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Effects of trimesic acid-Ni based metal organic framework on the hydrogen sorption performances of MgH2
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Effects of trimesic acid-Ni based metal organic framework on the hydrogen sorption performances of MgH2 Zhewen Ma a, Jianxin Zou a,b,*, Chuanzhu Hu a, Wen Zhu a, Darvaish Khan a, Xiaoqin Zeng a,b, Wenjiang Ding a,b a
National Engineering Research Center of Light Alloys Net Forming & State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, PR China b Shanghai Engineering Research Center of Mg Materials and Applications & School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, PR China
article info
abstract
Article history:
A metal-organic framework based on Ni (II) as metal ion and trimasic acid (TMA) as organic
Received 6 December 2018
linker was synthesized and introduced into MgH2 to prepare a Mg-(TMA-Ni MOF)-H com-
Received in revised form
posite through ball-milling. The microstructures, phase changes and hydrogen storage
26 January 2019
behaviors of the composite were systematically studied. It can be found that Ni ion in TMA-
Accepted 29 January 2019
Ni MOF is attracted by Mg to form nano-sized Mg2Ni/Mg2NiH4 after de/rehydrogenation.
Available online 23 March 2019
The hydriding and dehydriding enthalpies of the Mg-MOF-H composite are evaluated to be
Keywords:
remains unchanged. The absorption kinetics of the Mg-MOF-H composite is improved by
Hydrogen storage
showing an activation energy of 51.2 kJ mol1 H2. The onset desorption temperature of the
74.3 and 78.7 kJ mol1 H2, respectively, which means that the thermodynamics of Mg
Mg hydride
composite is 167.8 K lower than that of the pure MgH2 at the heating rate of 10 K/min. Such
TMA-Ni MOF
a significant enhancement on the sorption kinetic properties of the composite is attributed
Kinetics
to the catalytic effects of the nanoscale Mg2Ni/Mg2NiH4 derived from TMA-Ni MOF by providing gateways for hydrogen diffusion during re/dehydrogenation processes. © 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Increasing demand of conventional energy and the increasing environmental contamination caused by the consumption of traditional fuels have hampered the worldwide sustainable development [1,2]. “Hydrogen energy”, as one kind of sustainable and clean energy with zero carbon emission, has attracted numerous of attention to replace long-standing traditional energy. Nevertheless, the lack of reliable and
efficient hydrogen storage medium is now inhibiting the wide application of hydrogen energy. As a solid state hydrogen carrier, MgH2 has been regarded one of prospective candidates for hydrogen storage, which is attributed to its high H2 uptake capacity (7.6 wt %), environmental friendly nature, low cost as well as the high abundance of Mg on earth [3,4]. Nevertheless, the hydrogenation enthalpy value (74.7 kJ mol1 H2) of MgH2 indicates the high thermodynamic stability of magnesium hydride, leading to a high operation temperature (>623 K) [3].
* Corresponding author. National Engineering Research Center of Light Alloys Net Forming & State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, PR China. E-mail address: [email protected] (J. Zou). https://doi.org/10.1016/j.ijhydene.2019.01.288 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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Additionally, the sluggish hydrogen sorption kinetics is another obstacle for the practical application of Mg based hydrogen storage materials [4,5]. Despite the diminution of the particle size of Mg to nano-scale [6e9], addition of catalytic elements e.g. Ti [10,11], Fe [10e13], Co [12e14], V [15,16], Cu [17,18], Mn [10,12] and Ni [10,13,14,17e20] etc., metal oxides including CeO2 [21], Cr2O3 [22], Nb2O5 [23], TiO2 [24] or other kinds of metal hydrides e.g. LiBH4 [25] into Mg to form Mg based composites is another approach to solve the issues mentioned above. At the same time, like other kinds of solid state carriers including metallic compounds, pure metals, carbon nanostructures and chemical hydride, metal organic frameworks (MOFs) have been applied as hydrogen storage materials [1,26]. For instance, Zn4O(BDC)3 named MOF-5 and Zn4O(BTB)2 (MOF177) can uptake 10 wt % and 7.0 wt % of H2 at 77 K, respectively [27,28]. Nonetheless, numerous works related to MOFs mainly focus on the physical absorption performances of hydrogen at 77 K, while the beneficial effects of MOFs on the hydrogen storage behaviors in Mg/MgH2 have rarely been reported yet. In this study, trimesic acid (TMA) and Ni (II) based MOF (TMA-Ni MOF) with no toxic solvents was firstly synthesized in deionized water (DI water) at room temperature. Considering the catalytic role of the TMA-Ni MOF played in MgH2, the overall hydrogen absorption/desorption performances and corresponding thermodynamic and kinetic parameters of pure MgH2 and Mg-(TMA-Ni MOF)-H composite were carefully studied.
