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Anelasticity in nanostructured MgH2–Mg: Correlation with hydrogen sorption kinetics
Materials Science and Engineering A 521–522 (2009) 151–154
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Anelasticity in nanostructured MgH2 –Mg: Correlation with hydrogen sorption kinetics L. Pasquini a,∗ , S. Amadori a , E. Bonetti a , E. Callini a , A.L. Fiorini a , A. Montone b , M. Vittori Antisari b a b
Dipartimento di Fisica, Università di Bologna and CNISM, v.le Berti-Pichat 6/2, I-40127 Bologna, Italy Materials and Technology Unit, ENEA C.R. Casaccia, v. Anguillarese 301, I-00060 S. Maria di Galeria, Roma, Italy
a r t i c l e
i n f o
Article history: Received 13 June 2008 Accepted 22 September 2008 Keywords: Magnesium hydride Hydrogen storage Mechanical spectroscopy Nanostructure
a b s t r a c t Mechanical spectroscopy is employed, in correlation with differential scanning calorimetry and hydrogen sorption analysis, to investigate the hydrogen desorption reaction in nanostructured magnesium hydride prepared by ball milling. The hydrogen desorption process is marked by a dramatic drop of the specimen’s resonance frequency and by an internal friction peak that occurs – under high vacuum – at about 580 K. After full hydrogen desorption, a broad anelastic relaxation peak similar to commercial Mg is observed at 470 K, with activation energy close to the one for self-diffusion in Mg. © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Metal hydrides that undergo reversible hydrogen sorption reactions are the subject of intensive research activities owing to their application as hydrogen storage materials. In particular, MgH2 is very attractive thanks to its high storage capacity (7.6 wt.%) coupled with the natural abundance and low cost of Mg. However, critical features for the technological exploitation of Mg-based materials are related to the strong Mg–H interaction, reflected by the relatively high desorption temperature, and to the sluggish reaction kinetics due to the poor hydrogen transport both at the surface and in the bulk of magnesium. A possible solution to the slow reaction problem lies in the refinement of the MgH2 crystallite size down to the nanometer range [1–3]. In this article we employ mechanical spectroscopy to investigate hydrogen desorption in nanostructured MgH2 . Mechanical spectroscopy has a long tradition in the study of hydrogen-related processes, such as the short-range hydrogen mobility in amorphous and crystalline materials [4] or the investigation of long-range diffusion in nanocrystalline solids [5]. Here we correlate mechanical spectroscopy with results of thermal analysis techniques to gain more insight into the physics underlying the desorption process.
2.1. Sample preparation
∗ Corresponding author. Tel.: +39 0512095149; fax: +39 0512095113. E-mail address: [email protected] (L. Pasquini). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.114
Nanostructured (n-) materials were synthesized by ball milling in a Spex Mixer/Mill 8000 equipped with hardened steel vials and balls (ball-to-powder weight ratio = 10). Commercial magnesium hydride MgH2 from Th. Goldschmidt (60 m particle size, 5 wt.% residual Mg) was employed as starting material for the synthesis of n-MgH2 , while Mg powder from Alfa Aesar (−325 mesh, 99.8% purity) was used as starting material for the preparation of n-Mg. The vials were filled under Ar atmosphere and ball milling was carried out for 10 h. To obtain specimens suitable for anelastic measurements the ball milled powders were compacted in vacuum (1 Pa) at room temperature under a uniaxial pressure of 1 GPa. Reference specimens were also prepared by compacting the starting powders under identical conditions. Four specimen classes are studied in this paper: n-MgH2 , n-Mg, MgH2 , and Mg. 2.2. Characterization The microstructure of the specimens was characterized by Rietveld analysis of X-ray diffraction (XRD) profiles taken with a Rigaku-DMAX IIIC Bragg–Brentano diffractometer, equipped with Cu K␣ radiation and a graphite monochromator in the diffracted beam. Mechanical spectroscopy in the acoustic regime was carried out with a vibrating reed analyzer model VRA 1604 by Cantil [6], operated under high vacuum conditions (10−3 Pa) at a constant heating
152
L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
Table 1 Results of Rietveld analysis on XRD spectra of n-ball milled materials and starting commercial powders. Material
d (nm)
-MgH2 (wt.%)
␥-MgH2 (wt.%)
Mg (wt.%)
MgO (wt.%)
n-MgH2 n-Mg MgH2 Mg
10 (1) 60 (5) >100 >100
62 (2) 0 95 0
30 (2) 0 0 0
1 94 (2) 5 99
7 (2) 6 (2) 0 1
Crystallite size d of the majority phase and relative abundance of the Mg, MgH2 , and MgO phases.
