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Changes in microstructure, ductility and hydrogen permeability of Nb–(Ti, Hf)Ni alloy membranes by the substitution of Ti by Hf
Journal of Membrane Science 484 (2015) 47–56
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Changes in microstructure, ductility and hydrogen permeability of Nb–(Ti, Hf)Ni alloy membranes by the substitution of Ti by Hf Xinzhong Li a,n,1, Dongmei Liu b,1, Ruirun Chen a, Erhu Yan a, Xiao Liang a, Markus Rettenmayr b, Yanqing Su a, Jingjie Guo a, Hengzhi Fu a a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Otto Schott Institute of Materials Research, Friedrich Schiller University, Jena 07743, Germany
art ic l e i nf o
a b s t r a c t
Article history: Received 11 September 2014 Received in revised form 13 January 2015 Accepted 1 March 2015 Available online 10 March 2015
Microstructure, ductility and hydrogen permeability of as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 x HfxNi30 (x¼0…30) alloys are investigated. The two series of alloys mainly consist of primary BCC-(Nb, Ti, Hf) and eutectic structure [BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni/Orthorhombic-(Hf, Ti)Ni]. The substitution of Ti by Hf induces higher fraction of primary phase and less eutectic. When the Hf content in an alloy is close to or higher than that of Ti, the ductile B2-(Ti, Hf)Ni in eutectic turns to brittle Orthorhombic-(Hf, Ti)Ni and some other brittle phases appear which distinctly reduce ductility. The hydrogen permeability of both alloy systems increases with increasing Hf content, which results from the simultaneous increase of hydrogen solubility and diffusivity. The Hf content in practically applicable alloys should be lower than that of Ti for high tolerance to hydrogen embrittlement. Nb40Ti20Hf10Ni30 exhibits a high cold-rolling reduction ratio of 63%, a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (which is 1.85 times that of pure Pd) and excellent resistance against hydrogen embrittlement. This alloy is thus promising to fabricate membranes for hydrogen separation. & 2015 Elsevier B.V. All rights reserved.
Keywords: Hydrogen permeable membrane Microstructure Nb–(Ti, Hf)Ni alloy
1. Introduction Hydrogen permeable metal membranes are widely used for various industrial applications, including ultrapure hydrogen production for semiconductor manufacturing and nuclear fusion engineering [1–4], carbon monoxide removal from reformed gas for fuel cells [5] and membrane reactors to improve and control chemical reactions [6]. At present, Pd-based alloys exhibit the best combination of high hydrogen selectivity and permeability, high thermal stability and especially high ductility which is the key property for the fabrication of thin membranes [7]. However, Pd is expensive and scarce. There is a high demand for low-cost metal membranes with minimal content or complete absence of Pd, but exhibiting similarly high hydrogen permeability, ductility and resistance to the hydrogen embrittlement. Both crystalline alloys with body-centered-cubic (BCC) structure and amorphous alloys are promising alternatives to Pd-based membranes for hydrogen separation because of their lower cost and comparable hydrogen permeability. During the last decade, a series of amorphous alloys derived from Ni–Nb–Zr system [8–14] were developed. These materials exhibit high hydrogen permeability with mechanical stability
n
Corresponding author. Tel./fax: þ 86 451 86418754. E-mail address: [email protected] (X. Li). 1 These authors contributed equally to this work.
http://dx.doi.org/10.1016/j.memsci.2015.03.002 0376-7388/& 2015 Elsevier B.V. All rights reserved.
in hydrogen atmosphere in appropriate temperature regimes. However, their applicability at high temperatures is limited by the thermal stability of the amorphous alloys. Crystallization of the amorphous membranes reduces both hydrogen permeability and strength significantly [10]. In contrast, crystalline BCC solid-solution alloys comprising Group V metals such as V, Nb and Ta exhibit much higher hydrogen permeability and thermal stability [15–19]. The high hydrogen solubility provides them a high driving force for hydrogen permeation, which however aggravates the problem of hydrogen embrittlement. To address this problem, a new type of hydrogen permeable crystalline alloys has been developed, such as Nb–TiNi [20–24], Nb–TiCo [25,26], Ta–TiNi [27,28], V–Ti–Ni [29–31], etc. These alloys contain two phases, particularly a primary BCC solid-solution phase that provides the major contribution to hydrogen permeation, and in addition a secondary eutectic phase that suppresses the hydrogen embrittlement. Consequently, such alloys are promising to achieve a sound combination of high hydrogen permeability, excellent resistance against hydrogen embrittlement and high ductility. During the past several years, hydrogen permeable Nb–TiNi alloys consisting of primary BCC-(Nb, Ti) and eutectic {BCC-(Nb, Ti)þB2TiNi} phases have attracted intensive attention. This alloy system exhibits high ductility for fabrication of thin membranes and high tolerance to hydrogen embrittlement, but the hydrogen permeability is not quite satisfactory. It is found that appropriate substitution of Ti by Zr can induce a much higher hydrogen permeability (because of the higher hydrogen affinity of Zr than that of Ti) and comparable
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other properties [32,33]. Nevertheless, excessive addition of Zr leads to low ductility and serious hydrogen embrittlement due to the microstructural change and the inappropriately high hydrogen solubility [34]. Therefore, much of current research focuses on compositional modifications in these alloys to tailor the microstructures and hydrogen solubility and diffusivity for improved performance. The element Hf, in the same IVA group as Zr, also possesses a higher hydrogen affinity than that of Ti but lower than that of Zr [7,12]. It is most likely that appropriate substitution of Ti by Hf in Nb–TiNi alloys results in new hydrogen permeable alloys with high performance. In this work, microstructure and its influence on ductility and hydrogen permeability of as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys are investigated. The former alloys are based on eutectic Nb20Ti40Ni40, so that the microstructural changes induced by the addition of Hf in this system may be observed clearly. The latter alloys are based on the hypereutectic Nb40Ti30Ni30 with the aim to develop new hydrogen permeable alloys with high performance by appropriate substitution of Ti by Hf. A maximum coldrolling reduction ratio is used to evaluate ductility by measuring thickness changes in the samples. The solubility and diffusivity of hydrogen in the two series of alloys are characterized for revealing the detailed response of hydrogen permeability to the Hf content.
Table 1 Summary of constituting phases and their respective chemical compositions (EDX results) in the as-cast Nb20Ti40 xHfxNi40 (x ¼0…40) and Nb40Ti30 xHfxNi30 (x¼ 0… 30) alloys. Samples
Nb20Ti40Ni40 Nb20Ti35Hf5Ni40 Nb20Ti30Hf10Ni40 Nb20Ti25Hf15Ni40 Nb20Ti20Hf20Ni40 Nb20Ti15Hf25Ni40 Nb20Ti10Hf30Ni40 Nb20Ti5Hf35Ni40 Nb20Hf40Ni40 Nb40Ti30Ni30 Nb40Ti25Hf5Ni30 Nb40Ti20Hf10Ni30 Nb40Ti15Hf15Ni30 Nb40Ti10Hf20Ni30 Nb40Ti5Hf25Ni30 Nb40Hf30Ni30
Chemical composition of the constituting phases (mol%) Primary BCC-(Nb, Ti, Hf) phase
{BCC-(Nb, Ti, Hf) þ B2-(Ti, Hf) Ni} eutectic
Nb
Ti
Hf
Ni
Nb
– – 77.18 76.94 75.83 76.98 77.82 77.01 77.58 82.92 80.94 81.65 82.38 81.59 82.25 82.81
– – 10.19 8.45 8.22 6.79 3.09 2.42 0 12.65 11.72 9.59 5.83 4.46 2.98 0
– – 5.75 7.12 8.41 9.86 11.07 13.45 14.99 0 2.93 5.11 7.86 9.57 11.28 13.51
– 20.03 – 19.93 6.88 19.81 7.49 18.45 7.54 17.12 6.37 16.28 8.02 15.31 7.12 14.75 7.43 14.67 4.43 19.83 4.41 19.35 3.65 19.11 3.93 18.14 4.38 17.62 3.49 18.03 3.68 17.88
Ti
Hf
Ni
39.82 35.37 31.02 24.89 19.43 14.86 9.65 5.13 0 38.51 31.69 22.34 18.12 14.78 7.42 0
0 4.87 13.29 19.94 25.71 31.19 38.82 44.13 50.84 0 9.28 17.21 23.15 28.87 33.26 39.35
40.15 39.83 35.88 36.72 37.74 37.67 36.22 35.99 34.49 41.66 39.68 41.34 40.59 38.73 41.29 42.77
2. Experimental Ingots of about 20 g of Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼0…30) (all concentrations in at% if not specified differently) alloys were prepared by arc melting in a purified argon atmosphere using pure Nb, Ti, Hf and Ni (99.9 mass% purity for all). The crystal structures of the alloys were identified by powder X-ray diffractometry (XRD). Microstructural observations were carried out using a Scanning Electron Microscope (SEM) equipped with energy dispersive X-ray (EDX) spectrometry. The volume fraction of the constituent phases was analyzed using the public domain NIH Image program. Samples of 7 3 3 mm3 in dimension were cut from Nb40T i30 xHfxNi30 (x¼0…30) alloy ingots using a spark erosion wirecutting machine for the cold-rolling experiment at room temperature. Disk samples of 12 mm in diameter and 0.6 mm in thickness were cut from two series of alloys for the measurement of hydrogen permeation. Both sides of the disks were ground, polished and then coated with Pd of 190 nm in thickness by a radio frequency (RF) sputtering machine to prevent oxidation during permeation experiments and to act as catalyzer for hydrogen uptake into the material. The disks were sealed by copper gaskets. Both sides of the disks were evacuated using a diffusion pump to below 3 10 3 Pa, then heated to 673 K and kept for 20 min. Hydrogen gas (99.99999 mass% purity) of 0.2–0.4 MPa was introduced at the upstream side (Pu) and of 0.1 MPa at the downstream side (Pd). The hydrogen flux, J, passing through the disks was measured by a hydrogen flow meter. The measurements were repeated at 623 K, 573 K and 523 K. The pressure-composition-temperature (PCT) curves for some representative alloys were measured using a Sieverts-type apparatus at 673 K and in the pressure range from 0.01 to 0.9 MPa to investigate the hydrogen solubility. The experimental procedure was identical with those in Refs. [35–37]. The amount of absorbed hydrogen was calculated from the pressure drop in a constant inner volume chamber.
