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Size reduction in Mg rich intermetallics via hydrogen decrepitation
Journal of Alloys and Compounds xxx (2015) xxx–xxx
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Size reduction in Mg rich intermetallics via hydrogen decrepitation Alptekin Aydınlı, Burak Aktekin, Tayfur Öztürk ⇑ Dept. of Metallurgical and Materials Engineering, Middle East Technical University, Turkey
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Hydrogen decrepitation Size reduction Mg2Ni Mg2Cu
a b s t r a c t A study was carried out into hydrogen decrepitation of Mg rich intermetallics, namely Mg2Ni and Mg2Cu. These intermetallics are quite similar to each other except for the fact that Mg2Ni hydrides directly forming Mg2NiH4, whereas Mg2Cu when hydrided disproportionates into a two-phase structure. A total of ten sorption cycles was applied to the alloys and the resulting size reductions were monitored. The results showed that Mg2Ni decrepitate quite fast with cycling, the greatest size reduction occurring within the first three cycles. The size reduction in Mg2Cu, on the other hand, was quite sluggish. This was attributed to the disproportionation of the alloy which involve more extensive diffusion of the metallic species, counteracting some decrepitating effect of cycling due to ensuing particle sintering and growth. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Hydrogen decrepitation is an efficient size reduction method used in the material processing. The method can be used in metals, alloys and intermetallic compounds that react with hydrogen [1– 6]. It typically consists of several cycles of hydrogenation and dehydrogenation treatments. The size reduction that occurs as a result of this treatment originates from volume changes associated with hydriding and dehydriding reactions. The process of hydrogen decrepitation was of considerable interest in the 90s within the context of high performance magnets [2,3]. The so-called ‘‘HDDR’’ treatment leads to a substantial reduction in grain/particle size of sintered/bonded magnets with highly improved coercive properties. In this treatment the alloy, e.g., NdFeB, upon hydrogenation disproportionates into a multiphase structure which, when brought to elevated temperature under vacuum, dehydrogenates, and the phases recombine together re-forming the original phase with a drastic reduction in particle size. Size reduction brought about by hydrogen decrepitation was also of interest in hydrogen storage alloys [4–6]. With coarse powders, the reaction rate is often sluggish in such alloys. This makes it necessary to refine the powders. Thus, several cycles of hydrogenations are applied to the powders so as to improve their sorption kinetics. This is often referred as activation treatment where freshly generated surfaces arising from decrepitation reacts with hydrogen more easily. A more deliberate use of decrepitation treatment was made in the production of hydrogen storage tanks. Rather than processing ⇑ Corresponding author. Tel.: +90 312 210 5935; fax: +90 312 210 2518. E-mail address: [email protected] (T. Öztürk).
the powders to a fine size, the alloys are packed in storage tanks in the form of relatively coarse powders and they are then in-situ decrepitated by applying several cycles of hydrogenation and dehydrogenation. This ability of in-situ particle refinement is a considerable advantage in the production of hydrogen storage tanks. Although size reduction resulting from hydrogen decrepitation can be usefully employed to activate/improve the sorption kinetics, the same may also be associated with some adverse effects. With continued use, as a result of decrepitation, particles in storage tanks are more efficiently packed which leads to a creation of space between the storage material and the tank wall resulting in some loss of thermal conductivity [7]. A more serious adverse effect occurs in negative electrode material in alkaline batteries. Metal hydride in the negative electrode as it decrepitates may lead to enhanced corrosion and thus lead to a loss of capacity with cycling [8]. From a brief review above, it is apparent that the nature of size reduction as it occurs in cyclic hydrogenation is important not only as a way producing fine particle size, but also for developing reliable materials for extended use. The current study was therefore undertaken to investigate the decrepitation characteristics of two Mg rich intermetallics; Mg2Ni which hydrides directly and Mg2Cu which is subject to a more complex disproportionation–recombination reaction.
