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dehydriding properties of mechanically milled Mg–30 wt% TiMn1.5 composite
Journal of Alloys and Compounds 354 (2003) 296–302
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Preparation and hydriding / dehydriding properties of mechanically milled Mg–30 wt% TiMn 1.5 composite Y.Q. Hu, H.F. Zhang*, A.M. Wang, B.Z. Ding, Z.Q. Hu Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China Received 10 September 2002; received in revised form 5 December 2002; accepted 5 December 2002
Abstract A Mg–30 wt% TiMn 1.5 hydrogen storage composite was successfully synthesized by mechanical milling of a mixture of magnesium powder and amorphous TiMn 1.5 powder. The absorption / desorption rates and storage capacity under different temperatures were evaluated. The composite possesses high hydrogen storage capacity and exhibits excellent absorption / desorption kinetic properties and is activated in situ due to reaction ball milling (RBM). The hydrogen capacity is over 2.7 wt% at 373 K, and one absorption / desorption hydrogen cycle can be finished in 20 min at 523–573 K. Interestingly, the desorption temperature begins at about 500 K, decreasing about 40 K in contrast to that of the milled pure MgH 2 . Mechanical milling produces fine powder with nanometer-scaled grains and introduces many imperfections into the Mg matrix, which contribute to the enhanced hydrogen absorption / desorption rates with high capacity catalyzed by TiMn 1.5 (amorphous). 2002 Elsevier Science B.V. All rights reserved. Keywords: Transition metal alloys; Mechanical alloying; Nanostructured materials; Gas–solid reactions
1. Introduction Magnesium and Mg-based hydrogen storage materials have attracted many investigations owing to their high hydrogen storage capacity, low price and the fact that it is friendly to the environment [1]. However, the high work temperature, slow reaction kinetics and hard activation processes limit the practical application of Mg-based hydride system. Usually, hydriding / dehydriding of magnesium requests temperature above 573 K, and one hydrogen absorption / desorption cycle needs tens of hours. Some authors found that hydrogen molecules cannot be dissociated efficiently, as well as oxide-membrane exists on the surfaces of magnesium particles and diffusion of hydrogen atoms is slow through interfaces between metal and oxide, which leads to poor absorption / desorption kinetics [2]. At present, three main methods have been used to improve the work temperature and the hydriding / dehydriding properties of Mg-based hydrogen storage materials [3]: (1) The second phases, which possess favorable absorption / desorption kinetics at room temperature, are added to Mgbased alloys to form composites. (2) Transition metal *Corresponding author. Tel.: 186-24-2397-1783. E-mail address: [email protected] (H.F. Zhang).
elements substitute partial for magnesium and form multielement alloys. (3) Surfaces of Mg-based alloys are treated by F-ions, acidification, or by adding organic solution. The development of Mg-based composites has made some progress, including Mg–Mg 2 Ni 0.75 Fe 0.25 [4], Mg–LaNi 5 [5], Mg–FeTi [6]. Zaluski et al. [2] pointed out that the dissociation energy for hydrogen molecules is reduced drastically on the surfaces of some composites, and hydrogen molecules could dissociate to atoms relatively more easily. Furthermore, the particles of composites may provide paths for diffusion of hydrogen atoms to the matrix of Mg. Several authors [2,7,8] confirmed that the absorption / desorption of hydrogen properties are enhanced greatly when grains of Mg reach nanometer size. Iwakura and co-workers [9,10] considered some nano-materials including a part of amorphous phase which possessed favorable absorption / desorption kinetics. Wang et al. [11] thought that the metallic glass could reduce the grain size of Mg powder due to its brittleness during ball milling, and nano-scaled grains of Mg could be prepared. A large amount of defects exist in metallic glass, where hydrogen atoms may stay, and to some extent blocking of hydrogen atoms to move is reduced, even eliminated, which leads to a higher diffusion rate of hydrogen atoms [11].
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. doi:10.1016 / S0925-8388(02)01363-4
Y.Q. Hu et al. / Journal of Alloys and Compounds 354 (2003) 296–302
Recently the reaction ball milling (RBM) method was successfully introduced to prepare hydrogen storage materials [12,13]. It combines the courses of sample preparation, activation and hydrogenation into one step. Pure magnesium is ductile, while particles of Mg can be transformed into brittle hydride and nano-scaled grains of Mg (or MgH 2 ) may be produced by RBM. Furthermore, activation in situ due to reaction ball milling (RBM) in Mg-based hydrogen storage materials can overcome some problems caused by hard activation for hydrogen storage materials. In this case, referring to TiMn 1.5 , which possesses excellent hydriding / dehydriding properties at room temperature [14], we produce amorphous TiMn 1.5 powder by ball milling, then add the amorphous powder to the Mg powder and mill them under hydrogen atmosphere. Finally the nano / amorphous hydrogen storage composite, of which the fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg particles, is produced.
