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Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen
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Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen Junichiro Yamabe a,b,c,*, Osamu Takakuwa b,**, Hisao Matsunaga b,c,d, Hisatake Itoga b,1, Saburo Matsuoka b a
International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan b Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan c International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan d Department of Mechanical Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan
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abstract
Article history:
The effects of external and internal hydrogen on the slow-strain-rate tensile (SSRT)
Received 30 January 2017
properties at room temperature were studied for ten types of solution-treated austenitic
Received in revised form
stainless steels containing a small amount of additive elements. The hydrogen diffusivity
4 April 2017
and solubility of the steels were measured with high-pressure hydrogen gas. The
Accepted 7 April 2017
remarkable tensile-ductility loss observed in the SSRT tests was attributed to hydrogen-
Available online xxx
induced successive crack growth (HISCG) and was successfully quantified according to the nickel-equivalent content (Nieq), which represents the stability of the austenitic phase.
Keywords:
The relative reduction in area (RRA) of the steels with a larger Nieq was influenced by the
Hydrogen embrittlement
hydrogen distribution, whereas that of the steels with a smaller Nieq was not. This unique
Slow-strain-rate tensile test
trend was interpreted with regard to the hydrogen distribution and fracture morphology
Austenitic stainless steel
(HISCG or microvoid coalescence).
Nickel equivalent
© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Additive element
Introduction Hydrogen degrades the tensile and fatigue properties of various metals. This phenomenon is well known as hydrogen
embrittlement (HE). For ensuring the safety and reliability of the components used in high-pressure hydrogen gas, it is necessary to perform the strength design of the components in consideration of the detrimental effect of the hydrogen
* Corresponding author. International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan. ** Corresponding author. E-mail addresses: [email protected] (J. Yamabe), [email protected] (O. Takakuwa). 1 Present address: Hydrogen Energy Test and Research Center (HyTReC), 915-1 Tomi, Itoshima-shi, Fukuoka 819-1133, Japan. http://dx.doi.org/10.1016/j.ijhydene.2017.04.055 0360-3199/© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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[1e4]. The HE of metals is often determined using the relative reduction in area (RRA), which is obtained by slow-strain-rate tensile (SSRT) testing [5e21]. The RRA is defined as the ratio of the reduction in area (RA) for hydrogen gas to that for inert gas. Kobayashi et al. [17] analyzed a series of RRA data for conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldin high-pressure gaseous hydrogen, demonstrating a clear relationship between the RRA and the nickel-equivalent content (Nieq) based on Hirayama's equation [23]. However, Hirayama's equation does not include nitrogen, which improves the stability of the austenitic phase; therefore, an application of an alternative nickel-equivalent equation containing nitrogen based on Sanga's equation has been recently proposed [24]. The determination of the hydrogen compatibility according to the nickel-equivalent equations is applicable for solution-treated austenitic stainless steels showing degradations in the RRA that are caused by hydrogen-assisted surface crack growth (HASCG) [22]. In order to elucidate the mechanism of the HASCG, we performed elasto-plastic fracture toughness (JIC), fatigue crack growth and fatigue life tests of Types 304, 316 and 316L in high-pressure hydrogen gas [22]. As a result, it was demonstrated that, in high-pressure hydrogen gas, the SSRT surface crack grew via the same mechanism as for JIC crack and fatigue crack, i.e., these crack growths could be uniformly explained on the basis of the hydrogen-induced successive crack growth (HISCG) model, which considers that cracks successively grow with a sharp shape under the loading process, due to local slip deformations near the crack tip by hydrogen [22]. Accordingly, the HASCG can be explained based on the HISCG model. In contrast, when degradations in the RRA are dominated by a mechanism different from the HISCG, the hydrogen compatibility is not necessarily determined by the nickelequivalent equations. It is reported that the RRA of stable austenitic stainless steels with a large amount of manganese and nitrogen is sometimes degraded by hydrogen, in spite of the high stability of the austenitic phase [16]. As a substitute for the SSRT test of non-charged specimens in gaseous hydrogen, SSRT tests of hydrogen-charged steels prepared by exposure to gaseous hydrogen are performed in inert gas to determine the RRA [14,19e21] because an expensive high-pressure hydrogen tensile tester is not needed. In addition, the hydrogen charging can be performed electrochemically. However, the relationship between the RRA of specimens with internal hydrogen (hydrogen-charged specimens tested in inert gas) and that of specimens with external hydrogen (non-charged specimens tested in gaseous hydrogen) is unclear. Therefore, the correlations between the RRAs for specimens with internal and external hydrogen should be clarified for interpreting experimental data obtained under various environmental conditions. In this study, we investigated the effects of external and internal hydrogen on the SSRT properties of solution-treated austenitic stainless steels containing a small amount of additive elements. Whereas degradations of the RRA for various stable austenitic stainless steels have been reported [16], we targeted various metastable austenitic stainless steels rather than stable ones because metastable austenitic stainless steels have a higher tensile strength (TS) and lower cost than
stable austenitic stainless steels and may be suitable for use in hydrogen-gas environments with a reasonable strength design in consideration of the detrimental effect of hydrogen [22]. The SSRT tests were conducted using ten types of austenitic stainless steels. Five types of 300-series austenitic stainless steels containing additive elements (nitrogen, niobium, vanadium, and titanium) were used. For comparison, three types of conventional 300-series stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM19dwere also tested. The following SSRT tests were performed with both internal hydrogen and external hydrogen at room temperature (RT): (i) tests of non-charged specimens in air, (ii) tests of hydrogen-charged specimens in air, and (iii) tests of non-charged specimens in high-pressure hydrogen gas. To estimate the hydrogen distribution in the SSRT specimens, the hydrogen diffusivity and solubility of the steels were determined by a desorption method with high-pressure hydrogen gas.
Experimental procedures Materials Ten types of solution-treated austenitic stainless steels were used. Table 1 shows the chemical composition along with the average Vickers hardness measured under a load of 9.8 N. Table 2 shows the values of Nieq for the steels calculated by using Eq. (4) mentioned later. Five types of 300-series austenitic stainless steels containing a small amount of additive elementsdTypes 304N2, 321, 304(V), 304(N), and 304(Nb)dwere prepared. Types 304N2 and 321 were commercial materials regulated by the Japanese Industrial Standards. Types 304(V), 304(N), and 304(Nb) were specially prepared for this study. Types 321, 304(V), 304(N), and 304(Nb) contained a small amount of titanium, vanadium, nitrogen, and niobium, respectively. Type 304N2 contained both nitrogen and niobium. For comparison, three types of conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM-19dwere also tested.
