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Influence of thermal annealing on the hydrogenation properties of mechanically milled AB5-type alloys
Materials Science and Engineering B108 (2004) 76–80
Influence of thermal annealing on the hydrogenation properties of mechanically milled AB5-type alloys J.R. Ares∗ , F. Cuevas, A. Percheron-Guégan Laboratoire de Chimie Métallurgique des Terres Rares, ISCSA-CNRS, 2-8 Rue Henri Dunant, 94320 Thiais Cedex, Paris, France
Abstract The influence of mechanical milling and posterior annealing treatment on kinetic and thermodynamic H-properties of AB5 -based alloys are here reported. Short milling process yields non-homogeneous microstructural (wide distribution of nano-crystallite sizes and amorphous phase) and morphological alloy powders. Annealing treatment of long-milled alloy produces a homogeneous microstructural alloy. Annealed alloy exhibits a unique nano-crystallite size without presence of amorphous structure. H-capacity in short-milled alloy is higher than for annealed alloy due to the contribution of both amorphous and well crystallised domains. The short-milled alloy exhibits extra H-trapping sites in amorphous domains as detected by thermal desorption. Absorption hydrogen kinetic in short-milled alloy is faster than for annealed alloy mainly due to fast H-transport through amorphous domains. © 2003 Elsevier B.V. All rights reserved. Keywords: Hydrides; Intermetallic compounds; Mechanical milling; Nano-crystalline
1. Introduction In the last 30 years, LaNi5 has been the most investigated intermetallic compound as hydrogen absorber [1]. However, despite of its high hydrogen capacity (one hydrogen atom to metal atom (H/M)) and easy activation, the binary compound can not be used in battery electrodes due to its high plateau pressure (0.17 MPa at room temperature) and short cycle-life [2]. These problems can be solved by substitutions of La (A-type element) and Ni (B-type element) for another rare earths (A = Ce) [3] and transition metals (B = Mn, Co, Al) [4], respectively. Thus, multisubstituted AB5 -based alloys are commercially used in Ni–MH batteries [5]. Unfortunately, these substitutions produce a diminution of H-capacity as well as a reduction of the hydrogenation kinetics [6]. As result of this, AB5 alloys are not well adapted for battery applications where fast discharge kinetics is required. Recently, in order to improve hydrogenation kinetics of intermetallic compounds, non-equilibrium preparation methods such as melt–spinning, sputtering and mechanical ∗ Corresponding author. Tel.: +33-1-4978-1225; fax: +33-1-4978-1203. E-mail address: [email protected] (J.R. Ares).
0921-5107/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2003.10.083
milling have been used [7,8]. The resulting alloys exhibit particular structural characteristics unachieved by equilibrium preparation methods: nano-crystalline grain size with a high density of grain boundaries and even lacking of long-range order (amorphous state). These microstructures currently provide fast hydrogenation kinetics and better cycle-life behaviour [9,10]. Among all non-equilibrium preparation methods, mechanical milling (MM) has become a popular technique because of its simplicity, relative inexpensive equipment and applicability to most intermetallic compounds [11]. MM has been used in several hydrogen storage alloys, such as FeTi and Mg2 Ni [10], with an increasing of their absorption–desorption hydrogenation rates. As concerns AB5 -type compounds, MM has been applied in binary LaNi5 [12,13]. A good improvement in hydrogen activation and kinetics is observed. On the other hand, a significant loss of hydrogen capacity occurs. Capacity loss has been attributed to amorphous formation [14]. In this paper, we report on the hydrogenation properties of a short-milled AB5 -based alloy as compared to those of a long-milled alloy after annealing treatment. The differences in their H-properties will be discussed in terms of their distinct microstructures.
