Preparation and structural, thermal and hydriding characteristics of melt-spun Mg–Ni alloys

Journal of Alloys and Compounds 292 (1999) 107–117

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Preparation and structural, thermal and hydriding characteristics of melt-spun Mg–Ni alloys a ,1 a, a b G. Friedlmeier , M. Arakawa *, T. Hirai , E Akiba a

NTT Integrated Information & Energy Systems Laboratories, 3 -9 -11, Midori-cho, Musashino-shi, Tokyo 180 -8585, Japan b National Institute of Materials and Chemical Research, 1 -1 Higashi, Tsukuba, Ibaraki 305 -8565, Japan Received 1 June 1999; received in revised form 5 June 1999; accepted 19 June 1999

Abstract Glassy, nanostructured and crystalline hydride-forming Mg 12x Ni x alloys with five different compositions (0#x#0.33) were prepared by rapid quenching of the melt with the melt-spinning method. Glassy Mg–Ni alloys (Mg 0.87 Ni 0.13 and Mg 0.84 Ni 0.16 ) crystallize at around 438 K. During crystallization, an intermediate metastable Mg–Ni phase is formed. Hydrogen activation of these alloys requires temperatures and pressures above 663 K and about 1.5 MPa, respectively. Thereafter a very high storage capacity of nearly 6 mass% H was achieved with Mg 0.87 Ni 0.13 .  1999 Elsevier Science S.A. All rights reserved. Keywords: Hydrogen; Metal hydrides; Magnesium–nickel; Metallic glasses; Melt spinning

1. Introduction Nanostructured and amorphous hydride-forming alloys have gained a great deal of attention in recent years [1–7]. They often show improved electrochemical characteristics over their crystalline counterparts, such as good kinetics and corrosion resistivity, features which make them especially interesting for battery applications [4,6,8–10]. A number of works have been published concerning hydrogen in nanostructured and amorphous Mg–Ni alloys prepared by ball milling (mechanical alloying) [11–19]. Even though mechanical alloying has several merits for amorphous alloy preparation, it also has intrinsic problems associated with it, such as contamination with materials from bowls and balls and a difficult handling of the alloys which are usually highly reactive due to their large specific surface area. In addition, mechanical alloying is not always

*Corresponding author. Tel.: 181-422-59-3883; fax: 181-422-594350. E-mail address: [email protected] (M. Arakawa) 1 Present address: DaimlerChrysler AG, Fuel Cell Project House (EP/ BZ), D-73230 Kirchheim / Teck-Nabern, Germany.

suitable for mass production. We used melt-spinning technique [20–22] as an alternative with the purpose of determining its suitability for preparing amorphous (or, more properly, ‘glassy’) Mg–Ni alloys with improved hydriding characteristics and, in a future step, to study their electrochemical behavior. The very high cooling rates of the melt on the order of 10 5 K s 21 or higher which can be achieved with the melt-spinning method often lead to the formation of metastable metallic glasses by ‘freezing’ the structure of the liquid [20]. The thermodynamic conditions required for the formation of metallic glasses have been studied in detail in the literature [23,24]. Under equilibrium conditions, the Mg–Ni alloys studied in this work form two crystalline phases: Mg and Mg 2 Ni, both hexagonal [25]. Sommer et al. [26] and Masumoto et al. [27] reported the formation of amorphous Mg–Ni alloys in the range of 8 to 25 at.% Ni but, to our knowledge, the hydriding characteristics of melt-spun binary Mg–Ni alloys have not yet been studied. In the following we shall describe the alloy preparation method, the morphology, the chemical composition, the microstructure, the thermal behavior, and the activation and hydrogen absorption characteristics of melt-spun Mg– Ni alloy ribbons. The focus will be on those alloys which formed glassy structures.