Specimens and test methods SSRT testing The SSRT specimens machined from the steels had a diameter of 6 mm and a reduced-section length of 28.6 mm, in accordance with ASTM G142-98 [25]. The specimen surface was carefully finished according to ASTM G142-98, as the surface roughness substantially affects the SSRT properties in hydrogen gas. The SSRT tests were conducted by using a servo-hydraulic testing machine equipped with a highpressure vessel. The crosshead speed in the SSRT tests was 0.002 mm/s, which accords with ASTM G142-98 [25]. Some SSRT specimens were charged via exposure to 100-MPa hydrogen gas at 543 K for 200 h. As mentioned later, this hydrogen exposure provided a uniform distribution of hydrogen in the SSRT specimens of all the steels. The
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Table 1 e Chemical composition [mass %] and Vickers hardness.a Material 304N2 321 304(V) 304(N) 304(Nb) 304 316 316L XM-19 HP160 a b c
Heat
C
Si
Mn
P
S
Ni
Cr
Mo
Cu
V
Nb
Ti
N
HV
A A F G H Ab Cc Bb Cc Bb Cc B B
0.060 0.030 0.049 0.047 0.047 0.056 0.059 0.020 0.046 0.010 0.021 0.040 0.030
0.67 0.48 0.50 0.50 0.51 0.43 0.45 0.53 0.63 0.53 0.54 0.41 0.47
1.84 0.83 1.49 1.50 1.50 1.72 0.85 0.98 0.93 0.77 0.88 4.56 3.98
0.028 0.025 0.010 0.010 0.010 0.030 0.028 0.021 0.031 0.023 0.021 0.017 0.016
0.001 0.001 0.0098 0.0098 0.010 0.025 0.002 0.001 0.001 0.001 <0.001 0.001 0.0015
7.92 9.21 10.00 10.10 10.10 8.87 8.15 10.15 10.12 12.13 12.22 12.50 9.71
18.54 17.21 19.04 18.80 18.80 18.26 18.16 16.21 16.89 17.16 17.64 22.40 20.70
e e e e e e 0.32 2.08 2.06 2.86 2.19 2.04 2.29
e e e e e e e e e e e e 0.11
e e 0.22 e e e e e e e e 0.19 e
0.10 e e e 0.20 e e e e e e 0.12 0.28
e 0.35 e e e e e e e e e e e
0.22 e e 0.11 e e 0.052 e 0.028 e 0.034 0.34 0.39
212 139 162 176 159 e 174 e 157 e 148 201 238
e: not assigned or not measured. Used only for determination of hydrogen-diffusion properties. Used only for SSRT testing.
Table 2 e Nickel-equivalent content [mass %].a Material Heat Nieq a
304N2
321
304(V)
304(N)
304(Nb)
304
316
316L
XM-19
HP160
A 25.6
A 22.8
F 25.9
G 26.6
H 25.8
C 23.5
C 25.8
C 28.2
B 38.5
B 34.3
The nickel-equivalent content was calculated by using Eq. (4).
following specimens were used: (i) non-charged specimens tested in air at RT, (ii) hydrogen-charged specimens tested in air at RT, and (iii) non-charged specimens tested in hydrogen gas at a pressure of 78e115 MPa at RT. In test (ii), internal hydrogen was used, whereas in test (iii), external hydrogen was used. For the non-charged specimens tested in hydrogen gas, no kept time was generally given before the SSRT tests. The purity of the source hydrogen gas to fill the chamber was 99.999% and the guaranteed values of oxygen and water were less than 1 and 2.6 vol. ppm, respectively [3]. In contrast, the measured contents of oxygen and water after tests are always less than 1 and 10 vol. ppm [3].
Determination of hydrogen diffusivity and solubility To determine the temperature dependence of the hydrogen diffusivity and solubility, cylindrical specimens with 2r0 ¼ 5e7 mm and z0 ¼ 0.1e5 mm, where r0 is the radius of the specimen and z0 is its thickness, were sampled from each of the steels. The surfaces of these specimens were finished with #600 emery paper. Based on our preliminary investigations with carbon, low-alloy, and austenitic stainless steels, the surfaces of specimens finished by buff polishing provided the same measured hydrogen solubility and diffusivity as those finished with #600 and #2000 emery papers, revealing that, within the present measurement method, the surface condition hardly affects the measured values. The specimens were exposed to hydrogen gas at pressures ranging from 10 to 100 MPa and temperatures ranging from 383 to 543 K for the specific time needed to obtain the uniform distribution of hydrogen. After the exposure, the hydrogen contents of the specimens were measured under constant or rising temperatures by gas chromatographyemass spectroscopy. The hydrogen diffusivity was determined by fitting the solution of
a diffusion equation to the experimental hydrogen contents measured at various constant temperatures [26e28]. This method for determining the hydrogen diffusivity and solubility is a desorption method [28], which differs from the conventional permeation method [29].
Results and discussion Hydrogen diffusivity and solubility Fig. 1 shows Arrhenius plots of the hydrogen diffusivity and solubility of Type 304 obtained by the desorption method with high-pressure hydrogen. Data from the literature for the hydrogen diffusivity and solubility measured by permeation tests at a low pressure are shown in Fig. 1 [30e34]. The hydrogen diffusivity and solubility determined by the desorption method at a high pressure were nearly consistent with the literature data. The results shown in Fig. 1 suggest that reliable data for the hydrogen diffusivity and hydrogen solubility of austenitic stainless steels can be obtained by the desorption method with high-pressure hydrogen [28]. Fig. 2 shows the hydrogen diffusivity and hydrogen solubility of ten types of austenitic stainless steels determined by the desorption method at a high pressure. Interestingly, the hydrogen diffusivity of all the stainless steels was the same; thus, its temperature dependence was fitted by the following Arrhenius-type equation: D ¼ D0 expð ED =RTÞ; 6
(1) 2
1
1
where D0 ¼ 3.91 10 m s , ED ¼ 60.4 kJ mol , R is the gas constant, and T is the absolute temperature. At RT, the hydrogen diffusivity was ~1 1016 m2 s1, which is six orders
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 1 e Hydrogen diffusivity and solubility of Type 304: a hydrogen diffusivity; b hydrogen solubility.
