- Email: [email protected]
Effect of amorphous Mg50Ni50 on hydriding and dehydriding behavior of Mg2Ni alloy
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
available at www.sciencedirect.com
www.elsevier.com/locate/matchar
Effect of amorphous Mg50Ni50 on hydriding and dehydriding behavior of Mg2Ni alloy D. Guzmána,⁎, S. Ordoñezb , J.F. Fernándezc , C. Sánchezc , D. Serafinid , P.A. Rojase , C. Aguilar f , P. Tapiag a
Departamento de Ingeniería en Metalurgia, Facultad de Ingeniería, Universidad de Atacama y Centro Regional de Investigación y Desarrollo Sustentable de Atacama (CRIDESAT), Av. Copayapu 485, Copiapó, Chile b Departamento de Ingeniería Metalúrgica, Facultad de Ingeniería, Universidad de Santiago de Chile, Av. Lib. Bernardo O'Higgins 3363, Santiago, Chile c Departamento de Física de Materiales, Facultad de Ciencias, Universidad Autónoma de Madrid, Cantoblanco 28049, Madrid, Spain d Departamento de Física, Facultad de Ciencias, Universidad de Santiago de Chile and Center for Interdisciplinary Research in Materials, CIMAT, Av. Lib. Bernardo O'Higgins 3363, Santiago, Chile e Escuela de Ingeniería Mecánica, Facultad de Ingeniería, Av. Los Carrera 01567, Quilpué, Pontificia Universidad Católica de Valparaíso, PUCV, Chile f Departamento de Ingeniería Metalúrgica y Materiales, Universidad Técnica Federico Santa María, Av. España 1680, Valparaíso, Chile g Departamento de Ingeniería en Metalurgia, Facultad de Ingeniería, Universidad de Atacama, Av. Copayapu 485, Copiapó, Chile
AR TIC LE D ATA
ABSTR ACT
Article history:
Composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 was prepared by mechanical milling
Received 23 September 2010
starting with nanocrystalline Mg2Ni and amorphous Mg50Ni50 powders, by using a SPEX 8000
Received in revised form
D mill. The morphological and microstructural characterization of the powders was
14 February 2011
performed via scanning electron microscopy and X-ray diffraction. The hydriding
Accepted 23 February 2011
characterization of the composite was performed via a solid gas reaction method in a Sievert's-type apparatus at 363 K under an initial hydrogen pressure of 2 MPa. The
Keywords:
dehydriding behavior was studied by differential thermogravimetry. On the basis of the
Hydrogen storage materials
results, it is possible to conclude that amorphous Mg50Ni50 improved the hydriding and
Amorphous materials
dehydriding kinetics of Mg2Ni alloy upon cycling. A tentative rationalization of experimental
Mechanical alloying
observations is proposed. © 2011 Elsevier Inc. All rights reserved.
Hydriding and dehydriding processes X-ray diffraction
1.
Introduction
Hydrogen storage in metal hydrides offers a high storage density per unit volume and is relatively safer than other hydrogen storage methods, such as the compression and liquefaction of hydrogen gas and absorption by carbon [1]. Magnesium is a
suitable material for hydrogen storage because of its light weight, high hydrogen capacity (7.6 wt.% H for pure MgH2), and low cost. Nevertheless, MgH2 is very stable (equilibrium pressure at 552 K is 0.1 MPa), and its hydrogen absorption and desorption kinetics are considerably slow [2,3] for technological applications. Several attempts have been made to overcome these limitations by
⁎ Corresponding author. Tel.: +56 52206646; fax: +56 52206619. E-mail address: [email protected] (D. Guzmán). 1044-5803/$ – see front matter © 2011 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2011.02.007
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
alloying magnesium with other elements [4–7], producing magnesium-based composites [8–11], and modifying its morphology and microstructure [12]. Two stable intermetallics, Mg2Ni and MgNi2, have been reported in the Mg–Ni system [13]. The former absorbs hydrogen in a solid solution up to Mg2NiH0.3 and forms a hydride, Mg2NiH4, at higher hydrogen concentrations. The hydride has a lower storage capacity than MgH2, but faster hydriding and dehydriding kinetics [4,14,15]. Below 503 K, it crystallizes in a monoclinic lowtemperature phase, transforming to a cubic high-temperature phase above that temperature [16–18]. On the other hand, MgNi2 does not react with hydrogen at pressures up to 2.8 MPa and a temperature of 623 K [4]. In relation to nonequilibrium hydrides, Orimo et al. [19–21] worked on the amorphous Mg50Ni50–H2 system. The results obtained in these investigations showed that the amorphous phase can absorb approximately 2.2 wt.% H without changing its amorphous structure. On the other hand, Guzmán et al. [22] reported the formation of oversaturated solid solution in the nanocrystalline Mg2Ni–H2 system (corresponding to Mg2NiH2.2). The mechanical alloying (MA) process developed by Benjamin et al. [23,24] in the early 1970s has been widely used to produce hydrogen storage materials, particularly Mg-based alloys. Considering the large differences in melting points and
Fig. 1 – XRD patterns of Mg and Ni in 2:1 atomic proportion: (a) milled for 10 h and (b) milled as (a) and annealed at 673 K for 0.5 h.
443
vapor pressures between Mg and Ni, the production of Mg–Ni based alloys via the conventional casting method is more difficult than the MA process [25–27]. In addition, MA produces a microstructural refinement [19,20,22,28–31] and increases the specific surface area [32], which improve the hydriding and dehydriding kinetics [33–35]. Thus, amorphous Mg50Ni50, produced by MA, has showed a marked improvement in the hydriding and dehydriding kinetics compared to Mg2Ni and ball-milled Mg2Ni [28,35,36]. Although many researches have been performed on the hydriding and dehydriding behavior of the nanocrystalline/ amorphous composites in the Mg–Ni system [19,21,22,29, 30,36,37], the interaction between amorphous and nanocrystalline phases is not yet completely understood. In this study, it is proposed to establish the effect of the amorphous Mg50Ni50 on the hydriding and dehydriding properties of nanocrystalline Mg2Ni.
2.
Experimental Procedures
2.1.
Preparation of Composite Mg2Ni–Amorphous Mg50Ni50
On the basis of the results obtained in previous studies [25–27], intermetallic Mg2Ni was synthesized by combining the MA and heat treatment. The MA process of Mg turnings (98 wt.% pure, particle size below 4.7 mm; Aldrich) and Ni powders
Fig. 2 – XRD patterns of Mg and Ni in 1:1 atomic proportion: (a) initial mixture and (b) milled for 20 h.
444
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
(99.99 wt.% pure, particle size below 61 μm; Fluka) in a 2:1 atomic proportion was carried out in a SPEX 8000 D mill for 10 h, under a controlled atmosphere of Ar (99.999 wt.% pure) using a ball to powder weight ratio of 20:1. The heat treatment was accomplished under vacuum at 673 K for 0.5 h. According to the data reported by Guzmán et al. [38], amorphous Mg50Ni50 was produced via MA of an equiatomic mixture of Mg turnings and Ni powders in a SPEX 8000 D mill for 20 h with a ball to powder weight ratio of 20:1 under inert Ar atmosphere. In order to produce the composite material, a mixture consisting of 75 wt.% Mg2Ni and 25 wt.% amorphous Mg50Ni50 was milled for 10 min in a SPEX 8000 D mill under inert Ar atmosphere using a ball to powder weight ratio of 10:1. For comparison, a Mg2Ni sample was obtained by the previously mentioned procedure but without adding any amorphous phase. This sample will help to establish the effect of the amorphous phase on the hydriding and dehydriding proper-
ties of Mg2Ni without any influence of the microstructural refinement of Mg2Ni during the composite synthesis (milling).
2.2.
Characterization
The powders produced were microstructurally characterized by X-ray diffraction (XRD) in a Shimadzu XRD-600 diffractometer using Cu Kα radiation. The average crystallite size was calculated using the Scherrer method [39]. In order to obtain the full width at half maximum peak intensity, the experimental peaks were fitted using a Lorentz function. The sizes, morphologies, and phase distributions of the powders were studied using a Jeol 5410 scanning electron microscope (SEM) equipped with X-ray energy dispersive spectroscopy (EDS). In order to determine the particle size distribution, four micrographs were recorded using approximately 400 measurements in each analysis. Hydrogen absorption of the composite was investigated in a Sievert's-type apparatus at 363 K and an initial hydrogen
Fig. 3 – SEM micrographs obtained in SE mode, showing general and magnified views of (a) Mg2Ni and (b) amorphous Mg50Ni50 powders. Particle size distributions are showed in each case.