rate of 1 K/min. Throughout this paper, the internal friction Q−1 and normalized squared resonance frequency (f/f0 )2 will be reported, where f0 is the resonance frequency of the as-compacted specimen at ambient temperature. Additional Q−1 data at low frequency were obtained with a dynamic mechanical analyzer (DMA) model DMA 2980 by TA Instruments operating in flowing Ar at a constant heating rate of 1 K/min. The hydrogen desorption process was studied with a hydrogen sorption analyzer (HSA) by Cantil and a differential scanning calorimeter (DSC) model Q10 by TA Instruments. Both instruments were operated under 0.1 MPa Ar pressure in thermal scan mode at 1 K/min heating rate. In the HSA, the pressure variation during hydrogen release is converted into relative mass change m by using a calibrated volume. 3. Results 3.1. X-ray diffraction analysis Table 1 summarizes the microstructure parameters obtained from Rietveld analysis of the XRD profiles shown in Fig. 1. The commercial MgH2 displays narrow peaks pertaining to the tetragonal -MgH2 phase and residual traces of Mg (5 wt.%). The n-MgH2 specimen is characterized by very broad peaks owing to the fine crystallite size of 10 nm. Further, the presence of orthorhombic ␥MgH2 phase (30 wt.% amount) is evident from Fig. 1. This phase can be considered as a distortion of tetragonal -MgH2 with the same packing type and coordination number, and its appearance is a consequence of the large local strain induced by high energy ball milling [3]. The crystallite size of n-Mg is larger (about 60 nm) due to the ductility of Mg which makes the milling process less efficient in comparison with brittle MgH2 . Both n-materials contain about 6–7 wt.% MgO, which in the case of n-MgH2 is due to partial oxidation of the residual Mg in the commercial hydride.
Fig. 1. XRD profiles of investigated specimens, showing also the effects of one heating run on n-MgH2 . The positions of the ␥-MgH2 reflections are shown at the bottom.
3.2. Mechanical spectroscopy and thermal analysis Fig. 2(a) shows the first VRA run on n-MgH2 and MgH2 specimens. Two remarkable features are visible: (i) a Q−1 peak at 350 K for n-MgH2 and 420 K for MgH2 ; (ii) a dramatic drop of (f/f0 )2 which ends at 600 K and 650 K for n-MgH2 and MgH2 respectively. Notably, in n-MgH2 , the Q−1 peak at 350 K couples with an anomalous trend of (f/f0 )2 , while the (f/f0 )2 drop at high temperature is accompanied by a fast decrease of Q−1 to below the value of MgH2 , resulting in a Q−1 peak. In the second run (Fig. 2(b)), the above features disappear: both specimens exhibit a broad Q−1 peak centred at about 470 K and superposed to an increasing background. Fig. 3 reports a series of second runs, carried out by VRA on Mg, n-Mg and n-MgH2 . All specimens display a broad Q−1 peak whose temperature ranges from 440 K of n-Mg to 490 K of Mg. The temperature shift in DMA measures at 1 Hz proves that the peak is a thermally activated relaxation process. Fig. 4 displays DSC and HSA heating runs carried out on the same hydride materials. The hydrogen desorption is marked by the DSC
Fig. 2. First (a) and second (b) VRA heating runs at 1 K/min on hydride specimens.
L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
Fig. 3. Mechanical spectroscopy on Mg, n-Mg and n-MgH2 after full transformation to metal. DMA measures at 1 Hz (open symbols) were carried out in the 300–520 K range only. The other curves are VRA data taken at the indicated resonance frequency.