Fig. 2 e Hydrogen diffusivity and solubility of all stainless steels: a hydrogen diffusivity; b hydrogen solubility. of magnitude lower than that of quenched and tempered Cre Mo steels [28]. In contrast to the hydrogen diffusivity, the hydrogen solubility differed slightly according to the type of steel; i.e., the hydrogen solubility of the high-strength, nitrogen-added steelsdXM-19 and HP160dwas slightly higher than that of the 300-series stainless steels, although the activation energy for the hydrogen solubility (ES) was independent of the type of steel. Under the assumption that ES was independent of the type of steel, the temperature dependence of the hydrogen solubility was fitted by the Arrhenius-type equation: S ¼ S0 expð ES =RTÞ;
(2) 1/2
1
where S0 ¼ 13.3 mass ppm MPa and ES ¼ 2.16 kJ mol for the 300-series stainless steels; S0 ¼ 20.5 mass ppm MPa1/2 and ES ¼ 2.16 kJ mol1 for HP160; and S0 ¼ 24.4 mass ppm MPa1/2 and ES ¼ 2.16 kJ mol1 for XM-19. According to the measured hydrogen diffusivity and solubility, the distributions of the hydrogen concentration in the specimens with external hydrogen and internal hydrogen
were estimated by using a solution of Fick's second law. Especially, when hydrogen exists only near the specimen surface, its distribution can be approximately estimated by using a solution of Fick's second law for hydrogen entry into a semi-infinite plate, as follows: oi h n c ¼ c0 þ ðcs c0 Þ 1 erf z=ð2ðDtÞÞ1=2 ;
(3)
where c is the hydrogen concentration of the specimen at z and t, c0 is the initial hydrogen concentration of the specimen, cs is the saturated hydrogen concentration, z is the penetration depth of hydrogen in the specimen, t is the exposure time or the tested time, and erf is the error function. Fig. 3 shows the estimated hydrogen concentrations of the 300-series austenitic stainless steel near the surface of the specimens with external hydrogen and internal hydrogen. Hydrogen exposure for 200 h at a pressure of 100 MPa and a temperature of 543 K (internal hydrogen) yielded a nearly uniform distribution of hydrogen in the SSRT specimen. The validity of the estimated hydrogen concentration was experimentally verified by secondary ion mass spectroscopy [18].
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 3 e Estimated distribution of hydrogen concentration near surface of specimens with internal hydrogen (100 MPa and 543 K for 200 h) and external hydrogen (100 MPa and RT for 3 h).
On the other hand, in the case of hydrogen exposure for 3 h at a pressure of 100 MPa at RT (external hydrogen), there was a high concentration of hydrogen only at the surface of the SSRT specimen; i.e., a uniform distribution of hydrogen was not obtained. The distributions of the hydrogen concentration in the specimens and related boundary conditions differed significantly between the cases of internal and external hydrogen.
SSRT properties in presence of hydrogen Fig. 4 shows the nominal (engineering) stressestrain curves of the 300-series stainless steels tested at RT. Among the hydrogen-charged specimens tested in air, all the stainless steels except for Types 304(N) and 316L failed before reaching the TS of the non-charged specimens tested in air. Similarly, among the non-charged specimens tested in high-pressure hydrogen gas, all the stainless steels except for Types 304(N), 316, and 316L failed before reaching the TS of the non-charged specimens tested in air. The RRAs of Types 304, 316, and 316L have already been investigated by using the Nieq [17]. Although the concept of Nieq is discussed controversially in the community, this study mainly uses 300-series austenitic stainless steels (Nieq ¼ 22.8e38.5%); therefore, the RRA was investigated based on the concept of Nieq. Fig. 5 shows the relationship between the RRA and Nieq for all the stainless steel. The Nieq was calculated by a modified Sanga's equation [24,35]: Nieq ¼ Ni þ 12:93 C þ 1:11 Mn þ 0:72 Cr þ 0:88 Mo 0:27 Si þ 7:55 N
(4)
The stainless steels with a smaller Nieq (Nieq z 23%) were not influenced by the hydrogen distribution in the specimen (external or internal hydrogen). In contrast, the stainless steels with a larger Nieq (Nieq 28%) were influenced by the hydrogen distribution. The effects of external and internal hydrogen on the RRAs depended on the Nieq. This interesting
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phenomenon was interpreted according to the observation of the tensile fracture morphologies, which is discussed later. Fig. 5 also contains the literature data of 300-series austenitic stainless steels with external and internal hydrogen [14]. The RRAs at 223 K were lower in 40-MPa external hydrogen compared to internal hydrogen; however, the RRAs at RT were almost the same between 70-MPa external hydrogen and internal hydrogen [14]. The RRAs for external and internal hydrogen in the literature reasonably agreed with those for external hydrogen in this study; however, seemed to be greater than those for internal hydrogen in this study. For internal hydrogen, the literature used 34.5-MPa hydrogen gas at 573 K (cs ~ 60 mass ppm), whereas this study used 100-MPa hydrogen gas at 543 K (cs ~ 100 mass ppm). Since the saturated hydrogen concentration in this study was 1.7 times higher than that in the literature for internal hydrogen, the RRAs for internal hydrogen in this study were presumed to be lower than those in the literature.
Tensile fracture morphologies by scanning electron microscopy (SEM) Steel with a smaller Nieq: Type 304 Fig. 6 shows SEM micrographs of the tensile fracture morphologies of Type 304 at RT. In the non-charged specimen tested in air, the specimen failed because of an ordinary cupand-cone fracture dominated by microvoid coalescence (MVC). The fracture surface consisted of a normal stressfracture region covered with equiaxial dimple (A in Fig. 6(a)) and a shear stress-fracture region covered with elongated dimple (B in Fig. 6(a)). In the hydrogen-charged specimen tested in air, numerous surface cracks were observed, and the specimen failed because of the extension of these cracks. The fracture surface was dominated by quasi-cleavage, differing completely from that of the non-charged specimen tested in air. We consider that the internal hydrogen at the crack tip underwent stressinduced diffusion, extending the surface cracks. Our previous study demonstrates that, for austenitic stainless steels used in hydrogen gas, the SSRT surface crack grew via the same mechanism as for JIC and fatigue cracks based on the HISCG model [22]. Accordingly, the HIS crack is ductile, not brittle. The results shown in Fig. 6 suggest that the degradation in the RRA of Type 304 with internal hydrogen was caused by extension of surface cracks due to hydrogen associated with the HISCG [22]. The non-charged specimen tested in hydrogen gas also failed because of the extension of surface cracks, which was assisted by external hydrogen, and the fracture surface was dominated by quasi-cleavage. This morphology reveals that the degradation in the RA of Type 304 with external hydrogen, as well as that of Type 304 with internal hydrogen, was caused by the HISCG.
