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Hydrogenation of a TiFe-based alloy at high pressures and temperatures
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 4 0 ( 2 0 1 5 ) 3 2 8 3 e3 2 8 7
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Short Communication
Hydrogenation of a TiFe-based alloy at high pressures and temperatures Naruki Endo a,*, Itoko Saita b, Yumiko Nakamura b, Hiroyuki Saitoh c, Akihiko Machida c a
Renewable Energy Research Center, National Institute of Advanced Industrial Science and Technology (AIST), 2-2-9 Machiike-dai, Koriyama, Fukushima 963-0215, Japan b Energy Technology Research Institute, National Institute of Advanced Industrial Science and Technology (AIST), AIST Central-5, 1-1-1, Higashi, Tsukuba, Ibaraki 305-8565, Japan c Quantum Beam Science Center, Japan Atomic Energy Agency, 1-1-1 Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5148, Japan
article info
abstract
Article history:
We investigated the hydrogenation of a ternary TiFe0.8Mn0.2 alloy under high pressure (up
Received 19 November 2014
to 6.5 GPa) and at high temperature (up to 600 C) by in situ synchrotron radiation X-ray
Received in revised form
diffraction measurements. A chemical orderedisorder phase transition and the formation
30 December 2014
of body-centered cubic (BCC) hydride were observed. The phase-transition temperature
Accepted 3 January 2015
was 400 C, which is 200 C lower than that of a binary TiFe alloy. A disproportionation
Available online 29 January 2015
reaction, which was observed for the binary TiFe alloy, did not occur at this temperature. The BCC hydride existed stably, and new hydrides were not formed after the BCC hydride
Keywords:
became a single phase. Since the unit cell volume of the BCC hydride was almost equal to
TiFe-based alloy
that of the g hydride, TiFe0.8Mn0.2H1.8, under the same pressureetemperature conditions,
Hydrogen storage
the hydrogen content of the former was roughly estimated as the same as that of the latter
High hydrogen pressure
(hydrogen to metal atom ratio ¼ 0.9).
Body-centered cubic
Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
In situ X-ray diffraction measurement
Introduction High hydrogen capacity is one of the most important property for hydrogen storage materials. We have been investigating the hydrogenation of CsCl-type TiFe alloys [1] under high hydrogen pressure to realize a novel hydride with high hydrogen content. In a recent study, we reported that when a
binary TiFe alloy was hydrogenated at 5 GPa and 600 C for 150 min, a chemical orderedisorder phase transition occurred, after which a BCC hydride was formed [2]. The unit cell volume of the hydride expanded by 21.0%, which is larger than that of the g hydride, TiFeH1.9, prepared under nearambient conditions (~18%) [3]. Since many tetrahedral (T) and octahedral (O) interstitial sites were still vacant in the
* Corresponding author. Tel.: þ81 29 861 2647; fax: þ81 24 963 0828. E-mail address: [email protected] (N. Endo). http://dx.doi.org/10.1016/j.ijhydene.2015.01.015 0360-3199/Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
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i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 4 0 ( 2 0 1 5 ) 3 2 8 3 e3 2 8 7
metal lattice of the BCC hydride, we believe that the hydride might have been able to absorb more hydrogen. As recently reported, results of first-principle calculations suggested that it is possible to form a novel TiFe hydride with high hydrogen content [4], a tetragonal structure (P4/mmm; Z ¼ 1), and a hydrogen atom (H) to metal atom (M) ratio of H/ M ¼ 2.0. The hydride was calculated to be unstable under near-ambient conditions. The calculations also pointed out that some alterations in the composition of the alloy or in the reaction pressure and temperature are necessary to synthesize the hydride. Since it has been reported that hydrogenation below 100 C and several tens of MPa does not increase the hydrogen content by Ref. [5], the hydride with H/M ¼ 2.0 is expected to be formed by hydrogenation under high pressure (>1 GPa) and temperature (>100 C). In our earlier study in which the TiFe alloy was hydrogenated at 5 GPa and 600 C, a disproportionation reaction occurred and the BCC hydride immediately decomposed after the formation of the single BCC phase [2]. It is thought that if the disproportionation reaction had been suppressed, the TiFe hydride with H/M ¼ 2.0 or other novel TiFe hydrides with high hydrogen content could have been formed by further hydrogenation of the BCC hydride. The disproportionation reaction occurred because a high temperature was necessary to hydrogenate the TiFe alloy. However, the temperature of 600 C did not represent the thermodynamic condition of hydrogenationdthe kinetic effects of the surface oxides raised the reaction temperature, which was the temperature at which the sample was completely hydrogenated [2]. As the kinetics effects cannot be avoided when the binary alloy is used, a more reactive sample should be chosen for the synthesis of novel hydrides. In the present study, we chose a ternary TiFe0.8Mn0.2 alloy because it is highly reactive for hydrogenation comparable to LaNi5 [6,7]. Previous studies have reported that substituting Fe for Mn improves the activation properties of the TiFe alloy and the kinetics of absorption and desorption of the ternary alloy is faster than that of the binary alloy [6e8]. We investigated the hydrogenation of the TiFe0.8Mn0.2 alloy at high pressures and temperatures to verify the formation of novel hydrides with high hydrogen content. From the investigation of the hydrogenation processes by in situ synchrotron radiation (SR) X-ray diffraction (XRD) measurements, we confirmed that they were similar to those of the binary alloy: a
chemical orderedisorder phase transition and the formation of a BCC hydride were observed. The hydrogenation temperature was lower than that of the binary TiFe alloy, and the disproportionation reaction was suppressed. We estimated the hydrogen content of the BCC hydride from a comparison of the unit cell volumes of the BCC and g hydrides.