2. Experimental
2.1. Sample preparation Mixture (atom ratio is 1:1.5) of Ti (99.9 mass%, 200 mesh) and Mn (99.98 mass%, 100–200 mesh) powder was sealed in a stainless steel vial (65350 mm [) with stainless steel balls under argon atmosphere. The ball-topowder ratio (BPR) was 16:1. TiMn 1.5 (amorphous) was produced after 16 h of milling. Then the amorphous powder (30 wt%) and Mg (99 mass%, 100 mesh) were put into the stainless steel vial with stainless steel balls (BPR is 30:1) and the vial was evacuated to 10 22 Torr. Hydrogen (99.999%, over 0.5 MPa) was filled into the vial. The mixture was analyzed at different milling time. Hydrogen gas was introduced in the vial every half an hour again to keep the hydrogen pressure at 0.5 MPa. The powder was prepared by ball milling in a commercial SPEX8000.
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3. Results and discussion
3.1. Influence of TiMn1.5 (amorphous) on Mg-based hydride during the ball milling The evolution of the XRD patterns as a function of milling time for Mg1TiMn 1.5 (amorphous) milled under hydrogen atmosphere is shown in Fig. 1. Tetragonal bMgH 2 begins to form after 1 h. Two hours later, the relative intensity of Mg-XRD patterns decrease and the peaks broaden drastically. The reduction of the relative intensity for Mg rests on two reasons: one is hydriding of Mg; the other is reduction of granularity of the Mg grains. After 6 h, XRD patterns of Mg almost disappear and nanometer-scaled b-MgH 2 becomes a principal part of the mixture. Interestingly, orthorhombic g-MgH 2 also appears apparently after 5 h. Schulz et al. [15] observed a similar phenomenon in the milling course of MgH 2 . Five hours later, the XRD patterns of b-MgH 2 and g-MgH 2 become very broadened, which is associated with the reduction of the granularity. Another possible reason is the concentrated micro-strain caused by high-energy ball milling. The patterns of Fe and TiFe also appear after 5 h milling, which indicates that iron is introduced in the mixture and some titanium is separated out from the amorphous TiMn 1.5 , then the phase of TiFe forms during the ball milling process. Some iron combines with Ti and forms TiFe, the other dissolves into amorphous TiMn 1.5 (or TiFe) and forms a solid solution under the mechanical driving force of the high-energy milling, so the patterns of Fe disappear after 6 h of milling. The calculated crystal size of b-MgH 2 milled for 6 h, using the Scherrer equation, is about 10 nm. According to Bortz et al. [16], formation of b-MgH 2 by gas–solid reaction needs high temperature (.600 K) and high hydrogen pressure (.5 MPa). Furthermore, the synthesis of g-MgH 2 structure from b-MgH 2 needs higher temperature (523–1173 K) and higher hydrogen pressure
2.2. Sample characterization The structural evolution during milling was detected by X-ray diffraction (XRD) analysis in a Rigaku X-ray diffractometer equipped with graphite monochromator and Cu Ka radiation. The starting desorption temperature was tested by DSC, which was carried out in a Perkin-Elmer DSC under a flowing argon atmosphere at a heating rate of 20 K min 21 . The hydriding / dehydriding properties of the composite were measured at different temperatures using a conventional volumetric method. Microstructure studies were made on a SEM (Jeol JSM6301F) equipped with an energy dispersive X-ray analysis system and a TEM (Philips EM420).
Fig. 1. XRD patterns for powder of Mg and TiMn 1.5 (amorphous) after milling for various times under hydrogen atmosphere.
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(2.5–8 GPa) [16]. In our case, the formation of b-MgH 2 and g-MgH 2 under so low a hydrogen pressure at room temperature is attributed to three points below: (1) Generally, the smaller the grain size of Mg, the bigger the surface area, the easier the formation of Mg-based hydride. Toughness of magnesium is good, so the particles of magnesium are hardly slenderized only by ball milling. When the brittle particles of TiMn 1.5 (amorphous) are added, the brittle particles promote the crushing of the particles of magnesium into pieces and the tough particles are shattered easily. When the mixture has been milling for 6 h, the granularity of magnesium is near 10 nm. (2) Hydrogen molecules can be decomposed to atoms efficiently on the surfaces of amorphous TiMn 1.5 at room temperature, and the structural defects and the mechanical deformation occurring during ball milling produce plenty of sites that hydrogen atoms may stay in, so to some extent the blocking of hydrogen molecules to move is reduced, even eliminated, and at the same time the diffusion of hydrogen atoms is faster than that of conventional Mgbased alloys. All features mentioned above may improve the kinetics of absorption / desorption of hydrogen. (3) During the high-energy ball milling, collisions between balls and powder, or balls and balls, or balls and wall of vial provide instant high temperature and high pressure conditions under which b-MgH 2 , even g-MgH 2 can come into being. Furthermore, fresh interfaces of magnesium particles have been exposed continuously under the hydrogen atmosphere. Hydrogen atoms reach the fresh interfaces quickly without long diffusion paths and finish the process of hydriding in a short time. While for pure magnesium, after milling under hydrogen atmosphere for 4 h, no hydride phase is detected, and even the grain size is almost unchanged. The catalytic action of TiMn 1.5 (amorphous) for the hydrogenation of Mg is confirmed obviously from the above comparison.
Fig. 2. Hydrogen absorption kinetics for Mg–30 wt% TiMn 1.5 (amorphous) at different temperatures (starting 2 MPa).
