macadre A, Yamabe J, Matsunaga H, Matsuoka S. Determination of hydrogen compatibility for solution-treated austenitic stainless steels based on a newly proposed nickel-equivalent equation. Int J Hydrogen Energy 2016;41:15095e100. [9] Wada Y. Material selection and safety variation of steel tanks for hydrogen stations (PVP2014-28677). In: Proceedings of PVP-2014: ASME pressure vessels and piping division conference (California, 2014). New York: American Society of Mechanical Engineers; 2014. [10] San Marchi C, Michler T, Nibur KA, Somerday BP. On the physical differences between tensile testing of type 304 and 316 austenitic stainless steels with internal hydrogen and in external hydrogen. Int J Hydrogen Energy 2010;35:9736e45. [11] Takakuwa O, Yamabe J, Matsunaga H, Furuya Y, Matsuoka S. Comprehensive understanding of ductility loss mechanisms in various steels with external and internal hydrogen. Metall Mater Trans A 2017;48:5717e32. [12] Takakuwa O, Yamabe J, Matsunaga H, Furuya Y, Matsuoka S. Recent progress on interpretation of tensile ductility loss for various austenitic stainless steels with external and internal hydrogen (PVP2017-65671). In: Proceedings of PVP-2017:
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ASME pressure vessels and piping division conference (Hawaii, 2017). New York: American Society of Mechanical Engineers; 2017. Zhang L, Wen M, Imade M, Fukuyama S, Yokogawa K. Effect of nickel equivalent on hydrogen gas embrittlement of austenitic stainless steels based on type 316 at low temperatures. Acta Mater 2008;56:3414e21. Wada Y. Effect of cycle frequency on fatigue crack propagation behavior for steels in hydrogen storage (PVP2013-97485). In: Proceedings of PVP-2013: ASME pressure vessels and piping division conference (Paris, 2013). New York: American Society of Mechanical Engineers; 2013. San Marchi C, Somerday BP, Tang X, Schiroky GH. Effects of alloy composition and strain hardening on tensile fracture of hydrogen-precharged type 316 stainless steels. Int J Hydrogen Energy 2008;33:889e904. Han G, He J, Fukuyama S, Yokogawa K. Effect of straininduced martensite on hydrogen environment embrittlement of sensitized austenitic stainless steels at low temperatures. Acta Mater 1998;46:4559e70. Mine Y, Horita Z, Murakami Y. Effect of hydrogen on martensite formation in austenitic stainless steels in highpressure torsion. Acta Mater 2009;57:2993e3002. Mine Y, Narazaki C, Murakami K, Matsuoka S, Murakami Y. Hydrogen transport in solution-treated and pre-strained austenitic stainless steels and its role in hydrogen-enhanced fatigue crack growth. Int J Hydrogen Energy 2009;34:1097e107. Michler T, Naumann J. Hydrogen environment embrittlement of austenitic stainless steels at low temperatures. Int J Hydrogen Energy 2008;33:2111e22. Sun S, Han G, Vaodee S, Fukuyama S, Yokogawa K. Tensile behaviour of type 304 austenitic stainless steels in hydrogen atmosphere at low temperatures. Mater Sci Technol 2001;17:302e8. Michler T, Yukhimchuk AA, Naumann J. Hydrogen environment embrittlement testing at low temperatures and high pressures. Corros Sci 2008;50:3519e26. Xu K, Rana M. Tensile and fracture properties of carbon and low alloy steels in high pressure hydrogen. In: Proceedings of the 2008 international hydrogen conference; 2009. p. 349e56. Matsuoka S, Takakuwa O, Okazaki S, Yoshikawa M, Yamabe J, Matsunaga H. Peculiar temperature dependence of hydrogen-enhanced fatigue crack growth of low-carbon steel in gaseous hydrogen. Scr Mater 2018;154:101e5. Ogawa Y, Birenis D, Matsunaga H, Takakuwa O, Yamabe J, Prytz O, Thogersen A. The role of intergranular fracture on hydrogen-assisted fatigue crack propagation in pure iron at a low stress intensity range. Mater Sci Eng A 2018;733:316e28. Birnbaum HK, Sofronis P. Hydrogen-enhanced localized plasticity-a mechanism for hydrogen-related fracture. Mater Sci Eng A 1994;176:191e202. Ulmer DG, Altstetter CJ. Hydrogen-induced strain localization and failure of austenitic stainless steels at high hydrogen concentrations. Acta Metall 1991;39:1237e48. Nagao A, Smith CD, Dadfarnia M, Sofronis P, Robertson IM. The role of hydrogen in hydrogen embrittlement fracture of lath martensitic steel. Acta Mater 2012;60:5182e9. ASTM. A240/A240M-17, standard specification for chromium and chromium-nickel stainless steel plate, sheet, and strip for pressure vessels and for general applications. West Conshohocken, PA Std: ASTM International; 2017. Nguyen TT, Nahm SH, Ryu KS, Baek UB. Ductility and fatigue strength loss of a hydrogen-charged 316L austenitic stainless steel (PVP2019-93180). In: Proceedings of PVP-2019: ASME pressure vessels and piping division conference (San Antonio, 2019). New York: American Society of Mechanical Engineers; 2019.
