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Effects of tungsten addition on hydrogen absorption and permeation properties of Nb40Ti30Ni30 alloy
Journal of Alloys and Compounds 749 (2018) 634e639
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effects of tungsten addition on hydrogen absorption and permeation properties of Nb40Ti30Ni30 alloy Kazuhiro Ishikawa*, Kazuho Yonehara Faulty of Mechanical Engineering, Institute of Science and Engineering, Kanazawa University, Kanazawa, 920-1192, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 18 July 2017 Received in revised form 8 March 2018 Accepted 12 March 2018 Available online 13 March 2018
Tungsten was added to Nb40Ti30Ni30 alloys to investigate its effects on microstructure, hydrogen permeability and hydrogen absorption properties. Tungsten-added alloys consisted of a primary (Nb, Ti) phase and TiNiþ(Nb, Ti) eutectic structure, although tungsten concentrated at the center of the primary phase. Tungsten addition did not change their hydrogen permeability. However, the hydrogen capacity decreased with increasing tungsten content, and 20% reduction of hydrogen capacity was attained by adding 5% tungsten. After annealing, tungsten was dissolved homogeneously in the primary phase, and segregation disappeared. The hydrogen permeability and capacity of the annealed alloy were 50% higher and 10% lower, respectively, than those of tungsten-free alloy. Thus, tungsten can reduce hydrogen capacity without degrading hydrogen permeability, which is expected to improve the resistance to hydrogen embrittlement. © 2018 Elsevier B.V. All rights reserved.
Keywords: Hydrogen permeability Hydrogen absorption Hydrogen diffusivity Microstructure Tungsten
1. Introduction The membrane reforming technique is one of the most promising candidates for industrial hydrogen production [1,2]. In this technique, a hydrogen permeation alloy is installed in a reformer for separation of pure hydrogen from steam-reformed gas containing impurities. Pd and its alloys are used for commercial hydrogen separation and purification. They are, however, very expensive and scarce in nature, so alternative hydrogen permeation alloys are strongly desired. The 5A elements, such as V, Nb and Ta, have higher hydrogen permeability than pure Pd [3]. However, they absorb too much hydrogen in their lattice and show severe hydrogen embrittlement. To enable practical application of hydrogen permeation alloys, compatible hydrogen permeability and resistance to hydrogen embrittlement are required. Nishimura et al. first developed hydrogen-permeable V-Ni alloys having a higher hydrogen permeability than that of pure Pd [4]. Furthermore, they found that adding Al to V improved resistance to hydrogen embrittlement [5]. These alloys were designed on the basis of the concept that substitutional elements having low affinity for hydrogen lead to reduction of hydrogen capacity of the alloy. Nambu et al.
* Corresponding author. E-mail address: [email protected] (K. Ishikawa). https://doi.org/10.1016/j.jallcom.2018.03.146 0925-8388/© 2018 Elsevier B.V. All rights reserved.
investigated hydrogen absorption and mechanical properties in Nb-based alloys. They concluded that a ductile-brittle transition hydrogen concentration exists in Nb-based alloys. If the hydrogen concentration is below 0.2 H/M (hydrogen to metal ratio), hydrogen embrittlement does not occur [6]. Recently, we proposed a new concept of Nb-based alloys for hydrogen permeation. The Nb-TiNi alloys, for example, have a two-phase structure consisting of the (Nb, Ti) solid solution and TiNi intermetallic compound [7]. They are designed on the basis of the concept that hydrogen embrittlement of Nb can be suppressed by coexistence of TiNi having resistance to hydrogen embrittlement. This concept was applied to other alloy systems, and hydrogen embrittlement was successfully suppressed in alloys such as Nb-TiCo [8], Nb-ZrNi [9], Nb-HfNi [10] and Nb-HfCo [11] systems. The hydrogen flux J permeating through a sample is expressed using the Fick's first diffusion law as
J ¼ DðCu Cd Þ=L ¼ DDC=L
(1)
where Cu and Cd are the hydrogen concentrations at the surface of the sample in the upstream and downstream sides, respectively; D is the hydrogen diffusion coefficient; and L is the sample thickness. If the relationship between hydrogen concentration and hydrogen pressure obeys the Sieverts' law, then hydrogen concentration C can be expressed using hydrogen pressure P as
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639
C ¼ KP 0:5
(2)
635
(10 cc/min). Hydrogen pressure in the upstream side was increased to 0.5 MPa in 0.1 increments. After measurements, temperature was decreased to 523 K in 50 K increments.
where K is the hydrogen solubility coefficient. Combining Eqs. (1) and (2) gives
. . L ¼ FD P 0:5 L J ¼ DK Pu0:5 Pd0:5
3. Results and discussion
(3)
where F, the product of D and K, is defined as hydrogen permeability. From Eq. (1), we can see that hydrogen flux J depends on the hydrogen concentration differential between both sides of the sample, but it is not affected by the absolute value of hydrogen concentration C. If we can increase DC and decrease C simultaneously, both hydrogen permeability and resistance to hydrogen embrittlement would improve. Yukawa et al. achieved this by adding W, Re and Mo to Nb [12,13] or V [14], which reduced its hydrogen concentration and increased the hydrogen concentration differential. They showed large hydrogen flux without hydrogen embrittlement. In Nb-TiNi alloys, the (Nb, Ti) phase contributes hydrogen permeability, and equilibrium Nb content in this phase is about 85 mol % [15]. Therefore, it is expected that the addition of W to Nb-TiNi alloys also reduces hydrogen concentration without degrading hydrogen permeability. In this study, microstructure, hydrogen absorption and permeation properties of W-added Nb40Ti30Ni30 alloys were investigated both before and after annealing. Effects of W addition are discussed on the basis of the experimental data.
