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Journal of Nanomaterials
Volume 2012 (2012), Article ID 592147, 8 pages
Preparation and Hydrogen Storage Properties of Mg-Rich Mg-Ni Ultrafine Particles
1Shanghai Engineering Research Center of Magnesium Materials and Applications and National Engineering Research Center of Light Alloy Net Forming, Shanghai Jiao Tong University, Shanghai 200240, China
2State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
3Unité Matériaux Et Transformations (UMET), CNRS UMR 8207, Université Lille 1, 59655 Villeneuve d'Ascq, France
Received 29 August 2012; Accepted 16 October 2012
Academic Editor: Kemin Zhang
Copyright © 2012 Jianxin Zou et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.
In the present work, Mg-rich Mg-Ni ultrafine powders were prepared through an arc plasma method. The phase components, microstructure, and hydrogen storage properties of the powders were carefully investigated. It is found that Mg2Ni and MgNi2 could be obtained directly from the vapor state reactions between Mg and Ni, depending on the local vapor content in the reaction chamber. A nanostructured MgH2 + Mg2NiH4 hydrogen storage composite could be generated after hydrogenation of the Mg-Ni ultrafine powders. After dehydrogenation, MgH2 and Mg2NiH4 decomposed into nanograined Mg and Mg2Ni, respectively. Thermogravimetry/differential scanning calorimetry (TG/DSC) analyses showed that Mg2NiH4 phase may play a catalytic role in the dehydriding process of the hydrogenated Mg ultrafine particles.
Mg-based alloys and composites are believed to be promising candidates as hydrogen storage carriers due to their high hydrogen storage capacity (7.6 wt% for MgH2 and 3.6 wt% for Mg2NiH4), great abundance, and low cost . However, the sluggish hydrogen sorption kinetics and high operating temperature of Mg (>573 K) limit its practical applications. Different methods, such as alloying, catalyst addition, and nanocrystallization, have been used to improve the hydrogen thermodynamics and sorption kinetics of Mg [2–4]. Consequently, various Mg-based alloys, composites, and compounds have been developed with superior hydrogen sorption kinetics over pure Mg. Among those Mg-based materials, Mg2Ni intermetallic compound is well known to form Mg2NiH4 hydride with high reaction rate, which has a hydrogen storage capacity of 3.6 wt% . However, it is difficult to obtain Mg alloys with accurately desirable compositions by conventional melt-cast methods due to following reasons: Mg has low melting point (923 K) and high vapor pressure (133 Pa at 878 K); Mg is quite active in the presence of oxygen. Especially in the case of Mg-Ni binary system, casting is quite difficult because of the large difference in melting point between Ni (1728 K) and Mg (923 K). Furthermore, based on the binary phase diagram [5, 6], single phase Mg2Ni cannot be simply obtained by casting as a phase separation occurs during solidification.
Mechanical alloying (MA) method has been widely used to produce nanostructured alloys or composites, especially to produce amorphous and nanocrystalline Mg2Ni alloys. However, it usually takes a long time to prepare alloys by the MA method, the sample can be contaminated during the milling process, oxidation occurs even in a protective atmosphere . Moreover, it is difficult to obtain a homogenous sample by the MA method . It is known that nano- or ultrafine particles have a lowered melting point because of the small size and large specific surface area of the particles, which implies a possibility of preparing the Mg2Ni compound by using magnesium and nickel nanoparticles. Lio et al. synthesized Mg2Ni intermetallic nanoparticles by a two-step method . First, Mg and Ni nanoparticles were prepared by hydrogen plasma-metal reaction. Second, the mixture of Mg and Ni nanoparticles in a 2 : 1 molar ratio was compressed into a pellet and was heated up to 623 K under a hydrogen pressure of about 40 bar to react with hydrogen. It is established that at 553 K and 3 MPa hydrogen pressure, Mg2NiH4 can be generated, and Mg2Ni intermetallic phase was obtained with high productivity after hydrogen desorption. The obtained Mg2Ni compound has excellent hydrogen storage properties and can absorb hydrogen to its maximum capacity in the first cycle at a high speed without any activation. Its excellent performance is due to the ultrafine particle size, large surface area, and internal defects.
In this paper, we have attempted the direct preparation of Mg-Ni composite ultrafine particles through vapor state reactions by using a DC arc plasma method. This process is similar to the in-flight plasma processes, which injected the metal powders into the plasma flame leaving the gasification. In this process, the metal vapor cooled down rapidly after leaving the plasma flame, reached saturation, then condensed and formed nano- or ultrafine powders .
