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Influence of boron on the hydriding of nanocrystalline Mg2Ni
Intermetallics 34 (2013) 63e68
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Influence of boron on the hydriding of nanocrystalline Mg2Ni Murad Redzeb, Zlatina Zlatanova, Tony Spassov* Department of Chemistry and Pharmacy, University of Sofia “St.Kl.Ohridski”, 1 J. Bourchier Str., 1164 Sofia, Bulgaria
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 August 2012 Received in revised form 20 October 2012 Accepted 23 October 2012 Available online 7 December 2012
Nanocrystalline Mg2Ni(B) alloys (Mg64Ni32B4 and Mg58Ni29B13) were synthesized by ball milling and their hydrogen sorption properties were characterized by three different methods: Sievert’s type apparatus, High-pressure DSC and by electrochemical hydrogen charge/discharge. The main idea was to elucidate the influence of boron on the hydriding behavior of Mg2Ni, an intermetallic compound largely investigated for hydrogen storage, but still needing improvement. Boron was selected because of its small atomic radius and large electronegativity, as well as chemical resistance in solutions of alkaline basics. Hydrogen capacity and hydriding/dehydriding cycle life of the alloys were found to depend on the presence of B. Whereas small amounts of boron (4 at.%) increase the maximum hydrogen capacity of Mg2Ni, at higher B concentration (13 at.%) the capacity and cycle life deteriorate significantly due to the Mg2Ni(B) solid solution decomposition during hydriding/dehydriding. Rapid worsening was found to take place during the first 2e3 electrochemical charge/discharge cycles to a much larger extent for Mg2Ni(B), compared to pure Mg2Ni. Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: A. Nanostructured intermetallics B. Hydrogen storage C. Mechanical alloying and milling
1. Introduction There is considerable interest in the use of metal hydrides as rechargeable hydrogen storage media. This interest stems from the fact that this is the safest means of storing hydrogen (compared to hydrogen liquefaction and hydrogen storage at high pressure), with a large hydrogen storage capacity, both in mass and in volume, and a range of absorption/desorption temperatures and pressures that can fit numerous applications. The binary MgeNi system is unique from the standpoint of hydrogen storage, as it contains two hydrogen absorption phases: a hexagonal (hp2) Mg metal phase and hexagonal (hp18) intermetallic line compound, Mg2Ni, whose reported maximum gravimetric hydrogen capacities are 7.66 and 3.62 wt.%, respectively. Mg2Ni is the most intensively studied due to its high hydrogen absorption capacity, low specific weight and low cost [1e21]. The Mg absorption kinetics is improved by the Ni addition, because of its good catalytic activity [1]. However, under normal conditions (i.e. at room temperature and atmospheric pressure) Mg2Ni does not absorb hydrogen. The usual hydrogenation conditions are: 250e350 C and hydrogen pressure of 15e 50 bar [2,3]. Considering the high vapor pressure and generally small solubility of alloying elements in Mg alloys, MgeNi alloys can be produced more easily and reliably by mechanical alloying (MA)
* Corresponding author. E-mail address: [email protected]fia.bg (T. Spassov). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.10.016
than by conventional melting methods [4]. It is suggested that the heat generated due to collision of the balls resulting in continuous fragmentation, coalescence events and the interdiffusion of ingredients of the alloy at the collision sites are the main steps of the process of mechanical alloying [2]. Another advantage of MA is the formation of nanocrystalline Mg2Ni with very active surfaces of the powder and grain sizes of about 20e30 nm which absorbs hydrogen at lower temperature [3,4]. It has good activation properties and shows a large discharge capacity at room temperature. Nevertheless, it loses 80% of its maximum capacity within ten cycles. The main cause of the rapid degradation of the electrode is the formation of the passive layer at the surface, which hinders the charge transfer reaction. Recently, many efforts have been made to improve Mg-based alloys’ electrochemical characteristics. In particular, element substitution is proved to be very effective method for improving the high discharge capacity. Platinum group metals (PGMs) have been employed to superficially modify Mg2Ni-based compounds. Small amount of Pd catalyst enhances hydrogen absorption kinetics at 200 C. Nanocrystalline Mg2Ni with Pd absorbs hydrogen even at room temperature, without activation and with relatively good kinetics [3]. Reactively milled Mg2Ni co-milled with Ru shows an onset temperature of hydrogen desorption as low as 80 C [5]. Ball-milling Mg2Ni with metallic Ni (70 wt.% vs. Mg2Ni) leads to the formation of a homogeneous amorphous alloy, which exhibits a maximum discharge capacity of ca. 870 mA h g1 (Mg2Ni) at 30 C [6]. The ternary hydrogen-storage alloys Mg2Ni1xZrx (0 < x 0.3),
