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Hydrogen storage properties of 2Mg–Fe after the combined processes of hot extrusion and cold rolling
Journal of Alloys and Compounds xxx (2013) xxx–xxx
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Hydrogen storage properties of 2Mg–Fe after the combined processes of hot extrusion and cold rolling G.F. Lima ⇑, M.R.M. Triques, C.S. Kiminami, W.J. Botta, A.M. Jorge Jr. Department of Materials Engineering, Federal University of São Carlos, São Carlos, SP, Brazil
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Hydrogen storage properties 2 Mg–Fe Hot-extrusion Cold-rolling
a b s t r a c t Bulk consolidated samples of 2 Mg–Fe were produced by hot extrusion and by hot extrusion followed by cold rolling. The starting 2 Mg–Fe mixture was obtained by high-energy ball milling. The cold-rolled samples, which presented a favorable (0 0 2) texture in Mg, absorbed 5.23 wt.% H2, almost the full capacity of Mg2FeH6, without any incubation time. This level of hydrogen absorption capacity was higher than the ones observed for the milled powders or the extruded samples. The hydrogen desorption temperatures were very close for all types of samples. The good results observed in the hot-extruded plus cold-rolled sample were explained by the smaller grain sizes, along with a favorable texture and the catalytic effect of residual Fe. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium alloys are adequate materials for hydrogen storage in the solid state; magnesium hydride, MgH2, has the highest storage capacity (7.6 wt.% H) among the reversible hydrides. The main drawbacks to its extensive use are the poor kinetics of the Mg/ MgH2 system and the high H-sorption temperatures. Another unfavorable aspect of MgH2 when in the powder form is the poor air resistance, which results in the formation of a highly stable oxide layer, lowering absorption rates. In such cases, activation cycles at high temperatures and hydrogen pressures are necessary to reach full capacity [1,2]. To minimize the effects of these unfavorable aspects, investigations have been focused on the reduction of the crystallite size to destabilize MgH2 and to shorten diffusion paths, on the addition of ‘‘catalytic’’ elements or compounds to favor hydrogen dissociation and recombination, and on the improvement of surface properties to protect against surface contamination and alloying MgH2 to weak Mg–H bonds [2–8]. More recently, severe plastic deformation (SPD) methods have been employed to obtain Mg-hydrides with refined grain sizes, controlled textures, and attractive hydrogen storage properties [9–18]. The SPD-processed materials resulted in easier activation, rapidly reaching full capacity without any incubation time. In addition, these materials have a good air resistance due to the lower surface-volume ratio in comparison with powders [10]. The desorption kinetics of these materials were also enhanced by the presence of catalysts. In fact, good hydrogen storage properties have been observed in hot-extruded 2 Mg–Fe samples [19–21],
⇑ Corresponding author. Tel.: +55 16 33518531. E-mail address: [email protected] (G.F. Lima).