Steel with a larger Nieq: Type 316L Fig. 7 shows SEM micrographs of the tensile fracture morphologies of Type 316L. All of the specimens in Fig. 7 failed because of cup-and-cone fractures dominated by the MVC. In contrast to Type 304 with internal and external hydrogen, the HISCG did not dominate the fracture of Type 316L with
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 4 e Effect of external or internal hydrogen on SSRT properties of austenitic stainless steels at RT: a Type 304N2; b Type 321; c Type 304(V); d Type 304(N); e Type 304(Nb); f Type 304; g Type 316; h Type 316L.
internal and external hydrogen. As shown in Fig. 2, the hydrogen diffusivity and hydrogen solubility of Types 304 and 316L were the same; accordingly, the difference in the fracture-surface morphologies between Types 304 and 316L
cannot be explained according to the hydrogen distribution in the specimen. It is speculated that the HISCG is attributed to the combined effect of the strain-induced martensitic transformation and the hydrogen distribution near the crack tip. In
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Comparing the normal stress-fracture regions of Type 316L between the cases of external and internal hydrogen, a reduction in the dimple size was observed for internal hydrogen, whereas such a reduction was not observed for external hydrogen. Matsuo and Yamabe et al. [19] attributed the small dimple of Type 316L with internal hydrogen to void sheets produced by local shear strain [36,37]. The localization of this shear strain was enhanced by internal hydrogen, resulting in a slight reduction in the RA. In contrast, in Type 316L with external hydrogen, hydrogen did not exist in the normal stress-fracture region, as suggested by Fig. 3. Therefore, the external hydrogen did not affect the dimple morphology of the normal stress-fracture region, causing a negligible reduction in the RA.
Steel with a medium Nieq: Type 304(Nb)
Fig. 5 e Relationship between RRA and nickel-equivalent content of austenitic stainless steels at RT, together with literature data [14].
Type 316L, the strain-induced martensitic transformation was prevented by the high stability of the austenitic phase; as a result, the HISCG was not produced during the SSRT test, and the specimen failed because of the cup-and-cone fracture dominated by the MVC.
Fig. 8 shows SEM micrographs of the tensile fracture morphologies of Type 304(Nb) at RT. The non-charged specimen tested in air failed because of the cup-and-cone fracture, consisting of the normal stress-fracture region associated with equiaxial dimple and the shear stress-fracture region associated with elongated dimple. In the hydrogen-charged specimen tested in air, the fracture surface was covered with a mixture of equiaxial dimple and quasi-cleavage. In the non-charged specimen tested in hydrogen gas, the fracture surface was covered with a mixture of elongated dimple and quasi-cleavage. The fracture of the stainless steels having an RRA in the transition region competed with MVC and HISCG in the presence of hydrogen.
Fig. 6 e SEM micrographs of tensile fracture morphologies of Type 304 at RT. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 7 e SEM micrographs of tensile fracture morphologies of Type 316L at RT.
Fig. 8 e SEM micrographs of tensile fracture morphologies of Type 304(Nb) at RT. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 9 e Tensile fracture processes of austenitic stainless steels with external and internal hydrogen: a Fracture dominated by HISCG; b Ordinary cup-and-cone fracture.
Effects of external and internal hydrogen on RRA As shown in Fig. 5, the RRAs of the steels with a smaller Nieq were not influenced by the hydrogen distribution, whereas those of the steels with a larger Nieq were influenced by the hydrogen distribution. In this section, we discuss the effect of the hydrogen distribution on the RRAs. In the stainless steels with a smaller Nieq, the degradation in the RRA was dominated by the HISCG. Localized hydrogen at the crack tip was not directly identified in this study; however, it is speculated that both internal hydrogen and external hydrogen diffused at the crack tip during the SSRT test and contributed to the HISCG, as shown in Fig. 9. Thus, the RRA of the steels with a smaller Nieq was barely influenced by the hydrogen distribution. In the steels with a larger Nieq, the specimens failed because of the cup-and-cone fracture dominated by the MVC. The hydrogen affected the void formation in the normal stress-fracture region, slightly degrading the RRA. Only internal hydrogen contributed to the void formation in the normal stress-fracture region, because the SSRT specimen with internal hydrogen had a uniform distribution of hydrogen. Understanding the hydrogen distribution in the specimen and the tensile-fracture process (MVC or HISCG) is important for interpreting the effect of the hydrogen distribution on solution-treated austenitic stainless steels.
Conclusion SSRT tests for ten types of solution-treated austenitic stainless steels with external and internal hydrogen were performed at RT. Five types of 300-series austenitic stainless steels containing a small amount of additive elementsdTypes 304N2, 321, 304(V), 304(N), and 304(Nb)dwere prepared. For comparison, three types of conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM-19dwere also tested. The hydrogen diffusivity and solubility of these steels were determined by the desorption method with high-pressure gaseous hydrogen. Our conclusions are summarized as follows: 1. The measured hydrogen diffusivity of all the steels was the same, which is consistent with literature data obtained by conventional permeation tests at a low pressure. In contrast, the measured hydrogen solubility was found to slightly depend on the type of steel, and the hydrogen solubility of the high-strength, nitrogen-added austenitic stainless steels HP160 and XM-19 was slightly higher than that of the 300-series austenitic stainless steels. The hydrogen solubility of the 300-series steels was not affected by a small amount of additive elements.
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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2. All the 300-series stainless steels except for Types 304(N), 316, and 316L failed before reaching the TS of the non-charged specimens tested in air with internal or external hydrogen. 3. All the stainless steels failed because of ductile cup-andcone fracture dominated by MVC in the absence of hydrogen, and several stainless steels failed because of quasi-cleavage fracture dominated by HISCG in the presence of hydrogen. In the presence of hydrogen, a significant reduction in the RRA was caused by the HISCG. 4. The stainless steels with a higher nickel-equivalent content (Nieq) had a higher RRA. This suggests that the RRA of the present austenitic stainless steels can be quantified according to the stability of the austenitic phase. 5. The RRA of the steels with a larger Nieq was influenced by the hydrogen distribution in the specimen (external or internal hydrogen), whereas that of the steels with a smaller Nieq was not. These unique trends can be interpreted according to the hydrogen distribution in the SSRT specimen and the tensile-fracture morphology (HISCG or MVC).
[10]
[11]
[12]
[13]
[14]
[15]
[16]
Acknowledgment [17]
This work was partially supported by the New Energy and Industrial Technology Development Organization (NEDO), Hydrogen Utilization Technology (2013e2018).
references
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[32] Sun XK, Xu J, Li YY. Hydrogen permeation behaviour in austenitic stainless steels. Mat Sci Eng A 1989;114:179e87. [33] Perng T-P, Altstetter CJ. Effects of deformation on hydrogen permeation in austenitic stainless steels. Acta Metall 1986;34:1771e87. [34] Yamabe J, Awane T, Matsuoka S. Elucidating the hydrogenentry-obstruction mechanism of a newly developed aluminum-based coating in high-pressure gaseous hydrogen. Int J Hydrogen Energy 2015;40:10329e39. [35] Sanga M, Yukawa N, Ishikawa T. Influence of chemical composition on d-induced transformation in austenitic stainless steel. JSTP 2000;41:64e8. [36] Roger HC. The tensile, fracture of ductile metals. T Metall Soc AIME 1960;218:498e506. [37] Cox TB, Low Jr JR. An investigation of the plastic fracture of AISI 44340 and 18 nickel-200 grade maraging steels. Metall Trans 1974;5:1457e70.
