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Morphology and hydrogen absorption properties of an AB2 type alloy ball milled with Mg2Ni
Journal of Alloys and Compounds 268 (1998) 248–255
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Morphology and hydrogen absorption properties of an AB 2 type alloy ball milled with Mg 2 Ni ´ D. Cracco*, A. Percheron-Guegan ´ Laboratoire de Chimie Metallurgique des Terres Rares 1, Pl. A. Briand, 92195 Meudon Cedex, France Received 6 October 1997; received in revised form 3 November 1997
Abstract The influence of ball milling an AB 2 -type Laves phase pre-alloy with Mg 2 Ni was investigated from a structural and a morphological point of view. The influence of such a milling on the hydrogen uptake properties was also investigated. It was found out that magnesium did not enter the pre-alloy cell, at least for low energy millings. However, the magnesium did intimately mix with the pre-alloy. Regarding the weakest milling, significant improvement of hydrogen absorption capacity and kinetics were observed. Referring to previous work, this enhancement was attributed to the magnesium enrichment of the surface consequent to ball milling with Mg 2 Ni. 1998 Elsevier Science S.A. Keywords: Hydrogen absorption; Laves phase; Magnesium–nickel alloys; Ball milling
1. Introduction The AB 2 -type Laves phase compounds containing zirconium have attracted a great deal of attention as hydrogen storage materials since the studies of Shaltiel et al. [1]. The binary alloy ZrMn 2 exhibited an interesting hydrogen absorption capacity (3.6 H / mol), but its hydride was too stable to be of practical significance. Studies were therefore carried out in order to increase the plateau pressure. One of the successful ways explored was to substitute manganese with V, Cr, Ni or Co [2–4]. The aim of this paper is to decrease the weight of a Zr-based AB 2 -type Laves phase compound studied by Gamo et al. [4], namely Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ), by partially substituting Mg for Zr and to study its hydrogen absorption properties. The samples were characterized and their hydrogen absorption properties were measured.
2. Experimental details As magnesium has a high vapor pressure (2 MPa at 1833 K) and Zr has high melting point (2125 K), the high-temperature elaboration of an alloy containing Zr and Mg is troublesome. An original method in three steps was therefore adopted. First, a magnesium-free pre-alloy was *Corresponding author. E-mail: [email protected] 0925-8388 / 98 / $19.00 1998 Elsevier Science S.A. All rights reserved. PII S0925-8388( 97 )00552-5
prepared by melting metals together: Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ). Magnesium was too ductile to be added by ball milling, therefore Mg 2 Ni was chosen. Mg 2 Ni was prepared by ball-milling. Finally, the pre-alloy and Mg 2 Ni were milled together in stoichiometric proportions in order to obtain Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). The pre-alloy Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) was prepared from elements by melting in a r.f. levitation furnace under an argon atmosphere. The elements had the following purity: zirconium 99.9%, manganese 99.99%, nickel 99.9%, vanadium 99.99% and chromium 99.9%. It was annealed for 11 days at 1323 K. It was analyzed by X-ray diffraction and electronic microprobe analysis (EMPA) to ensure its composition and homogeneity. Mg 2 Ni was prepared by mechanical alloying using a Fritsch ‘Pulverisette 7’ planetary ball mill, starting from elementary powders of magnesium (325 mesh and purity 99.6%) and nickel (2 mm average size and purity 99.9%). It was analysed by X-ray diffraction. The pre-alloy was roughly crushed under argon and 1 mol of pre-alloy was combined with 0.05 mol of Mg 2 Ni, in order to elaborate Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). The mixture was then placed in a cylindrical tempered steel container which was tightly closed under argon atmosphere. The alloys were milled for various times at different speeds using a Fritsch ‘Pulverisette 7’ planetary ball mill (Table 1). The samples were named after the following pattern: aSb, with a being
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Table 1 Milling times and intensities used for the mechanical alloying of Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) Setting intensity
Disc rotation speed (rpm)
Vial rotation speed (rpm)
Milling time (sample name)
3 5 7
229 371 510
458 742 1020
1 h (1S3) 1 h (1S5) 1 h (1S7), 3 h (3S7), 6 h (6S7), 12 h (12S7), 24 h (24S7)
the milling time and b being the milling intensity. The powder to ball mass ratio was kept constant: 1 / 13. Pressure–composition–temperature (PCT) isotherm measurements were performed at 343 K using the Sieverts method, after the samples were heated at 573 K under vacuum for 1 h. The capacities were measured via absorption and the isothermal curves reported are desorption measurements. X-ray diffraction was performed using Cu Ka on a Philips diffractometer, EMPA using a Camebax, scanning electron microscopy using a Jeol JSM-840, transmission electron microscopy using a Phillips 2000FX, and electron dispersive X-ray analysis (EDX) with a Link system. Diffractograms were fitted using the Rietveld method [5,6], with the Fullprof software [7].
3. Results and discussion of structural and morphology studies
3.1. X-ray diffraction patterns Fig. 1 shows the pre-alloy diffractogram after being annealed for 11 days at 1323 K. It presents a C15 ˚ After structure, and its cell parameter is equal to 7.054 A. being ball milled with Mg 2 Ni, the samples still have the C15 structure. However, the longer the milling times and intensities the broader the diffraction lines will be, as shown in Fig. 2. After milling at intensity 7, the lines were so broad that
Fig. 1. Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) diffractogram, after being annealed for 11 days at 1323 K.
Fig. 2. Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) diffractograms depending on milling time at intensity 7.
a pseudo-Voigt function was not satisfying to fit a line. In order to be able to compare the characteristics of the samples, two phases were considered: a ‘crystalline’ phase and an ‘amorphous’ phase. The ‘amorphous’ phase was set as the phase of which the diffracting domain size was ˚ Therefore, each line was smaller than three cells (21 A). fitted with two pseudo-Voigt functions. This method, despite considering a raw two-phase model, allowed us to fit the patterns of the samples milled during up to 9 h and also to compare the percentages of ‘amorphous’ phase between the different samples milled at intensity 7. As shown in Fig. 3, the longer the samples were milled the more amorphous they became, which results from the
Fig. 3. Evolution of amorphous phase percentage with milling time, at intensity 7.
