Hydriding behavior of Mg–Al and leached Mg–Al compounds prepared by high-energy ball-milling

Journal of Alloys and Compounds 297 (2000) 282–293

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Hydriding behavior of Mg–Al and leached Mg–Al compounds prepared by high-energy ball-milling a a a, b b b S. Bouaricha , J.P. Dodelet , D. Guay *, J. Huot , S. Boily , R. Schulz b

a ´ ´ ´ , Canada J3 X 1 S2 , 1650 Blvd. Lionel-Boulet, C. P. 1020, Varennes, Quebec INRS-Energie et Materiaux ´ ´ , 1800 Blvd. Lionel-Boulet, C. P. 1000, Varennes, Technologies Emergentes de production et de stockage, Institut de Recherche d’ Hydro-Quebec ´ , Canada J3 X 1 S1 Quebec

Received 13 August 1999; accepted 1 September 1999

Abstract The structure and hydrogen absorption properties of Mg:Al alloys prepared by high-energy ball milling were studied over the whole compositional range. These materials were prepared in their as-milled and Al-leached forms. The latter are obtained from the former materials by leaching out Al in a 1 N NaOH solution. The structure of the various alloys was determined by X-ray diffraction. The structure of the material in the hydrided state was also determined in some cases. In the as-milled state, hcp Mg(Al) with a small proportion of Mg 17 Al 12 and fcc Al(Mg) are formed at Mg:Al (90:10) and (20:80) compositions, respectively. At intermediate (58:42) and (37:63) compositions, the intermetallic Mg 17 Al 12 and Mg 3 Al 2 phases are formed, respectively. Following leaching, the Al content of Mg:Al (90:10) and (20:80) varies from 10.4 and 77.0 to 3.0 and 51.0 at.%, respectively. In both cases, noticeable change in the XRD pattern confirms that bulk dissolution of Al has been achieved. There is a two-fold increase in the specific surface area of Mg:Al (90:10) following leaching of Al. In the case of Mg:Al with intermediate compositions, dissolution of Al, if any, does not lead to discernable modification in the structure of the material. The measured hydrogen capacity of the as milled material decreases with Al content, from H / M51.74 for pure un-milled Mg, to 1.38 for Mg:Al (90:10), and then to 1.05 for Mg:Al (75:25). In each case, there is a further 10–15% decline of the hydrogen absorption capacity after leaching. In the case of Mg:Al (58:42), which basically only contains a nanocrystalline Mg 17 Al 12 intermetallic phase, hydriding leads to the formation of MgH 2 and Al. This reaction is totally reversible and Mg 17 Al 12 is recovered upon de-hydriding. In each case, there is an increase in the kinetics of hydrogen absorption and desorption following Al leaching.  2000 Elsevier Science S.A. All rights reserved. Keywords: Hydrogen storage; Metal hydrides; Nanocrystalline materials; Lixiviation; Ball milling

1. Introduction Hydrogen is certainly one of the best alternatives to replace petroleum products as a clean energy carrier for both transportation and stationary applications. However, there remains a number of problems to be solved before large scale introduction of hydrogen in day-to-day human activities could occur. One of them is the development of compact, lightweight, cost-effective and safe means of storing hydrogen. One of the most attractive means of hydrogen storing is in the form of metal hydrides. To this end, magnesium and magnesium based alloys are particularly interesting, since they can store more hydrogen by weight than most of the other currently know metal hydrides (MgH 2 is equivalent *Corresponding author. E-mail address: [email protected] (D. Guay)

to 7.6 wt.% hydrogen). But there remains a number of problems to be solved with these systems since Mg and Mg-based alloys have: (i) very poor hydrogen absorption and desorption kinetics, and (ii) are too stable for most applications. This study will address the first of these problems. In recent years, there has been an increasing interest in the use of ball-milling to prepare several Mg-based hydrogen absorbing materials with improved properties. This interest stems mainly from the fact that Mg has a low miscibility with most transition metals, a low melting temperature and a high vapor pressure, which reduce the efficiency of more conventional processes to prepare modified Mg-based materials. Among the benefits of using ball-milling to prepare improved Mg-based materials, there are: i) the ease of formation of several alloys [1–3] and hydride phases [4]; ii) the obtainment of Mg-based materials in the nanocrystalline and / or amorphous state, with