3. Results and discussion 3.1. Microstructures The constituting phases and their respective chemical compositions in the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) and Nb40Ti30 x
HfxNi30 (x¼0…30) alloys as determined by XRD, SEM and EDS are tabulated in Table 1. The XRD patterns of the two series of alloys are shown in Figs. 1 and 2. Nb–(Ti, Hf)Ni alloys with an initial Hf content below that of Ti consist of BCC-(Nb, Ti, Hf) solid solution and B2-(Ti, Hf)Ni, and the substitution of Ti by Hf does not lead to a change of the constituting phases. However, if the Hf content is close to or higher than that of Ti, some intermetallic phases such as HfNi, Hf3Ni7, Hf8Ni21, Hf2Ni, HfNi3, Hf2Ni7, etc. appear additionally. With increasing Hf content, the diffraction peak intensity of these intermetallic phases increases, especially that of Orthorhombic-HfNi phase. For the Nb20Ti40 xHfxNi40 (x¼ 25…40) alloys, the intensity of the main diffraction peak of the HfNi is comparable to or even higher than that of BCC-Nb (see Fig. 1). This indicates that a large amount of HfNi phase is present in these alloys. In comparison, for the Nb40 Ti30 xHfxNi30 (x¼15…30) alloys, the diffraction peak intensity of the HfNi phase is much lower than that of BCC-Nb (see Fig. 2). In addition, the main diffraction peak of B2-(Ti, Hf)Ni corresponding to 431 decreases in intensity with the gradual substitution of Ti by Hf in both series of alloys. This peak vanishes finally when the Hf content in alloys is close to or higher than that of Ti, and simultaneously the diffraction peak of Orthorhombic-(Hf, Ti)Ni corresponding to 271 becomes obvious and grows in intensity with further increasing the Hf content. It should be noted that B2-(Ti, Hf)Ni results from the replacement of Ti by Hf in B2-TiNi lattice and thus retains the ductile characteristics. However, the Orthorhombic-(Hf, Ti)Ni results from the replacement of Hf by Ti in Orthorhombic-HfNi lattice, which is intrinsically brittle [32]. SEM micrographs of the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) alloys are shown in Fig. 3. A fully eutectic microstructure, {BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni}, is observed for Nb20Ti40Ni40 (x¼ 0) and Nb20Ti35Hf5Ni40 (x¼ 5) alloy, Fig. 3(a) and (b). With further increase of Hf content (x), primary BCC-(Nb, Ti, Hf) dendrites embedded in eutectic are found in Nb20Ti30Hf10Ni40 (x¼10) and Nb20Ti25Hf15Ni40 (x¼15) alloys, Fig. 3(c) and (d). For alloys with xZ20, a similar hypereutectic microstructure appears, Fig. 3(e)–(i); but the B2-(Ti, Hf) Ni in eutectic turns to Orthorhombic-(Hf, Ti)Ni (see Fig. 1). The eutectic morphology turns to be rod-like, compared to the lamellar one in the alloys with xo20, see the higher magnification inset image in Fig. 5. Fig. 4 shows SEM micrographs for the as-cast Nb40Ti30 xHfxNi30 (x¼0…30) alloys. Primary BCC-(Nb, Ti, Hf) dendrites embedded in
X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
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Fig. 1. XRD patterns of the as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) alloys. Fig. 2. XRD patterns of the as-cast Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys.
eutectic [BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni/Orthorhombic-(Hf, Ti)Ni] are mainly found in these alloys. When the Hf content is close to or higher than that of Ti, the B2-(Ti, Hf)Ni in eutectic turns to Orthorhombic(Hf, Ti)Ni (see Fig. 2), and the lamellar eutectic changes to rod-like one (see insets in Fig. 5) as that in Nb20Ti40 xHfxNi40 alloys. For the above two series of alloys, the substitution of Ti by Hf induces higher volume fractions of the primary phase and less eutectic, as shown in Fig. 5. This can be related to the fact that the eutectic Nb20Ti40Ni40 is closer to the Nb corner of the ternary phase diagram than that of eutectic Nb16Hf42Ni42 [38]. Correspondingly, the effect of the above substitution on the solidification behavior is equivalent to increasing the initial Nb content of ternary Nb–TiNi alloys. The solid-solution Hf in primary BCC-Nb dendrites increases with more Hf addition, as illustrated in Table 1. When the initial Hf content in alloys is close to or higher than that of Ti, the ductile B2-(Ti, Hf)Ni in eutectic turns to brittle Orthorhombic-(Hf, Ti)Ni and some intermetallic phases (HfNi, HfNi3, etc.) appear additionally at the boundaries of eutectic grains (see the insets in Fig. 5). This corresponds to the XRD results in Figs. 1 and 2. It should be noted that these intermetallic phases are inherently brittle and cannot be eliminated by annealing [39]. Further evaluation of ductility is needed for alloys containing these phases. 3.2. Ductility Considering the workability of hydrogen permeable alloy membranes, e.g. mechanically sealed with copper gaskets, these alloys must have at least some ductility. Most simply, a material that can be broken easily by hammering is qualitatively defined as brittle [32–34], while non-broken materials are defined as ductile. Here, the maximum cold-rolling reduction ratio is used to more reproducibly evaluate the ductility of alloys. A qualitative ductility
index for Nb20Ti40 xHfxNi40 (x ¼0…40) and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys is tabulated in Tables 2 and 3, respectively. Overall, alloys with Hf content lower than that of Ti are found to be ductile, because their constituting phases of BCC-(Nb, Ti, Hf) and B2-(Ti, Hf)Ni are ductile. Alloys with Hf content close to or higher than that of Ti are brittle due to the existence of brittle intermetallic phases. Fig. 6 shows the maximum cold-rolling reduction ratio of the as-cast Nb40Ti30 xHfxNi30 alloys as a function of the Hf content (x). The cold-rolling reduction ratio (rt) is defined as rt ¼(t0 t)/t0, where t0 and t are the original and the final thicknesses of the sample, respectively [32]. It can be seen that 60% or higher reduction ratio (rt) was obtained for Hf contents of 10 at% or less. In contrast, if the Hf content is 15% or higher, rt is reduced significantly. This result appropriately corresponds to the above XRD results and microstructural observations. Therefore, the Hf content in alloys should be limited to that of Ti for good ductility in the as-cast state. 3.3. Hydrogen permeability 0:5 Fig. 7 shows the relation between J L and ΔP0.5 ( ¼ P 0:5 u Pd ) for the as-cast Nb20Ti30Hf10Ni40 and Nb40Ti20Hf10Ni30 alloys from 523 to 673 K, where J is the measured hydrogen flux, L is the thickness of the membrane and ΔP is the difference of the hydrogen pressure at the two sides of the membrane. In the plot, the experimental data are quite precisely positioned on straight lines passing through the origin, and the correlation coefficient R2, obtained by linear regression, is higher than 0.99. A similar relation between J L and ΔP0.5 appears in other alloys. However, this does not indicate that these alloys follow Sieverts' law at the
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Fig. 3. SEM micrographs of as-cast Nb20Ti40 xHfxNi40 alloys with increasing x (x¼ 0…40). (a) Nb20Ti40Ni40, (b) Nb20Ti35Hf5Ni40, (c) Nb20Ti30Hf10Ni40, (d) Nb20Ti25Hf15Ni40, (e) Nb20Ti20Hf20Ni40, (f) Nb20Ti15Hf25Ni40, (g) Nb20Ti10Hf30Ni40, (h) Nb20Ti5Hf35Ni40, and (i) Nb20Hf40Ni40.