Experimental The starting materials were TiFe0.8Mn0.2 alloy and a g hydride, TiFe0.8Mn0.2H1.8 (H/M ¼ 0.9) powder. The particles were less than 100 mm in size. The lattice parameters of the alloy and the ˚ , and hydride before the experiments were a ¼ 2.986(1) A ˚ , b ¼ 2.8338(3) A ˚ , c ¼ 4.719(3) A ˚ , and b ¼ 97.31(2), a ¼ 4.7024(3) A which are slightly larger than those of TiFe and g-phase TiFeH1.9 [9,10] because of the Mn substitution [11]. The crystal structure of the g hydride has been reported as monoclinic (P2/m; Z ¼ 2) [9e11] and orthorhombic (Cmmm; Z ¼ 4) [3,12]. In the present study, we nominally chose the monoclinic structure to represent the g hydride since the same results were obtained from the monoclinic and the orthorhombic structures. High pressureetemperature conditions were generated with a cubic-type multi-anvil high-pressure apparatus installed in the beamline BL14B1 at SPring-8, Japan [13]. Powder XRD profiles of the samples were obtained in situ by using high-brilliance SR X-ray as the incident beam. Here, the incident beam was white X-ray from SR light source and the energy-dispersive XRD method was employed [14], which is the standard technique for high-pressure experiments using a multi-anvil high-pressure press [15]. The diffracted X-rays were measured by a Ge solid state detector mounted on a goniometer. The diffraction angle 2q was fixed to 6 in all ˚ ) was calculated from measurements. The lattice spacing d (A the energies of the diffracted X-rays E (keV) based on Bragg's equation: 2d sin q ¼ 12:4=E
(1)
The measurement time was 1.0 min. Fig. 1 shows the schematics of the high-pressure cell assemblies with and without hydrogen sources. In both highpressure cell assemblies, the powder samples were filled in pyrolytic boron nitride (PBN) capsules. The high-pressure cell
Fig. 1 e Schematic of the high-pressure cell assembly used for the present study: (a) with internal hydrogen sources; (b) without internal hydrogen sources.
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 4 0 ( 2 0 1 5 ) 3 2 8 3 e3 2 8 7
assembly for hydrogenation was firstly developed by Fukai and Okuma [16] and has been used by other groups for highpressure synthesis of metal hydrides [2,14,15,17]. We modified the high-pressure cell assembly to determine the high hydrogen pressure. Fig. 1(a) shows the high-pressure cell assembly for hydrogenation used in the present study. NaCl powder, which was mixed homogeneously with BN powder to prevent grain growth at high temperature, was used as a pressure marker. The pressure was calculated from the unit cell volume of the pressure marker based on Decker's equation of state [18]. The hydrogenation reactions were carried out as follows: The samples were pressurized to ~6 GPa at room temperature (RT). Next, the sample and the hydrogen sources were initially heated to 600 C at a rate of 100 C/min to evolve hydrogen by decomposition of the hydrogen sources [16]. Upon reaching 600 C, the sample was immediately cooled to 300 C for 0.1 min. The reaction temperatures were set in the range of 300e600 C. In the case of the reaction without the hydrogen sources [Fig. 1(b)], the temperature was raised to a prescribed temperature of 300e600 C at a heating rate of 100 C/min and held at the conditions. The unit cell volumes of the g hydride before and after the experiments were the same, indicating that hydrogen was not released during the experiments. Details of the sample preparations, the experimental apparatus, and experimental procedures are described elsewhere [6,14,19].
Results and discussion The hydrogenation processes of the TiFe0.8Mn0.2 alloy at 6.3(2) GPa were investigated in the temperature range of 300e600 C. The hydrogenation reaction did not proceed completely at 300 C, while the entire hydrogenation proceeded at 400 C without the disproportionation reaction.