0925-8388 / 00 / $ – see front matter  2000 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 99 )00612-X

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various amounts of dislocations and special defects with high binding energy for hydrogen [5,6]; iii) the generation of fresh and highly reactive surfaces during the milling operation which increase the hydrogen absorption rate [7]; iv) the ease of formation of nanocomposite, whereby a Mg-based compound is intimately mixed with an other element or compound that can act as a catalyst to improve the hydrogen absorption kinetic [8,9]. Despite these numerous advantages, it remains that most of the Mg-based materials obtained by ball-milling have a very low effective surface area, in the range of 1 m 2 g 21 [10]. For example, Chen et al. [11] have measured 0.48 m 2 g 21 for Mg 2 Ni synthesized by mechanical alloying. It would thus be highly desirable to be able to increase the effective surface area of Mg-based compound produced by ball-milling in order to improve further their hydrogen absorption / desorption rates. This can be achieved by first adding an element that is intimately mixed with the Mgbased compound, and then selectively removing it through dissolution [12,13]. In this paper, the hydriding behavior of Mg–Al and leached Mg–Al compounds prepared by high-energy ballmilling has been investigated. Mg–Al compounds spanning the whole compositional range were prepared and the hydriding properties of both as-milled and leached materials were studied.

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where l 50.15406 nm, t is the crystallite size and ´ is the strain. b is determined from the full width at half maximum (FWHM), according to the relation b 2 5 FWHM 2 2 B 2s [16], where B 2s is defined by the resolution of the apparatus. Thus, from the slope and the ordinate at origin of a plot of b cos u vs. sin u, one can find the values of t and ´, respectively. Scanning electron microscopy (SEM) micrographs were obtained with a Hitachi H-4700 electron microscope. The chemical composition (Mg and Al) of the samples were determined by Energy Dispersive X-ray (EDX) measurements. The powder specific surface areas were measured by N2 adsorption (multipoint BET) using a Quantachrome Autosorb Automated Gas Sorption System. The pressure–composition (P–c) isotherm and the kinetics of hydrogen sorption were measured on a specifically designed gas titration apparatus [17]. The samples were activated at 350 or 4008C, by exposure to hydrogen at high pressure. This was done until the hydrogen capacity of the material reached a maximum and did not evolve anymore with time. Generally, this state was reached after 12 h. After that, absorption and desorption cycles were made until the kinetics of hydrogen sorption did not vary from one cycle to the other. The P–c isotherms reported herein were obtained after activation of the Mg–Al compounds.

2. Experimental details The alloys were prepared from mixtures of pure elemental Mg (Alpha, 325 mesh, 99.99%) and Al (Alpha, 325 mesh, 99.99%). The milling was carried out in a SPEX 8000 shaker mill. The vial and the balls were in stainless steel. The vial has an internal diameter of 38.1 mm and a length of 47.6 mm. Three balls were used, one with a diameter of 14 mm and two with a diameter of 11 mm. The ball to powder ratio was 10:1. The handling of the samples was always done in an argon-filled glove box to avoid contamination with air. The leaching of Al was made by pouring |1 g of powder in |200 ml of a 1 N NaOH solution at room temperature. The suspension was stirred continuously in an ultrasonic bath during 1 h unless otherwise stated. After that period, the powder was filtered out, washed several times with de-ionized water and dried overnight at 708C in air. X-ray powder diffractograms were recorded on a Siemens D500 diffractometer with Cu K a radiation. The crystallite size and strain were calculated from the full width at half maximum of the reflection peaks using a Williamson–Hall style plot [14]. This was done using the following equation [15]

b cos u 5 l /t 1 4´ sin u

Fig. 1. X-ray diffraction patterns of Mg:Al (90:10) after several milling times: (A) 1 h, (B) 3 h, (C) 10 h, and (D) 30 h. The position of the diffraction peaks of Mg (1), Al (2), and Mg 17 Al 12 (3) are indicated.

