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Investigation of hydrogen transport behavior of various low-alloy steels with high-pressure hydrogen gas
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 5 ) 1 e1 2
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Investigation of hydrogen transport behavior of various low-alloy steels with high-pressure hydrogen gas Junichiro Yamabe a,b,c,*, Tohru Awane b,d, Saburo Matsuoka b a
International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 8190395, Japan b Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan c International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan d Kobe Material Testing Laboratory Co., Ltd., 47-13 Nii-jima, Harima-cho, Kako-gun, Hyogo 675-0155, Japan
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abstract
Article history:
The apparent hydrogen diffusivity and the saturated hydrogen content of CreMo and Ni
Received 22 May 2015
eCreMo steels were determined with high-pressure hydrogen gas. Surface effects on
Received in revised form
hydrogen entry and exit were also investigated by using palladium-coated samples and by
29 June 2015
diffusion analysis using the finite-element method. Hydrogen contents of hydrogen-
Accepted 3 July 2015
exposed cylindrical specimens of various sizes were measured by means of gas chroma-
Available online xxx
tographyemass spectrometry to obtain the saturated hydrogen content. The diffusivity was determined by fitting the solution of a diffusion equation to the experimental
Keywords:
hydrogen contents determined by desorption at various constant temperatures. In the
High-pressure hydrogen gas
specimens examined, surface effects were significant at room temperature. The temper-
Diffusivity
ature dependences of the diffusivity were reasonably consistent with reference data
Solubility
mainly measured with electrochemical charging. These results were interpreted in terms
Surface effect
of hydrogen trapping. Ordinary electrochemical charging represents a more severe con-
Low-alloy steel
dition than exposure to high-pressure hydrogen, for example, at 100 MPa. Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Fuel-cell vehicles have recently become commercialized, and the necessary hydrogen filling stations are gradually being built. In terms of safety, these fuel-cell vehicles and their
associated hydrogen stations have been designed to use metallic materials with a high resistance to hydrogen embrittlement, such as type 316L austenitic stainless steel and A6061-T6 aluminum alloy. However, such metals have lower strengths and are more expensive than conventional steels such as carbon or low-alloy steels. For the widespread
* Corresponding author. International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka, Japan Zip code: 819-0395. Tel.: þ81 92 802 3247. E-mail address: [email protected] (J. Yamabe). http://dx.doi.org/10.1016/j.ijhydene.2015.07.006 0360-3199/Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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commercialization of hydrogen-energy systems, expected to occur in the near future, conventional steels that have higher strengths and lower costs will have to be widely used. Consequently, many experiments on the tensile, fatigue, and hydrogen-diffusion properties of such materials in highpressure hydrogen gas need to be performed, and appropriate design methods will have to be established on the basis of the resulting scientific understanding [1e3]. Yamabe et al. [3] investigated the effect of hydrogen on the fatigue-crack growth (FCG) and fracture-toughness properties of CreMo steels. We found that the effects of hydrogen on these properties were dependent on the tensile strength and microstructure of the materials. Furthermore, we showed that the rate of FCG was accelerated by hydrogen, but there was an upper limit on the hydrogen-accelerated FCG rate for CreMo steels with a tensile strength of less than about 900 MPa. In contrast, Matsuo et al. [4] and Nishiguchi et al. [5] reported that the relative reduction in area of carbon steels decreased with increasing hydrogen content. The susceptibility of steels to hydrogen embrittlement is therefore believed to be related to the distribution of dissolved hydrogen, irrespective of the fact that precharged specimens were tested in air and noncharged specimens were tested in gaseous hydrogen [6,7]. The macroscopic hydrogen distribution in the steels under nonloaded conditions can be calculated from their hydrogen diffusivity and their saturated hydrogen content, provided that hydrogen entry and exit occur by a diffusion-controlled process. It has also been reported that the hydrogen diffusivity of various steels decreases with increasing tensile strength [8]. There have been many studies on the hydrogen-diffusion behavior of metals that are relevant to the practical significance of this behavior [9e26]. Studies have also been made on the hydrogen-diffusion behavior of low-alloy steels [21e26]; however, in these studies, only hydrogen diffusivity was determined mainly with electrochemical charging [21e24]. Consequently, the hydrogen content of such steels exposed to high-pressure hydrogen gas has not been determined directly. It is also unclear whether or not the state of steels subjected to electrochemical charging is identical to that of steels exposed to hydrogen gas [27,28]. To interpret the results of tensile and fatigue tests in high-pressure hydrogen gas precisely, it is essential to understand the distribution of hydrogen dissolved under practical conditions involving gaseous hydrogen, rather than that produced by electrochemical charging. Mine et al. [29,30] reported a method for the determination of hydrogen diffusivity and saturated hydrogen content by using discshaped specimens of various thicknesses exposed to highpressure hydrogen gas. They applied this method to austenitic stainless steels (g-stainless steel) and they found that the hydrogen diffusivity and solubility were reasonably consistent with the reference values determined by a permeation method with low-pressure hydrogen gas. However, their method mainly used precharged disc-shaped specimens with a thickness of less than 1 mm. Consequently, because the hydrogen diffusivity of carbon and low-alloy steels is much larger than that of g-stainless steels, almost all the hydrogen desorbed from the specimens before their hydrogen content could be measured. Mine's method is therefore unsuitable for determining the hydrogen diffusivity and solubility of carbon and low-alloy steels.
Our current study focused on the desorption of hydrogen from samples exposed to high-pressure hydrogen gas. We used a desorption-based technique to determine the temperature dependence of the hydrogen diffusivity as well as the saturated hydrogen content of CreMo and NieCreMo steel samples, most of which were taken from storage cylinders used for hydrogen filling stations. Five low-alloy steels with tensile strengths of around 900 MPa were used, and cylindrical specimens with 2r0 ¼ z0 ¼ 7, 10, 15, and 19 mm, where r0 was the radius of the specimen and z0 was its thickness, were sampled from the steels. These cylindrical specimens were larger than those used by Mine et al.; consequently, rapid desorption of hydrogen from the specimens was suppressed. The specimens were exposed to highpressure hydrogen gas at 100 MPa and at room temperature (RT; ~25 C), 50 C, or 85 C for 200e300 h. After the hydrogen exposure, the hydrogen that desorbed from the specimens was measured by means of gas chromatographyemass spectroscopy (GCeMS) at a rising or constant temperature. The hydrogen contents of specimens of different sizes were measured to identify the specimen size necessary to achieve saturation with hydrogen. The hydrogen diffusivity was determined by fitting the solution of the nonsteady diffusion equation to the experimental hydrogen-desorption profile by the least-squares method. Surface effects on the entry and exit of hydrogen were also investigated by using palladium-sputtered low-alloy steel [14,19]. To interpret these surface effects, we used a finiteelement method (FEM) to perform a diffusion analysis in which the surface effects were taken into account. The results that we obtained are useful in understanding tensile and fatigue data for high-pressure hydrogen gas; moreover, they permit the estimation of the amount of hydrogen that will permeate into storage cylinders at hydrogen stations or into pipes used to carry high-pressure hydrogen gas.
Experimental procedures Materials Four JISeSCM435 steels (CreMo steels, Materials A to D) and one JISeSNCM439 steel (NieCreMo steel, Material E) were used in this study. Table 1 lists the chemical composition and Vickers hardness for each of these steels. Materials A, B, C, and E were sampled from real storage cylinders for hydrogen stations, whereas Material D was sampled from a round bar. These steels were quenched and tempered.
Specimens and hydrogen exposure To clarify the effects of the specimen size on the hydrogen contents of steels exposed to hydrogen, four cylindrical specimens with 2r0 ¼ z0 ¼ 7, 10, 15, and 19 mm were sampled from each of the steels. The surfaces of these specimens were finished with #2000 emery paper. Identical experimental results were obtained when the surfaces were finished with #600 emery paper. Based on our previous study [31], a native oxide layer with a few nanometers thickness is formed on the surface and surface reactions of hydrogen, such as
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Table 1 e Chemical composition (mass%) and Vickers hardness (HV) of low-alloy steels. Material CreMoa
Symbol
A B C D Requirement (JIS G 3441) NieCreMob E Requirement (JIS G 4053) a b c d
C
Si
Mn
P
S
Cr
Mo
Ni
HV
0.38 0.37 0.35 0.36 0.33e0.38 0.43 0.36e0.43
0.22 0.21 0.25 0.18 0.15e0.35 0.27 0.15e0.35
0.79 0.77 0.74 0.78 0.60e0.85 0.82 0.60e0.90
0.006 0.012 0.011 0.013 0.030 0.005 0.030
0.004 0.007 0.004 0.005 0.030 0.002 0.030
1.10 1.07 1.08 1.04 0.90e1.20 0.91 0.60e1.00
0.23 0.23 0.26 0.20 0.15e0.30 0.23 0.15e0.30
ec ec ec ec 0.25 1.95 1.60e2.00
256 289 275 258 ed 292 ed
JISeSCM435. JISeSNCM439. Not measured. Not regulated.
dissociation and recombination, may affects the measured hydrogen diffusivity. To examine effects of such surface reactions on the entry and exit of hydrogen gas, a palladium layer approximately 10 nm thick was deposited by ion sputtering on the surface of some of the specimens. It is noted that the native oxide layer remains at the subsurface of the palladium-sputtered specimens, whereas the surface is covered with the palladium layer [31]. All specimens were exposed to high-pressure hydrogen gas at 100 MPa and a temperature of RT, 50 C, or 85 C for around 250 h. To prevent unwanted hydrogen desorption, the treated samples were kept under liquid nitrogen until their hydrogen content was measured.