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[30] ANSI/CSA CHMC 1-2014 Test methods for evaluating material compatibility in compressed hydrogen applications e Metals. CSA group; 2014. [31] Coffin LF. A study of the effects of cyclic thermal stresses on a ductile metal. Trans Am Soc Mech Eng 1954;76:931e50. [32] Manson SS. Behavior of materials under conditions of thermal stress. National Advisory Committee for Aeronautics; 1953. NACA TN-2933. [33] Murakami Y, Kanezaki T, Mine Y. Hydrogen effect against hydrogen embrittlement. Metall Mater Trans A 2010;41:2548e62. [34] Gibbs PJ, San Marchi C, Nibur KA, Tang X. Comparison of internal and external hydrogen on fatigue-life of austenitic stainless steels (PVP2016-63563). In: Proceedings of PVP-2016: ASME pressure vessels and piping division conference (Vancouver, 2016). New York: American Society of Mechanical Engineers; 2016. [35] Michler T, Naumann J, Sattler E. Influence of high-pressure gaseous hydrogen on S-N fatigue in two austenitic stainless steels. Int J Fatigue 2013;51:1e7. [36] Michler T, Naumann J, Weber S, Martin M, Pargeter RS-. N fatigue properties of a stable high aluminum austenitic stainless steel for hydrogen applications. Int J Hydrogen Energy 2013;38:9935e41. [37] Itoga H, Matsunaga H, Yamabe J, Matsuoka S. Effects of external and internal hydrogen on tensile properties of austenitic stainless steels containing additive elements. (PVP2015-45740. In: Proceedings of PVP-2015: ASME pressure vessels and piping division conference (Boston, 2015). New York: American Society of Mechanical Engineers; 2015.
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Available online at www.sciencedirect.com
ScienceDirect journal homepage: www.elsevier.com/locate/he
Ductility and fatigue properties of low nickel content type 316L austenitic stainless steel after gaseous thermal pre-charging with hydrogen Thanh Tuan Nguyen, Jaeyeong Park, Seung Hoon Nahm, Naehyung Tak, Un Bong Baek* Center for Energy Materials Measurement, Division of Industrial Metrology, Korea Research Institute of Standard and Science (KRISS), 267 Gajeong-Ro, Yuseoung-Gu, Daejeon 34113, Republic of Korea
highlights Effect of hydrogen on the mechanical properties of low Ni content type 316L was evaluated. HE was more severe in gaseous hydrogen environment than pre-charged hydrogen. Loss of fatigue strength was only observed in the plastic component of fatigue life. Fatigue limit was not degraded in the presence of hydrogen environment. Pronounced strain induced-martensite transformation occurred at high strain amplitude regime.
article info
abstract
Article history:
The susceptibility of low nickel content type 316L austenitic stainless steel to hydrogen was
Received 7 June 2019
quantified using low strain rate tensile tests and strain-controlled low-cycle fatigue life
Received in revised form
measurements. Both tests were performed under air condition after charging with
25 August 2019
high-pressure 10-MPa hydrogen gas at 300 C for eight days. No significant influence of
Accepted 28 August 2019
hydrogen was recognized in 0.2% proof stress, but the strain at fracture and reduction area
Available online 28 September 2019
was decreased significantly in both hydrogen pre-charged and in gaseous hydrogen conditions compared to companion tests conducted in air. The decrease of fatigue life in the
Keywords:
high strain amplitude region was related to a significant decrease in the plastic component
316L
while the effect of hydrogen on the elastic component was negligible. Highly localized
Hydrogen embrittlement
deformation and a pronounced martensite transformation occurred near the site of the
Strain-controlled fatigue
fracture surface in the high strain amplitude regime, resulting in the early formation of
Crack initiation
abundant micro-surface cracks in this regime of the hydrogen pre-charged samples.