Fig. 4. DSC and HSA runs at 1 K/min under 0.1 MPa Ar atmosphere. The inset displays a magnified view of the low-temperature DSC traces.
Table 2 Figures of the hydrogen desorption process for the two hydride specimens. Onset temperature TON and integral area AP of the DSC endothermic peak, relative mass change m measured by HSA, and variation of (f/f0 )2 measured at ambient temperature after the first heating run. Material
TON (K)
AP (J/g)
m (wt.%)
(f/f0 )2
n-MgH2 MgH2
624 658
1920 1980
−6.1 −6.6
−0.67 −0.74
endothermic process and by the mass decrease m measured in the HSA. Table 2 reports the main figures of the desorption process together with the (f/f0 )2 variation measured at ambient temperature after the first heating run. 4. Discussion To investigate in more detail the two main features reported in Section 3.2 for the anelastic spectrum of as-compacted hydride specimens, we have carried out intermediate heating runs on nMgH2 . The comparison between the 1st and 2nd run (Fig. 5) proves unambiguously that the Q−1 peak at 350 K is associated with a non-
153
Fig. 5. Successive VRA runs at 1 K/min on n-MgH2 . In the 2nd run also data taken on cooling are shown. Inset: Q−1 measured by DMA in frequency multiplexing on an identical specimen subjected to the first two heating runs.
reversible phenomenon. The inset in Fig. 4, which magnifies the low temperature tail of the DSC traces, reveals that an exothermic process occurs in the same temperature range of the Q−1 peak, and is more pronounced for the n-specimen, as are the Q−1 peak and anomalous (f/f0 )2 trend. The XRD analysis after the 1st run of Fig. 5 (not shown) reveals a modest increase of the MgO content by about 1–2%. In addition, we have experienced that when the n-MgH2 powder is left in air, MgH2 progressively transforms into Mg(OH)2 , and a similar exothermic reaction coupled with MgO formation appears in DSC scans. Therefore we attribute the Q−1 peak at 350 K to the formation of MgO mediated by physisorbed or chemisorbed H2 O at the surface of hydride particles. Since MgO has an elastic modulus much larger than MgH2 , this interpretation also explains the anomalously increasing (f/f0 )2 trend in n-MgH2 where the reaction is enhanced. The second feature, i.e., the (f/f0 )2 drop at high temperature, clearly marks the hydride–metal phase transformation. The first evidence along this direction is given by the strong parallelism between (f/f0 )2 decrease and thermal analysis results (Fig. 4), with n-MgH2 anticipating MgH2 by about 40 K both in DSC peak and mass decrease. This is due to the faster hydrogen desorption kinetics generally shown by nanostructured hydrides with respect to conventional ones [1–3]. The temperature of (f/f0 )2 drop is lower than the onset temperature TON reported in Table 2 for the endothermic DSC peak because VRA runs are carried out in high dynamic vacuum, a condition which favours the hydrogen release, in contrast with the static Ar atmosphere of DSC and HSA runs. The values of the DSC peak area AP and of the mass variation due to hydrogen release m (Table 2), indicate full hydrogen desorption and are consistent with the slightly lower hydride weight fraction in the n-MgH2 sample (Table 1). The complete hydride–metal phase transformation is proved also by XRD profiles of the specimen after one VRA run to 650 K (Fig. 1): the peaks of the hydride completely disappear leaving the characteristics hexagonal Mg reflections with some additional MgO. The large (f/f0 )2 variation between 2nd and 1st run (Table 2) can not be ascribed entirely to the lower elastic modulus of Mg with respect to MgH2 (by about 28% according to calculations [7]). In fact, since the volume mismatch between Mg and MgH2 is about 30%, the transformation is accompanied by an increased porosity which lowers the resonance frequency. The intermediate heating runs on n-MgH2 reported in Fig. 5 were performed also with the aim to study the anelastic behaviour
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L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