Observation of microstructures
Measurement of hydrogen contents
Fig. 1 shows the microstructures of Materials A to E, as analyzed by SEMeEBSD. Material B had the coarsest microstructure among the CreMo steels (Materials A to D). The NieCreMo steel (Material E) had a finer microstructure than any of the CreMo steels (Materials A to D).
The hydrogen contents were measured by GCeMS at an increasing or a constant temperature. In the measurement performed with an increasing temperature, the hydrogen that desorbed from the specimens at temperatures from RT to 520 C was detected as the temperature was increased at a rate of 100 C per hour. In the measurements at constant temperature, the residual hydrogen content was determined after various elapsed times at a constant temperature. The time between hydrogen exposure and the start of the hydrogen-content measurement was defined as the preparation time. To determine the amount of hydrogen that desorbed from the specimens during the preparation time, cylindrical specimens with 2r0 ¼ z0 ¼ 19 mm sampled from Material C were exposed to high-pressure hydrogen gas at 100 MPa and 85 C for 240 h as a preliminary investigation. After the exposure, the specimens were immediately removed from the pressure vessel and their hydrogen contents were measured by GCeMS (Test 1). It took seven minutes to remove the specimen, and its hydrogen content was 0.51 mass ppm. In a second test (Test 2), the specimen was removed after cooling the pressure vessel. It took three hours to remove the specimen and its hydrogen content was 0.49 mass ppm. The hydrogen contents determined in Tests 1 and 2 were almost equal; accordingly, we adopted the method used in Test 2 for our subsequent studies. Because all our preliminary investigation were performed at temperatures of 85 C or less, the saturated hydrogen contents of specimens exposed to hydrogen gas were also measured at temperatures of 85 C or below.
We analyzed the microstructures of the materials by means of scanning electron microscopy (SEM; Hitachi, FB-2000A) with electron backscatter diffraction (EBSD). The acceleration voltage for the SEMeEBSD was 20 kV. The specimen surface was polished by buffing with colloidal silica with a particle size of 0.04 mm.
Results and discussion Microstructures
Thermal desorption analysis of control and hydrogenexposed specimens Fig. 2(a) shows the profile of the hydrogen desorption rate for Material C (CreMo steel) after exposure to high-pressure hydrogen at 100 MPa and 85 C for 240 h. The units for the hydrogen desorption rate are mass ppm per minute, and the size of the cylindrical specimen was 2r0 ¼ z0 ¼ 19 mm. For comparison, the profile of the hydrogen desorption rate of a nonexposed (control) specimen is also provided in this figure. For the hydrogen-exposed specimen, two peaks were observed at temperatures of ~100 C and ~400 C, respectively, whereas only the second of these peaks was observed for the nonexposed specimen. The hydrogen content of the hydrogen-exposed specimen was 0.47 mass ppm, which was nearly equal to that corresponding to the first peak (0.44 mass ppm). The hydrogen content of the nonexposed specimen, which was equivalent to the hydrogen content associated with the second peak, was 0.01 mass ppm. Fig. 2(b) shows the profile of the rate of hydrogen desorption for Material E (NieCreMo steel) after exposure to highpressure hydrogen at 100 MPa and 85 C for 240 h. The size of the cylindrical specimen was 2r0 ¼ z0 ¼ 19 mm. The hydrogen-exposed specimen sampled from Material E, like that of Material C, showed the first and second peaks, and its
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Fig. 1 e Microstructures analyzed by SEMeEBSD: (a) Material A, (b) Material B, (c) Material C, (d) Material D, (e) Material E.
Fig. 2 e Profiles of the hydrogen desorption rates of Materials C and E exposed to high-pressure gaseous hydrogen: (a) Material C, (b) Material E.
measured hydrogen content was 0.45 mass ppm. The nonexposed specimen of Material E similarly showed only the second peak and the associated hydrogen content was 0.01 mass ppm. The other materials showed similar hydrogen desorption rate profiles. Novak et al. [32] reported the profile of the hydrogen desorption rate for AISI4340 steel quenched at 800 C for 1 h and tempered at 200 C for 2 h after exposure to high-pressure hydrogen at 97 MPa and 85 C for 210 h; this material was similar to Material E and, consequently, showed similar peaks. The activation energies for the first and second peaks, estimated from the profiles of the hydrogen desorption rate at various heating rates, indicated that these peaks were associated with the elastic strain field of dislocations and with metal carbides, respectively [32].
Effect of specimen size on hydrogen content Next, we compared the hydrogen contents associated with the first peak CH1 for specimens of various sizes. Fig. 3 is a plot of
the value of CH1 against the thickness of the specimen z0 for several cylindrical specimens of Materials A, B, C, and E exposed to high-pressure hydrogen at 100 MPa and 85 C for 240 h. Irrespective of the material, CH1 decreased with decreasing z0 for specimens with 2r0 ¼ z0 10 mm, whereas the values of CH1 for specimens with 2r0 ¼ z0 15 mm were nearly constant. This result shows that the hydrogen desorbed from the specimens during the preparation time hardly affected the measured CH1 value for specimens with 2r0 ¼ z0 15 mm. The experimental result that CH1 was independent of the specimen size showed that the hydrogen concentration in the specimen was uniform. A similar result was obtained with specimens exposed to high-pressure hydrogen gas at 100 MPa and 50 C. For the case in which CH1 >> CH2, where CH2 is the hydrogen content associated with the second peak, the value of CH1 obtained from specimens with 2r0 ¼ z0 15 mm after exposure to hydrogen at 50 or 85 C was regarded as being equivalent to the saturated hydrogen content, CHS, in the present study. The values of CHS at 100 MPa and 85 C were 0.43, 0.56, 0.49, 0.43, and
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Temperature dependence of the hydrogen diffusivity We will consider hydrogen-exposed cylindrical specimens saturated with hydrogen. When exposure to hydrogen ceases, the amount of hydrogen that is dissolved in the specimen decreases with elapsed time as the hydrogen desorbs. In our case, the residual hydrogen content CHR(tR) at an elapsed time tR can be calculated by solving the following nonsteady diffusion equation [33,34]: h i) ( 2 ∞ exp ð2n þ 1Þ p2 DtR z2 0 CHR ðtR Þ 32 X ¼ 2 CHS p ð2n þ 1Þ2 n¼0 ( ) ∞ X exp Db2m tR r20 $ b2m m¼1
Fig. 3 e Relationship between hydrogen content and specimen size for cylindrical specimens exposed to highpressure hydrogen gas.
where CHS is the saturated hydrogen content, D is the (effective) diffusivity of hydrogen, r0 is the radius of the specimen, z0 is the thickness of the specimen, and bm is the root of the zeroorder Bessel function. When tR >> 0, Eq. (1) can expressed by in the following simplified form: 2 CHR ðtR Þ 32 p b2 z 2 2 exp þ 21 DtR 2 CHS z0 r0 p b1
0.48 mass ppm for Materials A, B, C, D, and E, respectively. Similarly, the values of CHS at 100 MPa and 50 C were 0.42, 0.52, 0.48, 0.42, and 0.46 mass ppm for Materials A, B, C, D, and E, respectively. The resolution of GCeMS for hydrogencontent measurement is of the order of ppm and CHS tends to be slightly greater for materials having a higher tensile strength and a coarser microstructure. As discussed later, the entry and exit of hydrogen at RT are significantly affected by surface reactions of hydrogen, such as dissociation and recombination. Thus, a palladiumsputtered cylindrical specimen with 2r0 ¼ z0 ¼ 19 mm was prepared from Material D, and its hydrogen contents were measured after various exposure times to determine its saturated hydrogen content at RT as shown in Fig. 6.