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compared to a Mg2Ni alloy, have a larger specific surface (z1.20 m2/g), more promising dehydriding kinetics, lower enthalpy of formation of hydrides than that of Mg2NiH4, and lower decomposition temperatures in an open system. An optimum desorption storage capacity of about 3.3 wt.% is observed [7]. The discharge capacity of a (Mg1xZrx)2Ni electrode increases with Zr content and reaches the highest capacity of 530 mA h g1 at x ¼ 0.3, then decreases to 230 mA h g1 at x ¼ 0.4, i.e., it is lower than that of pure Mg2Ni. Cyclic stability and rate capability vary with Zr addition, but they are both not greatly improved [8]. The nanocrystalline Mg1.9Ti0.1Ni shows better kinetics compared to nanocrystalline Mg2Ni, absorbs more than 3 wt.% H2 in 2000 s at 423 K and also destabilizes the hydride phase slightly compared to nanocrystalline Mg2Ni [9]. Graphite addition inhibits the formation of a new oxide layer on the surface of the materials once the native oxide layer is broken during the milling process and has a positive effect on the absorption kinetics [10,11]. The hydriding kinetics of Mg2Ni is improved by adding Ag [12]. Few studies have focused, however, on the evaluation of boron addition to hydrogen storage alloys. Anik et al. have concluded that Mg0.9(B)0.1Ni and Mg0.8(B)0.2Ni alloys don’t possess better properties compared to Mg50Ni50 [22]. More extensive research related to the boron influence shows that small amount of boron added to Mg50Ni50 alloy improves the hydrogen storage properties [23]. Therefore, in the present work boron was chosen as an alloying element in the nanostructured Mg2Ni (Mg64Ni32B4 and Mg58Ni29B13), because of its non-metal nature, chemical resistance in solutions of alkaline solutions, small atomic radius (thus forming interstitial solid solutions in Mg2Ni) and large electronegativity e 2.04 (Pauling). In this way, greater changes in the electronic structure of Mg2Ni compared to the metal substituted compositions can be expected, which might result in different hydriding properties as well. 2. Experimental details Mg2Ni(B) nanocrystalline alloys (Mg58Ni29B13 and Mg64Ni32B4) were synthesized by ball milling in a planetary type ball mill (Fritsch P6) under pure argon gas atmosphere with a rotation speed of 450 rpm for 15 h, ball to powder mass ratio (B/P) of 12/1, with 10 min rest after each 1 h of milling. Pure elemental powders of magnesium, nickel, and boron (amorphous) were used as starting materials. Stainless steel vials and balls were used. The average composition of the boron containing alloys was determined using EDS analyzer and was found to correspond to their nominal composition. The samples were handled in a glove box under Ar. Small amount of the powder was taken from the mill at certain periods of time for structural, morphological and thermal analysis. The powder morphology and microstructure of the alloys were analyzed by x-ray diffraction (XRD) using a Cu-Ka radiation and scanning electron microscopy (JEOL 5510 SEM). Possible phase transformations during annealing of the as-milled materials and the hydriding/dehydriding processes were studied by high-pressure differential scanning calorimetry (HPDSC), SETARAM Sensys Evo TG-DSC. The hydrogen sorption behavior (capacity and kinetics) was investigated using Sievert’s type (PCT) apparatus with samples from 100 to 150 mg. The electrochemical measurements were realized by a three electrode cell with Hg/HgO as a reference electrode and a counter electrode prepared from Ni mesh. The working electrodes were made by mixing the alloy powder (100 mg) with 30 mg teflonyzed carbon black (VULCAN 72 10%PTFE) and pressing the as-prepared mixture with a pressure of 15 MPa. In the charge/discharge cycles each electrode was charged for 5e10 h at 5 mA/cm2 and discharged to e 500 mV vs. Hg/HgO at 2 mA/cm2 in a 6 M KOH water solution.
3. Results and discussion 3.1. Morphology and microstructure of the as-milled alloys The x-ray diffraction analysis shows that alloying with even 13 at.% boron does not result in a new phase formation and due to the small atomic radius of B the positions of the diffractions peaks do not shift noticeably, Fig. 1. It can be seen that after sufficient time of milling (15 h) the powders reveal nanocrystalline microstructure, consisting of the hexagonal Mg2Ni phase. For both boron containing alloys, Mg64Ni32B4 and Mg58Ni29B13, an average size of the nanocrystallites of 10.0 0.2 nm was determined. The crystal lattice parameters show some small differences compared A and c ¼ 13.98 A), revealing large to the pure Mg2Ni (a ¼ 5.85 solid state solubility of boron into the nanocrystalline hexagonal Mg2Ni. SEM observations reveal that after continuous milling the powder particles are in the range of 1e10 mm and are relatively homogeneous in size, Fig. 2. Some smaller particles (<1 mm), cold welded together, can also be seen. 3.2. Hydrogen sorption properties The hydriding behavior of the as-milled Mg2Ni(B) alloys was analyzed by high-pressure DSC under hydrogen pressure of 15 bar, Fig. 3. In the same figure the hydriding curves for the pure nanocrystalline Mg2Ni is also plotted. From the thermograms it is evident that the nanocrystalline Mg2Ni(B) alloys absorb hydrogen already at about 150 C, proved by XRD analysis of samples, hydrided in the HPDSC at low temperature of 200 C (Fig. 1). The hydrided alloy reveals (Fig. 1) only the Mg2NiH4 phase with lattice parameters a ¼ 14,324 A and c ¼ 2701 A. Comparing the initial temperature of hydriding between the three nanocrystalline alloys it is seen that the boron containing alloys start to absorb hydrogen at slightly lower temperature than the pure Mg2Ni, as for Mg64Ni32B4 the hydriding process starts at about 130 C. For pure nanocrystalline Mg2Ni the hydrogen sorption takes place in the range 170e220 C, which is noticeably narrower temperature range than those for the boron alloys. The shape of the H-absorption peak shows that larger distribution of hydrogen sites has been observed for the B-containing alloys, especially for Mg64Ni32B4. Substantially smaller hydriding peak at about 300e330 C is also observed for the intermetallic compound
Fig. 1. XRD patterns of the ball milled Mg2Ni(B) alloys.
M. Redzeb et al. / Intermetallics 34 (2013) 63e68
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Fig. 2. SEM micrograph of Mg2Ni(B) alloys: Mg64Ni32B4 (a, b) and Mg58Ni29B13 (c, d).