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2. Experimental A commercial-type Mm(Ni, Mn, Al, Co)5.2 alloy ingot, 10 g in mass, was prepared by induction melting. A part of this ingot (3 g) was sieved under 63 m, introduced in a vial with stainless-steel balls (7 mm in diameter) and mechanically grounded in a Fritsch P7 planetary mill. Milling was conducted at vial rotation speed of 400 rpm for 200 min under argon atmosphere, with ball to powder ratio 12:1. This milled sample is labelled as MS. A second sample was obtained by the same procedure but with a much longer milling time (4000 min) to obtain an amorphous-like alloy. It was then annealed at 600 K during 14 h under argon atmosphere. This annealed sample is labelled as AS. Alloy structural properties were studied by X-ray diffraction in a – D5000 Siemens Diffractometer using Cu K␣ radiation. Diffraction patterns were analysed by the Rietveld Method using the Fullproff software [15]. Sample morphology was analysed by Scanning Electron Microscopy (SEM). Hydrogenation thermodynamics were studied from Pressure–composition–isotherm (P–C–I) curves obtained at 298 K, with an experimental set-up based in Sievert’s Method. H-kinetic properties were studied by hydrogenation curves at hydrogen quasi-constant pressure. Thermal desorption analyses have been carried out within the temperature range 300–800 K at a heating rate 10 K/min under primary vacuum.
3. Results and discussion 3.1. Structural characterisation Fig. 1 shows XRD patterns of the MS and AS samples together with those of the precursor samples, i.e., the starting alloy ingot and the long-milled alloy, respectively. All patterns are single-phase and can be indexed to the hexagonal CaCu5 -type structure. As expected, diffraction peaks broaden with milling time (Fig. 1a and b) and narrow on annealing (Fig. 1c and d). Cell parameters and average value of coherent diffraction domains (average crystallite size D) for all patterns are given in Table 1. Crystallite size is evaluated using a “pseudo-Voigt” profile function [16] as will be discussed in more detail in a forthcoming publication.
Table 1 Structural parameters (a, c, V) and crystallite sizes D of starting, MS, long milled and AS alloys
Starting alloy MS Long-milled alloy AS
a (Å)
c (Å)
V (Å3 )
D (nm)
4.9862(1) 4.988(1) 4.979(6) 4.951(1)
4.0552(1) 4.057(1) 4.131(6) 4.076(1)
87.31(1) 87.43(3) 88.70(20) 86.57(2)
68(6) 11(1) 4(1) 7(2)
Numbers between brackets refer to standard deviations.
Fig. 1. XRD patterns of (a) starting alloy, (b) short-milled alloy, (c) longmilled alloy and (d) annealed-treated alloy. Peak indexation to CaCu5 -type structure is given for the starting alloy.
D-values are very close for MS and AS samples: 11 ± 1 and 7 ± 2 nm, respectively. However, as shown in Fig. 2, a more detailed inspection of their XRD peak profiles reveals strong differences between them. For the annealed sample AS, a reasonable Rietveld fitting (Rwp = 11.6%) with a unique crystallite size (D = 7 ± 2 nm) is achieved. On the contrary, for the milled sample MS, several crystallite sizes should be considered to account for large broadening at the base of the peaks and narrow profiles at their top. Besides, MS exhibits a diffuse diffraction halo between 30 and 50◦ denoting a certain amount of amorphous phase. Thus, a good Rietveld fitting of MS (Rwp = 8.21%) is obtained summing up the contributions of two crystallite sizes (D1 = 68 ± 10 nm and D2 = 9 ± 3 nm) and an amorphous background. The relative contribution between the two crystallite sizes is 1:4, respectively. Fig. 3 shows SEM micrographs for MS(a) and AS(b) alloys. MS exhibits a wide particle size distribution comprised between 5 and 60 m. Particles are bounded by cleavage facets. On the contrary, for AS alloy, particles are rounded and possess a narrow particle size distribution (15 ± 5 m). The particular morphology of AS alloy is related to particle agglomeration at long-milling times, a phenomenon already observed in other intermetallic alloys such as Mg2 Ni [17]. To summarise, AS alloy exhibits a homogeneous microstructure with crystallite and particle size of 7 nm and
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Fig. 4. P–C–I absorption curves at 298 K of un-milled starting alloy, MS and AS alloys.
on AB5 -type alloy is a very inhomogeneous process at short milling times. 3.2. Hydrogenation properties
Fig. 2. Rietveld analysis of (a) MS and (b) AS alloys. Dashed lines in (a) stands for contributions of distinct crystallite-size domains D1 , D2 and amorphous background. Continuous lines in (a) and (b) stands for the global Rietveld fitting.
