Indications of the formation of an oversaturated solid solution during hydrogenation of Mg–Ni based nanocomposite produced by mechanical alloying

international journal of hydrogen energy 34 (2009) 5429–5438

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Indications of the formation of an oversaturated solid solution during hydrogenation of Mg–Ni based nanocomposite produced by mechanical alloying D. Guzma´na,*, S. Ordon˜ezb, J.F. Ferna´ndezc, C. Sa´nchezc, D. Serafinid, P.A. Rojase, C. Aguilarf a

Departamento de Ingenierı´a en Metalurgia, Facultad de Ingenierı´a, Universidad de Atacama y Centro Regional de Investigacio´n y Desarrollo Sustentable de Atacama, CRIDESAT, Av. Copayapu 485, Copiapo´, Chile b Departamento de Ingenierı´a Metalu´rgica, Facultad de Ingenierı´a, Universidad de Santiago de Chile, Av. Lib. Bernardo O’Higgins 3363, Santiago, Chile c Departamento de Fı´sica de Materiales, Facultad de Ciencias, Universidad Auto´noma de Madrid, Cantoblanco 28049, Madrid, Spain d Departamento de Fı´sica, 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 Ingenierı´a Meca´nica, Facultad de Ingenierı´a, Av. Los Carrera 01567, Quilpue´, Pontificia Universidad Cato´lica de Valparaı´so, PUCV, Chile f Instituto de Materiales y Procesos Termomeca´nicos, Facultad de Ciencias de la Ingenierı´a, Universidad Austral de Chile, Av. General Lagos 2086, Valdivia, Chile

article info

abstract

Article history:

An oversaturated solid solution of H in a nanocomposite material formed mainly by

Received 24 November 2008

nanocrystalline Mg2Ni, some residual nanocrystalline Ni and an Mg rich amorphous phase

Received in revised form

has been found for the first time. The nanocomposite was produced by mechanical alloying

23 April 2009

starting from Mg and Ni elemental powders, using a SPEX 8000D mill. The hydriding

Accepted 23 April 2009

characterization of the nanocomposite was carried out by solid–gas reaction method in

Available online 4 June 2009

a Sievert’s type apparatus. The maximum hydrogen content reached in a period of 21 Ks without prior activation was 2.00 wt.% H under hydrogen pressure of 2 MPa at 363 K. The

Keywords:

X-ray diffraction analysis showed the presence of an oversaturated solid solution between

Hydrogen absorbing materials

nanocrystalline Mg2Ni and H without any sign of Mg2NiH4 hydride formation. The dehy-

Nanostructured materials

driding behaviour was studied by differential scanning calorimetry and thermogravimetry.

Mechanical alloying

The results showed the existence of two desorption peaks, the first one associated with the

Thermal analysis

transformation of the oversaturated solid solution into Mg2NiH4, and the second one with

X-ray diffraction

the Mg2NiH4 desorption. ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved.

1.

Introduction

Magnesium appears as 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, the MgH2 is very stable (the equilibrium pressure at 552 K is 0.1 MPa) and its hydrogen absorption–desorption kinetics is quite low [1,2] for technological applications. In this way,

* Corresponding author. Tel.: þ56 52 206646. E-mail address: [email protected] (D. Guzma´n). 0360-3199/$ – see front matter ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2009.04.070