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When the hydrogenation reaction took place at temperatures above 500 C, the disproportionation reaction occurred immediately after the formation of the single BCC phase; the hydrogen absorption process was completed within 10 min at 600 C (cf. 150 min for TiFe [2]). Fig. 2(a) shows selected in situ SR-XRD profiles of the TiFe0.8Mn0.2 alloy at around 6.3(2) GPa. The bottom pattern represents the XRD profile of the alloy after pressurizing at RT. Hydrogen was not emitted from the internal hydrogen sources at this temperature. The diffraction peaks were slightly broadened by strain, which was induced during the pressurizing process. The 110, 200, and 220 peaks of the alloy phase shifted toward the low energy side just after the alloy was kept in the hydrogen fluid at 400 C. The lattice parameter of the ˚ , which sample after it was heated to 400 C was a ¼ 2.975(2) A ˚ is larger than that of the alloy (a ¼ 2.934(3) A) under the same conditions without hydrogen. The alloy was hydrogenated at this stage and a solid solution with hydrogen was formed. After the sample had been held in the hydrogen fluid for 30 min, shoulders appeared on the lower energy side of the diffraction peaks from the solid-solution phase, indicating that a new hydride phase was formed. The peaks from the hydride were indexed by a cubic unit cell. After 170 min of the hydrogenation treatment, the Bragg peaks from the solid solution vanished and the hydride phase became a single phase. A continuous lattice expansion of the hydride occurred and ˚ after 530 min the lattice parameter increased to a ¼ 3.116(4) A of the hydrogenation treatment. Fig. 2(b) shows an enlargement of the XRD profiles in Fig. 2(a). The small peaks correspond to the Bragg peaks from NaCl and the C15 Laves phase [7]. The arrows indicate the 111 and 210 peaks from the CsCl-type structure, whose intensities decreased with increasing volume of the hydride during hydrogenation; these peaks were hardly observable when the hydride became a single phase [top profile, Fig. 2(b)]. These peaks represent the fundamental reflections from the CsCl
Fig. 2 e In situ SR-XRD patterns of (a) the TiFe0.8Mn0.2 alloy with hydrogen and (c) the g hydride, TiFe0.8Mn0.2H1.8, with and without hydrogen at high pressures. (b) Enlargement of the XRD profiles in (a). The arrows indicate the 111 and 210 peaks from the CsCl-type structure.
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structure, but such reflections are forbidden from the BCC structure. This means that a chemical orderedisorder phase transition from the CsCl structure to the BCC structure occurred, and the reaction process was similar to that of the binary alloy [2], which showed the formation of BCC TiFe0.8Mn0.2 hydride. For the alloy that was kept in the hydrogen fluid at 400 C and 6.19(4) GPa, the XRD profiles of the BCC hydride did not change at all over 200 min after the completion of the hydrogenation reaction. As we will describe later, the BCC hydride also existed stably when the pressure was changed at 400 C. Therefore, we conclude that the hydride was the most stable phase under the high pressureetemperature conditions, which indicates that a novel hydride with high hydrogen content was unlikely to form after formation of the BCC hydride under our experimental conditions. Kinaci and Aydinol reported that first-principles calculations predicted the possible formation of a tetragonal hydride with H/M ¼ 2.0 in the TiFe system [4]. However, we did not observe any change in the XRD profiles corresponding to the formation of a tetragonal or other structure after the BCC hydride became a single phase. The disproportionation reaction occurred at temperatures above 400 C, whereas the hydrogenation did not proceed completely at temperatures below 400 C. Therefore, the hydride with H/M ¼ 2.0 predicted by the theoretical study could not be synthesized at a pressure below 6.5 GPa and a temperature below 600 C. In order to synthesize the hydride, it might be necessary to carry out hydrogenation under higher pressure. Next, we tried to estimate the hydrogen content of the BCC hydride at high pressures and temperatures using in situ SRXRD measurements as follows. The unit cell volume of the BCC hydride was compared with that of the g hydride (H/ M ¼ 0.9) under the same pressureetemperature conditions. If the BCC hydride was larger, we could assume that its hydrogen content was larger than that of the g hydride because the unit cell volume of an interstitial hydride generally increases with increasing hydrogen content [5]. We began by examining the structure of the g hydride at high pressure. Fig. 2(c) shows selected in situ SR-XRD profiles around 6 GPa. The bottom pattern shows the XRD profile of the sample observed at 6.64(5) GPa and RT, which corresponds to a monoclinic structure. The middle pattern is the XRD profile of the g hydride observed at 6.43(3) GPa and 400 C. Although the 111 and 111 peak intensities near 58 keV changed after heating, each d value of the Bragg peaks was unchanged. No difference in the diffraction profile of the g hydride was observed after the sample was held under these conditions for 3 h, which indicates that the phase transition from the monoclinic structure to the BCC structure did not occur. A similar result was obtained when the sample was kept in hydrogen [top profile, Fig. 2(c)]. In the next step, we measured the unit cell volumes of both hydrides in a pressure range of 6.5e4.2 GPa during depressurization at a constant temperature of 400 C. While the g hydride desorbed hydrogen at 4.2 GPa, no hydrogen was desorbed from the BCC hydride. Fig. 3 shows a plot of the unit cell volumes of the BCC hydride and the g hydride observed at 400 C as functions of pressure. We obtained the unit cell volumes of the BCC hydride with hydrogen, while the g-
Fig. 3 e Unit cell volumes of BCC and g hydrides, observed at 400 C, plotted as functions of pressure. For comparison of the g hydrides, the plotted values of BCC hydride were two times of the unit cell volume.
