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Direct synthesis of Mg2FeH6 by mechanical alloying
Journal of Alloys and Compounds 280 (1998) 306–309
L
Direct synthesis of Mg 2 FeH 6 by mechanical alloying a, a b a J. Huot *, S. Boily , E. Akiba , R. Schulz a
Technologies Emergentes de production et de stockage, Institut de Recherche d’ Hydro-Quebec, 1800 Boul. Lionel-Boulet, Varennes, Quebec, Canada, J3 X 1 S1 b National Institute of Materials and Chemical Research, Higashi, Tsukuba, Ibaraki 305, Japan Received 22 June 1998
Abstract The hydride Mg 2 FeH 6 was synthesized by high-energy ball milling of MgH 2 and Fe under argon atmosphere without subsequent sintering. After 60 h of milling, 56% wt. of Mg 2 FeH 6 was synthesized. This yield was deduced from Rieltveld analysis of the X-ray powder measurements and confirmed by pressured differential scanning calorimeter (PDSC). Hydrogen capacity measurements indicated that the loss of capacity with cycling is minimal. 1998 Elsevier Science S.A. All rights reserved. Keywords: Metal hydrides; Mechanical alloying; Nanocrystalline material
1. Introduction Because of their high weight efficiencies for hydrogen storage, magnesium-based transition metal hydrides are interesting hydrogen storage materials [1]. The classic example is Mg 2 Ni, which has high hydrogen capacity (3.6 wt.%), good absorption kinetics and durability. However, its desorption temperature is too high for wide practical applications. Mg 2 Ni is a member of the electronic structural series Mg 2 FeH 6 –Mg 2 CoH 5 –Mg 2 NiH 4 . It is well known that these compounds are not easy to synthesize as single phase materials [2]. In this paper we present a method which dramatically increases the efficiency of hydride synthesis. As a demonstration of the technique, we have selected the compound Mg 2 FeH 6 because it is much harder to synthesize than Mg 2 NiH 4 . Mg 2 FeH 6 is cubic with K 2 PtCl 6 type structure. It 24 contains octahedral [FeH 6 ] complexes of anions surrounded by Mg in eight-fold cubic configuration [3]. Because magnesium and iron are immiscible, the conventional way of preparing the hydride Mg 2 FeH 6 is by
*Corresponding author.
sintering at elevated temperature (723–793K) under high hydrogen pressure (20–120 bar) [4]. However, the as prepared sample contains up to 50% of non-reacted elements and need to be purified. In a more recent work, Selvam and Yvon [2] showed that the formation of single phase Mg 2 FeH 6 required pressures of at least 90 bars and temperatures of at least 723 K. Konstanchuk et al. [5] showed that the hydride Mg 2 FeH 6 was synthesized by mechanical alloying a composition Mg–25%Fe for 5 min and successive hydrogenation and dehydrogenation at high temperature. Hightower et al. [6] have shown that for Mg concentration up to about 20 at.% mechanical alloying can ¨ produce single-phase bcc alloys. From Mossbauer spectroscopy, they found Mg-rich regions in iron lattice. In the same way, iron atoms form clusters in hcp Mg-rich alloys. Recently, it was shown that by milling the composition 2Mg1Fe in a planetary mill for 20 h, it was possible to synthesize Mg 2 FeH 6 by sintering at 623 K under hydrogen pressure of 50 bars [7]. The next step was to try to synthesize Mg 2 FeH 6 without sintering. As shown in the work of Hightower et al. [6], milling of Mg and Fe produces very finely dispersed clusters of one phase in the other. The Mg 2 FeH 6 phase being more stable than MgH 2 , the conjecture was that by finely dispersing MgH 2 in Fe by ball milling, the ternary hydride Mg 2 FeH 6 would be directly formed without the need of sintering. For this study, we selected the stoichiometry 2MgH 2 1
0925-8388 / 98 / $ – see front matter 1998 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 98 )00725-7
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
Fe. In such composition, a complete formation of Mg 2 FeH 6 will leave some unreacted Mg and Fe.