3.3. Hydriding /dehydriding properties After evacuation at 373 K for about 1 h, the system is heated up to 623 K. It is found that the composite has been dehydrided at once. When the hydrogen (2 MPa) is introduced firstly, the sample is hydrided rapidly. It is well known that the sample is activated in situ due to reaction ball milling (RBM). The first absorption hydrogen capacity reaches 2.6 wt% (623 K). After three absorption / desorption cycles, the sample is activated completely. The work temperature for Mg-based hydrogen storage composites usually is above 573 K and it needs many hours to finish one hydrogen absorption / desorption cycle. The shortcomings mentioned above can be overcome efficiently to some extent in our sample. Fig. 2 shows hydrogen absorption kinetics curves for the composite at different temperatures. At 573 K, the highest hydrogen capacity reaches 4.4 wt%, and it needs only 2 min in finishing 90 wt% of the final capacity. Fig. 3 indicates the temperature dependence of hydrogen capacity and t 80% for
3.2. Stability for TiMn1.5 (amorphous) during the ball milling and the absorption /desorption of hydrogen cycles According to the binary phase diagrams, no compound exists between Ti and Mg, as does between Mn and Mg. The XRD patterns (Figs. 1 and 6) in our case also indicate that no interaction occurs between Mg and TiMn 1.5 (amorphous) during the ball milling and the absorption / desorption of hydrogen cycles. A little of titanium precipitates from the amorphous and forms TiFe phase with iron. In the mass, the amorphous phase remains relatively stable and only a small amount of nano-crystals precipitate. Toshio et al. [14] pointed out that the hydrogen plateau pressure of TiMn 1.5 above 373 K is above 2 MPa. Below 2 MPa, the main hydrided phase in Mg–30 wt% TiMn 1.5 is magnesium powder. The XRD patterns for the milled and the hydrided mixture also confirm that there is no appreciable hydride formation except for the Mg-hydride.
Fig. 3. Temperature dependence of the hydrogen capacity and t 80% for Mg–30 wt% TiMn 1.5 (amorphous) under 2.0 MPa hydrogen pressure.
Y.Q. Hu et al. / Journal of Alloys and Compounds 354 (2003) 296–302
Fig. 4. Hydrogen desorption kinetics for Mg–30 wt% TiMn 1.5 at different temperatures (ending 0.005 MPa).
PH 2 52.0 MPa. Here, t 80% is defined, as the time required to reach 80 wt% of the final capacity. It is important for Mg-based hydrogen storage materials to work at lower temperature. It should be noticed that even at a moderate temperature of 373 K, the composite still possesses relatively high hydrogen storage capacity (2.7 wt%, 2 MPa). It can absorb in 70 wt% of the final capacity in 1 min and finishes the whole process completely in 10 min. The property of desorption also is improved greatly. Fig. 4 shows that it needs only 6 min in finishing the desorption process. Fig. 5 manifests DSC curve of Mg after milling for 6 h. Endothermic peak means desorption of hydrogen. It is seen that the temperature of desorption of hydrogen starts at about 500 K. Compared with other authors, it descends a lot, which proves indirectly that desorption of hydrogen of the composite is improved greatly due to addition of TiMn 1.5 (amorphous) [8,10,15]. The absorption / desorption kinetic properties of the sample do not decrease markedly after 20 cycles. The XRD patterns after absorption of hydrogen at 573 and 373 K are shown in Fig. 6. The pattern a is that of hydro-
Fig. 5. DSC curve of Mg–30 wt% TiMn 1.5 (amorphous) after milling for 6 h.
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Fig. 6. XRD patterns for Mg–30 wt% TiMn 1.5 (amorphous) before and after hydriding / dehydriding cycles. (a) Hydrogenated composite at 373 K after 20 cycles; (b) hydrogenated composite at 573 K after 10 cycles; (c) composite milled under hydrogen atmosphere for 6 h.
genated sample at 373 K after 20 cycles. The pattern b is that of hydrogenated sample at 573 K after 10 cycles. Compared to that of pattern c (milled mixture in hydrogen atmosphere for 6 h), the accumulation of strain has been released and the grain size of MgH 2 increases a little in patterns a and b. High patterns of magnesium exist in pattern a because about 50 wt% of Mg is not transformed into hydride due to only 2.7 wt% of capacity at 373 K. At 573 K, most magnesium is already hydrogenated, so the patterns intensity of Mg is weak. Because of the effect of thermal during the hydrogenated cycles, the iron dissolved in the amorphous phase (or TiFe) precipitates, so the patterns of Fe appear in patterns a and b. The grain size of Mg is about 30 nm after 20 cycles. To sum up, the composite exhibits satisfactory stability during the absorption / desorption of hydrogen cycles.
