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Hydrogen permeation through amorphous-Zr36−xHfxNi64-alloy membranes
Journal of Membrane Science 211 (2003) 149–156
Hydrogen permeation through amorphous-Zr36−x Hfx Ni64 -alloy membranes S. Hara a,∗ , N. Hatakeyama a , N. Itoh a , H.-M. Kimura b , A. Inoue b a
National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba Central 5, Tsukuba 305-8565, Japan b Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 4 February 2002; accepted 9 September 2002
Abstract Amorphous-alloy membranes of a series of Zr36−x Hfx Ni64 (0 ≤ x ≤ 36) were successfully prepared by the rapid-quenching method using a single roller; their amorphous structure and thermal behavior were characterized by XRD and DSC. All of the membranes, coated by palladium to provide activity to hydrogen dissociation and recombination, were sufficiently robust in a hydrogen atmosphere and showed stable permeability only to hydrogen at least in the range of 473–573 K. On the other hand, over 573 K, the permeation rate slowly decreased over time. Permeability was found to decrease with Hf substitute amount for Zr mainly due to increase in activation energy for permeation. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Hydrogen permeation; Amorphous-alloys; Zr–Hf–Ni alloys; Hydrogen diffusion; Hydrogen embrittlement
1. Introduction Hydrogen-permeable metal membranes generally have high selectivity, permeability, and thermal stability, so that they are expected to be used for various industrial applications, including ultra-pure hydrogen production for semiconductor manufacture [1], carbon-monoxide removal from reformed gas for fuel cells [2], and membrane reactors to improve and control chemical reactions [3]. However, metal membranes employed now in industry are strictly palladium or palladium-alloy membranes, which are extremely expensive. It is, therefore, desirable to develop cheaper membrane materials substitutable for them. ∗ Corresponding author. Tel.: +81-298-61-9336; fax: +81-298-61-4674. E-mail address: [email protected] (S. Hara).
Metals used for gaseous-hydrogen separation have to meet at least three requirements: high permeability, activity to dissociate hydrogen molecules, and resistance to hydrogen embrittlement [4,5]. Of these requirements, finding metals with high hydrogen permeability is not difficult because the permeation rate can be estimated from hydrogen diffusivity and solubility in the metal if hydrogen permeation is controlled by atomic-hydrogen diffusion in the membrane. For example, hydrogen is known to diffuse easily between metal atoms in a body-centered cubic (bcc) lattice; so, high permeability is expected for bcc metals with high solubility such as vanadium, niobium, and tantalum [6]. However, because hydrogen in metals generally weakens the metals (hydrogen embrittlement), most such metals which are expected to have high permeability absorb a large amount of hydrogen in a hydrogen atmosphere; they then break due to stress from differential pressure across the membrane
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and/or internal stress caused by absorbed hydrogen atoms. Therefore, the developmental obstacle for metal membrane materials is not improved permeability but resolution of the two conflicting requirements of high permeability and resistance to hydrogen embrittlement. Nevertheless, much research has avoided challenging the hydrogen embrittlement issue. For instance, electrochemical methods have been employed to find hydrogen permeability through metal membranes [7]. In this case, the driving force for hydrogen permeation is differential electrochemical potential for hydrogen ions in solutions on both sides of the membrane, so that no mechanical stress need be applied to the membrane during measurement. Therefore, the same membranes cannot be said to also be useful for gaseous-hydrogen separation without breaking when some stress is applied. Many experiments have been reported on metal membranes prepared on supports to improve mechanical strength. Such a composite structure is one effective way to increase permeation flux; indeed, it is one which may popularize use of metal membranes. However, making composite membranes is rather difficult, as reported by many authors for palladium composite membranes [8–10]. Sakaguchi et al. [11] deposited an amorphous LaNi5 alloy layer on a support by the sputtering method, but they never attained complete selectivity to hydrogen due to metal layer defects. A composite membrane composed of amorphous FeTi and a palladium support by Amano et al. [12] showed complete selectivity; but, as a practical matter, such a membrane is not substitutable for palladium membranes because dense palladium sheet was employed as a support. As described above, not a few researchers have reported on hydrogen permeability through metal membranes other than palladium-alloys, but many of them have made various devices to avoid hydrogen embrittlement which are not always useful for industrial application. Amorphous-alloys are promising in this aspect because amorphous-alloys are known to possess both high strength and good ductility compared to crystalline structures [13]. Therefore, we focused our search on amorphous-alloys for new membrane materials. As a result, it was found that amorphous Zr36 Ni64 alloy, which was quenched from the liquid using a single copper roller, was useful for gaseous-
hydrogen separation as a self-supported membrane [5]. Moreover, it was also found that the membrane spontaneously exhibited the activity of dissociating hydrogen molecules to atoms by exposure of both sides of the membrane to a hydrogen atmosphere at elevated temperatures. That is, we successfully demonstrated for the first time that there was a metal other than palladium-alloys which met all three requirements necessary for hydrogen-permeable membrane materials. Of course this alloy is much cheaper than palladium-alloys, but permeability was insufficient for industrial applications; for such use, it must be improved 5–10 times. However, because no other metals with similar characteristics have been found, there is no information about how to look for other membrane materials with permeability higher than those of amorphous Zr36 Ni64 . In this study, to widen the variety of alloys which meet two of the three requirements, high permeability and resistance to hydrogen embrittlement, Hf-substituted alloys of Zr in Zr36 Ni64 , we attempted quenching and then estimated permeability; the remaining requirement, activity to dissociate hydrogen, was provided by the palladium coating. The Hf is in the same IVA group as Zr and Hf–Ni system is very similar in the phase diagram to Zr–Ni system. That is why Hf was selected as a first element to substitute for Zr.