Striation
© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Hydrogen has become a choice as an alternative energy source for the replacement of traditional fuels, and is expected as a
primary source for future energy generation to address the challenges of global warming. However, the limited number of hydrogen fueling stations and the relatively high cost of storing methods have been restricting the development of
* Corresponding author. Korea Research Institute of Standard and Science, Republic of Korea. E-mail address: [email protected] (U.B. Baek). https://doi.org/10.1016/j.ijhydene.2019.08.233 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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hydrogen-fueled vehicles. In recent years, research and development on the practical industrial use and storage of hydrogen energy have been accelerated to promote the commercialization of fuel cell vehicles. Furthermore, it is generally accepted that hydrogen degrades the mechanical properties of nearly every structural component, a phenomenon known as hydrogen embrittlement (HE) [1e4]. To ensure the long-term safety and the reliability of hydrogen facilities, suitable structural materials are important in the commercialization of such technologies. In recent decades, several criteria and regulations have been developed for determining the hydrogen compatibility of materials [5e9]. In general, mechanical testing provides the fundamental requirements for material selection and for the design of structural components in the presence of hydrogen. Typically, the mechanical properties in the presence of hydrogen are compared with those in air to understand the differences in fracture behaviors. The testing of the susceptibility of materials to HE can be performed in a hydrogen gas environment or using hydrogen pre-charged material. The main difference in the two testing methods is the distribution of hydrogen within the specimen, resulting in a difference in the fracture mechanisms [10e12]. The susceptibility of materials to HE often manifests in the decrease of the tensile strength, ductility, and fatigue life, or in the acceleration of fatigue crack growth properties. Considering structural materials, group 300 series austenitic stainless steels are typically less susceptible to HE for the application in hydrogen facilities, due to their excellent HE resistance and based on their satisfactory historical service experience. It has been reported that the grades of susceptibility of austenitic stainless steels strongly depend of various factors such as environmental conditions (gaseous conditions and temperature), alloying components, and loading conditions [13e16]. Among these, nickel content and temperature conditions are two critically important parameters that control the HE effects. Experimentally, it was found that materials with lower Ni content are more susceptible to HE than alloys with greater Ni content [8,15,17e19]. This can be attributed to the difference in the crystal structure of materials and to the stability of the austenite phase. Nevertheless, the temperature factor can be related to the interaction between hydrogen and the metal surface, and to the solubility and diffusivity of hydrogen. The maximum susceptibility of austenitic stainless steel has been reported to be in the range of 200e250 K for austenitic stainless steels. During the testing of different material classes, the individual temperature dependence needs to be considered [15,16,19e21]. Pressure dependency of HE should be also established for the testing under the hydrogen atmosphere. It is reported that HE increased with increasing gas pressure and tend to reach a saturation at a specific pressure [22,23]. No further reduction in relative reduction area is occurred when exceeding the saturation value. It is also reported that increased hydrogen pressure enhanced a larger fraction of intergranular fracture and accelerating fatigue crack growth [24]. Several mechanisms are proposed to describe HE degradation; however, there are two generally accepted mechanisms. The first explains the premature failure due to hydrogen by the increased dislocation mobility at a high hydrogen concentration location (HELP) [25,26] while the
second attributes it to a weakening of interatomic bond at the grain boundaries, which results in brittle fracture, known as hydrogen-enhanced decohesion (HEDE) [27]. In several cases, these mechanisms are likely to be combined, where the dominant mechanism depends on the fracture mode. This distinction can be discussed based on the fractographic observation of the fracture surface. So far, important studies have been conducted to analyze hydrogen susceptibility of materials exposed to hydrogen gas and the effect of hydrogen on particular materials under a specific condition have been investigated, especially the degradation of fatigue properties. In this study, the effects of hydrogen on the mechanical properties of a 316L type austenitic stainless steel were quantified based on low strain rate tensile test results and strain-controlled low-cycle fatigue (LCF) life measurements. Both tests were conducted under air condition with specimens charged with high-pressure hydrogen gas at elevated temperature. The strain at the fracture and reduction area decreased significantly under both hydrogen pre-charged and gaseous hydrogen conditions compared with the corresponding tests conducted in air. A considerable effect of hydrogen pre-charging on the fatigue life properties was only observed in the high strain amplitude regime, while there was a negligible difference in fatigue life in the low strain deformation region. Fractographic analyses were performed to determine the effect of hydrogen on the failure mechanism during the fatigue life test (see Fig. 1).
Experimental procedures The material used in this study was a hot-rolled plate of commercial type 316L austenitic stainless steel with thickness of 25 mm. Its nickel content was approximately 10 wt% according to the specification of the American Society for Testing and Materials (ASTM) standard A240 [28], which was
Fig. 1 e Geometry and dimensions of (a) SSRT test specimen and (b) low-cycle strain-controlled fatigue life test specimen.