Mg64Ni32B4, which can be due to some microstructural inhomogeneity after ball milling, resulting in the formation of hydrogen sites with different energy. This is further confirmed by the observed two endothermic peaks of hydride(s) decomposition, taking place in the temperature range 400e450 C, which however transforms to a single-peak hydrogen absorption effect during the second hydriding experiment (see Fig. 4). It is obtained that the enthalpy change of Mg58Ni29B13 hydriding (DHhydr, [J/g]), Fig. 3, is larger than those of the other two materials, meaning that the first hydriding reaction of this alloy is connected with larger amount of the formed hydride. The multi-peaks exothermic effects, observed for the boron containing nanocrystalline compounds (Fig. 3), obviously have to be associated with the occupation of different sites by the hydrogen atoms, incl. for example inner volume of the nanocrystals and intercrystalline regions. The latter microcrystalline component is rather large in the materials studied due to the extremely fine nanocrystals formed during the milling. After the complete alloys hydriding only hydrides of Mg2Ni could be found (Fig. 1), which means that the different exothermic peaks have to be mostly associated with hydrogen occupation of different sites, rather than hydrogenation of different phases. At this stage it can be resumed that the boron containing Mg2Nibased alloys reveal somehow improved hydriding behavior, compared to pure Mg2Ni, expressed in small increase of the amount of absorbed hydrogen (assessed by the area of the exothermic peaks of hydriding) and certain lowering of the initial temperature of the hydriding process, accompanied by an increase of the temperature range of hydrogen absorption. The hydrided alloys undergo hydrides decomposition at about 400e420 C when further annealed in the DSC under hydrogen pressure of 15 bars, Fig. 3. Single peak endothermic effect
connected with Mg2NiH4 decomposition is seen for Mg2Ni and Mg58Ni29B13 and two overlapped peaks, associated with the decomposition of different hydrides and/or release of hydrogen occupying sites with different energy, characterize the Mg64Ni32B4 alloy. It is important to be mentioned that the dehydrogenation of the alloy with lower B content takes place at about 50 C lower temperature compared to that of the pure Mg2Ni alloy and Mg58Ni29B13. This result reveals some weakening of the metale hydrogen interaction, caused by the addition of small amount of B to Mg2Ni. Such improvement of the thermodynamics of the Mg64Ni32B4 hydriding/dehydriding has been observed at all hydriding/dehydriding experiments in the present study. During subsequent cooling of the dehydrided materials under hydrogen atmosphere from 500 C down to room temperature (Fig. 3b) they undergo again hydriding at about 380e400 C, as the enthalpy of this process corresponds exactly to that of dehydriding. The thermal peaks this time are well shaped and narrow due to the fast hydriding reaction at these temperatures. At temperatures of about 220 C small exothermic effects are also observed. These effects correspond to the transition of the high temperature (fcc) to the low temperature (monoclinic) Mg2NiH4, described comprehensively by Hayakawa et al. [24]. While the first hydriding reaction for the Mg58Ni29B13 reveals definitely higher enthalpy change compared to pure Mg2Ni and the alloy with lower B content and has to be associated with larger amount of the formed hydride, during the second hydriding/dehydriding cycle Mg58Ni29B13 deteriorates dramatically. Practically, this alloy shows lack of any stability against hydriding/dehydriding cycling, Fig. 4, due to decomposition of the alloy during hydriding/dehydriding. Magnesium nickel borides were detected after repeated hydriding of the alloy. This is not the case for the Mg64Ni32B4 alloy, where we observe repetition of the enthalpy values measured during the
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a
Dehydriding (0.1 bar H2) 50 mW
exo >
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Mg64Ni32B4
Heat flow
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Mg2Ni
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Temperature, C Fig. 5. Decomposition (0.1 bar and 5 K/min) of the hydrided Mg2Ni(B) alloys.
50 mW
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Mg2Ni cooling
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Temperature, C Fig. 3. HPDSC plots of Mg2Ni(B) alloys at hydrogen pressure of 15 bar (heating (a) and cooling (b) with a rate of 5 K/min).
first hydriding/dehydriding cycle, which is also valid for the pure Mg2Ni [25]. The hydrided materials were also subjected to hydrides decomposition by annealing in the DSC with a constant heating rate under low vacuum (0.1 bar H2), Fig. 5. The intermetallic compounds reveal also clear endothermic peaks, associated with a decomposition of the hydride phases. The composites dehydriding is associated with large and well-defined endothermic peaks at about 200e220 C. The enthalpies of dehydriding under vacuum agree very well with those determined from the DSC experiments under hydrogen pressure for the corresponding materials. Hydriding kinetic data for the nanocrystalline Mg2Ni-based powders obtained by Sievert’s type apparatus at 200 C are presented in Fig. 6. The nanocrystalline powders absorb hydrogen without activation, as the maximum capacity is reached during the first absorptionedesorption cycle. The maximum capacity of the powders is about 3.0 wt.%, as for Mg64Ni32B4 the hydrogen absorption process starts before its registration at 200 C, due to the higher rate of hydriding at this temperature, compared to Mg58Ni29B13. Mg58Ni29B13 shows noticeable deterioration in capacity and kinetics during the second hydriding experiment, which was not observed for the Mg64Ni32B4 alloy, possessing
Hydriding/dehydriding (second run) 13,3
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Hydrogen absorption o 200 C
Mg58Ni29B13
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Mg64Ni32B4
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Temperature, C Fig. 4. Second hydriding/dehydriding HPDSC cycle.
500
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time, s Fig. 6. PeC isotherms at 200 C for the Mg2Ni(B) alloys.
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Mg58Ni29B13
M. Redzeb et al. / Intermetallics 34 (2013) 63e68
reproducible hydrogen sorption behavior. This fact confirms quite well the results obtained by HPDSC. In general, the initial hydrogen sorption kinetics is fast for the alloys studied, particularly for this relatively low temperature for Mg2Ni-based alloys (200 C), reaching about 50% of hydriding
a
Mg58Ni29B13
1000
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within the first 500e600 s. The initial absorption rate obtained from the slope of the fitted straight line on the sorption curves in the range of 100e500 s indicates a favorable kinetics for the boron richer alloy, but as already mentioned its hydriding kinetics becomes worse immediately after the first hydriding cycle. This effect was already explained by drastic microstructural changes (phase decomposition of the B richer alloy) observed by XRD, result of the hydriding/dehydriding process. Obviously the solid solution of boron into the hexagonal Mg2Ni intermetallics at low B content (Mg64Ni32B4) is stable and therefore shows reproducible hydriding behavior.