2. Experimental Mg2Ni was prepared from Mg (>99.8%, 325 mesh) and Ni (99.9%, 300 mesh) powders. They were first mixed in stoichiometric proportions and then pressed (250 MPa) into pellets of 30 mm diameter and approx. 10 mm height. The pellets were placed in a cold-crucible and melted under argon atmosphere (8 bar). Mg2Cu
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Please cite this article in press as: A. Aydınlı et al., J. Alloys Comp. (2015), http://dx.doi.org/10.1016/j.jallcom.2015.01.260
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A. Aydınlı et al. / Journal of Alloys and Compounds xxx (2015) xxx–xxx
was prepared from the bulk. For this purpose, pieces cut from Mg ingot (>99%) and an electrolytic copper rod (10 mm in diameter) were mixed in stoichiometric proportions and melted in a graphite crucible under argon atmosphere (8 bar). Having removed the alloys from the furnace they were first broken into pieces inside a plastic bag and then hand-crushed using a ceramic mortar. Decrepitation experiments were carried out in a Sievert type apparatus. For initial experiments, 1 g powder was used. This was loaded into the reactor which was then subjected to several vacuum-pressure (argon) cycles. The reactor was then charged with argon to 1 bar then heated up to 350 °C. Having reached the temperature, the reactor was charged to 35 bar hydrogen and allowed to cool down to room temperature. This first cycle ended when the reactor reached the room temperature. The second cycle started with re-heating the reactor to 350 °C. Having reached 350 °C, the reactor was let to cool down as before. This process of heating and cooling under hydrogen atmosphere, each with a rate of approximately 7 °C/ min, was repeated at least ten times. Values of pressure and temperature are monitored and pressure drop (hydriding) and pressure rise (dehydriding) values were determined. These refer to differences in pressure between room temperature and 350 °C during cooling and heating respectively. In order to follow size reduction of powders caused by the cycling, separate experiments were carried out using 12 g of alloy powder. These experiments were carried out in the same manner as above, except for the fact that having reached the room temperature, the pressure acting on the powders was noted and, the reactor was removed from the system. A sample of 1–1.5 g was removed from the main batch in a glove-box and the reactor, having been closed, was returned to the system. The system was then taken under vacuum at room temperature and charged with hydrogen to the same pressure and the experiment was continued. The cooling and heating rates were again 7 °C/min. Preliminary experiments with Mg2Cu showed that sorption reactions were not as easy to apply as was the case in Mg2Ni. Samples could be hydrogenated under hydrogen pressure of 35 bar as before, but the rate was rather slow. For this reason experiments were conducted with a slow rate together with a pressure change during absorption and desorption. Heating/cooling rate was 0.5 °C/min. Absorption was initiated at 350 °C by applying/increasing the pressure to 35 bar and letting the sample to cool down to room temperature. Desorption cycle was started at room temperature by decreasing the hydrogen pressure to 1 bar and heating the sample to 350 °C. Both hand-crushed and cycled samples were characterized with respect to their particle size using a multi-point BET analyzer. Samples were characterized structurally by X-ray diffraction using Cu Ka radiation and morphologically using a field emission SEM. Crystallite size in hydride and dehydrided samples were determined from X-ray patterns via Rietveld refinement [9]. The degree of fit in this refinement was always better than, Rw < 15.
3. Results and discussion Cast alloys, both Mg2Ni and Mg2Cu, were homogenous, but slightly away from the exact stoichiometry. XRD pattern recorded from Mg2Ni (see below) had the expected peaks plus a trace of MgNi2 peaks indicating that the alloy was slightly lean in its Mg content. In the case of Mg2Cu, reverse was the case. Thus, there was a small amount of free Mg (2.6 wt.%) in the alloy, i.e., the alloy was slightly richer in its Mg content. Hand crushing was quite easy with Mg2Ni as it pulverized easily producing a range of particle sizes. This was in contrast to Mg2Cu where a great effort was needed to pulverize the alloy with hand crushing. This was attributed to the presence of an Mg phase, which in the form of eutectics, envelop the intermetallics along the grain boundaries making the alloy less friable. Crushed powders were then sieved and particles between 140 and +170 mesh, i.e., 105 lm to +88 lm, were selected for the experiments. Powders were subjected to ten hydrogenation and dehydrogenation cycles. Fig. 1(a) shows pressure drop and pressure rise values determined as a function of cycling. It is seen that pressure change which was initially rather small improves with cycling, greatest increase occurring within the first three cycles. Samples taken from the starting powder and the powder after the tenth cycle, which was in hydrided form, were examined with X-ray diffraction. XRD patterns recorded from these samples are given Fig. 2(a) and (b) in the respective order. It is seen that the pattern from the starting powders are compatible with Mg2Ni phase. This phase was replaced by Mg2NiH4 in the hydrided powder as would be expected [10]. The patterns, do contain a trace amount of MgNi2 which, as pointed out above, indicate that the
Fig. 1. Variation of pressure change as a function of cycling (a) for Mg2Ni and (b) for Mg2Cu. Note that pressure change both in absorption and desorption increases with cycling. The change saturates after three-four cycle in Mg2Ni but it continues to increase in Mg2Cu without saturation.