Experimental Sample preparation Synthesis of TMA-Ni MOF in DI water NiSO4 $6H2O (99%) was chosen as the metal ion source, and trimesic acid (TMA, 99%) was used as the organic linker for the preparation of TMA-Ni MOF. Sodium hydroxide (99%) was used to adjust the pH value of the TMA solution. A NiSO4 solution was obtained by dissolving 0.009 mol NiSO4 $6H2O into 60 mL DI water. Meanwhile, 0.0045 mol TMA was dissolved into same amount of DI water and the pH of the solution was adjusted to 7 by dropping wise addition of 0.5 M NaOH. Then the prepared metal ion solution was added into the TMA solution under a constant stirring rate of 1200 rpm. Seconds later, green-colored precipitate (TMA-Ni MOF) appeared in the solution. Ten hours of mixing time was set for the complete formation of MOF. The synthesized precipitate was washed with excess DI water to remove residual reactants. Finally, the prepared TMA-Ni MOF was dried at 333 K for 8 h, followed with another 24 h at 373 K. Then the dried MOF was stored in glove box for further measurements.
container (material: SUS304, capacity ¼ 100 mL) under Argon atmosphere (1 bar) and ball milled at 150 rpm for 4, 8 and 16 h. For the preparation of the as milled TMA-Ni MOF, about 2.0 g pure TMA-Ni MOF powder was added into the container and ball milled with same parameters. Same amount (2.0 g) of pure MgH2 was ball milled in another container at 150 rpm for 8 h for comparison. The rotation speed of 150 rpm for ball milling was selected for the homogenization of TMA-Ni MOF powder with MgH2 in the Mg-MOF-H composite without severe destruction of the MOF structure.
Characterization The pore parameters of the TMA-Ni MOF were determined by N2 sorption on a 3H-2000PS2 (Beishide Co. Ltd). The structural analysis of the TMA-Ni MOF was performed via Fourier transform infrared spectroscopy (FTIR) by a Nicolet iS5 (Thermo Fisher Scientific Inc.) equipped with a horizontal ATR accessory. The thermal stability of the MOF was studied using thermolgravimetry analysis (TG, Netzsch STA449F3 Jupiter) by heating samples from 293 K to 823 K with the heating rate of 5 K/min under 1 bar of Ar atmosphere. X-ray diffraction (XRD, D/max 2550VL/PCX) measurements were carried out for the phase identification of the TMA-Ni MOF, Mg-MOF-H and the pure MgH2 powders at different states. A scotch tape was used on the XRD holder to prevent the samples from oxygen and H2O contamination. The collected data were analyzed using Jade 6.5 software. The microstructures and morphology of the Mg-MOF-H sample were identified by a transmission electron microscope (TEM, JEM-2100F) and a scanning electron microscope (SEM, Phenom-XL) equipped with an energy dispersive spectrometer (EDS). The hydriding/dehydriding performances of the Mg-MOF-H composite (ball milled for 8 h) and the pure MgH2 (ball milled for 8 h) were evaluated using a Sievert type pressure-composition-temperature (PCT) apparatus. The auto-PCT sorption tests at various temperatures including 598, 623 and 648 K were performed with minimum and maximum hydrogen pressures of 0.01 and 4 MPa, respectively. Before the PCT measurements, the Mg-MOF-H powder and the as milled pure MgH2 were both activated at 693 K for 2 h by operating a dehydrogenation process in vacuum. Dehydrogenation properties of the samples were examined by PCT at different temperatures and differential scanning calorimetry (DSC, Netzsch STA449F3 Jupiter) under 1 bar argon atmosphere at different heating rates from 298 K to 773 K. The charging/discharging performance of the Mg-MOF-H composite was measured at 623 K for 10 cycles. Each cycle was performed under 3 MPa H2 for absorption and in vacuum for desorption within 60 min. The sample powders were both hydrogenated at 623 K under 3 MPa hydrogen pressure for 2 h and dehydrogenated at 673 K in vacuum for 2 h in the PCT apparatus for XRD, SEM, TEM as well as DSC measurements.
