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Mechanochemical hydrogenation of ZrCuAlNi based powders
Journal of Alloys and Compounds 454 (2008) 83–88
Mechanochemical hydrogenation of ZrCuAlNi based powders Part I. Structural reconstruction and chemical properties G. Mulas ∗ , L. Schiffini, S. Scudino 1 , G. Cocco Department of Chemistry, University of Sassari, Via Vienna 2, 07100 Sassari, Italy Received 20 October 2006; accepted 20 December 2006 Available online 8 January 2007
Abstract We report on the structural evolution and thermodynamic features of Zr-based amorphous and nanostructured powders subjected to mechanically induced hydrogen absorption. Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions were studied. The process began with a milling induction period, during which neither gaseous uptake nor structural modifications were observed, then, according to an s-shaped kinetics, sorption reached the maximum rate, approaching finally the asymptotic value of 1.5 H/M wt.%. Specific kinetic features were depending on the starting materials. As a result of the absorption, the amorphous alloys of both the compositions promoted the segregation of crystalline ␦-ZrH1.66 phase, which was paralleled by the formation of an amorphous hydride of depleted Zr composition, with relative abundance of the two hydrides 45:55. The different thermodynamic stabilities of the two hydrides were evidenced by calorimetric analyses. While desorption started below 400 K in the metastable one, conversely the crystalline ␦-ZrH1.66 decomposed only above 1070 ◦ C. Faster absorption kinetics and different structural reconstruction path characterised the hydrogenation of the Zr55 Cu30 Al10 Ni5 nanocrystalline alloy. Starting from a mixture of two nanosized fcc phases, with similar (Cu, Al, Ni)Zr2 composition, a single fcc H–(CuAlNi)Zr2 phase was formed, which reversibly desorbed H2 , restoring the two initial phases. © 2007 Elsevier B.V. All rights reserved. Keywords: Energy storage materials; Metal hydrides; Mechanical alloying; Mechanochemical processing; Kinetics
1. Introduction Hydrogen storage alternatives have been the subject of intensive research for many years. Hydride-forming alloys – common hydrides are combination of a metal getter (rare earth elements such as La, or transition metals such as Ti, Zr, and also Mg and Ca), with a non-absorbing metal partner (Fe, Ni, Mn, Co) – have been recognised as a convenient and safe material for reversible hydrogen storage [1,2]. Furthermore, recent works have proved that nanocrystalline hydrides can combine fast kinetics in the absorption–desorption processes with a good storing capability [3,4] and a great interest has been directed towards bi- or multi-component alloys with an highly refined microstructure. Recent studies also exist concerning the hydriding qualities of
∗
Corresponding author. Tel.: +39 079 229524; fax: +39 079 212069. E-mail address: [email protected] (G. Mulas). 1 Present address: IFW Dresden, Institute fur Komplexe Materialien, Postfach 27 01 16, D-01171 Dresden, Germany. 0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2006.12.109
rapidly quenched alloys (RQ), particularly of Zr-based systems with an extended super liquid region and forming bulk metallic glasses. Ismail et al. focused on the micro structural evolution of RQ Zr55 Cu30 Al10 Ni5 and Zr65 Cu12 Al7.5 Ni11 amorphous ribbons during electrolytic hydrogen charging [5,6] and were able to relate the formation of several crystalline hydrides to different interstitial sites available for hydrogen atoms. Koster et al. [7–9] observed that quasicrystal phases – obtained by controlled crystallisation of amorphous RQ ribbons of similar composition – are characterised by absorption–desorption kinetics even faster than the amorphous alloys. Shoji and Inoue focused on the gaseous H2 absorption on amorphous Zr55 Cu30 Al10 Ni5 ribbons and showed that the occurrence of a plateau in the P–C isotherm was due to the formation of an amorphous hydride with the approximate composition (Zr0.55 Al0.10 Ni0.05 Cu0.30 )49 H51 [10]. In addition to a good thermodynamic stability and satisfactory mechanical properties, the interesting hydrogen storage properties observed in Zr-based glassy ribbons prompted us to test the hydrogen absorption characteristics of amorphous and nanostructured powders of a similar composition. As we pointed
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out in a preliminary communication, no absorption phenomena were observed in the as-prepared powders unless mechanically treated under hydrogen atmosphere [11]. Furthermore a significant enhancement was observed in the hydrogen absorption rate, and remarkable differences in the hydriding end products, depending on the intensity of the mechanical treatment [11]. In the attempt to clarify the possible mechanism of the mechanochemical absorbing reactions in these quaternary systems, we studied the overall features of the process in more details. In this work we focus on the progressive structural changes brought about by the hydrogen diffusion under milling and on their correlations with the absorbing properties. A comparison with the absorbing features of nanocrystalline powders was also performed. In a companion paper [12] we report on the effects of the milling intensity, i.e. of the rate of the mechanical energy transfer, on the conversion behaviour. Powders of Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 composition were studied. 