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several attempts have been made to overcome these limitations as alloying magnesium with other elements [3–8], production of magnesium base composites [9–12] and modifications of its morphology and microstructure [13]. Mg2Ni can be considered as a very attractive material for applications where a high hydrogen storage capacity is required. First of all, this material has a high hydrogen capacity (3.6 wt.% H) compared to the well established lowtemperature metal hydrides as LaNi5 (1.4 wt.%) or TiFe and TiMn (1.8 wt.%) [14,15] at a lower cost. On the other hand, although other complex hydrides like LiBH4 and NaAlH4 have a higher theoretical hydrogen capacity than Mg2NiH4 [14,16], the use of these complex hydrides for hydrogen storage is challenging because of kinetics and thermodynamics limitations [14]. In the recent years, some research have been devoted to the area of the hydrogen storage in metal-organic frameworks (MOFs) [17]. The maximum hydrogen storage reported in these materials is about 7 wt.% [18]. Nevertheless, the principal drawback of MOFs-compared with Mg2Ni is the low hydrogenation temperature (z77 K) necessary to maintain high adsorption efficiency. Two stable intermetallics, Mg2Ni and MgNi2, have been reported in the Mg–Ni system [19]. The first one absorbs hydrogen in solid solution (SS) up to Mg2NiH0.3 and forms a hydride, Mg2NiH4, at higher hydrogen concentration. The hydride has a lower storage capacity than MgH2 but faster hydriding–dehydriding kinetics [3,20,21]. Below 503 K, it crystallises in a monoclinic low-temperature phase transforming to a cubic high-temperature one above that temperature [22–24]. On the other hand, MgNi2 does not react with hydrogen at pressure up to 2.8 MPa and a temperature of 623 K [3]. Considering the large differences in melting points and vapour pressure between Mg and Ni, the production of high quality Mg2Ni by conventional fusion process is complicated. However, it is possible to avoid the inherent difficulties of the fusion technique by using mechanical alloying process (MA) developed by Benjamin et al. [25,26] in the early 1970s. In the last decade, MA has been widely used to produce hydrogen storage materials, especially Mg-based alloys [27–33], because MA produces a microstructural refinement [34–39] and increase of the specific surface area [40], which improve the hydriding– dehydriding kinetics. In relation to the production of Mg2Ni by MA, Rojas et al. [32] have reported that the phase transformation sequence during the milling of pure Mg and Ni could be: Mgc þ Nic / Mgnc þ Nic / amorphous þ Ninc þ Mgnc / amorphous þ Ninc þ Mg2Ninc / Mg2Ninc, where c and nc stand for crystalline and nanocrystalline, respectively. Homogeneous Mg2Ninc can be also obtained by a combination of MA and heat treatment [32,41,42]. A reduced number of investigations [21,41,42] have analyzed the hydrogen sorption properties by solid–gas reaction of nanocrystalline Mg2Ni produced by MA of pure Mg and Ni, and none has studied the hydrogen behaviour of nanocrystalline/amorphous phases obtained at intermediate milling time. On the other hand, a lot of work has been done by Orimo and co-workers on the reactive milling of crystalline Mg2Ni under hydrogen atmosphere [34–38]. The results obtained in these investigations showed that the hydrogen content in the nanostructured Mg2Ni–H system synthesized by reactive milling reaches 1.6 wt.% (Mg2NiH1.8) without

changing the crystal structure of the Mg2Ni matrix. They proposed that the large amount of dissolved hydrogen in the nanostructured Mg2Ni–H system is due to the formation of the standard SS Mg2NiH0.3 and the high solubility of hydrogen in the Mg2Ni grain boundary regions, which can reach up to 4.0 wt.% H. Considering previous investigations which showed that Ni particles [43–46] and amorphous phase [47,48] have a catalytic effect on the hydrogen storage properties of several Mg-based alloys, in this work the hydriding and dehydriding behaviours of a nanocomposite formed mainly by Mg2Ninc some residual Ninc and amorphous phase, produced in the intermediate milling stage of elemental Mg and Ni in a 2:1 atomic proportion have been studied.

2.