(1)
(2)
where b1 ¼ 2.405. Furthermore, when 2r0 ¼ z0, which was the case for the cylindrical specimens used in the present study, the following equation is obtained: 2 CHR ðtR Þ 32 p þ 4b21 z 2 2 exp DtR : 2 CHS z0 p b1
(3)
Fig. 4(a) shows the relationship between CHR and tR for the cylindrical specimen with 2r0 ¼ z0 ¼ 15 mm sampled from Material E exposed to hydrogen at 100 MPa at a constant temperature of 85 C for 240 h. The first two plotted results were obtained during heating of the specimen from RT to 85 C. The experimental data were fitted to Eq. (1) by the least-squares method, with CHS and D as unknown parameters. The experimental results, apart from for the first
Fig. 4 e Hydrogen desorption behavior of cylindrical specimens sampled from Material E exposed to high-pressure hydrogen gas at (a) 85 C or (b) 50 C. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Fig. 5 e Arrhenius plot of hydrogen diffusivity of various low-alloy steels.
two or three measured results, were consistent with the curve fitted to Eq. (1). The hydrogen diffusivity D can therefore be determined by this method, whereas the saturated hydrogen content CHS cannot be determined by it, because at t z 0, CHR obtained during heating of the specimen was inconsistent with the fitted curve. In other words, the value of D can be determined by using the following h i) equation: ( ∞ 2 X exp ð2n þ 1Þ p2 DtR z20 CHR ðtR Þ ¼A ð2n þ 1Þ2 n¼0 (4) ( ) ∞ X exp Db2m tR r20 $ b2m m¼1
Fig. 6 e Plot of the hydrogen content against the exposure time for cylindrical specimens of Material D with and without palladium sputtering after exposure to hydrogen gas at 100 MPa and RT.
where A is the fitted constant. The hydrogen diffusivity determined by fitting with Eq. (4) was 1.4 109 m2/s. When tR >> 0, the experimental results were fitted by the following equation: 2 p þ 4b21 Dt CHR ðtR Þ ¼ B exp R z20
(5)
where B is the fitted constant. When tR 90 min, Eq. (5) agreed with Eq. (4) and it was shown that D ¼ 1.3 109 m2/s by fitting with Eq. (5); this value was almost equal to that obtained by fitting with Eq. (4). The curves fitted by Eqs. (4) and (5) are shown in the figure. Fig. 4(b) shows the relationship between CHR and tR for the cylindrical specimen of Material E with 2r0 ¼ z0 ¼ 19 mm exposed to hydrogen at 100 MPa at a constant temperature of 50 C for 240 h. The first two plotted results were obtained during the process of heating the specimen from RT to 50 C. The curves fitted with Eqs. (4) and (5) are also shown in the figure. As in the case of Fig. 4(a), the experimental results, apart for the first two or three values, were consistent with the fitted curve. We found that D ¼ 4.6 1010 m2/s by fitting with Eq. (4). For tR 280 min, Eq. (5) agreed with Eq. (4) and the value D ¼ 4.7 1010 m2/s was obtained by fitting with Eq. (5); this value was nearly equal to that obtained by fitting with Eq. (4). The hydrogen diffusivities of the other materials at various temperatures were similarly determined by this method. Fig. 5 shows an Arrhenius plot of the hydrogen diffusivities for Materials A to E. The values of the hydrogen diffusivities are the mean values obtained for specimens with 2r0 ¼ z0 ¼ 15 mm and 19 mm. The lattice hydrogen diffusivity in bcc iron without any traps [see Eqs. (6) and (7), below], as reported by Kiuchi et al. [18], and the reference hydrogen diffusivity of low-alloy steels determined by charging without high-pressure hydrogen gas, for example, by electrochemical charging [21e24], are also shown in the figure.
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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233 353 K ð40 80 o CÞ 5690 DL ¼ 7:28 108 exp RT
(6)
323 823K ð50 550 o CÞ
6700 7120 DL ¼ ð1 2:52Þ 107 exp RT
(7)
Here, DL is the lattice hydrogen diffusivity in bcc iron without any traps (m2/s), R is the gas constant (8.314 J mol1 K1), and T is the absolute temperature (K). The hydrogen diffusivities of the present low-alloy steels agreed closely with that reported by Fujii et al. [24] and were 5e15 times greater than the lattice hydrogen diffusivity calculated by using Eq. (6) or Eq. (7). The values of D for Materials A to E showed no large differences, although they tended to decrease for materials having a higher tensile strength and a coarser microstructure.
Entry and exit of high-pressure hydrogen gas in steels at room temperature As shown in Fig. 5, the diffusivities of hydrogen in Materials A to E at RT were approximately one order of magnitude less than that reported by Yamakawa et al. [23]. The value of the hydrogen diffusivity determined by Yamakawa et al. was obtained by an electrochemical permeation method in which the effects of surface reactions were carefully removed. The discrepancy between our results and those of Yamakawa implies that the hydrogen desorption at RT in our experiments did not occur by a diffusion-controlled process. We therefore determined the hydrogen diffusivity for palladiumsputtered cylindrical specimens of Material D with 2r0 ¼ z0 ¼ 19 mm. The hydrogen diffusivity of palladiumsputtered specimens of Material D at temperatures of RT, 50 C, and 85 C is shown in Fig. 5. The values of D at 50 C and 85 C were consistent with those for the nonsputtered specimens, whereas the value of D at RT for the palladiumsputtered specimen was larger than that for the nonsputtered one. However, the value for the sputtered specimen agreed with that reported by Yamakawa et al. Therefore, the hydrogen diffusivity at RT cannot be determined precisely by the present method using nonsputtered cylindrical specimens, because of the presence of a significant surface effect on the exit of hydrogen at RT. As well as the exit of hydrogen, entry of gas is also likely to be affected by surface reactions at RT. Therefore, palladiumsputtered and nonsputtered cylindrical specimens of Material D with 2r0 ¼ z0 ¼ 19 mm were exposed to high-pressure hydrogen at 100 MPa and RT for various times and their hydrogen contents were subsequently measured. Fig. 6 shows a plot of the hydrogen content against the exposure time for palladium-sputtered and nonsputtered cylindrical specimens exposed to hydrogen. The relationship between the hydrogen content and the exposure time, as calculated from the hydrogen diffusivity at RT (D ¼ 2.0 1010 m2/s) for a diffusion-controlled process, is also shown in the figure. For the palladium-sputtered cylindrical specimens, the exposure time required to obtain saturation by hydrogen was slightly
longer than the calculated value; however, a cylindrical specimen with 2r0 ¼ z0 ¼ 19 mm became saturated with hydrogen (0.48 mass ppm) on exposure to the high-pressure gas at RT for approximately 90 h. In contrast, even we exposed the nonsputtered specimen to high-pressure hydrogen at RT for approximately 300 h, we could not attain saturation. This confirmed that the entry of hydrogen into nonsputtered specimens, which contain native oxide layers on the surface, at RT does not occur by a diffusion-controlled process and that there is a significant surface effect on the entry of hydrogen. The hydrogen-entry behavior in highpressure hydrogen gas therefore differs from that in ordinary electrochemical charging [35].
Temperature dependence of the saturated hydrogen content Fig. 7 shows an Arrhenius plot of the saturated hydrogen content CHS for Materials A to E. The value of CHS at RT was also determined for a palladium-sputtered specimen of Material D. The results confirmed that palladium sputtering does not affect the value of CHS for specimens exposed to hydrogen at 100 MPa and 85 C. The lattice hydrogen content in bcc iron without any traps, as reported by Hirth [15], is also shown in the figure: qffiffiffi 3440 CLS ¼ 104:47 f exp T
(8)
where CLS is the lattice hydrogen content in bcc iron (mass ppm) and f is the fugacity (MPa). According to San Marchi et al. [19], the fugacity f can be expressed by using the hydrogen-gas pressure p as follows: bp RT
f ¼ p exp
(9)
where p is the pressure of the hydrogen gas (MPa) and b is a constant (1.584 105 m3 mol1) [19]. In the range RT to 85 C, the values of CHS for our low-alloy steels were not dependent on the temperature and they were 5e10 times larger than the values of CLS.
Interpretation of the temperature dependences of the hydrogen diffusivity and the saturated hydrogen content Fig. 8(a) and (b) show Arrhenius plots of the mean values of D and CHS, respectively, for Materials A to E. The experimental hydrogen diffusivity and saturated hydrogen content for Materials A to E showed no large differences; consequently, the temperature dependencies of the hydrogen diffusivity and the saturated hydrogen content were investigated by using the mean values of these properties. Figs. 5 and 7 suggest that the temperature dependence of D and CHS of the present low-alloy steels is affected by hydrogen trapping. The experimental values of D and CHS were therefore fitted in terms of Oriani's equilibrium theory [36] as follows: D¼
7:28 108 5690 ¼ exp RT EB EB 1 þ NNXL exp RT 1 þ NNXL exp RT DL
(10)
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Fig. 7 e Arrhenius plot of the saturated hydrogen contents of various low-alloy steels.
NX EB CLS ¼ 1þ exp NL RT
qffiffiffi NX EB 3440 f exp ¼ 104:47 1 þ exp T NL RT
CHS
(11)
where NL is the number of lattice sites per unit volume, NX is the number of trap sites per unit volume, and EB is the binding energy. Fitting by the least-squares method was performed because NX/NL and EB were unknown parameters. The experimental results for D and CHS were successfully fitted to Eqs. (10) and (11), giving the following parameters: NX/ NL ¼ 2.1 104 and EB ¼ 29.6 kJ/mol. The binding energy determined by the least-squares method has shown that the
hydrogen dissolved in low-alloy steels is mainly trapped at dislocations [37]. On the basis of Eq. (11), we compared the saturated hydrogen content of a CreMo steel in high-pressure hydrogen gas with that obtained by ordinary electrochemical charging. Matsuoka et al. [35] and Takeuchi et al. [38] reported, respectively, that the hydrogen contents of a CreMo steel saturated by hydrogen through electrochemical charging were 0.85 mass ppm for a sample charged by immersion in aqueous NH4SCN solution at 40 C for 12 h and 1.20 mass ppm for a sample cathodically charged in a mixed aqueous solution of NH4SCN and NaCl at a current density of 0.3 mA/cm2 for 24 h. From Eq. (11), the corresponding pressures of gaseous hydrogen at RT can be estimated to be 190 MPa and 250 MPa,
Fig. 8 e Interpretation of the temperature dependences of (a) the hydrogen diffusivity and (b) the saturated hydrogen content. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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respectively. It therefore follows that ordinary electrochemical charging is a more severe condition than exposure to hydrogen gas at a high pressure, for example, 100 MPa. In summary, the hydrogen diffusivity (more precisely, the effective hydrogen diffusivity) and the saturated hydrogen content of the present low-alloy steels at near-ambient temperatures can be determined by using Eqs. (10) and (11); however, it should be noted that significant surface effects on the entry and exit of hydrogen may occur in high-pressure hydrogen gas at RT in specimens containing native oxide layers on the surface.