alloy was slightly over stoichiometric. This phase was not affected by cycling. Size reduction brought about by cycling was followed both by morphological observations by SEM and by the surface area measurement by BET analyzer. SEM images before and after cycling are given in Fig. 3(a) and (b). It is seen that powders were decrepitated as a result of cycling. The appearance of SEM images is such that decrepitated particles are in the form of colonies that seem to have been derived from the starting powders. These colonies are quite large, have a crumbled form and seem to be made up of collection of smaller particles. Surface area values measured as a function of cycling are shown in Fig. 4. As pointed out above, these measurements were obtained from a separate experiment where a sample was removed from the reactor after each cycle. The variation of the surface area with cycling is quite similar to that depicted in Fig. 1 in that the greatest reduction occurs within the first 3–4 cycles. The surface area values after several cycling are at the level of 2.9–3 m2/g which corresponds to particle size values of approx. 700–800 nm. This size is quite compatible with the microstructure depicted in Fig. 3(b). Similar experiments were carried out with Mg2Cu. However, as given in the experimental section, heating and cooling were carried out with a much slower rate. Moreover pressure acting on the powders was changed before cooling and heating so as to sorb the powders successfully. Having determined the conditions of cycling, a continuous experiment was carried out for a total of 10 cycles. The XRD pattern of the powders before and after the cycling are given in Fig. 2(c) and (d) in the respective order. It is seen that the sample before cycling is effectively single phase Mg2Cu except for a trace amount of Mg which was present in the alloy, indicating that it was slightly rich in its Mg content. The sample after cycling is made up of two phases; MgH2 and MgCu2. The formation of these phases are compatible with a disproportionation reaction [11], i.e.,
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Fig. 2. X-ray diffractograms for (a) Mg2Ni, starting powder, (b) Mg2Ni, hydrided powder after the tenth cycle. Note that the starting powder is essentially Mg2Ni and the hydrided powder is Mg2NiH4. (c) Mg2Cu, starting powder, (d) Mg2Cu, hydrided after the tenth cycle, (e) sample in (d) after dehydrogenation. Note that the dehydrogenation is not complete in (e).
Fig. 3. SEM micrographs for; (a) Mg2Ni, starting powder, (b) Mg2Ni, hydrided powder after the tenth cycle, (c) Mg2Cu, starting powder and (d) Mg2Cu, hydrided powder after the tenth cycle. Note that original particle morphology in Mg2Cu is essentially the same before and after cycling.
Mg2 Cu þ H2 ¼ MgH2 þ MgCu2
Fig. 4. Variation of surface area of powders with cycling. Note that the decrease in the surface area is almost complete in Mg2Ni after the third cycle.
ð1Þ
SEM micrographs of the sample before and after cycling are given in Fig. 3(c) and (d) respectively. Here cycling seems to produce very little effect on the morphology of the Mg2Cu powders. Also it is seen that greater portion of powders remained as coarse as the starting material except for the fact that they appear to be rounded. Surface area values obtained with BET analysis are included in Fig. 4. Although there is some scatter in the data, it is seen the size reduction is not as pronounced as Mg2Ni. To investigate the dehydrogenation in Mg2Cu, a powder sample, after cycling, was subjected to a special dehydrogenation treatment. For this purpose after the tenth cycle, the hydrided powder was put into the reactor and heated up to 350 °C under 1 bar argon. Having reached the temperature the reactor was taken under vacuum (101 mbar) for a duration of 10 min. Then the reactor was
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charged with 5 bar argon and cooled down to room temperature. The sample was removed from the reactor and exposed to air. XRD pattern recorded from this sample is given in Fig. 2(e). The pattern contains Mg2Cu which is the expected phase, but there are also other phases in the pattern, namely MgCu2 and MgH2. Thus it appears that dehydrogenation of the powders was not complete. Since the conditions employed in this special treatment were more severe than those employed in the cycling, the same should also be true for powders during cycling. In fact, the pattern of pressure change in Fig. 1 implies that the powders even after tenth cycle have yet reached their full capacity. To investigate the (de)hydrogenation process in detail, the samples both in hydrogenated and dehydrogenated state were prepared metallographically by polishing and etching them using standard techniques. Fig. 5(a) is a section from the dehydrided sample where inner portion of the large fragment appears rather bright as compared to its surrounding. It is believed that the inner portion is the volume which had not dehydrogenated. MgH2 and MgCu2 observed in the XRD pattern probably originates from these inner portions. With continued cycling, it is expected that the volume of these inner core would reduce in size allowing powders to reach their full reversible capacity. No quantification was made with regard to particles containing such inner cores, but they were present almost in all particles larger than approx. 40 lm. Fig. 5(b) shows an example of a smaller fragment which had decrepitated quite successfully, individual grains loosened by this process has a size of around 5–6 lm. It is interesting to note that these small grains, though loosened from the fragments, rather than falling apart remain together in the form of colonies. Results reported above indicate that size reduction occurring in Mg2Ni and Mg2Cu are quite different from each other. Several sorption cycles are sufficient to refine Mg2Ni to submicron sizes, but the same is not true for Mg2Cu where the process appear to be more gradual, and not complete even after the tenth cycle. It may be pointed out that volume expansion/contraction that occur during hydriding/dehydriding is quite high in both intermetallics. Mg2Ni has a volume of 0.0519 nm3 per formula unit, the