2. Experimental We used elemental Zr, Cu, Al and Ni powders—3N purity, 325 mesh from Aldrich. Milling treatments were carried out in a SPEX 8000 Mixer-Mill equipped with a hardened steel vial and ball. An Ar filled glove box was used for vial charging and powder handling (H2 O and O2 level lower than 2 ppm). Nanostructured crystalline samples were obtained from the parent amorphous powders by crystallisation under controlled heating. The reactor vial, manufactured to fit the jaws of the mill clamp, consists of an hardened steel cylinder and two caps supplied with tapered thread male connectors for the hydrogen inlet and pressure monitoring. Filter houses were machined inside the caps for preventing fine powder from obstructing the gas supplying pipes during milling. We followed two hydriding procedures. To study the progressive structural changes brought about by the hydrogen, the charging process was carried out in a stepwise mode, i.e. the vial was charged with 8 g of amorphous powder, pressurised at 0.4 MPa, corresponding to about 0.01 mol of H2 , and subjected to grinding up to pressure equilibration. After the run, a small amount of powder was withdrawn from the vial, the vial was pressurised again and the grinding restarted. The procedure was repeated up to hydrogen saturation. The process kinetics however was studied in a continuous mode under isobaric conditions. To this the vial was connected to an hydrogen reservoir and the hydrogen pressure inside the vial was kept constant by a pressure reducing device valve upstream of the vial. The absorbing process was followed by monitoring the continuous pressure drop in the hydrogen reservoir. Pressure transmitters – PTX 1400, 0–1 MPa operating range, supplied by Druck Limited – or a capacitance sensor 750B from MKS, were employed and data were registered trough a 6034E data acquisition board coupled with a SCXI 1121 signal conditioning system from National Instruments. The temperatures of vial and reservoir were controlled during the whole milling period. The volumes of the different parts of the circuit were measured by helium gas. The same gas was used for seal tests. Otherwise stated, MA and hydriding processes were carried out at constant milling intensity, with the mill running at 875 rpm, with a single ball of 8.3 g at work and with a powder charge of 8 g. For these trials the impact energy was evaluated at around 0.07 J hit−1 , according to a previously developed procedure [13]. The structure of as milled and hydrogenated powders was characterised by Xray diffraction (XRD) using a Siemens D500 diffractometer and Cu K␣ radiation. The thermal stability of the powders in the amorphous state and after hydrogenation treatment, was analysed by differential scanning calorimetry (DSC) using a Perkin-Elmer DSC7 calorimeter, by thermogravimetry and differential thermal analysis using a TG-DTA Setaram Labsys. Heating was carried out at a constant rate of 40 K/min under pure Ar flow. Measurements of absorbed hydrogen content were performed, through hot extraction method, by a ELTRA OH 900 gas analyser (R&D and ELTRA courtesy).
Fig. 1. XRD patterns of mechanically alloyed amorphous Zr55 Cu30 Al10 Ni5 powders and of samples heated up to the quoted temperatures. A diffraction peak due to residual Zr is visible in the as milled sample.
3. Results and discussion 3.1. Structural and thermal characteristics of the amorphous powders The XRD pattern sequences in Figs. 1 and 2 refer to Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions, respectively. The lowest patterns in the figures correspond to the amorphous structures, resulting after about 60 h of mechanical treatment, whereas the upper patterns refer to the powders subjected to progressive heating up to crystallisation. For both the compositions the main features in the patterns agree with previous published data [14]. The 2θ positions of the broad main maxima, with subsequent modulations in the
Fig. 2. XRD patterns of mechanically alloyed and thermally treated Zr69.5 Cu12 Al7.5 Ni11 powders.
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Fig. 3. DSC traces of the Zr55 Cu30 Al10 Ni5 (lower curve) and Zr69.5 Cu12 Al7.5 Ni11 alloys.
scattered intensity, characterise a difference in the short-range order of the two amorphous structures, which is due to the large variation in the composition. Differences also exist in the thermal stability of the two amorphous alloys as evidenced by the crystallisation behaviour and phase transformation under heating [14]. Increasing the zirconium content decreases the temperature of the glass transition, Tg , and the onset of the crystallisation, Tx , as it can be observed in DSC traces quoted in Fig. 3. Furthermore, in the thermogram relevant to the Zr55 Cu30 Al10 Ni5 alloy, lower curve, a sharp crystallisation peak (at T = 776 K) results in the primary crystallisation of a fcc (Cu, Al, Ni)Zr2 phase, cell parameter a = 1.195 nm, NiTi2 type, Fd3m [15], and to the tetragonal (Cu, Al, Ni)Zr3 phase, a = 0.455 nm, c = 0.370 nm, AuCu type, P4/mmm [16]. Two distinct peaks are present in the trace of the Zr69.5 Cu12 Al7.5 Ni11 alloy, the former, Tx1 , relevant to the crystallisation into fcc-NiZr2 (JCPDS 41-898), and the latter, Tx2 , to the formation of tetragonal CuZr2 (JCPDS 38-1170) and hexagonal Zr6 NiAl2 , a = 0.792 nm, c = 0.334 nm, P62m [16]. Supplementing these data, the upper patterns in Figs. 1 and 2 refer to samples subjected to thermal rises up to relevant points of the thermograms and quenched in the DSC cell. The crystallisation process of amorphous quaternary alloys is often influenced by experimental conditions and by contaminant traces, however the phases here observed, including the big-cube NiTi2 type, agree with the literature data relevant to the crystallisation of such amorphous materials [14]. In addition to this, the elemental microanalysis revealed a high homogeneity of the matrix composition with a narrow particle size distribution centred around 15 m.
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Fig. 4. The XRD patterns illustrates the progressive formation of a zirconium hydride together with an amorphous hydride phase for the Zr55 Cu30 Al10 Ni5 alloy.
tural and the thermochemical evolution. Due to the expansion of the amorphous structure, a progressive shift of the main halo toward lower diffraction angles is evident in the XRD pattern sequences shown in Figs. 4 and 5. The main point here, however, is the [1 1 1] reflection of the crystalline ZrH1.66 ␦ phase (JCPDS 34-649), which progressively forms on the rising side of the main halo. The hydride demixes directly from the amorphous matrix and this phase separation is very rapid in the case of Zr-rich alloy. Hydrogen uptake induced powder embrittlement with a consequent decrease in the particles size, which values are now distributed around 3 m.