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
pressure of 2 MPa, whereas hydrogen desorption was investigated by thermogravimetry under flowing N2, using a heating rate of 5 K min− 1 in a SDT 2960 TA instrument. Thus, the apparent activation energy for the hydrogen desorption process was obtained via the Kissinger method [40] using heating rates of 5, 10, 15, and 20 K min− 1.
3.
Results and Discussion
3.1.
Microstructural and Morphological Characterization
Fig. 1 (a) shows the XRD pattern of a mixture of Mg and Ni in a 2:1 atomic proportion milled for 10 h in a SPEX mill. It can be seen that the sample consisted of a mixture of Mg2Ni and residual Ni with an average crystallite size of 11 and 15 nm, respectively. Additional Mg needed to maintain the milling stoichiometry formed an amorphous phase, which could be subsequently transformed in a mechanical and/or thermal manner into Mg2Ni [25–27]. Fig. 1 (b) shows the XRD pattern of a sample annealed at 673 K for 0.5 h under vacuum. As can be seen, the sample consisted mainly of Mg2Ni and a small amount of Ni because of the loss of Mg by the MgO formation. The reason for the decrease of the Ni content can be understood by considering that the amorphous transformation into Mg2Ni (423 K) and the solid reaction between Mg and residual Ni to form Mg2Ni (473–523 K) occur below 673 K [25–27]. After annealing (673 K), the Mg2Ni crystallite size increased to 30 nm. Further details about the Mg2Ni production via MA and MA and heat treatment can be found in Ref. [27]. Fig. 2 (a) shows the XRD pattern of the initial mixture of Mg and Ni in a 1:1 atomic proportion, whereas Fig. 2 (b) shows the XRD pattern of the powders after 20 h of milling. The milling conditions employed were sufficient to produce an amorphous Mg50Ni50 alloy. The latter result is in agreement with a previous report [38] in which more information about phase evolution during the amorphization process can be obtained. Fig. 3 shows SEM images of Mg2Ni and amorphous Mg50Ni50 obtained in the secondary electron mode (SE). It can be seen that Mg2Ni was composed of large irregular agglomerates with an apparent diameter of approximately 30 μm and several smaller particles with flake-like shapes around 1.5 μm welded together (Fig. 3 (a)). On the other hand, the amorphous Mg50Ni50 was composed of large irregular agglomerates with a maximum apparent size of 500 μm and small round particles with an apparent size of 3 μm (Fig. 3 (b)). A previous work [29] confirmed that the microstructural refinement in the nanometric range of Mg2Ni by milling increases its hydriding and dehydriding kinetics compared with crystalline Mg2Ni. The latter behavior can be understood by considering the decrease in the diffusion path and the increase in the inter-grain region because of the appearance of the nanocrystalline structure. In order to eliminate the influence of the microstructural refinement of Mg2Ni during the composite synthesis (milling) and establish the effects of the amorphous phase on the hydriding and dehydriding properties of Mg2Ni, the results obtained from the composite were compared with those obtained from a Mg2Ni sample milled for 10 min.
445
Fig. 4 (a) shows the XRD pattern of composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 produced by milling. The amorphous phase was detected in the XRD pattern due to an increase in the diffraction intensity over the background level in the regions close to 2θ ≈ 20, 40, and 65° [38]. The average crystallite size of Mg2Ni in the composite was 25 nm. On the other hand, Fig. 4 (b) shows the XRD pattern of the Mg2Ni prepared for comparison. In this case, the calculated crystallite size of Mg2Ni was 23 nm. On the basis of the results, it is possible to conclude that Mg2Ni in the composite and the milled Mg2Ni presented similar microstructural refinements; therefore, any difference in the hydriding and dehydriding properties between the milled Mg2Ni and the composite material might be related to the presence of the amorphous Mg50Ni50 phase. Fig. 5 (a) shows the SEM image of the composite material obtained in the SE mode. It can be seen that the particle size distribution of the composite is similar to that of the Mg2Ni alloy (Fig. 3 (a)). EDS analysis revealed that the brighter particles in the backscattered electron view (BSE) image (Fig. 5 (b)) correspond to amorphous Mg50Ni50 and the darker areas to the Mg2Ni matrix (Fig. 5 (c)). On the basis of the results, it is possible to conclude that the milling time employed (10 min) in the composite synthesis was sufficient to produce an intimate contact between Mg2Ni and the amorphous Mg50Ni50 particles. Finally, another important conclusion from the XRD (Fig. 4) and SEM (Fig. 5) results is that no alloy formation (between Mg2Ni and the amorphous phase) occurred during milling.
Fig. 4 – XRD patterns of (a) Mg2Ni (25 wt.%) amorphous Mg50Ni50 composite and (b) milled Mg2Ni.
446
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
Fig. 6 – Hydriding kinetics of (a) composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 and (b) milled Mg2Ni at 363 K under an initial hydrogen pressure of 2 MPa.
3.2.
Hydriding and Dehydriding Characterization
The hydriding and dehydriding characteristics of the prepared composite were studied and compared to that of the milled Mg2Ni to establish the possible utility of the composite for hydrogen storage. The samples were activated after one hydriding (initial hydrogen pressure of 2 MPa at 363 K) and dehydriding (under vacuum of 27 Pa at 453 K) cycle. The dehydriding temperature was selected to avoid the crystallization process of amorphous Mg50Ni50 that starts at approximately 573 K [41]. The low equilibrium H pressure of Mg2Ni below 473 K [42] precludes the ability to obtain the desorption data by the volumetric method. In this first hydriding cycle, the Mg2Ni sample absorbed approximately 1.60 wt.% H, whereas the composite material only absorbed 1.47 wt.% H. Fig. 6 shows the hydrogen absorption curves of the composite and Mg2Ni samples during four hydriding cycles.
Fig. 5 – SEM micrographs showing (a) SE view and (b) BSE of composite material (arrows indicate the position of amorphous Mg50Ni50). (c) Elemental EDS analysis patterns of the selected area in (b).
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
The reversible hydrogen storage capacity and hydriding kinetics of the composite material were greater than those of Mg2Ni. The initial hydrogen absorption rate for both samples showed no obvious dependence on the hydriding cycles. Thus, the average initial rate of hydrogen absorption for the composite material (0.60 ± 0.06 wt.% H min− 1) was approximately four times higher than that for the Mg2Ni sample (0.15 ± 0.02 wt.% H min− 1). On the basis of the results, it is possible to conclude that the amorphous phase improved the hydriding kinetics and hydrogen storage capacity of Mg2Ni upon cycling under the experimental conditions. The latter behavior could be attributed to two factors. First, it has been reported that on the surface of amorphous Mg50Ni50, Mg is combined with oxygen whereas Ni is present in metallic state [43,44]. This Ni, which formed enriched layers, served as an effective catalyst for dissociating hydrogen molecules. Second, Mg–Ni based amorphous alloys showed superior oxidation resistances to their crystalline counterparts [45], which, added to the existence of metallic Ni underneath the top surface, provided “passageways” to the composite for diffusing hydrogen atoms in atmospheres containing oxygen. Fig. 7 shows XRD patterns of the composite and Mg2Ni samples after five hydriding cycles. It is possible to observe that both materials consisted of a mixture of Mg2NiH4, Mg2NiH0.3, residual Ni, and MgO. The only difference between the two XRD patterns was the increase in the diffraction intensity over the
Fig. 7 – XRD patterns of hydrogenated (a) composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 and (b) milled Mg2Ni.