U, mV (vs. Hg/HgO)
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The observed in some way superior low-temperature hydriding behavior of the boron containing alloys (esp. Mg64Ni32B4) compared to pure Mg2Ni gives some reasons to expect improved electrochemical hydriding/dehydriding performance of these materials as well. Charge/discharge curves at galvanostatic conditions were presented in Fig. 7 for both alloys. The B containing Mg2Ni alloys show however worse electrochemical hydrogen storage properties in comparison with the pure intermetallic compound [26]. The cycle life of the electrodes is presented in Fig. 7c. It can be concluded that the amount of absorbed hydrogen decreases immediately after the first cycle. Both alloys show comparable corrosion stability, substantially lower than that of pure Mg2Ni [26]. Although the alloy with lower B content reveals slightly better electrochemical hydrogen charge/discharge behavior compared to the B richer alloy the discharge capacity and cycle life as a whole are rather low and unsatisfactory. Obviously, alloying with B deteriorates drastically the discharge capacity of Mg2Ni, accelerating the oxidation of magnesium on the alloy surface, forming Mg(OH)2.
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4. Conclusion 600
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Nanocrystalline solid solutions of boron into hexagonal Mg2Ni with overall composition of Mg64Ni32B4 and Mg58Ni29B13 were synthesized by ball milling and their hydrogen sorption properties were characterized. The main idea was to reveal the influence of boron, a non-metal element with small atomic radius and large electronegativity as well as chemically resistant in solutions of alkaline basics on the hydriding behavior of Mg2Ni. Relatively small amount of boron (4 at.% B) generally slightly increases the gas-phase hydrogen sorption capacity of Mg2Ni and improves the hydriding thermodynamics and kinetics. Additional increase of the B content (13 at.% B), however, results in hydriding properties deterioration due to boron solid solution decomposition during hydriding/dehydriding. In respect to the hydriding/dehydriding cycling stability in hydrogen gas phase Mg64Ni32B4 is comparable to pure Mg2Ni. Both boron containing Mg2Ni-based nanocrystalline alloys, however, reveal substantially worse electrochemical charging/ discharging performance compared to pure Mg2Ni, although it could be expected that B would result in Mg oxidation impediment due to the decrease of the electron density around the Mg atoms.
10 1
2
3
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number of cycles Fig. 7. Charge/discharge curves of Mg58Ni29B13 (a) and Mg64Ni32B4 (b) and cycle life of the alloys (c).
Acknowledgment The work was supported by the Bulgarian Scientific Research Fund under grant 226/2008. The authors are grateful to the project UNION (DCVP 02-2/2009).
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References [1] Ebrahimi-Purkani A, Kashani-Bozorg SF. Nanocrystalline Mg2Ni-based powders produced by high-energy ball milling and subsequent annealing. J Alloys Compd 2008;456:211e5. [2] Singh Arvind Kumar, Singh Ajay Kumar, Srivastava ON. On the synthesis of the Mg2Ni alloy by mechanical alloying. J Alloys Compd 1995;227:63e8. [3] Zaluski L, Zaluska A, Strom-Olsen JO. Hydrogen absorption in nanocrystalline Mg2Ni formed by mechanical alloying. J Alloys Compd 1995;217:245e9. [4] Orimo Shin-ichi, Fujii Hironohu. Effects of nanometer-scale structure on hydriding properties of MgeNi alloys: a review. Intermetallics 1998;6(3): 185e92. [5] Gutfleisch O, Schlorke-de Boer N, Ismail N, Herrich M, Walton A, Speight J, et al. Hydrogenation properties of nanocrystalline Mg- and Mg2Ni-based compounds modified with platinum group metals (PGMs). J Alloys Compd 2003;356e357:598e602. [6] Nohara Shinji, Fujita Naoya, Zhang Shu Guo, Inoue Hiroshi, Iwakura Chiaki. Electrochemical characteristics of a homogeneous amorphous alloy prepared by ball-milling Mg2Ni with Ni. J Alloys Compd 1998;267:76e8. [7] Zhang Yunshi, Yang Huabin, Yuan Huatang, Yang Endong, Zhou Zuoxiang, Song Deying. Dehydriding properties of ternary Mg2Ni1xZrx hydrides synthesized by ball milling and annealing. J Alloys Compd 1998;269: 278e83. [8] Goo Nam Hoon, Jeong Woon Tae, Lee Kyung Sub. Effects of Zr addition on discharge properties of mechanically alloyed Mg2Ni hydrogen-storage alloy electrode. J Power Sources 2000;87:118e24. [9] Liang G, Huot J, Boily S, Van Neste A, Schulz R. Hydrogen storage properties of nanocrystalline Mg1.9Ti0.1Ni made by mechanical alloying. J Alloys Compd 1999;282:286e90. [10] Bouaricha S, Dodelet JP, Guay D, Huot J, Schulz R. Activation characteristics of graphite modified hydrogen absorbing materials. J Alloys Compd 2001;325: 245e51. [11] Bobet J-L, Grigorova E, Khrussanova M, Khristov M, Stefanov P, Peshev P, et al. Hydrogen sorption properties of graphite-modified magnesium nanocomposites prepared by ball-milling. J Alloys Compd 2004;366:298e302. [12] Li Qian, Chou Kuo-Chih, Lin Qin, Jiang Li-Jun, Zhan Feng. Hydrogen absorption and desorption kinetics of AgeMgeNi alloys. Int J Hydrogen Energy 2004;29: 843e9.