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Fig. 4. Evolution of cell parameter and average size of diffracting domains with milling time.
accumulation of point and lattice defects, leading to the destabilization of the crystalline phases towards an amorphous phase of the same composition [8]. Fig. 4 shows the evolution of cell parameter with milling time. In the first ˚ for the hours, the cell parameter increased, from 7.054 A ˚ after 6 h of milling. Further unmilled pre-alloy to 7.065 A milling then decreased the value of the cell parameter. The increase of the cell volume resulted from the disorder created by milling, which lowered the packing efficiency. Regarding longer milling times, Lutterotti et al. [9] observed the same evolution and proposed that the creation of vacancies and anti-sites was responsible for the cell volume decrease. Fig. 4 shows the evolution of average size of the diffracting domains with milling time. At intensity 7, the crystallites size decreased when milling ˚ was reached time was increased. A critical size of 72 A ˚ after milling for 12 h, and a slightly smaller size (68 A) was obtained after milling for 24 h. The two-stage evolution observed is the same as the one mentioned by Li et al. [10]. First, the milling introduces shear bands which are followed by fractures, therefore decreasing the size of the crystallites. When the particle size becomes smaller, the shear strength becomes higher, and at a critical point the shear band cannot pass through the crystallite. This will therefore prevent the size refinement, as was observed after 12 h of milling.
Table 2 sums up the cell parameters and average diffracting domain sizes measured for different milling intensities and times. After having milled the pre-alloy at intensity 7 for 1 h, with or without Mg 2 Ni, cell parameters ˚ respecmeasured were very similar: 7.055 and 7.057 A, tively. In order to determine if magnesium entered the pre-alloy cell during the milling, the following is considered. Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) and Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) should have the same cell parameter, as Mg and Zr have the same atomic radius. Yet, Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ), milled for 1 h at intensity 7, has ˚ which is 0.007 A ˚ more than a cell parameter of 7.064 A, what was measured for Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled for 1 h at intensity 7 with Mg 2 Ni. This therefore proves that magnesium did not enter the pre-alloy cell after 1 h of milling at intensity 7. Regarding milling carried out at lower intensities, the cell parameters remained very similar to the pre-alloy cell parameter: these intensities are therefore too low to allow the magnesium to enter the cell.
3.2. EMPA First, it is to be remembered that, in this analysis, the grains studied measured more than 1 mm (because the size of the beam does not allow to analyze smaller grains) and that only the core of these grains could be analyzed with this technique. Regarding 1S3, magnesium was not detected and the pre-alloy composition was measured. It therefore showed that magnesium was not in the core of grains with sizes over 1 mm. This was confirmed by looking at the EDXA image of magnesium, as seen in Fig. 5. Indeed, the EDXA image of magnesium did not correspond with the image of the grains obtained by collecting the secondary electrons. Regarding 1S5, magnesium was detected in some grains and not in others, its distribution is very heterogeneous. Regarding alloys obtained by milling at intensity 7 for more than 1 h, magnesium was detected in all the grains, and Table 3 shows the measured composition. In order to quantify the heterogeneity of the element distributions in the alloys, the following ratio, R, was calculated: standard deviationof element X in wt.% R 5 ]]]]]]]]]]] average wt.% of element X
Table 2 Cell parameters and crystallite sizes of the alloys elaborated Sample
˚ a (A)
˚ Crystallite sizes (A)
Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ), non-milled Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled for 1 h (intensity 7) without Mg 2 Ni Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled with Mg 2 Ni 1 h (intensity 3) 6 h (intensity 3) 1 h (intensity 5) 2 h 36 min (intensity 5) 1 h (intensity 7) Zr 1.0 (Mn 0.6 Ni 1.17 V0.2 Cr 0.1 ) milled 1 h (intensity 7) without Mg 2 Ni
7.054 7.057
1550 194
7.051 7.052 7.050 7.052 7.055 7.064
455 353 290 239 180 580
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Fig. 5. Grains of alloy milled for 1 h at intensity 3: (a) EDXA image of magnesium; (b) image of secondary electrons.
Fig. 7. Grains of alloy milled for 12 h at intensity 7: (a) EDXA image of magnesium; (b) image of secondary electrons.
Table 3 Compositions analyzed by EMPA
3.3. SEM analysis
Milling times (h) at intensity 7
Composition
1 3 6 12 24
Zr 0.92 Zr 0.93 Zr 0.90 Zr 0.92 Zr 0.90
Mg 0.07 (Mn 0.59 Ni 1.18 V0.22 Cr 0.11 ) Mg 0.10 (Mn 0.57 Ni 1.18 V0.21 Cr 0.10 ) Mg 0.15 (Mn 0.57 Ni 1.18 V0.21 Cr 0.10 ) Mg 0.07 (Mn 0.60 Ni 1.17 V0.20 Cr 0.10 ) Mg 0.11 (Mn 0.60 Ni 1.15 V0.20 Cr 0.10 )
The bigger R, the more heterogeneous is the distribution. This ratio was calculated for each element for different milling times at intensity 7, the result is shown in Fig. 6. Looking at the figure, it becomes obvious that the longer the milling, the more homogeneous the distribution of magnesium. Furthermore, the EDXA image of magnesium in sample 12S7 confirmed that magnesium was very well distributed in this alloy. Indeed, as seen in Fig. 7, the EDXA image of magnesium was the same as that obtained when the secondary electrons were collected. Besides, it is interesting to note that milling for more than 1 h also increased the homogeneity of the other elements, in comparison with the annealed pre-alloy.
Fig. 6. Heterogeneity of the elements as a function of milling time, at intensity 7.