present pressures of hydrogen permeation. In fact, it has been found that hydrogen dissolution in niobium and its alloys does not follow Sieverts' law in the high hydrogen content region [40–43]. Therefore, further information about the pressure dependence of hydrogen solubility in the present alloys is needed [22,43]. Figs. 8 and 9 show the PCT curves for the as-cast Nb20 Ti40 xHfxNi40 (x ¼0, 10, 20, 30, 40) alloys and Nb40Ti30 xHfxNi30 (x ¼0, 10, 20, 30) alloys, respectively, measured at 673 K. They are plotted in the form of Sieverts' plot, i.e., hydrogen content C vs. the square root of the hydrogen pressure P0.5. This relation is crucial to define the hydrogen permeability for the hydrogen permeable alloys [41]. It can be seen that Sieverts' plot for each alloy shows straight line between 0.1 and 0.4 MPa, but does not pass through the origin. This indicates that Sieverts' law does not hold for these alloys under the present conditions. At pressures lower than 0.1 MPa, C vs. P0.5 is close to a straight line, but its slope is larger than that at 0.1–0.4 MPa. At pressures higher than 0.4 MPa, C vs. P0.5 deviates from a linear relation, and its slope tends to reduce with increasing pressures, which is also lower than that at low pressures. Actually, this slope corresponds to the hydrogen solubility coefficient (K). It is obvious that the values of K gradually reduce from low pressures (o0.1 MPa) to intermediate pressures
(0.1–0.4 MPa) and then high pressures ( 40.4 MPa). This hydrogen absorption behavior is a reflection of changes to the progressive filling of interstitial sites with increasing hydrogen content. At low pressures, the sites with the lowest energy are preferentially occupied by hydrogen atoms. A broad distribution of site energies for these alloys results in a significantly enhanced occupation with slightly increasing pressures. Correspondingly, C vs. P0.5 exhibits a relatively steep rise of line, which corresponds to a large K value. With further increasing pressures, the filling of the most favorable sites will be gradually finished and occupation of less favored sites occurs. This induces a lower K value, and the pressure required to achieve a given hydrogen content increases. This situation is enhanced particularly at high pressures, because most of hydrogen atoms have to stay among less favored sites. Therefore, the curve of C vs. P0.5 tends to slope gently at high pressures, which corresponds to the lowest K value. The hydrogen content in the pressure range of 0.1–0.4 MPa for each alloy, C, can be expressed as C ¼ K U P 0:5 þ α
ð1Þ
where K is the hydrogen solubility coefficient and α is a constant. Tables 2 and 3 show the values of K and α for each alloy, determined
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Fig. 4. SEM micrographs of as-cast Nb40Ti30 xHfxNi30 alloys with increasing x (x ¼0…30). (a) Nb40Ti30Ni30, (b) Nb40Ti25Hf5Ni30, (c) Nb40Ti20Hf10Ni30, (d) Nb40Ti15Hf15Ni30, (e) Nb40Ti5Hf25Ni30, and (f) Nb40Hf30Ni30.
by linear regression. It can be seen that the values of K for two series of alloys increase with the gradual substitution of Ti by Hf. The steady-state hydrogen flux, J, through a metal membrane is defined by Fick's first law ∂C J ¼ D ð2Þ ∂L where D is the hydrogen diffusion coefficient and ∂C/∂L is the gradient of the hydrogen content across the membrane. If C in Eq. (2) is substituted by C from Eq. (1), the hydrogen flux, J, can be
linearized as follows: J¼
D KðΔP 1=2 Þ ΦðΔP 1=2 Þ ¼ L L
ð3Þ
This implies that we can still define the hydrogen permeability (Φ) for the present alloys in a limited hydrogen pressure range, even if Sieverts' law does not hold. The slope of each straight line in the plot of J L vs. ΔP0.5 in Fig. 7 corresponds to the hydrogen permeability (Φ). This method is commonly used to determine Φ [22,43].
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Fig. 10 shows the temperature dependence of Φ in the form of an Arrhenius plot for the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) alloys and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys. The Φ values of pure Pd [44] are also included in the figure for comparison. The variations of Φ appear to be straight lines in the Arrhenius plot in all cases, and the Φ values of each alloy evidently increase with increasing temperature. The effect of temperature on the Φ values for present alloys is larger than that of pure Pd, i.e., the activation energy for hydrogen permeation (E) through these alloys is larger than that of pure Pd. The values of E for Nb–(Ti, Hf)Ni alloys range from 24.5 to 33.2 kJ mol 1 (see Tables 1 and 2), while that for pure Pd is 12.21 kJ mol 1. It can also be seen that the Φ values of both series of alloys increase with the gradual substitution of Ti by Hf. Most alloys in the Nb20Ti40 xHfxNi40 (x ¼0…40) system exhibit a lower Φ value than that of pure Pd (see Fig. 10(a)). Nb20Ti15Hf25Ni40 and Nb20Ti10Hf30Ni40 are prone to embrittlement and broke during hydrogen permeation at 523 K, so that the Φ value could not be measured. The problem of hydrogen embrittlement is aggravated for Nb20Ti5Hf35Ni40 and Nb20Hf40Ni40 at 523 K and 573 K. In contrast, most alloys in the Nb40Ti30 xHfxNi30 (x ¼0…30) system exhibit much higher Φ values than pure Pd (see Fig. 10(b)), especially at high temperatures. No cracks were found after the hydrogen permeability measurements at all temperatures for these alloys, indicating that they are less susceptible to hydrogen embrittlement during hydrogen permeation. According to the above defined K and Φ values, the hydrogen diffusion coefficient (D) in each alloy can be evaluated on the basis of Φ ¼D K, which is tabulated in Tables 2 and 3. In this way, a relationship of permeability, solubility and diffusivity of hydrogen
Fig. 5. Relationship of Hf content and volume fraction of primary BCC-(Nb, Ti, Hf) in the as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys, respectively. The inset images show the eutectic morphology and formation of additional new intermetallic phases at the boundaries of eutectic grains (labeled by arrows) in three representative alloys.
can be built for the as-cast Nb20Ti40 xHfxNi40 (x ¼0…40) alloys and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys. The hydrogen permeability of both alloys increasing with Hf content is found to be induced by the simultaneous increase of the hydrogen solubility and diffusivity (see Tables 2 and 3). The hydrogen solubility in metals is governed mainly by the size and the electronic structure of the interstitial sites associated with the change of the interstitial binding energy [45,46]. For the BCC-Nb, octahedral and tetrahedral interstices exist. The occupation of tetrahedral sites by hydrogen atoms in Group V BCC lattices is more favored than the octahedral ones [47–50]. Due to the larger atomic radius of Hf (159 pm) as compared to that of Ti (147 pm), the substitution of Ti by Hf in BCC-Nb will increase the size of the tetrahedral interstitial sites. This may potentially increase the tolerance to the lattice expansion for accommodating more hydrogen atoms and thus the hydrogen solubility. From the view of the electronic effect, the interstitial binding energy is related to the hydrogen charge where the stability of the site is reduced when hydrogen gains less charge for the group V materials [47]. In the pure niobium lattice, hydrogen gains charge from each of the nearest neighbor niobium atoms, because the electronegativity of hydrogen (2.2) is higher than that of niobium (1.6). When the neighboring niobium atoms are replaced by elements with a similar or higher electronegativity than that of hydrogen, less charge is shared between the lattice and hydrogen, thereby decreasing the binding energy. In contrast, for the present Nb–(Ti, Hf)Ni alloys, Ti and Hf in solid solution in BCC-Nb have a lower electronegativity than that of hydrogen. The electronegativity of Hf (1.3) is lower than that of Ti (1.54). It is expected that the substitution of Ti by Hf may increase the charge shared between the lattice and hydrogen and thus increase the binding energy and the hydrogen solubility. It also should be noted that alloys with Hf content close to or beyond that of Ti contain a certain amount of HfNi, especially for the Nb20Ti40 xHfxNi40 (x ¼25…40) alloys (as discussed in Section 3.1). This phase is known to absorb hydrogen for a long time [51–53], and its hydrogenation contributes partially to the high hydrogen solubility for these alloys. This provides a positive contribution to hydrogen permeability, but also induces a low tolerance to hydrogen embrittlement. In comparison, Nb40Ti30 x HfxNi30 (x ¼15…30) alloys contain a less HfNi phase, and they are less susceptible to hydrogen embrittlement during hydrogen permeation. Hydrogen diffusion in metals is described by classical hopping between interstitial sites, which is related to the hydrogen content in metals [54,55]. At dilute contents, hydrogen atoms are essentially found in the sites with the strongest binding energies, so that longrange diffusion is limited by the large barriers that exist to moving hydrogen atoms away from these sites. At higher contents, however, the most favorable sites are occupied by some population of hydrogen atoms, which allows other hydrogen atoms to stay and hop among less favored sites. The latter population moves far more rapidly than atoms trapped in the most favored sites, so they dominate the overall diffusion coefficient [54]. Therefore, the diffusivity of hydrogen increases markedly as the hydrogen content increases. This has been
Table 2 Hydrogen permeability (Ф), hydrogen solubility coefficient (K), hydrogen diffusion coefficient (D), fitting parameter (α) at 673 K and activation energy for hydrogen permeation (E) along with the ductility for the as-cast Nb20Ti40 xHfxNi40 alloys. Samples
Ф [10 8 mol H2 m 1 s 1 Pa 0.5]
K [mol H2 m 3 Pa 0.5]
D [10 10 m2 s 1]
α [mol H2 m 3]
E [kJ mol 1]
Nb20Ti40Ni40 Nb20Ti30Hf10Ni40 Nb20Ti20Hf20Ni40 Nb20Ti10Hf30Ni40 Nb20Hf40Ni40
0.64 0.84 1.19 1.63 2.01
21.36 23.56 26.7 27.88 31.91
3.0 3.57 4.64 5.85 6.30
2339 3200 3472 4291 5043
26.8 28.7 32.2 33.1 31.7
Ductile Ductile Brittle Brittle Brittle
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Table 3 Hydrogen permeability (Ф), hydrogen solubility coefficient (K), hydrogen diffusion coefficient (D), fitting parameter (α) at 673 K and activation energy for hydrogen permeation (E) along with the ductility for the as-cast Nb40Ti30 xHfxNi30 alloys. Samples
Ф [10 8 mol H2 m 1 s 1 Pa 0.5]
K [mol H2 m 3 Pa 0.5]
D [10 10 m2 s 1]
α [mol H2 m 3]
E [kJ mol 1]
Nb40Ti30Ni30 Nb40Ti20Hf10Ni30 Nb40Ti10Hf20Ni30 Nb40Hf30Ni30
1.76 2.96 3.62 4.33
24.29 25.57 23.37 29.09
7.25 11.58 13.23 14.89
5450 6350 7790 9860
24.5 32.1 33.2 31.8
Ductile Ductile Brittle Brittle
Fig. 6. Maximum cold-rolling reduction ratio of the as-cast Nb40Ti30 xHfxNi30 alloys as a function of the Hf content.