2. Experimental In a previous study [5], 64 at.% was adopted as a Ni composition because it is an eutectic composition of the Zr–Ni alloy system; eutectic alloys are generally known to become amorphous easily. Fortunately, the Hf–Ni system has an eutectic alloy with the same composition, so that a series of Zr–Hf–Ni alloys with 64 at.% Ni, Zr36−x Hfx Ni64 , were the focus in this study. They also seem to be made amorphous easily. The Zr (98%), Hf (99%), and Ni (>99.9%) were weighed, then melted in an arc furnace with argon atmosphere to produce alloys with desired compositions. They were crushed into pieces a few millimeters in size for the following rapid quenching. Rapid quenching was carried out in an argon atmosphere by using Makabe RQM-T-20. Alloy pieces were put into a fused quartz nozzle with a 6 mm by
S. Hara et al. / Journal of Membrane Science 211 (2003) 149–156
0.35 mm slit, melted by means of high frequency induction, and sprayed onto a rotating copper roller. The rotation rate was approximately 2000 rpm corresponding to 20 m/s in velocity of the cooling copper surface. Thus obtained alloy membranes were 5–6 mm wide and 30–40 m thick, characterized by X-ray diffraction (XRD) method and differential scanning calorimetry (DSC) to confirm them as amorphous. The DSC was carried out in an argon atmosphere at a heating rate of 10 K/min. Before hydrogen permeation experimentation, 100 nm palladium was deposited on both sides of the membrane by radio-frequency sputtering. The palladium coating was thick enough to give the membrane activity to hydrogen dissociation and recombination for permeation. Resistance of hydrogen permeation through coatings was neglected for estimation of permeability of amorphous-alloys because the coatings were much thinner than the amorphous-alloys and their permeability was higher. Hydrogen-permeation properties of thus prepared palladium-coated membranes were investigated using sweep gas. The measurement system is the same system as that in the previous study [5,14]. A membrane was mounted in a gas-permeation cell, where pure hydrogen or argon–hydrogen mixtures were introduced on one side of the membrane and pure argon on the other side as a sweep gas at a feed rate of 10−5 mol/s. The feed-side pressures ranged from 0.1 to 0.3 MPa and permeate-side pressure was atmospheric throughout. Concentration of effluent from the permeate-side was analyzed by gas chromatography to determine hydrogen-permeation rate.
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Fig. 1. XRD patterns for as-quenched Zr36−x Hfx Ni64 membranes obtained from the surfaces freely solidified (a) and contacted to the roller (b).
3. Results and discussion 3.1. Structure of as-quenched alloy membranes As expected, all of the as-quenched alloy membranes were ductile and strong in spite of thickness only 30–40 m so that they could be bent 180◦ to close contact, which is a typical characteristic of amorphous-alloys. Moreover, amorphous structure was confirmed by XRD. Fig. 1 shows XRD patterns obtained from each membrane surface. Every pattern has a broad peak (P2 ) around 2θ = 40◦ , indicating membranes had no
long-range order on atomic distribution; that is, they were amorphous. However, they were not always completely amorphous. With respect to freely solidified surfaces for x = 3, 6, and 9, a sharp peak showing existence of a crystalline phase was also superimposed. To prepare uniform amorphous membranes, it is necessary to quench them more rapidly or to customize alloys in terms of composition and additives for amorphous forming ability. From the peak angles, d-spacing can be calculated by Bragg’s equation, shown in Fig. 2. The figure shows that d-spacing corresponding to the main peak P2 was
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3.2. Thermal stability of amorphous-alloy membranes
Fig. 2. Change with Hf content in d-spacing corresponding to the peaks in Fig. 1 (b).
almost independent of Hf content. This is not surprising because the size of Hf, 0.3197 nm, approximates that of Zr, 0.3232 nm. Also, no relationship to concentration can be obtained with respect to full width at half maximum. On the other hand, some Hf-rich alloys have a peak P1 around 29◦ , whose d-spacing increasing with Hf. However, peak height was too low to indicate significant structural change. We conclude here that as-quenched Zr36−x Hfx Ni64 alloy membranes were confirmed as amorphous in most of the membranes and structure was little changed by Hf substitution.