2. Experimental According to binary phase diagrams of W-Nb, W-Ti and W-Ni systems [16], it is considered that Nb should be substituted by W in Nb-TiNi alloys because no intermetallic compounds are formed and no phase separation occurs only in Nb-W system. Thus, alloys formed by adding W to Nb-TiNi alloys were selected as (Nb40-xWx) Ti30Ni30 (x ¼ 0e7). The alloy ingots were prepared using arc melting under an Ar atmosphere. Disk samples 12 mm in diameter and 0.7 mm thick were cut from the ingots using an electrical discharge machine. Samples were enclosed in a quartz tube filled with Ar gas for heat treatment. The capsules were heated to 1373 K for 168 h (1 week) in an electric furnace, and then quenched in water. The surfaces of the samples were polished with 0.6 mm Al2O3 particles. Microstructural observation was performed using a scanning electron microscope (SEM). Chemical compositions in each phase were measured by energy dispersive X-ray spectrometry (EDS). Structure of the samples was identified by an X-ray diffractometer (XRD). Both surfaces of each sample were coated by about 200 nm Pd using a DC sputtering machine in order to enhance dissociation of molecular hydrogen and to avoid oxidation. Equilibrium pressurecomposition-temperature (PCT) properties were measured by a Sieverts' type apparatus at 573 K. The sample contained in a stainless steel tube was set in a reactor chamber, then the chamber was evacuated to 5 103 Pa using a diffusion pump. The chamber was heated to 573 K in an electric furnace and kept for 20 min. High purity hydrogen (99.99999%) was introduced in the chamber, and kept 5 min for equilibrium hydrogen absorption. Hydrogen pressure in the chamber was measured by a pressure gauge. Each disk sample was fixed by gaskets, and both sides of the sample were evacuated using a diffusion pump, and heated to 673 K. Hydrogen gas at 0.2 MPa and 0.1 MPa was introduced to the upstream and downstream sides of the sample, respectively. Hydrogen flux permeating through the sample was measured using a mass flow meter. The accuracy of this mass flow meter is 1% against full scale
3.1. Effects of W addition on hydrogenation properties of the as-cast alloys Fig. 1(a) shows XRD patterns of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys. The diffraction peaks of the samples can be indexed on the basis of the bcc-(Nb, Ti) and B2-TiNi phases, respectively. The lattice parameters of the bcc and B2 phases were 0.329 nm and 0.301 nm, respectively. A portion of the XRD pattern, surrounded by the dotted rectangle in Fig. 1(a), was selected and enlarged in Fig. 1(b) to show that a small peak was observed on the higher-angle side overlapping with the bcc phase when W content was increased. This new diffraction peak can also be indexed as the bcc phase. That is, the original bcc phase separated into two bcc phases after W addition. Their lattice parameters were 0.329 nm and 0.325 nm, respectively. The new bcc phase having a small lattice parameter was considered to be W-based phases, because the lattice parameters of pure Nb and W are reported as 0.329 nm and 0.315 nm, respectively [17]. Fig. 2 shows SEM micrographs of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys. Before W addition (x ¼ 0), shown in Fig. 2(a), the primary (Nb, Ti) phase (white) was surrounded by a lamellar eutectic {TiNiþ(Nb, Ti)} structure. However, we can see that morphology of the primary phase changed from “granular” to “rugged” with increasing W content. Furthermore, bright and dark areas were clearly observed within the primary phase in the alloys containing W. EDS analysis showed that the W concentrations in center - peripheral positions of the primary phase were 14e6% for x ¼ 3, 17e7% for x ¼ 5 and 29e15% for x ¼ 7 alloys, which indicates that W was segregated in the center of the primary phase. This phase segregation was confirmed by XRD measurements. In contrast, W was not detected in the eutectic structure. Therefore, it
Fig. 1. (a) XRD patterns of the as-cast (Nb40-xWx)Ti30Ni30 alloys. (b) Enlarged 200 peak of the bcc phase.
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K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639
Fig. 2. SEM micrographs of the as-cast (Nb40-xWx)Ti30Ni30 alloys for (a) x ¼ 0, (b) x ¼ 3, (c) x ¼ 5 and (d) x ¼ 7.
is considered that W was preferentially consumed for formation of the primary phase in the early stage of solidification; in other words, all the W was captured in the primary phase. Afterword, the eutectic structure was formed from W-free molten alloy. The relationship between hydrogen pressure and composition of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys at 573 K is shown in Fig. 3(a) as a function of P0.5 versus C. It is clearly seen that their hydrogen concentrations cannot be fitted by straight lines passing through the origin. That is, these alloys did not show Sieverts' type hydrogen absorption, and Eq. (2) cannot be used to describe their hydrogen absorption behavior. Thus, hydrogen permeability F cannot also be defined as Eq. (3). However, if the hydrogen pressure region was limited to between 0.1 and 0.5 MPa, which agrees with the pressure region for hydrogen permeation measurements, then hydrogen concentration C was in proportion to the square root of hydrogen pressure P0.5. We obtain a pseudo-Sieverts’ equation, instead of Eq. (2) as
C ¼ KP 0:5 þ a
(4)
where a is the virtual hydrogen concentration at P ¼ 0, which is considered excess hydrogen that does not contribute to hydrogen permeation. The gradient of this straight line is defined as the hydrogen solubility coefficient K. Substituting Eq. (4) into Eq. (1), hydrogen permeability can be defined in Eq. (3). However, hydrogen absorption properties of these alloys deviated on the lower side of the approximately straight line when hydrogen pressure was below 0.1 MPa and above 0.5 MPa. At these pressures, hydrogen permeability cannot be defined by Eq. (3). Table 1 lists their hydrogenation properties such as hydrogen absorption, hydrogen permeability and hydrogen diffusion. As shown in Table 1, the values of K were almost constant regardless of W content. However, a decreased monotonically with increasing W content. Fig. 3(b) shows the hydrogen concentrations at 0.1 MPa and 0.5 MPa (C0.1 and C0.5) and their difference DC (¼ C0.5 - C0.1) as a function of W content. Both C0.1 and C0.5 decreased with increasing W content. For the x ¼ 5 alloy, reduction of the hydrogen concentration was 26% at 0.1 MPa and 20% at 0.5 MPa in a hydrogen
Fig. 3. (a) Hydrogen concentration of the as-cast (Nb40-xWx)Ti30Ni30 alloys as a function of the square root of hydrogen pressure. (b) Relationship between W content and hydrogen concentration at P ¼ 0.1 MPa (C0.1), P ¼ 0.5 MPa (C0.5) and DC (¼ C0.5 - C0.1).