0925-8388 / 99 / $ – see front matter  1999 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 99 )00285-6

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2. Experimental details

2.1. Sample preparation Mg–Ni alloy ribbons with five different compositions (Mg, Mg 0.87 Ni 0.13 , Mg 0.84 Ni 0.16 , Mg 0.74 Ni 0.26 and Mg 2 Ni) were prepared by rapid quenching of the melt with the melt-spinning technique. Starting materials were a 20mesh Mg powder (assay 99.9% Mg, purchased from Furuuchi Chemical Co.) and a Mg 2 Ni alloy powder (48.2 mass% Mg, 50.0 mass% Ni, with S, Pb, P, Fe and Mn as main impurities, purchased from Japan Metals and Chemicals Co.). The desired proportions of Mg and Mg 2 Ni were carefully weighed, mixed and pressed into pellets. A graphite crucible containing irregularly shaped pieces of these pellets (approx. 10 g per batch) was inductively heated above the respective melting points and then maintained at selected temperatures for about 15 min in order to assure melting. Thereafter the melt was quenched by applying a gas overpressure into the crucible and pressing it through a nozzle onto a rapidly rotating, cooled Cu wheel (f 20.0 cm), in an inert gas atmosphere at room temperature. The main parameters affecting the cooling rate and ribbon morphology are the wheel speed, the overheating of the melt, the atmosphere, the gas overpressure, the geometry of the nozzle and the nozzle–wheel distance. The crucibles had a f 0.8-mm nozzle bore. After several preliminary experiments, the following parameters were selected as the best compromise between a high cooling rate, the ribbon uniformity and a low Mg evaporation rate: A wheel rotational speed of 3000 rpm (equivalent to a wheel edge velocity of 31.4 m s 21 ); heating temperatures of 973 K (Mg), 1123 K (Mg 0.87 Ni 0.13 ), 1153 K (Mg 0.84 Ni 0.16 and Mg 0.74 Ni 0.26 ) and 1173 K(Mg 2 Ni); a 30-kPa He atmosphere; an ejecting He gas overpressure of 50 kPa; and a nozzle–wheel distance of 1 mm. We used He instead of Ar as an inert gas atmosphere because He leads to higher cooling rates (i.e. higher degrees of amorphization). Once cooled down to room temperature, the obtained ribbons were collected and handled in air.

2.2. Characterization methods A morphological characterization of the ribbons was carried out by optical microscopy. After determining the elements present in each sample with X-ray fluorescent analysis (XFA), a quantitative elementary analysis of the melt-spun ribbons was performed by inductively coupled plasma atomic emission spectrometry (ICP). In order to determine the degree of contamination from the crucible, the concentration of C in one sample (Mg) was measured by combustion-gas infrared absorption analysis (the C content can not be measured by ICP). The degree of oxidation was not determined, but the presence of O in Mg 0.87 Ni 0.13 ribbons could be clearly observed by wave-

length-dispersive X-ray spectrometry (WDX) using a Microspec WDX400 apparatus attached to a Hitachi S3200 scanning electron microscope (SEM). These apparatus were also used to measure the Ni distribution in a Mg 0.87 Ni 0.13 ribbon by (Ni Ka) X-ray pulse mapping, an associated technique. Structural analyses were carried out with a Rigaku Geigerflex X-ray powder diffractometer (XRD) using the Bragg–Brentano geometry and Cu Ka radiation. XRD samples were prepared by attaching ribbon pieces with silicone grease onto glass sample holders. The structural characteristics of one alloy (Mg 0.87 Ni 0.13 ) were additionally investigated in the as-spun and annealed (at 463 K) conditions by transmission electron microscopy (TEM) and its associated technique electron diffraction spectrometry. A Hitachi H-9000 high-resolution TEM apparatus with an acceleration voltage of 300 kV (equivalent to l51.97 pm) was used for this purpose; its camera length is L51 m. The TEM samples were prepared directly out of ribbons by low-energy ion milling. In some cases, the surface layer was removed after ion milling by chemical etching using a 0.6-vol.% HNO 3 solution in ethanol. The thermal characteristics of the melt-spun ribbons were studied by differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA). These experiments were carried out in N 2 and Ar gas atmospheres at ambient pressure, respectively. Shimadzu DSC-50 and TGA-51 apparatus with Al and Pt sample pans and about 3- and 10-mg samples, respectively, were used. The samples were heated to selected temperatures (#723 K) at a constant rate of 5 K min 21 and, after an isothermal period of 5 min, cooled down to about 343 K at the same rate. Reference TGA experiments were carried out in order to detect possible sample mass changes caused by undesired chemical reactions of the sample with its environment, especially by oxidation. For Mg 0.87 Ni 0.13 a series of DSC experiments with different maximum temperatures and followed by XRD analyses was carried out in order to study the temperature-dependence of the structural changes induced by annealing. Also in this case an N 2 atmosphere as well as heating and cooling rates of 5 K min 21 were selected. The activation and hydrogen absorption characteristics of Mg 0.87 Ni 0.13 , Mg 0.84 Ni 0.16 and Mg 2 Ni were studied by pressure differential scanning calorimetry (PDSC) using a TA Instruments 2910 DSC apparatus, and (for Mg 0.87 Ni 0.13 with a fully automated Sieverts-type static [28] pressure–composition–temperature (PCT) apparatus (manufactured by Lesca Co., Ltd.). PDSC experiments were carried out either in an isochoric or in an isobaric mode at H 2 pressures varying between 0.3 and 3 MPa. The samples were cyclically heated to preselected temperatures and cooled down to about 353 K, both at a nearly constant rate of 10 (sometimes 5) K min 21 . PCT experiments were run at different (constant) temperatures. Sample activation was carried out in situ by cycling at the usual experimental conditions in both cases, without any surface pretreatment.