hydride unit cell volumes were obtained without hydrogen. The volumes at each pressure were evaluated within an error ˚ 3. The estimated error of the hydrogen content would of 0.6 A be H/M 0.1, which is negligibly small. As seen in Fig. 3, almost all the volumes of the BCC hydride were within the error bars of the g-hydride volumes. It appears that there were no significant differences between both volumes at each pressure. The result was also the same for the g hydride that was kept in hydrogen fluid (g with H2, Fig. 3). We thus estimated roughly that the hydrogen content of the BCC hydride was equal to that of the g hydride (H/M ¼ 0.9). All O sites in the BCC hydride are crystallographically equivalent, and in addition, all T sites are also equivalent. In the case of the g hydride, the O sites are partially occupied by H atoms [3,9]. Although we had expected H atoms in the BCC hydride to occupy more interstitial sites than in the g hydride, the roughly-estimated hydrogen content of the former turned out to be the same as that of the latter (H/M ¼ 0.9). Finally, we considered a possibility of formation of the superabundant vacancies (SAVs) induced by hydrogen at high pressure and temperature [5,16]. SAVs are formed in metal lattices when metals or alloys are held under hydrogen fluid at several giga-pascals. In general, gradual lattice contractions are observed for several minutes (in BCC structures) or hours (in face-centered cubic (FCC) structures) during the formation of the SAVs. In the present study, however, the lattice parameter of the BCC hydride remained constant over several hours after lattice expansion, as described earlier. Therefore, we conclude that the formation of SAVs was negligible.
Conclusions The hydrogenation of the ternary TiFe0.8Mn0.2 alloy at high pressures and temperatures was investigated via in situ
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 4 0 ( 2 0 1 5 ) 3 2 8 3 e3 2 8 7
synchrotron radiation X-ray diffraction measurements. The chemical orderedisorder phase transition and the formation of BCC hydride were observed to be similar to those for the binary TiFe alloy. The phase-transition temperature was below that of the binary alloy and the disproportionation reaction was suppressed. Novel hydrides with high hydrogen content were not formed after the formation of the single BCC phase. The hydrogen content of the BCC hydride was roughly estimated to be equal to that of the g hydride.
[7]
[8]
[9]
[10]
Acknowledgments This work was partly supported by JSPS KAKENHI (Nos. 24760586 and 25420725) and by the Photon and Quantum Basic Research Coordinated Development Program from MEXT. The synchrotron radiation experiments were performed at beamline BL14B1 of SPring-8 with the approval of the Japan Atomic Energy Agency (Proposal Nos. 2012B3602, 2013A3603 and 2013B3603).
references
[1] Reilly JJ, Wiswall Jr RH. Formation and properties of iron titanium hydride. Inorg Chem 1974;13(1):218e22. [2] Endo N, Saitoh H, Machida A, Katayama Y. Formation of BCC TiFe hydride under high hydrogen pressure. Int J Hydrogen Energy 2013;38:6726e9. [3] Hydrogen in intermetallic compounds I: electronic, thermodynamic, and crystallographic properties, preparation. In: Schlapbach L, editor. Springer series, topics in applied physics, vol. 63. Berlin, Heidelberg: Springer; 1998. [4] Kinaci A, Aydinol MK. Ab initio investigation of FeTi-H system. Int J Hydrogen Energy 2007;32:2466e74. [5] The metal hydrogen system: basic bulk properties. In: Fukai Y, editor. Springer series in materials science. 2nd ed., vol. 21. Berlin, Heidelberg: Springer; 2005. Berlin. [6] Johnson JR, Reilly JJ. The use of manganese substituted ferrotitanium alloys for energy storage. In: Proceedings of the