3.4. Microstructure characteristics In order to characterize the distribution of TiMn 1.5 (amorphous), TEM and SEM observations were performed for the milled composite particles. Fig. 7 shows three back scattering electron (BSE) images of the particles surfaces. EDS analysis verifies that the bright phase corresponds to TiMn 1.5 and the dark phase to Mg. The fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg particles. Such ideal microstructure of the composite can be rationally understood as a result of the continuous cold welding and fraction in the milling course. In Fig. 7a, Mg and TiMn 1.5 can still be distinguished. In Fig. 7b, Mg and TiMn 1.5 particles have been dispersed completely and the zone, where TiMn 1.5 exists separately, cannot be identified. From Fig. 7b and c, the samples before and after absorption / desorption do not change evidently. It proves that the composite remains relatively stable during the absorption / desorption cycles. The composite nature and the nanostructure of Mg and
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Fig. 7. BSE images of the composite and EDS spectra of the selected area. (a) Milling for 2 h; (b) milling for 6 h; (c) after 20 hydriding / dehydriding cycles; (d) EDS spectra of point a; (e) EDS spectra of point b.
TiMn 1.5 are further confirmed by TEM observation of the particles (shown in Fig. 8). In dark field, the white grains are TiMn 1.5 (amorphous), about 5–10 nm, distributed uniformly on the surface of the Mg matrix. In bright field,
the grain size of Mg is about 10 nm (the phase of MgH 2 is decomposed because of focalized electron beam [17], so there is only the crystal of magnesium), agreeing with the results of XRD.
Y.Q. Hu et al. / Journal of Alloys and Compounds 354 (2003) 296–302
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Fig. 8. TEM result of the composite Mg–30 wt% TiMn 1.5 (amorphous) particle. (a) Bright field image; (b) dark field image.
3.5. Analysis of hydrogenation mechanism
4. Conclusion
The main hydrided phase is magnesium powder in Mg–30 wt% TiMn 1.5 hydrogen storage composite, and only small amounts of hydrogen dissolves in TiMn 1.5 . TiMn 1.5 (amorphous) acts as catalyst mainly in the composite. As described above, the fine TiMn 1.5 (amorphous) particles are dispersed uniformly on the surfaces of Mg particles, thus playing a critical role in accelerating the initial hydrogenation. During the absorption / desorption cycles, TiMn 1.5 provides the essential catalytic activity for the dissociation of hydrogen molecules by degrading the dissociation energy and also acts as ‘paths’ for the diffusion of hydrogen atoms from the surfaces to the underlying Mg. Hydrogen atoms and Mg are combined into Mg-based hydride relatively easily, when TiMn 1.5 dissociates hydrogen molecules, because the energy level of forming hydride is decreased. So it becomes possible that Mg-based hydride is formed at lower temperature. A great amount of interphase for amorphous and nanometercrystal exists, which provides paths of lower activation energy, so the dissociated hydrogen atoms can pass through the surfaces of Mg particles and the layers of hydride quickly and reach the Mg matrix. This step is thought to be one of the most important steps that control the absorption / desorption of hydrogen kinetics of Mgbased hydrogen storage materials [2]. Additionally, as discussed above, nanometer-scaled grains of Mg (or MgH 2 ) are formed during the RBM process. The presence of much phase boundary will increase the active surface area of the composite, thus shortening greatly the diffusion distance of hydrogen atoms. The broad energy distribution of the available sites for hydrogen atoms in the disordered nanometer-grain boundary regions, associated with the enhanced diffusivity and solubility in nano-crystal, accelerate the process of hydrogenation reaction of Mg.
1. A nanostructured Mg–30 wt% TiMn 1.5 hydrogen storage composite was produced by high-energy reaction ball milling. Fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg (or MgH 2 ) particles and the grain size of MgH 2 is about 10 nm. 2. The Mg–30 wt% TiMn 1.5 hydrogen storage composite is activated in situ due to RBM. The sample exhibits the following overall hydrogenation properties: the highest hydrogen storage capacity (4.4 wt%) combined with excellent kinetics. Even at 373 K, the saturation capacity still reaches 2.7 wt%, and the process of absorption of hydrogen can be finished in 10 min. 3. TiMn 1.5 (amorphous) powder pulverizes Mg grains, dissociates hydrogen molecules and provides rapid passageways, which improve the hydriding / dehydriding kinetics properties of the Mg–30 wt% TiMn 1.5 hydrogen storage composite.
Acknowledgements This work was carried out under the financial support of National High Technical Research and Development Program of China (grant 2001AA331010) and National Key Basic Research and Development Program of China (grant G2000067201).
References [1] P. Wang, H.F. Zhang, B.Z. Ding, Z.Q. Hu, Acta Mater. 49 (2001) 921. ¨ [2] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Comp. 253 (1997) 70. [3] L.B. Wang, Y.J. Wang, H.T. Yuan, J. Mater. Sci. Technol. 17 (2001) 590.
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[4] H.T. Yuan, H.B. Yang, Z.X. Zhou, D.Y. Song, Y.S. Zhang, J. Alloys Comp. 260 (1997) 256. [5] G. Liang, S. Boily, J. Huot, A.V. Neste, R. Schulz, J. Alloys Comp. 268 (1998) 302. [6] P. Mandal, O.N. Srivastava, J. Alloys Comp. 205 (1994) 111. [7] M. Zhu, W.H. Zhu, Y. Gao, X.Z. Che, J.H. Ahn, Mater. Sci. Eng. A 286 (2000) 130. ¨ [8] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, J. Alloys Comp. 288 (1999) 217. [9] C. Iwakura, S. Nohara, S.G. Zhang, H. Inoue, J. Alloys Comp. 285 (1999) 246. [10] S. Orimo, H. Fujii, K. Ikeda, Acta Mater. 45 (1997) 331. [11] P. Wang, A.M. Wang, H.F. Zhang, J. Mater. Sci. Lett. 20 (8) (2001) 753.