which, because of better air resistance, also presented a higher hydrogen capacity in the first cycle involving powder samples [22]. An important aspect of the presence of Fe is the possibility of maintaining the nanograin size after many hydrogenation cycles; the pinning effect of Fe at the a-Mg grain boundary is predominant even when the sample is maintained at high temperatures for long periods [19]. Fe participates in the formation of the complex hydride Mg2FeH6 [23]; in excess, it will act as a catalyst favoring the dissociation of the H2 molecule and acting in the desorption process, reducing both the decomposition temperature and the energy barrier for decomposition [24,25]. In the present work, we used two combined processes, hot extrusion and repetitive cold rolling, with a mixture of 2 Mg–Fe powders to produce textured bulk samples with nanoscale grain sizes. 2. Experimental Pure elements Mg (+20–100 mesh, 99.98%, Alfa Aesar) and Fe (20 mesh, 99.998%, Alfa Aesar) were mixed to produce stoichiometric 2 Mg–Fe using high-energy ball milling (HEBM) in a Fritsch P7 planetary ball mill. The ball-to-powder ratio was 40:1, and milling was carried out under an argon atmosphere for 4 h. Milled samples (1 g) were cold-pressed into pre-forms and then hot-extruded at 300 °C using a ram speed of 1 mm/min and an extrusion ratio of 3/1. The extruded samples, inserted between two AISI 304 stainless steel plates, were processed by cold rolling in a duo-reversible conventional rolling mill (FENN). The samples were submitted to several passes (20–30), with a 50% reduction in each pass. Both the hot-extrusion and cold-rolling processes were performed in air. The final foil thickness was approximately 0.150 mm. The hydrogen absorption properties were measured in a Sieverts-type apparatus at 673 K and 15 bar of H2(g). The hydrogen desorption behaviors were investigated by simultaneous differential scanning calorimetry–thermogravimetry–STA (TG-DSC) coupled with quadrupole mass spectrometry (QMS) in a Netzsch STA 449C + QMS 403C apparatus, with a constant heating rate of 10 °C/min under purified argon in an overflow regime.
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The structural characterization of the samples was carried out by X-ray diffraction (XRD) using monochromatic Cu Ka radiation in a Rigaku diffractometer equipped with a C-monochromator and by transmission electron microscopy (TEM, FEI-CM-120). The XRD profiles were treated by Rietveld refinement [26] using the MAUD software [27].
Table 1 Average crystallite sizes (±5 nm) of Mg and Fe estimated by Rietveld refinement using the MAUD software.
Mg Fe
3. Results and discussion Fig. 1 shows XRD patterns of the 2 Mg–Fe powder mixtures in comparison with the bulk obtained by hot extrusion and by hot extrusion followed by cold rolling. The XRD patterns revealed only the presence of hcp-Mg and bcc-Fe; there is no indication of the presence of oxides in any of the patterns. The XRD patterns corresponding to the powders and to the samples that were only extruded are quite similar, with no indication of a strongly preferred orientation. However, in the rolled sample, a strong (0 0 2) texture is observed for Mg. As first stated by Léon et al. [28] and after by Dufour and Huot [10,11] and by Singh et al. [29], the (0 0 2) orientation is a type of texture that favors hydrogen absorption during activation. Singh et al. [29] explained this behavior in an investigation of the structural and hydrogen storage properties in nanostructured thin films of Mg deposited on Si (0 0 1) substrates, X-ray diffraction showed that the conversion of Mg to MgH2 follows a martensitic-like orientation relationship with Mg 0//MgH2. (0 0 2)//MgH2 (1 1 0) and Mg ½1 2 Table 1 shows the average crystallite sizes of the powder and the extruded samples determined from Rietveld refinement. Because the Rietveld method is not adequate for samples presenting texture, the crystallite sizes of the cold-rolled samples were determined by TEM (Fig. 2). After 4 h of ball milling, the average crystallite sizes were in the nanometric scale, very close to that for the powder milled for 12 h, as already reported in a previous work [19], indicating that only 4 h is enough to produce nanostructured alloys. As expected, the extruded samples presented larger crystallite sizes than the ball-milled ones. This difference can be attributed to the annealing effect during hot extrusion, although the crystallite sizes were still kept in the nanoscale range because of the pinning effect of Fe particles at the a-Mg grain boundaries, as previously observed [19]. Fig. 2 shows bright-field (BF) and dark-field (DF) images and the corresponding selected-area electron diffraction patterns (SAEDPs) of the cold-rolled sample. A refined nanostructure is observed with most crystallites in the size range of 10 nm, although it is possible to observe a few grains that grow to approximately 60 nm. These images also show that the microstructure is homogeneous, with
Fig. 1. XRD patterns of 2 Mg–Fe mixtures in powder form obtained by milling, in bulk form as extruded, and in foil form as-extruded/rolled.
v2
2 Mg–Fe – milled (nm ± 5 nm)
2 Mg–Fe – extruded (nm ± 5 nm)
37 39 1.03
64 56 1.01
Fig. 2. Images obtained by TEM of extruded, cold-rolled 2 Mg–Fe: (a) bright-field and (b) dark-field images.