2. Experimental details Pure magnesium hydride (98%) was provided by Th. Goldschmidt. It was mixed with iron powder (Atomet 1001 from Quebec Fer / Titane, 99.5%) inside an argon filled glove box. An oxygen content of 0.12 wt.% in the iron powder was measured with a LECO Nitrogen / Oxygen detector, model TC-136. The milling was carried out with a Spex mill 8000 and a vial and balls of stainless steel. The ball to powder weight ratio was 10:1. Milling was carried out for up to 60 h. At regular intervals, small amount of powder was taken for analysis. The X-ray powder diffraction were performed on a Philips X’pert system with Cu K(radiation. The X-ray diffraction pattern was analyzed by the Rietveld method using RIETAN software [8]. Lattice parameters, phase abundance, crystallite size and strain were extracted from the Rietveld refinement analyses. The thermal behavior was studied by pressurized differential scanning calorimeter (PDSC) using a TA 2190 calorimeter at heating rate of 20 K min 21 under 10 bar of hydrogen. The pressure–composition isotherms and kinetics of hydrogen absorption / desorption were measured on a specially designed gas titration apparatus.
3. Results and discussion
3.1. Crystallographic parameters In Fig. 1 we present the X-ray powder diffraction patterns as a function of milling time. After only 1 h of
307
milling, there is evidence of formation of an orthorhombic metastable g-MgH 2 [9]. The origin and condition of appearance of this metastable phase by ball milling will be discussed elsewhere. The 10 h of milling diffraction pattern showed the first appearance of Mg 2 FeH 6 phase with the MgH 2 and Fe phases still present. After 30 h of milling, the Mg 2 FeH 6 phase is more abundant and the MgH 2 phase has disappeared. Moreover, the Mg peaks are now present. This is an indication that all hydrogen initially in the MgH 2 phase has been used to synthesized Mg 2 FeH 6 , leaving some unreacted magnesium. This is confirmed by the fact that further milling up to 60 h does not significantly change the phase distribution. A Rietveld refinement was performed on the 60 h sample. By simple inspection of the powder diffraction pattern, it is very hard to identify the presence of MgO. However, the residue of a Rietveld refinement using only the phases Mg 2 FeH 6 , Mg and Fe clearly shows the presence of diffraction peaks attributed to MgO. This is an example of the power of Rietveld refinement for minor phases identification [10]. To judge the quality of Rietveld refinement, the most meaningful indicator is the ‘‘R-weighted pattern’’ R wp , which measures the weighted difference between the calculated and measured intensities. The ‘‘R-expected’’ value R e is an estimation of the minimum value of R wp . The ‘‘R-Bragg’’ R B measures the difference between calculated and ‘‘experimental’’ intensities of the Bragg reflections. The mathematical description of these quantities can be found in [11]. A reliable indicator of the goodness of the fit is the ratio R wp /R e 5S. An S value less than 1.3 is usually considered satisfactory [11]. A very good fit was obtained by fitting the phases Mg 2 FeH 6 , Mg, Fe, and MgO. After refinement, S51.20 and the Bragg factors for Mg 2 FeH 6 , Mg, Fe and MgO were respectively 2.70%, 2.79%, 1.34% and 3.06%. For each phase, the phase abundance, lattice parameter, crystallite size and strain determined from Rietveld refinement is shown in Table 1. The strain value is almost the same for each phase. Similarly, the crystallite size is of the same order of magnitude for all phases. Despite the fact that broad peaks are not ideal for lattice parameter determination, the value for Mg 2 FeH 6 is surprisingly good, as shown in Table 2. The value found in this work is in good
Table 1 Crystallographic parameters deduced from Rietveld refinement of the X-ray diffraction pattern of the 60 h milled sample
Fig. 1. X-ray powder diffraction pattern of 2MgH 2 1Fe mixture as a function of milling time m, MgH 2 ; j, g MgH 2 ; d, Fe; ♦, Mg 2 FeH 6 ; ., Mg; 3, MgO.