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3. Results

3.1. Physico–chemical characterizations The X-ray diffraction patterns of Mg:Al (90:10) are shown in Fig. 1 as a function of the milling time. After 3 h of milling, the characteristic diffraction peaks of Al are no longer observed. Instead, a new series of peaks are seen, which belong to Mg 17 Al 12 . There is also a shift in the peak position of Mg toward the larger 2u angle values, along with an increase of the peak width. Further milling up to 40 h does not change the X-ray diffraction patterns to any significant extent. As far as it can be evaluated from Fig. 1, the ratio between the phase proportions of Mg and Mg 17 Al 12 does not evolve with the milling time for t larger than 3 h. The change in the lattice constants of Mg with respect to the milling time is shown in Fig. 2(A). After about 10 h of

milling, the values of a and c have decreased from their ˚ and c55.185 A. ˚ Further initial values to a53.193 A milling up to 30 h does not lead to any further reduction of the lattice constants. According to the Mg–Al binary phase diagram [18], Al is soluble in the hcp structure of Mg up to |11 at.% at 4378C. From the work of Busk [19], who studied the variation of the lattice constants a and c of Mg with the Al content of the hexagonal phase, the following relations can be found a 5 2.807 1 4.0234E 23 (100 2 XAl )

(1)

c 5 4.672 1 5.3864E 23 (100 2 XAl )

(2)

where XAl is the fraction of Al dissolved in magnesium. The lattice parameters of the hcp phase of Mg(Al) found after 10 h or more of milling correspond to XAl |4.5 at.%,

Fig. 2. Variation of the lattice parameters of Mg:Al (90:10) as a function of (A) the milling time, and (B) the leaching period.

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which is the solubility limit of Al in hcp Mg at |2508C. According to the Mg–Al binary phase diagram, higher value of XAl lead to the formation of Mg 17 Al 12 , which is consistent with the observation made previously in Fig. 1. The crystallite size and the strain were evaluated from the full width at half maximum of the reflection peaks. The crystallite size and the strain reach |60 nm and 0.20%, respectively, during the first 10 h of milling. Further milling up to 30 h does not cause any more variation in these values. Considering that there is no more change in the structural properties of this compound after 30–40 h of milling, a similar period was also used for the preparation of the other samples. The X-ray diffraction patterns after 40 h of milling of several Mg:Al mixtures spanning the whole compositional range are shown in Fig. 3. The most important features to be emphasized are: (i) As the Al content of the sample increases, the phase composition evolves from a mixture of Mg(Al)1 Mg 17 Al 12 (curves A and B), to almost pure Mg 17 Al 12 (curve C), and then to a mixture of Mg 2 Al 3 1Al(Mg) (curve D). In the case of Mg:Al (20:80), the characteristic diffraction peaks of fcc Al are observed, al-

Fig. 3. X-ray diffraction patterns of Mg:Al after 40 h of milling: (A) 90:10, (B) 75:25, (C) 58:42, (D) 37:63, and (E) 20:80. The position of the diffraction peaks of Mg (1), Mg 17 Al 12 (2), Mg 2 Al 3 (3), and Al (4) are indicated.