The DSC curve for an as-quenched Zr36 Ni64 membrane had two close exothermic peaks around 840 K, suggesting a two-step crystallization. The first peak at lower temperature was always smaller than the second peak. Temperatures of the two peaks are shown in Fig. 3 as functions of Hf content as well as crystallization temperature determined as the onset of the first peak and heat during two-step crystallization. Two peaks moved toward higher temperatures with Hf substitute and approached each other to unite at complete substitution. Crystallization temperature also became higher with Hf, indicating that amorphous structure was more stabilized by Hf. This is reasonable because the 1463 K melting point of Hf36 Ni64 is higher than the 1343 K of Zr36 Ni64 . Crystallization heat is also known as a useful measure of amorphous stability but no relationship between the heat and Hf content can be recognized in Fig. 3. As shown by XRD, some membranes have small crystalline phases, which may hinder estimation of the intrinsic heat for uniform amorphous-alloys. 3.3. Hydrogen permeation rate Although, as-quenched alloy membranes did not have enough activity for hydrogen, not only amorphous Zr36 Ni64 , but also the Hf-containing
Fig. 3. Temperature of the two main peaks in DSC curves (䉱), crystallization temperature (䊉), and heat (䊐) of as-quenched Zr36−x Hfx Ni64 membranes.
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Fig. 4. Time profile of hydrogen flow rate in the outlet of the permeation cell at 573 K for a palladium-coated Zr27 Hf9 Ni64 membrane. Feed gas and total pressure are also designated.
amorphous-alloys could be activated by exposure to hydrogen at elevated temperatures. However, the effect for activation was not equal for all membranes. Even after hydrogen exposure, activity was often insufficient, so the subsequent permeation rate increased with time. On the other hand, membranes coated with palladium on both sides exhibited rather stable hydrogen permeation. This was why permeation was investigated in detail only for palladium-coated membranes in this work to discuss it without concern for surface activity. Fig. 4 shows a typical time profile of hydrogen flow rate in the sweep side outlet of the permeation cell during measurement at 573 K for a palladium-coated amorphous-Zr27 Hf9 Ni64 membrane, which is given as a product of total flow rate and hydrogen composition determined by gas chromatography. This figure shows that every time feed gas species or pressure changed, the hydrogen flow rate jumped to almost the same value as in the steady state within 5 min, the interval of composition analysis; this indicates fast response and high stability in permeation rate at least on the order of an hour. Furthermore, membranes were confirmed to keep hydrogen selectivity by introducing helium as feed gas instead of hydrogen after hydrogen permeation. Because the lower measurement limit to helium of this system was 10−11 mol/s, ideal hydrogen selectivity for all membranes was beyond 104 . It is rather difficult
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Fig. 5. Hydrogen permeation rate at 573 K through palladium√ coated Zr36−x Hfx Ni64 membranes as a function of p.
to achieve such high selectivity without using dense metallic membranes. 3.4. Hydrogen permeability of amorphous Zr36−x Hfx Ni64 alloys Hydrogen-permeation rate at 573 K is plotted in Fig. 5 against difference of square roots of partial pres√ sures, p. The permeation rate was clearly propor√ tional to p. If the membrane has defects less than 100 nm in size through which molecular hydrogen can pass, several permeation mechanisms can be considered, including the molecular sieve, Knudsen diffusion, surface diffusion, etc. In any case, permeation is usually linear to p in the pressure range examined. If the membrane has larger defects, permeation must obey the viscous flow mechanism to increase with pressure. Permeation behavior obtained was different from any case above, suggesting that permeation did not result from molecular hydrogen passing through any defect; in other words, we achieved complete hydrogen selectivity. Generally speaking, when it is controlled by diffusion of hydrogen atoms in the membrane, hydrogen permeation through a metal membrane, such as √ palladium-alloy membranes, is proportional to p. This directly relates to Sieverts’ law, where hydrogen concentration in metal is proportional to the square
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Fig. 6. Arrhenius plot of permeability of amorphous Zr36−x Hfx Ni64 alloys.
√ root of hydrogen partial pressure, C = K p. It is known, however, that in amorphous-alloys it does not always obey this law but can be approximately described by C = K p n , so that permeation rate, J, should be written for amorphous-alloys as follows [14]: J =
P¯ p n d
(1)
where d is membrane thickness. P¯ for n = 1/2 is widely used as permeability for metal membranes.
As seen from Fig. 5, permeation seems linear √ against p. Of course n might not be just 1/2 for amorphous Zr36−x Hfx Ni64 . It is, however, rather difficult to determine n more precisely due to measurement error and small deviation from 1/2, if any. Additionally, permeability determined using n = 1/2 is more useful than that using other n because it can be compared with many other metal membranes obeying Sieverts’ law. This is why permeability was determined assuming n = 1/2, which is plotted against 1000/T in Fig. 6. In this study, permeation rate was measured several times at 573 K to determine stability. For example, a permeation test for amorphous Hf36 Ni64 was carried out in the order of 573, 523, 548, 498, 473, 573, 598, and 573 K. Though 3 days elapsed between the first two measurements at 573 K, permeability remained almost unchanged, suggesting that permeation rate time dependency was very small in the range of 473–573 K. On the contrary, permeability clearly decreased between the latter two measurements at 573 K during only 1 day. Similarly, permeability for x = 0 and 9 was linear at temperatures between 573 and 473 K, but lower at 598 and 623 K measured after 573 K than the line. Therefore, it can be concluded that permeability decreased over time at high temperatures over 573 K and that this was a common tendency to all of alloys studied. For this reason, permeability at 573 K or less, which was always examined before higher temperatures, was
Fig. 7. Activation energy (䊉) and pre-exponential factor (䊐) for hydrogen permeation.