Phase
Weight %
Lattice parameter nm
Crystallite size (nm)
Strain (%)
Mg 2 FeH 6 Fe Mg
52(1) 31(1) 10(2)
a50.6444(3) a50.2874(1) a50.322(2) c50.520(4) a50.423(1)
27.0(6) 19.0(2) 9(1)
1.62(1) 1.13(1) 1.7(1)
12(4)
1 (1)
MgO
7(1)
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
308 Table 2 Lattice parameter (nm) of Mg 2 FeH 6 This work
Earlier work [7]
Didisheim et al. [4]
Selvam and Yvon [2]
0.6444(3)
0.64419(2)
0.6443(1)
0.64287(7)
agreement with the previous results, except for the more recent study of Selvam and Yvon [2]. The origin of the MgO phase has to be explained. In the system Mg–Ni50 wt.% it was shown that MgO originates from the oxide film on magnesium particles [12]. It has also been shown that in the synthesis of Al–AlN composites, oxygen comes from the elemental powders [13]. Before milling, the oxygen content of iron and magnesium hydride were respectively 0.12 wt.% and 0.6 wt.%, giving a total oxygen content of 0.35 wt.%. Because the reduction potential of magnesium is very low, the formation of MgO was favored. It has already been shown that oxides of iron can be reduced by magnesium during mechanical treatment [5]. After milling, the oxygen content from MgO phase was 0.437 wt.%52.8 wt.%. This high oxygen content could be explained by the high reactivity of magnesium which can react with oxides from the vial wall and balls to form MgO. The powder diffraction pattern of the desorbed sample is shown on Fig. 2. The peaks are narrow, indicating a certain degree of grain growth induced by hydrogen absorption / desorption. A Rietveld refinement was performed, giving the following R-factors: R wp /R e 5S51.24. The Bragg factors of Mg, Fe and MgO were respectively 13.46%, 9.95% and 7.55%. As seen in Table 3, for Mg and Fe phases, there is an important increase of crystallite size associated with strain relaxation. Except for a certain increase in abundance, the strain of MgO phase did not change. This confirms the inertia of this phase under hydrogen cycling.
Fig. 2. X-ray powder diffraction pattern of desorbed sample. Short vertical bars indicate the Bragg reflections of each fitted phase: upper row is Mg, middle row is Fe and lower row is MgO. The bottom line is the difference between observed pattern and Rietveld refinement.
Table 3 Crystallographic parameters deduced from Rietveld refinement of X-ray diffraction pattern of dehydrided sample Phase
Weight %
Lattice parameter nm
Crystallite size (nm)
Strain (%)
Mg
34(4)
202(47)
0.314(5)
Fe MgO
55(3) 11(5)
a50.3211(1) c50.5215(2) a50.28662(3) a50.423(3)
88(3) 6(1)
0.324(1) 0.9(6)
3.2. Calorimetric measurements In Fig. 3, we present the PDSC trace of 2MgH 2 1Fe ball-milled for 60 h. The endotherm at 700 K is associated to the desorption of Mg 2 FeH 6 . From the evolved heat of this peak and the heat of formation of Mg 2 FeH 6 [4], an abundance of 55 wt.% was calculated. This is in agreement with the value derived from X-ray measurements.
3.3. Hydrogen absorption From the phase abundance, Table 4, of the 60 h milled sample, the expected hydrogen capacity is: 0.5235.4%5 2.8 wt.% from the Mg 2 FeH 6 phase and 0.1037.6%50.76 wt.% from magnesium. The as-ball-milled sample was dehydrided at 623 K and pressure–composition isotherms at 623 K and 673 K were registered. The results are shown on Fig. 4. At 623 K, on the absorption side, the first plateau at 6 atm is the hydriding of magnesium up to a capacity of 2.4 wt.%. The second plateau indicates the formation of Mg 2 FeH 6 with a capacity of 1.8 wt.%. From these values, one may conclude that at the end of absorption, the phase composition is 33 wt.% of Mg 2 FeH 6 phase and 32 wt.% of magnesium phase. However, on the desorption side, the magnesium plateau has a capacity of only 1 wt.% while the Mg 2 FeH 6 plateau has 3.2 wt.%, giving a phase abundance of Mg 2 FeH 6 and magnesium of 59 wt.% and 13 wt.% respectively. This is very close to the
Fig. 3. PDSC trace, under 2 bar of hydrogen, of 2MgH 2 1Fe milled for 60 h.