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Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen Junichiro Yamabe a,b,c,*, Osamu Takakuwa b,**, Hisao Matsunaga b,c,d, Hisatake Itoga b,1, Saburo Matsuoka b a
International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan b Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan c International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan d Department of Mechanical Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan
article info
abstract
Article history:
The effects of external and internal hydrogen on the slow-strain-rate tensile (SSRT)
Received 30 January 2017
properties at room temperature were studied for ten types of solution-treated austenitic
Received in revised form
stainless steels containing a small amount of additive elements. The hydrogen diffusivity
4 April 2017
and solubility of the steels were measured with high-pressure hydrogen gas. The
Accepted 7 April 2017
remarkable tensile-ductility loss observed in the SSRT tests was attributed to hydrogen-
Available online xxx
induced successive crack growth (HISCG) and was successfully quantified according to the nickel-equivalent content (Nieq), which represents the stability of the austenitic phase.
Keywords:
The relative reduction in area (RRA) of the steels with a larger Nieq was influenced by the
Hydrogen embrittlement
hydrogen distribution, whereas that of the steels with a smaller Nieq was not. This unique
Slow-strain-rate tensile test
trend was interpreted with regard to the hydrogen distribution and fracture morphology
Austenitic stainless steel
(HISCG or microvoid coalescence).
Nickel equivalent
© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Additive element
Introduction Hydrogen degrades the tensile and fatigue properties of various metals. This phenomenon is well known as hydrogen
embrittlement (HE). For ensuring the safety and reliability of the components used in high-pressure hydrogen gas, it is necessary to perform the strength design of the components in consideration of the detrimental effect of the hydrogen
* Corresponding author. International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka-shi, Fukuoka 819-0395, Japan. ** Corresponding author. E-mail addresses: [email protected] (J. Yamabe), [email protected] (O. Takakuwa). 1 Present address: Hydrogen Energy Test and Research Center (HyTReC), 915-1 Tomi, Itoshima-shi, Fukuoka 819-1133, Japan. http://dx.doi.org/10.1016/j.ijhydene.2017.04.055 0360-3199/© 2017 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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[1e4]. The HE of metals is often determined using the relative reduction in area (RRA), which is obtained by slow-strain-rate tensile (SSRT) testing [5e21]. The RRA is defined as the ratio of the reduction in area (RA) for hydrogen gas to that for inert gas. Kobayashi et al. [17] analyzed a series of RRA data for conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldin high-pressure gaseous hydrogen, demonstrating a clear relationship between the RRA and the nickel-equivalent content (Nieq) based on Hirayama's equation [23]. However, Hirayama's equation does not include nitrogen, which improves the stability of the austenitic phase; therefore, an application of an alternative nickel-equivalent equation containing nitrogen based on Sanga's equation has been recently proposed [24]. The determination of the hydrogen compatibility according to the nickel-equivalent equations is applicable for solution-treated austenitic stainless steels showing degradations in the RRA that are caused by hydrogen-assisted surface crack growth (HASCG) [22]. In order to elucidate the mechanism of the HASCG, we performed elasto-plastic fracture toughness (JIC), fatigue crack growth and fatigue life tests of Types 304, 316 and 316L in high-pressure hydrogen gas [22]. As a result, it was demonstrated that, in high-pressure hydrogen gas, the SSRT surface crack grew via the same mechanism as for JIC crack and fatigue crack, i.e., these crack growths could be uniformly explained on the basis of the hydrogen-induced successive crack growth (HISCG) model, which considers that cracks successively grow with a sharp shape under the loading process, due to local slip deformations near the crack tip by hydrogen [22]. Accordingly, the HASCG can be explained based on the HISCG model. In contrast, when degradations in the RRA are dominated by a mechanism different from the HISCG, the hydrogen compatibility is not necessarily determined by the nickelequivalent equations. It is reported that the RRA of stable austenitic stainless steels with a large amount of manganese and nitrogen is sometimes degraded by hydrogen, in spite of the high stability of the austenitic phase [16]. As a substitute for the SSRT test of non-charged specimens in gaseous hydrogen, SSRT tests of hydrogen-charged steels prepared by exposure to gaseous hydrogen are performed in inert gas to determine the RRA [14,19e21] because an expensive high-pressure hydrogen tensile tester is not needed. In addition, the hydrogen charging can be performed electrochemically. However, the relationship between the RRA of specimens with internal hydrogen (hydrogen-charged specimens tested in inert gas) and that of specimens with external hydrogen (non-charged specimens tested in gaseous hydrogen) is unclear. Therefore, the correlations between the RRAs for specimens with internal and external hydrogen should be clarified for interpreting experimental data obtained under various environmental conditions. In this study, we investigated the effects of external and internal hydrogen on the SSRT properties of solution-treated austenitic stainless steels containing a small amount of additive elements. Whereas degradations of the RRA for various stable austenitic stainless steels have been reported [16], we targeted various metastable austenitic stainless steels rather than stable ones because metastable austenitic stainless steels have a higher tensile strength (TS) and lower cost than
stable austenitic stainless steels and may be suitable for use in hydrogen-gas environments with a reasonable strength design in consideration of the detrimental effect of hydrogen [22]. The SSRT tests were conducted using ten types of austenitic stainless steels. Five types of 300-series austenitic stainless steels containing additive elements (nitrogen, niobium, vanadium, and titanium) were used. For comparison, three types of conventional 300-series stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM19dwere also tested. The following SSRT tests were performed with both internal hydrogen and external hydrogen at room temperature (RT): (i) tests of non-charged specimens in air, (ii) tests of hydrogen-charged specimens in air, and (iii) tests of non-charged specimens in high-pressure hydrogen gas. To estimate the hydrogen distribution in the SSRT specimens, the hydrogen diffusivity and solubility of the steels were determined by a desorption method with high-pressure hydrogen gas.
Experimental procedures Materials Ten types of solution-treated austenitic stainless steels were used. Table 1 shows the chemical composition along with the average Vickers hardness measured under a load of 9.8 N. Table 2 shows the values of Nieq for the steels calculated by using Eq. (4) mentioned later. Five types of 300-series austenitic stainless steels containing a small amount of additive elementsdTypes 304N2, 321, 304(V), 304(N), and 304(Nb)dwere prepared. Types 304N2 and 321 were commercial materials regulated by the Japanese Industrial Standards. Types 304(V), 304(N), and 304(Nb) were specially prepared for this study. Types 321, 304(V), 304(N), and 304(Nb) contained a small amount of titanium, vanadium, nitrogen, and niobium, respectively. Type 304N2 contained both nitrogen and niobium. For comparison, three types of conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM-19dwere also tested.