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Pure H 2 gas (assay 99.999% H 2 ) was used. PDSC and PCT samples were finally analyzed by XRD in the hydrided condition.

3. Results and discussion

3.1. Morphological characterization All ribbons have a metallic appearance, are stable in air and only etched slowly by conventional agents (like a 0.6 vol.% solution of HNO 3 in ethanol). Except for Mg 2 Ni and Mg 0.74 Ni 0.26 , which are brittle, extremely ductile ribbons with up to several meters of length were obtained. A typical ribbon cross-section measures about 2–4 mm3 25 mm for all five alloys. Fig. 1(a) and (b) show optical micrographs of both sides of a Mg 0.87 Ni 0.13 ribbon: the side which was in contact with the Cu wheel while spinning (from now on: bottom side), and the opposite (top) side, respectively. The bottom-side surface is microscopically rough and relatively opaque, with many longitudinal stripes (replica of the Cu wheel surface, which is periodically polished with abrasive paper for cleaning purposes) and small scratches (caused by He gas bubble trapping [29]). The topside surface (Fig. 1(b)) is smoother and appears very shiny under the optical microscope and could be clearly observed only with polarized light. Ribbons of the other alloys show similar characteristics.

3.2. Chemical composition and homogeneity of as-spun ribbons Table 1 shows an elementary analysis for ribbons of all five melt-spun Mg–Ni alloys investigated in the present work, together with the corresponding measured and starting (nominal) Mg / Ni weight ratios. All alloys containing Ni show lower Mg / Ni weight ratios than the starting nominal values. This expected result is due to Mg losses related to the high vapor pressure of this metal. The degree of contamination is relatively low (#1.56 mass%). Very small amounts of Cu from the spinning wheel were found. Also the contamination from the graphite crucible is negligible: only 0.02 mass% of C was measured for a Mg ribbon. Table 2 shows the chemical composition of five randomly selected Mg–Ni ribbons (A–E), all from the same melt-spinning batch. For ribbon A, ICP analyses were carried out on three neighboring samples (A1–A3), each result representing an average over a ribbon length of about 1 cm. These results show that the ribbons are highly homogeneous: both the Mg and Ni contents fluctuate in a range smaller than 1% of the total mass. Fig. 2 shows the Ni distribution in a Mg 0.87 Ni 0.13 ribbon as determined by Ni Ka X-ray pulse mapping. A very homogeneous distribution of Ni was confirmed.