value was derived from the unit cell volume reported in [12]. The corresponding value for Mg2NiH4 is 0.0685 nm3 [13]. This implies a volume expansion/contraction of 31.9%. Similar evaluation for Mg2Cu, using data reported in [14–16] for Mg2Cu, MgH2 and MgCu2 respectively yields a value of 24%. Though this value is less than that for Mg2Ni, it is unlikely that this would be the cause of the difference in the observed behavior. The sluggish decrepitation in Mg2Cu might have its origin in the fact this intermetallic follow a different pathway in hydriding/dehydriding reaction. The disproportionation which occurs during hydriding and recombination that follows during dehydriding are complex processes. Unlike Mg2Ni which hydrides directly through crystallographic rearrangements of atoms, metallic species in Mg2Cu during disproportionation and in MgH2 and MgCu2 during recombination are subject to extensive diffusion which could make the process rather slow. Thus the size reduction taking place as a result of decrepitation may be counteracted by re-sintering of particles and their growth due to a relatively more prolonged exposure of the alloy at elevated temperatures. To check this, X-ray diffractograms of Mg2Ni and Mg2Cu were Rietveld refined to obtain the crystallite size values. As seen in Table 1, in a hydrided sample, i.e., Fig. 2(b), the value is 62.3 nm for Mg2NiH4. In the hydrided Mg2Cu, Fig. 2(d), though the size of MgCu2 phase is comparable to the previous size, that of MgH2 is significantly larger, i.e., 91.8 nm. Thus, the crystallites in Mg2Cu seem to be somewhat larger than those in Mg2Ni which seem to support the above argument. It should be noted that crystallite size in Mg2Cu phase in the dehydrided sample is even higher, i.e., 154 nm. This phase forms upon dehydrogenation via recombination of MgH2 and MgCu2. The sluggish decrepitation of Mg2Cu may well have its origin in disproportionation–recombination processes, but clearly the fact that the cycling produced little changes in the morphology of the initial particles might have other influencing factors. In this context it should be mentioned that Mg2Cu is slightly substoichiometric, i.e., the alloy contain free Mg in its structure. This Mg is expected to be present in the form of eutectic phase which will
Fig. 5. SEM micrographs of cycled Mg2Cu powders following dehydriding treatment. Micrographs refer to sections in powders polished metallographically. (a) A Mg2Cu particle with an inner core and (b) a smaller Mg2Cu particle which had decrepitated into smaller grains.
Table 1 Crystallite size of phases in Mg2Ni and Mg2Cu after tenth cycle. Mg2Ni
Mg2Cu
Hydrided (Rw = 13.6)
Hydrided (Rw = 8.4)
Dehydrided (Rw = 6.5)
Phases
Mg2NiH4
MgNi2
MgH2
MgCu2
Mg2Cu
MgH2
MgCu2
Crystallite size (nm)
62.3
242.0
91.8
54.5
154
86.3
72.6
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solidify last and form a binding layer along the grain boundaries. It is possible that this binding layer could hold the decrepitated grains/fragments together keeping the original morphology of the particles. Whether the persistence of original particle morphology arises from in-situ re-sintering of the decrepitated particles or Mg rich eutectic layer wrapping the decrepitated particles together cannot be ascertained with the current data. A further study would be needed to clarify this issue in which the same intermetallics in slightly different compositions both in a sub- and over-stoichiometry could be examined with regard to their decrepitation behavior. 4. Conclusion In the current study, hydrogen decrepitation in Mg2Ni and Mg2Cu were studied by subjecting them up to ten sorption cycles. The followings can be concluded from the current study; 1. Mg2Ni decrepitate quite fast with cycling, the greatest size reduction occurring within the first three cycles. 2. Size reduction in Mg2Cu is quite gradual and, despite reduction, keeps the original particle morphology throughout the cycling. This sluggishness in the decrepitation of Mg2Cu may be attributed to the disproportionation of this alloy which involve more extensive diffusion of the metallic species counteracting some decrepitating effect of cycling due to ensuing particle sintering and growth.
5
Acknowledgements Support for this work was provided by TUBITAK (Project number 112M193) in conjunction with COST MP1103, which, we gratefully acknowledge. References [1] M. Okada, J. Nakahigashi, A. Fujita, M. Yamauchi, A. Kamegawa, J. Alloys Comp. 580 (2013) S401. [2] I.R. Harris, P.J. McGuiness, J. Less-Common Met. 172–174 (1991) 1273. [3] T. Takeshita, J. Alloys Comp. 231 (1995) 51. [4] S. Semboshi, N. Masahashi, S. Hanada, J. Alloys Comp. 359 (2003) 236. [5] S. Semboshi, N. Masahashi, S. Hanada, Met. Mater. Int. 10 (2004) 45. [6] N. Eliaz, D. Eliezer, D.L. Olson, Mater. Sci. Eng., A 289 (2000) 41. [7] A. Chaise, P. de Rango, P. Marty, D. Fruchart, S. Miraglia, R. Olivès, S. Garrier, Int. J. Hydrogen Energy 34 (2009) 8589. [8] A. Etiemble, H. Idrissi, L. Roué, J. Power Sources 196 (2011) 5168. [9] L. Lutterotti, S. Matthies, H.R. Wenk, Maud: a friendly java program for materials analysis using diffraction, Int. U. Crystallogr. Comm. Powder Diffr. Newslett. 21 (1999) 14 (MAUD version 2.072.). [10] J.J. Reilly, R.H. Wiswall, Inorg. Chem. 7 (1968) 2254. [11] J.J. Reilly, R.H. Wiswall, Inorg. Chem. 6 (1967) 2220. [12] M.L. Post, J.J. Murray, G.J.G. Despault, J.B. Taylor, Mater. Res. Bull. 20 (1985) 337. [13] D. Noréus, P.E. Werner, Mater. Res. Bull. 16 (1981) 199. [14] F. Gingl, P. Selvam, K. Yvon, Acta Crystallogr. B 49 (1993) 201. [15] S. Ono, Y. Ishido, K. Imanari, T. Tabata, Y.K. Cho, R. Yamamoto, M. Doyama, J. Less-Common Met. 88 (1982) 57. [16] T. Ohba, Y. Kitano, Y. Komura, Acta Crystallogr. C 40 (1984) 1.