[12] J.L. Bobet, B. Chevalier, M.Y. Song, B. Darriet, J. Etourneau, J. Alloys Comp. 336 (2002) 292. [13] P. Wang, A.M. Wang, H.F. Zhang, B.Z. Ding, Z.Q. Hu, J. Alloys Comp. 313 (2000) 218. [14] T. Yamashita, T. Gamo, Y. Moriwaki, M. Fukuda, J. Jpn. Inst. Metals 41 (1977) 148. [15] R. Schulz, J. Huot, G. Liang, S. Boily, G. Lalande, M.C. Denis, J.P. Dodelet, Mater. Sci. Eng. A 267 (1999) 240. ¨ [16] M. Bortz, B. Bertheville, G. Bottger, K. Yvon, J. Alloys Comp. 287 (1999) L4. ¨ [17] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, J. Alloys Comp. 289 (1999) 197.
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Preparation and hydriding / dehydriding properties of mechanically milled Mg–30 wt% TiMn 1.5 composite Y.Q. Hu, H.F. Zhang*, A.M. Wang, B.Z. Ding, Z.Q. Hu Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China Received 10 September 2002; received in revised form 5 December 2002; accepted 5 December 2002
Abstract A Mg–30 wt% TiMn 1.5 hydrogen storage composite was successfully synthesized by mechanical milling of a mixture of magnesium powder and amorphous TiMn 1.5 powder. The absorption / desorption rates and storage capacity under different temperatures were evaluated. The composite possesses high hydrogen storage capacity and exhibits excellent absorption / desorption kinetic properties and is activated in situ due to reaction ball milling (RBM). The hydrogen capacity is over 2.7 wt% at 373 K, and one absorption / desorption hydrogen cycle can be finished in 20 min at 523–573 K. Interestingly, the desorption temperature begins at about 500 K, decreasing about 40 K in contrast to that of the milled pure MgH 2 . Mechanical milling produces fine powder with nanometer-scaled grains and introduces many imperfections into the Mg matrix, which contribute to the enhanced hydrogen absorption / desorption rates with high capacity catalyzed by TiMn 1.5 (amorphous). 2002 Elsevier Science B.V. All rights reserved. Keywords: Transition metal alloys; Mechanical alloying; Nanostructured materials; Gas–solid reactions
1. Introduction Magnesium and Mg-based hydrogen storage materials have attracted many investigations owing to their high hydrogen storage capacity, low price and the fact that it is friendly to the environment [1]. However, the high work temperature, slow reaction kinetics and hard activation processes limit the practical application of Mg-based hydride system. Usually, hydriding / dehydriding of magnesium requests temperature above 573 K, and one hydrogen absorption / desorption cycle needs tens of hours. Some authors found that hydrogen molecules cannot be dissociated efficiently, as well as oxide-membrane exists on the surfaces of magnesium particles and diffusion of hydrogen atoms is slow through interfaces between metal and oxide, which leads to poor absorption / desorption kinetics [2]. At present, three main methods have been used to improve the work temperature and the hydriding / dehydriding properties of Mg-based hydrogen storage materials [3]: (1) The second phases, which possess favorable absorption / desorption kinetics at room temperature, are added to Mgbased alloys to form composites. (2) Transition metal *Corresponding author. Tel.: 186-24-2397-1783. E-mail address: [email protected] (H.F. Zhang).
elements substitute partial for magnesium and form multielement alloys. (3) Surfaces of Mg-based alloys are treated by F-ions, acidification, or by adding organic solution. The development of Mg-based composites has made some progress, including Mg–Mg 2 Ni 0.75 Fe 0.25 [4], Mg–LaNi 5 [5], Mg–FeTi [6]. Zaluski et al. [2] pointed out that the dissociation energy for hydrogen molecules is reduced drastically on the surfaces of some composites, and hydrogen molecules could dissociate to atoms relatively more easily. Furthermore, the particles of composites may provide paths for diffusion of hydrogen atoms to the matrix of Mg. Several authors [2,7,8] confirmed that the absorption / desorption of hydrogen properties are enhanced greatly when grains of Mg reach nanometer size. Iwakura and co-workers [9,10] considered some nano-materials including a part of amorphous phase which possessed favorable absorption / desorption kinetics. Wang et al. [11] thought that the metallic glass could reduce the grain size of Mg powder due to its brittleness during ball milling, and nano-scaled grains of Mg could be prepared. A large amount of defects exist in metallic glass, where hydrogen atoms may stay, and to some extent blocking of hydrogen atoms to move is reduced, even eliminated, which leads to a higher diffusion rate of hydrogen atoms [11].