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639 Table 1 Hydrogen absorption, permeation and diffusion in the as-cast (Nb40-xWx)Ti30Ni30 alloys. x
0
3
5
7
K (molH2 m3 Pa0.5) a (103 molH2 m3) C0.1 (103 molH2 m3) C0.5 (103 molH2 m3) DC (103 molH2 m3)
13.5 17.2 21.5 26.8 5.3
15.2 14.5 19.3 25.2 5.9
14.1 11.5 16.0 21.5 5.5
13.7 10.6 15.0 20.3 5.3
F573 (109 molH2 m1 s1 Pa0.5) F0 (106 molH2 m1 s1 Pa0.5)
Ea (kJ mol1)
6.8 5.7 32
6.7 3.5 30
6.3 1.6 27
5.9 1.3 26
D (109 m2 s1)
0.49
0.45
0.45
0.42
Experimental error: hydrogen capacity: ±0.2 (103 molH2 m3); hydrogen permeation: ±0.3 (109 molH2 m1 s1 Pa0.5).
atmosphere. Therefore, we conclude that W can reduce hydrogen capacity of the Nb40Ti30Ni30 alloy. On the other hand, DC was almost constant at about 5e6 molH2 m3 regardless of W substitution. The temperature dependence of hydrogen permeability (F) of these alloys is shown in Fig. 4 as an Arrhenius plot. It is seen that hydrogen permeability decreased with increasing W content. However, the maximum and minimum hydrogen permeabilities at 573 K were 6.8 109 for x ¼ 0 and 5.9 109 (mol H2 m1 s1 Pa0.5) for x ¼ 7, respectively. That is, adding W slightly reduced the hydrogen permeability of Nb40Ti30Ni30 alloys. Their hydrogen permeability increased linearly with reciprocal temperature in the Arrhenius plot, so that hydrogen permeation data were fitted by the Arrhenius equation:
F ¼ F0 expðEa =RTÞ
(5)
637
independent of W content, the slight decrease of hydrogen permeability was caused by the reduced hydrogen diffusivity. In conclusion, we demonstrated that hydrogen capacity of 20% or more can be reduced without large degradation of hydrogen permeability by W substitution. 3.2. Effects of heat treatment on hydrogen permeation of (Nb35W5) Ti30Ni30 alloy It has been reported that the eutectic structure disappears and is replaced by a granulare (N, Ti) phase surrounded by TiNi matrix after rolling and annealing [18]. Therefore, it was also interesting and important to investigate the effects of annealing on hydrogen absorption and permeation properties of the W-added alloys. Fig. 5(a) shows an SEM micrograph of the (Nb35W5)Ti30Ni30 alloy annealed at 1373 K for 168 h (1 week). Small (Nb, Ti) particles were observed in the position where the eutectic structure originally existed. The contrast within the primary phase, which is observed in the as-cast alloy as shown in Fig. 2(b), disappeared. Distribution of W in the primary phase of heat-treated (Nb35W5)Ti30Ni30 alloy is shown in Fig. 5(b). It is seen that W distributed homogeneously only in the primary phase. Furthermore, peak separation (Fig. 1(b)) was not observed in the XRD pattern of this alloy (not shown here). These experimental data suggest that the W segregation found in the as-cast alloy was eliminated after annealing. Fig. 6 shows the relationship between hydrogen concentration and pressure in the form of DP0.5 versus C. As for the as-cast alloys, the hydrogen concentration in the annealed alloys increased linearly with increasing the square root of hydrogen pressure in the range of 0.1e0.5 MPa, and they deviated on the lower side of the straight line below and above these pressures. Thus, hydrogen solubility coefficients K in these two alloys can also be defined by
where F0 is the frequency factor, Ea is the activation energy for hydrogen permeation. R and T are the gas constant and temperature, respectively. Both F0 and Ea were reduced monotonically with increasing W content, which means that W substitution reduced the thermal barrier for hydrogen diffusion and low-frequency jumping of hydrogen atoms to next the interstitial site. Hydrogen diffusion coefficients for hydrogen permeation were estimated using Eq. (3). Hydrogen diffusivity of these alloys decreased slightly with increasing W content. Because K was
Fig. 4. Arrhenius plot of temperature dependence of hydrogen permeability of the ascast (Nb40-xWx)Ti30Ni30 alloys.
Fig. 5. (a) SEM micrograph of the (Nb35W5)Ti30Ni30 alloy annealed at 1373 K for 168 h. (b) Distribution of W in the primary phase obtained by EDS analysis.
638
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639
Fig. 6. Hydrogen concentration of the Nb40Ti30Ni30 and (Nb35W5)Ti30Ni30 alloys annealed at 1373 K for 168 h as a function of the square root of hydrogen pressure.
Fig. 7. Arrhenius plot of temperature dependence of hydrogen permeability of the Nb340Ti30Ni30 and (Nb35W5)Ti30Ni30 alloys annealed at 1373 K for 168 h.