S. Hara et al. / Journal of Membrane Science 211 (2003) 149–156
used for fit by P¯ = P¯0 exp(−EP /RT), where R is the gas constant (8.314 J/(K mol)). Determined parameters are shown in Fig. 7. With respect to P¯0 , it was difficult to find any tendency with Hf content because values included rather large error due to long extrapolation from the measured temperature range to the y-axis in the Arrhenius plot. On the other hand, activation energy for permeation, EP , was found to become larger with Hf-substitute amount. For metal membranes, activation energy for permeation is generally given as a sum of the activation energy for hydrogen diffusion and the enthalpy change for hydrogen-solution in the membrane. Because the enthalpy change for pure Hf, −37 kJ/mol, is larger than −64 kJ/mol for pure Zr [15], that for the alloys was expected to rise with Hf substitution to Zr in the amorphous Zr36 Ni64 alloy. This may explain why the Hf substitute increased in activation energy for permeation, resulting in decreased permeability. Needless to say, Hf content may also affect diffusivity. An attempt to investigate hydrogen-solution properties of this series of amorphous-alloys is now being made to elucidate that influence.
4. Conclusions A series of Zr36−x Hfx Ni64 (0 ≤ x ≤ 36) alloys were successfully prepared by the rapid quenching method using a copper roller. The XRD showed that as-quenched membranes were amorphous and almost the same in atomic distribution without respect to Hf content; DSC showed crystallization temperature increased with increasing Hf. Palladium-coated membranes were never broken during permeation testing in the range of 473–623 K. The permeation rate was stable for a few days at 573 K or less and approximately proportional to the square root difference of hydrogen partial pressures across the membrane, suggesting that permeation was controlled by the diffusion process of hydrogen atoms in the membrane; it was also inferred that membranes had no defect which all other gases could pass through. Permeability was found to decrease with Hf-substitute amount to Zr due to increased activation energy for permeation. Although, as-quenched amorphous Zr36−x Hfx Ni64 membranes did not always show enough surface
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activity to hydrogen necessary for permeation, this study widened the variety of hydrogen-permeable amorphous-alloys. This is greatly significant because amorphous Zr36 Ni64 had been the only amorphous-alloy useful for gaseous-hydrogen separation without any support. Further main works will clarify reasons for increased activation energy for permeation with Hf content and then make guidelines to improve of amorphous-alloy permeability.
Acknowledgements Portions of this work were carried under the Visiting Researcher’s Program of Laboratory for Developmental Research of Advanced Materials, Institute for Materials Research, Tohoku University, and supported by Industrial Technology Research Grant Program in 2001 from New Energy and Industrial Technology Development Organization (NEDO) of Japan. References [1] J.E. Philpott, Hydrogen diffusion technology, Platinum Met. Rev. 29 (1985) 12. [2] Y. Shirasaki, Y. Ohta, K. Kobayashi, K. Kuroda, Membrane Reformer Technical Bulletin, Mitsubishi Heavy Industries Ltd., Yokohama, 1996. [3] N. Itoh, A membrane reactor using palladium, AIChE J. 33 (1987) 1576. [4] C. Nishimura, M. Komaki, M. Amano, Hydrogen permeation characteristics of vanadium–nickel alloys, Mater. Trans. JIM 32 (1991) 501. [5] S. Hara, K. Sakaki, N. Itoh, H.-M. Kimura, A. Inoue, An amorphous alloy membrane without noble metals for gaseous hydrogen separation, J. Membr. Sci. 164 (2000) 289. [6] R.E. Buxbaum, P.C. Hsu, Hydrogen transport and embrittlement for palladium coated vanadium–chromium–titanium alloys, J. Nucl. Mater. 189 (1992) 183. [7] J.O. Ström-Olsen, Y. Zhao, D.H. Ryan, Y. Huai, R.W. Cochrane, Hydrogen diffusion in amorphous Ni–Zr, J. Less-Common Met. 172-174 (1991) 922. [8] S. Uemiya, T. Matsuda, E. Kikuchi, Hydrogen permeable palladium–silver alloy membrane supported on porous ceramics, J. Membr. Sci. 56 (1991) 315. [9] V. Jayaraman, Y.S. Lin, M. Pakala, R.Y. Lin, Fabrication of ultrathin metallic membranes on ceramic supports by sputter deposition, J. Membr. Sci. 99 (1995) 89. [10] S. Akiyama, H. Mizuta, H. Anzai, K. Kusakabe, S. Morooka, Preparation of gas separation membranes supported on ␣-alumina porous hollow fibers, in: Proceedings of the 5th
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International Conference on Inorganic Membrane, Nagoya, Japan, 1998, p. 398. [11] H. Sakaguchi, G. Adachi, J. Shiokawa, Hydrogen separation using LaNi5 films deposited on polymer membranes, Bull. Chem. Soc. Jpn. 61 (1988) 521. [12] M. Amano, Y. Sasaki, T. Yoshioka, Effect oxide phase and Ti phase on hydriding characteristics of FeTi phase, J. Jpn. Inst. Met. 45 (1981) 957.