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
309
Table 4 Phase abundance, in weight %, as determined by different methods a Phase
Mg 2 FeH 6 Fe Mg MgO a
Initial powder
52 44
Total transformation (calculation)
X-ray (as-milled)
PDSC
PCT (673 K)
PCT (623 K)
68 17 15
52(1) 31(1) 10(2) 7(1)
55
48
59
18
13
X-ray (dehydrided)
55(3) 34(2) 11(5)
In the initial powder, the abundance of hydrogen is 4 wt.%.
abundance determined by X-ray diffraction. This is a proof that, during the measurement, Mg 2 FeH 6 was formed from the MgH 2 phase. The same phenomenon is seen on the 673 K isotherm. This formation of Mg 2 FeH 6 phase at low temperature and pressure is favored by the fine mixing of Mg and Fe. This interpretation is supported by the finding of Hightower et al. [6]. In the case of hydrides synthesized by sintering, it has been reported that after only four consecutive cycles of absorption / desorption, the hydrogen capacity decreased by about 35% [4]. In this work, no extensive cycling has been performed. However, after five cycles, the hydrogen capacity has decreased by only 5%. This is another evidence that the fine microstructure enhanced the ternary alloy formation.
at room temperature with excellent yield. Moreover, because of the very fine microstructure, the ball-milled hydride shows better reversibility than the hydride produced by sintering. The loss of hydride capacity upon cycling was reduced. It has also been shown that oxygen impurities in the starting powder and oxides from the vial wall and balls can produce a relatively high content of detrimental MgO phase.
Acknowledgements We wish to thank Dr. T. Klassen of GKSS for helpful discussion.
References 4. Conclusion It has been shown that by milling magnesium hydride with iron, the ternary hydride Mg 2 FeH 6 can be synthesized
Fig. 4. Pressure–composition isotherm of nanocrystalline 2MgH 2 1Fe milled for 60 h.
[1] P. Selvam, B. Viswanathan, C.S. Swamy, V. Srinivasan, Int. J. Hydrogen Energy 11 (1986) 169. [2] P. Selvam, K. Yvon, Int. J. Hydrogen Energy 16 (1991) 615. [3] K. Yvon, Hydrides: Solid State Transition Metal Complexes, Wiley, New-York, 1994, p. 1401. [4] J.-J. Didisheim, P. Zolliker, K. Yvon, P. Fisher, J. Schefer, M. Gubelmann, A.F. Williams, Inorg. Chem. 23 (1984) 1953. [5] I.G. Konstanchuk, E.Y. Ivanov, M. Pezat, B. Darriet, V.V. Boldyrev, P. Hagenmuller, J. Less-Common Met. 131 (1987) 181. [6] A. Hightower, B. Fultz, R.C. Bowmann, J. Alloys Comp. 252 (1997) 238. [7] J. Huot, H. Hayakawa, E. Akiba, J. Alloys Comp. 248 (1997) 164. [8] F. Izumi, in: R.A. Young (Ed.), The Rietveld Method, Oxford University Press, Oxford, 1993, Chap. 13. [9] J.-P. Bastide, B. Bonnetot, J.-M. Letoffe, P. Claudy, Mater. Res. Bull. 15 (1980) 1215. [10] J. Huot, E. Akiba, Y. Ishido, J. Alloys Comp. 231 (1995) 85. [11] R.A. Young, in: R.A. Young (Ed.), The Rietveld Method, Oxford University Press, Oxford, 1993, Chap. 1. [12] E. Ivanov, I. Konstanchuk, A. Stepanov, V. Boldyrev, J. LessCommon Met. 131 (1987) 25. [13] A. Calka, J.I. Nikolov, J. Appl. Phys. 75 (1994) 4953.
L
Direct synthesis of Mg 2 FeH 6 by mechanical alloying a, a b a J. Huot *, S. Boily , E. Akiba , R. Schulz a
Technologies Emergentes de production et de stockage, Institut de Recherche d’ Hydro-Quebec, 1800 Boul. Lionel-Boulet, Varennes, Quebec, Canada, J3 X 1 S1 b National Institute of Materials and Chemical Research, Higashi, Tsukuba, Ibaraki 305, Japan Received 22 June 1998
Abstract The hydride Mg 2 FeH 6 was synthesized by high-energy ball milling of MgH 2 and Fe under argon atmosphere without subsequent sintering. After 60 h of milling, 56% wt. of Mg 2 FeH 6 was synthesized. This yield was deduced from Rieltveld analysis of the X-ray powder measurements and confirmed by pressured differential scanning calorimeter (PDSC). Hydrogen capacity measurements indicated that the loss of capacity with cycling is minimal. 1998 Elsevier Science S.A. All rights reserved. Keywords: Metal hydrides; Mechanical alloying; Nanocrystalline material
1. Introduction Because of their high weight efficiencies for hydrogen storage, magnesium-based transition metal hydrides are interesting hydrogen storage materials [1]. The classic example is Mg 2 Ni, which has high hydrogen capacity (3.6 wt.%), good absorption kinetics and durability. However, its desorption temperature is too high for wide practical applications. Mg 2 Ni is a member of the electronic structural series Mg 2 FeH 6 –Mg 2 CoH 5 –Mg 2 NiH 4 . It is well known that these compounds are not easy to synthesize as single phase materials [2]. In this paper we present a method which dramatically increases the efficiency of hydride synthesis. As a demonstration of the technique, we have selected the compound Mg 2 FeH 6 because it is much harder to synthesize than Mg 2 NiH 4 . Mg 2 FeH 6 is cubic with K 2 PtCl 6 type structure. It 24 contains octahedral [FeH 6 ] complexes of anions surrounded by Mg in eight-fold cubic configuration [3]. Because magnesium and iron are immiscible, the conventional way of preparing the hydride Mg 2 FeH 6 is by
*Corresponding author.