of the transforming material, characterized by the coexistence of metal and hydride phases. The 2nd run was stopped at 600 K, after a (f/f0 )2 decrease of about 50%, and the cooling process was monitored. It is worth to point out that the frequency does not rise immediately upon cooling, but falls down until the temperature corresponding to the onset of the (f/f0 )2 drop is reached, because in this temperature interval the hydride–metal transformation continues. At the 3rd run start, hydrogen desorption is not completed: the Q−1 peak height at about 580 K and the (f/f0 )2 drop appear reduced with respect to the 2nd run. The Q−1 peak seems to be of structural – as opposed to relaxational – nature, i.e., it originates with the competition between the increasing Q−1 vs. temperature trend and the Q−1 reduction which evidently accompanies the hydride–metal transformation. This picture is confirmed by a low frequency DMA measurement on a specimen subjected to the same two heating runs (inset of Fig. 5). In fact, an anelastic relaxation peak located at 580 K for f ≈1000 Hz should shift to lower temperature for f = 1 Hz, but no Q−1 peaks are detected below 580 K in the DMA run. In this temperature range and at ambient Ar pressure no hydrogen desorption takes place (Fig. 4), so we can be sure that hydrogen desorption does not affect the anelastic spectrum. This increasing Q−1 vs. temperature trend observed below the desorption temperature is indeed thermally activated, as proved by the temperature shift between the curves at 1 and 20 Hz. Finally, after full hydrogen desorption (Figs. 3 and 4) a relaxation peak with position and activation energy similar to reference Mg is observed. From the temperature shift the activation energy is estimated as 1.5 ± 0.1 eV. This value as well as the position of the peak are in good agreement with results by Delaplace et al. [8] on 99.8% pure magnesium, where a peak with an activation energy of 1.4 eV (equal to the one for volume self-diffusion) was interpreted as a grain boundary peak. The reduced peak height in n-MgH2 after desorption with respect to Mg might be due to the effect of oxidation which hinders interface sliding. The lower temperature of the n-Mg peak can be attributed to the reduced crystallite size, as already observed and suggested in n-Al [9].
5. Conclusions The hydrogen desorption process in magnesium hydride is marked by a dramatic drop of the resonance frequency and, in n-specimen, by a Q−1 peak at 580 K. This peak originates with a competition between the increasing Q−1 vs. temperature trend and the Q−1 reduction which accompanies the hydride–metal transformation. Mechanical spectroscopy also reveals an oxidation reaction at mild temperatures during the first heating run. In the stable state below the hydrogen desorption temperature, the increasing Q−1 vs. temperature trend (higher for n-MgH2 ) is thermally activated. The determination of the corresponding activation energy spectrum and the attribution to a microscopic mechanism requires further investigation. The anelastic spectrum of both n-MgH2 and MgH2 after full transformation to metal resembles the one of commercial Mg, with a broad relaxation peak whose activation energy is close to the one for self-diffusion in Mg. Acknowledgement The financial support in the framework of project FISR-TEPSI from the Italian Ministry of Research (MUR) is gratefully acknowledged. References [1] J. Huot, G. Liang, S. Boily, A. Van Neste, R. Schulz, J. Alloys Compd. 293–295 (1999) 495–500. [2] A. Zaluska, L. Zaluski, J.O. Ström-Olsen, J. Alloys Compd. 288 (1999) 217–225. [3] A. Bassetti, E. Bonetti, L. Pasquini, A. Montone, J. Grbovic, M. Vittori Antisari, Eur. Phys. J. B 45 (2005) 19–27. [4] B.S. Berry, W.C. Pritchet, in: G.E. Murch, H.K. Birnbaum, J.R. Cost (Eds.), Non Traditional Methods in Diffusion, The Metallurgical Society of AIME, 1984, pp. 83–110. [5] H.R. Sinning, Phys. Rev. Lett. 85 (2000) 3201–3204. [6] S. Amadori, E.G. Campari, A.L. Fiorini, R. Montanari, L. Pasquini, L. Savini, E. Bonetti, Mater. Sci. Eng. A 442 (2006) 543–546. [7] R. Yu, P.K. Lam, Phys. Rev. B 37 (1988) 8730–8737. [8] J. Delaplace, J.C. Nicoud, L. Trabut, J. Nucl. Mater. 35 (1970) 167–176. [9] E. Bonetti, L. Pasquini, J. Electron. Mater. 28 (1999) 1055–1061.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Anelasticity in nanostructured MgH2 –Mg: Correlation with hydrogen sorption kinetics L. Pasquini a,∗ , S. Amadori a , E. Bonetti a , E. Callini a , A.L. Fiorini a , A. Montone b , M. Vittori Antisari b a b