Please cite this article in press as: A. Aydınlı et al., J. Alloys Comp. (2015), http://dx.doi.org/10.1016/j.jallcom.2015.01.260
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Size reduction in Mg rich intermetallics via hydrogen decrepitation Alptekin Aydınlı, Burak Aktekin, Tayfur Öztürk ⇑ Dept. of Metallurgical and Materials Engineering, Middle East Technical University, Turkey
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Hydrogen decrepitation Size reduction Mg2Ni Mg2Cu
a b s t r a c t A study was carried out into hydrogen decrepitation of Mg rich intermetallics, namely Mg2Ni and Mg2Cu. These intermetallics are quite similar to each other except for the fact that Mg2Ni hydrides directly forming Mg2NiH4, whereas Mg2Cu when hydrided disproportionates into a two-phase structure. A total of ten sorption cycles was applied to the alloys and the resulting size reductions were monitored. The results showed that Mg2Ni decrepitate quite fast with cycling, the greatest size reduction occurring within the first three cycles. The size reduction in Mg2Cu, on the other hand, was quite sluggish. This was attributed to the disproportionation of the alloy which involve more extensive diffusion of the metallic species, counteracting some decrepitating effect of cycling due to ensuing particle sintering and growth. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Hydrogen decrepitation is an efficient size reduction method used in the material processing. The method can be used in metals, alloys and intermetallic compounds that react with hydrogen [1– 6]. It typically consists of several cycles of hydrogenation and dehydrogenation treatments. The size reduction that occurs as a result of this treatment originates from volume changes associated with hydriding and dehydriding reactions. The process of hydrogen decrepitation was of considerable interest in the 90s within the context of high performance magnets [2,3]. The so-called ‘‘HDDR’’ treatment leads to a substantial reduction in grain/particle size of sintered/bonded magnets with highly improved coercive properties. In this treatment the alloy, e.g., NdFeB, upon hydrogenation disproportionates into a multiphase structure which, when brought to elevated temperature under vacuum, dehydrogenates, and the phases recombine together re-forming the original phase with a drastic reduction in particle size. Size reduction brought about by hydrogen decrepitation was also of interest in hydrogen storage alloys [4–6]. With coarse powders, the reaction rate is often sluggish in such alloys. This makes it necessary to refine the powders. Thus, several cycles of hydrogenations are applied to the powders so as to improve their sorption kinetics. This is often referred as activation treatment where freshly generated surfaces arising from decrepitation reacts with hydrogen more easily. A more deliberate use of decrepitation treatment was made in the production of hydrogen storage tanks. Rather than processing ⇑ Corresponding author. Tel.: +90 312 210 5935; fax: +90 312 210 2518. E-mail address: [email protected] (T. Öztürk).
the powders to a fine size, the alloys are packed in storage tanks in the form of relatively coarse powders and they are then in-situ decrepitated by applying several cycles of hydrogenation and dehydrogenation. This ability of in-situ particle refinement is a considerable advantage in the production of hydrogen storage tanks. Although size reduction resulting from hydrogen decrepitation can be usefully employed to activate/improve the sorption kinetics, the same may also be associated with some adverse effects. With continued use, as a result of decrepitation, particles in storage tanks are more efficiently packed which leads to a creation of space between the storage material and the tank wall resulting in some loss of thermal conductivity [7]. A more serious adverse effect occurs in negative electrode material in alkaline batteries. Metal hydride in the negative electrode as it decrepitates may lead to enhanced corrosion and thus lead to a loss of capacity with cycling [8]. From a brief review above, it is apparent that the nature of size reduction as it occurs in cyclic hydrogenation is important not only as a way producing fine particle size, but also for developing reliable materials for extended use. The current study was therefore undertaken to investigate the decrepitation characteristics of two Mg rich intermetallics; Mg2Ni which hydrides directly and Mg2Cu which is subject to a more complex disproportionation–recombination reaction.