[13] T. Masumoto, R. Maddin, The mechanical properties of palladium 20 a/o silicon alloy quenched from the liquid state, Acta Metall. 19 (1971) 725. [14] N. Itoh, W.-C. Xu, S. Hara, H.-M. Kimura, T. Masumoto, Permeability of hydrogen in amorphous Pd(1−x) Six alloys at elevated temperatures, J. Membr. Sci. 139 (1998) 29. [15] Y. Fukai, Properties of hydrogen in metal lattice, Shyokubai 33 (1991) 254 (Japanese).
Hydrogen permeation through amorphous-Zr36−x Hfx Ni64 -alloy membranes S. Hara a,∗ , N. Hatakeyama a , N. Itoh a , H.-M. Kimura b , A. Inoue b a
National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba Central 5, Tsukuba 305-8565, Japan b Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 4 February 2002; accepted 9 September 2002
Abstract Amorphous-alloy membranes of a series of Zr36−x Hfx Ni64 (0 ≤ x ≤ 36) were successfully prepared by the rapid-quenching method using a single roller; their amorphous structure and thermal behavior were characterized by XRD and DSC. All of the membranes, coated by palladium to provide activity to hydrogen dissociation and recombination, were sufficiently robust in a hydrogen atmosphere and showed stable permeability only to hydrogen at least in the range of 473–573 K. On the other hand, over 573 K, the permeation rate slowly decreased over time. Permeability was found to decrease with Hf substitute amount for Zr mainly due to increase in activation energy for permeation. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Hydrogen permeation; Amorphous-alloys; Zr–Hf–Ni alloys; Hydrogen diffusion; Hydrogen embrittlement
1. Introduction Hydrogen-permeable metal membranes generally have high selectivity, permeability, and thermal stability, so that they are expected to be used for various industrial applications, including ultra-pure hydrogen production for semiconductor manufacture [1], carbon-monoxide removal from reformed gas for fuel cells [2], and membrane reactors to improve and control chemical reactions [3]. However, metal membranes employed now in industry are strictly palladium or palladium-alloy membranes, which are extremely expensive. It is, therefore, desirable to develop cheaper membrane materials substitutable for them. ∗ Corresponding author. Tel.: +81-298-61-9336; fax: +81-298-61-4674. E-mail address: [email protected] (S. Hara).
Metals used for gaseous-hydrogen separation have to meet at least three requirements: high permeability, activity to dissociate hydrogen molecules, and resistance to hydrogen embrittlement [4,5]. Of these requirements, finding metals with high hydrogen permeability is not difficult because the permeation rate can be estimated from hydrogen diffusivity and solubility in the metal if hydrogen permeation is controlled by atomic-hydrogen diffusion in the membrane. For example, hydrogen is known to diffuse easily between metal atoms in a body-centered cubic (bcc) lattice; so, high permeability is expected for bcc metals with high solubility such as vanadium, niobium, and tantalum [6]. However, because hydrogen in metals generally weakens the metals (hydrogen embrittlement), most such metals which are expected to have high permeability absorb a large amount of hydrogen in a hydrogen atmosphere; they then break due to stress from differential pressure across the membrane
0376-7388/02/$ – see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 3 7 6 - 7 3 8 8 ( 0 2 ) 0 0 4 1 6 - 7
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and/or internal stress caused by absorbed hydrogen atoms. Therefore, the developmental obstacle for metal membrane materials is not improved permeability but resolution of the two conflicting requirements of high permeability and resistance to hydrogen embrittlement. Nevertheless, much research has avoided challenging the hydrogen embrittlement issue. For instance, electrochemical methods have been employed to find hydrogen permeability through metal membranes [7]. In this case, the driving force for hydrogen permeation is differential electrochemical potential for hydrogen ions in solutions on both sides of the membrane, so that no mechanical stress need be applied to the membrane during measurement. Therefore, the same membranes cannot be said to also be useful for gaseous-hydrogen separation without breaking when some stress is applied. Many experiments have been reported on metal membranes prepared on supports to improve mechanical strength. Such a composite structure is one effective way to increase permeation flux; indeed, it is one which may popularize use of metal membranes. However, making composite membranes is rather difficult, as reported by many authors for palladium composite membranes [8–10]. Sakaguchi et al. [11] deposited an amorphous LaNi5 alloy layer on a support by the sputtering method, but they never attained complete selectivity to hydrogen due to metal layer defects. A composite membrane composed of amorphous FeTi and a palladium support by Amano et al. [12] showed complete selectivity; but, as a practical matter, such a membrane is not substitutable for palladium membranes because dense palladium sheet was employed as a support. As described above, not a few researchers have reported on hydrogen permeability through metal membranes other than palladium-alloys, but many of them have made various devices to avoid hydrogen embrittlement which are not always useful for industrial application. Amorphous-alloys are promising in this aspect because amorphous-alloys are known to possess both high strength and good ductility compared to crystalline structures [13]. Therefore, we focused our search on amorphous-alloys for new membrane materials. As a result, it was found that amorphous Zr36 Ni64 alloy, which was quenched from the liquid using a single copper roller, was useful for gaseous-
hydrogen separation as a self-supported membrane [5]. Moreover, it was also found that the membrane spontaneously exhibited the activity of dissociating hydrogen molecules to atoms by exposure of both sides of the membrane to a hydrogen atmosphere at elevated temperatures. That is, we successfully demonstrated for the first time that there was a metal other than palladium-alloys which met all three requirements necessary for hydrogen-permeable membrane materials. Of course this alloy is much cheaper than palladium-alloys, but permeability was insufficient for industrial applications; for such use, it must be improved 5–10 times. However, because no other metals with similar characteristics have been found, there is no information about how to look for other membrane materials with permeability higher than those of amorphous Zr36 Ni64 . In this study, to widen the variety of alloys which meet two of the three requirements, high permeability and resistance to hydrogen embrittlement, Hf-substituted alloys of Zr in Zr36 Ni64 , we attempted quenching and then estimated permeability; the remaining requirement, activity to dissociate hydrogen, was provided by the palladium coating. The Hf is in the same IVA group as Zr and Hf–Ni system is very similar in the phase diagram to Zr–Ni system. That is why Hf was selected as a first element to substitute for Zr.