[13] Iwakura C, Hazui S, Inoue H. Effect of temperature on electrochemical characteristics of Mg2Ni alloy. Electrochem Acta 1996;41(3):471e2. [14] Liang G, Boily S, Huot J, Van Neste A, Schulz R. Mechanical alloying and hydrogen absorption properties of the MgeNi system. J Alloys Compd 1998; 267:302e6. [15] Yuan Huatang, Li Qiudi, Song Henan, Wang Yijing, Liu Jingwang. Electrochemical characteristics of Mg2Ni-type alloys prepared by mechanical alloying. J Alloys Compd 2003;353:322e6. [16] Vijay R, Sundaresan R, Maiya MP, Srinivasa Murthy S. Comparative evaluation of MgeNi hydrogen absorbing materials prepared by mechanical alloying. Int J Hydrogen Energy 2005;30:501e8. [17] Gennari FC, Esquivel MR. Structural characterization and hydrogen sorption properties of nanocrystalline Mg2Ni. J Alloys Compd 2008;459:425e32. [18] Takahashi Y, Yukawa H, Morinaga M. Alloying effects on the electronic structure of Mg2Ni intermetallic hydride. J Alloys Compd 1996;242:98e107. [19] Spassov T, Köster U. Thermal stability and hydriding properties of nanocrystalline melt-spun Mg63Ni30Y7 alloy. J Alloys Compd 1998;279:279e86. [20] Gasiorowski A, Iwasieczko W, Skoryna D, Drulis H, Jurczyk M. Hydriding properties of nanocrystalline Mg2xMxNi alloys synthesized by mechanical alloying (M ¼ Mn, Al). J Alloys Compd 2004;364:283e8. [21] Lee Sheng-Long, Hsu Fu-Kai, Chen Wen-Chi, Lin Chih-Kuang, Lin Jing-Chie. Influence of Mg3AlNi2 content on the cycling stability of Mg2NieMg3AlNi2 hydrogen storage alloy electrodes. Intermetallics 2011;19:1953e8. [22] Anik M, Özdemir G, Küçükdeveci N, Baksan B. Effect of Al, B, Ti and Zr additive elements on the electrochemical hydrogen storage performance of MgNi alloy. Int J Hydrogen Energy 2011;36:1568e77. [23] Shao Huaiyu, Asano Kohta, Enoki Hirotoshi, Akiba Etsuo. Fabrication and hydrogen storage property study of nanostructured MgeNieB ternary alloys. J Alloys Compd 2009;479:409e13. [24] Haykawa H, Ishido Y, Nomura K, Uruno H, Ono S. Phase transformations among three polymorphs of Mg2NiH4. J Less Common Metals 1984;103:277e 83. [25] Zlatanova Z, Spassov T, Eggeler G, Spassova M. Synthesis and hydriding/ dehydriding properties of Mg2NieAB (AB ¼ TiNi or TiFe) nanocomposites. Int J Hydrogen Energy 2011;36(13):7559e66. [26] Drenchev N, Spassov T, Bliznakov St. Influence of tin on the electrochemical and gas phase hydrogen sorption in Mg2xSnxNi (x ¼ 0, 0.1, 0.3). J Alloys Compd 2008;450(1e2):288e92.
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Influence of boron on the hydriding of nanocrystalline Mg2Ni Murad Redzeb, Zlatina Zlatanova, Tony Spassov* Department of Chemistry and Pharmacy, University of Sofia “St.Kl.Ohridski”, 1 J. Bourchier Str., 1164 Sofia, Bulgaria
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 August 2012 Received in revised form 20 October 2012 Accepted 23 October 2012 Available online 7 December 2012
Nanocrystalline Mg2Ni(B) alloys (Mg64Ni32B4 and Mg58Ni29B13) were synthesized by ball milling and their hydrogen sorption properties were characterized by three different methods: Sievert’s type apparatus, High-pressure DSC and by electrochemical hydrogen charge/discharge. The main idea was to elucidate the influence of boron on the hydriding behavior of Mg2Ni, an intermetallic compound largely investigated for hydrogen storage, but still needing improvement. Boron was selected because of its small atomic radius and large electronegativity, as well as chemical resistance in solutions of alkaline basics. Hydrogen capacity and hydriding/dehydriding cycle life of the alloys were found to depend on the presence of B. Whereas small amounts of boron (4 at.%) increase the maximum hydrogen capacity of Mg2Ni, at higher B concentration (13 at.%) the capacity and cycle life deteriorate significantly due to the Mg2Ni(B) solid solution decomposition during hydriding/dehydriding. Rapid worsening was found to take place during the first 2e3 electrochemical charge/discharge cycles to a much larger extent for Mg2Ni(B), compared to pure Mg2Ni. Ó 2012 Elsevier Ltd. All rights reserved.