Diffusion analysis by the finite-element method taking the surface effect into account As mentioned, the palladium-sputtered and the non-sputtered specimens both contain the native oxide layers. In the palladium-sputtered specimen, the native oxide layer remains at the subsurface of the specimen, whereas the surface is covered with the palladium layer. Nevertheless, the entry and exit of hydrogen at RT was significantly different between the palladium-sputtered and non-sputtered specimens, presumably caused by the surface reaction of hydrogen. Such surface reactions of hydrogen significantly affect the hydrogen diffusivity of low-alloy steels at RT as determined by the desorption method, whereas such surface reactions hardly affected the hydrogen diffusivity of g-stainless steel as determined by the same method. In the latter determination, we used a hydrogen-exposed cylindrical specimen with 2r0 ¼ 5 mm and z0 ¼ 0.5 mm. As a result, the value of D for g-stainless steel at RT was estimated to be approximately 1 1016 m2/s for both palladium-sputtered and nonsputtered specimens. In this case, it is presumed that the time required for hydrogen diffusion is much greater than that required for the surface reaction, because the D value of the g-stainless steel at RT (D ¼ 1 1016 m2/s) is six orders of magnitude greater than that for the low-alloy steels at RT (D ¼ 2 1010 m2/s). To interpret the difference in the surface effect between the low-alloy steels and the g-stainless steels, we performed a diffusion analysis by FEM, taking the surface reaction of hydrogen into account. Fig. 9 shows the specimen geometry and the boundary conditions for the FEM analysis. For the sake of simplicity, we performed a one-dimensional diffusion analysis for the cylindrical specimen shown Fig. 9 by using commercial FEM software (ANSYS, Ver. 14.5). In our analysis, we used a six-node triangular element (PLANE 35) with a mesh size of 0.02 mm. In diffusion analysis involving surface reaction of hydrogen, the local hydrogen content at the surface of the exit side csurf is assumed to be not zero [39], unlike the case of diffusion analysis without consideration of the surface reaction of hydrogen. Therefore, the value of csurf was determined so that the following equation was satisfied: kcsurf ¼ D
vcsurf vr
(12)
where k is the rate of the surface reaction, which is a constant. In the experiment results shown in Fig. 5, the values of D at RT estimated by using nonsputtered specimens ranged from 1.4 1011 m2/s to 7.1 1011 m2/s, whereas that estimated by
9
using a palladium-sputtered specimen was 2.0 1010 m2/s. The FEM analysis was performed under various k values to obtain 1.4 1011 m2/s and 7.1 1011 m2/s; consequently, it was obtained that k ¼ 1.0 108 m/s and 4.0 108 m/s for D ¼ 1.4 1011 m2/s and 7.1 1011 m2/s, respectively. As a reference, Fig. 10 shows the relationship between CHR/CHS and tR when D ¼ 2 1010 m2/s and k ¼ 3 108 m/s, as calculated by FEM. The broken lines in the figure were those fitted by Eq. (1), as the values of CHS and D were unknown. Fig. 11 shows the hydrogen distribution when D ¼ 2 1010 m2/s and k ¼ 3 108 m/s as calculated by FEM with and without consideration of the surface reaction of hydrogen. Because of the surface reaction of hydrogen, the local hydrogen content at the surface of the exit side was not zero. These analytical results show that the surface reaction of hydrogen results in an underestimation of the hydrogen diffusivity of the present low-alloy steels; moreover, the calculated results were consistent with the experimental ones shown in Fig. 5. Different from low-alloy steel, the value of D for g-stainless steel at RT was estimated to be approximately 1 1016 m2/s for both palladium-sputtered and non-sputtered specimens, though the g-stainless steel possesses a native oxide layer of a few nanometers thickness, similar to that in low-alloy steel [31]. To interpret the experimental results, the FEM analysis taking account of the surface reaction of hydrogen was performed for the conditions D ¼ 1.0 1016 m2/s and k ¼ 1.0 108 m/s. Unlike the analytical results shown in Fig. 10, the value of D was estimated to be 1.0 1016 m2/s, even when the surface reaction of hydrogen was taken into consideration. This analytical result was consistent with the experimental one for the g-stainless steel; it therefore follows that surface reaction of hydrogen does not necessarily affect the hydrogen diffusivity of materials with low hydrogen diffusivity at RT determined by the desorption method, even if a native oxide layer is formed on the surface.
Conclusions The hydrogen diffusivity D and the saturated hydrogen content CHS of various low-alloy steels, potentially useful in the widespread commercialization of hydrogen-energy systems, were determined over a wide range of temperatures, including near-ambient temperatures, by exposing cylindrical specimens of four JISeSCM435 steels (CreMo steels) and one JISeSNCM439 steel (Nie CreMo steel) to high-pressure hydrogen gas at 100 MPa and a temperature of RT, 50 C, or 85 C for about 250 h. The size of the specimens was 2r0 ¼ z0 ¼ 7, 10, 15, or 19 mm, where r0 is the radius of the specimen and z0 is its thickness. After the hydrogen exposure, the hydrogen contents of the specimens were measured by gas chromatography (GCeMS) with a rising or constant temperature. The hydrogen diffusivity of the materials was determined by fitting the solution of a nonsteady diffusion equation to the experimental hydrogen-release profiles at constant temperature by means of the least-squares method. The effects of surface reactions, such as hydrogen dissociation and recombination, on the entry and exit of hydrogen were also investigated by using palladium-sputtered cylindrical specimens and a FEM diffusion analysis in which surface
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Fig. 9 e Specimen geometry and boundary conditions for the diffusion analysis by FEM taking the surface effect into account.
(2)
(3)
Fig. 10 e Calculated relationship between CHR/CHS and tR when D ¼ 2.0 £ 10¡10 m2/s and k ¼ 3.0 £ 10¡8 m/s.
(4)
effects were taken into consideration. Our conclusions can be summarized as follows. (1) In the profiles for the hydrogen desorption rates of both CreMo and NieCreMo steels under conditions of rising temperature, two peaks were observed at about 100 C (first peak) and 400 C (second peak). Regardless of the material, the hydrogen content associated with the first
(5)
(6)
Fig. 11 e Hydrogen distribution when D ¼ 2.0 £ 10¡10 m2/s and k ¼ 3.0 £ 10¡8 m/s, as calculated by the FEM.
peak (CH1) was much larger than that associated with the second peak (CH2); in other words, CH1 >> CH2. Under our exposure conditions at 50 C and 85 C, the hydrogen content associated with the first peak (CH1) decreased with decreasing specimen size for specimens with 2r0 ¼ z0 10 mm, whereas the values of CH1 were independent of the specimen size for specimens with 2r0 ¼ z0 15 mm. Because CH1 >> CH2, we regarded the hydrogen content for the first peak of specimens with 2r0 ¼ z0 15 mm as being equivalent to the saturated hydrogen content at 50 C or 85 C. The hydrogen diffusivity determined with the nonsputtered cylindrical specimens at temperatures other than RT showed no large differences among the various materials and it showed a similar tendency to reference data determined without high-pressure hydrogen gas, for example, with electrochemical charging. At RT, a significant reduction was observed in the hydrogen diffusivity determined with nonsputtered specimens in comparison to the corresponding reference date determined by the electrochemical permeation method. By contrast, the hydrogen diffusivity at RT determined with the palladium-sputtered specimens agreed closely with the reference data. Similarly, the hydrogen-entry behavior was markedly changed by palladium sputtering. It was therefore obvious that surface effects on the entry and exit of hydrogen were significant at RT. Consequently, the saturated hydrogen content at RT was determined by using palladiumsputtered cylindrical specimens. The saturated hydrogen content of the present lowalloy steels was 5e10 times larger than the lattice hydrogen content in bcc iron without any traps, whereas the hydrogen diffusivity was 5e15 time less than the lattice hydrogen diffusion in bcc iron without any traps. The temperature dependencies of the present low-alloy steels were quantitatively explained in terms of hydrogen trapping. The binding energy of the trap site was estimated to be 29.6 kJ/mol, implying that the hydrogen dissolved in the steels was mainly trapped at dislocations. No underestimation of the hydrogen diffusivity at RT was observed in g-stainless steel. This is because of the longer time required for hydrogen diffusion in comparison to that required for surface reactions in this steel. The effects of surface reactions on the hydrogen diffusivity determined for low-alloy and g-stainless
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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 5 ) 1 e1 2
steels can be interpreted by means of diffusion analysis by FEM taking account of surface effects.
Acknowledgment This work was partially supported by the New Energy and Industrial Technology Development Organization(NEDO), Hydrogen Utilization Technology (2013e2018).
references
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[34] Matsuo T, Yamabe J, Furukawa H, Seki K, Shimizu K, Watanabe S, et al. Development of new strain gage for highpressure hydrogen gas use. Exp Mech 2014;54:431e42. [35] Matsuoka S, Homma N, Tanaka H, Fukushima Y, Murakami Y. Effect of hydrogen on tensile properties of 900MPa-class JIS-SCM435 low-alloy-steel for use in storage cylinder of hydrogen station. Nippon Kinzoku Gakkaishi 2006;70:1002e11. [36] Oriani RA. The diffusion and trapping of hydrogen in steel. Acta Metall 1970;18:147e57.