the gradient of this straight line. Hydrogen permeation and absorption properties of these two alloys are summarized in Table 2. They showed almost the same value. However, virtual hydrogen concentrations at P ¼ 0 MPa (a), 0.1 MPa (C0.1) and 0.5 MPa (C0.5) in W-added alloy were about 10% lower than those in W-free alloy after annealing. Therefore, W has an ability to reduce the hydrogen concentration in Nb-TiNi alloy after annealing. Hydrogen capacities of the (Nb35W5)Ti30Ni30 alloys before and after annealing were compared. In this alloy, annealing increased hydrogen concentration a, C0.1 and C0.5, and decreased DC. Furthermore, annealing reduced the value of K. Temperature dependence of hydrogen permeability of annealed (Nb35W5)Ti30Ni30 alloy is shown in Fig. 7. The hydrogen permeability of annealed W-free Nb40Ti30Ni30 alloy is also plotted in this figure for reference. The (Nb35W5)Ti30Ni30 alloy showed higher hydrogen permeability than that of W-free alloy after annealing. Comparison of Figs. 4 and 7 shows that hydrogen permeability of the W-free alloy did not change after annealing, although activation energy for hydrogen permeation was slightly smaller in the annealed alloy. We can see that annealing increased the hydrogen permeability of the W-added alloy. However, the difference between the frequency factor and activation energy for hydrogen permeation was not large enough to explain the drastic improvement of hydrogen permeability due to annealing. Hydrogen diffusion coefficient D was estimated using experimentally determined F and K. Hydrogen diffusivity in the W-added
alloy after annealing was much higher than that in the annealed Wfree alloy and as-cast alloy. Therefore, improvement of hydrogen permeability in the W-added alloy after annealing was caused by high hydrogen diffusivity. It is considered that hydrogen diffusion is strongly affected by microstructure. Hydrogen atoms must pass through the eutectic structure in the as-cast alloy, and jump across many phase boundaries of the bcc and B2 phases, which would act as diffusion barriers for hydrogen atoms [19]. It is interesting that hydrogen permeability is decreased in the as-cast alloy but increased in the annealed alloy by addition of W, respectively. It was reported that hydrogen permeability in Nb-TiNi alloys is strongly affected by the primary phase [19]. W concentrated in the center of the primary phase and its composition was 17% in the ascast (Nb35W5)Ti30Ni30 alloy. It is considered that this high concentration of W suppresses hydrogen permeability in the as-cast alloy. However, segregation of W was disappeared and homogeneous primary phase was formed after annealing which improve hydrogen permeability. It is expected that hydrogen permeability of these alloys with/without W addition and annealing is affected by presence of the eutectic structure and chemical composition of W in the primary phase. However, their relationship and the exact mechanism are unclear and currently under investigation. In this study, we demonstrated that the addition of W reduces the hydrogen concentration of Nb40Ti30Ni30 alloys without large degradation of hydrogen permeability in the as-cast state, and hydrogen permeability increases due to improved hydrogen diffusivity. If the relationship between hydrogen permeability and concentration is controlled appropriately, substitution of elements having low affinity for hydrogen is expected to produce industrial hydrogen separation and purification alloys compatible with hydrogen permeability and resistance to hydrogen embrittlement.
Table 2 Hydrogen absorption, permeation and diffusion in the Nb40Ti30Ni30 and (Nb35W5) Ti30Ni30 alloys annealed at 1373 K for 168 h. x
0
5
K (molH2 m3 Pa0.5) a (103 molH2 m3) C0.1 (103 molH2 m3) C0.5 (103 molH2 m3) DC (103 molH2 m3)
10.5 15.8 19.1 23.2 4.1
10.3 13.8 17.1 21.0 3.9
F573 (109 molH2 m1 s1 Pa0.5) F0 (106 molH2 m1 s1 Pa0.5)
Ea (kJ mol1)
7.5 2.3 27
11.7 1.7 24
D (109 m2 s1)
0.71
1.1
Experimental error: hydrogen capacity: ±0.2 (103 molH2 m3); hydrogen permeation: ±0.3 (109 molH2 m1 s1 Pa0.5).
4. Summary Microstructure, hydrogen content and hydrogen permeability in Nb40Ti30Ni30 and (Nb40-xWx)Ti30Ni30 alloys were investigated. Hydrogen permeability F of these alloys slightly decreased with increasing W content. Hydrogen concentrations in these alloys were linearly proportional to the square root of hydrogen pressure in the pressure range of 0.1e0.5 MPa. Hydrogen solubility coefficient K was almost independent of W content. Hydrogen diffusivity D decreased slightly with W content. However, 20% or more
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639
reduction of hydrogen concentration was attained by 5% addition of W without a decrease of the hydrogen concentration differential at 0.1 MPa and 0.5 MPa. After heat treatment, hydrogen permeability of the alloy with 5% W increased due to improved hydrogen diffusivity. Thus, addition of W is a useful technique for producing industrial hydrogen separation and purification alloys compatible with hydrogen permeability and resistance to embrittlement. Acknowledgement This research was supported in part by a Grant-in-Aid for Scientific Research (C) (26420719) and (B) (17H0340) from the Japan Society for the Promotion of Science (JSPS). References [1] E. Kikuchi, Membrane reactor application to hydrogen production, Catal. Today 56 (2000) 97e101. [2] S. Tosti, Supported and laminated Pd-based metallic membranes, Int. J. Hydrogen Energy 28 (2003) 1445e1454. [3] R.E. Buxbaum, A.B. Kinney, Hydrogen transport through tubular membranes of palladium-coated tantalum and niobium, Ind. Eng. Chem. Res. 35 (1996) 530e537. [4] C. Nishimura, M. Komaki, M. Amano, Hydrogen permeation characteristics of vanadium-nickel alloys, Mater. Trans. JIM 32 (1991) 501e507. [5] Y. Zhang, T. Ozaki, M. Komaki, C. Nishimura, Hydrogen permeation characteristics of vanadium-aluminium alloys, Scripta Mater. 47 (2002) 601e606. [6] T. Nambu, K. Shimizu, Y. Matsumoto, R. Rong, N. Watanabe, H. Yukawa, M. Morinaga, I. Yasuda, Enhanced hydrogen embrittlement of Pd-coated niobium metal membrane detected by in situ small punch test under hydrogen permeation, J. Alloys Compd. 446e447 (2007) 588e592.