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though they are shifted to lower 2u values, which indicate the presence of the larger Mg atoms dissolved in that phase. The presence of a small amount of Mg 2 Al 3 cannot be totally ruled out. (ii) The various phases obtained in the Mg:Al compounds are close to those expected from the Mg–Al binary phase diagram [18]. (iii) The diffraction peaks of all Mg:Al samples are quite large, indicating that the crystallite sizes of the various phases are small. Since several samples contain two different phases, no attempt has been made to evaluate more precisely the crystallite size and the strain from the width of the diffraction peaks. The leaching of Al from the various Mg:Al compounds was performed by pouring them in a 1 N NaOH solution at room temperature. An evaluation of the efficiency of the leaching procedure could be obtained by looking at the amount of H 2 gas evolved during the operation. After rinsing, filtering and drying, the X-ray diffraction patterns of the powders were recorded to detect structural changes due to the leaching of Al. In the case of Mg:Al (58:42) and (37:63), the leaching treatment does not lead to any discernible change in the XRD histogram of the materials. This indicates that the phase structure and phase composition of the materials are not affected by the leaching treatment. This observation is consistent with the fact that there is very little gas evolution during the leaching operation, indicating that Al is not oxidized to any large extent. In the case of Mg:Al (75:25), there is some gas evolution, suggesting that Al is partially dissolved. This gas evolution is, however, much less important than in the case of Mg:Al (90:10) and (20:80), which suggests that a larger fraction of Al is oxidized and dissolved in the case of these latter two materials. Consistently, the XRD pattern of leached Mg:Al (75:25) is similar to that of the as milled material, while those of Mg:Al (90:10) and (20:80) are markedly different from the histograms of the corresponding as milled compounds. This would indicate that, in the case of Mg:Al (75:25), the oxidation and dissolution of Al is restricted to the surface of the powder particle. Considering that this latter material is made of a large fraction of non leachable Mg 17 Al 12 , this result is not so surprising. The XRD patterns of Mg:Al (20:80) after leaching for various periods of time are shown in Fig. 4. From t50 to t52 h, there is a steady displacement of the position of the Al diffraction peaks toward the larger 2u values, corresponding to a change in the value of the lattice parameter a ˚ to 4.1027 A, ˚ and then to 4.0762 A. ˚ At the from 4.1294 A same time, new phases are formed, which might consist of Mg 2 Al 3 , MgAl(OH) 14 (H 2 O) x or Mg 2 Al(OH) 7 depending on the leaching period. Considering the material obtained after 8 h of leaching (Fig. 4, curve D), the new phase formed upon leaching of Al is identified as Mg 2 Al(OH) 7 . The XRD patterns of Mg:Al (90:10) after various

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Fig. 4. X-ray diffraction patterns of milled Mg:Al (20:80) after various periods of leaching: (A) as-milled, (B) 10 min, (C) 2 h, and (D) 8 h. The position of the diffraction peaks of Al (1), MgAl(OH) 14 (H 2 O) x (2), Mg 2 Al 3 (3), and Mg 2 Al(OH) 7 (4) are indicated.

Fig. 5. X-ray diffraction patterns of milled Mg:Al (90:10) sample after various periods of leaching: (A) as-milled, (B) 10 min, (C) 2 h, and (D) 8 h. The position of the diffraction peaks of Mg (1), Mg 17 Al 12 (2), and Mg(OH) 2 (3) are indicated.

periods of leaching are shown in Fig. 5. Again, there is a steady displacement of the characteristic diffraction peaks of Mg towards smaller 2u angle values, indicating that Al atoms are dissolved from the hcp structure. For t52 and 8 h, the diffraction peaks of Mg(OH) 2 are also observed. The relative intensity of the (101) peak of Mg(OH) 2 and Mg(Al) does not vary for leaching periods exceeding 2 h, indicating that the relative proportion of these two phases is constant. The change in the lattice constants of the Mg(Al) phase with respect to the leaching period is depicted in Fig. 2(B). As expected, both the lattice parameters a and c increase with the leaching period to ˚ and c55.204 A ˚ after 2.5 h. reach values of a53.205 A These values are slightly lower than those of pure Mg ˚ and c55.2108 A) ˚ [18], and would indicate (a53.2094 A that some Al atoms are still present in the hcp phase. Referring to Eqs. 1 and 2, it can be estimated that there is still 1 at.% of Al in the hcp structure of Mg at the end of the leaching treatment. The relative Mg /Al chemical composition of the various samples has been evaluated by EDX. In the case of the as-milled Mg:Al materials, the Mg /Al ratio is close to that

expected from the nominal composition of the powders (see Table 1). After leaching, the Al contents of Mg:Al (90:10) and (20:80) are lower than the corresponding as milled compounds, indicating that a large fraction of the initial Al content has been removed during the leaching operation. At the opposite, very little (if any) Al has been leached from Mg:Al (75:25), (58:42) and (37:63). These results are fully consistent with the observations made previously on the structural evolution of Mg:Al compounds following the leaching treatment. The influence of the leaching treatment on the morphology of the various Mg:Al materials was studied by SEM. Following Al leaching, most of the Mg:Al compounds, with the exception of Mg:Al (90:10), do not show any marked change in their surface morphologies, although there is in some cases a small decrease of the particle size. In the case of Mg:Al (90:10), whose micrographs are shown in Fig. 6, the as milled material is made of very large particles with featureless surface, some of them being larger than 30 mm. After leaching, the surface of these particles is covered with a needle-like structure. As far as it can be told by examination of the SEM micrograph, the