3.2. The hydriding process: structure characterisation and thermal behaviour of the hydride phases For explorative purposes, the as-prepared amorphous powders were subjected to discontinuous hydrogen charging. After each hydriding step, powders were sampled to check the struc-
Fig. 5. The XRD patterns in the case of Zr69.5 Cu12 Al7.5 Ni11 hydriding process. The [1 1 1] peak of crystalline ZrH1.66 develops from the very beginning of the hydriding process.
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Fig. 6. DCS traces of the progressively hydrogen charged samples. Thermograms refer to Zr55 Cu30 Al10 Ni5 alloy.
The DSC analysis confirms the progressive loss of structural homogeneity in the hydrided samples. The process is reported in detail for the Zr55 Cu30 Al10 Ni5 . As shown in Fig. 6, the original crystallisation single peak splits into two overlapping signals, and only the high-temperature component is observable after further hydrogenation. The residual crystallisation peak broadens and moves towards still higher temperatures. In the DSC trace of the fully charged sample a broad endothermic signal appears between about 390 and 780 K, followed by a relatively neat exothermic peak at around 800 K. The DTA analysis, not reported, confirms and complements these findings showing a further endothermic signal between 900 and 1000 K. The low temperature signals are due to the hydrogen evolution from the residual amorphous phase. It is proved by the shift-back of the amorphous halo toward its original 2θ position, which is observed in the X-ray patterns of fully hydrided samples, subjected to thermal desorption at progressively increased temperatures. The sequence is presented in Fig. 7. One also notes the progressive sharpening of the 1 1 1 peak of the ␦-ZrH1.66 . As for the hydrogen release above 1000 K, the TG analysis, reported in Fig. 8, indicates that about 45% of the total hydrogen uptake relates to the decomposition of the ZrH1.66 crystalline phase. The above analysis circumstantiates the existence of two distinct structural environments for hydrogen, properly related to amorphous and crystalline hydride, respectively. The amorphous structure relaxes continuously under desorption and it indicates that a noticeable powder fraction works reversibly under cyclic processes. Conversely, no hydrogen evolution was observed at low temperature from the stable crystalline hydride. Furthermore, as a consequence of the changed stoichiometry, fcc-AlCu2 Zr (JCPDS 28-18) forms under crystallisation instead of tetragonal (Cu, Al, Ni)Zr3 , observed in the uncharged sample. A reduced fraction of (Cu, Al, Ni)Zr2 is also still visible.
Fig. 7. The XRD intensity distribution around the main amorphous peak for the hydrogen charged Zr55 Cu30 Al10 Ni5 powders subjected to thermal desorption at the quoted temperatures. One notices the shift of the broad maximum approaching the angular position initially observed in the uncharged samples.
A similar hydriding–dehydriding behaviour was observed in the case of Zr69.5 Cu12 Al7.5 Ni11 . The amorphous hydride fraction crystallises into a fcc (instead of tetragonal) Cu–Zr2 and fcc-(Cu, Al, Ni)Zr2 , NiTi2 type, phases – with the missing formation of the Zr6 NiAl2 phase – and crystalline zirconium hydride. Thus also for this composition hydrogen induces the phase separation of a Zr-poor matrix.
Fig. 8. TG analysis of full hydride Zr55 Cu30 Al10 Ni5 powders. The two weight losses at low and high temperatures pertain to hydrogen desorption from the amorphous phase and from the more stable crystalline zirconium hydride, respectively.
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3.3. The kinetics behaviour The absorption process was studied in the course of milling by keeping isobaric hydrogen conditions inside the vial at 0.4 MPa, and recording the isothermal pressure decrease in the hydrogen reservoir. Typical absorption kinetic curves are reported in Fig. 9 for the two amorphous compositions. We also present the absorbing features of a nanocrystalline Zr55 Cu30 Al10 Ni5 sample, which was obtained from the as prepared powders by controlled heating and rapid quenching. Absorption data collection was carried out at the same milling intensity (0.09 W g−1 ). Curves are characterised by an evident sigmoid behaviour. The initial stage corresponds to an activation period during which no H2 absorption was detected. This activation period depends slightly on the composition, as suggested by the stepwise hydriding mode, and mainly on the structural conditions. A rapid increase of the conversion is then observed up to a relatively long stationary stage, where conversion maintains the maximum rate. The absorption rate fades out progressively approaching the absorption limit at around 5.5 × 10−2 to 6.0 × 10−2 mol corresponding to 1.45–1.55 H/M wt.%. The hydrogen diffusion slows down in the ultimate part of the reaction probably due to the reduced availability of low-energy absorption sites in the amorphous structure and the concurrent formation of the crystalline Zr hydride. In the case of nanostrutured samples, the absorption tests still result in an s-shaped kinetic trend, even if the whole process is accomplished more quickly. Less mechanical work is required for equivalent hydrogen uptakes and also the induction period is shorter with respect to those observed in the full amorphous matrix. Some additional information is in need here. At variance with the equilibrium end products presented above, quench treatment led to new compounds. Indeed pattern a shown in Fig. 10, can be fitted by two nanosized fcc phases, again NiTi2 type, with similar composition (Cu, Al, Ni)Zr2 , but with different cell parameters, 1.193 nm and 1.225 nm, respectively. As also assessed by full pattern fitting procedures according to Rietveld methods [17], the grain size of the main crystalline phases were evaluated, respectively, 12.2 nm and 7.0 nm. These findings suggest that the higher conversion rate, 6.9 × 10−6 mol g−1 s−1 , observed for this sample, probably relates to the high diffusivity paths such as the grain boundaries and the free surface
Fig. 10. XRD patterns of Zr55 Cu30 Al10 Ni5 nanostructured powders: (a) as prepared, (b) subjected to hydriding and (c) de-hydriding.