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Miami international conference on alternative energy sources, December 5e7; 1997. p. 553e4 (Miami Beach, Florida, USA). Chung HD, Lee JY. Hydriding and dehydriding reaction rate of FeTi intermetallic compound. Int J Hydrogen Energy 1985;10:537e42. Miller HI, Murray J, Laury E, Reinhardt J, Goudy AJ. The hydriding and dehydring kinetics of FeTi and Fe0.9TiMn0.1. J Alloys Compd 1995;231:670e4. € fer W, Will G, Schober T. Neutron and electron Scha diffraction of the FeTi-D(H)-g-phase. Mat Res Bull 1980;15:627e34. Reidinger F, Lynch JF, Reilly JJ. An X-ray diffraction examination of the FeTieH2 system. J Phys F e Met Phys 1982;12:L49e55. M, Sauvage D, Rouault A, Fruchart D, Commandre Tellgren R. Structural and activation process studies of Fe-Ti-like hydride compounds. J Less-Common met 1980;74:55e63. Fischer P, Schefer J, Yvon K, Schlapbach L, Riesterer T. Orthorhombic structure of g-TiFeDz2. J Less-Common Met 1987;129:39e45. Utsumi W, Funakoshi K, Katayama Y, Yamakata M, Okada T, Shimomura O. High-pressure science with a multi-anvil apparatus at SPring-8. J Phys Condens Matter 2002;14:10497e504. Saitoh H, Machida A, Aoki K. Synchrotron X-ray diffraction techniques for in situ measurement of hydride formation under several gigapascals of hydrogen pressure. Chin Sci Bull 2014;59:5290e301. Saitoh H, Machida A, Katayama Y, Aoki K. Formation and decomposition of AlH3 in the aluminum-hydrogen system. Appl Phys Lett 2008;93:15918e23. Fukai Y, Okuma N. Evidence of copious vacancy formation in Ni and Pd under a high hydrogen pressure. Jpn J Appl Phys 1993;32:L1256e9. Kamegawa A, Goto Y, Kataoka R, Takamura H, Okada M. High-pressure synthesis of novel compounds in an MgeNi system. Renew Energy 2008;33:221e5. Decker DL. High-pressure equation of state for NaCl, KCl, and CsCl. J Appl Phys 1971;42:3239e44. Endo N, Saitoh H, Machida A, Katayama Y, Aoki K. Phase diagram and equation of state of TiH2 at high pressures and high temperatures. J Alloy Compd 2013;546:270e4.
Available online at www.sciencedirect.com
ScienceDirect journal homepage: www.elsevier.com/locate/he
Short Communication
Hydrogenation of a TiFe-based alloy at high pressures and temperatures Naruki Endo a,*, Itoko Saita b, Yumiko Nakamura b, Hiroyuki Saitoh c, Akihiko Machida c a
Renewable Energy Research Center, National Institute of Advanced Industrial Science and Technology (AIST), 2-2-9 Machiike-dai, Koriyama, Fukushima 963-0215, Japan b Energy Technology Research Institute, National Institute of Advanced Industrial Science and Technology (AIST), AIST Central-5, 1-1-1, Higashi, Tsukuba, Ibaraki 305-8565, Japan c Quantum Beam Science Center, Japan Atomic Energy Agency, 1-1-1 Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5148, Japan
article info
abstract
Article history:
We investigated the hydrogenation of a ternary TiFe0.8Mn0.2 alloy under high pressure (up
Received 19 November 2014
to 6.5 GPa) and at high temperature (up to 600 C) by in situ synchrotron radiation X-ray
Received in revised form
diffraction measurements. A chemical orderedisorder phase transition and the formation
30 December 2014
of body-centered cubic (BCC) hydride were observed. The phase-transition temperature
Accepted 3 January 2015
was 400 C, which is 200 C lower than that of a binary TiFe alloy. A disproportionation
Available online 29 January 2015
reaction, which was observed for the binary TiFe alloy, did not occur at this temperature. The BCC hydride existed stably, and new hydrides were not formed after the BCC hydride
Keywords:
became a single phase. Since the unit cell volume of the BCC hydride was almost equal to
TiFe-based alloy
that of the g hydride, TiFe0.8Mn0.2H1.8, under the same pressureetemperature conditions,
Hydrogen storage
the hydrogen content of the former was roughly estimated as the same as that of the latter
High hydrogen pressure
(hydrogen to metal atom ratio ¼ 0.9).