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[7] K. Hashi, K. Ishikawa, T. Matsuda, K. Aoki, Hydrogen permeation characteristics of multi-phase Ni-Ti-Nb alloys, J. Alloys Compd. 368 (2004) 215e220. [8] K. Hashi, K. Ishikawa, T. Matsuda, K. Aoki, Microstructure and hydrogen permeability in Nb-Ti-Co multiphase alloys, J. Alloys Compd. 425 (2006) 284e290. [9] K. Ishikawa, T. Takano, T. Matsuda, K. Aoki, High hydrogen permeability in the Nb-Zr-Ni eutectic alloy containing the primary body-centered-cubic (Nb, Zr) phase, Appl. Phys. Lett. 87 (2005) 81906. [10] F. Shi, X. Song, Effect of niobium on the microstructure, hydrogen embrittlement, and hydrogen permeability of NbxHf(1-x)/2Ni(1-x)/2 ternary alloys, Int. J. Hydrogen Energy 35 (2010) 10620e10623. [11] E. Yan, X. Li, D. Liu, M. Rettenmayr, Y. Su, J. Guo, Nb-HfCo alloys with pronounced high hydrogen permeability: a new family of metallic hydrogen permeation membranes, Int. J. Hydrogen Energy 39 (2014) 8385e8389. [12] N. Watanabe, H. Yukawa, T. Nambu, Y. Matsumoto, G.X. Zhang, M. Morinaga, Alloying effects of Ru and W on the resistance to hydrogen embrittlement and hydrogen permeability of niobium, J. Alloys Compd. 477 (2009) 851e854. [13] Y. Awakura, T. Nambu, Y. Matsumoto, H. Yukawa, Hydrogen solubility and permeability of Nb-W-Mo alloy membrane, J. Alloys Compd. 509 (2011) S877eS880. [14] H. Yukawa, C. Tsukada, T. Nambu, Y. Matsumoto, Hydrogen solubility and permeability of V-W-Mo alloy membrane for hydrogen separation and purification, J. Alloys Compd. 580 (2013) S386eS390. [15] W. Wang, K. Ishikawa, K. Aoki, Microstructural change-induced lowering of hydrogen permeability in eutectic Nb-TiNi alloy, J. Membr. Sci. 351 (2010) 65e68. [16] T.B. Massalski, H. Okanoto, P.R. Subramanian, L. Kacprzak (Eds.), Binary Alloy Phase Diagrams, second ed., ASM International, 1990. [17] P. Villars (Ed.), Pearson's Handbook Desk Edition, ASM International, 1997. [18] K. Ishikawa, S. Tokui, K. Aoki, Microstructure and hydrogen permeation of cold rolled and annealed Nb40Ti30Ni30 alloy, Intermetallics 17 (2009) 109e114. [19] K. Ishikawa, S. Tokui, K. Aoki, Hydrogen permeation in anisotropic Nb-TiNi two-phase alloys formed by forging and rolling, Int. J. Hydrogen Energy 42 (2017) 11411e11421.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effects of tungsten addition on hydrogen absorption and permeation properties of Nb40Ti30Ni30 alloy Kazuhiro Ishikawa*, Kazuho Yonehara Faulty of Mechanical Engineering, Institute of Science and Engineering, Kanazawa University, Kanazawa, 920-1192, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 18 July 2017 Received in revised form 8 March 2018 Accepted 12 March 2018 Available online 13 March 2018
Tungsten was added to Nb40Ti30Ni30 alloys to investigate its effects on microstructure, hydrogen permeability and hydrogen absorption properties. Tungsten-added alloys consisted of a primary (Nb, Ti) phase and TiNiþ(Nb, Ti) eutectic structure, although tungsten concentrated at the center of the primary phase. Tungsten addition did not change their hydrogen permeability. However, the hydrogen capacity decreased with increasing tungsten content, and 20% reduction of hydrogen capacity was attained by adding 5% tungsten. After annealing, tungsten was dissolved homogeneously in the primary phase, and segregation disappeared. The hydrogen permeability and capacity of the annealed alloy were 50% higher and 10% lower, respectively, than those of tungsten-free alloy. Thus, tungsten can reduce hydrogen capacity without degrading hydrogen permeability, which is expected to improve the resistance to hydrogen embrittlement. © 2018 Elsevier B.V. All rights reserved.
Keywords: Hydrogen permeability Hydrogen absorption Hydrogen diffusivity Microstructure Tungsten
1. Introduction The membrane reforming technique is one of the most promising candidates for industrial hydrogen production [1,2]. In this technique, a hydrogen permeation alloy is installed in a reformer for separation of pure hydrogen from steam-reformed gas containing impurities. Pd and its alloys are used for commercial hydrogen separation and purification. They are, however, very expensive and scarce in nature, so alternative hydrogen permeation alloys are strongly desired. The 5A elements, such as V, Nb and Ta, have higher hydrogen permeability than pure Pd [3]. However, they absorb too much hydrogen in their lattice and show severe hydrogen embrittlement. To enable practical application of hydrogen permeation alloys, compatible hydrogen permeability and resistance to hydrogen embrittlement are required. Nishimura et al. first developed hydrogen-permeable V-Ni alloys having a higher hydrogen permeability than that of pure Pd [4]. Furthermore, they found that adding Al to V improved resistance to hydrogen embrittlement [5]. These alloys were designed on the basis of the concept that substitutional elements having low affinity for hydrogen lead to reduction of hydrogen capacity of the alloy. Nambu et al.
* Corresponding author. E-mail address: [email protected] (K. Ishikawa). https://doi.org/10.1016/j.jallcom.2018.03.146 0925-8388/© 2018 Elsevier B.V. All rights reserved.
investigated hydrogen absorption and mechanical properties in Nb-based alloys. They concluded that a ductile-brittle transition hydrogen concentration exists in Nb-based alloys. If the hydrogen concentration is below 0.2 H/M (hydrogen to metal ratio), hydrogen embrittlement does not occur [6]. Recently, we proposed a new concept of Nb-based alloys for hydrogen permeation. The Nb-TiNi alloys, for example, have a two-phase structure consisting of the (Nb, Ti) solid solution and TiNi intermetallic compound [7]. They are designed on the basis of the concept that hydrogen embrittlement of Nb can be suppressed by coexistence of TiNi having resistance to hydrogen embrittlement. This concept was applied to other alloy systems, and hydrogen embrittlement was successfully suppressed in alloys such as Nb-TiCo [8], Nb-ZrNi [9], Nb-HfNi [10] and Nb-HfCo [11] systems. The hydrogen flux J permeating through a sample is expressed using the Fick's first diffusion law as
J ¼ DðCu Cd Þ=L ¼ DDC=L
(1)
where Cu and Cd are the hydrogen concentrations at the surface of the sample in the upstream and downstream sides, respectively; D is the hydrogen diffusion coefficient; and L is the sample thickness. If the relationship between hydrogen concentration and hydrogen pressure obeys the Sieverts' law, then hydrogen concentration C can be expressed using hydrogen pressure P as
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639
C ¼ KP 0:5
(2)
635
(10 cc/min). Hydrogen pressure in the upstream side was increased to 0.5 MPa in 0.1 increments. After measurements, temperature was decreased to 523 K in 50 K increments.
where K is the hydrogen solubility coefficient. Combining Eqs. (1) and (2) gives
. . L ¼ FD P 0:5 L J ¼ DK Pu0:5 Pd0:5
3. Results and discussion
(3)
where F, the product of D and K, is defined as hydrogen permeability. From Eq. (1), we can see that hydrogen flux J depends on the hydrogen concentration differential between both sides of the sample, but it is not affected by the absolute value of hydrogen concentration C. If we can increase DC and decrease C simultaneously, both hydrogen permeability and resistance to hydrogen embrittlement would improve. Yukawa et al. achieved this by adding W, Re and Mo to Nb [12,13] or V [14], which reduced its hydrogen concentration and increased the hydrogen concentration differential. They showed large hydrogen flux without hydrogen embrittlement. In Nb-TiNi alloys, the (Nb, Ti) phase contributes hydrogen permeability, and equilibrium Nb content in this phase is about 85 mol % [15]. Therefore, it is expected that the addition of W to Nb-TiNi alloys also reduces hydrogen concentration without degrading hydrogen permeability. In this study, microstructure, hydrogen absorption and permeation properties of W-added Nb40Ti30Ni30 alloys were investigated both before and after annealing. Effects of W addition are discussed on the basis of the experimental data.