Observing polished samples through SEM showed that grains were bigger as milling was longer and more intense, as presented in Fig. 8. Observing the powders glued on a carbon film showed that these grains had a rough surface which was actually due to the agglomeration of very fine particles, with sizes smaller than 1 mm (Fig. 9c Fig. 10b). This therefore showed that the longer and the more intensively the powders were milled, the more the fine
Fig. 8. SEM images of polished samples: (a) alloy 1S3; (b) alloy 12S7.
Fig. 9. SEM images of 1S3, as a powder: (a) unmilled pre-alloy grain; (b) rough grain; (c) zoom on the surface of the rough grain.
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Fig. 10. SEM images of 12S7, as a powder: (a) rough grain; (b) small agglomeration.
particles agglomerated. The size of these agglomerates was as large as several hundreds of micrometers. Regarding 1S3, not only those grains were observed, but unmilled pre-alloy grains were also detected. Their average size was around 10 mm and they were covered with fine particles, as shown in Fig. 9a. The more the samples were milled, the less unmilled pre-alloy grains were observed. EDX analysis on polished samples gave the same results as the one obtained with EMPA: magnesium was only detected on the surface of most of the grains in 1S3 and was homogeneously distributed in 12S7. Besides, EDX analysis on the powders showed that magnesium was detected on every grain surface, including the pre-alloy grains which were not crushed by the milling. Collecting back-scattered electrons with the SEM did not show any evidence of phase contrast. It therefore seems that magnesium was intimately mixed with the pre-alloy, on the surface at least.
3.4. TEM analysis TEM studies confirmed that the powders were constituted of aggregates of nano-particles, as can be seen in Fig. 11a Fig. 12b. The selected area diffraction pattern, shown
Fig. 12. TEM bright field image of a grain surface of 12S7.
in Fig. 11b, also confirmed the nano-crystallinity of the analyzed grains. In sample 1S3, EDX analysis showed that magnesium was not present in all the nano-particles. Besides, studying nano-particles containing magnesium, a gradient of magnesium concentration could be observed: the surface was magnesium rich (54 weight%) and the percentage decreased when approaching the core, reaching 3 wt.%. This weight percentage (3%) is actually the average weight percentage of magnesium measured with EMPA in the alloy Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). EDX analysis also showed that Mg 2 Ni mostly did not remain as Mg 2 Ni in the nano-particles, and new phases appeared. The phases have yet to be identified. It is therefore appropriate to use the term ‘mechanical alloying’, even for the less intense and shortest millings. In sample 12S7, the magnesium concentration measured was homogeneous on a nanometer scale. Only the grain surfaces could be analyzed (the core was too thick) but the EDXA images of magnesium (Fig. 7a) lead us to conclude that magnesium was also homogeneously distributed inside the nano-particles.
4. Results and discussion of the hydrogen absorption studies
4.1. PCT isotherm measurements
Fig. 11. TEM images of 1S3: (a) bright field; (b) selected area diffraction.
PCT isotherm measurements were carried out at 343 K after three absorption–desorption cycles. Results obtained are presented in Fig. 13. Table 4 summarizes the absorption capacity of the samples milled with Mg 2 Ni for different times and intensities. The best absorption capacity was obtained for 1S3 (3.1 H / mol), which is the sample milled for the shortest time and at the lowest intensity. Furthermore, 12S7 exhibited the typical PCT isotherm of
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Fig. 13. PCT isotherms at 343 K of samples milled with Mg 2 Ni for different times and intensities: 1S3 (n), 1S7 (3), 3S7 (h), 12S7 (d). Table 4 Hydrogen absorption capacity of different samples Sample
Absorption capacity (H / mol)
1S3 1S7 3S7 12S7
3.06 2.77 2.45 2.29
an amorphous sample, with a poor absorption capacity (2.3 H / mol) and a plateau with a slope. It therefore showed that the less the samples were milled the better their absorption properties. This was attributed to the amorphization and disordering of the samples with milling. As in a study on hydrogen diffusion rates reported previously [11,12], our results confirmed that amorphization in C15 Laves phase compounds was altering hydrogen absorption properties. This is also true for kinetics as shown in Fig. 14: indeed, the more the samples were milled the more slowly their absorption took place. The influence of milling with and without Mg 2 Ni was also studied. The PCT isotherm measurements at 343 K after 29 absorption–desorption cycles are reported in Fig. 15 for samples of pre-alloy unmilled or milled, and with or without Mg 2 Ni. It shows that milling the pre-alloy alone
Fig. 14. Absorption kinetics at 343 K of: 1S5 (d), 1S7 (s), 3S7 (n) and 6S7 (3).
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Fig. 15. PCT isotherms at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (h), milled for 1 h at intensity 3 (m), milled for 1 h at intensity 3 with Mg 2 Ni (3).
for 1 h at intensity 3 decreased the absorption capacity of the pre-alloy from 2.7 to 2.6 H / mol. This reduction was due to disordering as mentioned in the previous paragraph. Besides, when the pre-alloy was milled with Mg 2 Ni for the same time and the same intensity, it was found that the absorption capacity of the pre-alloy was increased from 2.7 to 3.1 H / mol. Regarding the absorption kinetics, the results for the unmilled pre-alloy were in good agreement with the results Hsu and Perng obtained with similar compounds [3]. Furthermore, Fig. 16 shows that milling the pre-alloy alone reduced the absorption rate of hydrogen, whereas milling with Mg 2 Ni lead to the same absorption rate as the unmilled pre-alloy, with an increased absorption capacity. This showed the beneficial role played by the addition of Mg 2 Ni, regarding hydrogen absorption. It therefore seems that some new phases or interfaces generated by milling with Mg 2 Ni have a catalytic activity. As mentioned by Uchida et al. [13], the presence of Ni could be responsible for such enhancements. Yet, not only the kinetics are enhanced by the addition of Mg 2 Ni, but also the absorption capacity. Another interesting hypothesis, suggested by Fujii et al. [14], is that magnesium could
Fig. 16. Absorption kinetics at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (j), milled for 1 h at intensity 3 (s), milled for 1 h at intensity 3 with Mg 2 Ni (m).