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Table 1 e Chemical compositions of type 316L austenitic stainless steel using OES analysis. Mater.
C
Si
Mn
P
S
Cr
Ni
Mo
N
Present ASTM
0.017 0.03
0.48 0.75
0.85 2.0
0.03 0.045
0.001 0.03
16.17 16.0e18.0
10.08 10.0e14.0
2.07 2.0e3.0
0.071 0.1
lower than the 12 mass% lower limit of the Ni content specified by the Japanese Industrial Standards (JIS) for the type SUS316L steel. The chemical compositions were specified by optical emission spectroscopy (OES) analysis, as shown in Table 1 [29]. Slow strain rate tensile (SSRT) tests were performed using smooth specimens under air condition for both hydrogen pre-charged and uncharged specimens. Fig. 1a shows the shape and dimensions of the test specimens. The SSRT tests were performed using a servo-hydraulic testing equipment. A constant crosshead speed of 0.0004 mm/s was applied in the test under hydrogen gas while a constant speed of 0.002 mm/s was used for the pre-charged and uncharged test specimens. Since electrochemical charging can generate a high hydrogen concentration near the surface that can cause some damages to the specimen’s surface, thermal pre-charging in hydrogen gas was used in this study. The precharged specimens were exposed to high-pressure hydrogen gas of 10 MPa and temperature of 300 C for eight days. The charging time was determined by applying the instruction in CSA/CSCM 1e2014 [30]. Before the beginning of the tests, the pre-charged specimens were stored at 80 C to retain the hydrogen concentration in the specimen. In addition, an SSRT test in gaseous hydrogen of 10 MPa was also performed, to compare the results with those under the hydrogen precharging condition. The tests were performed in a special pressure vessel that can withstand hydrogen pressures of up to 100 MPa. Strain was measured using an extensometer inside the pressure vessel while load was measured using a load cell located outside the pressure vessel. The LCF tests were performed at room temperature for both the uncharged and hydrogen pre-charged specimens. The tests were carried out in a strain-controlled mode. Round specimens with a gauge portion of diameter of 6.35 mm and gauge length of 18 mm were used, as shown in Fig. 1b. In SSRT tests the specimen was charged under the conditions mentioned above. An extensometer with a gauge length of 12.5 mm, strain ratio of 1, and strain rate of 0.005 s1 was used in the LCF tests. All tests were performed in ambient air at room temperature. After the tests, the reduction of area (RA) was measured by a travel micrometer across the fracture location. The fracture surfaces of the rupture specimens were analyzed using scanning electron microscopy (SEM). For the phase analysis, an electron backscatter diffraction (EBSD) system was employed with a scanning step size of 1.6 mm at a voltage of 25 kV. The samples for the EBSD measurements were prepared by mechanical polishing using #2400 emery paper and subsequence by electrochemical polishing. To measure the hydrogen content after the SSRT and fatigue life test, the specimens were immediately stored in liquid nitrogen after removing from the testing equipment. Thermal desorption spectroscopy (TDS) was performed at heating rates in the range of 100e600 C/h. A TDS sample was placed in a TDS tube
furnace and the gas was analyzed at 3 min intervals using HE as carrier gas. A standard gas mixture of H2 and He (balance) was used for the calibration.
Results and discussion SSRT test results Fig. 2 shows the stressestrain curves for the uncharged and hydrogen pre-charged specimens and for specimens in 10 MPa hydrogen gas. In each condition, the test was repeated three times to obtain an indication of the reproducibility of the results. The first observation showed that hydrogen had negligible effect on the 0.2% proof strength, and only slightly increased in the hydrogen pre-charged specimen. This indicates that the material retains its stability under the elastic regime regardless of hydrogen conditions, both in the case of hydrogen pre-charged specimen and in the case of treatment in 10 MPa hydrogen gas. The tensile strength in the hydrogen pre-charged specimens were comparable to that in the uncharged specimen; however, it dropped dramatically in specimens in 10 MPa hydrogen gas. Both samples showed premature failure compared to the uncharged specimen results. The average values of strain at fracture were at least 79.4% in the test of the uncharged specimens while hydrogen pre-charged specimens and those in gaseous hydrogen were only 49.6% and 32.5%, respectively. A similar tendency of loss of the strain at fracture could be seen in the results of the reduction area. The average RA was approximately 80.9% when it was tested for the uncharged specimen; however, it significantly decreased to 69.1 for the test with the precharged specimen and decreased below 37.7% in 10-MPa H2 gas. It can also be seen that the loss of ductility (strain at
Fig. 2 e Comparison of nominal stressestrain curves for the uncharged and under hydrogen-charged and gaseous hydrogen treated conditions.