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Investigation of hydrogen transport behavior of various low-alloy steels with high-pressure hydrogen gas Junichiro Yamabe a,b,c,*, Tohru Awane b,d, Saburo Matsuoka b a
International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 8190395, Japan b Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan c International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan d Kobe Material Testing Laboratory Co., Ltd., 47-13 Nii-jima, Harima-cho, Kako-gun, Hyogo 675-0155, Japan
article info
abstract
Article history:
The apparent hydrogen diffusivity and the saturated hydrogen content of CreMo and Ni
Received 22 May 2015
eCreMo steels were determined with high-pressure hydrogen gas. Surface effects on
Received in revised form
hydrogen entry and exit were also investigated by using palladium-coated samples and by
29 June 2015
diffusion analysis using the finite-element method. Hydrogen contents of hydrogen-
Accepted 3 July 2015
exposed cylindrical specimens of various sizes were measured by means of gas chroma-
Available online xxx
tographyemass spectrometry to obtain the saturated hydrogen content. The diffusivity was determined by fitting the solution of a diffusion equation to the experimental
Keywords:
hydrogen contents determined by desorption at various constant temperatures. In the
High-pressure hydrogen gas
specimens examined, surface effects were significant at room temperature. The temper-
Diffusivity
ature dependences of the diffusivity were reasonably consistent with reference data
Solubility
mainly measured with electrochemical charging. These results were interpreted in terms
Surface effect
of hydrogen trapping. Ordinary electrochemical charging represents a more severe con-
Low-alloy steel
dition than exposure to high-pressure hydrogen, for example, at 100 MPa. Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Fuel-cell vehicles have recently become commercialized, and the necessary hydrogen filling stations are gradually being built. In terms of safety, these fuel-cell vehicles and their
associated hydrogen stations have been designed to use metallic materials with a high resistance to hydrogen embrittlement, such as type 316L austenitic stainless steel and A6061-T6 aluminum alloy. However, such metals have lower strengths and are more expensive than conventional steels such as carbon or low-alloy steels. For the widespread
* Corresponding author. International Research Center for Hydrogen Energy, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka, Japan Zip code: 819-0395. Tel.: þ81 92 802 3247. E-mail address: [email protected] (J. Yamabe). http://dx.doi.org/10.1016/j.ijhydene.2015.07.006 0360-3199/Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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commercialization of hydrogen-energy systems, expected to occur in the near future, conventional steels that have higher strengths and lower costs will have to be widely used. Consequently, many experiments on the tensile, fatigue, and hydrogen-diffusion properties of such materials in highpressure hydrogen gas need to be performed, and appropriate design methods will have to be established on the basis of the resulting scientific understanding [1e3]. Yamabe et al. [3] investigated the effect of hydrogen on the fatigue-crack growth (FCG) and fracture-toughness properties of CreMo steels. We found that the effects of hydrogen on these properties were dependent on the tensile strength and microstructure of the materials. Furthermore, we showed that the rate of FCG was accelerated by hydrogen, but there was an upper limit on the hydrogen-accelerated FCG rate for CreMo steels with a tensile strength of less than about 900 MPa. In contrast, Matsuo et al. [4] and Nishiguchi et al. [5] reported that the relative reduction in area of carbon steels decreased with increasing hydrogen content. The susceptibility of steels to hydrogen embrittlement is therefore believed to be related to the distribution of dissolved hydrogen, irrespective of the fact that precharged specimens were tested in air and noncharged specimens were tested in gaseous hydrogen [6,7]. The macroscopic hydrogen distribution in the steels under nonloaded conditions can be calculated from their hydrogen diffusivity and their saturated hydrogen content, provided that hydrogen entry and exit occur by a diffusion-controlled process. It has also been reported that the hydrogen diffusivity of various steels decreases with increasing tensile strength [8]. There have been many studies on the hydrogen-diffusion behavior of metals that are relevant to the practical significance of this behavior [9e26]. Studies have also been made on the hydrogen-diffusion behavior of low-alloy steels [21e26]; however, in these studies, only hydrogen diffusivity was determined mainly with electrochemical charging [21e24]. Consequently, the hydrogen content of such steels exposed to high-pressure hydrogen gas has not been determined directly. It is also unclear whether or not the state of steels subjected to electrochemical charging is identical to that of steels exposed to hydrogen gas [27,28]. To interpret the results of tensile and fatigue tests in high-pressure hydrogen gas precisely, it is essential to understand the distribution of hydrogen dissolved under practical conditions involving gaseous hydrogen, rather than that produced by electrochemical charging. Mine et al. [29,30] reported a method for the determination of hydrogen diffusivity and saturated hydrogen content by using discshaped specimens of various thicknesses exposed to highpressure hydrogen gas. They applied this method to austenitic stainless steels (g-stainless steel) and they found that the hydrogen diffusivity and solubility were reasonably consistent with the reference values determined by a permeation method with low-pressure hydrogen gas. However, their method mainly used precharged disc-shaped specimens with a thickness of less than 1 mm. Consequently, because the hydrogen diffusivity of carbon and low-alloy steels is much larger than that of g-stainless steels, almost all the hydrogen desorbed from the specimens before their hydrogen content could be measured. Mine's method is therefore unsuitable for determining the hydrogen diffusivity and solubility of carbon and low-alloy steels.
Our current study focused on the desorption of hydrogen from samples exposed to high-pressure hydrogen gas. We used a desorption-based technique to determine the temperature dependence of the hydrogen diffusivity as well as the saturated hydrogen content of CreMo and NieCreMo steel samples, most of which were taken from storage cylinders used for hydrogen filling stations. Five low-alloy steels with tensile strengths of around 900 MPa were used, and cylindrical specimens with 2r0 ¼ z0 ¼ 7, 10, 15, and 19 mm, where r0 was the radius of the specimen and z0 was its thickness, were sampled from the steels. These cylindrical specimens were larger than those used by Mine et al.; consequently, rapid desorption of hydrogen from the specimens was suppressed. The specimens were exposed to highpressure hydrogen gas at 100 MPa and at room temperature (RT; ~25 C), 50 C, or 85 C for 200e300 h. After the hydrogen exposure, the hydrogen that desorbed from the specimens was measured by means of gas chromatographyemass spectroscopy (GCeMS) at a rising or constant temperature. The hydrogen contents of specimens of different sizes were measured to identify the specimen size necessary to achieve saturation with hydrogen. The hydrogen diffusivity was determined by fitting the solution of the nonsteady diffusion equation to the experimental hydrogen-desorption profile by the least-squares method. Surface effects on the entry and exit of hydrogen were also investigated by using palladium-sputtered low-alloy steel [14,19]. To interpret these surface effects, we used a finiteelement method (FEM) to perform a diffusion analysis in which the surface effects were taken into account. The results that we obtained are useful in understanding tensile and fatigue data for high-pressure hydrogen gas; moreover, they permit the estimation of the amount of hydrogen that will permeate into storage cylinders at hydrogen stations or into pipes used to carry high-pressure hydrogen gas.
Experimental procedures Materials Four JISeSCM435 steels (CreMo steels, Materials A to D) and one JISeSNCM439 steel (NieCreMo steel, Material E) were used in this study. Table 1 lists the chemical composition and Vickers hardness for each of these steels. Materials A, B, C, and E were sampled from real storage cylinders for hydrogen stations, whereas Material D was sampled from a round bar. These steels were quenched and tempered.
Specimens and hydrogen exposure To clarify the effects of the specimen size on the hydrogen contents of steels exposed to hydrogen, four cylindrical specimens with 2r0 ¼ z0 ¼ 7, 10, 15, and 19 mm were sampled from each of the steels. The surfaces of these specimens were finished with #2000 emery paper. Identical experimental results were obtained when the surfaces were finished with #600 emery paper. Based on our previous study [31], a native oxide layer with a few nanometers thickness is formed on the surface and surface reactions of hydrogen, such as
Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Table 1 e Chemical composition (mass%) and Vickers hardness (HV) of low-alloy steels. Material CreMoa
Symbol
A B C D Requirement (JIS G 3441) NieCreMob E Requirement (JIS G 4053) a b c d
C
Si
Mn
P
S
Cr
Mo
Ni
HV
0.38 0.37 0.35 0.36 0.33e0.38 0.43 0.36e0.43
0.22 0.21 0.25 0.18 0.15e0.35 0.27 0.15e0.35
0.79 0.77 0.74 0.78 0.60e0.85 0.82 0.60e0.90
0.006 0.012 0.011 0.013 0.030 0.005 0.030
0.004 0.007 0.004 0.005 0.030 0.002 0.030
1.10 1.07 1.08 1.04 0.90e1.20 0.91 0.60e1.00
0.23 0.23 0.26 0.20 0.15e0.30 0.23 0.15e0.30
ec ec ec ec 0.25 1.95 1.60e2.00
256 289 275 258 ed 292 ed
JISeSCM435. JISeSNCM439. Not measured. Not regulated.
dissociation and recombination, may affects the measured hydrogen diffusivity. To examine effects of such surface reactions on the entry and exit of hydrogen gas, a palladium layer approximately 10 nm thick was deposited by ion sputtering on the surface of some of the specimens. It is noted that the native oxide layer remains at the subsurface of the palladium-sputtered specimens, whereas the surface is covered with the palladium layer [31]. All specimens were exposed to high-pressure hydrogen gas at 100 MPa and a temperature of RT, 50 C, or 85 C for around 250 h. To prevent unwanted hydrogen desorption, the treated samples were kept under liquid nitrogen until their hydrogen content was measured.