Dipartimento di Fisica, Università di Bologna and CNISM, v.le Berti-Pichat 6/2, I-40127 Bologna, Italy Materials and Technology Unit, ENEA C.R. Casaccia, v. Anguillarese 301, I-00060 S. Maria di Galeria, Roma, Italy
a r t i c l e
i n f o
Article history: Received 13 June 2008 Accepted 22 September 2008 Keywords: Magnesium hydride Hydrogen storage Mechanical spectroscopy Nanostructure
a b s t r a c t Mechanical spectroscopy is employed, in correlation with differential scanning calorimetry and hydrogen sorption analysis, to investigate the hydrogen desorption reaction in nanostructured magnesium hydride prepared by ball milling. The hydrogen desorption process is marked by a dramatic drop of the specimen’s resonance frequency and by an internal friction peak that occurs – under high vacuum – at about 580 K. After full hydrogen desorption, a broad anelastic relaxation peak similar to commercial Mg is observed at 470 K, with activation energy close to the one for self-diffusion in Mg. © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental
Metal hydrides that undergo reversible hydrogen sorption reactions are the subject of intensive research activities owing to their application as hydrogen storage materials. In particular, MgH2 is very attractive thanks to its high storage capacity (7.6 wt.%) coupled with the natural abundance and low cost of Mg. However, critical features for the technological exploitation of Mg-based materials are related to the strong Mg–H interaction, reflected by the relatively high desorption temperature, and to the sluggish reaction kinetics due to the poor hydrogen transport both at the surface and in the bulk of magnesium. A possible solution to the slow reaction problem lies in the refinement of the MgH2 crystallite size down to the nanometer range [1–3]. In this article we employ mechanical spectroscopy to investigate hydrogen desorption in nanostructured MgH2 . Mechanical spectroscopy has a long tradition in the study of hydrogen-related processes, such as the short-range hydrogen mobility in amorphous and crystalline materials [4] or the investigation of long-range diffusion in nanocrystalline solids [5]. Here we correlate mechanical spectroscopy with results of thermal analysis techniques to gain more insight into the physics underlying the desorption process.
2.1. Sample preparation
∗ Corresponding author. Tel.: +39 0512095149; fax: +39 0512095113. E-mail address: [email protected] (L. Pasquini). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.09.114
Nanostructured (n-) materials were synthesized by ball milling in a Spex Mixer/Mill 8000 equipped with hardened steel vials and balls (ball-to-powder weight ratio = 10). Commercial magnesium hydride MgH2 from Th. Goldschmidt (60 m particle size, 5 wt.% residual Mg) was employed as starting material for the synthesis of n-MgH2 , while Mg powder from Alfa Aesar (−325 mesh, 99.8% purity) was used as starting material for the preparation of n-Mg. The vials were filled under Ar atmosphere and ball milling was carried out for 10 h. To obtain specimens suitable for anelastic measurements the ball milled powders were compacted in vacuum (1 Pa) at room temperature under a uniaxial pressure of 1 GPa. Reference specimens were also prepared by compacting the starting powders under identical conditions. Four specimen classes are studied in this paper: n-MgH2 , n-Mg, MgH2 , and Mg. 2.2. Characterization The microstructure of the specimens was characterized by Rietveld analysis of X-ray diffraction (XRD) profiles taken with a Rigaku-DMAX IIIC Bragg–Brentano diffractometer, equipped with Cu K␣ radiation and a graphite monochromator in the diffracted beam. Mechanical spectroscopy in the acoustic regime was carried out with a vibrating reed analyzer model VRA 1604 by Cantil [6], operated under high vacuum conditions (10−3 Pa) at a constant heating
152
L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
Table 1 Results of Rietveld analysis on XRD spectra of n-ball milled materials and starting commercial powders. Material
d (nm)
-MgH2 (wt.%)
␥-MgH2 (wt.%)
Mg (wt.%)
MgO (wt.%)
n-MgH2 n-Mg MgH2 Mg
10 (1) 60 (5) >100 >100
62 (2) 0 95 0
30 (2) 0 0 0
1 94 (2) 5 99
7 (2) 6 (2) 0 1
Crystallite size d of the majority phase and relative abundance of the Mg, MgH2 , and MgO phases.