Specimens and test methods SSRT testing The SSRT specimens machined from the steels had a diameter of 6 mm and a reduced-section length of 28.6 mm, in accordance with ASTM G142-98 [25]. The specimen surface was carefully finished according to ASTM G142-98, as the surface roughness substantially affects the SSRT properties in hydrogen gas. The SSRT tests were conducted by using a servo-hydraulic testing machine equipped with a highpressure vessel. The crosshead speed in the SSRT tests was 0.002 mm/s, which accords with ASTM G142-98 [25]. Some SSRT specimens were charged via exposure to 100-MPa hydrogen gas at 543 K for 200 h. As mentioned later, this hydrogen exposure provided a uniform distribution of hydrogen in the SSRT specimens of all the steels. The
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Table 1 e Chemical composition [mass %] and Vickers hardness.a Material 304N2 321 304(V) 304(N) 304(Nb) 304 316 316L XM-19 HP160 a b c
Heat
C
Si
Mn
P
S
Ni
Cr
Mo
Cu
V
Nb
Ti
N
HV
A A F G H Ab Cc Bb Cc Bb Cc B B
0.060 0.030 0.049 0.047 0.047 0.056 0.059 0.020 0.046 0.010 0.021 0.040 0.030
0.67 0.48 0.50 0.50 0.51 0.43 0.45 0.53 0.63 0.53 0.54 0.41 0.47
1.84 0.83 1.49 1.50 1.50 1.72 0.85 0.98 0.93 0.77 0.88 4.56 3.98
0.028 0.025 0.010 0.010 0.010 0.030 0.028 0.021 0.031 0.023 0.021 0.017 0.016
0.001 0.001 0.0098 0.0098 0.010 0.025 0.002 0.001 0.001 0.001 <0.001 0.001 0.0015
7.92 9.21 10.00 10.10 10.10 8.87 8.15 10.15 10.12 12.13 12.22 12.50 9.71
18.54 17.21 19.04 18.80 18.80 18.26 18.16 16.21 16.89 17.16 17.64 22.40 20.70
e e e e e e 0.32 2.08 2.06 2.86 2.19 2.04 2.29
e e e e e e e e e e e e 0.11
e e 0.22 e e e e e e e e 0.19 e
0.10 e e e 0.20 e e e e e e 0.12 0.28
e 0.35 e e e e e e e e e e e
0.22 e e 0.11 e e 0.052 e 0.028 e 0.034 0.34 0.39
212 139 162 176 159 e 174 e 157 e 148 201 238
e: not assigned or not measured. Used only for determination of hydrogen-diffusion properties. Used only for SSRT testing.
Table 2 e Nickel-equivalent content [mass %].a Material Heat Nieq a
304N2
321
304(V)
304(N)
304(Nb)
304
316
316L
XM-19
HP160
A 25.6
A 22.8
F 25.9
G 26.6
H 25.8
C 23.5
C 25.8
C 28.2
B 38.5
B 34.3
The nickel-equivalent content was calculated by using Eq. (4).
following specimens were used: (i) non-charged specimens tested in air at RT, (ii) hydrogen-charged specimens tested in air at RT, and (iii) non-charged specimens tested in hydrogen gas at a pressure of 78e115 MPa at RT. In test (ii), internal hydrogen was used, whereas in test (iii), external hydrogen was used. For the non-charged specimens tested in hydrogen gas, no kept time was generally given before the SSRT tests. The purity of the source hydrogen gas to fill the chamber was 99.999% and the guaranteed values of oxygen and water were less than 1 and 2.6 vol. ppm, respectively [3]. In contrast, the measured contents of oxygen and water after tests are always less than 1 and 10 vol. ppm [3].
Determination of hydrogen diffusivity and solubility To determine the temperature dependence of the hydrogen diffusivity and solubility, cylindrical specimens with 2r0 ¼ 5e7 mm and z0 ¼ 0.1e5 mm, where r0 is the radius of the specimen and z0 is its thickness, were sampled from each of the steels. The surfaces of these specimens were finished with #600 emery paper. Based on our preliminary investigations with carbon, low-alloy, and austenitic stainless steels, the surfaces of specimens finished by buff polishing provided the same measured hydrogen solubility and diffusivity as those finished with #600 and #2000 emery papers, revealing that, within the present measurement method, the surface condition hardly affects the measured values. The specimens were exposed to hydrogen gas at pressures ranging from 10 to 100 MPa and temperatures ranging from 383 to 543 K for the specific time needed to obtain the uniform distribution of hydrogen. After the exposure, the hydrogen contents of the specimens were measured under constant or rising temperatures by gas chromatographyemass spectroscopy. The hydrogen diffusivity was determined by fitting the solution of
a diffusion equation to the experimental hydrogen contents measured at various constant temperatures [26e28]. This method for determining the hydrogen diffusivity and solubility is a desorption method [28], which differs from the conventional permeation method [29].
Results and discussion Hydrogen diffusivity and solubility Fig. 1 shows Arrhenius plots of the hydrogen diffusivity and solubility of Type 304 obtained by the desorption method with high-pressure hydrogen. Data from the literature for the hydrogen diffusivity and solubility measured by permeation tests at a low pressure are shown in Fig. 1 [30e34]. The hydrogen diffusivity and solubility determined by the desorption method at a high pressure were nearly consistent with the literature data. The results shown in Fig. 1 suggest that reliable data for the hydrogen diffusivity and hydrogen solubility of austenitic stainless steels can be obtained by the desorption method with high-pressure hydrogen [28]. Fig. 2 shows the hydrogen diffusivity and hydrogen solubility of ten types of austenitic stainless steels determined by the desorption method at a high pressure. Interestingly, the hydrogen diffusivity of all the stainless steels was the same; thus, its temperature dependence was fitted by the following Arrhenius-type equation: D ¼ D0 expð ED =RTÞ; 6
(1) 2
1
1
where D0 ¼ 3.91 10 m s , ED ¼ 60.4 kJ mol , R is the gas constant, and T is the absolute temperature. At RT, the hydrogen diffusivity was ~1 1016 m2 s1, which is six orders
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 1 e Hydrogen diffusivity and solubility of Type 304: a hydrogen diffusivity; b hydrogen solubility.