2. Experimental According to binary phase diagrams of W-Nb, W-Ti and W-Ni systems [16], it is considered that Nb should be substituted by W in Nb-TiNi alloys because no intermetallic compounds are formed and no phase separation occurs only in Nb-W system. Thus, alloys formed by adding W to Nb-TiNi alloys were selected as (Nb40-xWx) Ti30Ni30 (x ¼ 0e7). The alloy ingots were prepared using arc melting under an Ar atmosphere. Disk samples 12 mm in diameter and 0.7 mm thick were cut from the ingots using an electrical discharge machine. Samples were enclosed in a quartz tube filled with Ar gas for heat treatment. The capsules were heated to 1373 K for 168 h (1 week) in an electric furnace, and then quenched in water. The surfaces of the samples were polished with 0.6 mm Al2O3 particles. Microstructural observation was performed using a scanning electron microscope (SEM). Chemical compositions in each phase were measured by energy dispersive X-ray spectrometry (EDS). Structure of the samples was identified by an X-ray diffractometer (XRD). Both surfaces of each sample were coated by about 200 nm Pd using a DC sputtering machine in order to enhance dissociation of molecular hydrogen and to avoid oxidation. Equilibrium pressurecomposition-temperature (PCT) properties were measured by a Sieverts' type apparatus at 573 K. The sample contained in a stainless steel tube was set in a reactor chamber, then the chamber was evacuated to 5 103 Pa using a diffusion pump. The chamber was heated to 573 K in an electric furnace and kept for 20 min. High purity hydrogen (99.99999%) was introduced in the chamber, and kept 5 min for equilibrium hydrogen absorption. Hydrogen pressure in the chamber was measured by a pressure gauge. Each disk sample was fixed by gaskets, and both sides of the sample were evacuated using a diffusion pump, and heated to 673 K. Hydrogen gas at 0.2 MPa and 0.1 MPa was introduced to the upstream and downstream sides of the sample, respectively. Hydrogen flux permeating through the sample was measured using a mass flow meter. The accuracy of this mass flow meter is 1% against full scale
3.1. Effects of W addition on hydrogenation properties of the as-cast alloys Fig. 1(a) shows XRD patterns of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys. The diffraction peaks of the samples can be indexed on the basis of the bcc-(Nb, Ti) and B2-TiNi phases, respectively. The lattice parameters of the bcc and B2 phases were 0.329 nm and 0.301 nm, respectively. A portion of the XRD pattern, surrounded by the dotted rectangle in Fig. 1(a), was selected and enlarged in Fig. 1(b) to show that a small peak was observed on the higher-angle side overlapping with the bcc phase when W content was increased. This new diffraction peak can also be indexed as the bcc phase. That is, the original bcc phase separated into two bcc phases after W addition. Their lattice parameters were 0.329 nm and 0.325 nm, respectively. The new bcc phase having a small lattice parameter was considered to be W-based phases, because the lattice parameters of pure Nb and W are reported as 0.329 nm and 0.315 nm, respectively [17]. Fig. 2 shows SEM micrographs of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys. Before W addition (x ¼ 0), shown in Fig. 2(a), the primary (Nb, Ti) phase (white) was surrounded by a lamellar eutectic {TiNiþ(Nb, Ti)} structure. However, we can see that morphology of the primary phase changed from “granular” to “rugged” with increasing W content. Furthermore, bright and dark areas were clearly observed within the primary phase in the alloys containing W. EDS analysis showed that the W concentrations in center - peripheral positions of the primary phase were 14e6% for x ¼ 3, 17e7% for x ¼ 5 and 29e15% for x ¼ 7 alloys, which indicates that W was segregated in the center of the primary phase. This phase segregation was confirmed by XRD measurements. In contrast, W was not detected in the eutectic structure. Therefore, it
Fig. 1. (a) XRD patterns of the as-cast (Nb40-xWx)Ti30Ni30 alloys. (b) Enlarged 200 peak of the bcc phase.
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Fig. 2. SEM micrographs of the as-cast (Nb40-xWx)Ti30Ni30 alloys for (a) x ¼ 0, (b) x ¼ 3, (c) x ¼ 5 and (d) x ¼ 7.
is considered that W was preferentially consumed for formation of the primary phase in the early stage of solidification; in other words, all the W was captured in the primary phase. Afterword, the eutectic structure was formed from W-free molten alloy. The relationship between hydrogen pressure and composition of the as-cast (Nb40-xWx)Ti30Ni30 (x ¼ 0e7) alloys at 573 K is shown in Fig. 3(a) as a function of P0.5 versus C. It is clearly seen that their hydrogen concentrations cannot be fitted by straight lines passing through the origin. That is, these alloys did not show Sieverts' type hydrogen absorption, and Eq. (2) cannot be used to describe their hydrogen absorption behavior. Thus, hydrogen permeability F cannot also be defined as Eq. (3). However, if the hydrogen pressure region was limited to between 0.1 and 0.5 MPa, which agrees with the pressure region for hydrogen permeation measurements, then hydrogen concentration C was in proportion to the square root of hydrogen pressure P0.5. We obtain a pseudo-Sieverts’ equation, instead of Eq. (2) as
C ¼ KP 0:5 þ a
(4)
where a is the virtual hydrogen concentration at P ¼ 0, which is considered excess hydrogen that does not contribute to hydrogen permeation. The gradient of this straight line is defined as the hydrogen solubility coefficient K. Substituting Eq. (4) into Eq. (1), hydrogen permeability can be defined in Eq. (3). However, hydrogen absorption properties of these alloys deviated on the lower side of the approximately straight line when hydrogen pressure was below 0.1 MPa and above 0.5 MPa. At these pressures, hydrogen permeability cannot be defined by Eq. (3). Table 1 lists their hydrogenation properties such as hydrogen absorption, hydrogen permeability and hydrogen diffusion. As shown in Table 1, the values of K were almost constant regardless of W content. However, a decreased monotonically with increasing W content. Fig. 3(b) shows the hydrogen concentrations at 0.1 MPa and 0.5 MPa (C0.1 and C0.5) and their difference DC (¼ C0.5 - C0.1) as a function of W content. Both C0.1 and C0.5 decreased with increasing W content. For the x ¼ 5 alloy, reduction of the hydrogen concentration was 26% at 0.1 MPa and 20% at 0.5 MPa in a hydrogen
Fig. 3. (a) Hydrogen concentration of the as-cast (Nb40-xWx)Ti30Ni30 alloys as a function of the square root of hydrogen pressure. (b) Relationship between W content and hydrogen concentration at P ¼ 0.1 MPa (C0.1), P ¼ 0.5 MPa (C0.5) and DC (¼ C0.5 - C0.1).