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Table 1 Comparison between the relative elemental composition (Mg and Al) of nanocrystalline Mg:Al compounds as milled and after Al leaching Nominal Mg:Al composition

As milled

After leaching

Mg [at.%]

Al [at.%]

Surface area (m 2 g 21 )

Mg [at.%]

Al [at.%]

Surface area (m 2 g 21 )

90:10 75:25 58:42 37:63 20:80

89.6 75.4 58.0 40.5 23.0

10.4 24.6 42.0 59.5 77.0

0.6 0.9 – – –

97.0 77.8 58.3 39.8 49.0

3.0 22.2 41.7 62.2 51.0

1.3 1.4 – – –

Fig. 6. Scanning electron micrograph of Mg:Al (90:10) before (A) and after (B) leaching of Al.

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cavities formed between these needles are not very deep. In the case of Mg:Al (90:10) and (75:25), the specific surface area determined by BET measurements is increased by a factor of two after leaching. It is interesting to note that a much larger increase in the specific surface area of Mg was reached when Li was used as a sacrificial element [12]. It is possible to leach Al from Mg:Al compounds as long as their composition is such that they are either on the Mg- or Al-rich side of the compositional range. For intermetallic compounds like Mg 17 Al 12 and Mg 2 Al 3 , leaching of Al (if any) does not proceed to any significant extent. From an energetic point of view, the leaching of Al from Mg:Al alloys will depend on the partial excess heat of Mg:Al formation, the Gibbs energy of the Al oxidation reaction and the stability constant of the aluminum hydroxide ion. As shown by Brown and Pratt [20], the intermediate phases b and g in the Mg–Al phase diagram exhibit negative heats of formation, which pass through pronounced minima at compositions within their homogeneity ranges. It is believed that this extra stabilization energy is responsible for the absence of noticeable Al dissolution from Mg 17 Al 12 (g phase) and Mg 2 Al 3 (b phase). In the case of Mg:Al (20:80) and (90:10), the formation of Mg 2 Al(OH) 7 (see Fig. 4) and Mg(OH) 2 (see Fig. 5) can also contribute to slow the kinetics of Al dissolution.

3.2. Hydrogen absorption properties: P–c–T The pressure–composition (P–c) isotherms at 4008C of Mg:Al compounds before and after leaching are shown in Fig. 7. For comparison, the P–c isotherm of pure polycrystalline (un-milled) Mg is also shown. It exhibits plateau pressures of |15 and |12 atmospheres during absorption and desorption, respectively, with a maximum H / M ratio of 1.75, which is lower than the expected value of 2 for MgH 2 . This indicates that a small fraction of the polycrystalline Mg could not be properly activated. The free energy change of the hydriding reaction based on the variation of the plateau pressure with the temperature (Van’t Hoff equation) is 270.09 kJ / mole. This value is close to that found by Pedersen et al. [21]. From their data, the reversible formation of MgH 2 took place at a pressure of 20 bars, which is quite close to the value found in Fig. 7. The addition of Al to the initial powder mixture considerably reduced the sticking of Mg on the milling tools and Mg:Al compounds can be obtained with high yield. As shown in Fig. 7, the addition of Al reduces the reversible capacity of the hydride, from typically |1.7 for pure Mg to 1.25 for Mg:Al (90:10), and then to |1.0 for Mg:Al (75:25). In the case of Mg:Al (58:42), a well defined plateau is no longer visible, although there is a change in the slope of the curve at H / M|0.13, which is the H / M value at which the formation of hydride begins in