Fig. 2 e Hydrogen diffusivity and solubility of all stainless steels: a hydrogen diffusivity; b hydrogen solubility. of magnitude lower than that of quenched and tempered Cre Mo steels [28]. In contrast to the hydrogen diffusivity, the hydrogen solubility differed slightly according to the type of steel; i.e., the hydrogen solubility of the high-strength, nitrogen-added steelsdXM-19 and HP160dwas slightly higher than that of the 300-series stainless steels, although the activation energy for the hydrogen solubility (ES) was independent of the type of steel. Under the assumption that ES was independent of the type of steel, the temperature dependence of the hydrogen solubility was fitted by the Arrhenius-type equation: S ¼ S0 expð ES =RTÞ;
(2) 1/2
1
where S0 ¼ 13.3 mass ppm MPa and ES ¼ 2.16 kJ mol for the 300-series stainless steels; S0 ¼ 20.5 mass ppm MPa1/2 and ES ¼ 2.16 kJ mol1 for HP160; and S0 ¼ 24.4 mass ppm MPa1/2 and ES ¼ 2.16 kJ mol1 for XM-19. According to the measured hydrogen diffusivity and solubility, the distributions of the hydrogen concentration in the specimens with external hydrogen and internal hydrogen
were estimated by using a solution of Fick's second law. Especially, when hydrogen exists only near the specimen surface, its distribution can be approximately estimated by using a solution of Fick's second law for hydrogen entry into a semi-infinite plate, as follows: oi h n c ¼ c0 þ ðcs c0 Þ 1 erf z=ð2ðDtÞÞ1=2 ;
(3)
where c is the hydrogen concentration of the specimen at z and t, c0 is the initial hydrogen concentration of the specimen, cs is the saturated hydrogen concentration, z is the penetration depth of hydrogen in the specimen, t is the exposure time or the tested time, and erf is the error function. Fig. 3 shows the estimated hydrogen concentrations of the 300-series austenitic stainless steel near the surface of the specimens with external hydrogen and internal hydrogen. Hydrogen exposure for 200 h at a pressure of 100 MPa and a temperature of 543 K (internal hydrogen) yielded a nearly uniform distribution of hydrogen in the SSRT specimen. The validity of the estimated hydrogen concentration was experimentally verified by secondary ion mass spectroscopy [18].
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y x x x ( 2 0 1 7 ) 1 e1 1
Fig. 3 e Estimated distribution of hydrogen concentration near surface of specimens with internal hydrogen (100 MPa and 543 K for 200 h) and external hydrogen (100 MPa and RT for 3 h).
On the other hand, in the case of hydrogen exposure for 3 h at a pressure of 100 MPa at RT (external hydrogen), there was a high concentration of hydrogen only at the surface of the SSRT specimen; i.e., a uniform distribution of hydrogen was not obtained. The distributions of the hydrogen concentration in the specimens and related boundary conditions differed significantly between the cases of internal and external hydrogen.
SSRT properties in presence of hydrogen Fig. 4 shows the nominal (engineering) stressestrain curves of the 300-series stainless steels tested at RT. Among the hydrogen-charged specimens tested in air, all the stainless steels except for Types 304(N) and 316L failed before reaching the TS of the non-charged specimens tested in air. Similarly, among the non-charged specimens tested in high-pressure hydrogen gas, all the stainless steels except for Types 304(N), 316, and 316L failed before reaching the TS of the non-charged specimens tested in air. The RRAs of Types 304, 316, and 316L have already been investigated by using the Nieq [17]. Although the concept of Nieq is discussed controversially in the community, this study mainly uses 300-series austenitic stainless steels (Nieq ¼ 22.8e38.5%); therefore, the RRA was investigated based on the concept of Nieq. Fig. 5 shows the relationship between the RRA and Nieq for all the stainless steel. The Nieq was calculated by a modified Sanga's equation [24,35]: Nieq ¼ Ni þ 12:93 C þ 1:11 Mn þ 0:72 Cr þ 0:88 Mo 0:27 Si þ 7:55 N
(4)
The stainless steels with a smaller Nieq (Nieq z 23%) were not influenced by the hydrogen distribution in the specimen (external or internal hydrogen). In contrast, the stainless steels with a larger Nieq (Nieq 28%) were influenced by the hydrogen distribution. The effects of external and internal hydrogen on the RRAs depended on the Nieq. This interesting
5
phenomenon was interpreted according to the observation of the tensile fracture morphologies, which is discussed later. Fig. 5 also contains the literature data of 300-series austenitic stainless steels with external and internal hydrogen [14]. The RRAs at 223 K were lower in 40-MPa external hydrogen compared to internal hydrogen; however, the RRAs at RT were almost the same between 70-MPa external hydrogen and internal hydrogen [14]. The RRAs for external and internal hydrogen in the literature reasonably agreed with those for external hydrogen in this study; however, seemed to be greater than those for internal hydrogen in this study. For internal hydrogen, the literature used 34.5-MPa hydrogen gas at 573 K (cs ~ 60 mass ppm), whereas this study used 100-MPa hydrogen gas at 543 K (cs ~ 100 mass ppm). Since the saturated hydrogen concentration in this study was 1.7 times higher than that in the literature for internal hydrogen, the RRAs for internal hydrogen in this study were presumed to be lower than those in the literature.
Tensile fracture morphologies by scanning electron microscopy (SEM) Steel with a smaller Nieq: Type 304 Fig. 6 shows SEM micrographs of the tensile fracture morphologies of Type 304 at RT. In the non-charged specimen tested in air, the specimen failed because of an ordinary cupand-cone fracture dominated by microvoid coalescence (MVC). The fracture surface consisted of a normal stressfracture region covered with equiaxial dimple (A in Fig. 6(a)) and a shear stress-fracture region covered with elongated dimple (B in Fig. 6(a)). In the hydrogen-charged specimen tested in air, numerous surface cracks were observed, and the specimen failed because of the extension of these cracks. The fracture surface was dominated by quasi-cleavage, differing completely from that of the non-charged specimen tested in air. We consider that the internal hydrogen at the crack tip underwent stressinduced diffusion, extending the surface cracks. Our previous study demonstrates that, for austenitic stainless steels used in hydrogen gas, the SSRT surface crack grew via the same mechanism as for JIC and fatigue cracks based on the HISCG model [22]. Accordingly, the HIS crack is ductile, not brittle. The results shown in Fig. 6 suggest that the degradation in the RRA of Type 304 with internal hydrogen was caused by extension of surface cracks due to hydrogen associated with the HISCG [22]. The non-charged specimen tested in hydrogen gas also failed because of the extension of surface cracks, which was assisted by external hydrogen, and the fracture surface was dominated by quasi-cleavage. This morphology reveals that the degradation in the RA of Type 304 with external hydrogen, as well as that of Type 304 with internal hydrogen, was caused by the HISCG.