of the more-open nano-structure. However, the structural transformation path, observed for the nanocrystalline powder under absorption, gives possible additional explanations. As shown in Fig. 10, pattern b, hydrogen absorption results into a single homogeneous fcc H–(Cu, Al, Ni)Zr2 alloy, whose cell parameters is 1.236 nm. Neither the ZrH1.66 phase, observed in the hydriding of amorphous samples, nor other crystalline hydride species separates from the initial structure. As discussed above, the formation of ZrH1.66 implies important structural and compositional changes due to the topological environment of hydrogen with four zirconium atoms at the vertexes of tetrahedral sites. It requires an expenditure of an extra ordering energy. Conversely, in the case of the nanocrystalline texture, the hydrogen diffusion is a less expensive process, since involves only the expansion–contraction of the parent lattice. It therefore appears that thermodynamic constraint and structural characteristics concur in shortening the bulk hydriding process of the nanocrystalline alloy. It is also noticeable that cyclic hydrogen de-absorption processes in the nanocrystalline alloy are accomplished through reversible structural sequences, as demonstrated by the transformation fcc-(Cu, Al, Ni)Zr2 → fcc-H–(Cu, Al, Ni)Zr2 → fcc-(Cu, Al, Ni)Zr2 —upper pattern in Fig. 10. At variance, in the amorphous sample the separation of the stable ZrH1.66 from the matrix hinders the mobility of the atomic hydrogen in a large part of the powder. 4. Conclusions
Fig. 9. Hydrogen uptake data as a function of the elapsed milling time for the amorphous and nanocrystalline Zr55 Cu30 Al10 Ni5 and for the amorphous Zr69.5 Cu12 Al7.5 Ni11 compositions.
Results confirm the effectiveness of mechanically alloying and reactive milling techniques to synthesise novel hydride forming alloys and for hydrogen activation, respectively. The Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions studied here are able to absorb hydrogen up to a weight content of
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approximately 1.45–1.55%, close to the amount stored by more conventional techniques. With reference to these, however, reactive milling is more effective in speeding up the reaction due to the unique capability to refine the microstructure and to reduce the grain dimension of the absorbing alloys. Thus, the milling action goes with and matches the progressive embrittlement of the starting material due to the hydrogen uptake. The dynamic role of atomic hydrogen in reconstructing its own short-range order with four zirconium atoms is well documented by the progressive formation of ZrH1.66 ␦ phase. The formation of this stable hydride restricts the ab-desorption capability to the residual amorphous fraction and this is an important drawback for a practical exploitation of the Zr-based materials. Nanocrystalline powders show an enhanced absorbing rate with respect to the original amorphous structure. Possible reason for this was recognised in the lack of a structural reconstruction, which conversely was observed to characterise the absorption process in the amorphous matrix. This structural reconstruction leads to a phase separation with consequent homogeneity loss. The hydriding mechanism under reactive milling is in need of further inquiry. To this, in the parallel communication, attention is paid to work out some relationships relating the intrinsic hydriding kinetics to the intensity of the mechanical treatment. Acknowledgements This work was funded by Ministero della Ricerca Scientifica e Tecnologica, Rome, and University of Sassari. R&D Tecnologie dei Materiali and ELTRA are gratefully acknowledged for performing quantitative determination of hydrogen absorbed.
References [1] H. Wipf (Ed.), Hydrogen in Metals III, Properties and Applications, Springer-Verlag, Berlin Heidelberg, 1997. [2] H. Figiel, O.J. Zogal, V. Yartis (Eds.), Proceedings of the Ninth International Symposium on Metal–Hydrogen Systems MH2004, September 5–10, 2004, Cracow, Poland, J. Alloys Compd. (2005) 404–406. [3] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, Appl. Phys. A 72 (2001) 157. [4] J. Huot, in: H.S. Nalwa (Ed.), Nanoclusters and Nanocrystals, American Scientific Publishers, CA, USA, 2003, pp. 53–85. [5] N. Ismail, M. Uhlemann, A. Gebert, J. Eckert, J. Alloys Compd. 298 (2000) 146. [6] N. Ismail, A. Gebert, M. Uhlemann, J. Eckert, L. Schultz, J. Alloys Compd. 314 (2001) 170. [7] D. Zander, H. Leptien, U. Koster, N. Eliaz, D. Eliezer, J. Non-Crystal. Solids 250–252 (1999) 893. [8] D. Zander, U. Koster, N. Eliaz, D. Eliezer, Mater. Sci. Eng. A 294 (2000) 112. [9] U. Koster, D. Zander, Triwikantoro, J. Metastable Nanocrystal. Mater. 8 (2000) 203. [10] T. Shoji, A. Inoue, J. Alloys Compd. 292 (1999) 275. [11] G. Mulas, S. Scudino, G. Cocco, Mater. Sci. Eng. A 375–377 (2004) 961. [12] G. Mulas, L. Schiffini, S. Scudino, G. Cocco, J. Alloys Compd. (2007), doi:10.1016/j.jallcom.2007.01.025. [13] F. Delogu, M. Monagheddu, G. Mulas, L. Schiffini, G. Cocco, Int. J. NonEquilib. Process. 11 (2000) 235. [14] J. Eckert, M. Seidel, L. Schultz, Mater. Sci. Forum 225–227 (1996) 113. [15] M.H. Mueller, H.W. Knott, Trans. Met. Soc. AIME 227 (1963) 674. [16] P. Villars, L.D. Calvert (Eds.), Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, American Society for Metals, Metals Park, OH, 1985; Karls, J. Inst. Met. 79 (1951) 404. [17] L. Lutterotti, R. Ceccato, R. DalMaschio, E. Pagani, Mater. Sci. Forum 278–281 (1998) 87.