Preparation of Mg-(TMA-Ni MOF)-H composites and as milled TMA-Ni MOF The Mg-(TMA-Ni MOF)-H (Mg-MOF-H) composites and the as milled TMA-Ni MOF with different milling durations were all prepared in a planetary ball mill with a ball-to-powder weight ratio of 20:1. About 1.8 g MgH2 (produced by Shanghai Mg Power Technology Co. Ltd., purity >98%) [29] and 0.2 g TMA-Ni MOF powder were added into a sealed stainless steel
Results and discussions Characterization of the Ni-MOF Despite synthetic method, metal salt and solvent used in this work, the interaction of Ni2þ and trimesic acid could be
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Fig. 1 e Schematic diagram of the preparation of TMA-Ni MOF.
assumed to be electrostatic in essence. During preparation process, green-colored precipitates that indicated the formation of TMA-Ni MOF rapidly formed with the addition of NiSO4 into the TMA solution. The schematic diagram of the prepared TMA-Ni MOF is shown in Fig. 1. It can be seen that
every Ni2þ bonding with two carboxylic functional groups and every two TMA necessitates the use of three Ni ions under ideal conditions. Fig. 2(a) and (b) present the micrographs of the as prepared TMA-Ni MOF at different magnifications. It can be found that the diameter and length of
Fig. 2 e SEM micrographs of the TMA-Ni MOF (a, b) under different magnifications and the raw material (pure MH2) (c) used in this work.
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Table 1 e Pore structural parameters of the TMA-Ni MOF obtained using the N2 absorption-desorption isotherms. Metal salt NiSO4
BET surface area (m2/g)
Pore size (nm)
Pore volume (cm3/g)
346.22
7.40
0.1736
individual needle-shaped TMA-Ni MOF are about 0.5 and 5 mm, respectively. Meanwhile, as shown in Fig. 2 (c), the MgH2 powder (raw material) used in this work is comprised of spherical particles with average particle size of 10 mm. The surface area with other pore specifications of the TMA-Ni MOF is listed in Table 1. As seen in Table 1, the BET specific surface area of the TMA-Ni MOF is 346.22 m2 g1 and its pore size as well as pore volume are estimated to be 7.40 nm and 0.1736 cm3 g1, respectively. The FTIR spectra of the TMA and the TMA-Ni MOF are given in Fig. 3 for structural determination. It is obvious that the most crucial functional groups in the MOF synthesized here are CeO single stretching bonds at about 1437.3 and
1373.6 cm1 and C]O bonds at about 1700 cm1. It can be found that the carbonyl groups peak divided into two peaks between 1500 and 1650 cm1 for the as prepared TMA-Ni MOF, which can be put down to the interaction of nickel ion and carboxylate groups [30]. Compared with the peaks in FTIR spectra of TMA, the disappeared peak at about 930 cm1 in that of the TMA-Ni MOF is ascribed to the OeH bond breaking during preparation process. This can be explained by the linking of Ni ion and trimesic acid in Fig. 1. The thermogravimetry (TG) curve of the TMA-Ni MOF is given in Fig. 4 to evaluate the thermostability of the MOF. According to Fig. 4, the degradation profile of the TMA-Ni MOF possesses two steps at about the temperature ranges of 310e500 and 670e740 K, indicating the solvent (DI water) loss and decomposition of MOF, respectively. Moreover, the as prepared MOF shows approximately 54 wt % of weight loss upon heating to 773 K.