Experimental procedure

Nanocomposite material formed mainly by Mg2Ni, some residual Ni and amorphous phase were produced by MA. Mg turnings (98 wt.% pure, supplied by Aldrich) and Ni powders (99,99% wt. pure, supplied by Fluka) in a 2:1 atomic proportion were milled in a SPEX 8000D during 10 h, with a ball to powder weight ratio of 20:1 under controlled atmosphere of purified Ar (99.999 wt.% pure). The milling process was performed in discontinuous way consisting of 1 milling hour followed by rest period of 0.5 h. X-ray diffraction (XRD) measurements were carried out in a Shimadzu XRD-600 diffractometer with Cu Ka radiation. The thermal stability of 10 h milled powders was investigated by differential scanning calorimetry (DSC) using a heating rate of 20 K/min in flowing N2 in a DSC 2920 TA Instruments device. The internal microstructure of the powder was studied using a Philips CM-12 transmission electron microscopy (TEM) operated at 100 kV, whereas the powder size and morphology were studied using a Philips XL 30 S-FEG scanning electron microscopy (SEM) equipped with an EDS system. The hydrogen absorption was investigated in a Sievert type apparatus at 363 K and hydrogen pressure of 2 MPa, while the hydrogen desorption behaviour was investigated by thermogravimetry (TG) and differential scanning calorimetry (DSC) under flowing N2 at 20 K/min using a SDT 2960 and DSC 2920 TA instrument devices.

3.

Results

3.1. Structure, composition and morphology of Mg2Ni nanocrystalline material Fig. 1(a) shows the X-ray diffraction pattern of the initial mixture of Mg turnings and Ni powders. The initial average crystallite size was calculated using the Scherrer method [49] considering the (101) Mg plane and (200) Ni plane. In order to obtain the full width at the half maximum peak intensity, the experimental peaks were fitted using a Lorentz function. The results obtained were corrected considering the instrumental width. The calculated initial average crystallite sizes for Mg and Ni were 105 and 80 nm, respectively. Fig. 1(b) shows the X-ray diffraction pattern of powder milled for 10 h in a SPEX mill. It can be seen that the material

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Fig. 2 – DSC curve corresponding to sample milled for 10 h.

Fig. 1 – XRD patterns of (a) initial mixture of Mg and Ni in a 2:1 atomic proportion and (b) powder produced after 10 h of milling. The dotted curve is an estimation of the amorphous contribution to the XRD pattern.

consists of a mixture of Mg2Ni and residual Ni, with an average crystallite sizes of 11 and 15 nm, respectively. The Mg2Ni average crystallite size was calculated considering the (112) plane. According to previous work [30–32], the additional Mg needed to maintain the milling stoichiometry is forming an amorphous phase, which can be subsequently transformed by mechanical and/or thermal way into Mg2Ni. This amorphous phase can be detected in the XRD pattern showed in Fig. 1(b), by a rising in the diffraction intensity over the background level, in the regions near to 2q z 20, 40 and 65 [30–32]. The dotted curve in Fig. 1(b) represents an estimation for the contribution of the amorphous phase to the XRD pattern. The production of this amorphous phase rich in Mg during the intermediate stages of MA, can be understood considering the faster microstructural refinement to nanometric range of this element in comparison with Ni. At this stage, the disordered grain boundaries are enriched in Mg atoms compared to the grains. Therefore, these zones increase their reactivity and can react with elemental Ni to produce an amorphous phase [32].