Body-centered cubic
Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
In situ X-ray diffraction measurement
Introduction High hydrogen capacity is one of the most important property for hydrogen storage materials. We have been investigating the hydrogenation of CsCl-type TiFe alloys [1] under high hydrogen pressure to realize a novel hydride with high hydrogen content. In a recent study, we reported that when a
binary TiFe alloy was hydrogenated at 5 GPa and 600 C for 150 min, a chemical orderedisorder phase transition occurred, after which a BCC hydride was formed [2]. The unit cell volume of the hydride expanded by 21.0%, which is larger than that of the g hydride, TiFeH1.9, prepared under nearambient conditions (~18%) [3]. Since many tetrahedral (T) and octahedral (O) interstitial sites were still vacant in the
* Corresponding author. Tel.: þ81 29 861 2647; fax: þ81 24 963 0828. E-mail address: [email protected] (N. Endo). http://dx.doi.org/10.1016/j.ijhydene.2015.01.015 0360-3199/Copyright © 2015, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
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metal lattice of the BCC hydride, we believe that the hydride might have been able to absorb more hydrogen. As recently reported, results of first-principle calculations suggested that it is possible to form a novel TiFe hydride with high hydrogen content [4], a tetragonal structure (P4/mmm; Z ¼ 1), and a hydrogen atom (H) to metal atom (M) ratio of H/ M ¼ 2.0. The hydride was calculated to be unstable under near-ambient conditions. The calculations also pointed out that some alterations in the composition of the alloy or in the reaction pressure and temperature are necessary to synthesize the hydride. Since it has been reported that hydrogenation below 100 C and several tens of MPa does not increase the hydrogen content by Ref. [5], the hydride with H/M ¼ 2.0 is expected to be formed by hydrogenation under high pressure (>1 GPa) and temperature (>100 C). In our earlier study in which the TiFe alloy was hydrogenated at 5 GPa and 600 C, a disproportionation reaction occurred and the BCC hydride immediately decomposed after the formation of the single BCC phase [2]. It is thought that if the disproportionation reaction had been suppressed, the TiFe hydride with H/M ¼ 2.0 or other novel TiFe hydrides with high hydrogen content could have been formed by further hydrogenation of the BCC hydride. The disproportionation reaction occurred because a high temperature was necessary to hydrogenate the TiFe alloy. However, the temperature of 600 C did not represent the thermodynamic condition of hydrogenationdthe kinetic effects of the surface oxides raised the reaction temperature, which was the temperature at which the sample was completely hydrogenated [2]. As the kinetics effects cannot be avoided when the binary alloy is used, a more reactive sample should be chosen for the synthesis of novel hydrides. In the present study, we chose a ternary TiFe0.8Mn0.2 alloy because it is highly reactive for hydrogenation comparable to LaNi5 [6,7]. Previous studies have reported that substituting Fe for Mn improves the activation properties of the TiFe alloy and the kinetics of absorption and desorption of the ternary alloy is faster than that of the binary alloy [6e8]. We investigated the hydrogenation of the TiFe0.8Mn0.2 alloy at high pressures and temperatures to verify the formation of novel hydrides with high hydrogen content. From the investigation of the hydrogenation processes by in situ synchrotron radiation (SR) X-ray diffraction (XRD) measurements, we confirmed that they were similar to those of the binary alloy: a
chemical orderedisorder phase transition and the formation of a BCC hydride were observed. The hydrogenation temperature was lower than that of the binary TiFe alloy, and the disproportionation reaction was suppressed. We estimated the hydrogen content of the BCC hydride from a comparison of the unit cell volumes of the BCC and g hydrides.
Experimental The starting materials were TiFe0.8Mn0.2 alloy and a g hydride, TiFe0.8Mn0.2H1.8 (H/M ¼ 0.9) powder. The particles were less than 100 mm in size. The lattice parameters of the alloy and the ˚ , and hydride before the experiments were a ¼ 2.986(1) A ˚ , b ¼ 2.8338(3) A ˚ , c ¼ 4.719(3) A ˚ , and b ¼ 97.31(2), a ¼ 4.7024(3) A which are slightly larger than those of TiFe and g-phase TiFeH1.9 [9,10] because of the Mn substitution [11]. The crystal structure of the g hydride has been reported as monoclinic (P2/m; Z ¼ 2) [9e11] and orthorhombic (Cmmm; Z ¼ 4) [3,12]. In the present study, we nominally chose the monoclinic structure to represent the g hydride since the same results were obtained from the monoclinic and the orthorhombic structures. High pressureetemperature conditions were generated with a cubic-type multi-anvil high-pressure apparatus installed in the beamline BL14B1 at SPring-8, Japan [13]. Powder XRD profiles of the samples were obtained in situ by using high-brilliance SR X-ray as the incident beam. Here, the incident beam was white X-ray from SR light source and the energy-dispersive XRD method was employed [14], which is the standard technique for high-pressure experiments using a multi-anvil high-pressure press [15]. The diffracted X-rays were measured by a Ge solid state detector mounted on a goniometer. The diffraction angle 2q was fixed to 6 in all ˚ ) was calculated from measurements. The lattice spacing d (A the energies of the diffracted X-rays E (keV) based on Bragg's equation: 2d sin q ¼ 12:4=E
(1)
The measurement time was 1.0 min. Fig. 1 shows the schematics of the high-pressure cell assemblies with and without hydrogen sources. In both highpressure cell assemblies, the powder samples were filled in pyrolytic boron nitride (PBN) capsules. The high-pressure cell
Fig. 1 e Schematic of the high-pressure cell assembly used for the present study: (a) with internal hydrogen sources; (b) without internal hydrogen sources.