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3.3. Structural characterization of as-spun samples Fig. 3 shows XRD patterns for as-spun ribbons of all five alloys. From these patterns we can conclude that Mg 0.87 Ni 0.13 and Mg 0.84 Ni 0.16 are glassy (no sharp peaks) while the other three alloys are crystalline. These results agree with published data indicating that Mg–Ni metallic glasses can be prepared by rapid quenching of the melt in the range of 8 to 25 at.% Ni [25,26,29]. The broad peaks in the patterns of Mg 0.87 Ni 0.13 and Mg 0.84 Ni 0.16 at 2u ¯38 8(d¯0.236 nm) are typical for metallic glasses and indicate the existence of a certain short-range atomic ordering. The peak broadening that arises for Mg 0.74 Ni 0.26 (Fig. 3) indicates a nanocrystalline structure (compare with Mg 2 Ni); the average crystallite size was determined to be 20 nm with Rietveld analysis using RIETAN97 [30]. Comparison of the XRD patterns for (crystalline) meltspun Mg and Mg 2 Ni ribbons with ones of the conventionally prepared powders (not presented here) shows similar results for Mg 2 Ni but a noticeable difference in the relative peak intensities for Mg, indicating a preferred crystallite orientation in the [001] lattice direction in the ribbons. Fig. 4 shows a selected-area electron diffraction pattern for as-spun Mg 0.87 Ni 0.13 . The fact that no sharp diffraction spots or rings arise confirms the glassy condition of the sample. The diffuse diffraction ring is characteristic for metallic glasses and corresponds to the broad peak observed at 2u ¯388 in the XRD pattern (Fig. 3). In agreement with those XRD results, a d-spacing of about 0.236 nm was determined for this ring. Fig. 5 shows a TEM image for the same sample using the direct electron beam and the entire diffracted beam corresponding to the ring; the selected area that contributed to the diffraction pattern of Fig. 4 (about 0.05 mm 2 ) is located in the center. The appearance of this image is characteristic for glassy structures [31] and visualizes the lack of well-defined crystallites (compare with Fig. 9). Since our primary interest here is devoted to glassy alloys, the following results will focus mainly on Mg 0.87 Ni 0.13 and Mg 0.84 Ni 0.16 .

3.4. Thermal behavior and structural characterization of annealed samples Fig. 6 shows DSC and TGA experiments with Mg 0.87 Ni 0.13 which were both run up to 723 K (heating and cooling rates were 5 K min 21 ). A well-defined exothermic peak equivalent to DH¯83 J g 21 can be observed at 438 K. The TGA curve indicates that within the experimental precision the sample mass remains constant, eliminating the possibility of an undesired chemical reaction of the sample with its environment. The observed DSC peak indicates therefore the crystallization of the original glassy sample, which was confirmed by XRD and TEM results (see below). The measured crystallization temperature (438 K) agrees with the literature [27]. Similar

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Fig. 1. Optical micrograph of a melt-spun Mg 0.87 Ni 0.13 ribbon; (a) the bottom-side surface (side in contact with the Cu wheel during spinning; 1003); (b) the top-side surface taken with polarized light (303).

results were obtained for Mg 0.84 Ni 0.16 but no DSC peaks were observed for the other three (crystalline) alloys at temperatures up to 723 K. Fig. 7 compares XRD patterns for Mg 0.87 Ni 0.13 in the as-spun and various annealed conditions. Annealing was carried out in DSC experiments similar to the one shown in Fig. 6, but varying the maximum temperature as indicated in Fig. 7. These results deliver conclusive information: First, crystallization between 413 and 448 K

can be confirmed as sharp diffraction peaks first arise for the latter temperature. Second, a further, gradually progressing phase transformation occurs between about 523 and 623 K (which could not be observed in DSC experiments). This can be concluded since the diffraction patterns for 473 and 673 K are completely different. Careful analysis shows that, even though weak XRD peaks indicating the presence of a small amount of hexagonal Mg crystallites arise immediately after crystallization, only at

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Table 1 Elementary analysis and Mg / Ni weight ratio of melt-spun Mg–Ni alloy ribbons as determined by ICP Mass%

Mg

Mg 0.87 Ni 0.13

Mg 0.84 Ni 0.16

Mg 0.74 Ni 0.26

Mg 2 Ni

Mg Ni Cu Mn Fe Pb S P Ca Others Mg / Ni (measured) Mg / Ni (nominal)

98.44 – 0.05 0.01 – 0.24 0.30 – 0.02 0.94 8 8

72.62 26.70 0.01 0.02 0.01 0.13 0.13 0.01 ? 0.37 272 325

67.17 31.69 0.03 0.01 0.01 0.15 0.12 0.03 ? 0.79 215 233

54.81 44.92 0.03 0.02 0.01 – 0.12 0.02 ? 0.07 122 150

44.24 54.81

a

0.07 0.01 0.07 0.03 0.03 ? 0.74 81 83

Measured by combustion-gas infrared absorption analysis.