The DSC curve of the mechanically alloyed powders milled for 10 h is shown in Fig. 2. It can be seen that the DSC curve exhibits two exothermic peaks during the heating process. Based on previous works [30–32] it can be concluded that the first exothermic reaction at 423 K, is related with Mg2Ni crystallization, and the second exothermic event, which takes place between 473 and 523 K, is related to matrix relaxation and solid state reaction between Mg and Ni to form Mg2Ni. To obtain more information about the stability of the amorphous phase, two samples were annealed at 373 and 573 K for 24 h in flowing Ar. Fig. 3 shows the XRD patterns of the annealed samples. The amount of amorphous phase in each sample can be estimated from the ratio of the integrated intensity of the diffraction bump to that of the whole diffraction pattern [50]. The lattice parameters were calculated using the method developed by Cohen [49]. The microstructural parameters for the studied samples are summarized in Table 1. Based on the obtained results, it can be concluded that the sample annealed at 373 K is triphasic (nanocrystalline Mg2Ni þ nanocrystalline Ni þ amorphous) whereas the sample annealed at 573 K is biphasic (nanocrystalline Mg2Ni þ amorphous). After the high-temperature annealing (573 K) the Mg2Ni crystallite size increases slightly, the amount of amorphous phase is reduced and nanocrystalline Ni is not detected. This behaviour can be understood considering that below 573 K (Fig. 2) the amorphous transformation into Mg2Ni (423 K) and the solid reaction between Mg and residual Ni to form Mg2Ni (473–523 K) takes place [30–32]. As concerns the Mg2Ni lattice parameters, it can be seen that the ao and co parameters for the milled and annealed at 373 K samples are higher than those reported by Soubeyroux et al. [51] (ao ¼ 0.5205 nm, co ¼ 1.3235 nm, unit cell volume ¼ 0.3105 nm3) and Hirata et al. [52] (ao ¼ 0.521 nm co ¼ 1.323 nm, unit cell volume ¼ 0.3110 nm3). This effect could be attributed to atomic-site interchange between the relatively smaller Ni atoms and the relatively larger Mg atoms during the milling process [53]. After annealing treatment at 573 K, the Mg2Ni lattice parameters are reduced to the literature reported values.

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particles (Fig. 4(d)). Elemental EDS analyses reveal that the Ni content in the dark homogeneous regions was 46 wt.% while it increases to 50 wt.% when the analysis was carried out on the bright particles. Finally, the milled powders were observed by TEM. Fig. 5 shows the SAD diffraction pattern and the dark field image. It can be seen that the grain size of the intermetallic compound Mg2Ni is less than 20 nm, which confirms the result obtained by XRD. On the other hand, the SAD diffraction pattern shows very diffuse rings with two main contributions: nanocrystalline and amorphous phase.

3.2.

Fig. 3 – XRD patterns of (a) sample annealed at 373 K and (b) sample annealed 573 K for 24 h.

Fig. 4 shows SEM images of the powder after 10 h of milling. It can be seen that the powder has a bimodal size distribution composed of large irregular agglomerates with an apparent diameter around 150 mm (Fig. 4 (a)), and small particles with flake-like shape of 5 mm (Fig. 4 (b)). The large agglomerates are made of several smaller particles welded together. In order to observe the internal microstructure, backscattered electron (BSE) images were taken from the cross-section of the milled powder. It can be seen that the chemical composition, revealed by the difference in grey scale in BSE image, is practically homogeneous throughout the powder (Fig. 4 (c)), although it is possible to observe a small number of brighter

Hydrogen absorption in nanocrystalline Mg2Ni

Hydriding kinetics data for nanocrystalline Mg2Ni obtained at 363 K under hydrogen pressure of 2 MPa can be seen in Fig. 6. The as-produced powder readily absorbs hydrogen without any activation process, achieving hydrogen content of about 2.0 wt.% H in 21 Ks; this fact is an improvement with respect to crystalline Mg2Ni which only absorbs hydrogen at 523 K under hydrogen pressure of 1.5 MPa after previous activation [3]. The influence of Ni and amorphous phase on the improvement of the hydriding process will be discussed later. Fig. 7(a) shows XRD pattern of the hydrogenated nanocomposite sample and main position diffraction angles for Mg2NiH0.3 standard SS [54] and Ni [55]. No traces of hydride phase (Mg2NiH4) were found. It is possible to observe that the diffraction peaks of the hydrogenated Mg2Ni shift to lower angles compared with those of standard Mg2NiH0.3 while Ni peak positions remained unchanged. The last behaviour can be observed more clearly in Fig. 7(b). The shift to lower angles of the Mg2Ni XRD peaks in the hydrogenated sample is related with an increase in its volume unit cell and is an evidence of a new Mg2Ni–H oversaturated SS formation, which has not been reported previously. Table 2 shows the lattice parameters of the Mg2Ni–H oversaturated SS calculated from Cohen method [49]. It can be seen that the co reticular parameter suffers a greater increase than ao reticular parameter, which practically remained unchanged. The last behaviour is a evidence that the hydrogen atom inside Mg2Ni unit cell produces a preferential distortion along the [001] direction. The calculated unit cell volume for Mg2Ni–H oversaturated SS was 0.3303 nm3, which is 6.38% greater than the unit cell volume for Mg2Ni and 3.84% greater than the unit cell volume for standard SS Mg2NiH0.3. The hydrogen content of this oversaturated SS could not be accurately established due to the multiphase nature of this material. However, considering that Ni does not form