Keywords: A. Nanostructured intermetallics B. Hydrogen storage C. Mechanical alloying and milling
1. Introduction There is considerable interest in the use of metal hydrides as rechargeable hydrogen storage media. This interest stems from the fact that this is the safest means of storing hydrogen (compared to hydrogen liquefaction and hydrogen storage at high pressure), with a large hydrogen storage capacity, both in mass and in volume, and a range of absorption/desorption temperatures and pressures that can fit numerous applications. The binary MgeNi system is unique from the standpoint of hydrogen storage, as it contains two hydrogen absorption phases: a hexagonal (hp2) Mg metal phase and hexagonal (hp18) intermetallic line compound, Mg2Ni, whose reported maximum gravimetric hydrogen capacities are 7.66 and 3.62 wt.%, respectively. Mg2Ni is the most intensively studied due to its high hydrogen absorption capacity, low specific weight and low cost [1e21]. The Mg absorption kinetics is improved by the Ni addition, because of its good catalytic activity [1]. However, under normal conditions (i.e. at room temperature and atmospheric pressure) Mg2Ni does not absorb hydrogen. The usual hydrogenation conditions are: 250e350 C and hydrogen pressure of 15e 50 bar [2,3]. Considering the high vapor pressure and generally small solubility of alloying elements in Mg alloys, MgeNi alloys can be produced more easily and reliably by mechanical alloying (MA)
* Corresponding author. E-mail address: [email protected]fia.bg (T. Spassov). 0966-9795/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.10.016
than by conventional melting methods [4]. It is suggested that the heat generated due to collision of the balls resulting in continuous fragmentation, coalescence events and the interdiffusion of ingredients of the alloy at the collision sites are the main steps of the process of mechanical alloying [2]. Another advantage of MA is the formation of nanocrystalline Mg2Ni with very active surfaces of the powder and grain sizes of about 20e30 nm which absorbs hydrogen at lower temperature [3,4]. It has good activation properties and shows a large discharge capacity at room temperature. Nevertheless, it loses 80% of its maximum capacity within ten cycles. The main cause of the rapid degradation of the electrode is the formation of the passive layer at the surface, which hinders the charge transfer reaction. Recently, many efforts have been made to improve Mg-based alloys’ electrochemical characteristics. In particular, element substitution is proved to be very effective method for improving the high discharge capacity. Platinum group metals (PGMs) have been employed to superficially modify Mg2Ni-based compounds. Small amount of Pd catalyst enhances hydrogen absorption kinetics at 200 C. Nanocrystalline Mg2Ni with Pd absorbs hydrogen even at room temperature, without activation and with relatively good kinetics [3]. Reactively milled Mg2Ni co-milled with Ru shows an onset temperature of hydrogen desorption as low as 80 C [5]. Ball-milling Mg2Ni with metallic Ni (70 wt.% vs. Mg2Ni) leads to the formation of a homogeneous amorphous alloy, which exhibits a maximum discharge capacity of ca. 870 mA h g1 (Mg2Ni) at 30 C [6]. The ternary hydrogen-storage alloys Mg2Ni1xZrx (0 < x 0.3),
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compared to a Mg2Ni alloy, have a larger specific surface (z1.20 m2/g), more promising dehydriding kinetics, lower enthalpy of formation of hydrides than that of Mg2NiH4, and lower decomposition temperatures in an open system. An optimum desorption storage capacity of about 3.3 wt.% is observed [7]. The discharge capacity of a (Mg1xZrx)2Ni electrode increases with Zr content and reaches the highest capacity of 530 mA h g1 at x ¼ 0.3, then decreases to 230 mA h g1 at x ¼ 0.4, i.e., it is lower than that of pure Mg2Ni. Cyclic stability and rate capability vary with Zr addition, but they are both not greatly improved [8]. The nanocrystalline Mg1.9Ti0.1Ni shows better kinetics compared to nanocrystalline Mg2Ni, absorbs more than 3 wt.% H2 in 2000 s at 423 K and also destabilizes the hydride phase slightly compared to nanocrystalline Mg2Ni [9]. Graphite addition inhibits the formation of a new oxide layer on the surface of the materials once the native oxide layer is broken during the milling process and has a positive effect on the absorption kinetics [10,11]. The hydriding kinetics of Mg2Ni is improved by adding Ag [12]. Few studies have focused, however, on the evaluation of boron addition to hydrogen storage alloys. Anik et al. have concluded that Mg0.9(B)0.1Ni and Mg0.8(B)0.2Ni alloys don’t possess better properties compared to Mg50Ni50 [22]. More extensive research related to the boron influence shows that small amount of boron added to Mg50Ni50 alloy improves the hydrogen storage properties [23]. Therefore, in the present work boron was chosen as an alloying element in the nanostructured Mg2Ni (Mg64Ni32B4 and Mg58Ni29B13), because of its non-metal nature, chemical resistance in solutions of alkaline solutions, small atomic radius (thus forming interstitial solid solutions in Mg2Ni) and large electronegativity e 2.04 (Pauling). In this way, greater changes in the electronic structure of Mg2Ni compared to the metal substituted compositions can be expected, which might result in different hydriding properties as well. 2. Experimental details Mg2Ni(B) nanocrystalline alloys (Mg58Ni29B13 and Mg64Ni32B4) were synthesized by ball milling in a planetary type ball mill (Fritsch P6) under pure argon gas atmosphere with a rotation speed of 450 rpm for 15 h, ball to powder mass ratio (B/P) of 12/1, with 10 min rest after each 1 h of milling. Pure elemental powders of magnesium, nickel, and boron (amorphous) were used as starting materials. Stainless steel vials and balls were used. The average composition of the boron containing alloys was determined using EDS analyzer and was found to correspond to their nominal composition. The samples were handled in a glove box under Ar. Small amount of the powder was taken from the mill at certain periods of time for structural, morphological and thermal analysis. The powder morphology and microstructure of the alloys were analyzed by x-ray diffraction (XRD) using a Cu-Ka radiation and scanning electron microscopy (JEOL 5510 SEM). Possible phase transformations during annealing of the as-milled materials and the hydriding/dehydriding processes were studied by high-pressure differential scanning calorimetry (HPDSC), SETARAM Sensys Evo TG-DSC. The hydrogen sorption behavior (capacity and kinetics) was investigated using Sievert’s type (PCT) apparatus with samples from 100 to 150 mg. The electrochemical measurements were realized by a three electrode cell with Hg/HgO as a reference electrode and a counter electrode prepared from Ni mesh. The working electrodes were made by mixing the alloy powder (100 mg) with 30 mg teflonyzed carbon black (VULCAN 72 10%PTFE) and pressing the as-prepared mixture with a pressure of 15 MPa. In the charge/discharge cycles each electrode was charged for 5e10 h at 5 mA/cm2 and discharged to e 500 mV vs. Hg/HgO at 2 mA/cm2 in a 6 M KOH water solution.