Observation of microstructures
Measurement of hydrogen contents
Fig. 1 shows the microstructures of Materials A to E, as analyzed by SEMeEBSD. Material B had the coarsest microstructure among the CreMo steels (Materials A to D). The NieCreMo steel (Material E) had a finer microstructure than any of the CreMo steels (Materials A to D).
The hydrogen contents were measured by GCeMS at an increasing or a constant temperature. In the measurement performed with an increasing temperature, the hydrogen that desorbed from the specimens at temperatures from RT to 520 C was detected as the temperature was increased at a rate of 100 C per hour. In the measurements at constant temperature, the residual hydrogen content was determined after various elapsed times at a constant temperature. The time between hydrogen exposure and the start of the hydrogen-content measurement was defined as the preparation time. To determine the amount of hydrogen that desorbed from the specimens during the preparation time, cylindrical specimens with 2r0 ¼ z0 ¼ 19 mm sampled from Material C were exposed to high-pressure hydrogen gas at 100 MPa and 85 C for 240 h as a preliminary investigation. After the exposure, the specimens were immediately removed from the pressure vessel and their hydrogen contents were measured by GCeMS (Test 1). It took seven minutes to remove the specimen, and its hydrogen content was 0.51 mass ppm. In a second test (Test 2), the specimen was removed after cooling the pressure vessel. It took three hours to remove the specimen and its hydrogen content was 0.49 mass ppm. The hydrogen contents determined in Tests 1 and 2 were almost equal; accordingly, we adopted the method used in Test 2 for our subsequent studies. Because all our preliminary investigation were performed at temperatures of 85 C or less, the saturated hydrogen contents of specimens exposed to hydrogen gas were also measured at temperatures of 85 C or below.
We analyzed the microstructures of the materials by means of scanning electron microscopy (SEM; Hitachi, FB-2000A) with electron backscatter diffraction (EBSD). The acceleration voltage for the SEMeEBSD was 20 kV. The specimen surface was polished by buffing with colloidal silica with a particle size of 0.04 mm.
Results and discussion Microstructures
Thermal desorption analysis of control and hydrogenexposed specimens Fig. 2(a) shows the profile of the hydrogen desorption rate for Material C (CreMo steel) after exposure to high-pressure hydrogen at 100 MPa and 85 C for 240 h. The units for the hydrogen desorption rate are mass ppm per minute, and the size of the cylindrical specimen was 2r0 ¼ z0 ¼ 19 mm. For comparison, the profile of the hydrogen desorption rate of a nonexposed (control) specimen is also provided in this figure. For the hydrogen-exposed specimen, two peaks were observed at temperatures of ~100 C and ~400 C, respectively, whereas only the second of these peaks was observed for the nonexposed specimen. The hydrogen content of the hydrogen-exposed specimen was 0.47 mass ppm, which was nearly equal to that corresponding to the first peak (0.44 mass ppm). The hydrogen content of the nonexposed specimen, which was equivalent to the hydrogen content associated with the second peak, was 0.01 mass ppm. Fig. 2(b) shows the profile of the rate of hydrogen desorption for Material E (NieCreMo steel) after exposure to highpressure hydrogen at 100 MPa and 85 C for 240 h. The size of the cylindrical specimen was 2r0 ¼ z0 ¼ 19 mm. The hydrogen-exposed specimen sampled from Material E, like that of Material C, showed the first and second peaks, and its
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Fig. 1 e Microstructures analyzed by SEMeEBSD: (a) Material A, (b) Material B, (c) Material C, (d) Material D, (e) Material E.
Fig. 2 e Profiles of the hydrogen desorption rates of Materials C and E exposed to high-pressure gaseous hydrogen: (a) Material C, (b) Material E.
measured hydrogen content was 0.45 mass ppm. The nonexposed specimen of Material E similarly showed only the second peak and the associated hydrogen content was 0.01 mass ppm. The other materials showed similar hydrogen desorption rate profiles. Novak et al. [32] reported the profile of the hydrogen desorption rate for AISI4340 steel quenched at 800 C for 1 h and tempered at 200 C for 2 h after exposure to high-pressure hydrogen at 97 MPa and 85 C for 210 h; this material was similar to Material E and, consequently, showed similar peaks. The activation energies for the first and second peaks, estimated from the profiles of the hydrogen desorption rate at various heating rates, indicated that these peaks were associated with the elastic strain field of dislocations and with metal carbides, respectively [32].
Effect of specimen size on hydrogen content Next, we compared the hydrogen contents associated with the first peak CH1 for specimens of various sizes. Fig. 3 is a plot of
the value of CH1 against the thickness of the specimen z0 for several cylindrical specimens of Materials A, B, C, and E exposed to high-pressure hydrogen at 100 MPa and 85 C for 240 h. Irrespective of the material, CH1 decreased with decreasing z0 for specimens with 2r0 ¼ z0 10 mm, whereas the values of CH1 for specimens with 2r0 ¼ z0 15 mm were nearly constant. This result shows that the hydrogen desorbed from the specimens during the preparation time hardly affected the measured CH1 value for specimens with 2r0 ¼ z0 15 mm. The experimental result that CH1 was independent of the specimen size showed that the hydrogen concentration in the specimen was uniform. A similar result was obtained with specimens exposed to high-pressure hydrogen gas at 100 MPa and 50 C. For the case in which CH1 >> CH2, where CH2 is the hydrogen content associated with the second peak, the value of CH1 obtained from specimens with 2r0 ¼ z0 15 mm after exposure to hydrogen at 50 or 85 C was regarded as being equivalent to the saturated hydrogen content, CHS, in the present study. The values of CHS at 100 MPa and 85 C were 0.43, 0.56, 0.49, 0.43, and
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Temperature dependence of the hydrogen diffusivity We will consider hydrogen-exposed cylindrical specimens saturated with hydrogen. When exposure to hydrogen ceases, the amount of hydrogen that is dissolved in the specimen decreases with elapsed time as the hydrogen desorbs. In our case, the residual hydrogen content CHR(tR) at an elapsed time tR can be calculated by solving the following nonsteady diffusion equation [33,34]: h i) ( 2 ∞ exp ð2n þ 1Þ p2 DtR z2 0 CHR ðtR Þ 32 X ¼ 2 CHS p ð2n þ 1Þ2 n¼0 ( ) ∞ X exp Db2m tR r20 $ b2m m¼1
Fig. 3 e Relationship between hydrogen content and specimen size for cylindrical specimens exposed to highpressure hydrogen gas.
where CHS is the saturated hydrogen content, D is the (effective) diffusivity of hydrogen, r0 is the radius of the specimen, z0 is the thickness of the specimen, and bm is the root of the zeroorder Bessel function. When tR >> 0, Eq. (1) can expressed by in the following simplified form: 2 CHR ðtR Þ 32 p b2 z 2 2 exp þ 21 DtR 2 CHS z0 r0 p b1
0.48 mass ppm for Materials A, B, C, D, and E, respectively. Similarly, the values of CHS at 100 MPa and 50 C were 0.42, 0.52, 0.48, 0.42, and 0.46 mass ppm for Materials A, B, C, D, and E, respectively. The resolution of GCeMS for hydrogencontent measurement is of the order of ppm and CHS tends to be slightly greater for materials having a higher tensile strength and a coarser microstructure. As discussed later, the entry and exit of hydrogen at RT are significantly affected by surface reactions of hydrogen, such as dissociation and recombination. Thus, a palladiumsputtered cylindrical specimen with 2r0 ¼ z0 ¼ 19 mm was prepared from Material D, and its hydrogen contents were measured after various exposure times to determine its saturated hydrogen content at RT as shown in Fig. 6.