Fe being homogenously dispersed in the Mg matrix. The SAEDP confirms the presence of nanosized grains with a strong (0 0 2) texture, in agreement with the XRD pattern in Fig. 1. Fig. 3 shows the first hydrogen absorption curves for the 2 Mg– Fe samples processed by milling, extrusion, and extrusion followed by cold rolling. Absorption was measured at 673 K under a hydrogen pressure of 15 bar. The milled and cold-rolled samples presented faster kinetics than the extruded sample, did not have an incubation time, and reaching maximum capacity (saturation) after approximately 900 min (15 h). The time to reach maximum capacity is shorter than the times previously observed in coldrolled Mg–Pd – 1250 min [10] and 1700 min [11] – most likely because the samples in the present work had much smaller initial grain sizes. The 2 Mg–Fe composition chosen in this work was intended to form bulk textured samples of Mg2FeH6. The formation of this complex hydride was confirmed by the XRD pattern of the coldrolled sample after hydrogen absorption, as shown in Fig. 4. In addition to the Mg2FeH6 phase, we can observe peaks corresponding to residual Fe. The absorption curves of Fig. 3 also corroborates the efficient conversion of the initial 2 Mg–Fe mixtures to this
Please cite this article in press as: G.F. Lima et al., J. Alloys Comp. (2013), http://dx.doi.org/10.1016/j.jallcom.2013.01.115
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Fig. 3. Hydrogen absorption kinetics at 673 K (400 °C) and under a hydrogen pressure of 1500 kPa (15 bar) for 2 Mg–Fe milled, extruded and extruded/rolled samples.
complex phase; the cold-rolled samples could absorb 5.2% H, corresponding to 95% of the theoretical capacity. This value is significantly higher than the obtained in our previous work for the asextruded materials [19], which reached a maximum capacity of only 1.2% H. The good hydrogen storage capacity reached in the cold-rolled sample can then be attributed to the following factors: a favorable texture along the (0 0 2) direction in Mg, the additional deformation due to the cold-rolling process, and the catalytic action of residual Fe. In contrast, as shown in Fig. 3, the extruded sample presented an incubation time and very slow absorption kinetics. Consequently, a low hydrogen storage capacity was reached during the time of the experiment. A comparison of the two bulk nanostructured samples suggest that the texture along the (0 0 2) direction and the additional deformation caused by cold rolling are the most important factors contributing to the improved H-absorption kinetics and the capacity of the cold-rolled samples. The hydrogen desorption kinetics was estimated by a TG analysis, as shown in Fig. 5. The best desorption kinetics was observed for the powder sample, the one with the largest surface area. However, in agreement with Fig. 3, the cold-rolled sample presented the highest capacity, followed by the powder and the extruded samples. In contrast to our results, cold-rolled samples presented slower sorption kinetics and reduced capacities than corresponding samples obtained by milling [30]; however, in this case, there was an important difference in the crystallite sizes, which were smaller in the milled samples. The mass loss values are similar to
3
Fig. 5. Thermogravimetry measurements of 2 Mg–Fe milled, extruded, and extruded/rolled samples.
Fig. 6. DSC measurements of 2 Mg–Fe milled, extruded, and extruded/rolled samples.
but smaller than that found in Fig. 3. Most likely, these dissimilarities are due to differences between equipment and experimental analysis methodology. Additionally, the sample preparation was made in air for the DSC/TG analysis, which may result in a small level of oxidation that could not be detected by QMS analysis. Fig. 6 shows the DSC curves for a commercial MgH2 powder in comparison with the 2 Mg–Fe samples; the hydrogen desorption temperatures were lower for all the 2 Mg-Fe in all processing conditions. As expected, the lowest hydrogen desorption temperature was observed for the milled sample – the one with the highest surface area. Fig. 7 shows TEM images of hydrogenated samples. The nanograin sizes were maintained after absorption, and even the texture was not totally lost, as suggested by the electron diffraction pattern. Additional work is underway to investigate the texture stability in the cold-rolled sample. 4. Conclusions
Fig. 4. XRD patterns of hydrogenated 2 Mg–Fe as-extruded/rolled.