2. Experimental In a previous study [5], 64 at.% was adopted as a Ni composition because it is an eutectic composition of the Zr–Ni alloy system; eutectic alloys are generally known to become amorphous easily. Fortunately, the Hf–Ni system has an eutectic alloy with the same composition, so that a series of Zr–Hf–Ni alloys with 64 at.% Ni, Zr36−x Hfx Ni64 , were the focus in this study. They also seem to be made amorphous easily. The Zr (98%), Hf (99%), and Ni (>99.9%) were weighed, then melted in an arc furnace with argon atmosphere to produce alloys with desired compositions. They were crushed into pieces a few millimeters in size for the following rapid quenching. Rapid quenching was carried out in an argon atmosphere by using Makabe RQM-T-20. Alloy pieces were put into a fused quartz nozzle with a 6 mm by
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0.35 mm slit, melted by means of high frequency induction, and sprayed onto a rotating copper roller. The rotation rate was approximately 2000 rpm corresponding to 20 m/s in velocity of the cooling copper surface. Thus obtained alloy membranes were 5–6 mm wide and 30–40 m thick, characterized by X-ray diffraction (XRD) method and differential scanning calorimetry (DSC) to confirm them as amorphous. The DSC was carried out in an argon atmosphere at a heating rate of 10 K/min. Before hydrogen permeation experimentation, 100 nm palladium was deposited on both sides of the membrane by radio-frequency sputtering. The palladium coating was thick enough to give the membrane activity to hydrogen dissociation and recombination for permeation. Resistance of hydrogen permeation through coatings was neglected for estimation of permeability of amorphous-alloys because the coatings were much thinner than the amorphous-alloys and their permeability was higher. Hydrogen-permeation properties of thus prepared palladium-coated membranes were investigated using sweep gas. The measurement system is the same system as that in the previous study [5,14]. A membrane was mounted in a gas-permeation cell, where pure hydrogen or argon–hydrogen mixtures were introduced on one side of the membrane and pure argon on the other side as a sweep gas at a feed rate of 10−5 mol/s. The feed-side pressures ranged from 0.1 to 0.3 MPa and permeate-side pressure was atmospheric throughout. Concentration of effluent from the permeate-side was analyzed by gas chromatography to determine hydrogen-permeation rate.
151
Fig. 1. XRD patterns for as-quenched Zr36−x Hfx Ni64 membranes obtained from the surfaces freely solidified (a) and contacted to the roller (b).
3. Results and discussion 3.1. Structure of as-quenched alloy membranes As expected, all of the as-quenched alloy membranes were ductile and strong in spite of thickness only 30–40 m so that they could be bent 180◦ to close contact, which is a typical characteristic of amorphous-alloys. Moreover, amorphous structure was confirmed by XRD. Fig. 1 shows XRD patterns obtained from each membrane surface. Every pattern has a broad peak (P2 ) around 2θ = 40◦ , indicating membranes had no
long-range order on atomic distribution; that is, they were amorphous. However, they were not always completely amorphous. With respect to freely solidified surfaces for x = 3, 6, and 9, a sharp peak showing existence of a crystalline phase was also superimposed. To prepare uniform amorphous membranes, it is necessary to quench them more rapidly or to customize alloys in terms of composition and additives for amorphous forming ability. From the peak angles, d-spacing can be calculated by Bragg’s equation, shown in Fig. 2. The figure shows that d-spacing corresponding to the main peak P2 was
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3.2. Thermal stability of amorphous-alloy membranes
Fig. 2. Change with Hf content in d-spacing corresponding to the peaks in Fig. 1 (b).
almost independent of Hf content. This is not surprising because the size of Hf, 0.3197 nm, approximates that of Zr, 0.3232 nm. Also, no relationship to concentration can be obtained with respect to full width at half maximum. On the other hand, some Hf-rich alloys have a peak P1 around 29◦ , whose d-spacing increasing with Hf. However, peak height was too low to indicate significant structural change. We conclude here that as-quenched Zr36−x Hfx Ni64 alloy membranes were confirmed as amorphous in most of the membranes and structure was little changed by Hf substitution.