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Table 2 e SSRT test results of type 316L austenitic stainless steel under the three tested conditions. Testing Condition Uncharged Pre-charged 10-MPa H2
Yield Strength (MPa)
Tensile Strength (MPa)
Strain at Fracture (%)
Reduction of Area (%)
233.2 241.7 220.8
591.4 575.1 517.9
79.4 49.6 32.5
80.9 69.1 37.7
fracture and RA) is more severe than the loss of tensile strength. The SSRT results are summarized in Table 2. Figs. 3 and 4 show low- and high-magnification images for the comparison of the fracture surfaces and the longitudinal sections after the SSRT test. The low-magnification SEM images show macroscopic observation of the fracture surface and side-longitudinal view near the fracture location while the high-magnification images were taken from the central area of each rupture surface. As shown in Fig. 3a, a typical cup-cone fracture was observed for the uncharged specimens, which includes the purely normal and shear stress fracture region. The examined state of the hydrogen pre-charged specimen did not reveal noticeable signs of cup and cone fracture and the fracture surface was mixed brittle-ductile, with a predominantly transgranular fracture zone and fine dimples, as shown in Fig. 3b. Several internal and secondary surface cracks, forming from the outer surface, were also found at the rupture location of the hydrogen pre-charged specimen. The test result of the specimen in gaseous hydrogen clearly shows brittle fracture, covered by quasicleavage transgranular fracture surface. The minor necking shown Fig. 3c, indicates a minimal plastic deformation and that the propagation of the crack was nearly perpendicular to the applied tensile force. Compared to the hydrogen precharged specimens, a high density of rough surface cracks was also observed on the outer surface of the specimen treated with gaseous hydrogen. Microscopic observation of the at the center of the fracture surface are presented in Fig. 4. The differences in the ductility loss mentioned above indicate that the grade of susceptibility by hydrogen precharged condition was lower compared with the degraded level by gaseous hydrogen condition. These can be attributed to the significant difference in the hydrogen distributions in the specimen. As the thermal hydrogen charging was carried out by exposing the specimen to high-pressure hydrogen gas at an elevated temperature, an equilibrium hydrogen distribution was established over the entire specimen. Nevertheless, as the hydrogen diffusivity of austenitic stainless steels is in the order of only 1016 m2/s [17], hydrogen is not able to distribute uniformly in the whole specimen during the SSRT test. Thus, the distribution of hydrogen adsorbed near the surface layer was significantly higher than that in the interior. Takakuwa et al. reported that, for the group 300 series austenitic stainless steel, hydrogen diffusivity only occurs in a very thin surface layer that is less than 4 mm thick [12]. Thus, the micro-cracks formed on the outer surface of the gaseous hydrogen specimens can be attributed to the stress/strain incompatibility between the outer surface layer containing hydrogen and the inner material without hydrogen content. The hydrogen-enhanced surface crack occurred during the uniform plastic deformation before reaching ultimate tensile strength, thereby leading to a considerable decrease in the
fracture elongation and the reduction area. It is argued that gaseous hydrogen treatment has a stronger effect on the tensile strength, fracture elongation, and reduction area compared with hydrogen pre-charging. The observed surface cracks in the specimen due to the presence of hydrogen are also consistent with the features of the HELP mechanism reported by other researchers as well [26,27]. It should be noted that the difference in fracture mechanisms due to external and internal hydrogen can only be seen in low Nieq content austenitic stainless steels. For high Nieq content materials, hydrogen cannot enter the inner section of the specimen due to its extremely low diffusivity and the high stability of the austenite crystalline structure [10,11,20] (see Table 3).