Table 2 Chemical composition of five different Mg–Ni ribbons (A–E) melt-spun in one and the same batch (ICP results; nominal composition: Mg–23.5 mass% Ni) Ribbon→

A1

A2

A3

B

C

D

E

Mass% Mg Mass% Ni Impurities

75.90

75.98

75.47

75.82

75.90

75.33

76.37

23.00

23.07

23.09

23.24

23.08

22.97

23.26

1.10

0.95

1.44

0.94

1.02

1.70

0.37

673 K the sample has completely transformed into both equilibrium phases (hexagonal Mg and Mg 2 Ni). Third, Mg crystallites with a highly preferred orientation in the [001] lattice direction are formed during this second phase transformation. From these results we conclude that annealing glassy Mg–Ni alloys at temperatures between crystallization and about 523 K leads to the formation of

an intermediate metastable crystalline phase. A metastable phase has also been observed by Sommer et al. in annealed Mg–Ni alloys prepared by rapid quenching with a different technique [26]. Due to its complexity, we have not yet determined the crystal structure of the metastable phase. Fig. 8 shows a selected-area electron diffraction pattern for Mg 0.87 Ni 0.13 after annealing at 463 K in a DSC experiment. According to the XRD results shown in Fig. 7, this sample is expected to consist of the described intermediate metastable crystalline phase, together with small amounts of hexagonal Mg. The crystallization of the sample can be confirmed since sharp spots are now observed instead of diffuse diffraction rings (compare with Fig. 5). The fact that several, apparently randomly distributed diffraction spots arise, indicates that they are generated by several different crystallites. Since an area of only about 0.05 mm 2 contributed to the pattern, this result is an indication that quite small crystallites (f #200 nm) have

Fig. 2. Ni distribution in a melt-spun Mg 0.87 Ni 0.13 ribbon as determined by Ni Ka X-ray pulse mapping.

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Fig. 5. TEM image of the same Mg 0.87 Ni 0.13 sample as used in Fig. 4.

predominant metastable Mg–Ni phase was confirmed since electron diffraction spots corresponding to the Mg(101) and Mg(202) planes were identified. Fig. 9 shows a darkfield TEM image of the same sample; it was taken using the most intense diffracted beam observed in Fig. 8. Bright contrast arises wherever there is strong diffraction. The two brightest diffraction spots were most probably generated by the crystallite appearing very bright near the middle of Fig. 9, which should therefore correspond to the described metastable phase. The contrasts we observe in this micrograph allow an estimation of the crystallite size to f #150 nm. Fig. 3. XRD patterns of as-spun Mg–Ni alloy ribbons for all five compositions studied here.

been formed. The brightest diffraction spots in Fig. 8 agree well with the most intense XRD peaks observed for the same alloy annealed at 175 and 473 K (Fig. 8). The existence of some hexagonal Mg crystallites besides the

Fig. 4. Selected-area electron diffraction pattern for as-spun Mg 0.87 Ni 0.13 .

3.5. Activation and hydrogen absorption characteristics Fig. 10 shows the result of a PDSC experiment with Mg 0.87 Ni 0.13 run at constant volume in a H 2 gas atmosphere; the system pressure varied between about 1 and 2 MPa. Thermal cycling was carried out at a rate of 5 K min 21 between about 333 and 393 K for the first two cycles and afterwards between 333 and 473 K. Besides the exothermic peak owing to the crystallization of the sample, which was already observed in conventional DSC experiments run in a N 2 atmosphere (Fig. 6), no other peaks indicating H 2 absorption or desorption arise. From this and similar results, also with Mg 0.84 Ni 0.16 , we can conclude that repeated cycling of as-spun glassy Mg–Ni alloy ribbons between 333 and 393 to 423 K and H 2 gas pressures of up to 2 MPa at PDSC experimental conditions leads to no activation. Isothermal pressure cycling between 0.01 and 2.6 MPa at temperatures of up to 423 K at PDSC experimental conditions also leads to no activation (of Mg 0.87 Ni 0.13 ). The pressure and temperature conditions required for hydrogen gas activation of these alloys can be deduced from the PDSC result shown in Fig. 11. This experiment was run at 10 K min 21 between about 353 and 763 K, also isochorically, with the H 2 pressure varying