0925-8388 / 02 / $ – see front matter 2002 Elsevier Science B.V. All rights reserved. doi:10.1016 / S0925-8388(02)01363-4
Y.Q. Hu et al. / Journal of Alloys and Compounds 354 (2003) 296–302
Recently the reaction ball milling (RBM) method was successfully introduced to prepare hydrogen storage materials [12,13]. It combines the courses of sample preparation, activation and hydrogenation into one step. Pure magnesium is ductile, while particles of Mg can be transformed into brittle hydride and nano-scaled grains of Mg (or MgH 2 ) may be produced by RBM. Furthermore, activation in situ due to reaction ball milling (RBM) in Mg-based hydrogen storage materials can overcome some problems caused by hard activation for hydrogen storage materials. In this case, referring to TiMn 1.5 , which possesses excellent hydriding / dehydriding properties at room temperature [14], we produce amorphous TiMn 1.5 powder by ball milling, then add the amorphous powder to the Mg powder and mill them under hydrogen atmosphere. Finally the nano / amorphous hydrogen storage composite, of which the fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg particles, is produced.
2. Experimental
2.1. Sample preparation Mixture (atom ratio is 1:1.5) of Ti (99.9 mass%, 200 mesh) and Mn (99.98 mass%, 100–200 mesh) powder was sealed in a stainless steel vial (65350 mm [) with stainless steel balls under argon atmosphere. The ball-topowder ratio (BPR) was 16:1. TiMn 1.5 (amorphous) was produced after 16 h of milling. Then the amorphous powder (30 wt%) and Mg (99 mass%, 100 mesh) were put into the stainless steel vial with stainless steel balls (BPR is 30:1) and the vial was evacuated to 10 22 Torr. Hydrogen (99.999%, over 0.5 MPa) was filled into the vial. The mixture was analyzed at different milling time. Hydrogen gas was introduced in the vial every half an hour again to keep the hydrogen pressure at 0.5 MPa. The powder was prepared by ball milling in a commercial SPEX8000.
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3. Results and discussion
3.1. Influence of TiMn1.5 (amorphous) on Mg-based hydride during the ball milling The evolution of the XRD patterns as a function of milling time for Mg1TiMn 1.5 (amorphous) milled under hydrogen atmosphere is shown in Fig. 1. Tetragonal bMgH 2 begins to form after 1 h. Two hours later, the relative intensity of Mg-XRD patterns decrease and the peaks broaden drastically. The reduction of the relative intensity for Mg rests on two reasons: one is hydriding of Mg; the other is reduction of granularity of the Mg grains. After 6 h, XRD patterns of Mg almost disappear and nanometer-scaled b-MgH 2 becomes a principal part of the mixture. Interestingly, orthorhombic g-MgH 2 also appears apparently after 5 h. Schulz et al. [15] observed a similar phenomenon in the milling course of MgH 2 . Five hours later, the XRD patterns of b-MgH 2 and g-MgH 2 become very broadened, which is associated with the reduction of the granularity. Another possible reason is the concentrated micro-strain caused by high-energy ball milling. The patterns of Fe and TiFe also appear after 5 h milling, which indicates that iron is introduced in the mixture and some titanium is separated out from the amorphous TiMn 1.5 , then the phase of TiFe forms during the ball milling process. Some iron combines with Ti and forms TiFe, the other dissolves into amorphous TiMn 1.5 (or TiFe) and forms a solid solution under the mechanical driving force of the high-energy milling, so the patterns of Fe disappear after 6 h of milling. The calculated crystal size of b-MgH 2 milled for 6 h, using the Scherrer equation, is about 10 nm. According to Bortz et al. [16], formation of b-MgH 2 by gas–solid reaction needs high temperature (.600 K) and high hydrogen pressure (.5 MPa). Furthermore, the synthesis of g-MgH 2 structure from b-MgH 2 needs higher temperature (523–1173 K) and higher hydrogen pressure
2.2. Sample characterization The structural evolution during milling was detected by X-ray diffraction (XRD) analysis in a Rigaku X-ray diffractometer equipped with graphite monochromator and Cu Ka radiation. The starting desorption temperature was tested by DSC, which was carried out in a Perkin-Elmer DSC under a flowing argon atmosphere at a heating rate of 20 K min 21 . The hydriding / dehydriding properties of the composite were measured at different temperatures using a conventional volumetric method. Microstructure studies were made on a SEM (Jeol JSM6301F) equipped with an energy dispersive X-ray analysis system and a TEM (Philips EM420).
Fig. 1. XRD patterns for powder of Mg and TiMn 1.5 (amorphous) after milling for various times under hydrogen atmosphere.