2 Mg–Fe mixtures processed by hot-extrusion combined with additional cold rolling produced, after hydrogenation, bulk nanostructured samples of the complex hydride Mg2FeH6 containing residual Fe.
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The better results observed for the combined processed sample may be explained by the smaller grain sizes, along with a favorable (0 0 2) texture of Mg and the catalyst effect of residual Fe. Acknowledgements The authors acknowledge the financial support of the Brazilian agencies FAPESP, CAPES, and CNPq. References
Fig. 7. Images obtained by TEM (bright field, 300 kX) of hydrogenated 2 Mg–Fe samples: (a) milled, (b) extruded and (c) extruded/rolled.
[1] B. Sakintuna, F. Lamari-Darkrim, M. Hirscher, Int. J. Hydrogen Energy 32 (2007) 1121–1140. [2] I.P. Jain, C. Lal, A. Jain, Int. J. Hydrogen Energy 35 (2010) 5133–5144. [3] R. Schulz, J. Huot, G. Liang, S. Boily, G. Lalande, M.C. Denis, J.P. Dodelet, Mater. Sci. Eng. A 267 (1999) 240–245. [4] L. Zaluski, A. Zaluska, J.O. Strom-Olsen, J. Alloys Comp. 253–254 (1997) 70–79. [5] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, J. Alloys Comp. 288 (1999) 217–225. [6] J.F.R. Castro, A.R. Yavari, A. Lemoulec, T.T. Ishikawa, W.J. Botta, J. Alloys Comp. 389 (2005) 270–274. [7] A.R. Yavari, A. Lemoulec, J.F.R. Castro, S. Deledda, O. Friedrichs, W.J. Botta, G. Vaughan, T. Klassen, A. Fernandez, A. Kvick, Scr. Mater. 52 (2005) 719–724. [8] B. Bogdanovic, A. Reiser, K. Schlichte, B. Spliethoff, B. Tesche, J. Alloys Comp. 345 (2002) 77–89. [9] T.T. Ueda, M. Tsukahara, Y. Kamiya, S. Kikuchi, J. Alloys Comp. 386 (2005) 253– 257. [10] J. Dufour, J. Huot, J. Alloys Comp. 439 (2007) L5–L7. [11] J. Dufour, J. Huot, J. Alloys Comp. 446–447 (2007) 147–151. [12] N. Takeichi, K. Tanaka, H. Tanaka, T.T. Ueda, Y. Kamiya, M. Tsukahara, H. Miyamura, S. Kikuchi, J. Alloys Comp. 446–447 (2007) 543–548. [13] V.M. Skripnyuk, E. Rabkin, Y. Estrin, R. Lapovok, Acta Mater. 52 (2004) 405– 414. [14] V. Skripnyuk, E. Buchman, E. Rabkin, Y. Estrin, M. Popov, S. Jorgensen, J. Alloys Comp. 436 (2007) 99–106. [15] S. Loken, J.K. Solberg, J.P. Maehlen, R.V. Denys, M.V. Lototsky, B.P. Tarasov, V.A. Yartys, J. Alloys Comp. 446–447 (2007) 114–120. [16] D.R. Leiva, D. Fruchart, M. Bacia, G. Girard, N. Skriabina, A.C.S. Villela, S. Miraglia, D.S. Santos, W.J. Botta, Int. J. Mater. Res. 100 (2009) 1–8. [17] Y. Kusadome, K. Ikeda, Y. Nakamori, S. Orimo, Z. Horita, Scr. Mater. 57 (2007) 751–753. [18] G.F. Lima, A.M. Jorge Jr., D.R. Leiva, C.S. Kiminami, C. Bolfarini, W.J. Botta, J. Phys.: Conf. Ser. 144 (2009) 012015. [19] G.F. Lima, M.M. Peres, S. Garroni, M.D. Baró, S. Surinyach, C.S. Kiminami, T.T. Ishikawa, W.J. Botta, A.M. Jorge Jr., J. Alloys Comp. 504S (2010) S299–S301. [20] G.F. Lima, S. Garroni, M.D. Baró, S. Surinyach, C.S. Kiminami, W.J. Botta, M.M. Peres, A.M. Jorge Jr, J. Alloys Comp. 509S (2011) S460–S463. [21] R. Cerutti, G.F. Lima, C.S. Kiminami, W.J. Botta, A.M. Jorge Jr., J. Alloys Comp. 509S (2011) S464–S467. [22] G.F. de Lima, S. Garroni, M.D. Baró, S. Suriñach, C.S. Kiminami, W.J. Botta, A.M. Jorge Jr, Int. J. Hydrogen Energy 37 (2012) 15196–15203. [23] J.J. Didisheim, P. Zolliker, K. Yvon, P. Fischer, J. Schefer, M. Gubelmann, A.F. Williams, Inorg. Chem. 23 (1984) 1953–1957. [24] R. Checcheto, N. Bazzanella, A. Miotello, P. Mengucci, J. Alloys Comp. 446–447 (2007) 58–62. [25] F.C. Gennari, F.J. Castro, J.J. Andrade Gamboa, J. Alloys Comp. 339 (2002) 261– 267. [26] H.M. Rietveld, J. Appl. Cryst. 