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Fig. 17. Evolution of absorption capacities (after 17 h exposure to hydrogen) with cycling at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (m), milled for 1 h at intensity 3 (h), milled for 1 h at intensity 3 with Mg 2 Ni (s).
Fig. 18. Evolution of absorption capacities (after 1 h exposure to hydrogen) with cycling at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (n), milled for 1 h at intensity 3 (j), milled for 1 h at intensity 3 with Mg 2 Ni (s).
act as a binder and prevent the pre-alloy from oxidation. Further studies are planned in order to account for such enhancements.
¨ Zuttel et al. observed, using XPS, that after cycling ZrV0.4 Ni 1.6 (which is also a Zr-based AB 2 compound) its surface became Zr enriched, with most of the Zr oxidized, and also Ni deficient. The ZrO 2 layer and the Ni deficiency are both responsible for kinetic slowdown, as they prevent the dissociation of H 2 [13,16]. A similar phenomenon probably takes place with the unmilled pre-alloy, accounting for such a kinetic decrease in the first cycles. Complementary observations, such as XPS, are needed in order to be able to confirm such a hypothesis. The absorption kinetics of the milled samples did not vary as much: a 19% decrease was observed for the sample milled alone and a 14% decrease for the sample milled with Mg 2 Ni. In order to be able to explain properly the role played by the addition of Mg 2 Ni and ball milling, further investigations are needed. All we can say at present is that the magnesium-rich surface seems to enhance the kinetics. Indeed, as Hjort et al. reported [17], MgO x layers can increase hydrogen uptake rates. It is also interesting to note the advantage of milling with Mg 2 Ni, since it yielded the best average absorption capacity after 1 h exposure to hydrogen, over the first 29 cycles: 2.6 H / mol, as presented in Table 5. The average for the unmilled pre-alloy was 2.1 H / mol and for the pre-alloy milled alone it was 1.8 H / mol. If absorption after 17 h exposure to hydrogen is considered, 1S3 is still the best absorbing sample, but the difference with the two other samples is not that large: 2.9 H / mol for 1S3 and 2.8 H / mol for the unmilled pre-alloy. Even if magnesium did
4.2. The absorption evolution with cycling The absorption capacities of unmilled pre-alloy, of prealloy milled alone for 1 h at intensity 3, and pre-alloy milled with Mg 2 Ni for 1 h at intensity 3, were recorded at 343 K. They were then desorbed at 573 K for 1 h; a subsequent absorption was then carried out after the sample temperature was stabilized at 343 K. Twenty-nine absorption–desorption cycles were achieved following this procedure. Fig. 17 presents the absorption capacity evolutions after 17 h of exposure to hydrogen. It shows that the sample milled with Mg 2 Ni always kept a better absorption capacity than the other two. Another improvement due to the addition of Mg 2 Ni is in the evolution of the kinetics with cycling. Fig. 18 presents the evolution of absorption capacities of the same samples, but after only 1 h of exposure to hydrogen. As the absorption was not completed after 1 h exposure to hydrogen, these data will therefore yield information on the evolution of the kinetics with cycling. Regarding the unmilled pre-alloy, its absorption kinetics were slowed down in the first cycles. During the first 10 cycles, the absorption uptake in 1 h decreased by 43% (from 2.95 to 1.68 H / mol). This could be due to the fact that elements ¨ diffuse with cycling, as Zuttel et al. [15] reported. Indeed, Table 5 Average absorption capacities over the 29 absorption–desorption cycles
Average absorption capacity (H / mol):
Unmilled pre-alloy Pre-alloy milled alone for 1 h at intensity 3 Pre-alloy milled with Mg 2 Ni for 1 h at intensity 3
considering 1 h exposure to H 2
considering 17 h exposure to H 2
2.1 1.8 2.6
2.8 2.4 2.9
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not enter the cell of the pre-alloy, it increased its kinetics and absorption capacity at 343 K.
5. Conclusion From the structural data, we were not able to conclude if we succeeded in incorporating magnesium in the pre-alloy cell for samples milled over 1 h at intensity 7. However, these samples were too amorphous to be of practical interest regarding hydrogen absorption. Regarding samples milled at lower intensities, it was shown that magnesium did not enter the cell of the pre-alloy. Magnesium did, however, intimately mix with the other elements, as shown by the EDX analysis performed with TEM, and a gradient of magnesium concentration could be measured inside submicron grains. Furthermore, it was shown that magnesium did improve the hydrogen uptake rate as well as the absorption capacity. And this remained true even after 29 absorption–desorption cycles.
Acknowledgements The authors would like to express their thanks for financial support by the company Alcatel Alsthom Recherche.
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References [1] D. Shaltiel, I. Jacob, D. Davidov, J. Less-Common Met. 53 (1977) 117. [2] Y. Moriwaki, K. Gamo, H. Seri, T. Iwaki, J. Less-Common Met. 172,174 (1991) 1211. [3] Y.-S. Hsu, T.-P. Perng, J. Alloys Comp. 227 (1995) 180. [4] T. Gamo, Y. Tsuji, Y. Moriwaki, in: P.D. Benett, T. Sakai (Eds.), Proc. Symp. Hydrogen and Metal Hydride Batteries, The Electrochemical Society, Pennington, NJ, 1994, p. 155. [5] H.M. Rietveld, Acta Crystallogr. 22 (1967) 151. [6] H.M. Rietveld, J. Appl. Crystallogr. 2 (1969) 65. [7] J. Rodriguez-Carvajal, in: XVth Congr. Int. Union of Crystallography, Satellite Meeting on Powder Diffraction, Toulouse, 1990, pp. 1–127. [8] R.B. Schwartz, R.R. Petrich, C.K. Saw, J. Non-Cryst. Solids 76 (1986) 281. [9] L. Lutterotti, S. Gialanella, R. Caudron, Mater. Sci. Forum 228,231 (1996) 551. [10] S. Li, K. Wang, L. Sun, Z. Wang, Scripta Metall. Mater. 27 (1992) 437. [11] S.V. Rychkova, M.Yu Belyaev, A.V. Skripov, A.P. Stepanov, Sov. Phys. Solid State 30(7) (1988) 1285. [12] A.V. Skripov, S.V. Rychkova, M.Yu Belyaev, A.P. Stepanov, Solid State Commun. 71 (1989) 1119. [13] H.H. Uchida, H.-G. Wulz, E. Fromm, J. Less-Common Met. 172,174 (1991) 1076. [14] H. Fujii, S. Orimo, K. Yamamoto, K. Yoshimoto, T. Ogasawara, J. Less-Common Met. 175 (1991) 243. ¨ [15] A. Zuttel, F. Meli, L. Schlapbach, J. Alloys Comp. 200 (1993) 157. ¨ [16] J.O. Strom-Olsen, Y. Zhao, D.H. Ryan, Y. Huai, R.W. Cochrane, J. Less-Common Met. 172,174 (1991) 922. [17] P. Hjort, A. Kroser, B. Kasemo, J. Alloys Comp. 237 (1996) 74.