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(2.5–8 GPa) [16]. In our case, the formation of b-MgH 2 and g-MgH 2 under so low a hydrogen pressure at room temperature is attributed to three points below: (1) Generally, the smaller the grain size of Mg, the bigger the surface area, the easier the formation of Mg-based hydride. Toughness of magnesium is good, so the particles of magnesium are hardly slenderized only by ball milling. When the brittle particles of TiMn 1.5 (amorphous) are added, the brittle particles promote the crushing of the particles of magnesium into pieces and the tough particles are shattered easily. When the mixture has been milling for 6 h, the granularity of magnesium is near 10 nm. (2) Hydrogen molecules can be decomposed to atoms efficiently on the surfaces of amorphous TiMn 1.5 at room temperature, and the structural defects and the mechanical deformation occurring during ball milling produce plenty of sites that hydrogen atoms may stay in, so to some extent the blocking of hydrogen molecules to move is reduced, even eliminated, and at the same time the diffusion of hydrogen atoms is faster than that of conventional Mgbased alloys. All features mentioned above may improve the kinetics of absorption / desorption of hydrogen. (3) During the high-energy ball milling, collisions between balls and powder, or balls and balls, or balls and wall of vial provide instant high temperature and high pressure conditions under which b-MgH 2 , even g-MgH 2 can come into being. Furthermore, fresh interfaces of magnesium particles have been exposed continuously under the hydrogen atmosphere. Hydrogen atoms reach the fresh interfaces quickly without long diffusion paths and finish the process of hydriding in a short time. While for pure magnesium, after milling under hydrogen atmosphere for 4 h, no hydride phase is detected, and even the grain size is almost unchanged. The catalytic action of TiMn 1.5 (amorphous) for the hydrogenation of Mg is confirmed obviously from the above comparison.
Fig. 2. Hydrogen absorption kinetics for Mg–30 wt% TiMn 1.5 (amorphous) at different temperatures (starting 2 MPa).
3.3. Hydriding /dehydriding properties After evacuation at 373 K for about 1 h, the system is heated up to 623 K. It is found that the composite has been dehydrided at once. When the hydrogen (2 MPa) is introduced firstly, the sample is hydrided rapidly. It is well known that the sample is activated in situ due to reaction ball milling (RBM). The first absorption hydrogen capacity reaches 2.6 wt% (623 K). After three absorption / desorption cycles, the sample is activated completely. The work temperature for Mg-based hydrogen storage composites usually is above 573 K and it needs many hours to finish one hydrogen absorption / desorption cycle. The shortcomings mentioned above can be overcome efficiently to some extent in our sample. Fig. 2 shows hydrogen absorption kinetics curves for the composite at different temperatures. At 573 K, the highest hydrogen capacity reaches 4.4 wt%, and it needs only 2 min in finishing 90 wt% of the final capacity. Fig. 3 indicates the temperature dependence of hydrogen capacity and t 80% for
3.2. Stability for TiMn1.5 (amorphous) during the ball milling and the absorption /desorption of hydrogen cycles According to the binary phase diagrams, no compound exists between Ti and Mg, as does between Mn and Mg. The XRD patterns (Figs. 1 and 6) in our case also indicate that no interaction occurs between Mg and TiMn 1.5 (amorphous) during the ball milling and the absorption / desorption of hydrogen cycles. A little of titanium precipitates from the amorphous and forms TiFe phase with iron. In the mass, the amorphous phase remains relatively stable and only a small amount of nano-crystals precipitate. Toshio et al. [14] pointed out that the hydrogen plateau pressure of TiMn 1.5 above 373 K is above 2 MPa. Below 2 MPa, the main hydrided phase in Mg–30 wt% TiMn 1.5 is magnesium powder. The XRD patterns for the milled and the hydrided mixture also confirm that there is no appreciable hydride formation except for the Mg-hydride.
Fig. 3. Temperature dependence of the hydrogen capacity and t 80% for Mg–30 wt% TiMn 1.5 (amorphous) under 2.0 MPa hydrogen pressure.
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Fig. 4. Hydrogen desorption kinetics for Mg–30 wt% TiMn 1.5 at different temperatures (ending 0.005 MPa).
PH 2 52.0 MPa. Here, t 80% is defined, as the time required to reach 80 wt% of the final capacity. It is important for Mg-based hydrogen storage materials to work at lower temperature. It should be noticed that even at a moderate temperature of 373 K, the composite still possesses relatively high hydrogen storage capacity (2.7 wt%, 2 MPa). It can absorb in 70 wt% of the final capacity in 1 min and finishes the whole process completely in 10 min. The property of desorption also is improved greatly. Fig. 4 shows that it needs only 6 min in finishing the desorption process. Fig. 5 manifests DSC curve of Mg after milling for 6 h. Endothermic peak means desorption of hydrogen. It is seen that the temperature of desorption of hydrogen starts at about 500 K. Compared with other authors, it descends a lot, which proves indirectly that desorption of hydrogen of the composite is improved greatly due to addition of TiMn 1.5 (amorphous) [8,10,15]. The absorption / desorption kinetic properties of the sample do not decrease markedly after 20 cycles. The XRD patterns after absorption of hydrogen at 573 and 373 K are shown in Fig. 6. The pattern a is that of hydro-
Fig. 5. DSC curve of Mg–30 wt% TiMn 1.5 (amorphous) after milling for 6 h.