2 (1969) 65–71. [27] L. Lutterotti, R. Ceccato, R. Dal Maschio, E. Pagani, Mater. Sci. Forum 278–281 (1998) 87–92. [28] A. Léon, E.J. Knystautas, J. Huot, R. Schulz, Thin Solid Films 496 (2006) 683– 687. [29] S. Singh, S.W.H. Eijt, M.W. Zandbergen, W.J. Legerstee, V.L. Svetchnikov, J. Alloys Comp. 441 (2007) 344–351. [30] J. Bellemare, J. Huot, J. Alloys Comp. 512 (2012) 33–34.
This sample presented faster hydrogen absorption kinetics and a higher hydrogen capacity than the samples processed by only hot-extrusion or high-energy ball milling.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Hydrogen storage properties of 2Mg–Fe after the combined processes of hot extrusion and cold rolling G.F. Lima ⇑, M.R.M. Triques, C.S. Kiminami, W.J. Botta, A.M. Jorge Jr. Department of Materials Engineering, Federal University of São Carlos, São Carlos, SP, Brazil
a r t i c l e
i n f o
Article history: Available online xxxx Keywords: Hydrogen storage properties 2 Mg–Fe Hot-extrusion Cold-rolling
a b s t r a c t Bulk consolidated samples of 2 Mg–Fe were produced by hot extrusion and by hot extrusion followed by cold rolling. The starting 2 Mg–Fe mixture was obtained by high-energy ball milling. The cold-rolled samples, which presented a favorable (0 0 2) texture in Mg, absorbed 5.23 wt.% H2, almost the full capacity of Mg2FeH6, without any incubation time. This level of hydrogen absorption capacity was higher than the ones observed for the milled powders or the extruded samples. The hydrogen desorption temperatures were very close for all types of samples. The good results observed in the hot-extruded plus cold-rolled sample were explained by the smaller grain sizes, along with a favorable texture and the catalytic effect of residual Fe. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium alloys are adequate materials for hydrogen storage in the solid state; magnesium hydride, MgH2, has the highest storage capacity (7.6 wt.% H) among the reversible hydrides. The main drawbacks to its extensive use are the poor kinetics of the Mg/ MgH2 system and the high H-sorption temperatures. Another unfavorable aspect of MgH2 when in the powder form is the poor air resistance, which results in the formation of a highly stable oxide layer, lowering absorption rates. In such cases, activation cycles at high temperatures and hydrogen pressures are necessary to reach full capacity [1,2]. To minimize the effects of these unfavorable aspects, investigations have been focused on the reduction of the crystallite size to destabilize MgH2 and to shorten diffusion paths, on the addition of ‘‘catalytic’’ elements or compounds to favor hydrogen dissociation and recombination, and on the improvement of surface properties to protect against surface contamination and alloying MgH2 to weak Mg–H bonds [2–8]. More recently, severe plastic deformation (SPD) methods