Mechanochemical hydrogenation of ZrCuAlNi based powders Part I. Structural reconstruction and chemical properties G. Mulas ∗ , L. Schiffini, S. Scudino 1 , G. Cocco Department of Chemistry, University of Sassari, Via Vienna 2, 07100 Sassari, Italy Received 20 October 2006; accepted 20 December 2006 Available online 8 January 2007
Abstract We report on the structural evolution and thermodynamic features of Zr-based amorphous and nanostructured powders subjected to mechanically induced hydrogen absorption. Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions were studied. The process began with a milling induction period, during which neither gaseous uptake nor structural modifications were observed, then, according to an s-shaped kinetics, sorption reached the maximum rate, approaching finally the asymptotic value of 1.5 H/M wt.%. Specific kinetic features were depending on the starting materials. As a result of the absorption, the amorphous alloys of both the compositions promoted the segregation of crystalline ␦-ZrH1.66 phase, which was paralleled by the formation of an amorphous hydride of depleted Zr composition, with relative abundance of the two hydrides 45:55. The different thermodynamic stabilities of the two hydrides were evidenced by calorimetric analyses. While desorption started below 400 K in the metastable one, conversely the crystalline ␦-ZrH1.66 decomposed only above 1070 ◦ C. Faster absorption kinetics and different structural reconstruction path characterised the hydrogenation of the Zr55 Cu30 Al10 Ni5 nanocrystalline alloy. Starting from a mixture of two nanosized fcc phases, with similar (Cu, Al, Ni)Zr2 composition, a single fcc H–(CuAlNi)Zr2 phase was formed, which reversibly desorbed H2 , restoring the two initial phases. © 2007 Elsevier B.V. All rights reserved. Keywords: Energy storage materials; Metal hydrides; Mechanical alloying; Mechanochemical processing; Kinetics
1. Introduction Hydrogen storage alternatives have been the subject of intensive research for many years. Hydride-forming alloys – common hydrides are combination of a metal getter (rare earth elements such as La, or transition metals such as Ti, Zr, and also Mg and Ca), with a non-absorbing metal partner (Fe, Ni, Mn, Co) – have been recognised as a convenient and safe material for reversible hydrogen storage [1,2]. Furthermore, recent works have proved that nanocrystalline hydrides can combine fast kinetics in the absorption–desorption processes with a good storing capability [3,4] and a great interest has been directed towards bi- or multi-component alloys with an highly refined microstructure. Recent studies also exist concerning the hydriding qualities of
∗
Corresponding author. Tel.: +39 079 229524; fax: +39 079 212069. E-mail address: [email protected] (G. Mulas). 1 Present address: IFW Dresden, Institute fur Komplexe Materialien, Postfach 27 01 16, D-01171 Dresden, Germany. 0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2006.12.109
rapidly quenched alloys (RQ), particularly of Zr-based systems with an extended super liquid region and forming bulk metallic glasses. Ismail et al. focused on the micro structural evolution of RQ Zr55 Cu30 Al10 Ni5 and Zr65 Cu12 Al7.5 Ni11 amorphous ribbons during electrolytic hydrogen charging [5,6] and were able to relate the formation of several crystalline hydrides to different interstitial sites available for hydrogen atoms. Koster et al. [7–9] observed that quasicrystal phases – obtained by controlled crystallisation of amorphous RQ ribbons of similar composition – are characterised by absorption–desorption kinetics even faster than the amorphous alloys. Shoji and Inoue focused on the gaseous H2 absorption on amorphous Zr55 Cu30 Al10 Ni5 ribbons and showed that the occurrence of a plateau in the P–C isotherm was due to the formation of an amorphous hydride with the approximate composition (Zr0.55 Al0.10 Ni0.05 Cu0.30 )49 H51 [10]. In addition to a good thermodynamic stability and satisfactory mechanical properties, the interesting hydrogen storage properties observed in Zr-based glassy ribbons prompted us to test the hydrogen absorption characteristics of amorphous and nanostructured powders of a similar composition. As we pointed
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out in a preliminary communication, no absorption phenomena were observed in the as-prepared powders unless mechanically treated under hydrogen atmosphere [11]. Furthermore a significant enhancement was observed in the hydrogen absorption rate, and remarkable differences in the hydriding end products, depending on the intensity of the mechanical treatment [11]. In the attempt to clarify the possible mechanism of the mechanochemical absorbing reactions in these quaternary systems, we studied the overall features of the process in more details. In this work we focus on the progressive structural changes brought about by the hydrogen diffusion under milling and on their correlations with the absorbing properties. A comparison with the absorbing features of nanocrystalline powders was also performed. In a companion paper [12] we report on the effects of the milling intensity, i.e. of the rate of the mechanical energy transfer, on the conversion behaviour. Powders of Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 composition were studied. 2. Experimental We used elemental Zr, Cu, Al and Ni powders—3N purity, 325 mesh from Aldrich. Milling treatments were carried out in a SPEX 8000 Mixer-Mill equipped with a hardened steel vial and ball. An Ar filled glove box was used for vial charging and powder handling (H2 O and O2 level lower than 2 ppm). Nanostructured crystalline samples were obtained from the parent amorphous powders by crystallisation under controlled heating. The reactor vial, manufactured to fit the jaws of the mill clamp, consists of an hardened steel cylinder and two caps supplied with tapered thread male connectors for the hydrogen inlet and pressure monitoring. Filter houses were machined inside the caps for preventing fine powder from obstructing the gas supplying pipes during milling. We followed two hydriding procedures. To study the progressive structural changes brought about by the hydrogen, the charging process was carried out in a stepwise mode, i.e. the vial was charged with 8 g of amorphous