2.1. Sample Preparation
The Mg-Ni ultrafine powders were prepared using an arc plasma evaporation apparatus. Figure 1 shows a schematic illustration of the experimental equipment. It mainly consists of a reaction chamber and a collection room. Mg-Ni ultrafine particles were prepared by arc evaporation of the mixture of magnesium and nickel powders. Commercially available Mg and Ni powders with 99.9 wt% in purity and 100 μm in particle size were delicately mixed in a molar ratio of 1 : 2. The mixed powders were then compressed at room temperature to form cylinders with 10 mm in diameter and 7 mm in height by an uniaxial compressor under a pressure of 25 MPa. These cylinders, as anode materials, were put into the reaction chamber filled with mixed 0.75 atm Ar + 0.05 atm H2 gas after the chamber was evacuated to 5 × 10−2 Pa. The DC current is set at 140 A during arc evaporation. Before the Mg-Ni ultrafine particles were taken out from the reaction chamber, they were slowly passivated with a mixture of argon and air to prevent the particles from burning in air.
The phase identification of the Mg-Ni nanoparticles before and after hydrogen absorption was carried out by X-ray diffraction (XRD) using a D/max 2550VL/PCX apparatus equipped with Cu-Kα radiation source. The morphology and microstructure of the powders were observed by using a JEM-2100 transmission electron microscope (TEM), operated at 200 kV. The composition of the Mg-Ni powders was analyzed by inductive coupled plasma emission spectrometer (ICP). A conventional Sieviet-type pressure-composition-temperature (P-C-T) apparatus, which means measuring hydrogen content versus pressure by recording the change of gas pressure in a constant volume, was used to test the hydrogen sorption properties of the Mg-Ni ultrafine powders. The testing temperatures were set at 400°C, 375°C, 350°C and 325°C. The hydrogen desorption property of the prepared powders was analyzed by thermogravimetry/differential scanning calorimetry spectroscopy (TG/DSC) technique.
3. Results and Discussions
3.1. XRD Analysis
Figures 2(a), 2(b), and 2(c) show the XRD patterns of the as- prepared Mg-Ni ultrafine particles, the samples obtained after the hydrogen absorption and after desorption, respectively. From Figure 2(a), it is seen that the main phases of the Mg-Ni ultrafine particles prepared by DC arc plasma are Mg and Ni, with a small amount of MgO. Besides, the presence of Mg2Ni and MgNi2 phases is also detected. The peaks corresponding to MgNi2 phase are fairly weak, indicating that the volume content of Mg2Ni in the powders is very low. The MgO phase was formed when the Mg-Ni nanoparticles were passivated in the mixture of argon and air. The MgO layer at the surface of the magnesium particles can effectively prevent the further oxidation or even burning of the Mg particles when exposed to air. The existence of Mg2Ni and MgNi2 shows that the following reactions have occurred during the arc plasma evaporation: These reactions occurred in vapor state depending on the local vapor concentrations of Mg and Ni. Indeed, Mg has much higher vapor pressure than Ni at the same temperature, resulting in the generation of a vapor rich in Mg during arc evaporation. ICP analysis shows that the average Mg and Ni content in the Mg-Ni ultrafine powders is 7 : 1 in molar ratio, much higher than the ratio 1 : 2 in the original mixture. However, the local vapor content may not be homogeneous during arc evaporation of the Mg-Ni mixture. Ni concentration might be higher than Mg in local regions inside the plasma. As a result, both reaction (1) and reaction (2) could occur in the mixed vapor, leading to the formation of Mg2Ni and MgNi2 phases. As the vapor is highly rich in Mg, reaction (1) should be more important than reaction (2). As a result, the amount of Mg2Ni should be higher than the amount of MgNi2. The fact that the majority phases in the Mg-Ni ultrafine powders are pure Mg and Ni indicates that most of Mg and Ni vapors do not react with each other during evaporation. It has been established that the evaporation rate of metals plays a major role for the generation of the metal particles. According to the Ohno’s model, the vapor generation rate of a metal through hydrogen plasma reaction method is proximately proportional to its reaction parameter, , which can be expressed by the following equation: where is the reaction enthalpy between metal and hydrogen, is the heat of vaporization of the metal at the temperature , and () and (273) are densities of the hydrogen molecules in the metal at temperature and 273 K. By inputting parameters for Mg and Ni into (3), it is estimated that for Mg and for Ni. The large difference in values indicates that the evaporation rate for Ni is much lower than that for Mg. This is consistent with the ICP result for which the Ni content in the composite is lower than that in the mixed powders before arc evaporation.