The DSC curve for an as-quenched Zr36 Ni64 membrane had two close exothermic peaks around 840 K, suggesting a two-step crystallization. The first peak at lower temperature was always smaller than the second peak. Temperatures of the two peaks are shown in Fig. 3 as functions of Hf content as well as crystallization temperature determined as the onset of the first peak and heat during two-step crystallization. Two peaks moved toward higher temperatures with Hf substitute and approached each other to unite at complete substitution. Crystallization temperature also became higher with Hf, indicating that amorphous structure was more stabilized by Hf. This is reasonable because the 1463 K melting point of Hf36 Ni64 is higher than the 1343 K of Zr36 Ni64 . Crystallization heat is also known as a useful measure of amorphous stability but no relationship between the heat and Hf content can be recognized in Fig. 3. As shown by XRD, some membranes have small crystalline phases, which may hinder estimation of the intrinsic heat for uniform amorphous-alloys. 3.3. Hydrogen permeation rate Although, as-quenched alloy membranes did not have enough activity for hydrogen, not only amorphous Zr36 Ni64 , but also the Hf-containing
Fig. 3. Temperature of the two main peaks in DSC curves (䉱), crystallization temperature (䊉), and heat (䊐) of as-quenched Zr36−x Hfx Ni64 membranes.
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Fig. 4. Time profile of hydrogen flow rate in the outlet of the permeation cell at 573 K for a palladium-coated Zr27 Hf9 Ni64 membrane. Feed gas and total pressure are also designated.
amorphous-alloys could be activated by exposure to hydrogen at elevated temperatures. However, the effect for activation was not equal for all membranes. Even after hydrogen exposure, activity was often insufficient, so the subsequent permeation rate increased with time. On the other hand, membranes coated with palladium on both sides exhibited rather stable hydrogen permeation. This was why permeation was investigated in detail only for palladium-coated membranes in this work to discuss it without concern for surface activity. Fig. 4 shows a typical time profile of hydrogen flow rate in the sweep side outlet of the permeation cell during measurement at 573 K for a palladium-coated amorphous-Zr27 Hf9 Ni64 membrane, which is given as a product of total flow rate and hydrogen composition determined by gas chromatography. This figure shows that every time feed gas species or pressure changed, the hydrogen flow rate jumped to almost the same value as in the steady state within 5 min, the interval of composition analysis; this indicates fast response and high stability in permeation rate at least on the order of an hour. Furthermore, membranes were confirmed to keep hydrogen selectivity by introducing helium as feed gas instead of hydrogen after hydrogen permeation. Because the lower measurement limit to helium of this system was 10−11 mol/s, ideal hydrogen selectivity for all membranes was beyond 104 . It is rather difficult
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Fig. 5. Hydrogen permeation rate at 573 K through palladium√ coated Zr36−x Hfx Ni64 membranes as a function of p.
to achieve such high selectivity without using dense metallic membranes. 3.4. Hydrogen permeability of amorphous Zr36−x Hfx Ni64 alloys Hydrogen-permeation rate at 573 K is plotted in Fig. 5 against difference of square roots of partial pres√ sures, p. The permeation rate was clearly propor√ tional to p. If the membrane has defects less than 100 nm in size through which molecular hydrogen can pass, several permeation mechanisms can be considered, including the molecular sieve, Knudsen diffusion, surface diffusion, etc. In any case, permeation is usually linear to p in the pressure range examined. If the membrane has larger defects, permeation must obey the viscous flow mechanism to increase with pressure. Permeation behavior obtained was different from any case above, suggesting that permeation did not result from molecular hydrogen passing through any defect; in other words, we achieved complete hydrogen selectivity. Generally speaking, when it is controlled by diffusion of hydrogen atoms in the membrane, hydrogen permeation through a metal membrane, such as √ palladium-alloy membranes, is proportional to p. This directly relates to Sieverts’ law, where hydrogen concentration in metal is proportional to the square
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Fig. 6. Arrhenius plot of permeability of amorphous Zr36−x Hfx Ni64 alloys.
√ root of hydrogen partial pressure, C = K p. It is known, however, that in amorphous-alloys it does not always obey this law but can be approximately described by C = K p n , so that permeation rate, J, should be written for amorphous-alloys as follows [14]: J =
P¯ p n d
(1)
where d is membrane thickness. P¯ for n = 1/2 is widely used as permeability for metal membranes.
As seen from Fig. 5, permeation seems linear √ against p. Of course n might not be just 1/2 for amorphous Zr36−x Hfx Ni64 . It is, however, rather difficult to determine n more precisely due to measurement error and small deviation from 1/2, if any. Additionally, permeability determined using n = 1/2 is more useful than that using other n because it can be compared with many other metal membranes obeying Sieverts’ law. This is why permeability was determined assuming n = 1/2, which is plotted against 1000/T in Fig. 6. In this study, permeation rate was measured several times at 573 K to determine stability. For example, a permeation test for amorphous Hf36 Ni64 was carried out in the order of 573, 523, 548, 498, 473, 573, 598, and 573 K. Though 3 days elapsed between the first two measurements at 573 K, permeability remained almost unchanged, suggesting that permeation rate time dependency was very small in the range of 473–573 K. On the contrary, permeability clearly decreased between the latter two measurements at 573 K during only 1 day. Similarly, permeability for x = 0 and 9 was linear at temperatures between 573 and 473 K, but lower at 598 and 623 K measured after 573 K than the line. Therefore, it can be concluded that permeability decreased over time at high temperatures over 573 K and that this was a common tendency to all of alloys studied. For this reason, permeability at 573 K or less, which was always examined before higher temperatures, was
Fig. 7. Activation energy (䊉) and pre-exponential factor (䊐) for hydrogen permeation.