powder, pressurised at 0.4 MPa, corresponding to about 0.01 mol of H2 , and subjected to grinding up to pressure equilibration. After the run, a small amount of powder was withdrawn from the vial, the vial was pressurised again and the grinding restarted. The procedure was repeated up to hydrogen saturation. The process kinetics however was studied in a continuous mode under isobaric conditions. To this the vial was connected to an hydrogen reservoir and the hydrogen pressure inside the vial was kept constant by a pressure reducing device valve upstream of the vial. The absorbing process was followed by monitoring the continuous pressure drop in the hydrogen reservoir. Pressure transmitters – PTX 1400, 0–1 MPa operating range, supplied by Druck Limited – or a capacitance sensor 750B from MKS, were employed and data were registered trough a 6034E data acquisition board coupled with a SCXI 1121 signal conditioning system from National Instruments. The temperatures of vial and reservoir were controlled during the whole milling period. The volumes of the different parts of the circuit were measured by helium gas. The same gas was used for seal tests. Otherwise stated, MA and hydriding processes were carried out at constant milling intensity, with the mill running at 875 rpm, with a single ball of 8.3 g at work and with a powder charge of 8 g. For these trials the impact energy was evaluated at around 0.07 J hit−1 , according to a previously developed procedure [13]. The structure of as milled and hydrogenated powders was characterised by Xray diffraction (XRD) using a Siemens D500 diffractometer and Cu K␣ radiation. The thermal stability of the powders in the amorphous state and after hydrogenation treatment, was analysed by differential scanning calorimetry (DSC) using a Perkin-Elmer DSC7 calorimeter, by thermogravimetry and differential thermal analysis using a TG-DTA Setaram Labsys. Heating was carried out at a constant rate of 40 K/min under pure Ar flow. Measurements of absorbed hydrogen content were performed, through hot extraction method, by a ELTRA OH 900 gas analyser (R&D and ELTRA courtesy).
Fig. 1. XRD patterns of mechanically alloyed amorphous Zr55 Cu30 Al10 Ni5 powders and of samples heated up to the quoted temperatures. A diffraction peak due to residual Zr is visible in the as milled sample.
3. Results and discussion 3.1. Structural and thermal characteristics of the amorphous powders The XRD pattern sequences in Figs. 1 and 2 refer to Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions, respectively. The lowest patterns in the figures correspond to the amorphous structures, resulting after about 60 h of mechanical treatment, whereas the upper patterns refer to the powders subjected to progressive heating up to crystallisation. For both the compositions the main features in the patterns agree with previous published data [14]. The 2θ positions of the broad main maxima, with subsequent modulations in the
Fig. 2. XRD patterns of mechanically alloyed and thermally treated Zr69.5 Cu12 Al7.5 Ni11 powders.
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Fig. 3. DSC traces of the Zr55 Cu30 Al10 Ni5 (lower curve) and Zr69.5 Cu12 Al7.5 Ni11 alloys.
scattered intensity, characterise a difference in the short-range order of the two amorphous structures, which is due to the large variation in the composition. Differences also exist in the thermal stability of the two amorphous alloys as evidenced by the crystallisation behaviour and phase transformation under heating [14]. Increasing the zirconium content decreases the temperature of the glass transition, Tg , and the onset of the crystallisation, Tx , as it can be observed in DSC traces quoted in Fig. 3. Furthermore, in the thermogram relevant to the Zr55 Cu30 Al10 Ni5 alloy, lower curve, a sharp crystallisation peak (at T = 776 K) results in the primary crystallisation of a fcc (Cu, Al, Ni)Zr2 phase, cell parameter a = 1.195 nm, NiTi2 type, Fd3m [15], and to the tetragonal (Cu, Al, Ni)Zr3 phase, a = 0.455 nm, c = 0.370 nm, AuCu type, P4/mmm [16]. Two distinct peaks are present in the trace of the Zr69.5 Cu12 Al7.5 Ni11 alloy, the former, Tx1 , relevant to the crystallisation into fcc-NiZr2 (JCPDS 41-898), and the latter, Tx2 , to the formation of tetragonal CuZr2 (JCPDS 38-1170) and hexagonal Zr6 NiAl2 , a = 0.792 nm, c = 0.334 nm, P62m [16]. Supplementing these data, the upper patterns in Figs. 1 and 2 refer to samples subjected to thermal rises up to relevant points of the thermograms and quenched in the DSC cell. The crystallisation process of amorphous quaternary alloys is often influenced by experimental conditions and by contaminant traces, however the phases here observed, including the big-cube NiTi2 type, agree with the literature data relevant to the crystallisation of such amorphous materials [14]. In addition to this, the elemental microanalysis revealed a high homogeneity of the matrix composition with a narrow particle size distribution centred around 15 m.
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Fig. 4. The XRD patterns illustrates the progressive formation of a zirconium hydride together with an amorphous hydride phase for the Zr55 Cu30 Al10 Ni5 alloy.
tural and the thermochemical evolution. Due to the expansion of the amorphous structure, a progressive shift of the main halo toward lower diffraction angles is evident in the XRD pattern sequences shown in Figs. 4 and 5. The main point here, however, is the [1 1 1] reflection of the crystalline ZrH1.66 ␦ phase (JCPDS 34-649), which progressively forms on the rising side of the main halo. The hydride demixes directly from the amorphous matrix and this phase separation is very rapid in the case of Zr-rich alloy. Hydrogen uptake induced powder embrittlement with a consequent decrease in the particles size, which values are now distributed around 3 m.
3.2. The hydriding process: structure characterisation and thermal behaviour of the hydride phases For explorative purposes, the as-prepared amorphous powders were subjected to discontinuous hydrogen charging. After each hydriding step, powders were sampled to check the struc-
Fig. 5. The XRD patterns in the case of Zr69.5 Cu12 Al7.5 Ni11 hydriding process. The [1 1 1] peak of crystalline ZrH1.66 develops from the very beginning of the hydriding process.