rate of 1 K/min. Throughout this paper, the internal friction Q−1 and normalized squared resonance frequency (f/f0 )2 will be reported, where f0 is the resonance frequency of the as-compacted specimen at ambient temperature. Additional Q−1 data at low frequency were obtained with a dynamic mechanical analyzer (DMA) model DMA 2980 by TA Instruments operating in flowing Ar at a constant heating rate of 1 K/min. The hydrogen desorption process was studied with a hydrogen sorption analyzer (HSA) by Cantil and a differential scanning calorimeter (DSC) model Q10 by TA Instruments. Both instruments were operated under 0.1 MPa Ar pressure in thermal scan mode at 1 K/min heating rate. In the HSA, the pressure variation during hydrogen release is converted into relative mass change m by using a calibrated volume. 3. Results 3.1. X-ray diffraction analysis Table 1 summarizes the microstructure parameters obtained from Rietveld analysis of the XRD profiles shown in Fig. 1. The commercial MgH2 displays narrow peaks pertaining to the tetragonal -MgH2 phase and residual traces of Mg (5 wt.%). The n-MgH2 specimen is characterized by very broad peaks owing to the fine crystallite size of 10 nm. Further, the presence of orthorhombic ␥MgH2 phase (30 wt.% amount) is evident from Fig. 1. This phase can be considered as a distortion of tetragonal -MgH2 with the same packing type and coordination number, and its appearance is a consequence of the large local strain induced by high energy ball milling [3]. The crystallite size of n-Mg is larger (about 60 nm) due to the ductility of Mg which makes the milling process less efficient in comparison with brittle MgH2 . Both n-materials contain about 6–7 wt.% MgO, which in the case of n-MgH2 is due to partial oxidation of the residual Mg in the commercial hydride.
Fig. 1. XRD profiles of investigated specimens, showing also the effects of one heating run on n-MgH2 . The positions of the ␥-MgH2 reflections are shown at the bottom.
3.2. Mechanical spectroscopy and thermal analysis Fig. 2(a) shows the first VRA run on n-MgH2 and MgH2 specimens. Two remarkable features are visible: (i) a Q−1 peak at 350 K for n-MgH2 and 420 K for MgH2 ; (ii) a dramatic drop of (f/f0 )2 which ends at 600 K and 650 K for n-MgH2 and MgH2 respectively. Notably, in n-MgH2 , the Q−1 peak at 350 K couples with an anomalous trend of (f/f0 )2 , while the (f/f0 )2 drop at high temperature is accompanied by a fast decrease of Q−1 to below the value of MgH2 , resulting in a Q−1 peak. In the second run (Fig. 2(b)), the above features disappear: both specimens exhibit a broad Q−1 peak centred at about 470 K and superposed to an increasing background. Fig. 3 reports a series of second runs, carried out by VRA on Mg, n-Mg and n-MgH2 . All specimens display a broad Q−1 peak whose temperature ranges from 440 K of n-Mg to 490 K of Mg. The temperature shift in DMA measures at 1 Hz proves that the peak is a thermally activated relaxation process. Fig. 4 displays DSC and HSA heating runs carried out on the same hydride materials. The hydrogen desorption is marked by the DSC
Fig. 2. First (a) and second (b) VRA heating runs at 1 K/min on hydride specimens.
L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
Fig. 3. Mechanical spectroscopy on Mg, n-Mg and n-MgH2 after full transformation to metal. DMA measures at 1 Hz (open symbols) were carried out in the 300–520 K range only. The other curves are VRA data taken at the indicated resonance frequency.
Fig. 4. DSC and HSA runs at 1 K/min under 0.1 MPa Ar atmosphere. The inset displays a magnified view of the low-temperature DSC traces.