447
background level in the regions close to 2θ ≈ 20, 40, and 65° in the composite pattern (Fig. 7 (a)), which proved the presence of the hydrogenated amorphous phase [46]. Dehydrogenation of the samples was studied by thermogravimetric analysis after the last hydriding cycle. Fig. 8 (a) shows the results. The total hydrogen mass losses for the milled Mg2Ni (1.6 wt.%) and composite (1.4 wt.%) samples were similar to the amount of hydrogen absorbed during the first hydriding cycle, demonstrating that only partial dehydrogenation was achieved in the volumetric experiments under the experimental conditions. A slight rise of the thermogravimetric curves in the high temperature region was due to the oxidation of the alloys by the residual oxygen absorbed on the alloy surfaces. Iwakura et al. previously reported the latter behavior in Mg–Ni based alloys [37]. After five cycles, the hydrogen storage capacities of the milled Mg2Ni and composite samples were, respectively, 47% and 73% of their initial capacities. Considering that both samples were subjected to some desorption conditions and that the capacity loss could be attributed principally to partial
Fig. 8 – (a) Thermogravimetric and (b) derivative thermogravimetric profiles of the hydrogenated samples obtained at 5 K/min under flowing N2.
448
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
dehydrogenation, it is possible to conclude that the amorphous phase improved the desorption kinetics of hydrogenated Mg2Ni. Hydrogen absorption can produce nanocrystallization in the amorphous phase [47,48]. This fact could worsen the hydriding and dehydriding kinetics of the composite. A detailed investigation of how the hydrogen absorption process affects amorphous Mg50Ni50 stability is beyond the scope of this work. However, on the basis of the results of XRD of the hydrogenated composite (Fig. 7(a)), it is possible to establish that during the hydriding and dehydriding cycles, a complete crystallization of the amorphous phase did not occur. Fig. 8 (b) shows the derivative thermogravimetric curves of the milled Mg2Ni and composite materials. Instead of two separate desorption steps (desorption of the hydrogenated amorphous phase and Mg2NiH4), the entire composite released hydrogen in a single principal process. The temperature of the maximum hydrogen desorption rate of the composite material (481 K) was lower than that of the milled Mg2Ni (525 K). Furthermore, the starting temperature for the desorption of the composite (375 K) was similar to those reported by Orimo et al. [20] (373 K) and Guzmán et al. [46] (373 K) for the desorption of hydrogenated amorphous Mg50Ni50. On the basis of the results, it is possible to conclude that during the desorption process of the composite material, a catalytic effect of hydrogenated amorphous Mg50Ni50 on the hydrogen desorption of Mg2NiH4 was observed. Zaluska et al. [49] previously observed a similar behavior in the MgH2– Mg2NiH4 system. Catalytic desorption could be explained owing to the interaction between neighboring particles of Mg2NiH4 and hydrogenated amorphous Mg50Ni50. Thus, when hydrogenated amorphous Mg50Ni50 begins to desorb hydrogen, it undergoes a contraction that applies a strain to the attached Mg2NiH4 and facilitates its desorption. In addition, other dehydriding reactions could be observed in the derivative thermogravimetric curves of Mg2Ni and the composite material around 585 K (Fig. 8 (b)). This reaction was unaffected by the existence of the amorphous phase. According to Selvan et al. [50,51], this desorption process corresponded to the dehydriding of Mg2NiH, which is obtained according to the following equations: Mg2 NiH4 →Mg2 NiH þ 1:5 H2
ðaround 540–570 KÞ
Mg2 NiH→Mg2 NiH0:3 þ 0:35 H2
ðaround 560–590 KÞ:
The apparent activation energy for the principal dehydriding process was estimated with the Kissinger method [40]. According to the Kissinger theory, the slope of the linear plot of ln(β TP− 2) versus TP− 1 is the apparent activation energy of the process, where β is the heating rate and TP is the peak temperature. Fig. 9 shows the Kissinger plot of the dehydriding process of Mg2Ni and the composite materials. The apparent activation energy for the hydrogen desorption of the composite material (48 ± 6 kJ mol− 1) was lower than the calculated value for the dehydriding process of Mg2Ni (85 ± 4 kJ mol− 1). The latter results confirmed the existence of a catalytic effect during the composite desorption. On the other hand, the apparent activation energies for hydrogen desorption of the composite and Mg2Ni samples were lower than that reported by Selvan et al. [50] for the desorption of Mg2NiH4 (183 kJ mol− 1). The
Fig. 9 – Plot of ln(β TP− 2) versus 1000 TP− 1 for dehydriding of milled Mg2Ni and composite Mg2Ni (25 wt.%) amorphous Mg50Ni50.
discrepancy between the values reported by Selvan et al. [50] and those obtained in this study can be understood considering the differences in the experimental procedures. Selvan et al. [50] employed thermogravimetric analysis under a hydrogen pressure of 101.3 kPa, whereas in this study, the thermogravimetric analyses were carried out under flowing N2.
4.
Conclusions
In order to improve the hydriding and dehydriding properties of Mg2Ni, a composite of Mg2Ni with 25 wt.% of amorphous Mg50Ni50 was prepared via mechanical milling. On the basis of the results, it is possible to conclude that amorphous Mg50Ni50 improved the hydrogen storage capacity and hydriding kinetics of Mg2Ni alloy upon cycling. In addition, a catalytic effect of amorphous Mg50Ni50 on the hydrogen desorption of Mg2NiH4 was observed. Thus, it is proposed that when hydrogenated amorphous Mg50Ni50 begins to desorb hydrogen, it undergoes a contraction that applies a strain to the attached Mg2NiH4 and facilitates its desorption. Finally, although the addition of amorphous Mg50Ni50 improved the hydrogen absorption and desorption kinetics of Mg2Ni and its cyclability at low temperatures, the technical targets established by the U.S. Department of Energy (DOE) [52] for transport applications have not been achieved. The main problem of the composites based on Mg2Ni is the high stability of Mg2NiH4, which has an equilibrium H pressure of 91 Pa at 358 K (maximum delivery temperature recommended by DOE), which is very low for mobile applications.
Acknowledgments The authors gratefully acknowledge “Fondo Nacional Desarrollo Científico y Tecnológico de Chile”, FONDECYT, proyect No. 1070085 and “Comisión Nacional de Investigación Científica y Tecnológica de Chile”, CONICYT, for their economic support granted to realize this work. Some authors thank the
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
Spanish Ministry of Education and Science, MEC, for its financial support under contract No. MAT2008-06547-C02-01.
REFERENCES
[1] Züttel A. Materials for hydrogen storage. Mater Today 2003: 24–33. [2] Selvan P, Viswanathan B, Swamy CS, Srinivasan V. Magnesium and magnesium alloy hydrides. Int J Hydrogen Energ 1986;11: 169–92. [3] Fernández JF, Sánchez CR. Rate determining step in the absorption and desorption of hydrogen by magnesium. J Alloy Comp 2002;340:189–98. [4] Reilly JJ, Wiswall RH. Reaction of hydrogen with alloys of magnesium and nickel and the formation of Mg2NiH4. Inorg Chem 1968;7:2254–6. [5] Reilly JJ, Wiswall RH. The reaction of hydrogen with alloys of magnesium and copper. Inorg Chem 1967;6:2220–3. [6] Zaluski L, Zaluska A, Ström-Olsen J. Hydrogen absorption in nannocrystalline Mg2Ni formed by mechanical alloying. J Alloy Comp 1995;217:245–9. [7] Sahlberg M, Andersson Y. Hydrogen absorption in Mg–Y–Zn ternary compounds. J Alloy Comp 2007;446–447:134–7. [8] Fernández JF, Bodega J, Sánchez CR. Hydriding/dehydriding properties of magnesium–ZrCr2 composites. J Alloy Comp 2003;356–357:343–7. [9] Hanada N, Ichikawa T, Fujii H. Hydrogen absorption kinetics of the catalyzed MgH2 by niobium oxide. J Alloy Comp 2007;446–447: 67–71. [10] Wang P, Wang AM, Zhang HF, Ding BZ, Hu ZQ. Direct hydrogenation of Mg under the action of FeTi1.2 (amorphous) and mechanical driving force. J Mater Sci Lett 2001;20:753–4. [11] Wang P, Wang AM, Ding BZ, Hu ZQ. Mg–FeTi1.2 (amorphous) composite for hydrogen storage. J Alloy Comp 2002;334:243–8. [12] Zaluska A, Zaluski L, Ström-Olsen J. Structure, catalysis and atomic reactions on the nano-scale: a systematic approach to metal hydrides for hydrogen storage. Appl Phys A 2001;72:157–65. [13] Massalski T, Okamoto H, Subramanian P, Kacprzak L. Binary alloy phase diagram. 2nd ed. Metals Park, OH: American for Metals; 1990. [14] Nohara S, Inoe H, Fukumoto Y, Iwakura C. Compositional and structural characteristics of MgNi alloy prepared by mechanical alloying for use as negative electrodes in nickel-metal hydride batteries. J Alloy Comp 1997;259:183–5. [15] Abdellaoui M, Gracco D, Percheron-Guegan A. Structural characterization and reversible hydrogen absorption properties of Mg2Ni rich nanocomposite materials synthesized by mechanical alloying. J Alloy Comp 1998;268: 233–40. [16] Gavra Z, Mintz MH, Kimmel G, Hadari Z. Allotropic transitions of magnesium nickel hydride (Mg2NiH4). Inorg Chem 1979;18: 3595–7. [17] Noréus D, Werner PE. Structural studies of monoclinic Mg2NiD4. J Less Common Met 1984;97:215–22. [18] Soubeyroux JL, Fruchart D, Mikou A, Pezat M, Darriet B. Etude structurale du systeme Mg2Ni–H2 III La variete monoclinique de Mg2NiH4. Mater Res Bull 1984;19:1119–28. [19] Orimo S, Fujii H. Effects of nanometer-scale structure on hydriding properties of Mg–Ni alloys: a review. Intermetallics 1998;6:185–92. [20] Orimo S, Fujii H. Materials science of Mg–Ni-based new hydrides. Appl Phys A 2001;72:167–86. [21] Orimo S, Ikeda K, Fujii H, Fujikawa Y, Kitano Y, Yamamoto K. Structural and hydriding properties of Mg–Ni–H system with nano and/or amorphous structures. Acta Mater 1997;45: 2271–8.