Fatigue life test results Strainelife curve, ε-N Fig. 5 shows the relationship between the total strain amplitude and the number of cycles to failure (strainelife curves, ε-N) of the uncharged and hydrogen pre-charged specimens. As the fatigue life tests in this study are performed under the strain-controlled mode, the total strain amplitude can be divided into two distinct components: elastic and plastic srain. A well-known relationship between the strain amplitude and the fatigue life can be approximated by the equation proposed by Coffin and Manson [31,32], as εa ¼ εea þ εpa ¼
b c s0f 2Nf þ ε0f 2Nf E
(1)
where εa is the total strain amplitude, E is Young’s modulus, and Nf is the number of cycles. The relationship between the elastic strain component, εea and the stable stress amplitude can be written as εea ¼ sa/E while the plastic strain amplitude εpa is a measurement of the width of the stressestrain hysteresis loop near the half of the fatigue life. The parameters of fatigue strength coefficient (s0 f), fatigue strength exponent (b), ductility coefficient (εf), and fatigue ductility exponent (c), are considered as material properties. The obtained data of the uncharged and hydrogen-charged specimens shown in Fig. 5 are indicated by blue and red symbols, respectively. Although, there is a certain amount of statistical scatter in the fatigue life test result shown in Fig. 5, a clear tendency can be seen that hydrogen decreases dramatically the fatigue-life in
Table 3 e Cyclic parameters in the CoffineMason equation (Eq. (1)) obtained from the strain-controlled fatigue life test. Material Uncharged Hydrogen Pre-charged
s0 f (MPa)
b
εf
c
1198.2 717.8
0.145 0.0864
0.132 0.044
0.392 0.285
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Fig. 3 e Macroscopic fracture surfaces of type 316L austenitic stainless steel: (a) for uncharged specimen, (b) for hydrogen pre-charged specimen, and (c) for specimen in 10-MPa H2.
the high total strain amplitude regime, (εa 0.50%), while there is no noticeable difference in the fatigue lives in the low strain amplitude regime (εa 0.35%). In addition, it is wellknown that elastic strain is dominant in the low-amplitude regime while in the high strain amplitude level the plastic strains are large compared to the elastic strains. Thus, the
decrease of fatigue life in the LCF regime is related to a significant decrease in the plastic component while the effect of hydrogen on the elastic component was not significant. Furthermore, the strain as a function of the number of cycles to failure tends to decrease slowly at low strain amplitudes, which corresponds to the longest live regime. It can be
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Fig. 5 e Strain amplitude as a function of the number of cycles to failure of type 316L for the uncharged and hydrogen-charged conditions.
Fig. 4 e Comparison of the microscopic images at the center location of type 316L austenitic stainless steel: (a) for uncharged specimen, (b) for hydrogen pre-charged specimen, and (c) for specimen in 10-MPa H2.
concluded that the fatigue limit sw is apparently not degraded in the presence of hydrogen environment. Similar results for the fatigue life properties have also been reported for other alloys of group 300 series austenitic stainless steel, either in the case of hydrogen pre-charged specimens or in specimens treated in gaseous hydrogen [33e36]. Table 2 summarizes the cyclic parameters expressed in Eq. (1) obtained from the fatigue life results. The fatigue strength exponent s0 f and the fatigue ductility exponent εf are theoretically expected to be similar to the true fracture stress and fracture strain from the SSRT test, respectively. The inferior s0 f and εf values of the hydrogen pre-charged sample are in good agreement with the degradation of ductility and strength values in the SSRT results. Fig. 6 shows a comparison of the flow curves of the peak stress between the uncharged and hydrogen pre-charged conditions of the hysteresis loops at the specific strain amplitudes of 1.2% and 0.35%. Both materials exhibit a pronounced cyclic work hardening behavior in the high strain amplitude regime, while a cyclic work softening behavior can be observed in the low strain amplitude regime. In addition, as shown in Fig. 6, compared with the uncharged specimen, the hydrogen pre-charged specimen exhibits a slight increase in the stabilized stress amplitude, which is consistent with the hydrogen-induced yield strength in the SSRT test results. The hydrogen content of the selected specimens at each strain amplitude level was measured by TDS analysis following the fatigue life test for each sample, as shown in Fig. 7. The maximum desorption rate was found to be at approximately 560 C for all specimens, and the average value of the measured hydrogen content was 35.1 wppm and the standard deviation is only ±2.1 wppm. Since the outgassing of hydrogen from the specimen due to the low diffusivity under ambient air conditions during the fatigue tests was reported to be negligible in type 316L steels [12,17], the measured content could be considered the hydrogen content at the initial stage following the fatigue life test. However, the desorption rate of the low-strain amplitude specimen is a slightly higher than that of the high-strain amplitude specimens. It is attributed to
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Fig. 6 e Flow curves of peak stress of hysteresis loops at total strain amplitudes of 1.2% and 0.35% for the as-received and hydrogen-charged conditions.
the higher strain-induced martensite fraction when tested at the high-strain amplitude level, resulting in an increasing amount of hydrogen egress from the specimen during the fatigue life test. Additional errors in thermal desorption analysis results should be included since only one measurement was conducted for each strain amplitude level. Further study is needed to clarify this difference.