3. Results and discussion 3.1. Morphology and microstructure of the as-milled alloys The x-ray diffraction analysis shows that alloying with even 13 at.% boron does not result in a new phase formation and due to the small atomic radius of B the positions of the diffractions peaks do not shift noticeably, Fig. 1. It can be seen that after sufficient time of milling (15 h) the powders reveal nanocrystalline microstructure, consisting of the hexagonal Mg2Ni phase. For both boron containing alloys, Mg64Ni32B4 and Mg58Ni29B13, an average size of the nanocrystallites of 10.0 0.2 nm was determined. The crystal lattice parameters show some small differences compared A and c ¼ 13.98 A), revealing large to the pure Mg2Ni (a ¼ 5.85 solid state solubility of boron into the nanocrystalline hexagonal Mg2Ni. SEM observations reveal that after continuous milling the powder particles are in the range of 1e10 mm and are relatively homogeneous in size, Fig. 2. Some smaller particles (<1 mm), cold welded together, can also be seen. 3.2. Hydrogen sorption properties The hydriding behavior of the as-milled Mg2Ni(B) alloys was analyzed by high-pressure DSC under hydrogen pressure of 15 bar, Fig. 3. In the same figure the hydriding curves for the pure nanocrystalline Mg2Ni is also plotted. From the thermograms it is evident that the nanocrystalline Mg2Ni(B) alloys absorb hydrogen already at about 150 C, proved by XRD analysis of samples, hydrided in the HPDSC at low temperature of 200 C (Fig. 1). The hydrided alloy reveals (Fig. 1) only the Mg2NiH4 phase with lattice parameters a ¼ 14,324 A and c ¼ 2701 A. Comparing the initial temperature of hydriding between the three nanocrystalline alloys it is seen that the boron containing alloys start to absorb hydrogen at slightly lower temperature than the pure Mg2Ni, as for Mg64Ni32B4 the hydriding process starts at about 130 C. For pure nanocrystalline Mg2Ni the hydrogen sorption takes place in the range 170e220 C, which is noticeably narrower temperature range than those for the boron alloys. The shape of the H-absorption peak shows that larger distribution of hydrogen sites has been observed for the B-containing alloys, especially for Mg64Ni32B4. Substantially smaller hydriding peak at about 300e330 C is also observed for the intermetallic compound
Fig. 1. XRD patterns of the ball milled Mg2Ni(B) alloys.
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Fig. 2. SEM micrograph of Mg2Ni(B) alloys: Mg64Ni32B4 (a, b) and Mg58Ni29B13 (c, d).
Mg64Ni32B4, which can be due to some microstructural inhomogeneity after ball milling, resulting in the formation of hydrogen sites with different energy. This is further confirmed by the observed two endothermic peaks of hydride(s) decomposition, taking place in the temperature range 400e450 C, which however transforms to a single-peak hydrogen absorption effect during the second hydriding experiment (see Fig. 4). It is obtained that the enthalpy change of Mg58Ni29B13 hydriding (DHhydr, [J/g]), Fig. 3, is larger than those of the other two materials, meaning that the first hydriding reaction of this alloy is connected with larger amount of the formed hydride. The multi-peaks exothermic effects, observed for the boron containing nanocrystalline compounds (Fig. 3), obviously have to be associated with the occupation of different sites by the hydrogen atoms, incl. for example inner volume of the nanocrystals and intercrystalline regions. The latter microcrystalline component is rather large in the materials studied due to the extremely fine nanocrystals formed during the milling. After the complete alloys hydriding only hydrides of Mg2Ni could be found (Fig. 1), which means that the different exothermic peaks have to be mostly associated with hydrogen occupation of different sites, rather than hydrogenation of different phases. At this stage it can be resumed that the boron containing Mg2Nibased alloys reveal somehow improved hydriding behavior, compared to pure Mg2Ni, expressed in small increase of the amount of absorbed hydrogen (assessed by the area of the exothermic peaks of hydriding) and certain lowering of the initial temperature of the hydriding process, accompanied by an increase of the temperature range of hydrogen absorption. The hydrided alloys undergo hydrides decomposition at about 400e420 C when further annealed in the DSC under hydrogen pressure of 15 bars, Fig. 3. Single peak endothermic effect
connected with Mg2NiH4 decomposition is seen for Mg2Ni and Mg58Ni29B13 and two overlapped peaks, associated with the decomposition of different hydrides and/or release of hydrogen occupying sites with different energy, characterize the Mg64Ni32B4 alloy. It is important to be mentioned that the dehydrogenation of the alloy with lower B content takes place at about 50 C lower temperature compared to that of the pure Mg2Ni alloy and Mg58Ni29B13. This result reveals some weakening of the metale hydrogen interaction, caused by the addition of small amount of B to Mg2Ni. Such improvement of the thermodynamics of the Mg64Ni32B4 hydriding/dehydriding has been observed at all hydriding/dehydriding experiments in the present study. During subsequent cooling of the dehydrided materials under hydrogen atmosphere from 500 C down to room temperature (Fig. 3b) they undergo again hydriding at about 380e400 C, as the enthalpy of this process corresponds exactly to that of dehydriding. The thermal peaks this time are well shaped and narrow due to the fast hydriding reaction at these temperatures. At temperatures of about 220 C small exothermic effects are also observed. These effects correspond to the transition of the high temperature (fcc) to the low temperature (monoclinic) Mg2NiH4, described comprehensively by Hayakawa et al. [24]. While the first hydriding reaction for the Mg58Ni29B13 reveals definitely higher enthalpy change compared to pure Mg2Ni and the alloy with lower B content and has to be associated with larger amount of the formed hydride, during the second hydriding/dehydriding cycle Mg58Ni29B13 deteriorates dramatically. Practically, this alloy shows lack of any stability against hydriding/dehydriding cycling, Fig. 4, due to decomposition of the alloy during hydriding/dehydriding. Magnesium nickel borides were detected after repeated hydriding of the alloy. This is not the case for the Mg64Ni32B4 alloy, where we observe repetition of the enthalpy values measured during the
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a
Dehydriding (0.1 bar H2) 50 mW
exo >
Mg58Ni29B13
exo>
Mg58Ni29B13
Mg64Ni32B4
Heat flow
Heat flow
Mg64Ni32B4
Mg2Ni
20 mW
heating
50
100
150
200
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300
350
400
450
500
100
200
o
Temperature, C
b
300
400
o
Temperature, C Fig. 5. Decomposition (0.1 bar and 5 K/min) of the hydrided Mg2Ni(B) alloys.