Table 2 Figures of the hydrogen desorption process for the two hydride specimens. Onset temperature TON and integral area AP of the DSC endothermic peak, relative mass change m measured by HSA, and variation of (f/f0 )2 measured at ambient temperature after the first heating run. Material
TON (K)
AP (J/g)
m (wt.%)
(f/f0 )2
n-MgH2 MgH2
624 658
1920 1980
−6.1 −6.6
−0.67 −0.74
endothermic process and by the mass decrease m measured in the HSA. Table 2 reports the main figures of the desorption process together with the (f/f0 )2 variation measured at ambient temperature after the first heating run. 4. Discussion To investigate in more detail the two main features reported in Section 3.2 for the anelastic spectrum of as-compacted hydride specimens, we have carried out intermediate heating runs on nMgH2 . The comparison between the 1st and 2nd run (Fig. 5) proves unambiguously that the Q−1 peak at 350 K is associated with a non-
153
Fig. 5. Successive VRA runs at 1 K/min on n-MgH2 . In the 2nd run also data taken on cooling are shown. Inset: Q−1 measured by DMA in frequency multiplexing on an identical specimen subjected to the first two heating runs.
reversible phenomenon. The inset in Fig. 4, which magnifies the low temperature tail of the DSC traces, reveals that an exothermic process occurs in the same temperature range of the Q−1 peak, and is more pronounced for the n-specimen, as are the Q−1 peak and anomalous (f/f0 )2 trend. The XRD analysis after the 1st run of Fig. 5 (not shown) reveals a modest increase of the MgO content by about 1–2%. In addition, we have experienced that when the n-MgH2 powder is left in air, MgH2 progressively transforms into Mg(OH)2 , and a similar exothermic reaction coupled with MgO formation appears in DSC scans. Therefore we attribute the Q−1 peak at 350 K to the formation of MgO mediated by physisorbed or chemisorbed H2 O at the surface of hydride particles. Since MgO has an elastic modulus much larger than MgH2 , this interpretation also explains the anomalously increasing (f/f0 )2 trend in n-MgH2 where the reaction is enhanced. The second feature, i.e., the (f/f0 )2 drop at high temperature, clearly marks the hydride–metal phase transformation. The first evidence along this direction is given by the strong parallelism between (f/f0 )2 decrease and thermal analysis results (Fig. 4), with n-MgH2 anticipating MgH2 by about 40 K both in DSC peak and mass decrease. This is due to the faster hydrogen desorption kinetics generally shown by nanostructured hydrides with respect to conventional ones [1–3]. The temperature of (f/f0 )2 drop is lower than the onset temperature TON reported in Table 2 for the endothermic DSC peak because VRA runs are carried out in high dynamic vacuum, a condition which favours the hydrogen release, in contrast with the static Ar atmosphere of DSC and HSA runs. The values of the DSC peak area AP and of the mass variation due to hydrogen release m (Table 2), indicate full hydrogen desorption and are consistent with the slightly lower hydride weight fraction in the n-MgH2 sample (Table 1). The complete hydride–metal phase transformation is proved also by XRD profiles of the specimen after one VRA run to 650 K (Fig. 1): the peaks of the hydride completely disappear leaving the characteristics hexagonal Mg reflections with some additional MgO. The large (f/f0 )2 variation between 2nd and 1st run (Table 2) can not be ascribed entirely to the lower elastic modulus of Mg with respect to MgH2 (by about 28% according to calculations [7]). In fact, since the volume mismatch between Mg and MgH2 is about 30%, the transformation is accompanied by an increased porosity which lowers the resonance frequency. The intermediate heating runs on n-MgH2 reported in Fig. 5 were performed also with the aim to study the anelastic behaviour
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L. Pasquini et al. / Materials Science and Engineering A 521–522 (2009) 151–154