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used for fit by P¯ = P¯0 exp(−EP /RT), where R is the gas constant (8.314 J/(K mol)). Determined parameters are shown in Fig. 7. With respect to P¯0 , it was difficult to find any tendency with Hf content because values included rather large error due to long extrapolation from the measured temperature range to the y-axis in the Arrhenius plot. On the other hand, activation energy for permeation, EP , was found to become larger with Hf-substitute amount. For metal membranes, activation energy for permeation is generally given as a sum of the activation energy for hydrogen diffusion and the enthalpy change for hydrogen-solution in the membrane. Because the enthalpy change for pure Hf, −37 kJ/mol, is larger than −64 kJ/mol for pure Zr [15], that for the alloys was expected to rise with Hf substitution to Zr in the amorphous Zr36 Ni64 alloy. This may explain why the Hf substitute increased in activation energy for permeation, resulting in decreased permeability. Needless to say, Hf content may also affect diffusivity. An attempt to investigate hydrogen-solution properties of this series of amorphous-alloys is now being made to elucidate that influence.
4. Conclusions A series of Zr36−x Hfx Ni64 (0 ≤ x ≤ 36) alloys were successfully prepared by the rapid quenching method using a copper roller. The XRD showed that as-quenched membranes were amorphous and almost the same in atomic distribution without respect to Hf content; DSC showed crystallization temperature increased with increasing Hf. Palladium-coated membranes were never broken during permeation testing in the range of 473–623 K. The permeation rate was stable for a few days at 573 K or less and approximately proportional to the square root difference of hydrogen partial pressures across the membrane, suggesting that permeation was controlled by the diffusion process of hydrogen atoms in the membrane; it was also inferred that membranes had no defect which all other gases could pass through. Permeability was found to decrease with Hf-substitute amount to Zr due to increased activation energy for permeation. Although, as-quenched amorphous Zr36−x Hfx Ni64 membranes did not always show enough surface
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activity to hydrogen necessary for permeation, this study widened the variety of hydrogen-permeable amorphous-alloys. This is greatly significant because amorphous Zr36 Ni64 had been the only amorphous-alloy useful for gaseous-hydrogen separation without any support. Further main works will clarify reasons for increased activation energy for permeation with Hf content and then make guidelines to improve of amorphous-alloy permeability.
Acknowledgements Portions of this work were carried under the Visiting Researcher’s Program of Laboratory for Developmental Research of Advanced Materials, Institute for Materials Research, Tohoku University, and supported by Industrial Technology Research Grant Program in 2001 from New Energy and Industrial Technology Development Organization (NEDO) of Japan. References [1] J.E. Philpott, Hydrogen diffusion technology, Platinum Met. Rev. 29 (1985) 12. [2] Y. Shirasaki, Y. Ohta, K. Kobayashi, K. Kuroda, Membrane Reformer Technical Bulletin, Mitsubishi Heavy Industries Ltd., Yokohama, 1996. [3] N. Itoh, A membrane reactor using palladium, AIChE J. 33 (1987) 1576. [4] C. Nishimura, M. Komaki, M. Amano, Hydrogen permeation characteristics of vanadium–nickel alloys, Mater. Trans. JIM 32 (1991) 501. [5] S. Hara, K. Sakaki, N. Itoh, H.-M. Kimura, A. Inoue, An amorphous alloy membrane without noble metals for gaseous hydrogen separation, J. Membr. Sci. 164 (2000) 289. [6] R.E. Buxbaum, P.C. Hsu, Hydrogen transport and embrittlement for palladium coated vanadium–chromium–titanium alloys, J. Nucl. Mater. 189 (1992) 183. [7] J.O. Ström-Olsen, Y. Zhao, D.H. Ryan, Y. Huai, R.W. Cochrane, Hydrogen diffusion in amorphous Ni–Zr, J. Less-Common Met. 172-174 (1991) 922. [8] S. Uemiya, T. Matsuda, E. Kikuchi, Hydrogen permeable palladium–silver alloy membrane supported on porous ceramics, J. Membr. Sci. 56 (1991) 315. [9] V. Jayaraman, Y.S. Lin, M. Pakala, R.Y. Lin, Fabrication of ultrathin metallic membranes on ceramic supports by sputter deposition, J. Membr. Sci. 99 (1995) 89. [10] S. Akiyama, H. Mizuta, H. Anzai, K. Kusakabe, S. Morooka, Preparation of gas separation membranes supported on ␣-alumina porous hollow fibers, in: Proceedings of the 5th
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