From the XRD pattern in Figure 2(b), it can be seen that most of the Mg-Ni powders have transformed into MgH2 and Mg2NiH4 hydrides after the hydrogen absorption at 400°C and 4 MPa. The mechanism for the formation of Mg2NiH4 hydride is thought to be the reaction (4) followed by reaction (5) occurred during the isothermal period: As the atomic ratio of Mg : Ni is 7 : 1 in the Mg-Ni ultrafine powders, the reactions (4) and (5) can be combined together into reaction (6). That is, when Mg-Ni ultrafine powders are completely hydrogenated, the molar ratio of MgH2 to Mg2NiH4 should be around 5 : 1: There is still a small amount of Mg phase observed in the sample after hydrogenation at 400°C, as seen in Figure 2(b). This is due to the fact that some relatively large Mg particles cannot be completely hydrogenated. It is observed from Figure 2(c) that after hydrogen desorption, MgH2 decomposes completely and transforms into Mg, and Mg2NiH4 transforms into Mg2Ni compound, as described by the following reaction: This principle of the reactions (4), (5), and (7) is used to prepare Mg2Ni compound by Shao et al. [12, 13]. Comparing Figures 2(a) and 2(b), it is worth noting that the diffraction peaks of MgNi2 disappeared after hydrogen absorption. It is believed that MgNi2 should react with MgH2 to form Mg2NiH4 in the presence of high temperature and hydrogen atmosphere, as described by the following reaction:
3.2. TEM Observations of the Mg-Ni Ultrafine Particles
Figures 3(a), 3(b), and 3(c) show typical bright field TEM micrographs of Mg-Ni composite powders before hydrogen absorption, the samples obtained after hydrogen absorption and after hydrogen desorption, respectively, with corresponding selected area electron diffraction (SAED) patterns. From Figure 3(a), it is observed that most of the particles have quasispherical shape with the particle size ranges from 50 to 600 nm. The corresponding SAED pattern consists of rings and overlapped dispersed spots. The rings are identified to be Mg and MgO phases. However, the spots are not indexed and could correspond to the Ni, Mg2Ni, and MgNi2 phases based on the result of XRD analysis (Figure 2(a)). Normally, Mg can be directly evaporated into gaseous atoms before melting in the high temperature plasma. This is due to the low melting point and small latent heat of evaporation for Mg. The fast vapor generation rate of the Mg results in the formation of relatively large particles . After hydrogenation at 400°C in 4 MPa hydrogen atmosphere, MgH2 and Mg2NiH4 phases formed in the composite. The typical morphology of these hydrides is shown in Figure 3(b). The particles became transparent having size in the range between 50 and 200 nm, smaller than the Mg particles shown in Figure 3(a). In the corresponding SAED pattern, diffracted rings belonging to MgH2, Mg2NiH4, MgO, and Mg phases can be detected. It is found from Figure 3(c) that after 4 hydrogen absorption and desorption cycles, those large Mg particles have been cracked and broken into smaller particles. The corresponding SAED pattern is mainly composed of well-defined rings belonging to pure Mg. Besides this, MgO and Mg2Ni are also identified. Compared with that of the as-prepared particles obtained using the same SAED aperture (Figure 3(a)), it indicates the grain refinement of Mg particles after hydrogen sorption cycles. The grain refinement of Mg particles is expected to benefit from the hydrogen storage properties owing to the increased surface areas and reduced diffusion length for hydrogen. As a consequence of the reduction in particle size, fresh surface was generated during the hydriding/dehydriding cycles. The presence of new surfaces makes the Mg particles easier to be oxidized. As a result, the amount of MgO increases after hydrogen sorption cycles, as can be observed from the comparison between diffraction peaks of MgO phase in Figures 2(c) and 2(a).
3.3. Hydrogen Storage Properties
Figure 4 shows the P-C-T curves of the hydrogen absorption-desorption processes for the Mg-Ni ultrafine particles at 400°C, 375°C, 350°C, and 325°C. The data obtained from the P-C-T measurements are summarized in Table 1. Two plateaus of absorption and desorption are clearly visible on each profile. According to the XRD results given above, two types of hydrides, MgH2 and Mg2NiH4, formed after hydrogenation. Therefore, the two plateaus of the absorption and desorption are attributed to the formation and decomposition of MgH2 and Mg2NiH4. According to previous investigations on the Mg-rich Mg-Ni composite materials, the low plateaus are due to hydriding and dehydriding of Mg, while the high plateaus are the results of hydriding and dehydriding of Mg2Ni [15, 16]. The maximum hydrogen absorption capacities at 400, 375, 350, and 325°C are 3.16, 3.05, 2.97, and 2.96 wt%, respectively, which are even lower than the theoretical value of 3.6 wt% for Mg2Ni. This is due to the existence of MgO, the residual Mg and Ni in the hydrided Mg-Ni composite, as shown in Figure 2(b). Indeed, some large magnesium particles cannot be completely hydrogenated. From these PCT data, the van’t Hoff plots ( versus ) for the Mg-Ni composite are drawn in Figure 5. According to the linear fitting lines from the experimental data, the van’t Hoff equations for the hydrogenation are for Mg and for Mg2Ni. The van’t Hoff equations for the dehydrogenation are for Mg and for Mg2Ni. Therefore, the obtained values of the hydride formation enthalpies () are −83.1 kJ/mol for Mg and −64.6 kJ/mol for Mg2Ni, while the dehydrogenation enthalpies () are 77.4 kJ/mol for Mg and 66.2 kJ/mol for Mg2Ni. These hydrogenation and dehydrogenation enthalpies of Mg and Mg2Ni phases agree well with the values reported in the literature, for example, about −66.3 kJ/mol H2 for hydrogenation of Mg2Ni [10, 17] and −74.5 kJ/mol H2 for hydrogenation of Mg .