testified by a variety of experiments and first-principle calculations [56–58]. For the present Ni–(Ti, Hf)Ni alloys, the substitution of Ti by Hf induces an increase of hydrogen contents (see Figs. 8 and 9), which potentially leads to the increase of the hydrogen diffusivity. 3.4. Resistance against hydrogen embrittlement (HE) For hydrogen permeable metal membranes, the large resistance against HE is crucial to the durability with respect to long-term hydrogen permeation. The short-term tests in Fig. 10 indicate that ascast Nb20Ti40 xHfxNi40 (xr20) and Nb40Ti30 xHfxNi30 (x¼0…30) alloys exhibit an appropriate tolerance to HE. However, such information is limited as it does not consider long-term durability which is of greater interest. Fig. 11 shows the variations of hydrogen permeability with time in 72 h for four representative alloys at 673 K. It can be seen that Nb20Ti25Hf15Ni40 and Nb40Ti20Hf10Ni30 exhibit excellent durability in hydrogen permeation and their Φ values are almost stable after initial transient. The final membranes remain totally free from cracks or pores after a test of 72 h, which were confirmed by a room temperature permeation test with Ar. This indicates that both alloys show large resistance against HE. In contrast, Nb40Ti15Hf15Ni30 and Nb20Ti20Hf20Ni40 are subjected to HE failure after hydrogen permeation of 15 h and 23 h, respectively. This is not detectable in a short-time test of hydrogen permeation, because these two alloys also exhibit stable permeability during the early 10 or 20 h. For the present hydrogen permeable alloys, the eutectic in as-cast microstructures plays an important role in resisting HE. Once hydrogen permeates, primary BCC-Nb phases dilate with the extent depending on the solubility of hydrogen, and the surrounding eutectic may tolerate this dilation due to its high deformability [59]. A high fraction of eutectic is thus beneficial to the embrittlement resistance. However, it should be noted that BCC-Nb in eutectic also act as pathways for hydrogen permeation, and their dilation after
Fig. 7. Relation between J L and ΔP0.5 for the as-cast Nb20Ti30Hf10Ni40 (a) and Nb40Ti20Hf10Ni30 (b) in the temperature range of 523–673 K.
hydrogen absorption exerts force to the surrounding intermetallic phases. For Nb20Ti25Hf15Ni40 and Nb40Ti20Hf10Ni30, ductile B2-(Ti, Hf) Ni provides a high tolerance to the dilation of BCC-Nb in the eutectic. Therefore, the high fractions of eutectic ( 68 vol% or higher) in ascast microstructures of these two alloys and the ductile B2-type intermetallic phase in eutectic ensure their large resistance against HE. In comparison, there are also high fractions of eutectic in Nb20Ti20Hf20Ni40 and Nb40Ti15Hf15Ni30 which provide an appropriate resistance against HE. However, brittle Orthorhombic-(Hf, Ti)Ni in their eutectic provide a low tolerance to the dilation of surrounding BCC-Nb. Especially, a high hydrogen solubility in Orthorhombic-(Hf, Ti)Ni is expected as its parental HfNi, and its hydrogenation also induces a large lattice expansion. These result in a deterioration of embrittlement resistance. The positive contribution from the eutectic
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X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
Fig. 8. Pressure-composition-temperature (PCT) curves for the as-cast Nb20Ti40 xHfxNi40 (x ¼0, 10, 20, 30, 40) alloys at 673 K, and the hydrogen content C (mol H2 m 3) is plotted as a function of the square root of the hydrogen pressure P0.5 for each alloy, i.e., in form of a Sieverts' plot.
Fig. 10. Temperature dependence of hydrogen permeability (Ф) in the form of an Arrhenius plot for the as-cast alloys of (a) Nb20Ti40 xHfxNi40 (x¼ 0…40) and (b) Nb40Ti30 xHfxNi30 (x ¼ 0…30).
Fig. 9. Pressure-composition-temperature (PCT) curves for the as-cast Nb40Ti30 xHfxNi30 (x¼ 0, 10, 20, 30) alloys at 673 K, and the hydrogen content C (mol H2 m 3) is plotted as a function of the square root of the hydrogen pressure P0.5 for each alloy, i.e., in form of a Sieverts' plot.
and the negative one from the brittle Orthorhombic-(Hf, Ti)Ni appropriately explain the HE failure after some time of hydrogen permeation for these two alloys. This can be further extended to deduce that Nb20Ti40 xHfxNi40 (xo20) and Nb40Ti30 xHfxNi30 (xo15) may exhibit a large resistance against HE because of their high fractions of eutectic and ductile B2-(Ti, Hf)Ni in their eutectic structure. In a general view taking into account the microstructure, ductility, hydrogen permeability and the resistance against hydrogen embrittlement of the above Nb-(Ti, Hf)Ni alloys, Nb40Ti20Hf10Ni30 is a promising candidate to fabricate high-performance membranes for hydrogen separation. This alloy consists of the appropriate microstructure for hydrogen permeation (without any intermetallic phases besides the eutectic), exhibits a high cold-rolling reduction of 63%,
Fig. 11. Variations of hydrogen permeability with time in 72 h for the as-cast Nb20Ti25Hf15Ni40, Nb20Ti20Hf20Ni40, Nb40Ti20Hf10Ni30 and Nb40Ti15Hf15Ni30 at 673 K.
a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (which is 1.85 times that of pure Pd) and excellent resistance against hydrogen embrittlement. Table 4 shows the hydrogen permeability for some representative alloys, and their volume fractions of primary phase and eutectic and the tolerance to HE are also
X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
Table 4 Hydrogen permeability along with the volume fractions of primary phase and eutectic and the tolerance to HE for some representative alloys, obtained from the research group of Aoki using the same characterizing methods. Alloys
Nb19Ti40Ni41 [22] Nb40Ti30Ni30 [21] Nb56Ti23Ni21 [20] Nb30Ti35Co35 [25] Nb60Ti21Co19 [26] V41Ti30Ni29 [27] Nb40Zr30Ni30 [60] Nb40Hf30Ni30 [39] Nb40Ti18Zr12Ni30 [33] Nb40Ti18Zr12Ni25Co5 [33] Nb40Ti15Hf10Ni30
Volume fractions (vol%) Primary BCC-Nb phase
Eutectic
Hydrogen permeability (Ф) at 673 K (10 8 mol H2 m 1 s 1 Pa 0.5)
0 29 62 0 62 40 40 40 45
100 71 38 100 38 60 60 60 55
0.64 1.8 3.47 2.64 3.99 1.1 4.64 4.3 2.85
High High Low High Low High Low Low High
42
58
3.82
High
32
68
2.96
High
Tolerance to HE
included. Although containing a low fraction of primary BCC-Nb, Nb40Ti20Hf10Ni30 exhibits moderate permeability among these alloys. Simultaneously, the high fraction of eutectic ensures its high tolerance to HE. Therefore, Nb40Ti20Hf10Ni30 clearly reaches an excellent balance between permeability and embrittlement resistance. In comparison, Nb56Ti23Ni21 and Nb60Ti21Co19 exhibit higher permeability than Nb40Ti20Hf10Ni30, but their low fractions of eutectic induce a low tolerance to HE. The high permeability of Nb40Zr30Ni30 and Nb40Hf30Ni30 is originated from the high hydrogen solubility in BCC-Nb and also Orthorhombic-ZrNi/HfNi in their microstructures. The hydrogenation of Orthorhombic-ZrNi/HfNi results in severe HE. Nb30Ti35Co35, Nb40Ti18Zr12Ni30 and Nb40Ti18Zr12Ni25Co5 also reach an appropriate balance between permeability and embrittlement resistance. In contrast, Nb40Ti20Hf10Ni30 exhibits higher permeability than Nb30Ti35Co35 and Nb40Ti18Zr12Ni30, but lower one than Nb40Ti18Zr12Ni25Co5. It is expected that partial substitution of Ni by Co for Nb40Ti20Hf10Ni30 induces a higher permeability.
4. Conclusions Nb-based alloys are an emerging alternative to palladium alloys for hydrogen separation, while the key challenge remains to balance the ductility, hydrogen permeability and resistance against HE. This work has shown that substitution of Ti by Hf in Nb–(Ti, Hf)Ni alloys increases hydrogen solubility and diffusivity and thus hydrogen permeability, but greater hydrogen solubility brings the penalty of poorer resistance against HE. When the Hf content in an alloy is close to or higher than that of Ti, some undesired brittle phases appear in microstructures which distinctly reduce ductility. Correspondingly, the Hf content in practically applicable alloys should be lower than that of Ti for high ductility and high tolerance to HE. Nb40Ti20Hf10Ni30 reaches an excellent balance among ductility, hydrogen permeability and embrittlement resistance. This alloy exhibits a high cold-rolling reduction of 63%, a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (1.85 times that of pure Pd) and excellent durability for long-term permeation, which is promising to fabricate membranes for hydrogen separation.
Acknowledgments This project was supported by the National Natural Science Foundation of China (Grants nos. 51274077 and 51271068) and
55
the National Basic Research Program of China 973 (Grant no. 2011CB610406).