Table 1 – Microstructural parameters for MA Mg2Ni nanocrystalline studied samples. Sample

Milled for 10 h

Microstructure

Amorphous (23%) þ nanocrystalline (77%)

Annealed at 373 K Amorphous (21%) þ nanocrystalline (79%) Annealed at 573 K Amorphous (9%) þ nanocrystalline (91%)

Nanocrystalline Lattice parameters, nm Unit cell Crystallite size, nm component vol., nm3 ao co Mg2Ni Ni Mg2Ni Ni Mg2Ni

0.5236 0.3527 0.5229 0.3526 0.5216

1.3286 – 1.3280 – 1.3212

0.3154 0.0439 0.3145 0.0438 0.3113

11 15 13 23 20

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Fig. 4 – SEM micrographs of the powder after 10 h of milling showing (a) general and (b) magnified view. BSE cross-section micrographs showing (c) general and (d) magnified view.

a hydride phase under the experimental condition employed [56] and that the maximum hydrogen capacity reported for the amorphous alloys in the Mg–Ni system [37,38,57,58] is about 2.00 wt.% H, it is thought that the composition of the oversaturated SS between Mg2Ni and hydrogen should be close to 2.00 wt.% H, (Mg2NiH2.2).

It is interesting to analyse the improvement of the hydriding properties of the nanocrystalline Mg2Ni at the light of the Ni and/or amorphous phase content and, also, taken into account their larger unit cell volume. In order to obtain pure nanocrystalline Mg2Ni with a unit cell volume similar to the values reported in previous works [51,52] and without any sign of Ni or amorphous phases, a nanocomposite sample was heated in vacuum at 673 K and immediately cooled down and hydrogenated at 363 K and 2 MPa for 21 Ks. The results

Fig. 5 – Dark field TEM micrograph and SAD pattern showing the Mg2Ni nanostructure formed after 10 h of milling.

Fig. 6 – Hydrogen absorption curve of nanocrystalline Mg2Ni at 363 K and a hydrogen pressure of 2 MPa (without prior activation).

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Fig. 8 – XRD pattern of hydrogenated sample after the heat treatment at 673 K.

the Mg2NiH0.3 standard SS was detected. The lattice parameter of the SS is given in Table 2. Based on the results obtained, it is possible to conclude that the presence of Ni and/or amorphous in the nanocomposite added to the major unit volume cell of Mg2Ni are necessary conditions for the observed improvement of the Mg2Ni hydriding properties.

3.3.

Fig. 7 – (a) XRD pattern of the hydrogenated sample and main diffraction lines for Mg2NiH0.3 and Ni (b) magnified view of XRD pattern shown in (a).

obtained from solid–gas reaction analysis showed that the nanocrystalline Mg2Ni practically does not absorb hydrogen at the imposed condition. The XRD pattern from the hydrogenated sample can be observed in Fig. 8. It can be seen that only

Table 2 – Lattice parameters and unit cell volume comparison. Lattice parameters. nm

Mg2Ni [51] Mg2Ni milled 10 h (this work) Mg2NiH0.3a [54] Mg2NiH0.3 (this work)b Mg2Ni–H (this work)c

Unit cell vol., nm3

Relative change unit cell vol.%

ao

co

0.5205 0.5236

1.3235 1.3286

0.3105 0.3154

0.00 1.57

0.5231 0.5233 0.5264

1.3404 1.3429 1.3763

0.3176 0.3185 0.3303

2.29 2.58 6.38

a Standard SS. b After heat treatment at 673 K. c Oversaturated SS.