Fig. 7. Pressure–composition isotherm at 4008C of as milled (–h–) and leached (–s–) Mg:Al compounds. The curve of pure un-milled (–n–) polycrystalline Mg is also shown for comparison.

the preceeding two Mg:Al compounds. The P–c isotherm of Mg:Al (37:63) and (20:80) are not shown, hydrogen sorption in these compounds being negligible. The absence of a well defined plateau pressure in the as milled Mg:Al (58:42) material has prompted us to lower the temperature at which the P–c isotherms are measured. Thus, the P–c isotherms at 3508C of as-milled Mg:Al (75:25) and (58:42) are shown in Fig. 8. In the case of Mg:Al (58:42), two (sloping) plateaus are observed at |11 and |20 atmospheres, indicating that hydrogen sorption in Mg 17 Al 12 is a two step process. Very similar observations are made in the case of as-milled Mg:Al (75:25), although the second plateau is less clearly resolved than for Mg:Al (58:42). Also, as shown in Fig. 8, the maximum H / M ratio attainable in these compounds decreases as the Al content increases. However, when the maximum H / M ratio is normalized with respect to the Mg content of the sample, the values of [H / M] norm found for Mg:Al (75:25) and (58:42) are 1.69 and 1.81, respectively. These values are quite close to the maximum H / M ratio (1.75) of polycrystalline Mg at 4008C (see Fig. 7). This indicates that Al acts essentially as a diluent and does not change significantly the hydrogen absorption capacities of Mg in the materials.

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Fig. 8. Pressure–composition isotherm at 3508C of as-milled Mg:Al (75:25) and (58:42).

The XRD histograms of Mg:Al (58:42) at various stages of the hydriding process are compared in Fig. 9. In this experiment, the XRD histograms of the intermetallic phase is recorded after exposition to H 2 at high pressure for periods of time corresponding to H / M51.0 and 1.18 (curve B and C of Fig. 9, respectively). Upon partial hydriding (curve B), MgH 2, Al and Mg 2 Al 3 are observed, along with the starting Mg 17 Al 12 intermetallic phase. Following complete hydriding (curve C), MgH 2 and Al are observed, while the characteristic diffraction lines of the starting Mg 17 Al 12 intermetallic compound are no longer visible and that belonging to Mg 2 Al 3 are considerably reduced. Thus, hydrogen absorption into Mg 17 Al 12 yields to the following reactions: Mg 17 Al 12 1 (17 2 2y) H 2 → y Mg 2 Al 3 1 (17 2 2y) MgH 2 1 (12 2 3y) Al Mg 2 Al 3 1 2H 2 → 2MgH 2 1 3Al

(3) (4)

A more precise estimation of the value of y could be obtained through an evaluation of the various proportion of phases, which is beyond the scope of this work. It is interesting to note that these reactions are totally reversible and that the intermetallic Mg 17 Al 12 phase is recovered following complete de-hydriding (see curve D of Fig. 9). A similar phenomenon was also observed in the case of Mg 2 Al 3 [22,23], although in that later case the reaction was reversible only when the de-hydriding reaction was performed slowly and that sufficient time was allowed for the reactants to rearrange themselves [23,24]. In the present case complete reversibility could be observed even

Fig. 9. X-ray diffraction patterns of Mg:Al (58:42): (A) as-milled, (B) after partial hydriding, (C) after total hydriding, and (D) after dehydriding. The position of the diffraction peaks of Mg 17 Al 12 (1), Al (2), Mg 2 Al 3 (3), and MgH 2 (4) are indicated.

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in the case of fast de-hydriding at essentially zero pressure. As evidenced by the small width of the diffraction peaks, the crystallite size of the intermetallic phase recovered

after de-hydriding are larger than that of the starting material. The XRD histogram of Mg:Al (90:10) in the fully

Fig. 10. Variation of H / M with respect to time for as milled and leached Mg:Al compounds. These curves were obtained at 4008C, under 38 atmospheres of H 2 . The curve of pure un-milled polycrystalline Mg is also shown for comparison.