449
[22] Guzmán D, Ordoñez S, Fernández JF, Sánchez C, Serafini D, Rojas PA, et al. Indications of the formation of an oversaturated solid solution during hydrogenation of Mg–Ni based nanocomposite produced by mechanical alloying. Int J Hydrogen Energ 2009;34:5429–38. [23] Benjamin J. Dispersion strengthened superalloys by mechanical alloying. Metall Trans 1970;1:2943–51. [24] Benjamin J, Volin T. The mechanism of mechanical alloying. Metall Trans 1974;5:1929–34. [25] Ordoñez S, Rojas P, San Martín A, Serafini D, Bustos O. Stoichiometric quantification of the amorphous precursor in the Mg–Ni system. Mater Sci Forum 2003;416–418:138–43. [26] Ordoñez S, Rojas P, Bustos O, Martínez V, Serafíni D, San Martín A. Crystalline Mg2Ni obtained from Mg–Ni amorphous precursor produced by mechanical alloying. J Mater Sci Lett 2003;22:717–20. [27] Rojas P, Ordoñez S, Serafini D, Zúñiga A, Lavernia E. Microstructural evolution during mechanical alloying of Mg and Ni. J Alloy Comp 2005;391:267–76. [28] Orimo S, Fujii H. Hydriding properties of the Mg2Ni–H system synthesized by reactive mechanical grinding. J Alloy Comp 1996;232:L16–9. [29] OrimoS,FujjiH,IkedaK.Notablehydridingpropertiesofa nanostructuredcomposite materialoftheMg2Ni–Hsystem synthesizedbyreactivemechanicalgrinding.ActaMater1997;45: 331–41. [30] Fujii H, Orimo S, Ikeda K. Cooperative hydriding properties in a nanostructured Mg2Ni–H system. J Alloy Comp 1997;253–254: 80–3. [31] Zaluska L, Zaluski A, Ström-Olsen J. Nanocrystalline magnesium for hydrogen storage. J Alloy Comp 1999;288:217–25. [32] Huot J, Liang G, Schulz R. Mechanically alloyed metal hydride systems. Appl Phys A 2001;72:187–95. [33] Mulas G, Delogu F, Cocco G. Effects of mechanical processing on the kinetics of H2 adsorption in Mg2Ni alloys. J Alloy Comp 2009;473:180–4. [34] Ruggeri S, Lenain C, Roué L, Liang G, Huot J, Schulz R. Mechanically driven crystallization of amorphous MgNi alloy during prolonged milling: applications in Ni–MH batteries. J Alloy Comp 2002;339:195–201. [35] Vijay R, Sundaresan R, Maiya MP, Murthy SS. Comparative evaluation of Mg–Ni hydrogen absorbing materials prepared by mechanical alloying. Int J Hydrogen Energ 2005;30: 501–8. [36] Anik M. Electrochemical hydrogen storage capacities of Mg2Ni and MgNi alloys synthesized by mechanical alloying. J Alloy Comp 2010;491:565–70. [37] Iwakura C, Nohara S, Zhang S, Inoue H. Hydriding and dehydriding characteristics of an amorphous Mg2Ni–Ni composite. J Alloy Comp 1999;285:246–9. [38] Guzmán D, Ordoñez S, Serafini D, Rojas P, Bustos O. Effect of the milling energy on the production and thermal stability of amorphous Mg50Ni50. J Alloy Comp 2009;471: 435–41. [39] Cullity B. Elements of X ray diffraction. 2nd ed. USA: Addison-Wesley Publishing Co.; 1978. [40] Kissinger HE. Reaction kinetics in differential thermal analysis. Anal Chem 1957;29:1702–6. [41] Guzmán D, Ordoñez S, Serafini D, Rojas PA, Aguilar C, 2Santander M. Thermal stability of amorphous Mg50Ni50 alloy produced by mechanical alloying. J Non-Cryst Solids 2010;356: 120–3. [42] Fukai Y. The metal–hydrogen system. USA: Springer-Verlag; 1993. [43] Zhang S, Yorimitsu K, Nohara S, Morikawa T, Inoue H, Iwakura C. Surface analysis of an amorphous MgNi alloy prepared by mechanical alloying for use in nickel-metal hydride batteries. J Alloy Comp 1998;270:123–6.
450
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
[44] Zhang S, Hara Y, Morikawa T, Inoue H, Iwakura C. Electrochemical and structural characteristics of amorphous MgNix (x≥1) alloys prepared by mechanical alloying. J Alloy Comp 1999;293–295:552–5. [45] Zhang Y, Dong X, Guo S, Wang G, Ren J, Wang X. The cycle stabilities of the as-cast and quenched La2Mg(Ni0.85Co0.15)9Mx (M=B, Cr, Ti; x=0, 0.1) hydrogen storage alloys. J Alloy Comp 2005;298:178–83. [46] D. Guzmán, Estudio de la fabricación y comportamiento frente a la hidrogenación de aleaciones Mg–Ni, Doctoral Thesis Universidad de Santiago de Chile, Chile, 2008. [47] Li L, Wu DC, Liang GY, Sun ZB, Guo YL. The dehydrogenation process of amorphous (Mg60Ni25)92La8 alloy. J Alloy Comp 2009;474:378–81.