have been employed to obtain Mg-hydrides with refined grain sizes, controlled textures, and attractive hydrogen storage properties [9–18]. The SPD-processed materials resulted in easier activation, rapidly reaching full capacity without any incubation time. In addition, these materials have a good air resistance due to the lower surface-volume ratio in comparison with powders [10]. The desorption kinetics of these materials were also enhanced by the presence of catalysts. In fact, good hydrogen storage properties have been observed in hot-extruded 2 Mg–Fe samples [19–21],
⇑ Corresponding author. Tel.: +55 16 33518531. E-mail address: [email protected] (G.F. Lima).
which, because of better air resistance, also presented a higher hydrogen capacity in the first cycle involving powder samples [22]. An important aspect of the presence of Fe is the possibility of maintaining the nanograin size after many hydrogenation cycles; the pinning effect of Fe at the a-Mg grain boundary is predominant even when the sample is maintained at high temperatures for long periods [19]. Fe participates in the formation of the complex hydride Mg2FeH6 [23]; in excess, it will act as a catalyst favoring the dissociation of the H2 molecule and acting in the desorption process, reducing both the decomposition temperature and the energy barrier for decomposition [24,25]. In the present work, we used two combined processes, hot extrusion and repetitive cold rolling, with a mixture of 2 Mg–Fe powders to produce textured bulk samples with nanoscale grain sizes. 2. Experimental Pure elements Mg (+20–100 mesh, 99.98%, Alfa Aesar) and Fe (20 mesh, 99.998%, Alfa Aesar) were mixed to produce stoichiometric 2 Mg–Fe using high-energy ball milling (HEBM) in a Fritsch P7 planetary ball mill. The ball-to-powder ratio was 40:1, and milling was carried out under an argon atmosphere for 4 h. Milled samples (1 g) were cold-pressed into pre-forms and then hot-extruded at 300 °C using a ram speed of 1 mm/min and an extrusion ratio of 3/1. The extruded samples, inserted between two AISI 304 stainless steel plates, were processed by cold rolling in a duo-reversible conventional rolling mill (FENN). The samples were submitted to several passes (20–30), with a 50% reduction in each pass. Both the hot-extrusion and cold-rolling processes were performed in air. The final foil thickness was approximately 0.150 mm. The hydrogen absorption properties were measured in a Sieverts-type apparatus at 673 K and 15 bar of H2(g). The hydrogen desorption behaviors were investigated by simultaneous differential scanning calorimetry–thermogravimetry–STA (TG-DSC) coupled with quadrupole mass spectrometry (QMS) in a Netzsch STA 449C + QMS 403C apparatus, with a constant heating rate of 10 °C/min under purified argon in an overflow regime.
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The structural characterization of the samples was carried out by X-ray diffraction (XRD) using monochromatic Cu Ka radiation in a Rigaku diffractometer equipped with a C-monochromator and by transmission electron microscopy (TEM, FEI-CM-120). The XRD profiles were treated by Rietveld refinement [26] using the MAUD software [27].