K. Ishikawa, K. Yonehara / Journal of Alloys and Compounds 749 (2018) 634e639 Table 1 Hydrogen absorption, permeation and diffusion in the as-cast (Nb40-xWx)Ti30Ni30 alloys. x
0
3
5
7
K (molH2 m3 Pa0.5) a (103 molH2 m3) C0.1 (103 molH2 m3) C0.5 (103 molH2 m3) DC (103 molH2 m3)
13.5 17.2 21.5 26.8 5.3
15.2 14.5 19.3 25.2 5.9
14.1 11.5 16.0 21.5 5.5
13.7 10.6 15.0 20.3 5.3
F573 (109 molH2 m1 s1 Pa0.5) F0 (106 molH2 m1 s1 Pa0.5)
Ea (kJ mol1)
6.8 5.7 32
6.7 3.5 30
6.3 1.6 27
5.9 1.3 26
D (109 m2 s1)
0.49
0.45
0.45
0.42
Experimental error: hydrogen capacity: ±0.2 (103 molH2 m3); hydrogen permeation: ±0.3 (109 molH2 m1 s1 Pa0.5).
atmosphere. Therefore, we conclude that W can reduce hydrogen capacity of the Nb40Ti30Ni30 alloy. On the other hand, DC was almost constant at about 5e6 molH2 m3 regardless of W substitution. The temperature dependence of hydrogen permeability (F) of these alloys is shown in Fig. 4 as an Arrhenius plot. It is seen that hydrogen permeability decreased with increasing W content. However, the maximum and minimum hydrogen permeabilities at 573 K were 6.8 109 for x ¼ 0 and 5.9 109 (mol H2 m1 s1 Pa0.5) for x ¼ 7, respectively. That is, adding W slightly reduced the hydrogen permeability of Nb40Ti30Ni30 alloys. Their hydrogen permeability increased linearly with reciprocal temperature in the Arrhenius plot, so that hydrogen permeation data were fitted by the Arrhenius equation:
F ¼ F0 expðEa =RTÞ
(5)
637
independent of W content, the slight decrease of hydrogen permeability was caused by the reduced hydrogen diffusivity. In conclusion, we demonstrated that hydrogen capacity of 20% or more can be reduced without large degradation of hydrogen permeability by W substitution. 3.2. Effects of heat treatment on hydrogen permeation of (Nb35W5) Ti30Ni30 alloy It has been reported that the eutectic structure disappears and is replaced by a granulare (N, Ti) phase surrounded by TiNi matrix after rolling and annealing [18]. Therefore, it was also interesting and important to investigate the effects of annealing on hydrogen absorption and permeation properties of the W-added alloys. Fig. 5(a) shows an SEM micrograph of the (Nb35W5)Ti30Ni30 alloy annealed at 1373 K for 168 h (1 week). Small (Nb, Ti) particles were observed in the position where the eutectic structure originally existed. The contrast within the primary phase, which is observed in the as-cast alloy as shown in Fig. 2(b), disappeared. Distribution of W in the primary phase of heat-treated (Nb35W5)Ti30Ni30 alloy is shown in Fig. 5(b). It is seen that W distributed homogeneously only in the primary phase. Furthermore, peak separation (Fig. 1(b)) was not observed in the XRD pattern of this alloy (not shown here). These experimental data suggest that the W segregation found in the as-cast alloy was eliminated after annealing. Fig. 6 shows the relationship between hydrogen concentration and pressure in the form of DP0.5 versus C. As for the as-cast alloys, the hydrogen concentration in the annealed alloys increased linearly with increasing the square root of hydrogen pressure in the range of 0.1e0.5 MPa, and they deviated on the lower side of the straight line below and above these pressures. Thus, hydrogen solubility coefficients K in these two alloys can also be defined by
where F0 is the frequency factor, Ea is the activation energy for hydrogen permeation. R and T are the gas constant and temperature, respectively. Both F0 and Ea were reduced monotonically with increasing W content, which means that W substitution reduced the thermal barrier for hydrogen diffusion and low-frequency jumping of hydrogen atoms to next the interstitial site. Hydrogen diffusion coefficients for hydrogen permeation were estimated using Eq. (3). Hydrogen diffusivity of these alloys decreased slightly with increasing W content. Because K was
Fig. 4. Arrhenius plot of temperature dependence of hydrogen permeability of the ascast (Nb40-xWx)Ti30Ni30 alloys.
Fig. 5. (a) SEM micrograph of the (Nb35W5)Ti30Ni30 alloy annealed at 1373 K for 168 h. (b) Distribution of W in the primary phase obtained by EDS analysis.
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Fig. 6. Hydrogen concentration of the Nb40Ti30Ni30 and (Nb35W5)Ti30Ni30 alloys annealed at 1373 K for 168 h as a function of the square root of hydrogen pressure.