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Table 2 Characteristic T 90% and T 10% time of Mg:Al compounds (in s) Nominal Mg:Al composition

As milled Absorption T 0.9 a (s)

Desorption T 0.1 b (s)

Absorption T 0.9 a (s)

Desorption T 0.1 b (s)

90:10 75:25 58:42

1146 703 163

323 160 30

823 253 40

163 116 40

a b

After leaching

Characteristic time requested during absorption to reach a normalized H / M value of 0.9. Characteristic time requested during desorption to reach a normalized H / M value of 0.1.

hydrided state has been recorded to determine if the same reaction is also taking place. However, no evidence of the characteristic diffraction lines of Al has been found. In view of the low Al content (|4.5 at.%) of the Mg(Al) hcp phase, this may not be so surprising. Looking now at the P–c isotherms of the Mg:Al compounds after Al leaching, the following observations can be made: (i) In each case, there is a small reduction of the maximum H / M ratio. For Mg:Al (90:10) and (75:25), this reduction accounts for about 14 and 12% of the reversible hydrogen capacity of the as milled material (ii) In each case also, the plateau pressure of the leached materials are at a higher pressure than the corresponding plateau of the as milled material.

3.3. Hydrogen absorption properties: kinetics The H / M versus time curves of the various Mg:Al compounds are shown in Fig. 10. In order to compare the kinetic data of materials having different storage capacities, all these curves are expressed in term of the fraction of hydride, a, formed after time t, which is obtained by normalizing the H / M value by the maximum H / M value attained at the end of the absorption period. The time at which a 50.9 during absorption (denoted T 0.9 ), and a 5 0.1 during desorption (denoted T 0.1 ) are listed in Table 2. It is seen that T 0.9 and T 0.1 decrease as the Al content of

the as milled Mg:Al compounds increases. Also, the T 0.9 and T 0.1 values of all leached Mg:Al compounds are smaller than the corresponding values of the as milled materials. For comparison, the T 0.9 and T 0.1 value of pure polycrystalline Mg are 876 and 340 s, respectively, which is somewhat smaller than the values found for as milled Mg:Al (90:10), but larger than the values of all the other materials. Direct comparison between all these values is difficult since some samples are made of a mixture of phases (Mg:Al (75:25)) and others are not fully hydrided (Mg:Al (58:42)). For this reason, we will concentrate on the Mg:Al (90:10) samples, whose phase compositions bear close resemblance with pure un-milled Mg. In studying the kinetics of hydrogen absorption in materials, certain equations need to be considered. Some of these equations have been summarized and tabulated by Sharp et al. [25]. The fit of the experimental data to these equations have been tested. The various rate equations are summarized in Table 3, using the designation of Sharp et al. [25]. The maximum range of a, amax , within which a linear function is obtained for each of these equations are also given in Table 3. It is seen that the maximum range of linearity is obtained for F(a ) 5 D4 (a ), which represents a diffusion controlled reaction starting on the exterior of a spherical particle, and that 64% of the metal has reacted at amax . A similar analysis was performed on the kinetic data of both Mg:Al (90:10) samples. In each case, the maximum range of linearity at the beginning of the hydrogen

Table 3 Fit of various rate equations, F(a ), to the normalized kinetic absorption data of polycrystalline Mg at 4008C and initial H 2 pressure of 40 atmospheres F(a )a

Kinetic model

D4 (a ) 5 (1 2 2a / 3) 2 (1 2 a )2 / 3

Diffusion-controlled reaction starting on the exterior of a spherical particle One-dimensional diffusion Two-dimensional diffusion-controlled reaction into a cylinder Phase-boundary controlled reactions Phase-boundary controlled reactions First order reaction Nucleation and growth (Avrami–Erofe’ev) Nucleation and growth (Avrami–Erofe’ev)

D1 (a ) 5 a 2 D2 (a ) 5 (1 2 a ) ln(1 2 a ) 1 a R 2 (a ) 5 (1 2 (1 2 a )1 / 2 ) R 3 (a ) 5 (1 2 (1 2 a )1 / 3 ) F1 (a ) 5 ln(1 2 a ) A 2 (a ) 5 (2ln(1 2 a ))1 / 2 A 3 (a ) 5 (2ln(1 2 a ))1 / 3 a b