Table 1 Average crystallite sizes (±5 nm) of Mg and Fe estimated by Rietveld refinement using the MAUD software.
Mg Fe
3. Results and discussion Fig. 1 shows XRD patterns of the 2 Mg–Fe powder mixtures in comparison with the bulk obtained by hot extrusion and by hot extrusion followed by cold rolling. The XRD patterns revealed only the presence of hcp-Mg and bcc-Fe; there is no indication of the presence of oxides in any of the patterns. The XRD patterns corresponding to the powders and to the samples that were only extruded are quite similar, with no indication of a strongly preferred orientation. However, in the rolled sample, a strong (0 0 2) texture is observed for Mg. As first stated by Léon et al. [28] and after by Dufour and Huot [10,11] and by Singh et al. [29], the (0 0 2) orientation is a type of texture that favors hydrogen absorption during activation. Singh et al. [29] explained this behavior in an investigation of the structural and hydrogen storage properties in nanostructured thin films of Mg deposited on Si (0 0 1) substrates, X-ray diffraction showed that the conversion of Mg to MgH2 follows a martensitic-like orientation relationship with Mg 0//MgH2. (0 0 2)//MgH2 (1 1 0) and Mg ½1 2 Table 1 shows the average crystallite sizes of the powder and the extruded samples determined from Rietveld refinement. Because the Rietveld method is not adequate for samples presenting texture, the crystallite sizes of the cold-rolled samples were determined by TEM (Fig. 2). After 4 h of ball milling, the average crystallite sizes were in the nanometric scale, very close to that for the powder milled for 12 h, as already reported in a previous work [19], indicating that only 4 h is enough to produce nanostructured alloys. As expected, the extruded samples presented larger crystallite sizes than the ball-milled ones. This difference can be attributed to the annealing effect during hot extrusion, although the crystallite sizes were still kept in the nanoscale range because of the pinning effect of Fe particles at the a-Mg grain boundaries, as previously observed [19]. Fig. 2 shows bright-field (BF) and dark-field (DF) images and the corresponding selected-area electron diffraction patterns (SAEDPs) of the cold-rolled sample. A refined nanostructure is observed with most crystallites in the size range of 10 nm, although it is possible to observe a few grains that grow to approximately 60 nm. These images also show that the microstructure is homogeneous, with
Fig. 1. XRD patterns of 2 Mg–Fe mixtures in powder form obtained by milling, in bulk form as extruded, and in foil form as-extruded/rolled.
v2
2 Mg–Fe – milled (nm ± 5 nm)
2 Mg–Fe – extruded (nm ± 5 nm)
37 39 1.03
64 56 1.01
Fig. 2. Images obtained by TEM of extruded, cold-rolled 2 Mg–Fe: (a) bright-field and (b) dark-field images.
Fe being homogenously dispersed in the Mg matrix. The SAEDP confirms the presence of nanosized grains with a strong (0 0 2) texture, in agreement with the XRD pattern in Fig. 1. Fig. 3 shows the first hydrogen absorption curves for the 2 Mg– Fe samples processed by milling, extrusion, and extrusion followed by cold rolling. Absorption was measured at 673 K under a hydrogen pressure of 15 bar. The milled and cold-rolled samples presented faster kinetics than the extruded sample, did not have an incubation time, and reaching maximum capacity (saturation) after approximately 900 min (15 h). The time to reach maximum capacity is shorter than the times previously observed in coldrolled Mg–Pd – 1250 min [10] and 1700 min [11] – most likely because the samples in the present work had much smaller initial grain sizes. The 2 Mg–Fe composition chosen in this work was intended to form bulk textured samples of Mg2FeH6. The formation of this complex hydride was confirmed by the XRD pattern of the coldrolled sample after hydrogen absorption, as shown in Fig. 4. In addition to the Mg2FeH6 phase, we can observe peaks corresponding to residual Fe. The absorption curves of Fig. 3 also corroborates the efficient conversion of the initial 2 Mg–Fe mixtures to this
Please cite this article in press as: G.F. Lima et al., J. Alloys Comp. (2013), http://dx.doi.org/10.1016/j.jallcom.2013.01.115
G.F. Lima et al. / Journal of Alloys and Compounds xxx (2013) xxx–xxx
Fig. 3. Hydrogen absorption kinetics at 673 K (400 °C) and under a hydrogen pressure of 1500 kPa (15 bar) for 2 Mg–Fe milled, extruded and extruded/rolled samples.