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Fig. 6. DSC and TGA results for Mg 0.87 Ni 0.13 measured in inert atmospheres. Heating and cooling rates were 5 K min 21 .

between about 1 and 3 MPa. Since no exothermic peak besides the one due to crystallization arises in the first heating ramp, we conclude that temperatures above 663 K (the decomposition temperature of MgH 2 , the most stable

Fig. 8. Selected-area electron diffraction pattern for Mg 0.87 Ni 0.13 after annealing at 463 K.

Fig. 7. XRD patterns for Mg 0.87 Ni 0.13 in the as-spun condition and annealed to the indicated temperatures in DSC experiments.

Fig. 9. Dark-field TEM image of the same sample as used in Fig. 8. Bright contrast arises wherever there is strong diffraction.

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Fig. 10. PDSC result for Mg 0.87 Ni 0.13 . The experiment was run isochorically at relatively low temperatures; the H 2 gas pressure varied between about 1 and 2 MPa. Heating and cooling rates were 5 K min 21 .

hydride expected in this system at the present H 2 pressure conditions) are required for activation. At the following cooling ramp (Fig. 11), however, an exothermic peak due to hydrogen absorption first arises at about 663 K, indicating at least partial activation of the sample. (Since the system pressure could not be measured precisely, it is impossible to determine from these results if MgH 2 or Mg 2 NiH 4 , or both, are involved. However, we found indications that Mg 2 NiH 4 was only formed after six cycles in a PDSC experiment run at similar conditions with Mg 0.84 Ni 0.16 ). An endothermic peak indicating the decomposition of the previously formed hydride arises in the second heating ramp in Fig. 11. Since both the desorption and absorption peak areas (which are proportional to the amount of hydrogen involved) increase in the second cycle (a tendency which continues for several cycles), we

conclude that activation occurs gradually. For comparison, a further PDSC experiment carried out in the isobaric mode between the same temperatures as the ones shown in Fig. 11, but at a lower H 2 pressure of only 0.4 MPa, led to no H 2 absorption after seven cycles. At PCT experiments, activation of Mg 0.87 Ni 0.13 was achieved rather suddenly during the third pressure cycle at 673 K (shown above). Comparison of PDSC results for Mg 0.87 Ni 0.13 , Mg 0.84 Ni 0.16 and Mg 2 Ni (the last two not shown here) indicate that the activation characteristics improve with increasing Ni contents, in good agreement with the well known catalytic effect of this metal [32,33]. The above-described results indicate that, at least without a suitable surface pretreatment, it is impossible to activate – and therefore hydride and use – melt-spun glassy Mg–Ni alloys without previously inducing their

Fig. 11. PDSC result for Mg 0.87 Ni 0.13 . The experiment was run isochorically between about 353 and 763 K; the H 2 gas pressure varied between about 1 and 3 MPa. Heating and cooling rates were 10 K min 21 .

G. Friedlmeier et al. / Journal of Alloys and Compounds 292 (1999) 107 – 117

crystallization at temperatures which are high enough for activation (the crystallization temperature is about 438 K, see above). This stands in contrast to nanocrystalline and ‘amorphous’ Mg–Ni alloys prepared by ball milling, which absorb hydrogen at lower temperatures [13,14]. This difference is most probably related to the fact that activation is a surface phenomenon: While ball-milled alloys have a large specific surface associated with large amounts of active centers for hydrogen chemisorption, melt-spun alloys have a relatively low specific surface which is generally passivated by an oxide layer. For making it permeable to hydrogen gas, relatively high temperatures (in the order of 673 K) are often required [34]. A solution to the stated problem could be a suitable surface treatment, like electrochemical activation or fluorination [35,36], which have proven to be effective for reducing the activation temperature of various hydride materials. Fig. 12 shows PCT curves for Mg 0.87 Ni 0.13 measured at 598 K, 623 and 648 K. Activation was carried out at 673 K, as described above. The sample was hydrided and dehydrided three times before the shown experiments were carried out, so these results are expected to represent the thermodynamically stable phases Mg and Mg 2 Ni and their hydrides. Two plateaus can be recognized in each PCT curve. Comparison with literature data [37] confirms that the longer and lower plateau in each curve corresponds to the formation or decomposition of MgH 2 , while the other one corresponds to the more unstable Mg 2 NiH 4 hydride. The rather large plateau slope can be attributed to sample temperature inhomogeneities. In good agreement with published results for these phases [37], the plateau pressure hysteresis is rather small for Mg and larger for Mg 2 Ni. The measured hydrogen capacity of this alloy (about 5.9