(1)
(2)
where b1 ¼ 2.405. Furthermore, when 2r0 ¼ z0, which was the case for the cylindrical specimens used in the present study, the following equation is obtained: 2 CHR ðtR Þ 32 p þ 4b21 z 2 2 exp DtR : 2 CHS z0 p b1
(3)
Fig. 4(a) shows the relationship between CHR and tR for the cylindrical specimen with 2r0 ¼ z0 ¼ 15 mm sampled from Material E exposed to hydrogen at 100 MPa at a constant temperature of 85 C for 240 h. The first two plotted results were obtained during heating of the specimen from RT to 85 C. The experimental data were fitted to Eq. (1) by the least-squares method, with CHS and D as unknown parameters. The experimental results, apart from for the first
Fig. 4 e Hydrogen desorption behavior of cylindrical specimens sampled from Material E exposed to high-pressure hydrogen gas at (a) 85 C or (b) 50 C. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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Fig. 5 e Arrhenius plot of hydrogen diffusivity of various low-alloy steels.
two or three measured results, were consistent with the curve fitted to Eq. (1). The hydrogen diffusivity D can therefore be determined by this method, whereas the saturated hydrogen content CHS cannot be determined by it, because at t z 0, CHR obtained during heating of the specimen was inconsistent with the fitted curve. In other words, the value of D can be determined by using the following h i) equation: ( ∞ 2 X exp ð2n þ 1Þ p2 DtR z20 CHR ðtR Þ ¼A ð2n þ 1Þ2 n¼0 (4) ( ) ∞ X exp Db2m tR r20 $ b2m m¼1
Fig. 6 e Plot of the hydrogen content against the exposure time for cylindrical specimens of Material D with and without palladium sputtering after exposure to hydrogen gas at 100 MPa and RT.
where A is the fitted constant. The hydrogen diffusivity determined by fitting with Eq. (4) was 1.4 109 m2/s. When tR >> 0, the experimental results were fitted by the following equation: 2 p þ 4b21 Dt CHR ðtR Þ ¼ B exp R z20
(5)
where B is the fitted constant. When tR 90 min, Eq. (5) agreed with Eq. (4) and it was shown that D ¼ 1.3 109 m2/s by fitting with Eq. (5); this value was almost equal to that obtained by fitting with Eq. (4). The curves fitted by Eqs. (4) and (5) are shown in the figure. Fig. 4(b) shows the relationship between CHR and tR for the cylindrical specimen of Material E with 2r0 ¼ z0 ¼ 19 mm exposed to hydrogen at 100 MPa at a constant temperature of 50 C for 240 h. The first two plotted results were obtained during the process of heating the specimen from RT to 50 C. The curves fitted with Eqs. (4) and (5) are also shown in the figure. As in the case of Fig. 4(a), the experimental results, apart for the first two or three values, were consistent with the fitted curve. We found that D ¼ 4.6 1010 m2/s by fitting with Eq. (4). For tR 280 min, Eq. (5) agreed with Eq. (4) and the value D ¼ 4.7 1010 m2/s was obtained by fitting with Eq. (5); this value was nearly equal to that obtained by fitting with Eq. (4). The hydrogen diffusivities of the other materials at various temperatures were similarly determined by this method. Fig. 5 shows an Arrhenius plot of the hydrogen diffusivities for Materials A to E. The values of the hydrogen diffusivities are the mean values obtained for specimens with 2r0 ¼ z0 ¼ 15 mm and 19 mm. The lattice hydrogen diffusivity in bcc iron without any traps [see Eqs. (6) and (7), below], as reported by Kiuchi et al. [18], and the reference hydrogen diffusivity of low-alloy steels determined by charging without high-pressure hydrogen gas, for example, by electrochemical charging [21e24], are also shown in the figure.
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233 353 K ð40 80 o CÞ 5690 DL ¼ 7:28 108 exp RT
(6)
323 823K ð50 550 o CÞ
6700 7120 DL ¼ ð1 2:52Þ 107 exp RT
(7)
Here, DL is the lattice hydrogen diffusivity in bcc iron without any traps (m2/s), R is the gas constant (8.314 J mol1 K1), and T is the absolute temperature (K). The hydrogen diffusivities of the present low-alloy steels agreed closely with that reported by Fujii et al. [24] and were 5e15 times greater than the lattice hydrogen diffusivity calculated by using Eq. (6) or Eq. (7). The values of D for Materials A to E showed no large differences, although they tended to decrease for materials having a higher tensile strength and a coarser microstructure.
Entry and exit of high-pressure hydrogen gas in steels at room temperature As shown in Fig. 5, the diffusivities of hydrogen in Materials A to E at RT were approximately one order of magnitude less than that reported by Yamakawa et al. [23]. The value of the hydrogen diffusivity determined by Yamakawa et al. was obtained by an electrochemical permeation method in which the effects of surface reactions were carefully removed. The discrepancy between our results and those of Yamakawa implies that the hydrogen desorption at RT in our experiments did not occur by a diffusion-controlled process. We therefore determined the hydrogen diffusivity for palladiumsputtered cylindrical specimens of Material D with 2r0 ¼ z0 ¼ 19 mm. The hydrogen diffusivity of palladiumsputtered specimens of Material D at temperatures of RT, 50 C, and 85 C is shown in Fig. 5. The values of D at 50 C and 85 C were consistent with those for the nonsputtered specimens, whereas the value of D at RT for the palladiumsputtered specimen was larger than that for the nonsputtered one. However, the value for the sputtered specimen agreed with that reported by Yamakawa et al. Therefore, the hydrogen diffusivity at RT cannot be determined precisely by the present method using nonsputtered cylindrical specimens, because of the presence of a significant surface effect on the exit of hydrogen at RT. As well as the exit of hydrogen, entry of gas is also likely to be affected by surface reactions at RT. Therefore, palladiumsputtered and nonsputtered cylindrical specimens of Material D with 2r0 ¼ z0 ¼ 19 mm were exposed to high-pressure hydrogen at 100 MPa and RT for various times and their hydrogen contents were subsequently measured. Fig. 6 shows a plot of the hydrogen content against the exposure time for palladium-sputtered and nonsputtered cylindrical specimens exposed to hydrogen. The relationship between the hydrogen content and the exposure time, as calculated from the hydrogen diffusivity at RT (D ¼ 2.0 1010 m2/s) for a diffusion-controlled process, is also shown in the figure. For the palladium-sputtered cylindrical specimens, the exposure time required to obtain saturation by hydrogen was slightly
longer than the calculated value; however, a cylindrical specimen with 2r0 ¼ z0 ¼ 19 mm became saturated with hydrogen (0.48 mass ppm) on exposure to the high-pressure gas at RT for approximately 90 h. In contrast, even we exposed the nonsputtered specimen to high-pressure hydrogen at RT for approximately 300 h, we could not attain saturation. This confirmed that the entry of hydrogen into nonsputtered specimens, which contain native oxide layers on the surface, at RT does not occur by a diffusion-controlled process and that there is a significant surface effect on the entry of hydrogen. The hydrogen-entry behavior in highpressure hydrogen gas therefore differs from that in ordinary electrochemical charging [35].
Temperature dependence of the saturated hydrogen content Fig. 7 shows an Arrhenius plot of the saturated hydrogen content CHS for Materials A to E. The value of CHS at RT was also determined for a palladium-sputtered specimen of Material D. The results confirmed that palladium sputtering does not affect the value of CHS for specimens exposed to hydrogen at 100 MPa and 85 C. The lattice hydrogen content in bcc iron without any traps, as reported by Hirth [15], is also shown in the figure: qffiffiffi 3440 CLS ¼ 104:47 f exp T
(8)
where CLS is the lattice hydrogen content in bcc iron (mass ppm) and f is the fugacity (MPa). According to San Marchi et al. [19], the fugacity f can be expressed by using the hydrogen-gas pressure p as follows: bp RT
f ¼ p exp
(9)
where p is the pressure of the hydrogen gas (MPa) and b is a constant (1.584 105 m3 mol1) [19]. In the range RT to 85 C, the values of CHS for our low-alloy steels were not dependent on the temperature and they were 5e10 times larger than the values of CLS.
Interpretation of the temperature dependences of the hydrogen diffusivity and the saturated hydrogen content Fig. 8(a) and (b) show Arrhenius plots of the mean values of D and CHS, respectively, for Materials A to E. The experimental hydrogen diffusivity and saturated hydrogen content for Materials A to E showed no large differences; consequently, the temperature dependencies of the hydrogen diffusivity and the saturated hydrogen content were investigated by using the mean values of these properties. Figs. 5 and 7 suggest that the temperature dependence of D and CHS of the present low-alloy steels is affected by hydrogen trapping. The experimental values of D and CHS were therefore fitted in terms of Oriani's equilibrium theory [36] as follows: D¼
7:28 108 5690 ¼ exp RT EB EB 1 þ NNXL exp RT 1 þ NNXL exp RT DL
(10)
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Fig. 7 e Arrhenius plot of the saturated hydrogen contents of various low-alloy steels.
NX EB CLS ¼ 1þ exp NL RT
qffiffiffi NX EB 3440 f exp ¼ 104:47 1 þ exp T NL RT
CHS
(11)
where NL is the number of lattice sites per unit volume, NX is the number of trap sites per unit volume, and EB is the binding energy. Fitting by the least-squares method was performed because NX/NL and EB were unknown parameters. The experimental results for D and CHS were successfully fitted to Eqs. (10) and (11), giving the following parameters: NX/ NL ¼ 2.1 104 and EB ¼ 29.6 kJ/mol. The binding energy determined by the least-squares method has shown that the
hydrogen dissolved in low-alloy steels is mainly trapped at dislocations [37]. On the basis of Eq. (11), we compared the saturated hydrogen content of a CreMo steel in high-pressure hydrogen gas with that obtained by ordinary electrochemical charging. Matsuoka et al. [35] and Takeuchi et al. [38] reported, respectively, that the hydrogen contents of a CreMo steel saturated by hydrogen through electrochemical charging were 0.85 mass ppm for a sample charged by immersion in aqueous NH4SCN solution at 40 C for 12 h and 1.20 mass ppm for a sample cathodically charged in a mixed aqueous solution of NH4SCN and NaCl at a current density of 0.3 mA/cm2 for 24 h. From Eq. (11), the corresponding pressures of gaseous hydrogen at RT can be estimated to be 190 MPa and 250 MPa,
Fig. 8 e Interpretation of the temperature dependences of (a) the hydrogen diffusivity and (b) the saturated hydrogen content. Please cite this article in press as: Yamabe J, et al., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006
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respectively. It therefore follows that ordinary electrochemical charging is a more severe condition than exposure to hydrogen gas at a high pressure, for example, 100 MPa. In summary, the hydrogen diffusivity (more precisely, the effective hydrogen diffusivity) and the saturated hydrogen content of the present low-alloy steels at near-ambient temperatures can be determined by using Eqs. (10) and (11); however, it should be noted that significant surface effects on the entry and exit of hydrogen may occur in high-pressure hydrogen gas at RT in specimens containing native oxide layers on the surface.