50 mW
exo>
Mg58Ni29B13
Heat flow
Mg64Ni32B4
Mg2Ni cooling
450
400
350
300
250
200
150
100
o
Temperature, C Fig. 3. HPDSC plots of Mg2Ni(B) alloys at hydrogen pressure of 15 bar (heating (a) and cooling (b) with a rate of 5 K/min).
first hydriding/dehydriding cycle, which is also valid for the pure Mg2Ni [25]. The hydrided materials were also subjected to hydrides decomposition by annealing in the DSC with a constant heating rate under low vacuum (0.1 bar H2), Fig. 5. The intermetallic compounds reveal also clear endothermic peaks, associated with a decomposition of the hydride phases. The composites dehydriding is associated with large and well-defined endothermic peaks at about 200e220 C. The enthalpies of dehydriding under vacuum agree very well with those determined from the DSC experiments under hydrogen pressure for the corresponding materials. Hydriding kinetic data for the nanocrystalline Mg2Ni-based powders obtained by Sievert’s type apparatus at 200 C are presented in Fig. 6. The nanocrystalline powders absorb hydrogen without activation, as the maximum capacity is reached during the first absorptionedesorption cycle. The maximum capacity of the powders is about 3.0 wt.%, as for Mg64Ni32B4 the hydrogen absorption process starts before its registration at 200 C, due to the higher rate of hydriding at this temperature, compared to Mg58Ni29B13. Mg58Ni29B13 shows noticeable deterioration in capacity and kinetics during the second hydriding experiment, which was not observed for the Mg64Ni32B4 alloy, possessing
Hydriding/dehydriding (second run) 13,3
4
Hydrogen absorption o 200 C
Mg58Ni29B13
13,2
Mg64Ni32B4
50 mW
13,1
2
13,0
12,9
100
200
300
400 o
Temperature, C Fig. 4. Second hydriding/dehydriding HPDSC cycle.
500
0
12,8 0
500
1000
1500
2000
time, s Fig. 6. PeC isotherms at 200 C for the Mg2Ni(B) alloys.
2500
wt.% H2
pressure, bar
Mg64Ni32B4
Heat flow
exo>
Mg58Ni29B13
M. Redzeb et al. / Intermetallics 34 (2013) 63e68
reproducible hydrogen sorption behavior. This fact confirms quite well the results obtained by HPDSC. In general, the initial hydrogen sorption kinetics is fast for the alloys studied, particularly for this relatively low temperature for Mg2Ni-based alloys (200 C), reaching about 50% of hydriding
a
Mg58Ni29B13
1000
67
within the first 500e600 s. The initial absorption rate obtained from the slope of the fitted straight line on the sorption curves in the range of 100e500 s indicates a favorable kinetics for the boron richer alloy, but as already mentioned its hydriding kinetics becomes worse immediately after the first hydriding cycle. This effect was already explained by drastic microstructural changes (phase decomposition of the B richer alloy) observed by XRD, result of the hydriding/dehydriding process. Obviously the solid solution of boron into the hexagonal Mg2Ni intermetallics at low B content (Mg64Ni32B4) is stable and therefore shows reproducible hydriding behavior.
U, mV (vs. Hg/HgO)
900
3.3. Electrochemical hydriding/dehydriding 800
700
600
500 0
2
4
6
8
10
12
14
time, h
b
Mg64Ni32B4
1000
U, mV (vs. Hg/HgO)
900
800
The observed in some way superior low-temperature hydriding behavior of the boron containing alloys (esp. Mg64Ni32B4) compared to pure Mg2Ni gives some reasons to expect improved electrochemical hydriding/dehydriding performance of these materials as well. Charge/discharge curves at galvanostatic conditions were presented in Fig. 7 for both alloys. The B containing Mg2Ni alloys show however worse electrochemical hydrogen storage properties in comparison with the pure intermetallic compound [26]. The cycle life of the electrodes is presented in Fig. 7c. It can be concluded that the amount of absorbed hydrogen decreases immediately after the first cycle. Both alloys show comparable corrosion stability, substantially lower than that of pure Mg2Ni [26]. Although the alloy with lower B content reveals slightly better electrochemical hydrogen charge/discharge behavior compared to the B richer alloy the discharge capacity and cycle life as a whole are rather low and unsatisfactory. Obviously, alloying with B deteriorates drastically the discharge capacity of Mg2Ni, accelerating the oxidation of magnesium on the alloy surface, forming Mg(OH)2.
700
4. Conclusion 600
500
0
c
2
4
6
8
10
12
14
16
18
time, h 90
Mg58Ni29B13
80
Mg64Ni32B4
capacity, mAh/g
70 60 50 40 30 20
Nanocrystalline solid solutions of boron into hexagonal Mg2Ni with overall composition of Mg64Ni32B4 and Mg58Ni29B13 were synthesized by ball milling and their hydrogen sorption properties were characterized. The main idea was to reveal the influence of boron, a non-metal element with small atomic radius and large electronegativity as well as chemically resistant in solutions of alkaline basics on the hydriding behavior of Mg2Ni. Relatively small amount of boron (4 at.% B) generally slightly increases the gas-phase hydrogen sorption capacity of Mg2Ni and improves the hydriding thermodynamics and kinetics. Additional increase of the B content (13 at.% B), however, results in hydriding properties deterioration due to boron solid solution decomposition during hydriding/dehydriding. In respect to the hydriding/dehydriding cycling stability in hydrogen gas phase Mg64Ni32B4 is comparable to pure Mg2Ni. Both boron containing Mg2Ni-based nanocrystalline alloys, however, reveal substantially worse electrochemical charging/ discharging performance compared to pure Mg2Ni, although it could be expected that B would result in Mg oxidation impediment due to the decrease of the electron density around the Mg atoms.
10 1
2
3
4
5
6
7
number of cycles Fig. 7. Charge/discharge curves of Mg58Ni29B13 (a) and Mg64Ni32B4 (b) and cycle life of the alloys (c).
Acknowledgment The work was supported by the Bulgarian Scientific Research Fund under grant 226/2008. The authors are grateful to the project UNION (DCVP 02-2/2009).
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