Fractography analysis Fig. 8 shows a comparison of typical macroscopic images of the fracture surfaces after the fatigue life test for both hydrogen pre-charged and uncharged specimens. All fracture surfaces can be clearly divided into two separate regions: fatigue and overload fracture areas. On a macroscopic scale, the focused observations were concentrated on the location near the crack initiation sites as shown in Fig. 9. Although certain characteristic of the fracture surface could disappear during the compression mode, formation of striations marks was visually observed from the crack initiation site on the uncharged specimen. Initiation sites of the hydrogen pre-charged specimen surfaces were completely
Fig. 7 e Hydrogen thermal desorption profile of the hydrogen pre-charged samples after the fatigue life test.
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flat and decorated with facets. In general, the fracture process during the fatigue life test can be divided into three stages that can be distinguished based on the analysis of the appearance of the fracture surface. Crack initiation was typically formed from the specimen surface at the first stage and propagated stably in the direction of thickness perpendicular to the loading direction during the second stage. The fracture process in the first and second stages is generally related to the normal fatigue crack growth mechanism while the third stage of the fracture process involves shearing and ductile tearing after the crack has reached a sufficient length. Fig. 10 shows a comparison of SEM images of the longitudinal cross sections near the edge of the fracture surface. No micro-surface cracks can be observed in the uncharged specimen and the gauge length body is stretched. Nevertheless, a substantial surface crack formation can be seen on the outer surface near the fracture edge of the hydrogen-charged specimens and a transgranular internal crack can also be observed in inward direction on the fracture surface. A depth of the surface crack is in the range of approximately 8e10 mm. A detailed microscopic observation was performed at the central location of the fatigue fracture surface shown in Fig. 11. Striation can clearly be identified all over the fracture surface of the uncharged samples, which is a typical feature of fatigue fracture surfaces of ductile materials. This indicates the repeated blunting and resharpening of the crack tip during fatigue cycle loading. Each striation mark can represent one load cycle while every load cycle can result in the formation of a striation mark. On the other hand, the fatigue fracture surface of the hydrogencharged specimens are macroscopically flat, and the fatigue fracture surface is completely covered with quasi-cleavage transgranular fractures. A number of striation marks can be observed at the fatigue fracture region of the low strain amplitude regime of the specimen (εa ¼ 0.35%); however, the fraction area where striations appear is less distinct than those in the uncharged specimen. Fig. 12 shows images of the EBSD measurement performed at the crack initiation site of the hydrogen precharged samples after the fatigue life test. Measurement was also performed on the ruptured SSRT specimen of the hydrogen pre-charged specimen. As shown in Fig. 12a, a highly localized deformation clearly appeared along the fracture surface site in the high strain amplitude regime of the specimen. The fraction of stress-induced martensite (SIM) transformation along the boundaries of the deformed grains is also visually noticeable. By contrast, less evidence of SIM was observed in the low strain amplitude regime of the sample. In the case of SSRT samples, the phase transformation from austenite to martensite occurred extensively due to the severe damage from the deformation at the necking location, as shown in Fig. 12c. Although the EBSD measurement has not been performed at crack initiation site on the uncharged specimen, it is expected that the SIM fraction is higher in the uncharged sample than that in the hydrogen pre-charged specimens due to the abrupt failure. However, the magnitude of the plastic strain within the slip band deformation is higher, which results in an earlier crack formation and an accelerated fatigue crack growth in the hydrogen pre-charged specimens [16].
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Fig. 8 e Comparison of typical macroscopic images of fracture surfaces between as-received and hydrogen pre-charged specimen after the fatigue life test for (a) at strain amplitude of 0.35%, (b) at strain amplitude of 0.75%, and (c) at strain amplitude of 1.2%.
Fatigue failure mechanism The degradation of mechanical properties in austenitic stainless steel exposed to hydrogen is typically attributed to the phase transformation from austenite to martensite. In this aspect, Ni has been reported to be important in the stability of the austenite phase and in suppressing strain-induced
martensitic transformation. The higher Ni content stabilizes the austenitic matrix, suppress martensitic transformation, and increases the stacking fault energy [14,15]. The relative reduction area of the present material is nearly identical to the value of type 316 with nickel content of 10.12 wt%, but they are relatively lower than the value of type 316L, which contains
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Fig. 9 e Origin of fatigue crack, in which crack initiated at a strain amplitude of 0.35%: (a) uncharged specimen and (b) hydrogen pre-charged specimen.