2. Experimental Mg2Ni was prepared from Mg (>99.8%, 325 mesh) and Ni (99.9%, 300 mesh) powders. They were first mixed in stoichiometric proportions and then pressed (250 MPa) into pellets of 30 mm diameter and approx. 10 mm height. The pellets were placed in a cold-crucible and melted under argon atmosphere (8 bar). Mg2Cu
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was prepared from the bulk. For this purpose, pieces cut from Mg ingot (>99%) and an electrolytic copper rod (10 mm in diameter) were mixed in stoichiometric proportions and melted in a graphite crucible under argon atmosphere (8 bar). Having removed the alloys from the furnace they were first broken into pieces inside a plastic bag and then hand-crushed using a ceramic mortar. Decrepitation experiments were carried out in a Sievert type apparatus. For initial experiments, 1 g powder was used. This was loaded into the reactor which was then subjected to several vacuum-pressure (argon) cycles. The reactor was then charged with argon to 1 bar then heated up to 350 °C. Having reached the temperature, the reactor was charged to 35 bar hydrogen and allowed to cool down to room temperature. This first cycle ended when the reactor reached the room temperature. The second cycle started with re-heating the reactor to 350 °C. Having reached 350 °C, the reactor was let to cool down as before. This process of heating and cooling under hydrogen atmosphere, each with a rate of approximately 7 °C/ min, was repeated at least ten times. Values of pressure and temperature are monitored and pressure drop (hydriding) and pressure rise (dehydriding) values were determined. These refer to differences in pressure between room temperature and 350 °C during cooling and heating respectively. In order to follow size reduction of powders caused by the cycling, separate experiments were carried out using 12 g of alloy powder. These experiments were carried out in the same manner as above, except for the fact that having reached the room temperature, the pressure acting on the powders was noted and, the reactor was removed from the system. A sample of 1–1.5 g was removed from the main batch in a glove-box and the reactor, having been closed, was returned to the system. The system was then taken under vacuum at room temperature and charged with hydrogen to the same pressure and the experiment was continued. The cooling and heating rates were again 7 °C/min. Preliminary experiments with Mg2Cu showed that sorption reactions were not as easy to apply as was the case in Mg2Ni. Samples could be hydrogenated under hydrogen pressure of 35 bar as before, but the rate was rather slow. For this reason experiments were conducted with a slow rate together with a pressure change during absorption and desorption. Heating/cooling rate was 0.5 °C/min. Absorption was initiated at 350 °C by applying/increasing the pressure to 35 bar and letting the sample to cool down to room temperature. Desorption cycle was started at room temperature by decreasing the hydrogen pressure to 1 bar and heating the sample to 350 °C. Both hand-crushed and cycled samples were characterized with respect to their particle size using a multi-point BET analyzer. Samples were characterized structurally by X-ray diffraction using Cu Ka radiation and morphologically using a field emission SEM. Crystallite size in hydride and dehydrided samples were determined from X-ray patterns via Rietveld refinement [9]. The degree of fit in this refinement was always better than, Rw < 15.
3. Results and discussion Cast alloys, both Mg2Ni and Mg2Cu, were homogenous, but slightly away from the exact stoichiometry. XRD pattern recorded from Mg2Ni (see below) had the expected peaks plus a trace of MgNi2 peaks indicating that the alloy was slightly lean in its Mg content. In the case of Mg2Cu, reverse was the case. Thus, there was a small amount of free Mg (2.6 wt.%) in the alloy, i.e., the alloy was slightly richer in its Mg content. Hand crushing was quite easy with Mg2Ni as it pulverized easily producing a range of particle sizes. This was in contrast to Mg2Cu where a great effort was needed to pulverize the alloy with hand crushing. This was attributed to the presence of an Mg phase, which in the form of eutectics, envelop the intermetallics along the grain boundaries making the alloy less friable. Crushed powders were then sieved and particles between 140 and +170 mesh, i.e., 105 lm to +88 lm, were selected for the experiments. Powders were subjected to ten hydrogenation and dehydrogenation cycles. Fig. 1(a) shows pressure drop and pressure rise values determined as a function of cycling. It is seen that pressure change which was initially rather small improves with cycling, greatest increase occurring within the first three cycles. Samples taken from the starting powder and the powder after the tenth cycle, which was in hydrided form, were examined with X-ray diffraction. XRD patterns recorded from these samples are given Fig. 2(a) and (b) in the respective order. It is seen that the pattern from the starting powders are compatible with Mg2Ni phase. This phase was replaced by Mg2NiH4 in the hydrided powder as would be expected [10]. The patterns, do contain a trace amount of MgNi2 which, as pointed out above, indicate that the
Fig. 1. Variation of pressure change as a function of cycling (a) for Mg2Ni and (b) for Mg2Cu. Note that pressure change both in absorption and desorption increases with cycling. The change saturates after three-four cycle in Mg2Ni but it continues to increase in Mg2Cu without saturation.