complex phase; the cold-rolled samples could absorb 5.2% H, corresponding to 95% of the theoretical capacity. This value is significantly higher than the obtained in our previous work for the asextruded materials [19], which reached a maximum capacity of only 1.2% H. The good hydrogen storage capacity reached in the cold-rolled sample can then be attributed to the following factors: a favorable texture along the (0 0 2) direction in Mg, the additional deformation due to the cold-rolling process, and the catalytic action of residual Fe. In contrast, as shown in Fig. 3, the extruded sample presented an incubation time and very slow absorption kinetics. Consequently, a low hydrogen storage capacity was reached during the time of the experiment. A comparison of the two bulk nanostructured samples suggest that the texture along the (0 0 2) direction and the additional deformation caused by cold rolling are the most important factors contributing to the improved H-absorption kinetics and the capacity of the cold-rolled samples. The hydrogen desorption kinetics was estimated by a TG analysis, as shown in Fig. 5. The best desorption kinetics was observed for the powder sample, the one with the largest surface area. However, in agreement with Fig. 3, the cold-rolled sample presented the highest capacity, followed by the powder and the extruded samples. In contrast to our results, cold-rolled samples presented slower sorption kinetics and reduced capacities than corresponding samples obtained by milling [30]; however, in this case, there was an important difference in the crystallite sizes, which were smaller in the milled samples. The mass loss values are similar to
3
Fig. 5. Thermogravimetry measurements of 2 Mg–Fe milled, extruded, and extruded/rolled samples.
Fig. 6. DSC measurements of 2 Mg–Fe milled, extruded, and extruded/rolled samples.
but smaller than that found in Fig. 3. Most likely, these dissimilarities are due to differences between equipment and experimental analysis methodology. Additionally, the sample preparation was made in air for the DSC/TG analysis, which may result in a small level of oxidation that could not be detected by QMS analysis. Fig. 6 shows the DSC curves for a commercial MgH2 powder in comparison with the 2 Mg–Fe samples; the hydrogen desorption temperatures were lower for all the 2 Mg-Fe in all processing conditions. As expected, the lowest hydrogen desorption temperature was observed for the milled sample – the one with the highest surface area. Fig. 7 shows TEM images of hydrogenated samples. The nanograin sizes were maintained after absorption, and even the texture was not totally lost, as suggested by the electron diffraction pattern. Additional work is underway to investigate the texture stability in the cold-rolled sample. 4. Conclusions
Fig. 4. XRD patterns of hydrogenated 2 Mg–Fe as-extruded/rolled.
2 Mg–Fe mixtures processed by hot-extrusion combined with additional cold rolling produced, after hydrogenation, bulk nanostructured samples of the complex hydride Mg2FeH6 containing residual Fe.
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The better results observed for the combined processed sample may be explained by the smaller grain sizes, along with a favorable (0 0 2) texture of Mg and the catalyst effect of residual Fe. Acknowledgements The authors acknowledge the financial support of the Brazilian agencies FAPESP, CAPES, and CNPq. References
Fig. 7. Images obtained by TEM (bright field, 300 kX) of hydrogenated 2 Mg–Fe samples: (a) milled, (b) extruded and (c) extruded/rolled.
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This sample presented faster hydrogen absorption kinetics and a higher hydrogen capacity than the samples processed by only hot-extrusion or high-energy ball milling.
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