sintering at elevated temperature (723–793K) under high hydrogen pressure (20–120 bar) [4]. However, the as prepared sample contains up to 50% of non-reacted elements and need to be purified. In a more recent work, Selvam and Yvon [2] showed that the formation of single phase Mg 2 FeH 6 required pressures of at least 90 bars and temperatures of at least 723 K. Konstanchuk et al. [5] showed that the hydride Mg 2 FeH 6 was synthesized by mechanical alloying a composition Mg–25%Fe for 5 min and successive hydrogenation and dehydrogenation at high temperature. Hightower et al. [6] have shown that for Mg concentration up to about 20 at.% mechanical alloying can ¨ produce single-phase bcc alloys. From Mossbauer spectroscopy, they found Mg-rich regions in iron lattice. In the same way, iron atoms form clusters in hcp Mg-rich alloys. Recently, it was shown that by milling the composition 2Mg1Fe in a planetary mill for 20 h, it was possible to synthesize Mg 2 FeH 6 by sintering at 623 K under hydrogen pressure of 50 bars [7]. The next step was to try to synthesize Mg 2 FeH 6 without sintering. As shown in the work of Hightower et al. [6], milling of Mg and Fe produces very finely dispersed clusters of one phase in the other. The Mg 2 FeH 6 phase being more stable than MgH 2 , the conjecture was that by finely dispersing MgH 2 in Fe by ball milling, the ternary hydride Mg 2 FeH 6 would be directly formed without the need of sintering. For this study, we selected the stoichiometry 2MgH 2 1
0925-8388 / 98 / $ – see front matter 1998 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 98 )00725-7
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
Fe. In such composition, a complete formation of Mg 2 FeH 6 will leave some unreacted Mg and Fe.
2. Experimental details Pure magnesium hydride (98%) was provided by Th. Goldschmidt. It was mixed with iron powder (Atomet 1001 from Quebec Fer / Titane, 99.5%) inside an argon filled glove box. An oxygen content of 0.12 wt.% in the iron powder was measured with a LECO Nitrogen / Oxygen detector, model TC-136. The milling was carried out with a Spex mill 8000 and a vial and balls of stainless steel. The ball to powder weight ratio was 10:1. Milling was carried out for up to 60 h. At regular intervals, small amount of powder was taken for analysis. The X-ray powder diffraction were performed on a Philips X’pert system with Cu K(radiation. The X-ray diffraction pattern was analyzed by the Rietveld method using RIETAN software [8]. Lattice parameters, phase abundance, crystallite size and strain were extracted from the Rietveld refinement analyses. The thermal behavior was studied by pressurized differential scanning calorimeter (PDSC) using a TA 2190 calorimeter at heating rate of 20 K min 21 under 10 bar of hydrogen. The pressure–composition isotherms and kinetics of hydrogen absorption / desorption were measured on a specially designed gas titration apparatus.