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Contents lists available at ScienceDirect
Journal of Membrane Science journal homepage: www.elsevier.com/locate/memsci
Changes in microstructure, ductility and hydrogen permeability of Nb–(Ti, Hf)Ni alloy membranes by the substitution of Ti by Hf Xinzhong Li a,n,1, Dongmei Liu b,1, Ruirun Chen a, Erhu Yan a, Xiao Liang a, Markus Rettenmayr b, Yanqing Su a, Jingjie Guo a, Hengzhi Fu a a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China Otto Schott Institute of Materials Research, Friedrich Schiller University, Jena 07743, Germany
art ic l e i nf o
a b s t r a c t
Article history: Received 11 September 2014 Received in revised form 13 January 2015 Accepted 1 March 2015 Available online 10 March 2015
Microstructure, ductility and hydrogen permeability of as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 x HfxNi30 (x¼0…30) alloys are investigated. The two series of alloys mainly consist of primary BCC-(Nb, Ti, Hf) and eutectic structure [BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni/Orthorhombic-(Hf, Ti)Ni]. The substitution of Ti by Hf induces higher fraction of primary phase and less eutectic. When the Hf content in an alloy is close to or higher than that of Ti, the ductile B2-(Ti, Hf)Ni in eutectic turns to brittle Orthorhombic-(Hf, Ti)Ni and some other brittle phases appear which distinctly reduce ductility. The hydrogen permeability of both alloy systems increases with increasing Hf content, which results from the simultaneous increase of hydrogen solubility and diffusivity. The Hf content in practically applicable alloys should be lower than that of Ti for high tolerance to hydrogen embrittlement. Nb40Ti20Hf10Ni30 exhibits a high cold-rolling reduction ratio of 63%, a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (which is 1.85 times that of pure Pd) and excellent resistance against hydrogen embrittlement. This alloy is thus promising to fabricate membranes for hydrogen separation. & 2015 Elsevier B.V. All rights reserved.
Keywords: Hydrogen permeable membrane Microstructure Nb–(Ti, Hf)Ni alloy
1. Introduction Hydrogen permeable metal membranes are widely used for various industrial applications, including ultrapure hydrogen production for semiconductor manufacturing and nuclear fusion engineering [1–4], carbon monoxide removal from reformed gas for fuel cells [5] and membrane reactors to improve and control chemical reactions [6]. At present, Pd-based alloys exhibit the best combination of high hydrogen selectivity and permeability, high thermal stability and especially high ductility which is the key property for the fabrication of thin membranes [7]. However, Pd is expensive and scarce. There is a high demand for low-cost metal membranes with minimal content or complete absence of Pd, but exhibiting similarly high hydrogen permeability, ductility and resistance to the hydrogen embrittlement. Both crystalline alloys with body-centered-cubic (BCC) structure and amorphous alloys are promising alternatives to Pd-based membranes for hydrogen separation because of their lower cost and comparable hydrogen permeability. During the last decade, a series of amorphous alloys derived from Ni–Nb–Zr system [8–14] were developed. These materials exhibit high hydrogen permeability with mechanical stability
n
Corresponding author. Tel./fax: þ 86 451 86418754. E-mail address: [email protected] (X. Li). 1 These authors contributed equally to this work.
http://dx.doi.org/10.1016/j.memsci.2015.03.002 0376-7388/& 2015 Elsevier B.V. All rights reserved.
in hydrogen atmosphere in appropriate temperature regimes. However, their applicability at high temperatures is limited by the thermal stability of the amorphous alloys. Crystallization of the amorphous membranes reduces both hydrogen permeability and strength significantly [10]. In contrast, crystalline BCC solid-solution alloys comprising Group V metals such as V, Nb and Ta exhibit much higher hydrogen permeability and thermal stability [15–19]. The high hydrogen solubility provides them a high driving force for hydrogen permeation, which however aggravates the problem of hydrogen embrittlement. To address this problem, a new type of hydrogen permeable crystalline alloys has been developed, such as Nb–TiNi [20–24], Nb–TiCo [25,26], Ta–TiNi [27,28], V–Ti–Ni [29–31], etc. These alloys contain two phases, particularly a primary BCC solid-solution phase that provides the major contribution to hydrogen permeation, and in addition a secondary eutectic phase that suppresses the hydrogen embrittlement. Consequently, such alloys are promising to achieve a sound combination of high hydrogen permeability, excellent resistance against hydrogen embrittlement and high ductility. During the past several years, hydrogen permeable Nb–TiNi alloys consisting of primary BCC-(Nb, Ti) and eutectic {BCC-(Nb, Ti)þB2TiNi} phases have attracted intensive attention. This alloy system exhibits high ductility for fabrication of thin membranes and high tolerance to hydrogen embrittlement, but the hydrogen permeability is not quite satisfactory. It is found that appropriate substitution of Ti by Zr can induce a much higher hydrogen permeability (because of the higher hydrogen affinity of Zr than that of Ti) and comparable
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other properties [32,33]. Nevertheless, excessive addition of Zr leads to low ductility and serious hydrogen embrittlement due to the microstructural change and the inappropriately high hydrogen solubility [34]. Therefore, much of current research focuses on compositional modifications in these alloys to tailor the microstructures and hydrogen solubility and diffusivity for improved performance. The element Hf, in the same IVA group as Zr, also possesses a higher hydrogen affinity than that of Ti but lower than that of Zr [7,12]. It is most likely that appropriate substitution of Ti by Hf in Nb–TiNi alloys results in new hydrogen permeable alloys with high performance. In this work, microstructure and its influence on ductility and hydrogen permeability of as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys are investigated. The former alloys are based on eutectic Nb20Ti40Ni40, so that the microstructural changes induced by the addition of Hf in this system may be observed clearly. The latter alloys are based on the hypereutectic Nb40Ti30Ni30 with the aim to develop new hydrogen permeable alloys with high performance by appropriate substitution of Ti by Hf. A maximum coldrolling reduction ratio is used to evaluate ductility by measuring thickness changes in the samples. The solubility and diffusivity of hydrogen in the two series of alloys are characterized for revealing the detailed response of hydrogen permeability to the Hf content.
Table 1 Summary of constituting phases and their respective chemical compositions (EDX results) in the as-cast Nb20Ti40 xHfxNi40 (x ¼0…40) and Nb40Ti30 xHfxNi30 (x¼ 0… 30) alloys. Samples
Nb20Ti40Ni40 Nb20Ti35Hf5Ni40 Nb20Ti30Hf10Ni40 Nb20Ti25Hf15Ni40 Nb20Ti20Hf20Ni40 Nb20Ti15Hf25Ni40 Nb20Ti10Hf30Ni40 Nb20Ti5Hf35Ni40 Nb20Hf40Ni40 Nb40Ti30Ni30 Nb40Ti25Hf5Ni30 Nb40Ti20Hf10Ni30 Nb40Ti15Hf15Ni30 Nb40Ti10Hf20Ni30 Nb40Ti5Hf25Ni30 Nb40Hf30Ni30
Chemical composition of the constituting phases (mol%) Primary BCC-(Nb, Ti, Hf) phase
{BCC-(Nb, Ti, Hf) þ B2-(Ti, Hf) Ni} eutectic
Nb
Ti
Hf
Ni
Nb
– – 77.18 76.94 75.83 76.98 77.82 77.01 77.58 82.92 80.94 81.65 82.38 81.59 82.25 82.81
– – 10.19 8.45 8.22 6.79 3.09 2.42 0 12.65 11.72 9.59 5.83 4.46 2.98 0
– – 5.75 7.12 8.41 9.86 11.07 13.45 14.99 0 2.93 5.11 7.86 9.57 11.28 13.51
– 20.03 – 19.93 6.88 19.81 7.49 18.45 7.54 17.12 6.37 16.28 8.02 15.31 7.12 14.75 7.43 14.67 4.43 19.83 4.41 19.35 3.65 19.11 3.93 18.14 4.38 17.62 3.49 18.03 3.68 17.88
Ti
Hf
Ni
39.82 35.37 31.02 24.89 19.43 14.86 9.65 5.13 0 38.51 31.69 22.34 18.12 14.78 7.42 0
0 4.87 13.29 19.94 25.71 31.19 38.82 44.13 50.84 0 9.28 17.21 23.15 28.87 33.26 39.35
40.15 39.83 35.88 36.72 37.74 37.67 36.22 35.99 34.49 41.66 39.68 41.34 40.59 38.73 41.29 42.77
2. Experimental Ingots of about 20 g of Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼0…30) (all concentrations in at% if not specified differently) alloys were prepared by arc melting in a purified argon atmosphere using pure Nb, Ti, Hf and Ni (99.9 mass% purity for all). The crystal structures of the alloys were identified by powder X-ray diffractometry (XRD). Microstructural observations were carried out using a Scanning Electron Microscope (SEM) equipped with energy dispersive X-ray (EDX) spectrometry. The volume fraction of the constituent phases was analyzed using the public domain NIH Image program. Samples of 7 3 3 mm3 in dimension were cut from Nb40T i30 xHfxNi30 (x¼0…30) alloy ingots using a spark erosion wirecutting machine for the cold-rolling experiment at room temperature. Disk samples of 12 mm in diameter and 0.6 mm in thickness were cut from two series of alloys for the measurement of hydrogen permeation. Both sides of the disks were ground, polished and then coated with Pd of 190 nm in thickness by a radio frequency (RF) sputtering machine to prevent oxidation during permeation experiments and to act as catalyzer for hydrogen uptake into the material. The disks were sealed by copper gaskets. Both sides of the disks were evacuated using a diffusion pump to below 3 10 3 Pa, then heated to 673 K and kept for 20 min. Hydrogen gas (99.99999 mass% purity) of 0.2–0.4 MPa was introduced at the upstream side (Pu) and of 0.1 MPa at the downstream side (Pd). The hydrogen flux, J, passing through the disks was measured by a hydrogen flow meter. The measurements were repeated at 623 K, 573 K and 523 K. The pressure-composition-temperature (PCT) curves for some representative alloys were measured using a Sieverts-type apparatus at 673 K and in the pressure range from 0.01 to 0.9 MPa to investigate the hydrogen solubility. The experimental procedure was identical with those in Refs. [35–37]. The amount of absorbed hydrogen was calculated from the pressure drop in a constant inner volume chamber.