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Morphology and hydrogen absorption properties of an AB 2 type alloy ball milled with Mg 2 Ni ´ D. Cracco*, A. Percheron-Guegan ´ Laboratoire de Chimie Metallurgique des Terres Rares 1, Pl. A. Briand, 92195 Meudon Cedex, France Received 6 October 1997; received in revised form 3 November 1997
Abstract The influence of ball milling an AB 2 -type Laves phase pre-alloy with Mg 2 Ni was investigated from a structural and a morphological point of view. The influence of such a milling on the hydrogen uptake properties was also investigated. It was found out that magnesium did not enter the pre-alloy cell, at least for low energy millings. However, the magnesium did intimately mix with the pre-alloy. Regarding the weakest milling, significant improvement of hydrogen absorption capacity and kinetics were observed. Referring to previous work, this enhancement was attributed to the magnesium enrichment of the surface consequent to ball milling with Mg 2 Ni. 1998 Elsevier Science S.A. Keywords: Hydrogen absorption; Laves phase; Magnesium–nickel alloys; Ball milling
1. Introduction The AB 2 -type Laves phase compounds containing zirconium have attracted a great deal of attention as hydrogen storage materials since the studies of Shaltiel et al. [1]. The binary alloy ZrMn 2 exhibited an interesting hydrogen absorption capacity (3.6 H / mol), but its hydride was too stable to be of practical significance. Studies were therefore carried out in order to increase the plateau pressure. One of the successful ways explored was to substitute manganese with V, Cr, Ni or Co [2–4]. The aim of this paper is to decrease the weight of a Zr-based AB 2 -type Laves phase compound studied by Gamo et al. [4], namely Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ), by partially substituting Mg for Zr and to study its hydrogen absorption properties. The samples were characterized and their hydrogen absorption properties were measured.
2. Experimental details As magnesium has a high vapor pressure (2 MPa at 1833 K) and Zr has high melting point (2125 K), the high-temperature elaboration of an alloy containing Zr and Mg is troublesome. An original method in three steps was therefore adopted. First, a magnesium-free pre-alloy was *Corresponding author. E-mail: [email protected] 0925-8388 / 98 / $19.00 1998 Elsevier Science S.A. All rights reserved. PII S0925-8388( 97 )00552-5
prepared by melting metals together: Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ). Magnesium was too ductile to be added by ball milling, therefore Mg 2 Ni was chosen. Mg 2 Ni was prepared by ball-milling. Finally, the pre-alloy and Mg 2 Ni were milled together in stoichiometric proportions in order to obtain Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). The pre-alloy Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) was prepared from elements by melting in a r.f. levitation furnace under an argon atmosphere. The elements had the following purity: zirconium 99.9%, manganese 99.99%, nickel 99.9%, vanadium 99.99% and chromium 99.9%. It was annealed for 11 days at 1323 K. It was analyzed by X-ray diffraction and electronic microprobe analysis (EMPA) to ensure its composition and homogeneity. Mg 2 Ni was prepared by mechanical alloying using a Fritsch ‘Pulverisette 7’ planetary ball mill, starting from elementary powders of magnesium (325 mesh and purity 99.6%) and nickel (2 mm average size and purity 99.9%). It was analysed by X-ray diffraction. The pre-alloy was roughly crushed under argon and 1 mol of pre-alloy was combined with 0.05 mol of Mg 2 Ni, in order to elaborate Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). The mixture was then placed in a cylindrical tempered steel container which was tightly closed under argon atmosphere. The alloys were milled for various times at different speeds using a Fritsch ‘Pulverisette 7’ planetary ball mill (Table 1). The samples were named after the following pattern: aSb, with a being
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Table 1 Milling times and intensities used for the mechanical alloying of Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) Setting intensity
Disc rotation speed (rpm)
Vial rotation speed (rpm)
Milling time (sample name)
3 5 7
229 371 510
458 742 1020
1 h (1S3) 1 h (1S5) 1 h (1S7), 3 h (3S7), 6 h (6S7), 12 h (12S7), 24 h (24S7)
the milling time and b being the milling intensity. The powder to ball mass ratio was kept constant: 1 / 13. Pressure–composition–temperature (PCT) isotherm measurements were performed at 343 K using the Sieverts method, after the samples were heated at 573 K under vacuum for 1 h. The capacities were measured via absorption and the isothermal curves reported are desorption measurements. X-ray diffraction was performed using Cu Ka on a Philips diffractometer, EMPA using a Camebax, scanning electron microscopy using a Jeol JSM-840, transmission electron microscopy using a Phillips 2000FX, and electron dispersive X-ray analysis (EDX) with a Link system. Diffractograms were fitted using the Rietveld method [5,6], with the Fullprof software [7].