3. Results and discussion
3.1. Crystallographic parameters In Fig. 1 we present the X-ray powder diffraction patterns as a function of milling time. After only 1 h of
307
milling, there is evidence of formation of an orthorhombic metastable g-MgH 2 [9]. The origin and condition of appearance of this metastable phase by ball milling will be discussed elsewhere. The 10 h of milling diffraction pattern showed the first appearance of Mg 2 FeH 6 phase with the MgH 2 and Fe phases still present. After 30 h of milling, the Mg 2 FeH 6 phase is more abundant and the MgH 2 phase has disappeared. Moreover, the Mg peaks are now present. This is an indication that all hydrogen initially in the MgH 2 phase has been used to synthesized Mg 2 FeH 6 , leaving some unreacted magnesium. This is confirmed by the fact that further milling up to 60 h does not significantly change the phase distribution. A Rietveld refinement was performed on the 60 h sample. By simple inspection of the powder diffraction pattern, it is very hard to identify the presence of MgO. However, the residue of a Rietveld refinement using only the phases Mg 2 FeH 6 , Mg and Fe clearly shows the presence of diffraction peaks attributed to MgO. This is an example of the power of Rietveld refinement for minor phases identification [10]. To judge the quality of Rietveld refinement, the most meaningful indicator is the ‘‘R-weighted pattern’’ R wp , which measures the weighted difference between the calculated and measured intensities. The ‘‘R-expected’’ value R e is an estimation of the minimum value of R wp . The ‘‘R-Bragg’’ R B measures the difference between calculated and ‘‘experimental’’ intensities of the Bragg reflections. The mathematical description of these quantities can be found in [11]. A reliable indicator of the goodness of the fit is the ratio R wp /R e 5S. An S value less than 1.3 is usually considered satisfactory [11]. A very good fit was obtained by fitting the phases Mg 2 FeH 6 , Mg, Fe, and MgO. After refinement, S51.20 and the Bragg factors for Mg 2 FeH 6 , Mg, Fe and MgO were respectively 2.70%, 2.79%, 1.34% and 3.06%. For each phase, the phase abundance, lattice parameter, crystallite size and strain determined from Rietveld refinement is shown in Table 1. The strain value is almost the same for each phase. Similarly, the crystallite size is of the same order of magnitude for all phases. Despite the fact that broad peaks are not ideal for lattice parameter determination, the value for Mg 2 FeH 6 is surprisingly good, as shown in Table 2. The value found in this work is in good
Table 1 Crystallographic parameters deduced from Rietveld refinement of the X-ray diffraction pattern of the 60 h milled sample
Fig. 1. X-ray powder diffraction pattern of 2MgH 2 1Fe mixture as a function of milling time m, MgH 2 ; j, g MgH 2 ; d, Fe; ♦, Mg 2 FeH 6 ; ., Mg; 3, MgO.
Phase
Weight %
Lattice parameter nm
Crystallite size (nm)
Strain (%)
Mg 2 FeH 6 Fe Mg
52(1) 31(1) 10(2)
a50.6444(3) a50.2874(1) a50.322(2) c50.520(4) a50.423(1)
27.0(6) 19.0(2) 9(1)
1.62(1) 1.13(1) 1.7(1)
12(4)
1 (1)
MgO
7(1)
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
308 Table 2 Lattice parameter (nm) of Mg 2 FeH 6 This work
Earlier work [7]
Didisheim et al. [4]
Selvam and Yvon [2]
0.6444(3)
0.64419(2)
0.6443(1)
0.64287(7)
agreement with the previous results, except for the more recent study of Selvam and Yvon [2]. The origin of the MgO phase has to be explained. In the system Mg–Ni50 wt.% it was shown that MgO originates from the oxide film on magnesium particles [12]. It has also been shown that in the synthesis of Al–AlN composites, oxygen comes from the elemental powders [13]. Before milling, the oxygen content of iron and magnesium hydride were respectively 0.12 wt.% and 0.6 wt.%, giving a total oxygen content of 0.35 wt.%. Because the reduction potential of magnesium is very low, the formation of MgO was favored. It has already been shown that oxides of iron can be reduced by magnesium during mechanical treatment [5]. After milling, the oxygen content from MgO phase was 0.437 wt.%52.8 wt.%. This high oxygen content could be explained by the high reactivity of magnesium which can react with oxides from the vial wall and balls to form MgO. The powder diffraction pattern of the desorbed sample is shown on Fig. 2. The peaks are narrow, indicating a certain degree of grain growth induced by hydrogen absorption / desorption. A Rietveld refinement was performed, giving the following R-factors: R wp /R e 5S51.24. The Bragg factors of Mg, Fe and MgO were respectively 13.46%, 9.95% and 7.55%. As seen in Table 3, for Mg and Fe phases, there is an important increase of crystallite size associated with strain relaxation. Except for a certain increase in abundance, the strain of MgO phase did not change. This confirms the inertia of this phase under hydrogen cycling.