Diffusion analysis by the finite-element method taking the surface effect into account As mentioned, the palladium-sputtered and the non-sputtered specimens both contain the native oxide layers. In the palladium-sputtered specimen, the native oxide layer remains at the subsurface of the specimen, whereas the surface is covered with the palladium layer. Nevertheless, the entry and exit of hydrogen at RT was significantly different between the palladium-sputtered and non-sputtered specimens, presumably caused by the surface reaction of hydrogen. Such surface reactions of hydrogen significantly affect the hydrogen diffusivity of low-alloy steels at RT as determined by the desorption method, whereas such surface reactions hardly affected the hydrogen diffusivity of g-stainless steel as determined by the same method. In the latter determination, we used a hydrogen-exposed cylindrical specimen with 2r0 ¼ 5 mm and z0 ¼ 0.5 mm. As a result, the value of D for g-stainless steel at RT was estimated to be approximately 1 1016 m2/s for both palladium-sputtered and nonsputtered specimens. In this case, it is presumed that the time required for hydrogen diffusion is much greater than that required for the surface reaction, because the D value of the g-stainless steel at RT (D ¼ 1 1016 m2/s) is six orders of magnitude greater than that for the low-alloy steels at RT (D ¼ 2 1010 m2/s). To interpret the difference in the surface effect between the low-alloy steels and the g-stainless steels, we performed a diffusion analysis by FEM, taking the surface reaction of hydrogen into account. Fig. 9 shows the specimen geometry and the boundary conditions for the FEM analysis. For the sake of simplicity, we performed a one-dimensional diffusion analysis for the cylindrical specimen shown Fig. 9 by using commercial FEM software (ANSYS, Ver. 14.5). In our analysis, we used a six-node triangular element (PLANE 35) with a mesh size of 0.02 mm. In diffusion analysis involving surface reaction of hydrogen, the local hydrogen content at the surface of the exit side csurf is assumed to be not zero [39], unlike the case of diffusion analysis without consideration of the surface reaction of hydrogen. Therefore, the value of csurf was determined so that the following equation was satisfied: kcsurf ¼ D
vcsurf vr
(12)
where k is the rate of the surface reaction, which is a constant. In the experiment results shown in Fig. 5, the values of D at RT estimated by using nonsputtered specimens ranged from 1.4 1011 m2/s to 7.1 1011 m2/s, whereas that estimated by
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using a palladium-sputtered specimen was 2.0 1010 m2/s. The FEM analysis was performed under various k values to obtain 1.4 1011 m2/s and 7.1 1011 m2/s; consequently, it was obtained that k ¼ 1.0 108 m/s and 4.0 108 m/s for D ¼ 1.4 1011 m2/s and 7.1 1011 m2/s, respectively. As a reference, Fig. 10 shows the relationship between CHR/CHS and tR when D ¼ 2 1010 m2/s and k ¼ 3 108 m/s, as calculated by FEM. The broken lines in the figure were those fitted by Eq. (1), as the values of CHS and D were unknown. Fig. 11 shows the hydrogen distribution when D ¼ 2 1010 m2/s and k ¼ 3 108 m/s as calculated by FEM with and without consideration of the surface reaction of hydrogen. Because of the surface reaction of hydrogen, the local hydrogen content at the surface of the exit side was not zero. These analytical results show that the surface reaction of hydrogen results in an underestimation of the hydrogen diffusivity of the present low-alloy steels; moreover, the calculated results were consistent with the experimental ones shown in Fig. 5. Different from low-alloy steel, the value of D for g-stainless steel at RT was estimated to be approximately 1 1016 m2/s for both palladium-sputtered and non-sputtered specimens, though the g-stainless steel possesses a native oxide layer of a few nanometers thickness, similar to that in low-alloy steel [31]. To interpret the experimental results, the FEM analysis taking account of the surface reaction of hydrogen was performed for the conditions D ¼ 1.0 1016 m2/s and k ¼ 1.0 108 m/s. Unlike the analytical results shown in Fig. 10, the value of D was estimated to be 1.0 1016 m2/s, even when the surface reaction of hydrogen was taken into consideration. This analytical result was consistent with the experimental one for the g-stainless steel; it therefore follows that surface reaction of hydrogen does not necessarily affect the hydrogen diffusivity of materials with low hydrogen diffusivity at RT determined by the desorption method, even if a native oxide layer is formed on the surface.
Conclusions The hydrogen diffusivity D and the saturated hydrogen content CHS of various low-alloy steels, potentially useful in the widespread commercialization of hydrogen-energy systems, were determined over a wide range of temperatures, including near-ambient temperatures, by exposing cylindrical specimens of four JISeSCM435 steels (CreMo steels) and one JISeSNCM439 steel (Nie CreMo steel) to high-pressure hydrogen gas at 100 MPa and a temperature of RT, 50 C, or 85 C for about 250 h. The size of the specimens was 2r0 ¼ z0 ¼ 7, 10, 15, or 19 mm, where r0 is the radius of the specimen and z0 is its thickness. After the hydrogen exposure, the hydrogen contents of the specimens were measured by gas chromatography (GCeMS) with a rising or constant temperature. The hydrogen diffusivity of the materials was determined by fitting the solution of a nonsteady diffusion equation to the experimental hydrogen-release profiles at constant temperature by means of the least-squares method. The effects of surface reactions, such as hydrogen dissociation and recombination, on the entry and exit of hydrogen were also investigated by using palladium-sputtered cylindrical specimens and a FEM diffusion analysis in which surface
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Fig. 9 e Specimen geometry and boundary conditions for the diffusion analysis by FEM taking the surface effect into account.
(2)
(3)
Fig. 10 e Calculated relationship between CHR/CHS and tR when D ¼ 2.0 £ 10¡10 m2/s and k ¼ 3.0 £ 10¡8 m/s.
(4)
effects were taken into consideration. Our conclusions can be summarized as follows. (1) In the profiles for the hydrogen desorption rates of both CreMo and NieCreMo steels under conditions of rising temperature, two peaks were observed at about 100 C (first peak) and 400 C (second peak). Regardless of the material, the hydrogen content associated with the first
(5)
(6)
Fig. 11 e Hydrogen distribution when D ¼ 2.0 £ 10¡10 m2/s and k ¼ 3.0 £ 10¡8 m/s, as calculated by the FEM.
peak (CH1) was much larger than that associated with the second peak (CH2); in other words, CH1 >> CH2. Under our exposure conditions at 50 C and 85 C, the hydrogen content associated with the first peak (CH1) decreased with decreasing specimen size for specimens with 2r0 ¼ z0 10 mm, whereas the values of CH1 were independent of the specimen size for specimens with 2r0 ¼ z0 15 mm. Because CH1 >> CH2, we regarded the hydrogen content for the first peak of specimens with 2r0 ¼ z0 15 mm as being equivalent to the saturated hydrogen content at 50 C or 85 C. The hydrogen diffusivity determined with the nonsputtered cylindrical specimens at temperatures other than RT showed no large differences among the various materials and it showed a similar tendency to reference data determined without high-pressure hydrogen gas, for example, with electrochemical charging. At RT, a significant reduction was observed in the hydrogen diffusivity determined with nonsputtered specimens in comparison to the corresponding reference date determined by the electrochemical permeation method. By contrast, the hydrogen diffusivity at RT determined with the palladium-sputtered specimens agreed closely with the reference data. Similarly, the hydrogen-entry behavior was markedly changed by palladium sputtering. It was therefore obvious that surface effects on the entry and exit of hydrogen were significant at RT. Consequently, the saturated hydrogen content at RT was determined by using palladiumsputtered cylindrical specimens. The saturated hydrogen content of the present lowalloy steels was 5e10 times larger than the lattice hydrogen content in bcc iron without any traps, whereas the hydrogen diffusivity was 5e15 time less than the lattice hydrogen diffusion in bcc iron without any traps. The temperature dependencies of the present low-alloy steels were quantitatively explained in terms of hydrogen trapping. The binding energy of the trap site was estimated to be 29.6 kJ/mol, implying that the hydrogen dissolved in the steels was mainly trapped at dislocations. No underestimation of the hydrogen diffusivity at RT was observed in g-stainless steel. This is because of the longer time required for hydrogen diffusion in comparison to that required for surface reactions in this steel. The effects of surface reactions on the hydrogen diffusivity determined for low-alloy and g-stainless
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steels can be interpreted by means of diffusion analysis by FEM taking account of surface effects.
Acknowledgment 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., Investigation of hydrogen transport behavior of various low-alloy steels with highpressure hydrogen gas, International Journal of Hydrogen Energy (2015), http://dx.doi.org/10.1016/j.ijhydene.2015.07.006