Decomposition of the oversaturated SS

Fig. 9(a) shows the result obtained in the thermogravimetry experiment. It confirms that the nanocomposite can absorb about 2 wt.% H. Based on Fig. 9(b), which shows the derivative of the thermogravimetric curve and DSC desorption trace, it is established that dehydriding reaction begins at 450 K. This dehydriding temperature is much lower than that of Mg2NiH4 (550–573 K) [59,60] and similar to the one reported by Orimo et al. [34,35] for dehydriding of Mg2NiH0.3 standard SS (440 K). Moreover, it is possible to observe two stages in the dehydriding process; the first one, with a peak temperature of 493 K releases 0.27 wt.% H and it is associated with an exothermic reaction, while the second one, located at 566 K, releases 1.73 wt.% H and it is related with an endothermic reaction. The hydrogen amount desorbed in each stage was calculated considering the area under the derivative thermogravimetric curve. Generally, the dehydriding reactions in the Mg–Ni system are related with endothermic heat flow signals, therefore, the nature of the first exothermic event must be determined. To do so, a hydrogenated sample was heated at 506 K under N2 flowing and cooled quickly, using a DSC 2920 device, which is equipped with a cooler system. The XRD pattern of the sample after the heat treatment is shown in Fig. 10(a). It can be seen that the presence of diffractions peaks corresponding to the Mg2NiH4 hydride and to the standard SS Mg2NiH0.3, which was not present in the hydrogenated sample. Based on the results obtained, it is possible to conclude that the first exothermic reaction (493 K) is mainly related with crystallization of the Mg2Ni–H oversaturated SS into Mg2NiH4 and the standard SS Mg2NiH0.3. It must be considered that simultaneously or

international journal of hydrogen energy 34 (2009) 5429–5438

Fig. 9 – (a) Thermogravimetric profile and (b) DSC curve and derivative thermogravimetry trace of the hydrogenated sample obtained at 20 K/min under N2 atmosphere in a SDT 2960 TA device.

interacting with the Mg2Ni–H oversaturated SS transformation, the solid reaction between amorphous phase and residual Ni to form Mg2Ni [32] could take place. Fig. 10(b) shows the XRD pattern after the complete hydrogen desorption. As can be seen, the sample is formed mainly by Mg2Ni and a small amount of MgNi2 formed due to the lost of Mg by the MgO formation. Considering the last result, it is possible to conclude that the endothermic reaction shown in Fig. 9(b) is mainly related with the Mg2NiH4 desorption. Due to the high heating rate employed in this work, it is not possible to distinguish the Mg2NiH4 transition into Mg2NiH0.3 SS and its subsequent hydrogen desorption [61,62].

4.

Discussion

The experimental results in the previous sections show that the nanocrystalline Mg2Ni prepared in this work, with an expanded