Fig. 7. Arrhenius plot of temperature dependence of hydrogen permeability of the Nb340Ti30Ni30 and (Nb35W5)Ti30Ni30 alloys annealed at 1373 K for 168 h.
the gradient of this straight line. Hydrogen permeation and absorption properties of these two alloys are summarized in Table 2. They showed almost the same value. However, virtual hydrogen concentrations at P ¼ 0 MPa (a), 0.1 MPa (C0.1) and 0.5 MPa (C0.5) in W-added alloy were about 10% lower than those in W-free alloy after annealing. Therefore, W has an ability to reduce the hydrogen concentration in Nb-TiNi alloy after annealing. Hydrogen capacities of the (Nb35W5)Ti30Ni30 alloys before and after annealing were compared. In this alloy, annealing increased hydrogen concentration a, C0.1 and C0.5, and decreased DC. Furthermore, annealing reduced the value of K. Temperature dependence of hydrogen permeability of annealed (Nb35W5)Ti30Ni30 alloy is shown in Fig. 7. The hydrogen permeability of annealed W-free Nb40Ti30Ni30 alloy is also plotted in this figure for reference. The (Nb35W5)Ti30Ni30 alloy showed higher hydrogen permeability than that of W-free alloy after annealing. Comparison of Figs. 4 and 7 shows that hydrogen permeability of the W-free alloy did not change after annealing, although activation energy for hydrogen permeation was slightly smaller in the annealed alloy. We can see that annealing increased the hydrogen permeability of the W-added alloy. However, the difference between the frequency factor and activation energy for hydrogen permeation was not large enough to explain the drastic improvement of hydrogen permeability due to annealing. Hydrogen diffusion coefficient D was estimated using experimentally determined F and K. Hydrogen diffusivity in the W-added
alloy after annealing was much higher than that in the annealed Wfree alloy and as-cast alloy. Therefore, improvement of hydrogen permeability in the W-added alloy after annealing was caused by high hydrogen diffusivity. It is considered that hydrogen diffusion is strongly affected by microstructure. Hydrogen atoms must pass through the eutectic structure in the as-cast alloy, and jump across many phase boundaries of the bcc and B2 phases, which would act as diffusion barriers for hydrogen atoms [19]. It is interesting that hydrogen permeability is decreased in the as-cast alloy but increased in the annealed alloy by addition of W, respectively. It was reported that hydrogen permeability in Nb-TiNi alloys is strongly affected by the primary phase [19]. W concentrated in the center of the primary phase and its composition was 17% in the ascast (Nb35W5)Ti30Ni30 alloy. It is considered that this high concentration of W suppresses hydrogen permeability in the as-cast alloy. However, segregation of W was disappeared and homogeneous primary phase was formed after annealing which improve hydrogen permeability. It is expected that hydrogen permeability of these alloys with/without W addition and annealing is affected by presence of the eutectic structure and chemical composition of W in the primary phase. However, their relationship and the exact mechanism are unclear and currently under investigation. In this study, we demonstrated that the addition of W reduces the hydrogen concentration of Nb40Ti30Ni30 alloys without large degradation of hydrogen permeability in the as-cast state, and hydrogen permeability increases due to improved hydrogen diffusivity. If the relationship between hydrogen permeability and concentration is controlled appropriately, substitution of elements having low affinity for hydrogen is expected to produce industrial hydrogen separation and purification alloys compatible with hydrogen permeability and resistance to hydrogen embrittlement.
Table 2 Hydrogen absorption, permeation and diffusion in the Nb40Ti30Ni30 and (Nb35W5) Ti30Ni30 alloys annealed at 1373 K for 168 h. x
0
5
K (molH2 m3 Pa0.5) a (103 molH2 m3) C0.1 (103 molH2 m3) C0.5 (103 molH2 m3) DC (103 molH2 m3)
10.5 15.8 19.1 23.2 4.1
10.3 13.8 17.1 21.0 3.9
F573 (109 molH2 m1 s1 Pa0.5) F0 (106 molH2 m1 s1 Pa0.5)
Ea (kJ mol1)
7.5 2.3 27
11.7 1.7 24
D (109 m2 s1)
0.71
1.1
Experimental error: hydrogen capacity: ±0.2 (103 molH2 m3); hydrogen permeation: ±0.3 (109 molH2 m1 s1 Pa0.5).
4. Summary Microstructure, hydrogen content and hydrogen permeability in Nb40Ti30Ni30 and (Nb40-xWx)Ti30Ni30 alloys were investigated. Hydrogen permeability F of these alloys slightly decreased with increasing W content. Hydrogen concentrations in these alloys were linearly proportional to the square root of hydrogen pressure in the pressure range of 0.1e0.5 MPa. Hydrogen solubility coefficient K was almost independent of W content. Hydrogen diffusivity D decreased slightly with W content. However, 20% or more
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reduction of hydrogen concentration was attained by 5% addition of W without a decrease of the hydrogen concentration differential at 0.1 MPa and 0.5 MPa. After heat treatment, hydrogen permeability of the alloy with 5% W increased due to improved hydrogen diffusivity. Thus, addition of W is a useful technique for producing industrial hydrogen separation and purification alloys compatible with hydrogen permeability and resistance to embrittlement. Acknowledgement This research was supported in part by a Grant-in-Aid for Scientific Research (C) (26420719) and (B) (17H0340) from the Japan Society for the Promotion of Science (JSPS). References [1] E. Kikuchi, Membrane reactor application to hydrogen production, Catal. Today 56 (2000) 97e101. [2] S. Tosti, Supported and laminated Pd-based metallic membranes, Int. J. Hydrogen Energy 28 (2003) 1445e1454. [3] R.E. Buxbaum, A.B. Kinney, Hydrogen transport through tubular membranes of palladium-coated tantalum and niobium, Ind. Eng. Chem. Res. 35 (1996) 530e537. [4] C. Nishimura, M. Komaki, M. Amano, Hydrogen permeation characteristics of vanadium-nickel alloys, Mater. Trans. JIM 32 (1991) 501e507. [5] Y. Zhang, T. Ozaki, M. Komaki, C. Nishimura, Hydrogen permeation characteristics of vanadium-aluminium alloys, Scripta Mater. 47 (2002) 601e606. [6] T. Nambu, K. Shimizu, Y. Matsumoto, R. Rong, N. Watanabe, H. Yukawa, M. Morinaga, I. Yasuda, Enhanced hydrogen embrittlement of Pd-coated niobium metal membrane detected by in situ small punch test under hydrogen permeation, J. Alloys Compd. 446e447 (2007) 588e592.
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