12.22 wt% [37]. Michler et al. reported that the loss of RA decreased with decreasing Ni content from 9.07 wt% to 12.75 wt% and RA is inversely proportional to the straininduced martensite transformation fraction [38]. In that way, HE could be negligible when the Ni content is above 11.5 wt% [21]. Other studies in the literature have also presented an empirical parameter to relate the HE resistance to the Ni content and other alloying elements, known as Nickelequivalent, Nieq [8,9,15e19,39e41]. Also, based on the quantitative variation of Nieq from 23.1 to 35.7, the fracture morphologies of material with the presence of hydrogen can be
comprehensively classified into five mechanisms. The Nieq value of the material in this study, calculated using Takakuwa’s equation [11], is 25.11, which is lower than the saturation value of 27, at which the effect of hydrogen environment is eliminated for austenitic stainless steel. As shown by the comparison in Fig. 11a, b, in the high strain amplitude regime, the fraction of the deformed structures is higher and the SIM transformation appearing along the deformed grain on the initiation site is more pronounced than those in the low strain amplitude regime specimen. As the SIM transformation zone can form a hydrogen diffusion pathway, hydrogen atoms are
Fig. 10 e Images of the outer surface of the rupture specimen after the fatigue life test at the strain amplitude of 1.2%: (a) uncharged specimen and (b) hydrogen pre-charged specimen.
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Fig. 11 e Comparison of microscopic fatigue fracture surfaces between uncharged specimen and hydrogen pre-charged specimen for (a), (a) at strain amplitude of 0.35%, (b) at strain amplitude of 0.75%, and (c) at strain amplitude of 1.2%.
able to accumulate before the crack tip region during cyclic loading. This can result in a localized planar deformation and stress concentration, which results in the rapid formation of micro-cracks after the beginning of the fatigue life test. Once a micro-crack is formed, a stress concentration zone appears at the crack tip, which accumulates the high SIM transformation. As mentioned previously, due to the appearance of several micro-cracks from the beginning of the test process (as shown in Fig. 10), the failure mechanism in the hydrogen precharged specimen can be considered as a competition of crack growth processes under cyclic load among the crack initiations. Once a crack propagation at one surface crack dominates the others, the propagation at this crack proceeds
continuously while the others are nearly terminated. According to the hydrogen-assisted surface crack growth (HASCG) model [6,42], the cracks are generated successively and remain in a sharp shape because the localized slip deformation in the vicinity of the crack due to hydrogen, results in an accelerated crack growth rate and induces premature failure during fatigue life test. By contrast, the failure in the uncharged specimen is completely dominated by the propagation of a crack. As localized slip deformation did not occur, a large crack blunting resulted in clearly visible striation, as shown in Fig. 10. It should be noted, that the cyclic stressestrain behavior for hydrogen pre-charged specimen was similar to those of the uncharged specimens; however, in
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Fig. 12 e EBSD measurement results of the fracture edge: (a) after fatigue life test, εa ¼ 1.2%; (b) after fatigue life test, εa ¼ 0.35; and (c) after SSRT test of the hydrogen pre-charged specimen. In each figure, the inverse pole figure image is on the left side and the phase map is located on the right side.
the high strain amplitude regime a failure can occur sooner in the pre-charged specimens. Nagging a reported that, compared with hydrogen-induced surface crack initiation sites in the hydrogen pre-charged samples, the fatigue crack initiation mechanism for the uncharged specimens could be attributed to the irreversible plastic strain-induced surface roughening [43].
Conclusions An SSRT and strain-controlled fatigue life test of low nickel content type 316L austenitic stainless steels was performed and the results of uncharged and hydrogen pre-charged
specimens were compared, from which the following conclusions can be drawn: 1. The SSRT results clearly reveal that hydrogen changes the fracture mechanism of type 316L steel and results in significant loss in the reduction area and strain at the fracture. Treatment with gaseous hydrogen has a stronger effect than pre-charging specimens with hydrogen. 2. The reduction of fatigue life in the high strain amplitude region is related to a significant decrease of the plastic component while the effect of hydrogen on the elastic component was negligible. The fatigue limit of the precharged specimens was consistently equivalent with the results of the uncharged sample.
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3. The difference in fracture appearances is mainly related to the accelerated crack initiation due to hydrogen during the strain-controlled LCF life test. The highly localized deformation and pronounced martensite transformation occur near the side of the fracture surface in the high strain amplitude regime, resulting in an early formation of abundant surface cracks.
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Acknowledgements
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This research was supported by Development of Reliability Measurement & Standard Technology for Hydrogen Fueling Station Funded by Korea Research Institute of Standards and Science (KRISS-2019-GP2019-0012).
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