Phase compositions and microstructural characterization of the Mg-MOF-H composite and the pure MgH2 Fig. 5 (a) shows the XRD patterns of the as prepared TMA-Ni MOF and the as milled pure MOFs after different ball milling durations including 4, 8 and 16 h. The characteristic peaks of the TMA-Ni MOF are located at relatively low angles (2q), such as 13.92 , 15.34 and 18.74 . It can be found that peaks of MOF remain unchanged after different ball milling durations, demonstrating that neither crystal structure change nor phase transformation takes place during ball milling at 150 rpm for different durations. The corresponding SEM
Fig. 3 e FTIR spectrum of the trimesic acid (TMA) and the TMA-Ni MOF synthesized in DI water.
Fig. 4 e Thermogravimetric curve of the TMA-Ni MOF.
Fig. 5 e XRD patterns of the pure TMA-Ni MOF powders (a) and Mg-MOF-H composites (b) ball milled for different time.
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micrographs of the as milled pure TMA-Ni MOFs after different ball milling durations (4, 8 and 16 h) are given in Fig. 6 (a)-(c), respectively. It can be clearly observed from Fig. 6 (a) that, some of the TMA-Ni MOF needles crush into small particles after ball milling (4 h), while rest of the individual TMANi MOF remains to be needle-shaped with length of about 1.5 mm. According to Fig. 6 (b) and (c), all of the TMA-Ni MOF crushes into small particles with particle sizes range from 0.1 to 0.5 mm after ball milling for 8 h and the particle size of the as milled MOF cannot be further lowered down with increasing the ball milling duration (16 h). Fig. 5 (b) presents the XRD results of the as milled Mg-MOFH composites for different durations (4, 8 and 16 h). Despite the weak peaks from Mg, the patterns of composites are all indexed as MgH2 without MgO or other phases. It means that there is no reaction or phase transformation during ball milling process with the increase of ball milling duration from 4 h to 16 h at 150 rpm. The SEM images of the ball milled MgMOF-H composites for different ball milling times are shown in Fig. 6(def) and Fig. 6(gei), (j-l) and (m-o) display the EDS
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mapping of Mg, Ni and O, respectively. The distribution of Ni and O on the surface of MgH2 particles are more homogeneous in the Mg-MOF-H powders after ball milling for 8 and 16 h than those of the powder ball milled for only 4 h, implying that the ball milling duration of 8 h at 150 rpm is an optimum ball milling parameter for the preparation of the Mg-MOF-H composite. Fig. 7 (a)-(f) illustrate the XRD patterns of the TMA-Ni MOF, as milled Mg-MOF-H composite (8 h), as milled pure MgH2 (8 h) and the powders after hydriding/dehydriding, respectively. Despite a weak peak from Mg, the patterns of the as milled Mg-MOF composite (Fig. 7 (b)) and MgH2 powders (Fig. 7 (c)) can be identified as pure MgH2 without any other phases. Comparatively, same phases can also be identified in the pattern of the pure MgH2 powder after hydrogenation at 623 K for 2 h (Fig. 7 (d)). According to Fig. 7 (e), the re-hydrogenated Mg-MOF-H powder composed of b-MgH2 as major phase with small quantity of Mg2NiH4, Mg and MgO phases. The presence of remnant Mg demonstrates that some large-sized Mg cannot be completely hydrogenated at 623 K for 2 h.