[48] Huang LJ, Liang GY, Sun ZB, Zhou YF. Nanocrystallization and hydriding properties of amorphous mel-spun Mg65Cu25Nd10 alloy. J Alloy Comp 2007;432:171–6. [49] Zaluska A, Zaluski L, Ström-Olsen J. Synergy of hydrogen sorption in ball-milled hydrides of Mg and Mg2Ni. J Alloy Comp 1999;298:197–206. [50] Selvam P, Viswanathan B, Swamy CS, Srinivasan V. Thermal studies on Mg2NiH4: existence of additional hydride phase in the Mg2Ni–Hydrogen system. Thermochim Acta 1988;125:1–8. [51] Selvam P, Viswanathan B, Swamy CS, Srinivasan V. Mg2NiH: a new hydride phase in the Mg2Ni–H2 system. Int J Hydrogen Energ 1988;13:749–59. [52] Website:http://www1.eere.energy.gov/hydrogenandfuelcells/ storage/pdfs/targets_onboard_hydro_storage.pdf.
available at www.sciencedirect.com
www.elsevier.com/locate/matchar
Effect of amorphous Mg50Ni50 on hydriding and dehydriding behavior of Mg2Ni alloy D. Guzmána,⁎, S. Ordoñezb , J.F. Fernándezc , C. Sánchezc , D. Serafinid , P.A. Rojase , C. Aguilar f , P. Tapiag a
Departamento de Ingeniería en Metalurgia, Facultad de Ingeniería, Universidad de Atacama y Centro Regional de Investigación y Desarrollo Sustentable de Atacama (CRIDESAT), Av. Copayapu 485, Copiapó, Chile b Departamento de Ingeniería Metalúrgica, Facultad de Ingeniería, Universidad de Santiago de Chile, Av. Lib. Bernardo O'Higgins 3363, Santiago, Chile c Departamento de Física de Materiales, Facultad de Ciencias, Universidad Autónoma de Madrid, Cantoblanco 28049, Madrid, Spain d Departamento de Física, Facultad de Ciencias, Universidad de Santiago de Chile and Center for Interdisciplinary Research in Materials, CIMAT, Av. Lib. Bernardo O'Higgins 3363, Santiago, Chile e Escuela de Ingeniería Mecánica, Facultad de Ingeniería, Av. Los Carrera 01567, Quilpué, Pontificia Universidad Católica de Valparaíso, PUCV, Chile f Departamento de Ingeniería Metalúrgica y Materiales, Universidad Técnica Federico Santa María, Av. España 1680, Valparaíso, Chile g Departamento de Ingeniería en Metalurgia, Facultad de Ingeniería, Universidad de Atacama, Av. Copayapu 485, Copiapó, Chile
AR TIC LE D ATA
ABSTR ACT
Article history:
Composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 was prepared by mechanical milling
Received 23 September 2010
starting with nanocrystalline Mg2Ni and amorphous Mg50Ni50 powders, by using a SPEX 8000
Received in revised form
D mill. The morphological and microstructural characterization of the powders was
14 February 2011
performed via scanning electron microscopy and X-ray diffraction. The hydriding
Accepted 23 February 2011
characterization of the composite was performed via a solid gas reaction method in a Sievert's-type apparatus at 363 K under an initial hydrogen pressure of 2 MPa. The
Keywords:
dehydriding behavior was studied by differential thermogravimetry. On the basis of the
Hydrogen storage materials
results, it is possible to conclude that amorphous Mg50Ni50 improved the hydriding and
Amorphous materials
dehydriding kinetics of Mg2Ni alloy upon cycling. A tentative rationalization of experimental
Mechanical alloying
observations is proposed. © 2011 Elsevier Inc. All rights reserved.
Hydriding and dehydriding processes X-ray diffraction
1.
Introduction
Hydrogen storage in metal hydrides offers a high storage density per unit volume and is relatively safer than other hydrogen storage methods, such as the compression and liquefaction of hydrogen gas and absorption by carbon [1]. Magnesium is a
suitable material for hydrogen storage because of its light weight, high hydrogen capacity (7.6 wt.% H for pure MgH2), and low cost. Nevertheless, MgH2 is very stable (equilibrium pressure at 552 K is 0.1 MPa), and its hydrogen absorption and desorption kinetics are considerably slow [2,3] for technological applications. Several attempts have been made to overcome these limitations by
⁎ Corresponding author. Tel.: +56 52206646; fax: +56 52206619. E-mail address: [email protected] (D. Guzmán). 1044-5803/$ – see front matter © 2011 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2011.02.007
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
alloying magnesium with other elements [4–7], producing magnesium-based composites [8–11], and modifying its morphology and microstructure [12]. Two stable intermetallics, Mg2Ni and MgNi2, have been reported in the Mg–Ni system [13]. The former absorbs hydrogen in a solid solution up to Mg2NiH0.3 and forms a hydride, Mg2NiH4, at higher hydrogen concentrations. The hydride has a lower storage capacity than MgH2, but faster hydriding and dehydriding kinetics [4,14,15]. Below 503 K, it crystallizes in a monoclinic lowtemperature phase, transforming to a cubic high-temperature phase above that temperature [16–18]. On the other hand, MgNi2 does not react with hydrogen at pressures up to 2.8 MPa and a temperature of 623 K [4]. In relation to nonequilibrium hydrides, Orimo et al. [19–21] worked on the amorphous Mg50Ni50–H2 system. The results obtained in these investigations showed that the amorphous phase can absorb approximately 2.2 wt.% H without changing its amorphous structure. On the other hand, Guzmán et al. [22] reported the formation of oversaturated solid solution in the nanocrystalline Mg2Ni–H2 system (corresponding to Mg2NiH2.2). The mechanical alloying (MA) process developed by Benjamin et al. [23,24] in the early 1970s has been widely used to produce hydrogen storage materials, particularly Mg-based alloys. Considering the large differences in melting points and
Fig. 1 – XRD patterns of Mg and Ni in 2:1 atomic proportion: (a) milled for 10 h and (b) milled as (a) and annealed at 673 K for 0.5 h.
443
vapor pressures between Mg and Ni, the production of Mg–Ni based alloys via the conventional casting method is more difficult than the MA process [25–27]. In addition, MA produces a microstructural refinement [19,20,22,28–31] and increases the specific surface area [32], which improve the hydriding and dehydriding kinetics [33–35]. Thus, amorphous Mg50Ni50, produced by MA, has showed a marked improvement in the hydriding and dehydriding kinetics compared to Mg2Ni and ball-milled Mg2Ni [28,35,36]. Although many researches have been performed on the hydriding and dehydriding behavior of the nanocrystalline/ amorphous composites in the Mg–Ni system [19,21,22,29, 30,36,37], the interaction between amorphous and nanocrystalline phases is not yet completely understood. In this study, it is proposed to establish the effect of the amorphous Mg50Ni50 on the hydriding and dehydriding properties of nanocrystalline Mg2Ni.
2.
Experimental Procedures
2.1.
Preparation of Composite Mg2Ni–Amorphous Mg50Ni50
On the basis of the results obtained in previous studies [25–27], intermetallic Mg2Ni was synthesized by combining the MA and heat treatment. The MA process of Mg turnings (98 wt.% pure, particle size below 4.7 mm; Aldrich) and Ni powders
Fig. 2 – XRD patterns of Mg and Ni in 1:1 atomic proportion: (a) initial mixture and (b) milled for 20 h.
444
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
(99.99 wt.% pure, particle size below 61 μm; Fluka) in a 2:1 atomic proportion was carried out in a SPEX 8000 D mill for 10 h, under a controlled atmosphere of Ar (99.999 wt.% pure) using a ball to powder weight ratio of 20:1. The heat treatment was accomplished under vacuum at 673 K for 0.5 h. According to the data reported by Guzmán et al. [38], amorphous Mg50Ni50 was produced via MA of an equiatomic mixture of Mg turnings and Ni powders in a SPEX 8000 D mill for 20 h with a ball to powder weight ratio of 20:1 under inert Ar atmosphere. In order to produce the composite material, a mixture consisting of 75 wt.% Mg2Ni and 25 wt.% amorphous Mg50Ni50 was milled for 10 min in a SPEX 8000 D mill under inert Ar atmosphere using a ball to powder weight ratio of 10:1. For comparison, a Mg2Ni sample was obtained by the previously mentioned procedure but without adding any amorphous phase. This sample will help to establish the effect of the amorphous phase on the hydriding and dehydriding proper-
ties of Mg2Ni without any influence of the microstructural refinement of Mg2Ni during the composite synthesis (milling).
2.2.
Characterization
The powders produced were microstructurally characterized by X-ray diffraction (XRD) in a Shimadzu XRD-600 diffractometer using Cu Kα radiation. The average crystallite size was calculated using the Scherrer method [39]. In order to obtain the full width at half maximum peak intensity, the experimental peaks were fitted using a Lorentz function. The sizes, morphologies, and phase distributions of the powders were studied using a Jeol 5410 scanning electron microscope (SEM) equipped with X-ray energy dispersive spectroscopy (EDS). In order to determine the particle size distribution, four micrographs were recorded using approximately 400 measurements in each analysis. Hydrogen absorption of the composite was investigated in a Sievert's-type apparatus at 363 K and an initial hydrogen
Fig. 3 – SEM micrographs obtained in SE mode, showing general and magnified views of (a) Mg2Ni and (b) amorphous Mg50Ni50 powders. Particle size distributions are showed in each case.