15 m, respectively. In contrast, the microstructure of MS is rather inhomogeneous both in particle and crystallite sizes. This is particularly severe for crystallite-size as a wide distribution of amorphous and several nano-crystalline domains coexist in MS alloys. This proves that mechanical milling
Fig. 4 shows the P–C–I absorption curves at 298 K of MS and AS alloys as compared with that of the starting alloy. AS alloy does not exhibit any plateau pressure but an increased hydrogen solubility in the ␣-solute branch. This effect has been already reported in nano-crystalline materials [18]. As concerns MS alloy, the total hydrogen capacity is higher than AS. This increase is assigned to large crystallite domains, which remain in the AB5 -type alloy after short milling. In addition, a further increment of the solute phase is observed in MS. The expansion of the H-solubility region at low hydrogenation pressures can be ascribed to additional trapping sites caused by amorphous domains in the inhomogeneous MS alloy. To get further information on the H-trapping states, thermal desorption experiments were conducted for the MS, AS
Fig. 3. SEM micrographs of (a) MS alloy and (b) AS alloy.
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faster for MS than for AS. For instance, the time for reaching a reacted fraction of 0.8 is only 200 s for MS, whereas it is as long as 4000 s for AS. Taking into account the microstructural properties of both alloys, such remarkable difference is thought to be due to the occurrence of the amorphous phase for MS alloy. Amorphous domains act as hydrogen traps for small hydrogen concentrations but, at higher concentrations, H-diffusion may drastically increase with the saturation of deep-energy traps. This effect has been clearly demonstrated by Kirchheim et al. in Pd–Si alloys [20]. In addition, SEM micrographs show that AS alloys exhibits a smaller surface-to-volume ratio than MS as a result of particle agglomeration at long milling time. This also contributes to faster kinetic for MS alloy as compared to AS alloy. Fig. 5. Thermal desorption spectra of un-milled starting alloy, MS and AS alloys. The spectra are normalised to the same alloy mass.
4. Conclusion
and starting alloys (Fig. 5). H-desorption rate from AS and MS alloys are much higher than that for the starting alloy, in agreement with their higher H-solubility for the ␣-phase. In fact, the starting alloy only exhibits a small peak at 450 K, which can be ascribed to H-trapped in defects [19]. For nano-crystalline AS alloy, one desorption peak is observed at 410 K which extends up to 625 K. For the non-homogeneous MS alloy, the same trend is observed but extra desorption states appear at T > 550 K. These states can be attributed to H-trapped in tetrahedral sites with a high rare-earth concentration [7]. Hydriding absorption kinetics of MS and AS alloys were measured at 298 K with a quasi-constant pressure of 0.9 MPa. Reacted fraction curves (defined as the hydrogen moles absorbed at any arbitrary time divided by the number of moles which can be absorbed at the infinite time) were evaluated and are shown in Fig. 6. H-kinetics are much
Hydrogenation properties of two AB5 -type alloys with distinct microstructural properties but comparable average crystallite sizes have been investigated. One of them, MS, was short-milled in a planetary mill. This treatment produces a non-homogeneous microstructure composed of a wide distribution of amorphous and nano-crystallite sizes. The other, AS, was annealed after long mechanical milling, resulting in a homogeneous nano-crystallite microstructure. As concerns H-thermodynamics, both MS and AS alloys exhibit no significant plateau pressure and a higher H-solubility in the solute-␣ phase as compared with un-milled AB5 -type alloy. Widening of the ␣-branch is bigger for MS than for AS due to H-trapping in amorphous sites for the former alloy, as corroborated by thermal desorption measurements. As concerns H-kinetics, hydrogen absorption is much faster in MS than in AS alloy, which can be mainly ascribed to enhanced hydrogen diffusivity through amorphous domains.
Acknowledgements The authors wish to thank F. Briaucourt for technical assistance and J.L. Pastol for SEM measurements.
References
Fig. 6. Time–evolution of the alloy reacted fraction for MS and AS alloys during hydrogen absorption at 0.9 MPa, 298 K.