Fig. 2. X-ray powder diffraction pattern of desorbed sample. Short vertical bars indicate the Bragg reflections of each fitted phase: upper row is Mg, middle row is Fe and lower row is MgO. The bottom line is the difference between observed pattern and Rietveld refinement.
Table 3 Crystallographic parameters deduced from Rietveld refinement of X-ray diffraction pattern of dehydrided sample Phase
Weight %
Lattice parameter nm
Crystallite size (nm)
Strain (%)
Mg
34(4)
202(47)
0.314(5)
Fe MgO
55(3) 11(5)
a50.3211(1) c50.5215(2) a50.28662(3) a50.423(3)
88(3) 6(1)
0.324(1) 0.9(6)
3.2. Calorimetric measurements In Fig. 3, we present the PDSC trace of 2MgH 2 1Fe ball-milled for 60 h. The endotherm at 700 K is associated to the desorption of Mg 2 FeH 6 . From the evolved heat of this peak and the heat of formation of Mg 2 FeH 6 [4], an abundance of 55 wt.% was calculated. This is in agreement with the value derived from X-ray measurements.
3.3. Hydrogen absorption From the phase abundance, Table 4, of the 60 h milled sample, the expected hydrogen capacity is: 0.5235.4%5 2.8 wt.% from the Mg 2 FeH 6 phase and 0.1037.6%50.76 wt.% from magnesium. The as-ball-milled sample was dehydrided at 623 K and pressure–composition isotherms at 623 K and 673 K were registered. The results are shown on Fig. 4. At 623 K, on the absorption side, the first plateau at 6 atm is the hydriding of magnesium up to a capacity of 2.4 wt.%. The second plateau indicates the formation of Mg 2 FeH 6 with a capacity of 1.8 wt.%. From these values, one may conclude that at the end of absorption, the phase composition is 33 wt.% of Mg 2 FeH 6 phase and 32 wt.% of magnesium phase. However, on the desorption side, the magnesium plateau has a capacity of only 1 wt.% while the Mg 2 FeH 6 plateau has 3.2 wt.%, giving a phase abundance of Mg 2 FeH 6 and magnesium of 59 wt.% and 13 wt.% respectively. This is very close to the
Fig. 3. PDSC trace, under 2 bar of hydrogen, of 2MgH 2 1Fe milled for 60 h.
J. Huot et al. / Journal of Alloys and Compounds 280 (1998) 306 – 309
309
Table 4 Phase abundance, in weight %, as determined by different methods a Phase
Mg 2 FeH 6 Fe Mg MgO a
Initial powder
52 44
Total transformation (calculation)
X-ray (as-milled)
PDSC
PCT (673 K)
PCT (623 K)
68 17 15
52(1) 31(1) 10(2) 7(1)
55
48
59
18
13
X-ray (dehydrided)
55(3) 34(2) 11(5)
In the initial powder, the abundance of hydrogen is 4 wt.%.
abundance determined by X-ray diffraction. This is a proof that, during the measurement, Mg 2 FeH 6 was formed from the MgH 2 phase. The same phenomenon is seen on the 673 K isotherm. This formation of Mg 2 FeH 6 phase at low temperature and pressure is favored by the fine mixing of Mg and Fe. This interpretation is supported by the finding of Hightower et al. [6]. In the case of hydrides synthesized by sintering, it has been reported that after only four consecutive cycles of absorption / desorption, the hydrogen capacity decreased by about 35% [4]. In this work, no extensive cycling has been performed. However, after five cycles, the hydrogen capacity has decreased by only 5%. This is another evidence that the fine microstructure enhanced the ternary alloy formation.
at room temperature with excellent yield. Moreover, because of the very fine microstructure, the ball-milled hydride shows better reversibility than the hydride produced by sintering. The loss of hydride capacity upon cycling was reduced. It has also been shown that oxygen impurities in the starting powder and oxides from the vial wall and balls can produce a relatively high content of detrimental MgO phase.
Acknowledgements We wish to thank Dr. T. Klassen of GKSS for helpful discussion.
References 4. Conclusion It has been shown that by milling magnesium hydride with iron, the ternary hydride Mg 2 FeH 6 can be synthesized
Fig. 4. Pressure–composition isotherm of nanocrystalline 2MgH 2 1Fe milled for 60 h.
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