3. Results and discussion 3.1. Microstructures The constituting phases and their respective chemical compositions in the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) and Nb40Ti30 x
HfxNi30 (x¼0…30) alloys as determined by XRD, SEM and EDS are tabulated in Table 1. The XRD patterns of the two series of alloys are shown in Figs. 1 and 2. Nb–(Ti, Hf)Ni alloys with an initial Hf content below that of Ti consist of BCC-(Nb, Ti, Hf) solid solution and B2-(Ti, Hf)Ni, and the substitution of Ti by Hf does not lead to a change of the constituting phases. However, if the Hf content is close to or higher than that of Ti, some intermetallic phases such as HfNi, Hf3Ni7, Hf8Ni21, Hf2Ni, HfNi3, Hf2Ni7, etc. appear additionally. With increasing Hf content, the diffraction peak intensity of these intermetallic phases increases, especially that of Orthorhombic-HfNi phase. For the Nb20Ti40 xHfxNi40 (x¼ 25…40) alloys, the intensity of the main diffraction peak of the HfNi is comparable to or even higher than that of BCC-Nb (see Fig. 1). This indicates that a large amount of HfNi phase is present in these alloys. In comparison, for the Nb40 Ti30 xHfxNi30 (x¼15…30) alloys, the diffraction peak intensity of the HfNi phase is much lower than that of BCC-Nb (see Fig. 2). In addition, the main diffraction peak of B2-(Ti, Hf)Ni corresponding to 431 decreases in intensity with the gradual substitution of Ti by Hf in both series of alloys. This peak vanishes finally when the Hf content in alloys is close to or higher than that of Ti, and simultaneously the diffraction peak of Orthorhombic-(Hf, Ti)Ni corresponding to 271 becomes obvious and grows in intensity with further increasing the Hf content. It should be noted that B2-(Ti, Hf)Ni results from the replacement of Ti by Hf in B2-TiNi lattice and thus retains the ductile characteristics. However, the Orthorhombic-(Hf, Ti)Ni results from the replacement of Hf by Ti in Orthorhombic-HfNi lattice, which is intrinsically brittle [32]. SEM micrographs of the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) alloys are shown in Fig. 3. A fully eutectic microstructure, {BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni}, is observed for Nb20Ti40Ni40 (x¼ 0) and Nb20Ti35Hf5Ni40 (x¼ 5) alloy, Fig. 3(a) and (b). With further increase of Hf content (x), primary BCC-(Nb, Ti, Hf) dendrites embedded in eutectic are found in Nb20Ti30Hf10Ni40 (x¼10) and Nb20Ti25Hf15Ni40 (x¼15) alloys, Fig. 3(c) and (d). For alloys with xZ20, a similar hypereutectic microstructure appears, Fig. 3(e)–(i); but the B2-(Ti, Hf) Ni in eutectic turns to Orthorhombic-(Hf, Ti)Ni (see Fig. 1). The eutectic morphology turns to be rod-like, compared to the lamellar one in the alloys with xo20, see the higher magnification inset image in Fig. 5. Fig. 4 shows SEM micrographs for the as-cast Nb40Ti30 xHfxNi30 (x¼0…30) alloys. Primary BCC-(Nb, Ti, Hf) dendrites embedded in
X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
49
Fig. 1. XRD patterns of the as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) alloys. Fig. 2. XRD patterns of the as-cast Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys.
eutectic [BCC-(Nb, Ti, Hf)þB2-(Ti, Hf)Ni/Orthorhombic-(Hf, Ti)Ni] are mainly found in these alloys. When the Hf content is close to or higher than that of Ti, the B2-(Ti, Hf)Ni in eutectic turns to Orthorhombic(Hf, Ti)Ni (see Fig. 2), and the lamellar eutectic changes to rod-like one (see insets in Fig. 5) as that in Nb20Ti40 xHfxNi40 alloys. For the above two series of alloys, the substitution of Ti by Hf induces higher volume fractions of the primary phase and less eutectic, as shown in Fig. 5. This can be related to the fact that the eutectic Nb20Ti40Ni40 is closer to the Nb corner of the ternary phase diagram than that of eutectic Nb16Hf42Ni42 [38]. Correspondingly, the effect of the above substitution on the solidification behavior is equivalent to increasing the initial Nb content of ternary Nb–TiNi alloys. The solid-solution Hf in primary BCC-Nb dendrites increases with more Hf addition, as illustrated in Table 1. When the initial Hf content in alloys is close to or higher than that of Ti, the ductile B2-(Ti, Hf)Ni in eutectic turns to brittle Orthorhombic-(Hf, Ti)Ni and some intermetallic phases (HfNi, HfNi3, etc.) appear additionally at the boundaries of eutectic grains (see the insets in Fig. 5). This corresponds to the XRD results in Figs. 1 and 2. It should be noted that these intermetallic phases are inherently brittle and cannot be eliminated by annealing [39]. Further evaluation of ductility is needed for alloys containing these phases. 3.2. Ductility Considering the workability of hydrogen permeable alloy membranes, e.g. mechanically sealed with copper gaskets, these alloys must have at least some ductility. Most simply, a material that can be broken easily by hammering is qualitatively defined as brittle [32–34], while non-broken materials are defined as ductile. Here, the maximum cold-rolling reduction ratio is used to more reproducibly evaluate the ductility of alloys. A qualitative ductility
index for Nb20Ti40 xHfxNi40 (x ¼0…40) and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys is tabulated in Tables 2 and 3, respectively. Overall, alloys with Hf content lower than that of Ti are found to be ductile, because their constituting phases of BCC-(Nb, Ti, Hf) and B2-(Ti, Hf)Ni are ductile. Alloys with Hf content close to or higher than that of Ti are brittle due to the existence of brittle intermetallic phases. Fig. 6 shows the maximum cold-rolling reduction ratio of the as-cast Nb40Ti30 xHfxNi30 alloys as a function of the Hf content (x). The cold-rolling reduction ratio (rt) is defined as rt ¼(t0 t)/t0, where t0 and t are the original and the final thicknesses of the sample, respectively [32]. It can be seen that 60% or higher reduction ratio (rt) was obtained for Hf contents of 10 at% or less. In contrast, if the Hf content is 15% or higher, rt is reduced significantly. This result appropriately corresponds to the above XRD results and microstructural observations. Therefore, the Hf content in alloys should be limited to that of Ti for good ductility in the as-cast state. 3.3. Hydrogen permeability 0:5 Fig. 7 shows the relation between J L and ΔP0.5 ( ¼ P 0:5 u Pd ) for the as-cast Nb20Ti30Hf10Ni40 and Nb40Ti20Hf10Ni30 alloys from 523 to 673 K, where J is the measured hydrogen flux, L is the thickness of the membrane and ΔP is the difference of the hydrogen pressure at the two sides of the membrane. In the plot, the experimental data are quite precisely positioned on straight lines passing through the origin, and the correlation coefficient R2, obtained by linear regression, is higher than 0.99. A similar relation between J L and ΔP0.5 appears in other alloys. However, this does not indicate that these alloys follow Sieverts' law at the
50
X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
Fig. 3. SEM micrographs of as-cast Nb20Ti40 xHfxNi40 alloys with increasing x (x¼ 0…40). (a) Nb20Ti40Ni40, (b) Nb20Ti35Hf5Ni40, (c) Nb20Ti30Hf10Ni40, (d) Nb20Ti25Hf15Ni40, (e) Nb20Ti20Hf20Ni40, (f) Nb20Ti15Hf25Ni40, (g) Nb20Ti10Hf30Ni40, (h) Nb20Ti5Hf35Ni40, and (i) Nb20Hf40Ni40.