[1] J.H.N. Van Vucht, F.A. Kuijpers, H.C.A.M. Bruning, Philips Res. Rep. 25 (1970) 133. [2] H.F. Bittner, C.C. Badcock, J. Electrochem. Soc. 130 (1983) 193c. [3] J.-M. Joubert, M. Latroche, A. Percheron-Guégan, F. BouréeVigneron, J. Alloys Compd. 275 (1998) 118. [4] C. Lartigue, A. Percheron-Guégan, J.C. Achard, F. Tasset, J. Less-Common Metals 75 (1980) 23. [5] F. Feng, M. Geng, D.O. Northwood, Int. J. Hydrogen Energy 26 (2001) 725.
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[6] H. Ogawa, M. Ikoma, H. Kawano, I. Matsumoto, J. Power Sources 12 (1988) 393. [7] F. Cuevas, M. Hirscher, Acta Mater. 51 (2003) 701. [8] H. Niu, D.O. Northwood, Int. J. Hydrogen Energy 27 (2002) 69. [9] Y. Li, Y.-T. Cheng, M.A. Habib, J. Alloys Compd. 209 (1994) 7. [10] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Compd. 253 (1997) 70. [11] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1. [12] G. Liang, J. Hout, R. Schulz, J. Alloys Compd. 320 (2001) 133. [13] S. Corré, M. Bououdina, N. Kuriyama, D. Fruchart, G. Y Adachi, J. Alloys Compd. 29 (1999) 166.
[14] R. Kirchheim, F. Sommer, G. Schluckebier, Acta Metall. 30 (1982) 1059. [15] J. Rodr`ıguez-Carvajal, Physica B 192 (1993) 55. [16] P. Thompson, J.J. Reilly, J.M. Hastings, J. Less-Common Metals 129 (1987) 105. [17] D. Cracco, A. Percheron-Guégan, J. Alloys Compd. 268 (1998) 248. [18] N. Mommer, M. Hirscher, F. Cuevas, H. Kronmüller, J. Alloys Compd. 266 (1998) 255. [19] E.H. Kisi, E. Wu, M. Kemali, J. Alloys Compd. 330 (2002) 202. [20] R. Kirchheim, T. Mütschele, W. Kieninger, H. Gleiter, R. Birringer, T.D. Koblé, Mater. Sci. Eng. 99 (1988) 457.
Influence of thermal annealing on the hydrogenation properties of mechanically milled AB5-type alloys J.R. Ares∗ , F. Cuevas, A. Percheron-Guégan Laboratoire de Chimie Métallurgique des Terres Rares, ISCSA-CNRS, 2-8 Rue Henri Dunant, 94320 Thiais Cedex, Paris, France
Abstract The influence of mechanical milling and posterior annealing treatment on kinetic and thermodynamic H-properties of AB5 -based alloys are here reported. Short milling process yields non-homogeneous microstructural (wide distribution of nano-crystallite sizes and amorphous phase) and morphological alloy powders. Annealing treatment of long-milled alloy produces a homogeneous microstructural alloy. Annealed alloy exhibits a unique nano-crystallite size without presence of amorphous structure. H-capacity in short-milled alloy is higher than for annealed alloy due to the contribution of both amorphous and well crystallised domains. The short-milled alloy exhibits extra H-trapping sites in amorphous domains as detected by thermal desorption. Absorption hydrogen kinetic in short-milled alloy is faster than for annealed alloy mainly due to fast H-transport through amorphous domains. © 2003 Elsevier B.V. All rights reserved. Keywords: Hydrides; Intermetallic compounds; Mechanical milling; Nano-crystalline
1. Introduction In the last 30 years, LaNi5 has been the most investigated intermetallic compound as hydrogen absorber [1]. However, despite of its high hydrogen capacity (one hydrogen atom to metal atom (H/M)) and easy activation, the binary compound can not be used in battery electrodes due to its high plateau pressure (0.17 MPa at room temperature) and short cycle-life [2]. These problems can be solved by substitutions of La (A-type element) and Ni (B-type element) for another rare earths (A = Ce) [3] and transition metals (B = Mn, Co, Al) [4], respectively. Thus, multisubstituted AB5 -based alloys are commercially used in Ni–MH batteries [5]. Unfortunately, these substitutions produce a diminution of H-capacity as well as a reduction of the hydrogenation kinetics [6]. As result of this, AB5 alloys are not well adapted for battery applications where fast discharge kinetics is required. Recently, in order to improve hydrogenation kinetics of intermetallic compounds, non-equilibrium preparation methods such as melt–spinning, sputtering and mechanical ∗ Corresponding author. Tel.: +33-1-4978-1225; fax: +33-1-4978-1203. E-mail address: [email protected] (J.R. Ares).