Steel with a larger Nieq: Type 316L Fig. 7 shows SEM micrographs of the tensile fracture morphologies of Type 316L. All of the specimens in Fig. 7 failed because of cup-and-cone fractures dominated by the MVC. In contrast to Type 304 with internal and external hydrogen, the HISCG did not dominate the fracture of Type 316L with
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Fig. 4 e Effect of external or internal hydrogen on SSRT properties of austenitic stainless steels at RT: a Type 304N2; b Type 321; c Type 304(V); d Type 304(N); e Type 304(Nb); f Type 304; g Type 316; h Type 316L.
internal and external hydrogen. As shown in Fig. 2, the hydrogen diffusivity and hydrogen solubility of Types 304 and 316L were the same; accordingly, the difference in the fracture-surface morphologies between Types 304 and 316L
cannot be explained according to the hydrogen distribution in the specimen. It is speculated that the HISCG is attributed to the combined effect of the strain-induced martensitic transformation and the hydrogen distribution near the crack tip. In
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Comparing the normal stress-fracture regions of Type 316L between the cases of external and internal hydrogen, a reduction in the dimple size was observed for internal hydrogen, whereas such a reduction was not observed for external hydrogen. Matsuo and Yamabe et al. [19] attributed the small dimple of Type 316L with internal hydrogen to void sheets produced by local shear strain [36,37]. The localization of this shear strain was enhanced by internal hydrogen, resulting in a slight reduction in the RA. In contrast, in Type 316L with external hydrogen, hydrogen did not exist in the normal stress-fracture region, as suggested by Fig. 3. Therefore, the external hydrogen did not affect the dimple morphology of the normal stress-fracture region, causing a negligible reduction in the RA.
Steel with a medium Nieq: Type 304(Nb)
Fig. 5 e Relationship between RRA and nickel-equivalent content of austenitic stainless steels at RT, together with literature data [14].
Type 316L, the strain-induced martensitic transformation was prevented by the high stability of the austenitic phase; as a result, the HISCG was not produced during the SSRT test, and the specimen failed because of the cup-and-cone fracture dominated by the MVC.
Fig. 8 shows SEM micrographs of the tensile fracture morphologies of Type 304(Nb) at RT. The non-charged specimen tested in air failed because of the cup-and-cone fracture, consisting of the normal stress-fracture region associated with equiaxial dimple and the shear stress-fracture region associated with elongated dimple. In the hydrogen-charged specimen tested in air, the fracture surface was covered with a mixture of equiaxial dimple and quasi-cleavage. In the non-charged specimen tested in hydrogen gas, the fracture surface was covered with a mixture of elongated dimple and quasi-cleavage. The fracture of the stainless steels having an RRA in the transition region competed with MVC and HISCG in the presence of hydrogen.
Fig. 6 e SEM micrographs of tensile fracture morphologies of Type 304 at RT. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 7 e SEM micrographs of tensile fracture morphologies of Type 316L at RT.
Fig. 8 e SEM micrographs of tensile fracture morphologies of Type 304(Nb) at RT. Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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Fig. 9 e Tensile fracture processes of austenitic stainless steels with external and internal hydrogen: a Fracture dominated by HISCG; b Ordinary cup-and-cone fracture.
Effects of external and internal hydrogen on RRA As shown in Fig. 5, the RRAs of the steels with a smaller Nieq were not influenced by the hydrogen distribution, whereas those of the steels with a larger Nieq were influenced by the hydrogen distribution. In this section, we discuss the effect of the hydrogen distribution on the RRAs. In the stainless steels with a smaller Nieq, the degradation in the RRA was dominated by the HISCG. Localized hydrogen at the crack tip was not directly identified in this study; however, it is speculated that both internal hydrogen and external hydrogen diffused at the crack tip during the SSRT test and contributed to the HISCG, as shown in Fig. 9. Thus, the RRA of the steels with a smaller Nieq was barely influenced by the hydrogen distribution. In the steels with a larger Nieq, the specimens failed because of the cup-and-cone fracture dominated by the MVC. The hydrogen affected the void formation in the normal stress-fracture region, slightly degrading the RRA. Only internal hydrogen contributed to the void formation in the normal stress-fracture region, because the SSRT specimen with internal hydrogen had a uniform distribution of hydrogen. Understanding the hydrogen distribution in the specimen and the tensile-fracture process (MVC or HISCG) is important for interpreting the effect of the hydrogen distribution on solution-treated austenitic stainless steels.
Conclusion SSRT tests for ten types of solution-treated austenitic stainless steels with external and internal hydrogen were performed at RT. Five types of 300-series austenitic stainless steels containing a small amount of additive elementsdTypes 304N2, 321, 304(V), 304(N), and 304(Nb)dwere prepared. For comparison, three types of conventional 300-series austenitic stainless steelsdTypes 304, 316, and 316Ldand two types of high-strength, nitrogen-added austenitic stainless steelsdHP160 and XM-19dwere also tested. The hydrogen diffusivity and solubility of these steels were determined by the desorption method with high-pressure gaseous hydrogen. Our conclusions are summarized as follows: 1. The measured hydrogen diffusivity of all the steels was the same, which is consistent with literature data obtained by conventional permeation tests at a low pressure. In contrast, the measured hydrogen solubility was found to slightly depend on the type of steel, and the hydrogen solubility of the high-strength, nitrogen-added austenitic stainless steels HP160 and XM-19 was slightly higher than that of the 300-series austenitic stainless steels. The hydrogen solubility of the 300-series steels was not affected by a small amount of additive elements.
Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055
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2. All the 300-series stainless steels except for Types 304(N), 316, and 316L failed before reaching the TS of the non-charged specimens tested in air with internal or external hydrogen. 3. All the stainless steels failed because of ductile cup-andcone fracture dominated by MVC in the absence of hydrogen, and several stainless steels failed because of quasi-cleavage fracture dominated by HISCG in the presence of hydrogen. In the presence of hydrogen, a significant reduction in the RRA was caused by the HISCG. 4. The stainless steels with a higher nickel-equivalent content (Nieq) had a higher RRA. This suggests that the RRA of the present austenitic stainless steels can be quantified according to the stability of the austenitic phase. 5. The RRA of the steels with a larger Nieq was influenced by the hydrogen distribution in the specimen (external or internal hydrogen), whereas that of the steels with a smaller Nieq was not. These unique trends can be interpreted according to the hydrogen distribution in the SSRT specimen and the tensile-fracture morphology (HISCG or MVC).
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Acknowledgment [17]
This work was partially supported by the New Energy and Industrial Technology Development Organization (NEDO), Hydrogen Utilization Technology (2013e2018).
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Please cite this article in press as: Yamabe J, et al., Hydrogen diffusivity and tensile-ductility loss of solution-treated austenitic stainless steels with external and internal hydrogen, International Journal of Hydrogen Energy (2017), http://dx.doi.org/10.1016/ j.ijhydene.2017.04.055