3. Results and discussion of structural and morphology studies
3.1. X-ray diffraction patterns Fig. 1 shows the pre-alloy diffractogram after being annealed for 11 days at 1323 K. It presents a C15 ˚ After structure, and its cell parameter is equal to 7.054 A. being ball milled with Mg 2 Ni, the samples still have the C15 structure. However, the longer the milling times and intensities the broader the diffraction lines will be, as shown in Fig. 2. After milling at intensity 7, the lines were so broad that
Fig. 1. Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) diffractogram, after being annealed for 11 days at 1323 K.
Fig. 2. Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) diffractograms depending on milling time at intensity 7.
a pseudo-Voigt function was not satisfying to fit a line. In order to be able to compare the characteristics of the samples, two phases were considered: a ‘crystalline’ phase and an ‘amorphous’ phase. The ‘amorphous’ phase was set as the phase of which the diffracting domain size was ˚ Therefore, each line was smaller than three cells (21 A). fitted with two pseudo-Voigt functions. This method, despite considering a raw two-phase model, allowed us to fit the patterns of the samples milled during up to 9 h and also to compare the percentages of ‘amorphous’ phase between the different samples milled at intensity 7. As shown in Fig. 3, the longer the samples were milled the more amorphous they became, which results from the
Fig. 3. Evolution of amorphous phase percentage with milling time, at intensity 7.
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Fig. 4. Evolution of cell parameter and average size of diffracting domains with milling time.
accumulation of point and lattice defects, leading to the destabilization of the crystalline phases towards an amorphous phase of the same composition [8]. Fig. 4 shows the evolution of cell parameter with milling time. In the first ˚ for the hours, the cell parameter increased, from 7.054 A ˚ after 6 h of milling. Further unmilled pre-alloy to 7.065 A milling then decreased the value of the cell parameter. The increase of the cell volume resulted from the disorder created by milling, which lowered the packing efficiency. Regarding longer milling times, Lutterotti et al. [9] observed the same evolution and proposed that the creation of vacancies and anti-sites was responsible for the cell volume decrease. Fig. 4 shows the evolution of average size of the diffracting domains with milling time. At intensity 7, the crystallites size decreased when milling ˚ was reached time was increased. A critical size of 72 A ˚ after milling for 12 h, and a slightly smaller size (68 A) was obtained after milling for 24 h. The two-stage evolution observed is the same as the one mentioned by Li et al. [10]. First, the milling introduces shear bands which are followed by fractures, therefore decreasing the size of the crystallites. When the particle size becomes smaller, the shear strength becomes higher, and at a critical point the shear band cannot pass through the crystallite. This will therefore prevent the size refinement, as was observed after 12 h of milling.
Table 2 sums up the cell parameters and average diffracting domain sizes measured for different milling intensities and times. After having milled the pre-alloy at intensity 7 for 1 h, with or without Mg 2 Ni, cell parameters ˚ respecmeasured were very similar: 7.055 and 7.057 A, tively. In order to determine if magnesium entered the pre-alloy cell during the milling, the following is considered. Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) and Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ) should have the same cell parameter, as Mg and Zr have the same atomic radius. Yet, Zr 1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ), milled for 1 h at intensity 7, has ˚ which is 0.007 A ˚ more than a cell parameter of 7.064 A, what was measured for Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled for 1 h at intensity 7 with Mg 2 Ni. This therefore proves that magnesium did not enter the pre-alloy cell after 1 h of milling at intensity 7. Regarding milling carried out at lower intensities, the cell parameters remained very similar to the pre-alloy cell parameter: these intensities are therefore too low to allow the magnesium to enter the cell.
3.2. EMPA First, it is to be remembered that, in this analysis, the grains studied measured more than 1 mm (because the size of the beam does not allow to analyze smaller grains) and that only the core of these grains could be analyzed with this technique. Regarding 1S3, magnesium was not detected and the pre-alloy composition was measured. It therefore showed that magnesium was not in the core of grains with sizes over 1 mm. This was confirmed by looking at the EDXA image of magnesium, as seen in Fig. 5. Indeed, the EDXA image of magnesium did not correspond with the image of the grains obtained by collecting the secondary electrons. Regarding 1S5, magnesium was detected in some grains and not in others, its distribution is very heterogeneous. Regarding alloys obtained by milling at intensity 7 for more than 1 h, magnesium was detected in all the grains, and Table 3 shows the measured composition. In order to quantify the heterogeneity of the element distributions in the alloys, the following ratio, R, was calculated: standard deviationof element X in wt.% R 5 ]]]]]]]]]]] average wt.% of element X
Table 2 Cell parameters and crystallite sizes of the alloys elaborated Sample
˚ a (A)
˚ Crystallite sizes (A)
Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ), non-milled Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled for 1 h (intensity 7) without Mg 2 Ni Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ) milled with Mg 2 Ni 1 h (intensity 3) 6 h (intensity 3) 1 h (intensity 5) 2 h 36 min (intensity 5) 1 h (intensity 7) Zr 1.0 (Mn 0.6 Ni 1.17 V0.2 Cr 0.1 ) milled 1 h (intensity 7) without Mg 2 Ni
7.054 7.057
1550 194
7.051 7.052 7.050 7.052 7.055 7.064
455 353 290 239 180 580
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Fig. 5. Grains of alloy milled for 1 h at intensity 3: (a) EDXA image of magnesium; (b) image of secondary electrons.
Fig. 7. Grains of alloy milled for 12 h at intensity 7: (a) EDXA image of magnesium; (b) image of secondary electrons.