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lattice volume (0.3154 nm3) compared to the standard Mg2Ni (0.3105 nm3), absorbs 2.00 wt.% H under a H-pressure of 2 MPa at 363 K. Dehydriding data obtained from thermogravimetry is in agreement with the hydrogen absorption results. XRD pattern indicates the formation of a SS of H in Mg2Ni but without any sign of Mg2NiH4 hydride formation. The lattice volume of this SS is larger than the standard one reported for this compound. It is important to mention that the nanocrystalline material remains crystalline enough for a precise determination of the lattice parameters of the compound (see Fig. 1(b) and Fig. 7(a)–(b)). All those evidences point out to the formation of an oversaturated SS of H in nanocrystalline Mg2Ni at the conditions of our hydrogenation experiments. Previous investigations by Orimo and co-workers [34,35] in similar compounds have also reported H-absorption up to 1.6 wt.% H (Mg2NiH1.8) in nanocrystalline Mg2Ni. In these investigations Mg2Ni was mechanically ground under a hydrogen pressure of 1 MPa at room temperature. The XRD analysis only revealed the presence of a standard SS phase, Mg2NiH0.3, while the thermogravimetry analysis indicated that nanocrystalline sample could release about 1.6 wt.% H (Mg2NiH1.8). According to their results, they have proposed that the large amount of dissolved hydrogen in the nanostructured Mg2Ni–H system is due to the formation of the standard SS Mg2NiH0.3 and the high solubility of hydrogen in the Mg2Ni grain boundary regions, which can reach up to 4.0 wt.% H. The experimental processes employed in this work and those used by Orimo et al. [34,35] show clear differences. Orimo et al. employed reactive milling under hydrogen atmosphere of Mg2Ni obtained by conventional fusion method which presented residual Mg. In this work, a nanocomposite formed mainly by Mg2Ni, some residual Ni and amorphous phase, was produced by MA and hydrogenated by solid–gas reaction using a Sievert type apparatus. In our opinion the explanation given by Orimo and co-workers of a hydrogen rich (4 wt.% H) inter-grain region cannot satisfactorily explain our data. That inter-grain region is formed under the non-equilibrium conditions imposed during the reactive milling. The high amount of energy introduced in the system most probably leads to the amorphisation of the previously formed Mg2NiH4 hydride. A previous investigation has shown that amorphisation of Mg2NiH4 takes place more easily than that of the solid solution Mg2NiH0.3 [63]. In comparison, in our experiments hydrogenation takes place at relatively mild conditions of pressure and temperature which precludes the formation of any amorphous phase, apart from that already formed during the synthesis of the nanocomposites. In addition, the low amount of amorphous phase present in our sample cannot account for the total hydrogen content of the sample (2.0 wt.% H) if only 0.27 wt.% H is in the nanocrystalline grains. As mentioned previously, there are several reports [37,38,57,58] indicating that the maximum hydrogen capacity for the amorphous alloys in the Mg–Ni system is about 2.00 wt.% H, but not 4 wt.% H. Finally, the larger unit cell volume of the oversaturated SS compared to the standard one seems to point out to a larger amount of H in our samples compared to that in the standard Mg2NiH0.3 (0.27 wt.% H). We can conclude that an oversaturated SS between Mg2Ni and hydrogen is formed in the nanocrystalline

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5.

Conclusion

A nanocomposite formed mainly by nanocrystalline Mg2Ni, some residual nanocrystalline Ni and an Mg rich amorphous phase, was produced by mechanical alloying starting from elemental Mg and Ni. It was determined that the dissolved hydrogen content in the compound under the experimental conditions employed reaches about 2.0 wt.% H. The XRD pattern of hydrogenated sample revealed the presence of a new oversaturated SS between Mg2Ni and hydrogen, which had not been reported previously, with a unit cell volume of 0.3303 nm3, larger than the standard SS one, 0.3176 nm3. In relation to the nanocomposite dehydriding process, the results showed the existence of two desorption peaks, the first one associated mainly with the transformation of the SS Mg2NiH2.2 into Mg2NiH4, and the second one with the Mg2NiH4 desorption. Finally, based on the results obtained, it is possible to conclude that the presence of Ni and/or amorphous in the nanocomposite added to the larger unit cell volume of Mg2Ni are necessary conditions for the observed improvement of the Mg2Ni nanocomposite hydriding properties.

Acknowledgements

Fig. 10 – XRD patterns of (a) hydrogenated sample after the heat treatment at 506 K and (b) sample after the complete hydrogen desorption.

The authors gratefully acknowledge ‘‘Fondo Nacional Desarrollo Cientı´fico y Tecnolo´gico de Chile’’, FONDECYT proyect No 1070085 and ‘‘Comisio´n Nacional de Investigacio´n Cientı´fica y Tecnolo´gica de Chile’’, CONICYT, for the economical support granted to the realization of this work. Two of us (J.F. Ferna´ndez and C. Sa´nchez) thank the Spanish Minister of Education and Science, MEC, for financial support under contract No MAT2005-06738-C02-01 and to Mr. F. Moreno for technical assistance.

references

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