Figure 6 shows the TG/DSC curves of the hydrided Mg-Ni composite obtained in 4 MPa hydrogen atmosphere at 400°C. It is observed that three endothermic peaks appeared on the DSC curve measured from room temperature to 500°C. The small peak in the temperature range between 235°C and 250°C comes from the phase transformation of Mg2NiH4 from its low temperature form to its high temperature form . While in the same temperature range, TG curve does not show any mass loss. Two strong endothermic peaks in the high temperature range between 365°C and 425°C are observed on the DSC curve together with a mass loss of 3.3 wt%. They result from the dehydriding of Mg2NiH4 and MgH2. In order to understand the dehydriding processes of the hydrogenated Mg-Ni composite, a partially hydrogenated sample is prepared and studied by using DSC technique. The sample is hydrogenated at 350°C with 0.89 wt% of hydrogen absorption. Figure 7 shows the XRD pattern of the partially hydrogenated composite. From Figure 7, one can see that both MgH2 and Mg2NiH4 phases exist after absorbing 0.89 wt% of hydrogen. The intensity of peaks from MgH2 is relative higher than that of the Mg2NiH4 phase, indicating that the main product in this partially hydrogenated sample is MgH2. Therefore, the hydrogenation enthalpy calculated from the lower plateaus on the P-C-T curves actually comes from the formation of both MgH2 (majority) and Mg2NiH4 (minority) phases. Figure 6(b) shows the TG/DSC curves of the partially hydrided Mg-Ni composite. In Figure 6(b), two peaks are visible. A very weak endothermic peak appeared at 240°C, which corresponds to the phase transformation of Mg2NiH4. The other endothermic peak, with the onset point at 390°C and the peak point located at 420°C, must come from the dehydriding of MgH2. Therefore, the endothermic peak located at 370°C observed in Figure 6(a) is confirmed from the dehydriding reaction of Mg2NiH4 → Mg2Ni + 2H2 while that of 395°C is from the dehydriding reaction of MgH2 → Mg + H2. It is worth noting here that the dehydriding temperature of MgH2 in the fully hydrogenated powders is slightly lower than the partially hydrogenated powders. Also, the measured dehydrogenation enthalpy for MgH2 (83.1 kJ/mol H2) is slightly higher than its formation enthalpy (77.4 kJ/mol H2). XRD analysis showed that the difference between the partially hydrogenated Mg-Ni powders and fully hydrogenated powders lies in the amount of Mg2NiH4 phase. Therefore, the Mg2NiH4 phase may play a role for the dehydriding of MgH2 in the hydrogenated Mg-Ni composite powders. During dehydriding process, Mg2NiH4 decomposes prior to the MgH2. The decomposition of Mg2NiH4 into nanostructured Mg2Ni may act as “hydrogen channels” for the decomposition of MgH2. Similar phenomena were also observed in Mg-Nb2O5 system for which the Nb2O5 act as hydrogen channels for hydrogen sorption of Mg [20, 21].
Mg-rich Mg-Ni ultrafine powders were prepared through a DC arc plasma method. The results showed that Mg2Ni and MgNi2 phases could be obtained directly from the vapor state reactions between Mg and Ni vapors, depending on the local vapor content in the reaction chamber. After hydrogenation, a nanostructured MgH2+Mg2NiH4 hydrogen storage composite could be generated from the Mg-Ni ultrafine powders. After dehydrogenation, MgH2 and Mg2NiH4 decomposed into nanograined Mg and Mg2Ni, respectively. The hydrogenation enthalpy and dehydrogenation temperature of MgH2 in the composite are significantly reduced. TG/DSC and XRD analyses on the fully hydrogenated sample and partially hydrogenated sample showed that Mg2NiH4 phase may play a catalytic role in the dehydriding process of the hydrogenated Mg ultrafine particles.
Professor J. Zou would like to acknowledge the financial support from Research Funds for the Doctoral Program of Higher Education of China (no. 20100073120007) and from Shanghai Education Commission (no. 12ZZ017). This work is partly supported by projects from the Science and Technology Committee of Shanghai no. 10JC1407700, no. 11ZR1417600, and no. 10dz2211000 and “Pujiang” project no. 11PJ1406000. This work was also partially supported by the French-Chinese collaboration project “XU Guangqi” (convention no. 27847SC).
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