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assembly for hydrogenation was firstly developed by Fukai and Okuma [16] and has been used by other groups for highpressure synthesis of metal hydrides [2,14,15,17]. We modified the high-pressure cell assembly to determine the high hydrogen pressure. Fig. 1(a) shows the high-pressure cell assembly for hydrogenation used in the present study. NaCl powder, which was mixed homogeneously with BN powder to prevent grain growth at high temperature, was used as a pressure marker. The pressure was calculated from the unit cell volume of the pressure marker based on Decker's equation of state [18]. The hydrogenation reactions were carried out as follows: The samples were pressurized to ~6 GPa at room temperature (RT). Next, the sample and the hydrogen sources were initially heated to 600 C at a rate of 100 C/min to evolve hydrogen by decomposition of the hydrogen sources [16]. Upon reaching 600 C, the sample was immediately cooled to 300 C for 0.1 min. The reaction temperatures were set in the range of 300e600 C. In the case of the reaction without the hydrogen sources [Fig. 1(b)], the temperature was raised to a prescribed temperature of 300e600 C at a heating rate of 100 C/min and held at the conditions. The unit cell volumes of the g hydride before and after the experiments were the same, indicating that hydrogen was not released during the experiments. Details of the sample preparations, the experimental apparatus, and experimental procedures are described elsewhere [6,14,19].
Results and discussion The hydrogenation processes of the TiFe0.8Mn0.2 alloy at 6.3(2) GPa were investigated in the temperature range of 300e600 C. The hydrogenation reaction did not proceed completely at 300 C, while the entire hydrogenation proceeded at 400 C without the disproportionation reaction.
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When the hydrogenation reaction took place at temperatures above 500 C, the disproportionation reaction occurred immediately after the formation of the single BCC phase; the hydrogen absorption process was completed within 10 min at 600 C (cf. 150 min for TiFe [2]). Fig. 2(a) shows selected in situ SR-XRD profiles of the TiFe0.8Mn0.2 alloy at around 6.3(2) GPa. The bottom pattern represents the XRD profile of the alloy after pressurizing at RT. Hydrogen was not emitted from the internal hydrogen sources at this temperature. The diffraction peaks were slightly broadened by strain, which was induced during the pressurizing process. The 110, 200, and 220 peaks of the alloy phase shifted toward the low energy side just after the alloy was kept in the hydrogen fluid at 400 C. The lattice parameter of the ˚ , which sample after it was heated to 400 C was a ¼ 2.975(2) A ˚ is larger than that of the alloy (a ¼ 2.934(3) A) under the same conditions without hydrogen. The alloy was hydrogenated at this stage and a solid solution with hydrogen was formed. After the sample had been held in the hydrogen fluid for 30 min, shoulders appeared on the lower energy side of the diffraction peaks from the solid-solution phase, indicating that a new hydride phase was formed. The peaks from the hydride were indexed by a cubic unit cell. After 170 min of the hydrogenation treatment, the Bragg peaks from the solid solution vanished and the hydride phase became a single phase. A continuous lattice expansion of the hydride occurred and ˚ after 530 min the lattice parameter increased to a ¼ 3.116(4) A of the hydrogenation treatment. Fig. 2(b) shows an enlargement of the XRD profiles in Fig. 2(a). The small peaks correspond to the Bragg peaks from NaCl and the C15 Laves phase [7]. The arrows indicate the 111 and 210 peaks from the CsCl-type structure, whose intensities decreased with increasing volume of the hydride during hydrogenation; these peaks were hardly observable when the hydride became a single phase [top profile, Fig. 2(b)]. These peaks represent the fundamental reflections from the CsCl
Fig. 2 e In situ SR-XRD patterns of (a) the TiFe0.8Mn0.2 alloy with hydrogen and (c) the g hydride, TiFe0.8Mn0.2H1.8, with and without hydrogen at high pressures. (b) Enlargement of the XRD profiles in (a). The arrows indicate the 111 and 210 peaks from the CsCl-type structure.