of the transforming material, characterized by the coexistence of metal and hydride phases. The 2nd run was stopped at 600 K, after a (f/f0 )2 decrease of about 50%, and the cooling process was monitored. It is worth to point out that the frequency does not rise immediately upon cooling, but falls down until the temperature corresponding to the onset of the (f/f0 )2 drop is reached, because in this temperature interval the hydride–metal transformation continues. At the 3rd run start, hydrogen desorption is not completed: the Q−1 peak height at about 580 K and the (f/f0 )2 drop appear reduced with respect to the 2nd run. The Q−1 peak seems to be of structural – as opposed to relaxational – nature, i.e., it originates with the competition between the increasing Q−1 vs. temperature trend and the Q−1 reduction which evidently accompanies the hydride–metal transformation. This picture is confirmed by a low frequency DMA measurement on a specimen subjected to the same two heating runs (inset of Fig. 5). In fact, an anelastic relaxation peak located at 580 K for f ≈1000 Hz should shift to lower temperature for f = 1 Hz, but no Q−1 peaks are detected below 580 K in the DMA run. In this temperature range and at ambient Ar pressure no hydrogen desorption takes place (Fig. 4), so we can be sure that hydrogen desorption does not affect the anelastic spectrum. This increasing Q−1 vs. temperature trend observed below the desorption temperature is indeed thermally activated, as proved by the temperature shift between the curves at 1 and 20 Hz. Finally, after full hydrogen desorption (Figs. 3 and 4) a relaxation peak with position and activation energy similar to reference Mg is observed. From the temperature shift the activation energy is estimated as 1.5 ± 0.1 eV. This value as well as the position of the peak are in good agreement with results by Delaplace et al. [8] on 99.8% pure magnesium, where a peak with an activation energy of 1.4 eV (equal to the one for volume self-diffusion) was interpreted as a grain boundary peak. The reduced peak height in n-MgH2 after desorption with respect to Mg might be due to the effect of oxidation which hinders interface sliding. The lower temperature of the n-Mg peak can be attributed to the reduced crystallite size, as already observed and suggested in n-Al [9].
5. Conclusions The hydrogen desorption process in magnesium hydride is marked by a dramatic drop of the resonance frequency and, in n-specimen, by a Q−1 peak at 580 K. This peak originates with a competition between the increasing Q−1 vs. temperature trend and the Q−1 reduction which accompanies the hydride–metal transformation. Mechanical spectroscopy also reveals an oxidation reaction at mild temperatures during the first heating run. In the stable state below the hydrogen desorption temperature, the increasing Q−1 vs. temperature trend (higher for n-MgH2 ) is thermally activated. The determination of the corresponding activation energy spectrum and the attribution to a microscopic mechanism requires further investigation. The anelastic spectrum of both n-MgH2 and MgH2 after full transformation to metal resembles the one of commercial Mg, with a broad relaxation peak whose activation energy is close to the one for self-diffusion in Mg. Acknowledgement The financial support in the framework of project FISR-TEPSI from the Italian Ministry of Research (MUR) is gratefully acknowledged. References [1] J. Huot, G. Liang, S. Boily, A. Van Neste, R. Schulz, J. Alloys Compd. 293–295 (1999) 495–500. [2] A. Zaluska, L. Zaluski, J.O. Ström-Olsen, J. Alloys Compd. 288 (1999) 217–225. [3] A. Bassetti, E. Bonetti, L. Pasquini, A. Montone, J. Grbovic, M. Vittori Antisari, Eur. Phys. J. B 45 (2005) 19–27. [4] B.S. Berry, W.C. Pritchet, in: G.E. Murch, H.K. Birnbaum, J.R. Cost (Eds.), Non Traditional Methods in Diffusion, The Metallurgical Society of AIME, 1984, pp. 83–110. [5] H.R. Sinning, Phys. Rev. Lett. 85 (2000) 3201–3204. [6] S. Amadori, E.G. Campari, A.L. Fiorini, R. Montanari, L. Pasquini, L. Savini, E. Bonetti, Mater. Sci. Eng. A 442 (2006) 543–546. [7] R. Yu, P.K. Lam, Phys. Rev. B 37 (1988) 8730–8737. [8] J. Delaplace, J.C. Nicoud, L. Trabut, J. Nucl. Mater. 35 (1970) 167–176. [9] E. Bonetti, L. Pasquini, J. Electron. Mater. 28 (1999) 1055–1061.