0921-5107/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2003.10.083
milling have been used [7,8]. The resulting alloys exhibit particular structural characteristics unachieved by equilibrium preparation methods: nano-crystalline grain size with a high density of grain boundaries and even lacking of long-range order (amorphous state). These microstructures currently provide fast hydrogenation kinetics and better cycle-life behaviour [9,10]. Among all non-equilibrium preparation methods, mechanical milling (MM) has become a popular technique because of its simplicity, relative inexpensive equipment and applicability to most intermetallic compounds [11]. MM has been used in several hydrogen storage alloys, such as FeTi and Mg2 Ni [10], with an increasing of their absorption–desorption hydrogenation rates. As concerns AB5 -type compounds, MM has been applied in binary LaNi5 [12,13]. A good improvement in hydrogen activation and kinetics is observed. On the other hand, a significant loss of hydrogen capacity occurs. Capacity loss has been attributed to amorphous formation [14]. In this paper, we report on the hydrogenation properties of a short-milled AB5 -based alloy as compared to those of a long-milled alloy after annealing treatment. The differences in their H-properties will be discussed in terms of their distinct microstructures.
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2. Experimental A commercial-type Mm(Ni, Mn, Al, Co)5.2 alloy ingot, 10 g in mass, was prepared by induction melting. A part of this ingot (3 g) was sieved under 63 m, introduced in a vial with stainless-steel balls (7 mm in diameter) and mechanically grounded in a Fritsch P7 planetary mill. Milling was conducted at vial rotation speed of 400 rpm for 200 min under argon atmosphere, with ball to powder ratio 12:1. This milled sample is labelled as MS. A second sample was obtained by the same procedure but with a much longer milling time (4000 min) to obtain an amorphous-like alloy. It was then annealed at 600 K during 14 h under argon atmosphere. This annealed sample is labelled as AS. Alloy structural properties were studied by X-ray diffraction in a – D5000 Siemens Diffractometer using Cu K␣ radiation. Diffraction patterns were analysed by the Rietveld Method using the Fullproff software [15]. Sample morphology was analysed by Scanning Electron Microscopy (SEM). Hydrogenation thermodynamics were studied from Pressure–composition–isotherm (P–C–I) curves obtained at 298 K, with an experimental set-up based in Sievert’s Method. H-kinetic properties were studied by hydrogenation curves at hydrogen quasi-constant pressure. Thermal desorption analyses have been carried out within the temperature range 300–800 K at a heating rate 10 K/min under primary vacuum.
3. Results and discussion 3.1. Structural characterisation Fig. 1 shows XRD patterns of the MS and AS samples together with those of the precursor samples, i.e., the starting alloy ingot and the long-milled alloy, respectively. All patterns are single-phase and can be indexed to the hexagonal CaCu5 -type structure. As expected, diffraction peaks broaden with milling time (Fig. 1a and b) and narrow on annealing (Fig. 1c and d). Cell parameters and average value of coherent diffraction domains (average crystallite size D) for all patterns are given in Table 1. Crystallite size is evaluated using a “pseudo-Voigt” profile function [16] as will be discussed in more detail in a forthcoming publication.
Table 1 Structural parameters (a, c, V) and crystallite sizes D of starting, MS, long milled and AS alloys
Starting alloy MS Long-milled alloy AS
a (Å)
c (Å)
V (Å3 )
D (nm)
4.9862(1) 4.988(1) 4.979(6) 4.951(1)
4.0552(1) 4.057(1) 4.131(6) 4.076(1)
87.31(1) 87.43(3) 88.70(20) 86.57(2)
68(6) 11(1) 4(1) 7(2)
Numbers between brackets refer to standard deviations.
Fig. 1. XRD patterns of (a) starting alloy, (b) short-milled alloy, (c) longmilled alloy and (d) annealed-treated alloy. Peak indexation to CaCu5 -type structure is given for the starting alloy.