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structure, but such reflections are forbidden from the BCC structure. This means that a chemical orderedisorder phase transition from the CsCl structure to the BCC structure occurred, and the reaction process was similar to that of the binary alloy [2], which showed the formation of BCC TiFe0.8Mn0.2 hydride. For the alloy that was kept in the hydrogen fluid at 400 C and 6.19(4) GPa, the XRD profiles of the BCC hydride did not change at all over 200 min after the completion of the hydrogenation reaction. As we will describe later, the BCC hydride also existed stably when the pressure was changed at 400 C. Therefore, we conclude that the hydride was the most stable phase under the high pressureetemperature conditions, which indicates that a novel hydride with high hydrogen content was unlikely to form after formation of the BCC hydride under our experimental conditions. Kinaci and Aydinol reported that first-principles calculations predicted the possible formation of a tetragonal hydride with H/M ¼ 2.0 in the TiFe system [4]. However, we did not observe any change in the XRD profiles corresponding to the formation of a tetragonal or other structure after the BCC hydride became a single phase. The disproportionation reaction occurred at temperatures above 400 C, whereas the hydrogenation did not proceed completely at temperatures below 400 C. Therefore, the hydride with H/M ¼ 2.0 predicted by the theoretical study could not be synthesized at a pressure below 6.5 GPa and a temperature below 600 C. In order to synthesize the hydride, it might be necessary to carry out hydrogenation under higher pressure. Next, we tried to estimate the hydrogen content of the BCC hydride at high pressures and temperatures using in situ SRXRD measurements as follows. The unit cell volume of the BCC hydride was compared with that of the g hydride (H/ M ¼ 0.9) under the same pressureetemperature conditions. If the BCC hydride was larger, we could assume that its hydrogen content was larger than that of the g hydride because the unit cell volume of an interstitial hydride generally increases with increasing hydrogen content [5]. We began by examining the structure of the g hydride at high pressure. Fig. 2(c) shows selected in situ SR-XRD profiles around 6 GPa. The bottom pattern shows the XRD profile of the sample observed at 6.64(5) GPa and RT, which corresponds to a monoclinic structure. The middle pattern is the XRD profile of the g hydride observed at 6.43(3) GPa and 400 C. Although the 111 and 111 peak intensities near 58 keV changed after heating, each d value of the Bragg peaks was unchanged. No difference in the diffraction profile of the g hydride was observed after the sample was held under these conditions for 3 h, which indicates that the phase transition from the monoclinic structure to the BCC structure did not occur. A similar result was obtained when the sample was kept in hydrogen [top profile, Fig. 2(c)]. In the next step, we measured the unit cell volumes of both hydrides in a pressure range of 6.5e4.2 GPa during depressurization at a constant temperature of 400 C. While the g hydride desorbed hydrogen at 4.2 GPa, no hydrogen was desorbed from the BCC hydride. Fig. 3 shows a plot of the unit cell volumes of the BCC hydride and the g hydride observed at 400 C as functions of pressure. We obtained the unit cell volumes of the BCC hydride with hydrogen, while the g-
Fig. 3 e Unit cell volumes of BCC and g hydrides, observed at 400 C, plotted as functions of pressure. For comparison of the g hydrides, the plotted values of BCC hydride were two times of the unit cell volume.
hydride unit cell volumes were obtained without hydrogen. The volumes at each pressure were evaluated within an error ˚ 3. The estimated error of the hydrogen content would of 0.6 A be H/M 0.1, which is negligibly small. As seen in Fig. 3, almost all the volumes of the BCC hydride were within the error bars of the g-hydride volumes. It appears that there were no significant differences between both volumes at each pressure. The result was also the same for the g hydride that was kept in hydrogen fluid (g with H2, Fig. 3). We thus estimated roughly that the hydrogen content of the BCC hydride was equal to that of the g hydride (H/M ¼ 0.9). All O sites in the BCC hydride are crystallographically equivalent, and in addition, all T sites are also equivalent. In the case of the g hydride, the O sites are partially occupied by H atoms [3,9]. Although we had expected H atoms in the BCC hydride to occupy more interstitial sites than in the g hydride, the roughly-estimated hydrogen content of the former turned out to be the same as that of the latter (H/M ¼ 0.9). Finally, we considered a possibility of formation of the superabundant vacancies (SAVs) induced by hydrogen at high pressure and temperature [5,16]. SAVs are formed in metal lattices when metals or alloys are held under hydrogen fluid at several giga-pascals. In general, gradual lattice contractions are observed for several minutes (in BCC structures) or hours (in face-centered cubic (FCC) structures) during the formation of the SAVs. In the present study, however, the lattice parameter of the BCC hydride remained constant over several hours after lattice expansion, as described earlier. Therefore, we conclude that the formation of SAVs was negligible.
Conclusions The hydrogenation of the ternary TiFe0.8Mn0.2 alloy at high pressures and temperatures was investigated via in situ
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synchrotron radiation X-ray diffraction measurements. The chemical orderedisorder phase transition and the formation of BCC hydride were observed to be similar to those for the binary TiFe alloy. The phase-transition temperature was below that of the binary alloy and the disproportionation reaction was suppressed. Novel hydrides with high hydrogen content were not formed after the formation of the single BCC phase. The hydrogen content of the BCC hydride was roughly estimated to be equal to that of the g hydride.
[7]
[8]
[9]
[10]
Acknowledgments This work was partly supported by JSPS KAKENHI (Nos. 24760586 and 25420725) and by the Photon and Quantum Basic Research Coordinated Development Program from MEXT. The synchrotron radiation experiments were performed at beamline BL14B1 of SPring-8 with the approval of the Japan Atomic Energy Agency (Proposal Nos. 2012B3602, 2013A3603 and 2013B3603).
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