present pressures of hydrogen permeation. In fact, it has been found that hydrogen dissolution in niobium and its alloys does not follow Sieverts' law in the high hydrogen content region [40–43]. Therefore, further information about the pressure dependence of hydrogen solubility in the present alloys is needed [22,43]. Figs. 8 and 9 show the PCT curves for the as-cast Nb20 Ti40 xHfxNi40 (x ¼0, 10, 20, 30, 40) alloys and Nb40Ti30 xHfxNi30 (x ¼0, 10, 20, 30) alloys, respectively, measured at 673 K. They are plotted in the form of Sieverts' plot, i.e., hydrogen content C vs. the square root of the hydrogen pressure P0.5. This relation is crucial to define the hydrogen permeability for the hydrogen permeable alloys [41]. It can be seen that Sieverts' plot for each alloy shows straight line between 0.1 and 0.4 MPa, but does not pass through the origin. This indicates that Sieverts' law does not hold for these alloys under the present conditions. At pressures lower than 0.1 MPa, C vs. P0.5 is close to a straight line, but its slope is larger than that at 0.1–0.4 MPa. At pressures higher than 0.4 MPa, C vs. P0.5 deviates from a linear relation, and its slope tends to reduce with increasing pressures, which is also lower than that at low pressures. Actually, this slope corresponds to the hydrogen solubility coefficient (K). It is obvious that the values of K gradually reduce from low pressures (o0.1 MPa) to intermediate pressures
(0.1–0.4 MPa) and then high pressures ( 40.4 MPa). This hydrogen absorption behavior is a reflection of changes to the progressive filling of interstitial sites with increasing hydrogen content. At low pressures, the sites with the lowest energy are preferentially occupied by hydrogen atoms. A broad distribution of site energies for these alloys results in a significantly enhanced occupation with slightly increasing pressures. Correspondingly, C vs. P0.5 exhibits a relatively steep rise of line, which corresponds to a large K value. With further increasing pressures, the filling of the most favorable sites will be gradually finished and occupation of less favored sites occurs. This induces a lower K value, and the pressure required to achieve a given hydrogen content increases. This situation is enhanced particularly at high pressures, because most of hydrogen atoms have to stay among less favored sites. Therefore, the curve of C vs. P0.5 tends to slope gently at high pressures, which corresponds to the lowest K value. The hydrogen content in the pressure range of 0.1–0.4 MPa for each alloy, C, can be expressed as C ¼ K U P 0:5 þ α
ð1Þ
where K is the hydrogen solubility coefficient and α is a constant. Tables 2 and 3 show the values of K and α for each alloy, determined
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51
Fig. 4. SEM micrographs of as-cast Nb40Ti30 xHfxNi30 alloys with increasing x (x ¼0…30). (a) Nb40Ti30Ni30, (b) Nb40Ti25Hf5Ni30, (c) Nb40Ti20Hf10Ni30, (d) Nb40Ti15Hf15Ni30, (e) Nb40Ti5Hf25Ni30, and (f) Nb40Hf30Ni30.
by linear regression. It can be seen that the values of K for two series of alloys increase with the gradual substitution of Ti by Hf. The steady-state hydrogen flux, J, through a metal membrane is defined by Fick's first law ∂C J ¼ D ð2Þ ∂L where D is the hydrogen diffusion coefficient and ∂C/∂L is the gradient of the hydrogen content across the membrane. If C in Eq. (2) is substituted by C from Eq. (1), the hydrogen flux, J, can be
linearized as follows: J¼
D KðΔP 1=2 Þ ΦðΔP 1=2 Þ ¼ L L
ð3Þ
This implies that we can still define the hydrogen permeability (Φ) for the present alloys in a limited hydrogen pressure range, even if Sieverts' law does not hold. The slope of each straight line in the plot of J L vs. ΔP0.5 in Fig. 7 corresponds to the hydrogen permeability (Φ). This method is commonly used to determine Φ [22,43].
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Fig. 10 shows the temperature dependence of Φ in the form of an Arrhenius plot for the as-cast Nb20Ti40 xHfxNi40 (x¼0…40) alloys and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys. The Φ values of pure Pd [44] are also included in the figure for comparison. The variations of Φ appear to be straight lines in the Arrhenius plot in all cases, and the Φ values of each alloy evidently increase with increasing temperature. The effect of temperature on the Φ values for present alloys is larger than that of pure Pd, i.e., the activation energy for hydrogen permeation (E) through these alloys is larger than that of pure Pd. The values of E for Nb–(Ti, Hf)Ni alloys range from 24.5 to 33.2 kJ mol 1 (see Tables 1 and 2), while that for pure Pd is 12.21 kJ mol 1. It can also be seen that the Φ values of both series of alloys increase with the gradual substitution of Ti by Hf. Most alloys in the Nb20Ti40 xHfxNi40 (x ¼0…40) system exhibit a lower Φ value than that of pure Pd (see Fig. 10(a)). Nb20Ti15Hf25Ni40 and Nb20Ti10Hf30Ni40 are prone to embrittlement and broke during hydrogen permeation at 523 K, so that the Φ value could not be measured. The problem of hydrogen embrittlement is aggravated for Nb20Ti5Hf35Ni40 and Nb20Hf40Ni40 at 523 K and 573 K. In contrast, most alloys in the Nb40Ti30 xHfxNi30 (x ¼0…30) system exhibit much higher Φ values than pure Pd (see Fig. 10(b)), especially at high temperatures. No cracks were found after the hydrogen permeability measurements at all temperatures for these alloys, indicating that they are less susceptible to hydrogen embrittlement during hydrogen permeation. According to the above defined K and Φ values, the hydrogen diffusion coefficient (D) in each alloy can be evaluated on the basis of Φ ¼D K, which is tabulated in Tables 2 and 3. In this way, a relationship of permeability, solubility and diffusivity of hydrogen
Fig. 5. Relationship of Hf content and volume fraction of primary BCC-(Nb, Ti, Hf) in the as-cast Nb20Ti40 xHfxNi40 (x¼ 0…40) and Nb40Ti30 xHfxNi30 (x¼ 0…30) alloys, respectively. The inset images show the eutectic morphology and formation of additional new intermetallic phases at the boundaries of eutectic grains (labeled by arrows) in three representative alloys.
can be built for the as-cast Nb20Ti40 xHfxNi40 (x ¼0…40) alloys and Nb40Ti30 xHfxNi30 (x ¼0…30) alloys. The hydrogen permeability of both alloys increasing with Hf content is found to be induced by the simultaneous increase of the hydrogen solubility and diffusivity (see Tables 2 and 3). The hydrogen solubility in metals is governed mainly by the size and the electronic structure of the interstitial sites associated with the change of the interstitial binding energy [45,46]. For the BCC-Nb, octahedral and tetrahedral interstices exist. The occupation of tetrahedral sites by hydrogen atoms in Group V BCC lattices is more favored than the octahedral ones [47–50]. Due to the larger atomic radius of Hf (159 pm) as compared to that of Ti (147 pm), the substitution of Ti by Hf in BCC-Nb will increase the size of the tetrahedral interstitial sites. This may potentially increase the tolerance to the lattice expansion for accommodating more hydrogen atoms and thus the hydrogen solubility. From the view of the electronic effect, the interstitial binding energy is related to the hydrogen charge where the stability of the site is reduced when hydrogen gains less charge for the group V materials [47]. In the pure niobium lattice, hydrogen gains charge from each of the nearest neighbor niobium atoms, because the electronegativity of hydrogen (2.2) is higher than that of niobium (1.6). When the neighboring niobium atoms are replaced by elements with a similar or higher electronegativity than that of hydrogen, less charge is shared between the lattice and hydrogen, thereby decreasing the binding energy. In contrast, for the present Nb–(Ti, Hf)Ni alloys, Ti and Hf in solid solution in BCC-Nb have a lower electronegativity than that of hydrogen. The electronegativity of Hf (1.3) is lower than that of Ti (1.54). It is expected that the substitution of Ti by Hf may increase the charge shared between the lattice and hydrogen and thus increase the binding energy and the hydrogen solubility. It also should be noted that alloys with Hf content close to or beyond that of Ti contain a certain amount of HfNi, especially for the Nb20Ti40 xHfxNi40 (x ¼25…40) alloys (as discussed in Section 3.1). This phase is known to absorb hydrogen for a long time [51–53], and its hydrogenation contributes partially to the high hydrogen solubility for these alloys. This provides a positive contribution to hydrogen permeability, but also induces a low tolerance to hydrogen embrittlement. In comparison, Nb40Ti30 x HfxNi30 (x ¼15…30) alloys contain a less HfNi phase, and they are less susceptible to hydrogen embrittlement during hydrogen permeation. Hydrogen diffusion in metals is described by classical hopping between interstitial sites, which is related to the hydrogen content in metals [54,55]. At dilute contents, hydrogen atoms are essentially found in the sites with the strongest binding energies, so that longrange diffusion is limited by the large barriers that exist to moving hydrogen atoms away from these sites. At higher contents, however, the most favorable sites are occupied by some population of hydrogen atoms, which allows other hydrogen atoms to stay and hop among less favored sites. The latter population moves far more rapidly than atoms trapped in the most favored sites, so they dominate the overall diffusion coefficient [54]. Therefore, the diffusivity of hydrogen increases markedly as the hydrogen content increases. This has been
Table 2 Hydrogen permeability (Ф), hydrogen solubility coefficient (K), hydrogen diffusion coefficient (D), fitting parameter (α) at 673 K and activation energy for hydrogen permeation (E) along with the ductility for the as-cast Nb20Ti40 xHfxNi40 alloys. Samples
Ф [10 8 mol H2 m 1 s 1 Pa 0.5]
K [mol H2 m 3 Pa 0.5]
D [10 10 m2 s 1]
α [mol H2 m 3]
E [kJ mol 1]
Nb20Ti40Ni40 Nb20Ti30Hf10Ni40 Nb20Ti20Hf20Ni40 Nb20Ti10Hf30Ni40 Nb20Hf40Ni40
0.64 0.84 1.19 1.63 2.01
21.36 23.56 26.7 27.88 31.91
3.0 3.57 4.64 5.85 6.30
2339 3200 3472 4291 5043
26.8 28.7 32.2 33.1 31.7
Ductile Ductile Brittle Brittle Brittle
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Table 3 Hydrogen permeability (Ф), hydrogen solubility coefficient (K), hydrogen diffusion coefficient (D), fitting parameter (α) at 673 K and activation energy for hydrogen permeation (E) along with the ductility for the as-cast Nb40Ti30 xHfxNi30 alloys. Samples
Ф [10 8 mol H2 m 1 s 1 Pa 0.5]
K [mol H2 m 3 Pa 0.5]
D [10 10 m2 s 1]
α [mol H2 m 3]
E [kJ mol 1]
Nb40Ti30Ni30 Nb40Ti20Hf10Ni30 Nb40Ti10Hf20Ni30 Nb40Hf30Ni30
1.76 2.96 3.62 4.33
24.29 25.57 23.37 29.09
7.25 11.58 13.23 14.89
5450 6350 7790 9860
24.5 32.1 33.2 31.8
Ductile Ductile Brittle Brittle
Fig. 6. Maximum cold-rolling reduction ratio of the as-cast Nb40Ti30 xHfxNi30 alloys as a function of the Hf content.