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Fig. 6. DCS traces of the progressively hydrogen charged samples. Thermograms refer to Zr55 Cu30 Al10 Ni5 alloy.
The DSC analysis confirms the progressive loss of structural homogeneity in the hydrided samples. The process is reported in detail for the Zr55 Cu30 Al10 Ni5 . As shown in Fig. 6, the original crystallisation single peak splits into two overlapping signals, and only the high-temperature component is observable after further hydrogenation. The residual crystallisation peak broadens and moves towards still higher temperatures. In the DSC trace of the fully charged sample a broad endothermic signal appears between about 390 and 780 K, followed by a relatively neat exothermic peak at around 800 K. The DTA analysis, not reported, confirms and complements these findings showing a further endothermic signal between 900 and 1000 K. The low temperature signals are due to the hydrogen evolution from the residual amorphous phase. It is proved by the shift-back of the amorphous halo toward its original 2θ position, which is observed in the X-ray patterns of fully hydrided samples, subjected to thermal desorption at progressively increased temperatures. The sequence is presented in Fig. 7. One also notes the progressive sharpening of the 1 1 1 peak of the ␦-ZrH1.66 . As for the hydrogen release above 1000 K, the TG analysis, reported in Fig. 8, indicates that about 45% of the total hydrogen uptake relates to the decomposition of the ZrH1.66 crystalline phase. The above analysis circumstantiates the existence of two distinct structural environments for hydrogen, properly related to amorphous and crystalline hydride, respectively. The amorphous structure relaxes continuously under desorption and it indicates that a noticeable powder fraction works reversibly under cyclic processes. Conversely, no hydrogen evolution was observed at low temperature from the stable crystalline hydride. Furthermore, as a consequence of the changed stoichiometry, fcc-AlCu2 Zr (JCPDS 28-18) forms under crystallisation instead of tetragonal (Cu, Al, Ni)Zr3 , observed in the uncharged sample. A reduced fraction of (Cu, Al, Ni)Zr2 is also still visible.
Fig. 7. The XRD intensity distribution around the main amorphous peak for the hydrogen charged Zr55 Cu30 Al10 Ni5 powders subjected to thermal desorption at the quoted temperatures. One notices the shift of the broad maximum approaching the angular position initially observed in the uncharged samples.
A similar hydriding–dehydriding behaviour was observed in the case of Zr69.5 Cu12 Al7.5 Ni11 . The amorphous hydride fraction crystallises into a fcc (instead of tetragonal) Cu–Zr2 and fcc-(Cu, Al, Ni)Zr2 , NiTi2 type, phases – with the missing formation of the Zr6 NiAl2 phase – and crystalline zirconium hydride. Thus also for this composition hydrogen induces the phase separation of a Zr-poor matrix.
Fig. 8. TG analysis of full hydride Zr55 Cu30 Al10 Ni5 powders. The two weight losses at low and high temperatures pertain to hydrogen desorption from the amorphous phase and from the more stable crystalline zirconium hydride, respectively.
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3.3. The kinetics behaviour The absorption process was studied in the course of milling by keeping isobaric hydrogen conditions inside the vial at 0.4 MPa, and recording the isothermal pressure decrease in the hydrogen reservoir. Typical absorption kinetic curves are reported in Fig. 9 for the two amorphous compositions. We also present the absorbing features of a nanocrystalline Zr55 Cu30 Al10 Ni5 sample, which was obtained from the as prepared powders by controlled heating and rapid quenching. Absorption data collection was carried out at the same milling intensity (0.09 W g−1 ). Curves are characterised by an evident sigmoid behaviour. The initial stage corresponds to an activation period during which no H2 absorption was detected. This activation period depends slightly on the composition, as suggested by the stepwise hydriding mode, and mainly on the structural conditions. A rapid increase of the conversion is then observed up to a relatively long stationary stage, where conversion maintains the maximum rate. The absorption rate fades out progressively approaching the absorption limit at around 5.5 × 10−2 to 6.0 × 10−2 mol corresponding to 1.45–1.55 H/M wt.%. The hydrogen diffusion slows down in the ultimate part of the reaction probably due to the reduced availability of low-energy absorption sites in the amorphous structure and the concurrent formation of the crystalline Zr hydride. In the case of nanostrutured samples, the absorption tests still result in an s-shaped kinetic trend, even if the whole process is accomplished more quickly. Less mechanical work is required for equivalent hydrogen uptakes and also the induction period is shorter with respect to those observed in the full amorphous matrix. Some additional information is in need here. At variance with the equilibrium end products presented above, quench treatment led to new compounds. Indeed pattern a shown in Fig. 10, can be fitted by two nanosized fcc phases, again NiTi2 type, with similar composition (Cu, Al, Ni)Zr2 , but with different cell parameters, 1.193 nm and 1.225 nm, respectively. As also assessed by full pattern fitting procedures according to Rietveld methods [17], the grain size of the main crystalline phases were evaluated, respectively, 12.2 nm and 7.0 nm. These findings suggest that the higher conversion rate, 6.9 × 10−6 mol g−1 s−1 , observed for this sample, probably relates to the high diffusivity paths such as the grain boundaries and the free surface
Fig. 10. XRD patterns of Zr55 Cu30 Al10 Ni5 nanostructured powders: (a) as prepared, (b) subjected to hydriding and (c) de-hydriding.