alloy was slightly over stoichiometric. This phase was not affected by cycling. Size reduction brought about by cycling was followed both by morphological observations by SEM and by the surface area measurement by BET analyzer. SEM images before and after cycling are given in Fig. 3(a) and (b). It is seen that powders were decrepitated as a result of cycling. The appearance of SEM images is such that decrepitated particles are in the form of colonies that seem to have been derived from the starting powders. These colonies are quite large, have a crumbled form and seem to be made up of collection of smaller particles. Surface area values measured as a function of cycling are shown in Fig. 4. As pointed out above, these measurements were obtained from a separate experiment where a sample was removed from the reactor after each cycle. The variation of the surface area with cycling is quite similar to that depicted in Fig. 1 in that the greatest reduction occurs within the first 3–4 cycles. The surface area values after several cycling are at the level of 2.9–3 m2/g which corresponds to particle size values of approx. 700–800 nm. This size is quite compatible with the microstructure depicted in Fig. 3(b). Similar experiments were carried out with Mg2Cu. However, as given in the experimental section, heating and cooling were carried out with a much slower rate. Moreover pressure acting on the powders was changed before cooling and heating so as to sorb the powders successfully. Having determined the conditions of cycling, a continuous experiment was carried out for a total of 10 cycles. The XRD pattern of the powders before and after the cycling are given in Fig. 2(c) and (d) in the respective order. It is seen that the sample before cycling is effectively single phase Mg2Cu except for a trace amount of Mg which was present in the alloy, indicating that it was slightly rich in its Mg content. The sample after cycling is made up of two phases; MgH2 and MgCu2. The formation of these phases are compatible with a disproportionation reaction [11], i.e.,
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Fig. 2. X-ray diffractograms for (a) Mg2Ni, starting powder, (b) Mg2Ni, hydrided powder after the tenth cycle. Note that the starting powder is essentially Mg2Ni and the hydrided powder is Mg2NiH4. (c) Mg2Cu, starting powder, (d) Mg2Cu, hydrided after the tenth cycle, (e) sample in (d) after dehydrogenation. Note that the dehydrogenation is not complete in (e).
Fig. 3. SEM micrographs for; (a) Mg2Ni, starting powder, (b) Mg2Ni, hydrided powder after the tenth cycle, (c) Mg2Cu, starting powder and (d) Mg2Cu, hydrided powder after the tenth cycle. Note that original particle morphology in Mg2Cu is essentially the same before and after cycling.
Mg2 Cu þ H2 ¼ MgH2 þ MgCu2
Fig. 4. Variation of surface area of powders with cycling. Note that the decrease in the surface area is almost complete in Mg2Ni after the third cycle.
ð1Þ
SEM micrographs of the sample before and after cycling are given in Fig. 3(c) and (d) respectively. Here cycling seems to produce very little effect on the morphology of the Mg2Cu powders. Also it is seen that greater portion of powders remained as coarse as the starting material except for the fact that they appear to be rounded. Surface area values obtained with BET analysis are included in Fig. 4. Although there is some scatter in the data, it is seen the size reduction is not as pronounced as Mg2Ni. To investigate the dehydrogenation in Mg2Cu, a powder sample, after cycling, was subjected to a special dehydrogenation treatment. For this purpose after the tenth cycle, the hydrided powder was put into the reactor and heated up to 350 °C under 1 bar argon. Having reached the temperature the reactor was taken under vacuum (101 mbar) for a duration of 10 min. Then the reactor was
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charged with 5 bar argon and cooled down to room temperature. The sample was removed from the reactor and exposed to air. XRD pattern recorded from this sample is given in Fig. 2(e). The pattern contains Mg2Cu which is the expected phase, but there are also other phases in the pattern, namely MgCu2 and MgH2. Thus it appears that dehydrogenation of the powders was not complete. Since the conditions employed in this special treatment were more severe than those employed in the cycling, the same should also be true for powders during cycling. In fact, the pattern of pressure change in Fig. 1 implies that the powders even after tenth cycle have yet reached their full capacity. To investigate the (de)hydrogenation process in detail, the samples both in hydrogenated and dehydrogenated state were prepared metallographically by polishing and etching them using standard techniques. Fig. 5(a) is a section from the dehydrided sample where inner portion of the large fragment appears rather bright as compared to its surrounding. It is believed that the inner portion is the volume which had not dehydrogenated. MgH2 and MgCu2 observed in the XRD pattern probably originates from these inner portions. With continued cycling, it is expected that the volume of these inner core would reduce in size allowing powders to reach their full reversible capacity. No quantification was made with regard to particles containing such inner cores, but they were present almost in all particles larger than approx. 40 lm. Fig. 5(b) shows an example of a smaller fragment which had decrepitated quite successfully, individual grains loosened by this process has a size of around 5–6 lm. It is interesting to note that these small grains, though loosened from the fragments, rather than falling apart remain together in the form of colonies. Results reported above indicate that size reduction occurring in Mg2Ni and Mg2Cu are quite different from each other. Several sorption cycles are sufficient to refine Mg2Ni to submicron sizes, but the same is not true for Mg2Cu where the process appear to be more gradual, and not complete even after the tenth cycle. It may be pointed out that volume expansion/contraction that occur during hydriding/dehydriding is quite high in both intermetallics. Mg2Ni has a volume of 0.0519 nm3 per formula unit, the