D-values are very close for MS and AS samples: 11 ± 1 and 7 ± 2 nm, respectively. However, as shown in Fig. 2, a more detailed inspection of their XRD peak profiles reveals strong differences between them. For the annealed sample AS, a reasonable Rietveld fitting (Rwp = 11.6%) with a unique crystallite size (D = 7 ± 2 nm) is achieved. On the contrary, for the milled sample MS, several crystallite sizes should be considered to account for large broadening at the base of the peaks and narrow profiles at their top. Besides, MS exhibits a diffuse diffraction halo between 30 and 50◦ denoting a certain amount of amorphous phase. Thus, a good Rietveld fitting of MS (Rwp = 8.21%) is obtained summing up the contributions of two crystallite sizes (D1 = 68 ± 10 nm and D2 = 9 ± 3 nm) and an amorphous background. The relative contribution between the two crystallite sizes is 1:4, respectively. Fig. 3 shows SEM micrographs for MS(a) and AS(b) alloys. MS exhibits a wide particle size distribution comprised between 5 and 60 m. Particles are bounded by cleavage facets. On the contrary, for AS alloy, particles are rounded and possess a narrow particle size distribution (15 ± 5 m). The particular morphology of AS alloy is related to particle agglomeration at long-milling times, a phenomenon already observed in other intermetallic alloys such as Mg2 Ni [17]. To summarise, AS alloy exhibits a homogeneous microstructure with crystallite and particle size of 7 nm and
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Fig. 4. P–C–I absorption curves at 298 K of un-milled starting alloy, MS and AS alloys.
on AB5 -type alloy is a very inhomogeneous process at short milling times. 3.2. Hydrogenation properties
Fig. 2. Rietveld analysis of (a) MS and (b) AS alloys. Dashed lines in (a) stands for contributions of distinct crystallite-size domains D1 , D2 and amorphous background. Continuous lines in (a) and (b) stands for the global Rietveld fitting.
15 m, respectively. In contrast, the microstructure of MS is rather inhomogeneous both in particle and crystallite sizes. This is particularly severe for crystallite-size as a wide distribution of amorphous and several nano-crystalline domains coexist in MS alloys. This proves that mechanical milling
Fig. 4 shows the P–C–I absorption curves at 298 K of MS and AS alloys as compared with that of the starting alloy. AS alloy does not exhibit any plateau pressure but an increased hydrogen solubility in the ␣-solute branch. This effect has been already reported in nano-crystalline materials [18]. As concerns MS alloy, the total hydrogen capacity is higher than AS. This increase is assigned to large crystallite domains, which remain in the AB5 -type alloy after short milling. In addition, a further increment of the solute phase is observed in MS. The expansion of the H-solubility region at low hydrogenation pressures can be ascribed to additional trapping sites caused by amorphous domains in the inhomogeneous MS alloy. To get further information on the H-trapping states, thermal desorption experiments were conducted for the MS, AS
Fig. 3. SEM micrographs of (a) MS alloy and (b) AS alloy.
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faster for MS than for AS. For instance, the time for reaching a reacted fraction of 0.8 is only 200 s for MS, whereas it is as long as 4000 s for AS. Taking into account the microstructural properties of both alloys, such remarkable difference is thought to be due to the occurrence of the amorphous phase for MS alloy. Amorphous domains act as hydrogen traps for small hydrogen concentrations but, at higher concentrations, H-diffusion may drastically increase with the saturation of deep-energy traps. This effect has been clearly demonstrated by Kirchheim et al. in Pd–Si alloys [20]. In addition, SEM micrographs show that AS alloys exhibits a smaller surface-to-volume ratio than MS as a result of particle agglomeration at long milling time. This also contributes to faster kinetic for MS alloy as compared to AS alloy. Fig. 5. Thermal desorption spectra of un-milled starting alloy, MS and AS alloys. The spectra are normalised to the same alloy mass.