of the more-open nano-structure. However, the structural transformation path, observed for the nanocrystalline powder under absorption, gives possible additional explanations. As shown in Fig. 10, pattern b, hydrogen absorption results into a single homogeneous fcc H–(Cu, Al, Ni)Zr2 alloy, whose cell parameters is 1.236 nm. Neither the ZrH1.66 phase, observed in the hydriding of amorphous samples, nor other crystalline hydride species separates from the initial structure. As discussed above, the formation of ZrH1.66 implies important structural and compositional changes due to the topological environment of hydrogen with four zirconium atoms at the vertexes of tetrahedral sites. It requires an expenditure of an extra ordering energy. Conversely, in the case of the nanocrystalline texture, the hydrogen diffusion is a less expensive process, since involves only the expansion–contraction of the parent lattice. It therefore appears that thermodynamic constraint and structural characteristics concur in shortening the bulk hydriding process of the nanocrystalline alloy. It is also noticeable that cyclic hydrogen de-absorption processes in the nanocrystalline alloy are accomplished through reversible structural sequences, as demonstrated by the transformation fcc-(Cu, Al, Ni)Zr2 → fcc-H–(Cu, Al, Ni)Zr2 → fcc-(Cu, Al, Ni)Zr2 —upper pattern in Fig. 10. At variance, in the amorphous sample the separation of the stable ZrH1.66 from the matrix hinders the mobility of the atomic hydrogen in a large part of the powder. 4. Conclusions
Fig. 9. Hydrogen uptake data as a function of the elapsed milling time for the amorphous and nanocrystalline Zr55 Cu30 Al10 Ni5 and for the amorphous Zr69.5 Cu12 Al7.5 Ni11 compositions.
Results confirm the effectiveness of mechanically alloying and reactive milling techniques to synthesise novel hydride forming alloys and for hydrogen activation, respectively. The Zr55 Cu30 Al10 Ni5 and Zr69.5 Cu12 Al7.5 Ni11 compositions studied here are able to absorb hydrogen up to a weight content of
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approximately 1.45–1.55%, close to the amount stored by more conventional techniques. With reference to these, however, reactive milling is more effective in speeding up the reaction due to the unique capability to refine the microstructure and to reduce the grain dimension of the absorbing alloys. Thus, the milling action goes with and matches the progressive embrittlement of the starting material due to the hydrogen uptake. The dynamic role of atomic hydrogen in reconstructing its own short-range order with four zirconium atoms is well documented by the progressive formation of ZrH1.66 ␦ phase. The formation of this stable hydride restricts the ab-desorption capability to the residual amorphous fraction and this is an important drawback for a practical exploitation of the Zr-based materials. Nanocrystalline powders show an enhanced absorbing rate with respect to the original amorphous structure. Possible reason for this was recognised in the lack of a structural reconstruction, which conversely was observed to characterise the absorption process in the amorphous matrix. This structural reconstruction leads to a phase separation with consequent homogeneity loss. The hydriding mechanism under reactive milling is in need of further inquiry. To this, in the parallel communication, attention is paid to work out some relationships relating the intrinsic hydriding kinetics to the intensity of the mechanical treatment. Acknowledgements This work was funded by Ministero della Ricerca Scientifica e Tecnologica, Rome, and University of Sassari. R&D Tecnologie dei Materiali and ELTRA are gratefully acknowledged for performing quantitative determination of hydrogen absorbed.
References [1] H. Wipf (Ed.), Hydrogen in Metals III, Properties and Applications, Springer-Verlag, Berlin Heidelberg, 1997. [2] H. Figiel, O.J. Zogal, V. Yartis (Eds.), Proceedings of the Ninth International Symposium on Metal–Hydrogen Systems MH2004, September 5–10, 2004, Cracow, Poland, J. Alloys Compd. (2005) 404–406. [3] A. Zaluska, L. Zaluski, J.O. Strom-Olsen, Appl. Phys. A 72 (2001) 157. [4] J. Huot, in: H.S. Nalwa (Ed.), Nanoclusters and Nanocrystals, American Scientific Publishers, CA, USA, 2003, pp. 53–85. [5] N. Ismail, M. Uhlemann, A. Gebert, J. Eckert, J. Alloys Compd. 298 (2000) 146. [6] N. Ismail, A. Gebert, M. Uhlemann, J. Eckert, L. Schultz, J. Alloys Compd. 314 (2001) 170. [7] D. Zander, H. Leptien, U. Koster, N. Eliaz, D. Eliezer, J. Non-Crystal. Solids 250–252 (1999) 893. [8] D. Zander, U. Koster, N. Eliaz, D. Eliezer, Mater. Sci. Eng. A 294 (2000) 112. [9] U. Koster, D. Zander, Triwikantoro, J. Metastable Nanocrystal. Mater. 8 (2000) 203. [10] T. Shoji, A. Inoue, J. Alloys Compd. 292 (1999) 275. [11] G. Mulas, S. Scudino, G. Cocco, Mater. Sci. Eng. A 375–377 (2004) 961. [12] G. Mulas, L. Schiffini, S. Scudino, G. Cocco, J. Alloys Compd. (2007), doi:10.1016/j.jallcom.2007.01.025. [13] F. Delogu, M. Monagheddu, G. Mulas, L. Schiffini, G. Cocco, Int. J. NonEquilib. Process. 11 (2000) 235. [14] J. Eckert, M. Seidel, L. Schultz, Mater. Sci. Forum 225–227 (1996) 113. [15] M.H. Mueller, H.W. Knott, Trans. Met. Soc. AIME 227 (1963) 674. [16] P. Villars, L.D. Calvert (Eds.), Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, American Society for Metals, Metals Park, OH, 1985; Karls, J. Inst. Met. 79 (1951) 404. [17] L. Lutterotti, R. Ceccato, R. DalMaschio, E. Pagani, Mater. Sci. Forum 278–281 (1998) 87.