Table 3 Compositions analyzed by EMPA
3.3. SEM analysis
Milling times (h) at intensity 7
Composition
1 3 6 12 24
Zr 0.92 Zr 0.93 Zr 0.90 Zr 0.92 Zr 0.90
Mg 0.07 (Mn 0.59 Ni 1.18 V0.22 Cr 0.11 ) Mg 0.10 (Mn 0.57 Ni 1.18 V0.21 Cr 0.10 ) Mg 0.15 (Mn 0.57 Ni 1.18 V0.21 Cr 0.10 ) Mg 0.07 (Mn 0.60 Ni 1.17 V0.20 Cr 0.10 ) Mg 0.11 (Mn 0.60 Ni 1.15 V0.20 Cr 0.10 )
The bigger R, the more heterogeneous is the distribution. This ratio was calculated for each element for different milling times at intensity 7, the result is shown in Fig. 6. Looking at the figure, it becomes obvious that the longer the milling, the more homogeneous the distribution of magnesium. Furthermore, the EDXA image of magnesium in sample 12S7 confirmed that magnesium was very well distributed in this alloy. Indeed, as seen in Fig. 7, the EDXA image of magnesium was the same as that obtained when the secondary electrons were collected. Besides, it is interesting to note that milling for more than 1 h also increased the homogeneity of the other elements, in comparison with the annealed pre-alloy.
Fig. 6. Heterogeneity of the elements as a function of milling time, at intensity 7.
Observing polished samples through SEM showed that grains were bigger as milling was longer and more intense, as presented in Fig. 8. Observing the powders glued on a carbon film showed that these grains had a rough surface which was actually due to the agglomeration of very fine particles, with sizes smaller than 1 mm (Fig. 9c Fig. 10b). This therefore showed that the longer and the more intensively the powders were milled, the more the fine
Fig. 8. SEM images of polished samples: (a) alloy 1S3; (b) alloy 12S7.
Fig. 9. SEM images of 1S3, as a powder: (a) unmilled pre-alloy grain; (b) rough grain; (c) zoom on the surface of the rough grain.
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Fig. 10. SEM images of 12S7, as a powder: (a) rough grain; (b) small agglomeration.
particles agglomerated. The size of these agglomerates was as large as several hundreds of micrometers. Regarding 1S3, not only those grains were observed, but unmilled pre-alloy grains were also detected. Their average size was around 10 mm and they were covered with fine particles, as shown in Fig. 9a. The more the samples were milled, the less unmilled pre-alloy grains were observed. EDX analysis on polished samples gave the same results as the one obtained with EMPA: magnesium was only detected on the surface of most of the grains in 1S3 and was homogeneously distributed in 12S7. Besides, EDX analysis on the powders showed that magnesium was detected on every grain surface, including the pre-alloy grains which were not crushed by the milling. Collecting back-scattered electrons with the SEM did not show any evidence of phase contrast. It therefore seems that magnesium was intimately mixed with the pre-alloy, on the surface at least.
3.4. TEM analysis TEM studies confirmed that the powders were constituted of aggregates of nano-particles, as can be seen in Fig. 11a Fig. 12b. The selected area diffraction pattern, shown
Fig. 12. TEM bright field image of a grain surface of 12S7.
in Fig. 11b, also confirmed the nano-crystallinity of the analyzed grains. In sample 1S3, EDX analysis showed that magnesium was not present in all the nano-particles. Besides, studying nano-particles containing magnesium, a gradient of magnesium concentration could be observed: the surface was magnesium rich (54 weight%) and the percentage decreased when approaching the core, reaching 3 wt.%. This weight percentage (3%) is actually the average weight percentage of magnesium measured with EMPA in the alloy Zr 0.9 Mg 0.1 (Mn 0.6 Ni 1.2 V0.2 Cr 0.1 ). EDX analysis also showed that Mg 2 Ni mostly did not remain as Mg 2 Ni in the nano-particles, and new phases appeared. The phases have yet to be identified. It is therefore appropriate to use the term ‘mechanical alloying’, even for the less intense and shortest millings. In sample 12S7, the magnesium concentration measured was homogeneous on a nanometer scale. Only the grain surfaces could be analyzed (the core was too thick) but the EDXA images of magnesium (Fig. 7a) lead us to conclude that magnesium was also homogeneously distributed inside the nano-particles.
4. Results and discussion of the hydrogen absorption studies
4.1. PCT isotherm measurements
Fig. 11. TEM images of 1S3: (a) bright field; (b) selected area diffraction.
PCT isotherm measurements were carried out at 343 K after three absorption–desorption cycles. Results obtained are presented in Fig. 13. Table 4 summarizes the absorption capacity of the samples milled with Mg 2 Ni for different times and intensities. The best absorption capacity was obtained for 1S3 (3.1 H / mol), which is the sample milled for the shortest time and at the lowest intensity. Furthermore, 12S7 exhibited the typical PCT isotherm of
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Fig. 13. PCT isotherms at 343 K of samples milled with Mg 2 Ni for different times and intensities: 1S3 (n), 1S7 (3), 3S7 (h), 12S7 (d). Table 4 Hydrogen absorption capacity of different samples Sample
Absorption capacity (H / mol)
1S3 1S7 3S7 12S7
3.06 2.77 2.45 2.29
an amorphous sample, with a poor absorption capacity (2.3 H / mol) and a plateau with a slope. It therefore showed that the less the samples were milled the better their absorption properties. This was attributed to the amorphization and disordering of the samples with milling. As in a study on hydrogen diffusion rates reported previously [11,12], our results confirmed that amorphization in C15 Laves phase compounds was altering hydrogen absorption properties. This is also true for kinetics as shown in Fig. 14: indeed, the more the samples were milled the more slowly their absorption took place. The influence of milling with and without Mg 2 Ni was also studied. The PCT isotherm measurements at 343 K after 29 absorption–desorption cycles are reported in Fig. 15 for samples of pre-alloy unmilled or milled, and with or without Mg 2 Ni. It shows that milling the pre-alloy alone
Fig. 14. Absorption kinetics at 343 K of: 1S5 (d), 1S7 (s), 3S7 (n) and 6S7 (3).