4. Conclusion
and starting alloys (Fig. 5). H-desorption rate from AS and MS alloys are much higher than that for the starting alloy, in agreement with their higher H-solubility for the ␣-phase. In fact, the starting alloy only exhibits a small peak at 450 K, which can be ascribed to H-trapped in defects [19]. For nano-crystalline AS alloy, one desorption peak is observed at 410 K which extends up to 625 K. For the non-homogeneous MS alloy, the same trend is observed but extra desorption states appear at T > 550 K. These states can be attributed to H-trapped in tetrahedral sites with a high rare-earth concentration [7]. Hydriding absorption kinetics of MS and AS alloys were measured at 298 K with a quasi-constant pressure of 0.9 MPa. Reacted fraction curves (defined as the hydrogen moles absorbed at any arbitrary time divided by the number of moles which can be absorbed at the infinite time) were evaluated and are shown in Fig. 6. H-kinetics are much
Hydrogenation properties of two AB5 -type alloys with distinct microstructural properties but comparable average crystallite sizes have been investigated. One of them, MS, was short-milled in a planetary mill. This treatment produces a non-homogeneous microstructure composed of a wide distribution of amorphous and nano-crystallite sizes. The other, AS, was annealed after long mechanical milling, resulting in a homogeneous nano-crystallite microstructure. As concerns H-thermodynamics, both MS and AS alloys exhibit no significant plateau pressure and a higher H-solubility in the solute-␣ phase as compared with un-milled AB5 -type alloy. Widening of the ␣-branch is bigger for MS than for AS due to H-trapping in amorphous sites for the former alloy, as corroborated by thermal desorption measurements. As concerns H-kinetics, hydrogen absorption is much faster in MS than in AS alloy, which can be mainly ascribed to enhanced hydrogen diffusivity through amorphous domains.
Acknowledgements The authors wish to thank F. Briaucourt for technical assistance and J.L. Pastol for SEM measurements.
References
Fig. 6. Time–evolution of the alloy reacted fraction for MS and AS alloys during hydrogen absorption at 0.9 MPa, 298 K.
[1] J.H.N. Van Vucht, F.A. Kuijpers, H.C.A.M. Bruning, Philips Res. Rep. 25 (1970) 133. [2] H.F. Bittner, C.C. Badcock, J. Electrochem. Soc. 130 (1983) 193c. [3] J.-M. Joubert, M. Latroche, A. Percheron-Guégan, F. BouréeVigneron, J. Alloys Compd. 275 (1998) 118. [4] C. Lartigue, A. Percheron-Guégan, J.C. Achard, F. Tasset, J. Less-Common Metals 75 (1980) 23. [5] F. Feng, M. Geng, D.O. Northwood, Int. J. Hydrogen Energy 26 (2001) 725.
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[6] H. Ogawa, M. Ikoma, H. Kawano, I. Matsumoto, J. Power Sources 12 (1988) 393. [7] F. Cuevas, M. Hirscher, Acta Mater. 51 (2003) 701. [8] H. Niu, D.O. Northwood, Int. J. Hydrogen Energy 27 (2002) 69. [9] Y. Li, Y.-T. Cheng, M.A. Habib, J. Alloys Compd. 209 (1994) 7. [10] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Compd. 253 (1997) 70. [11] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1. [12] G. Liang, J. Hout, R. Schulz, J. Alloys Compd. 320 (2001) 133. [13] S. Corré, M. Bououdina, N. Kuriyama, D. Fruchart, G. Y Adachi, J. Alloys Compd. 29 (1999) 166.
[14] R. Kirchheim, F. Sommer, G. Schluckebier, Acta Metall. 30 (1982) 1059. [15] J. Rodr`ıguez-Carvajal, Physica B 192 (1993) 55. [16] P. Thompson, J.J. Reilly, J.M. Hastings, J. Less-Common Metals 129 (1987) 105. [17] D. Cracco, A. Percheron-Guégan, J. Alloys Compd. 268 (1998) 248. [18] N. Mommer, M. Hirscher, F. Cuevas, H. Kronmüller, J. Alloys Compd. 266 (1998) 255. [19] E.H. Kisi, E. Wu, M. Kemali, J. Alloys Compd. 330 (2002) 202. [20] R. Kirchheim, T. Mütschele, W. Kieninger, H. Gleiter, R. Birringer, T.D. Koblé, Mater. Sci. Eng. 99 (1988) 457.