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mass%) is very close to the maximum, assuming a complete conversion into MgH 2 and Mg 2 NiH 4 (6.1 mass%). Fig. 13 compares XRD patterns for Mg 0.87 Ni 0.13 in the as-spun and as-annealed (at 723 K in N 2 ) conditions as well as after PDSC and PCT experiments. The latter samples are expected to show their respective maximum ever reached hydrogen contents (cooling down to room temperature after the experiments were carried out in a H 2 atmosphere). The XRD pattern of the PCT sample shows peaks corresponding to the two expected hydride phases, MgH 2 and Mg 2 NiH 4 . The complete hydriding of the sample can be confirmed since no peaks corresponding to metallic Mg arise. The sample PDSC (1) in Fig. 13 is the same as the one used for the PDSC experiment shown in Fig. 11. We can see diffraction peaks both of the hydride and the metallic phases, confirming that the sample had been hydrided only partially. The sample PDSC (2) had been cycled seven times isobarically at 0.4 MPa H 2 between 80 and 773 K. No peaks corresponding to the hydride phases arise, confirming that the sample was not activated. The pattern is similar to the one corresponding to the sample that was annealed to 723 K in a N 2 atmosphere.

4. Conclusions We succeeded in preparing glassy and nanocrystalline Mg–Ni alloy ribbons by rapid quenching of the melt with the melt-spinning technique. The obtained ribbons are highly homogeneous, relatively pure, stable in air and corrosion-resistant. Glassy Mg–Ni alloys show crystalliza-

Fig. 12. PCT results for Mg 0.87 Ni 0.13 measured at 598, 623 and 648 K.

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electrochemical activation or a surface pretreatment like fluorination are candidates for solving the activation difficulty. The fact that intermediate metastable phases can be produced by melt-spinning in connection with annealing at intermediate temperatures offers an alternative way to search for new hydride phases.

Acknowledgements We thank Dr. H. Enoki, Dr. D.L. Sun and Dr. P. Tessier of the National Institute of Materials and Chemical Research, for their assistance regarding PDSC experiments and Rietveld analysis, and Dr. M. Tomita for his support concerning TEM.

References

Fig. 13. XRD patterns for Mg 0.87 Ni 0.13 in the as-spun and annealed (at 723 K in N 2 ) conditions as well as after PDSC and PCT experiments. PDSC (1): same sample as used in the experiment shown in Fig. 11. PDSC (2): experiment run isobarically at 0.4 MPa between about 353 and 773 K.

tion temperatures around 438 K (here Mg 0.87 Ni 0.13 and Mg 0.84 Ni 0.16 glassy alloys were studied). During crystallization, an intermediate metastable Mg–Ni phase is formed, which gradually transforms into the stable Mg and Mg 2 Ni phases if the alloys are annealed at temperatures above 523 K. Mg crystallites formed during this second phase transformation show a strongly preferred orientation in the [001] lattice direction. In this first work on the hydriding characteristics of melt-spun binary Mg–Ni alloys ever reported we found that hydrogen gas activation of glassy ribbons is impossible at temperatures below crystallization and H 2 pressures up to 2 MPa, at least without a surface pretreatment. Activation requires temperatures and pressures above 663 K and about 1.5 MPa, respectively. Thereafter a very high storage capacity of nearly 6 mass% H was achieved (with Mg 0.87 Ni 0.13 ). The activation behavior improves with increasing Ni content. Since the activation behavior is governed mainly by the surface condition,

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