testified by a variety of experiments and first-principle calculations [56–58]. For the present Ni–(Ti, Hf)Ni alloys, the substitution of Ti by Hf induces an increase of hydrogen contents (see Figs. 8 and 9), which potentially leads to the increase of the hydrogen diffusivity. 3.4. Resistance against hydrogen embrittlement (HE) For hydrogen permeable metal membranes, the large resistance against HE is crucial to the durability with respect to long-term hydrogen permeation. The short-term tests in Fig. 10 indicate that ascast Nb20Ti40 xHfxNi40 (xr20) and Nb40Ti30 xHfxNi30 (x¼0…30) alloys exhibit an appropriate tolerance to HE. However, such information is limited as it does not consider long-term durability which is of greater interest. Fig. 11 shows the variations of hydrogen permeability with time in 72 h for four representative alloys at 673 K. It can be seen that Nb20Ti25Hf15Ni40 and Nb40Ti20Hf10Ni30 exhibit excellent durability in hydrogen permeation and their Φ values are almost stable after initial transient. The final membranes remain totally free from cracks or pores after a test of 72 h, which were confirmed by a room temperature permeation test with Ar. This indicates that both alloys show large resistance against HE. In contrast, Nb40Ti15Hf15Ni30 and Nb20Ti20Hf20Ni40 are subjected to HE failure after hydrogen permeation of 15 h and 23 h, respectively. This is not detectable in a short-time test of hydrogen permeation, because these two alloys also exhibit stable permeability during the early 10 or 20 h. For the present hydrogen permeable alloys, the eutectic in as-cast microstructures plays an important role in resisting HE. Once hydrogen permeates, primary BCC-Nb phases dilate with the extent depending on the solubility of hydrogen, and the surrounding eutectic may tolerate this dilation due to its high deformability [59]. A high fraction of eutectic is thus beneficial to the embrittlement resistance. However, it should be noted that BCC-Nb in eutectic also act as pathways for hydrogen permeation, and their dilation after
Fig. 7. Relation between J L and ΔP0.5 for the as-cast Nb20Ti30Hf10Ni40 (a) and Nb40Ti20Hf10Ni30 (b) in the temperature range of 523–673 K.
hydrogen absorption exerts force to the surrounding intermetallic phases. For Nb20Ti25Hf15Ni40 and Nb40Ti20Hf10Ni30, ductile B2-(Ti, Hf) Ni provides a high tolerance to the dilation of BCC-Nb in the eutectic. Therefore, the high fractions of eutectic ( 68 vol% or higher) in ascast microstructures of these two alloys and the ductile B2-type intermetallic phase in eutectic ensure their large resistance against HE. In comparison, there are also high fractions of eutectic in Nb20Ti20Hf20Ni40 and Nb40Ti15Hf15Ni30 which provide an appropriate resistance against HE. However, brittle Orthorhombic-(Hf, Ti)Ni in their eutectic provide a low tolerance to the dilation of surrounding BCC-Nb. Especially, a high hydrogen solubility in Orthorhombic-(Hf, Ti)Ni is expected as its parental HfNi, and its hydrogenation also induces a large lattice expansion. These result in a deterioration of embrittlement resistance. The positive contribution from the eutectic
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Fig. 8. Pressure-composition-temperature (PCT) curves for the as-cast Nb20Ti40 xHfxNi40 (x ¼0, 10, 20, 30, 40) alloys at 673 K, and the hydrogen content C (mol H2 m 3) is plotted as a function of the square root of the hydrogen pressure P0.5 for each alloy, i.e., in form of a Sieverts' plot.
Fig. 10. Temperature dependence of hydrogen permeability (Ф) in the form of an Arrhenius plot for the as-cast alloys of (a) Nb20Ti40 xHfxNi40 (x¼ 0…40) and (b) Nb40Ti30 xHfxNi30 (x ¼ 0…30).
Fig. 9. Pressure-composition-temperature (PCT) curves for the as-cast Nb40Ti30 xHfxNi30 (x¼ 0, 10, 20, 30) alloys at 673 K, and the hydrogen content C (mol H2 m 3) is plotted as a function of the square root of the hydrogen pressure P0.5 for each alloy, i.e., in form of a Sieverts' plot.
and the negative one from the brittle Orthorhombic-(Hf, Ti)Ni appropriately explain the HE failure after some time of hydrogen permeation for these two alloys. This can be further extended to deduce that Nb20Ti40 xHfxNi40 (xo20) and Nb40Ti30 xHfxNi30 (xo15) may exhibit a large resistance against HE because of their high fractions of eutectic and ductile B2-(Ti, Hf)Ni in their eutectic structure. In a general view taking into account the microstructure, ductility, hydrogen permeability and the resistance against hydrogen embrittlement of the above Nb-(Ti, Hf)Ni alloys, Nb40Ti20Hf10Ni30 is a promising candidate to fabricate high-performance membranes for hydrogen separation. This alloy consists of the appropriate microstructure for hydrogen permeation (without any intermetallic phases besides the eutectic), exhibits a high cold-rolling reduction of 63%,
Fig. 11. Variations of hydrogen permeability with time in 72 h for the as-cast Nb20Ti25Hf15Ni40, Nb20Ti20Hf20Ni40, Nb40Ti20Hf10Ni30 and Nb40Ti15Hf15Ni30 at 673 K.
a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (which is 1.85 times that of pure Pd) and excellent resistance against hydrogen embrittlement. Table 4 shows the hydrogen permeability for some representative alloys, and their volume fractions of primary phase and eutectic and the tolerance to HE are also
X. Li et al. / Journal of Membrane Science 484 (2015) 47–56
Table 4 Hydrogen permeability along with the volume fractions of primary phase and eutectic and the tolerance to HE for some representative alloys, obtained from the research group of Aoki using the same characterizing methods. Alloys
Nb19Ti40Ni41 [22] Nb40Ti30Ni30 [21] Nb56Ti23Ni21 [20] Nb30Ti35Co35 [25] Nb60Ti21Co19 [26] V41Ti30Ni29 [27] Nb40Zr30Ni30 [60] Nb40Hf30Ni30 [39] Nb40Ti18Zr12Ni30 [33] Nb40Ti18Zr12Ni25Co5 [33] Nb40Ti15Hf10Ni30
Volume fractions (vol%) Primary BCC-Nb phase
Eutectic
Hydrogen permeability (Ф) at 673 K (10 8 mol H2 m 1 s 1 Pa 0.5)
0 29 62 0 62 40 40 40 45
100 71 38 100 38 60 60 60 55
0.64 1.8 3.47 2.64 3.99 1.1 4.64 4.3 2.85
High High Low High Low High Low Low High
42
58
3.82
High
32
68
2.96
High
Tolerance to HE
included. Although containing a low fraction of primary BCC-Nb, Nb40Ti20Hf10Ni30 exhibits moderate permeability among these alloys. Simultaneously, the high fraction of eutectic ensures its high tolerance to HE. Therefore, Nb40Ti20Hf10Ni30 clearly reaches an excellent balance between permeability and embrittlement resistance. In comparison, Nb56Ti23Ni21 and Nb60Ti21Co19 exhibit higher permeability than Nb40Ti20Hf10Ni30, but their low fractions of eutectic induce a low tolerance to HE. The high permeability of Nb40Zr30Ni30 and Nb40Hf30Ni30 is originated from the high hydrogen solubility in BCC-Nb and also Orthorhombic-ZrNi/HfNi in their microstructures. The hydrogenation of Orthorhombic-ZrNi/HfNi results in severe HE. Nb30Ti35Co35, Nb40Ti18Zr12Ni30 and Nb40Ti18Zr12Ni25Co5 also reach an appropriate balance between permeability and embrittlement resistance. In contrast, Nb40Ti20Hf10Ni30 exhibits higher permeability than Nb30Ti35Co35 and Nb40Ti18Zr12Ni30, but lower one than Nb40Ti18Zr12Ni25Co5. It is expected that partial substitution of Ni by Co for Nb40Ti20Hf10Ni30 induces a higher permeability.
4. Conclusions Nb-based alloys are an emerging alternative to palladium alloys for hydrogen separation, while the key challenge remains to balance the ductility, hydrogen permeability and resistance against HE. This work has shown that substitution of Ti by Hf in Nb–(Ti, Hf)Ni alloys increases hydrogen solubility and diffusivity and thus hydrogen permeability, but greater hydrogen solubility brings the penalty of poorer resistance against HE. When the Hf content in an alloy is close to or higher than that of Ti, some undesired brittle phases appear in microstructures which distinctly reduce ductility. Correspondingly, the Hf content in practically applicable alloys should be lower than that of Ti for high ductility and high tolerance to HE. Nb40Ti20Hf10Ni30 reaches an excellent balance among ductility, hydrogen permeability and embrittlement resistance. This alloy exhibits a high cold-rolling reduction of 63%, a high hydrogen permeability of 2.96 10 8 mol H2 m 1 s 1 Pa 0.5 at 673 K (1.85 times that of pure Pd) and excellent durability for long-term permeation, which is promising to fabricate membranes for hydrogen separation.
Acknowledgments This project was supported by the National Natural Science Foundation of China (Grants nos. 51274077 and 51271068) and
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the National Basic Research Program of China 973 (Grant no. 2011CB610406).
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