Designation of the various F(a ) as in Ref. [25]. amax is the maximum fraction reacted following a linear dependence between F(a ) and time.

amax b 0.64 0.47 0.45 0.27 0.27 0.21 ,0.09 ,0.09

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absorption is obtained for F(a ) 5 D4 (a ), which suggests that there is no change in the rate limiting step between all of these materials. The corresponding values of amax are given in Table 4. Also given are the rate constants K, obtained from the linear regions of the D4 (a ) vs. t curves. It is seen that there is almost no difference between the rate constant (K) of pure un-milled Mg and as milled Mg:Al (90:10), while K increases after Al leaching. The value of amax is almost constant through the samples. There have been a number of studies in the literature which have focused on the kinetic of hydrogen sorption in Mg and Mg alloys. For example, Stander [26] has shown that at the very beginning of the hydriding reaction, the reaction is controlled by the direct chemical reaction between hydrogen and the exposed metal surface. As expected [26], this expression is valid only at the very beginning of the hydriding process. After the metal particles have been covered by a hydride layer, the transport of material through the layer by diffusion becomes rate controlling, leading to kinetics for spherical particles that are governed by the expression denoted D4 (a ) in Table 3. According to Ref. [26], this expression describes adequately the kinetic of hydrogen sorption in Mg for values of a up to 0.50–0.55. The results shown in Table 3 for hydrogen absorption in Mg are fully consistent with these results, the kinetics of the process being best described by the rate equation D4 (a ) which is typical of the transport of material through a layer by diffusion. Moreover, the value of amax over which this expression is valid (amax 50.64) is close to that reported at lower temperature by Stander [26]. It would be tempting to attribute the almost three-fold increase in the rate constant of leached Mg:Al (90:10) to the increase in its surface area following removal of Al. However, as shown by Mintz et al. [27], the activation energies of hydrogenation of the Mg:Al alloys vary with the Al content. Thus, the ratio between the activation energy of hydrogenation, Ea , of Mg:Al alloys and pure Mg is 0.44 for [Al]50.4 at.% and increase up to 1 for [Al]54.5 at.%. In the case of as-milled Mg:Al (90:10), it was shown precedently that 4.5 at.% of Al is dissolved in the hcp phase of Mg. According to Mintz et al. [27], the ratio between the Ea values of Mg:Al (90:10) and pure Mg is one, which would translate in identical rate constant for both materials. This is consistent with what is found in Table 4. At the opposite, leached Mg:Al (90:10) contains only 1 at.% of Al, and the lower activation energy of Table 4 Values of rate constant (K) and amax of as milled and leached Mg:Al compounds Compounds Mg Mg:Al (90:10) as milled Mg:Al (90:10) after leaching a

Slope of the linear regression line.

Ka

amax 24

3.6E 3.2E 24 8.6E 24

0.64 0.68 0.72

hydrogenation would thus translate in larger rate constant, which is what is observed in Table 4. It would thus seem that, in a first approximation, the kinetics data of pure Mg and Mg:Al (90:10) materials can be accounted for by changes in the Al content of the hexagonal phase.

4. Conclusion It was shown that for Mg:Al (58:42), which is made of the Mg 17 Al 12 intermetallic phase, hydriding leads to the formation of MgH 2 and Al. This reaction is reversible and, upon de-hydriding, the intermetallic phase is formed back again. Significant leaching of Al from Mg:Al alloys occur only for samples whose compositions are at both ends of the compositional range. In the case of Mg:Al (90:10) powder, leaching of Al leads to a two-fold increase of its specific surface area. This is less important than the corresponding increase of surface area observed when Li is used as a sacrificial element. There is a small decrease of the hydrogen absorption capacity and an increase of the sorption kinetics of the leached Mg:Al (90:10) powder compared to the as-milled one. This change in the sorption kinetics is though to arise as a consequence of the presence of Al solutes in the hexagonal structure of Mg, rather than to be due to a purely geometric effect (increase of the surface area).

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