value was derived from the unit cell volume reported in [12]. The corresponding value for Mg2NiH4 is 0.0685 nm3 [13]. This implies a volume expansion/contraction of 31.9%. Similar evaluation for Mg2Cu, using data reported in [14–16] for Mg2Cu, MgH2 and MgCu2 respectively yields a value of 24%. Though this value is less than that for Mg2Ni, it is unlikely that this would be the cause of the difference in the observed behavior. The sluggish decrepitation in Mg2Cu might have its origin in the fact this intermetallic follow a different pathway in hydriding/dehydriding reaction. The disproportionation which occurs during hydriding and recombination that follows during dehydriding are complex processes. Unlike Mg2Ni which hydrides directly through crystallographic rearrangements of atoms, metallic species in Mg2Cu during disproportionation and in MgH2 and MgCu2 during recombination are subject to extensive diffusion which could make the process rather slow. Thus the size reduction taking place as a result of decrepitation may be counteracted by re-sintering of particles and their growth due to a relatively more prolonged exposure of the alloy at elevated temperatures. To check this, X-ray diffractograms of Mg2Ni and Mg2Cu were Rietveld refined to obtain the crystallite size values. As seen in Table 1, in a hydrided sample, i.e., Fig. 2(b), the value is 62.3 nm for Mg2NiH4. In the hydrided Mg2Cu, Fig. 2(d), though the size of MgCu2 phase is comparable to the previous size, that of MgH2 is significantly larger, i.e., 91.8 nm. Thus, the crystallites in Mg2Cu seem to be somewhat larger than those in Mg2Ni which seem to support the above argument. It should be noted that crystallite size in Mg2Cu phase in the dehydrided sample is even higher, i.e., 154 nm. This phase forms upon dehydrogenation via recombination of MgH2 and MgCu2. The sluggish decrepitation of Mg2Cu may well have its origin in disproportionation–recombination processes, but clearly the fact that the cycling produced little changes in the morphology of the initial particles might have other influencing factors. In this context it should be mentioned that Mg2Cu is slightly substoichiometric, i.e., the alloy contain free Mg in its structure. This Mg is expected to be present in the form of eutectic phase which will
Fig. 5. SEM micrographs of cycled Mg2Cu powders following dehydriding treatment. Micrographs refer to sections in powders polished metallographically. (a) A Mg2Cu particle with an inner core and (b) a smaller Mg2Cu particle which had decrepitated into smaller grains.
Table 1 Crystallite size of phases in Mg2Ni and Mg2Cu after tenth cycle. Mg2Ni
Mg2Cu
Hydrided (Rw = 13.6)
Hydrided (Rw = 8.4)
Dehydrided (Rw = 6.5)
Phases
Mg2NiH4
MgNi2
MgH2
MgCu2
Mg2Cu
MgH2
MgCu2
Crystallite size (nm)
62.3
242.0
91.8
54.5
154
86.3
72.6
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solidify last and form a binding layer along the grain boundaries. It is possible that this binding layer could hold the decrepitated grains/fragments together keeping the original morphology of the particles. Whether the persistence of original particle morphology arises from in-situ re-sintering of the decrepitated particles or Mg rich eutectic layer wrapping the decrepitated particles together cannot be ascertained with the current data. A further study would be needed to clarify this issue in which the same intermetallics in slightly different compositions both in a sub- and over-stoichiometry could be examined with regard to their decrepitation behavior. 4. Conclusion In the current study, hydrogen decrepitation in Mg2Ni and Mg2Cu were studied by subjecting them up to ten sorption cycles. The followings can be concluded from the current study; 1. Mg2Ni decrepitate quite fast with cycling, the greatest size reduction occurring within the first three cycles. 2. Size reduction in Mg2Cu is quite gradual and, despite reduction, keeps the original particle morphology throughout the cycling. This sluggishness in the decrepitation of Mg2Cu may be attributed to the disproportionation of this alloy which involve more extensive diffusion of the metallic species counteracting some decrepitating effect of cycling due to ensuing particle sintering and growth.
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Acknowledgements Support for this work was provided by TUBITAK (Project number 112M193) in conjunction with COST MP1103, which, we gratefully acknowledge. References [1] M. Okada, J. Nakahigashi, A. Fujita, M. Yamauchi, A. Kamegawa, J. Alloys Comp. 580 (2013) S401. [2] I.R. Harris, P.J. McGuiness, J. Less-Common Met. 172–174 (1991) 1273. [3] T. Takeshita, J. Alloys Comp. 231 (1995) 51. [4] S. Semboshi, N. Masahashi, S. Hanada, J. Alloys Comp. 359 (2003) 236. [5] S. Semboshi, N. Masahashi, S. Hanada, Met. Mater. Int. 10 (2004) 45. [6] N. Eliaz, D. Eliezer, D.L. Olson, Mater. Sci. Eng., A 289 (2000) 41. [7] A. Chaise, P. de Rango, P. Marty, D. Fruchart, S. Miraglia, R. Olivès, S. Garrier, Int. J. Hydrogen Energy 34 (2009) 8589. [8] A. Etiemble, H. Idrissi, L. Roué, J. Power Sources 196 (2011) 5168. [9] L. Lutterotti, S. Matthies, H.R. Wenk, Maud: a friendly java program for materials analysis using diffraction, Int. U. Crystallogr. Comm. Powder Diffr. Newslett. 21 (1999) 14 (MAUD version 2.072.
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