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Fig. 6. XRD patterns for Mg–30 wt% TiMn 1.5 (amorphous) before and after hydriding / dehydriding cycles. (a) Hydrogenated composite at 373 K after 20 cycles; (b) hydrogenated composite at 573 K after 10 cycles; (c) composite milled under hydrogen atmosphere for 6 h.
genated sample at 373 K after 20 cycles. The pattern b is that of hydrogenated sample at 573 K after 10 cycles. Compared to that of pattern c (milled mixture in hydrogen atmosphere for 6 h), the accumulation of strain has been released and the grain size of MgH 2 increases a little in patterns a and b. High patterns of magnesium exist in pattern a because about 50 wt% of Mg is not transformed into hydride due to only 2.7 wt% of capacity at 373 K. At 573 K, most magnesium is already hydrogenated, so the patterns intensity of Mg is weak. Because of the effect of thermal during the hydrogenated cycles, the iron dissolved in the amorphous phase (or TiFe) precipitates, so the patterns of Fe appear in patterns a and b. The grain size of Mg is about 30 nm after 20 cycles. To sum up, the composite exhibits satisfactory stability during the absorption / desorption of hydrogen cycles.
3.4. Microstructure characteristics In order to characterize the distribution of TiMn 1.5 (amorphous), TEM and SEM observations were performed for the milled composite particles. Fig. 7 shows three back scattering electron (BSE) images of the particles surfaces. EDS analysis verifies that the bright phase corresponds to TiMn 1.5 and the dark phase to Mg. The fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg particles. Such ideal microstructure of the composite can be rationally understood as a result of the continuous cold welding and fraction in the milling course. In Fig. 7a, Mg and TiMn 1.5 can still be distinguished. In Fig. 7b, Mg and TiMn 1.5 particles have been dispersed completely and the zone, where TiMn 1.5 exists separately, cannot be identified. From Fig. 7b and c, the samples before and after absorption / desorption do not change evidently. It proves that the composite remains relatively stable during the absorption / desorption cycles. The composite nature and the nanostructure of Mg and
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Fig. 7. BSE images of the composite and EDS spectra of the selected area. (a) Milling for 2 h; (b) milling for 6 h; (c) after 20 hydriding / dehydriding cycles; (d) EDS spectra of point a; (e) EDS spectra of point b.
TiMn 1.5 are further confirmed by TEM observation of the particles (shown in Fig. 8). In dark field, the white grains are TiMn 1.5 (amorphous), about 5–10 nm, distributed uniformly on the surface of the Mg matrix. In bright field,
the grain size of Mg is about 10 nm (the phase of MgH 2 is decomposed because of focalized electron beam [17], so there is only the crystal of magnesium), agreeing with the results of XRD.
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Fig. 8. TEM result of the composite Mg–30 wt% TiMn 1.5 (amorphous) particle. (a) Bright field image; (b) dark field image.
3.5. Analysis of hydrogenation mechanism
4. Conclusion
The main hydrided phase is magnesium powder in Mg–30 wt% TiMn 1.5 hydrogen storage composite, and only small amounts of hydrogen dissolves in TiMn 1.5 . TiMn 1.5 (amorphous) acts as catalyst mainly in the composite. As described above, the fine TiMn 1.5 (amorphous) particles are dispersed uniformly on the surfaces of Mg particles, thus playing a critical role in accelerating the initial hydrogenation. During the absorption / desorption cycles, TiMn 1.5 provides the essential catalytic activity for the dissociation of hydrogen molecules by degrading the dissociation energy and also acts as ‘paths’ for the diffusion of hydrogen atoms from the surfaces to the underlying Mg. Hydrogen atoms and Mg are combined into Mg-based hydride relatively easily, when TiMn 1.5 dissociates hydrogen molecules, because the energy level of forming hydride is decreased. So it becomes possible that Mg-based hydride is formed at lower temperature. A great amount of interphase for amorphous and nanometercrystal exists, which provides paths of lower activation energy, so the dissociated hydrogen atoms can pass through the surfaces of Mg particles and the layers of hydride quickly and reach the Mg matrix. This step is thought to be one of the most important steps that control the absorption / desorption of hydrogen kinetics of Mgbased hydrogen storage materials [2]. Additionally, as discussed above, nanometer-scaled grains of Mg (or MgH 2 ) are formed during the RBM process. The presence of much phase boundary will increase the active surface area of the composite, thus shortening greatly the diffusion distance of hydrogen atoms. The broad energy distribution of the available sites for hydrogen atoms in the disordered nanometer-grain boundary regions, associated with the enhanced diffusivity and solubility in nano-crystal, accelerate the process of hydrogenation reaction of Mg.
1. A nanostructured Mg–30 wt% TiMn 1.5 hydrogen storage composite was produced by high-energy reaction ball milling. Fine TiMn 1.5 particles are dispersed uniformly on the surfaces of Mg (or MgH 2 ) particles and the grain size of MgH 2 is about 10 nm. 2. The Mg–30 wt% TiMn 1.5 hydrogen storage composite is activated in situ due to RBM. The sample exhibits the following overall hydrogenation properties: the highest hydrogen storage capacity (4.4 wt%) combined with excellent kinetics. Even at 373 K, the saturation capacity still reaches 2.7 wt%, and the process of absorption of hydrogen can be finished in 10 min. 3. TiMn 1.5 (amorphous) powder pulverizes Mg grains, dissociates hydrogen molecules and provides rapid passageways, which improve the hydriding / dehydriding kinetics properties of the Mg–30 wt% TiMn 1.5 hydrogen storage composite.
Acknowledgements This work was carried out under the financial support of National High Technical Research and Development Program of China (grant 2001AA331010) and National Key Basic Research and Development Program of China (grant G2000067201).
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