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
pressure of 2 MPa, whereas hydrogen desorption was investigated by thermogravimetry under flowing N2, using a heating rate of 5 K min− 1 in a SDT 2960 TA instrument. Thus, the apparent activation energy for the hydrogen desorption process was obtained via the Kissinger method [40] using heating rates of 5, 10, 15, and 20 K min− 1.
3.
Results and Discussion
3.1.
Microstructural and Morphological Characterization
Fig. 1 (a) shows the XRD pattern of a mixture of Mg and Ni in a 2:1 atomic proportion milled for 10 h in a SPEX mill. It can be seen that the sample consisted of a mixture of Mg2Ni and residual Ni with an average crystallite size of 11 and 15 nm, respectively. Additional Mg needed to maintain the milling stoichiometry formed an amorphous phase, which could be subsequently transformed in a mechanical and/or thermal manner into Mg2Ni [25–27]. Fig. 1 (b) shows the XRD pattern of a sample annealed at 673 K for 0.5 h under vacuum. As can be seen, the sample consisted mainly of Mg2Ni and a small amount of Ni because of the loss of Mg by the MgO formation. The reason for the decrease of the Ni content can be understood by considering that the amorphous transformation into Mg2Ni (423 K) and the solid reaction between Mg and residual Ni to form Mg2Ni (473–523 K) occur below 673 K [25–27]. After annealing (673 K), the Mg2Ni crystallite size increased to 30 nm. Further details about the Mg2Ni production via MA and MA and heat treatment can be found in Ref. [27]. Fig. 2 (a) shows the XRD pattern of the initial mixture of Mg and Ni in a 1:1 atomic proportion, whereas Fig. 2 (b) shows the XRD pattern of the powders after 20 h of milling. The milling conditions employed were sufficient to produce an amorphous Mg50Ni50 alloy. The latter result is in agreement with a previous report [38] in which more information about phase evolution during the amorphization process can be obtained. Fig. 3 shows SEM images of Mg2Ni and amorphous Mg50Ni50 obtained in the secondary electron mode (SE). It can be seen that Mg2Ni was composed of large irregular agglomerates with an apparent diameter of approximately 30 μm and several smaller particles with flake-like shapes around 1.5 μm welded together (Fig. 3 (a)). On the other hand, the amorphous Mg50Ni50 was composed of large irregular agglomerates with a maximum apparent size of 500 μm and small round particles with an apparent size of 3 μm (Fig. 3 (b)). A previous work [29] confirmed that the microstructural refinement in the nanometric range of Mg2Ni by milling increases its hydriding and dehydriding kinetics compared with crystalline Mg2Ni. The latter behavior can be understood by considering the decrease in the diffusion path and the increase in the inter-grain region because of the appearance of the nanocrystalline structure. In order to eliminate the influence of the microstructural refinement of Mg2Ni during the composite synthesis (milling) and establish the effects of the amorphous phase on the hydriding and dehydriding properties of Mg2Ni, the results obtained from the composite were compared with those obtained from a Mg2Ni sample milled for 10 min.
445
Fig. 4 (a) shows the XRD pattern of composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 produced by milling. The amorphous phase was detected in the XRD pattern due to an increase in the diffraction intensity over the background level in the regions close to 2θ ≈ 20, 40, and 65° [38]. The average crystallite size of Mg2Ni in the composite was 25 nm. On the other hand, Fig. 4 (b) shows the XRD pattern of the Mg2Ni prepared for comparison. In this case, the calculated crystallite size of Mg2Ni was 23 nm. On the basis of the results, it is possible to conclude that Mg2Ni in the composite and the milled Mg2Ni presented similar microstructural refinements; therefore, any difference in the hydriding and dehydriding properties between the milled Mg2Ni and the composite material might be related to the presence of the amorphous Mg50Ni50 phase. Fig. 5 (a) shows the SEM image of the composite material obtained in the SE mode. It can be seen that the particle size distribution of the composite is similar to that of the Mg2Ni alloy (Fig. 3 (a)). EDS analysis revealed that the brighter particles in the backscattered electron view (BSE) image (Fig. 5 (b)) correspond to amorphous Mg50Ni50 and the darker areas to the Mg2Ni matrix (Fig. 5 (c)). On the basis of the results, it is possible to conclude that the milling time employed (10 min) in the composite synthesis was sufficient to produce an intimate contact between Mg2Ni and the amorphous Mg50Ni50 particles. Finally, another important conclusion from the XRD (Fig. 4) and SEM (Fig. 5) results is that no alloy formation (between Mg2Ni and the amorphous phase) occurred during milling.
Fig. 4 – XRD patterns of (a) Mg2Ni (25 wt.%) amorphous Mg50Ni50 composite and (b) milled Mg2Ni.
446
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
Fig. 6 – Hydriding kinetics of (a) composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 and (b) milled Mg2Ni at 363 K under an initial hydrogen pressure of 2 MPa.
3.2.
Hydriding and Dehydriding Characterization
The hydriding and dehydriding characteristics of the prepared composite were studied and compared to that of the milled Mg2Ni to establish the possible utility of the composite for hydrogen storage. The samples were activated after one hydriding (initial hydrogen pressure of 2 MPa at 363 K) and dehydriding (under vacuum of 27 Pa at 453 K) cycle. The dehydriding temperature was selected to avoid the crystallization process of amorphous Mg50Ni50 that starts at approximately 573 K [41]. The low equilibrium H pressure of Mg2Ni below 473 K [42] precludes the ability to obtain the desorption data by the volumetric method. In this first hydriding cycle, the Mg2Ni sample absorbed approximately 1.60 wt.% H, whereas the composite material only absorbed 1.47 wt.% H. Fig. 6 shows the hydrogen absorption curves of the composite and Mg2Ni samples during four hydriding cycles.
Fig. 5 – SEM micrographs showing (a) SE view and (b) BSE of composite material (arrows indicate the position of amorphous Mg50Ni50). (c) Elemental EDS analysis patterns of the selected area in (b).
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
The reversible hydrogen storage capacity and hydriding kinetics of the composite material were greater than those of Mg2Ni. The initial hydrogen absorption rate for both samples showed no obvious dependence on the hydriding cycles. Thus, the average initial rate of hydrogen absorption for the composite material (0.60 ± 0.06 wt.% H min− 1) was approximately four times higher than that for the Mg2Ni sample (0.15 ± 0.02 wt.% H min− 1). On the basis of the results, it is possible to conclude that the amorphous phase improved the hydriding kinetics and hydrogen storage capacity of Mg2Ni upon cycling under the experimental conditions. The latter behavior could be attributed to two factors. First, it has been reported that on the surface of amorphous Mg50Ni50, Mg is combined with oxygen whereas Ni is present in metallic state [43,44]. This Ni, which formed enriched layers, served as an effective catalyst for dissociating hydrogen molecules. Second, Mg–Ni based amorphous alloys showed superior oxidation resistances to their crystalline counterparts [45], which, added to the existence of metallic Ni underneath the top surface, provided “passageways” to the composite for diffusing hydrogen atoms in atmospheres containing oxygen. Fig. 7 shows XRD patterns of the composite and Mg2Ni samples after five hydriding cycles. It is possible to observe that both materials consisted of a mixture of Mg2NiH4, Mg2NiH0.3, residual Ni, and MgO. The only difference between the two XRD patterns was the increase in the diffraction intensity over the
Fig. 7 – XRD patterns of hydrogenated (a) composite Mg2Ni (25 wt.%) amorphous Mg50Ni50 and (b) milled Mg2Ni.