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Fig. 15. PCT isotherms at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (h), milled for 1 h at intensity 3 (m), milled for 1 h at intensity 3 with Mg 2 Ni (3).
for 1 h at intensity 3 decreased the absorption capacity of the pre-alloy from 2.7 to 2.6 H / mol. This reduction was due to disordering as mentioned in the previous paragraph. Besides, when the pre-alloy was milled with Mg 2 Ni for the same time and the same intensity, it was found that the absorption capacity of the pre-alloy was increased from 2.7 to 3.1 H / mol. Regarding the absorption kinetics, the results for the unmilled pre-alloy were in good agreement with the results Hsu and Perng obtained with similar compounds [3]. Furthermore, Fig. 16 shows that milling the pre-alloy alone reduced the absorption rate of hydrogen, whereas milling with Mg 2 Ni lead to the same absorption rate as the unmilled pre-alloy, with an increased absorption capacity. This showed the beneficial role played by the addition of Mg 2 Ni, regarding hydrogen absorption. It therefore seems that some new phases or interfaces generated by milling with Mg 2 Ni have a catalytic activity. As mentioned by Uchida et al. [13], the presence of Ni could be responsible for such enhancements. Yet, not only the kinetics are enhanced by the addition of Mg 2 Ni, but also the absorption capacity. Another interesting hypothesis, suggested by Fujii et al. [14], is that magnesium could
Fig. 16. Absorption kinetics at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (j), milled for 1 h at intensity 3 (s), milled for 1 h at intensity 3 with Mg 2 Ni (m).
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Fig. 17. Evolution of absorption capacities (after 17 h exposure to hydrogen) with cycling at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (m), milled for 1 h at intensity 3 (h), milled for 1 h at intensity 3 with Mg 2 Ni (s).
Fig. 18. Evolution of absorption capacities (after 1 h exposure to hydrogen) with cycling at 343 K of Zr 0.9 (Mn 0.6 Ni 1.15 V0.2 Cr 0.1 ): unmilled (n), milled for 1 h at intensity 3 (j), milled for 1 h at intensity 3 with Mg 2 Ni (s).
act as a binder and prevent the pre-alloy from oxidation. Further studies are planned in order to account for such enhancements.
¨ Zuttel et al. observed, using XPS, that after cycling ZrV0.4 Ni 1.6 (which is also a Zr-based AB 2 compound) its surface became Zr enriched, with most of the Zr oxidized, and also Ni deficient. The ZrO 2 layer and the Ni deficiency are both responsible for kinetic slowdown, as they prevent the dissociation of H 2 [13,16]. A similar phenomenon probably takes place with the unmilled pre-alloy, accounting for such a kinetic decrease in the first cycles. Complementary observations, such as XPS, are needed in order to be able to confirm such a hypothesis. The absorption kinetics of the milled samples did not vary as much: a 19% decrease was observed for the sample milled alone and a 14% decrease for the sample milled with Mg 2 Ni. In order to be able to explain properly the role played by the addition of Mg 2 Ni and ball milling, further investigations are needed. All we can say at present is that the magnesium-rich surface seems to enhance the kinetics. Indeed, as Hjort et al. reported [17], MgO x layers can increase hydrogen uptake rates. It is also interesting to note the advantage of milling with Mg 2 Ni, since it yielded the best average absorption capacity after 1 h exposure to hydrogen, over the first 29 cycles: 2.6 H / mol, as presented in Table 5. The average for the unmilled pre-alloy was 2.1 H / mol and for the pre-alloy milled alone it was 1.8 H / mol. If absorption after 17 h exposure to hydrogen is considered, 1S3 is still the best absorbing sample, but the difference with the two other samples is not that large: 2.9 H / mol for 1S3 and 2.8 H / mol for the unmilled pre-alloy. Even if magnesium did
4.2. The absorption evolution with cycling The absorption capacities of unmilled pre-alloy, of prealloy milled alone for 1 h at intensity 3, and pre-alloy milled with Mg 2 Ni for 1 h at intensity 3, were recorded at 343 K. They were then desorbed at 573 K for 1 h; a subsequent absorption was then carried out after the sample temperature was stabilized at 343 K. Twenty-nine absorption–desorption cycles were achieved following this procedure. Fig. 17 presents the absorption capacity evolutions after 17 h of exposure to hydrogen. It shows that the sample milled with Mg 2 Ni always kept a better absorption capacity than the other two. Another improvement due to the addition of Mg 2 Ni is in the evolution of the kinetics with cycling. Fig. 18 presents the evolution of absorption capacities of the same samples, but after only 1 h of exposure to hydrogen. As the absorption was not completed after 1 h exposure to hydrogen, these data will therefore yield information on the evolution of the kinetics with cycling. Regarding the unmilled pre-alloy, its absorption kinetics were slowed down in the first cycles. During the first 10 cycles, the absorption uptake in 1 h decreased by 43% (from 2.95 to 1.68 H / mol). This could be due to the fact that elements ¨ diffuse with cycling, as Zuttel et al. [15] reported. Indeed, Table 5 Average absorption capacities over the 29 absorption–desorption cycles
Average absorption capacity (H / mol):
Unmilled pre-alloy Pre-alloy milled alone for 1 h at intensity 3 Pre-alloy milled with Mg 2 Ni for 1 h at intensity 3
considering 1 h exposure to H 2
considering 17 h exposure to H 2
2.1 1.8 2.6
2.8 2.4 2.9
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not enter the cell of the pre-alloy, it increased its kinetics and absorption capacity at 343 K.
5. Conclusion From the structural data, we were not able to conclude if we succeeded in incorporating magnesium in the pre-alloy cell for samples milled over 1 h at intensity 7. However, these samples were too amorphous to be of practical interest regarding hydrogen absorption. Regarding samples milled at lower intensities, it was shown that magnesium did not enter the cell of the pre-alloy. Magnesium did, however, intimately mix with the other elements, as shown by the EDX analysis performed with TEM, and a gradient of magnesium concentration could be measured inside submicron grains. Furthermore, it was shown that magnesium did improve the hydrogen uptake rate as well as the absorption capacity. And this remained true even after 29 absorption–desorption cycles.
Acknowledgements The authors would like to express their thanks for financial support by the company Alcatel Alsthom Recherche.
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