447
background level in the regions close to 2θ ≈ 20, 40, and 65° in the composite pattern (Fig. 7 (a)), which proved the presence of the hydrogenated amorphous phase [46]. Dehydrogenation of the samples was studied by thermogravimetric analysis after the last hydriding cycle. Fig. 8 (a) shows the results. The total hydrogen mass losses for the milled Mg2Ni (1.6 wt.%) and composite (1.4 wt.%) samples were similar to the amount of hydrogen absorbed during the first hydriding cycle, demonstrating that only partial dehydrogenation was achieved in the volumetric experiments under the experimental conditions. A slight rise of the thermogravimetric curves in the high temperature region was due to the oxidation of the alloys by the residual oxygen absorbed on the alloy surfaces. Iwakura et al. previously reported the latter behavior in Mg–Ni based alloys [37]. After five cycles, the hydrogen storage capacities of the milled Mg2Ni and composite samples were, respectively, 47% and 73% of their initial capacities. Considering that both samples were subjected to some desorption conditions and that the capacity loss could be attributed principally to partial
Fig. 8 – (a) Thermogravimetric and (b) derivative thermogravimetric profiles of the hydrogenated samples obtained at 5 K/min under flowing N2.
448
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
dehydrogenation, it is possible to conclude that the amorphous phase improved the desorption kinetics of hydrogenated Mg2Ni. Hydrogen absorption can produce nanocrystallization in the amorphous phase [47,48]. This fact could worsen the hydriding and dehydriding kinetics of the composite. A detailed investigation of how the hydrogen absorption process affects amorphous Mg50Ni50 stability is beyond the scope of this work. However, on the basis of the results of XRD of the hydrogenated composite (Fig. 7(a)), it is possible to establish that during the hydriding and dehydriding cycles, a complete crystallization of the amorphous phase did not occur. Fig. 8 (b) shows the derivative thermogravimetric curves of the milled Mg2Ni and composite materials. Instead of two separate desorption steps (desorption of the hydrogenated amorphous phase and Mg2NiH4), the entire composite released hydrogen in a single principal process. The temperature of the maximum hydrogen desorption rate of the composite material (481 K) was lower than that of the milled Mg2Ni (525 K). Furthermore, the starting temperature for the desorption of the composite (375 K) was similar to those reported by Orimo et al. [20] (373 K) and Guzmán et al. [46] (373 K) for the desorption of hydrogenated amorphous Mg50Ni50. On the basis of the results, it is possible to conclude that during the desorption process of the composite material, a catalytic effect of hydrogenated amorphous Mg50Ni50 on the hydrogen desorption of Mg2NiH4 was observed. Zaluska et al. [49] previously observed a similar behavior in the MgH2– Mg2NiH4 system. Catalytic desorption could be explained owing to the interaction between neighboring particles of Mg2NiH4 and hydrogenated amorphous Mg50Ni50. Thus, when hydrogenated amorphous Mg50Ni50 begins to desorb hydrogen, it undergoes a contraction that applies a strain to the attached Mg2NiH4 and facilitates its desorption. In addition, other dehydriding reactions could be observed in the derivative thermogravimetric curves of Mg2Ni and the composite material around 585 K (Fig. 8 (b)). This reaction was unaffected by the existence of the amorphous phase. According to Selvan et al. [50,51], this desorption process corresponded to the dehydriding of Mg2NiH, which is obtained according to the following equations: Mg2 NiH4 →Mg2 NiH þ 1:5 H2
ðaround 540–570 KÞ
Mg2 NiH→Mg2 NiH0:3 þ 0:35 H2
ðaround 560–590 KÞ:
The apparent activation energy for the principal dehydriding process was estimated with the Kissinger method [40]. According to the Kissinger theory, the slope of the linear plot of ln(β TP− 2) versus TP− 1 is the apparent activation energy of the process, where β is the heating rate and TP is the peak temperature. Fig. 9 shows the Kissinger plot of the dehydriding process of Mg2Ni and the composite materials. The apparent activation energy for the hydrogen desorption of the composite material (48 ± 6 kJ mol− 1) was lower than the calculated value for the dehydriding process of Mg2Ni (85 ± 4 kJ mol− 1). The latter results confirmed the existence of a catalytic effect during the composite desorption. On the other hand, the apparent activation energies for hydrogen desorption of the composite and Mg2Ni samples were lower than that reported by Selvan et al. [50] for the desorption of Mg2NiH4 (183 kJ mol− 1). The
Fig. 9 – Plot of ln(β TP− 2) versus 1000 TP− 1 for dehydriding of milled Mg2Ni and composite Mg2Ni (25 wt.%) amorphous Mg50Ni50.
discrepancy between the values reported by Selvan et al. [50] and those obtained in this study can be understood considering the differences in the experimental procedures. Selvan et al. [50] employed thermogravimetric analysis under a hydrogen pressure of 101.3 kPa, whereas in this study, the thermogravimetric analyses were carried out under flowing N2.
4.
Conclusions
In order to improve the hydriding and dehydriding properties of Mg2Ni, a composite of Mg2Ni with 25 wt.% of amorphous Mg50Ni50 was prepared via mechanical milling. On the basis of the results, it is possible to conclude that amorphous Mg50Ni50 improved the hydrogen storage capacity and hydriding kinetics of Mg2Ni alloy upon cycling. In addition, a catalytic effect of amorphous Mg50Ni50 on the hydrogen desorption of Mg2NiH4 was observed. Thus, it is proposed that when hydrogenated amorphous Mg50Ni50 begins to desorb hydrogen, it undergoes a contraction that applies a strain to the attached Mg2NiH4 and facilitates its desorption. Finally, although the addition of amorphous Mg50Ni50 improved the hydrogen absorption and desorption kinetics of Mg2Ni and its cyclability at low temperatures, the technical targets established by the U.S. Department of Energy (DOE) [52] for transport applications have not been achieved. The main problem of the composites based on Mg2Ni is the high stability of Mg2NiH4, which has an equilibrium H pressure of 91 Pa at 358 K (maximum delivery temperature recommended by DOE), which is very low for mobile applications.
Acknowledgments The authors gratefully acknowledge “Fondo Nacional Desarrollo Científico y Tecnológico de Chile”, FONDECYT, proyect No. 1070085 and “Comisión Nacional de Investigación Científica y Tecnológica de Chile”, CONICYT, for their economic support granted to realize this work. Some authors thank the
M A TE RI A L S CH A RACT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2 –4 5 0
Spanish Ministry of Education and Science, MEC, for its financial support under contract No. MAT2008-06547-C02-01.
REFERENCES
[1] Züttel A. Materials for hydrogen storage. Mater Today 2003: 24–33. [2] Selvan P, Viswanathan B, Swamy CS, Srinivasan V. Magnesium and magnesium alloy hydrides. Int J Hydrogen Energ 1986;11: 169–92. [3] Fernández JF, Sánchez CR. Rate determining step in the absorption and desorption of hydrogen by magnesium. J Alloy Comp 2002;340:189–98. [4] Reilly JJ, Wiswall RH. Reaction of hydrogen with alloys of magnesium and nickel and the formation of Mg2NiH4. Inorg Chem 1968;7:2254–6. [5] Reilly JJ, Wiswall RH. The reaction of hydrogen with alloys of magnesium and copper. Inorg Chem 1967;6:2220–3. [6] Zaluski L, Zaluska A, Ström-Olsen J. Hydrogen absorption in nannocrystalline Mg2Ni formed by mechanical alloying. J Alloy Comp 1995;217:245–9. [7] Sahlberg M, Andersson Y. Hydrogen absorption in Mg–Y–Zn ternary compounds. J Alloy Comp 2007;446–447:134–7. [8] Fernández JF, Bodega J, Sánchez CR. Hydriding/dehydriding properties of magnesium–ZrCr2 composites. J Alloy Comp 2003;356–357:343–7. [9] Hanada N, Ichikawa T, Fujii H. Hydrogen absorption kinetics of the catalyzed MgH2 by niobium oxide. J Alloy Comp 2007;446–447: 67–71. [10] Wang P, Wang AM, Zhang HF, Ding BZ, Hu ZQ. Direct hydrogenation of Mg under the action of FeTi1.2 (amorphous) and mechanical driving force. J Mater Sci Lett 2001;20:753–4. [11] Wang P, Wang AM, Ding BZ, Hu ZQ. Mg–FeTi1.2 (amorphous) composite for hydrogen storage. J Alloy Comp 2002;334:243–8. [12] Zaluska A, Zaluski L, Ström-Olsen J. Structure, catalysis and atomic reactions on the nano-scale: a systematic approach to metal hydrides for hydrogen storage. Appl Phys A 2001;72:157–65. [13] Massalski T, Okamoto H, Subramanian P, Kacprzak L. Binary alloy phase diagram. 2nd ed. Metals Park, OH: American for Metals; 1990. [14] Nohara S, Inoe H, Fukumoto Y, Iwakura C. Compositional and structural characteristics of MgNi alloy prepared by mechanical alloying for use as negative electrodes in nickel-metal hydride batteries. J Alloy Comp 1997;259:183–5. [15] Abdellaoui M, Gracco D, Percheron-Guegan A. Structural characterization and reversible hydrogen absorption properties of Mg2Ni rich nanocomposite materials synthesized by mechanical alloying. J Alloy Comp 1998;268: 233–40. [16] Gavra Z, Mintz MH, Kimmel G, Hadari Z. Allotropic transitions of magnesium nickel hydride (Mg2NiH4). Inorg Chem 1979;18: 3595–7. [17] Noréus D, Werner PE. Structural studies of monoclinic Mg2NiD4. J Less Common Met 1984;97:215–22. [18] Soubeyroux JL, Fruchart D, Mikou A, Pezat M, Darriet B. Etude structurale du systeme Mg2Ni–H2 III La variete monoclinique de Mg2NiH4. Mater Res Bull 1984;19:1119–28. [19] Orimo S, Fujii H. Effects of nanometer-scale structure on hydriding properties of Mg–Ni alloys: a review. Intermetallics 1998;6:185–92. [20] Orimo S, Fujii H. Materials science of Mg–Ni-based new hydrides. Appl Phys A 2001;72:167–86. [21] Orimo S, Ikeda K, Fujii H, Fujikawa Y, Kitano Y, Yamamoto K. Structural and hydriding properties of Mg–Ni–H system with nano and/or amorphous structures. Acta Mater 1997;45: 2271–8.
449
[22] Guzmán D, Ordoñez S, Fernández JF, Sánchez C, Serafini D, Rojas PA, et al. Indications of the formation of an oversaturated solid solution during hydrogenation of Mg–Ni based nanocomposite produced by mechanical alloying. Int J Hydrogen Energ 2009;34:5429–38. [23] Benjamin J. Dispersion strengthened superalloys by mechanical alloying. Metall Trans 1970;1:2943–51. [24] Benjamin J, Volin T. The mechanism of mechanical alloying. Metall Trans 1974;5:1929–34. [25] Ordoñez S, Rojas P, San Martín A, Serafini D, Bustos O. Stoichiometric quantification of the amorphous precursor in the Mg–Ni system. Mater Sci Forum 2003;416–418:138–43. [26] Ordoñez S, Rojas P, Bustos O, Martínez V, Serafíni D, San Martín A. Crystalline Mg2Ni obtained from Mg–Ni amorphous precursor produced by mechanical alloying. J Mater Sci Lett 2003;22:717–20. [27] Rojas P, Ordoñez S, Serafini D, Zúñiga A, Lavernia E. Microstructural evolution during mechanical alloying of Mg and Ni. J Alloy Comp 2005;391:267–76. [28] Orimo S, Fujii H. Hydriding properties of the Mg2Ni–H system synthesized by reactive mechanical grinding. J Alloy Comp 1996;232:L16–9. [29] OrimoS,FujjiH,IkedaK.Notablehydridingpropertiesofa nanostructuredcomposite materialoftheMg2Ni–Hsystem synthesizedbyreactivemechanicalgrinding.ActaMater1997;45: 331–41. [30] Fujii H, Orimo S, Ikeda K. Cooperative hydriding properties in a nanostructured Mg2Ni–H system. J Alloy Comp 1997;253–254: 80–3. [31] Zaluska L, Zaluski A, Ström-Olsen J. Nanocrystalline magnesium for hydrogen storage. J Alloy Comp 1999;288:217–25. [32] Huot J, Liang G, Schulz R. Mechanically alloyed metal hydride systems. Appl Phys A 2001;72:187–95. [33] Mulas G, Delogu F, Cocco G. Effects of mechanical processing on the kinetics of H2 adsorption in Mg2Ni alloys. J Alloy Comp 2009;473:180–4. [34] Ruggeri S, Lenain C, Roué L, Liang G, Huot J, Schulz R. Mechanically driven crystallization of amorphous MgNi alloy during prolonged milling: applications in Ni–MH batteries. J Alloy Comp 2002;339:195–201. [35] Vijay R, Sundaresan R, Maiya MP, Murthy SS. Comparative evaluation of Mg–Ni hydrogen absorbing materials prepared by mechanical alloying. Int J Hydrogen Energ 2005;30: 501–8. [36] Anik M. Electrochemical hydrogen storage capacities of Mg2Ni and MgNi alloys synthesized by mechanical alloying. J Alloy Comp 2010;491:565–70. [37] Iwakura C, Nohara S, Zhang S, Inoue H. Hydriding and dehydriding characteristics of an amorphous Mg2Ni–Ni composite. J Alloy Comp 1999;285:246–9. [38] Guzmán D, Ordoñez S, Serafini D, Rojas P, Bustos O. Effect of the milling energy on the production and thermal stability of amorphous Mg50Ni50. J Alloy Comp 2009;471: 435–41. [39] Cullity B. Elements of X ray diffraction. 2nd ed. USA: Addison-Wesley Publishing Co.; 1978. [40] Kissinger HE. Reaction kinetics in differential thermal analysis. Anal Chem 1957;29:1702–6. [41] Guzmán D, Ordoñez S, Serafini D, Rojas PA, Aguilar C, 2Santander M. Thermal stability of amorphous Mg50Ni50 alloy produced by mechanical alloying. J Non-Cryst Solids 2010;356: 120–3. [42] Fukai Y. The metal–hydrogen system. USA: Springer-Verlag; 1993. [43] Zhang S, Yorimitsu K, Nohara S, Morikawa T, Inoue H, Iwakura C. Surface analysis of an amorphous MgNi alloy prepared by mechanical alloying for use in nickel-metal hydride batteries. J Alloy Comp 1998;270:123–6.
450
MA TE RI A L S CH A R A CT ER IZ A TI O N 62 ( 20 1 1 ) 4 4 2–4 5 0
[44] Zhang S, Hara Y, Morikawa T, Inoue H, Iwakura C. Electrochemical and structural characteristics of amorphous MgNix (x≥1) alloys prepared by mechanical alloying. J Alloy Comp 1999;293–295:552–5. [45] Zhang Y, Dong X, Guo S, Wang G, Ren J, Wang X. The cycle stabilities of the as-cast and quenched La2Mg(Ni0.85Co0.15)9Mx (M=B, Cr, Ti; x=0, 0.1) hydrogen storage alloys. J Alloy Comp 2005;298:178–83. [46] D. Guzmán, Estudio de la fabricación y comportamiento frente a la hidrogenación de aleaciones Mg–Ni, Doctoral Thesis Universidad de Santiago de Chile, Chile, 2008. [47] Li L, Wu DC, Liang GY, Sun ZB, Guo YL. The dehydrogenation process of amorphous (Mg60Ni25)92La8 alloy. J Alloy Comp 2009;474:378–81.
[48] Huang LJ, Liang GY, Sun ZB, Zhou YF. Nanocrystallization and hydriding properties of amorphous mel-spun Mg65Cu25Nd10 alloy. J Alloy Comp 2007;432:171–6. [49] Zaluska A, Zaluski L, Ström-Olsen J. Synergy of hydrogen sorption in ball-milled hydrides of Mg and Mg2Ni. J Alloy Comp 1999;298:197–206. [50] Selvam P, Viswanathan B, Swamy CS, Srinivasan V. Thermal studies on Mg2NiH4: existence of additional hydride phase in the Mg2Ni–Hydrogen system. Thermochim Acta 1988;125:1–8. [51] Selvam P, Viswanathan B, Swamy CS, Srinivasan V. Mg2NiH: a new hydride phase in the Mg2Ni–H2 system. Int J Hydrogen Energ 1988;13:749–59. [52] Website:http://www1.eere.energy.gov/hydrogenandfuelcells/ storage/pdfs/targets_onboard_hydro_storage.pdf.