Fig. 6 e SEM micrographs of the pure TMA-Ni MOF powders (aec) and Mg-MOF-H composites (def) ball milled for different time as well as the corresponding EDS mappings of Mg-MOF-H powders showing the distributions of Mg (gei), Ni (jel) and O (meo).
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Fig. 7 e XRD patterns of the TMA-Ni MOF (a), Mg-MOF-H composite and pure MgH2 in varying states: as milled (b), (c), re-hydrogenated at 623 K (d), (e), and dehydrogenated at 673 K (f), respectively.
Meanwhile, the formation of Mg2NiH4 phase in the hydrogenated Mg-MOF-H composite reveals a reaction of H2, Mg and Ni to generate Mg2NiH4 phase during hydrogenation process. The reaction can be expressed as follows: 2 Mg þ Ni þ 2H2 ¼ Mg2NiH4
(1)
Furthermore, when the re-hydrogenated Mg-MOF-H composite powder were dehydrogenated at 673 K, the XRD pattern shows the appearance of Mg and Mg2Ni phases, which indicates that all of the MgH2 and Mg2NiH4 transform to Mg and Mg2Ni, respectively (see Fig. 7 (f)). Additionally, small amount of MgO is present in re/dehydrogenated Mg-MOF-H composites. The existence of oxygen in MgO and nickel in Mg2NiH4/ Mg2Ni phases suggests the decomposition of the TMA-Ni MOF additive at high temperatures. It can be deduced from Figs. 4 and 7 that the TMA-Ni MOF in Mg-MOF-H composite is getting unstable with the increment of temperature. The structure of the TMA-Ni MOF collapses at 693 K during activation process. The chemical bonds among component elements in MOF structure were broken down and consequently, nickel, oxygen as well as amorphous carbon remain in the composite. The contact area between the MOF additive and MgH2 particles increases during the ball-milling process and part of MOF needles can penetrate the extremely thin MgO layer covered on MgH2 particles. In the following hydrogenation process at 623 K, the large contact area between the MOF and MgH2 makes it easy for nickel and oxygen from the TMA-Ni MOF to react with surrounding Mg under 3 MPa H2 atmosphere to produce Mg2NiH4 and MgO, respectively. The typical SEM images of the as milled, hydrogenated and dehydrogenated MgH2 powders are shown in Fig. 8 (a), (e) and (i), respectively. It can be observed that the average particle size of the pure MgH2 remains at 10 mm, which indicates that the ball-milling process on the given condition (150 rpm, 8 h) cannot prominently reduce the average particle size of the MgH2 powder. According to Fig. 8 (e) and (i), the particle sizes
of MgH2 and Mg both increase from about 10 mm to 30 mm, demonstrating a significant agglomeration of MgH2/Mg particles during re/dehydrogenation processes. Same phenomenon can be found on the Mg-MOF-H composite in varying states (see Fig. 8(bed)). The particle size of the composite powder goes up from about 12 mm (as milled state) to 30e40 mm (re/dehydrogenated states), suggesting that the addition of TMA-Ni MOF additive cannot reduce the average particle size of the Mg-MOF-H composite in varying states. Besides, Fig. 8(feh) and (j-l) display the EDS mapping of Ni and O, respectively. It can be found that, Ni and O are homogeneously distributed on MgH2 (as milled and hydrogenated) and Mg powders (dehydrogenated), indicating the uniform distribution of Mg2NiH4, Mg2Ni and MgO phases on the surface of Mg particles in varying states. It is mainly attributed to the appropriate ball-milling process during the preparation of the composite. The TEM observations were conducted in Fig. 9 to display the micro-morphology of the hydrogenated Mg-MOF-H composite. Fig. 9 (a) presents a bright field image of the Mg-MOF-H sample. It can be observed from Fig. 9(a) that two irregularshaped Mg-MOF-H particles aggregating together and the size of the composite is measured to be 500 nm. Based on the SAED pattern of the Mg-MOF-H composite in Fig. 9 (b), all the diffraction rings corresponding well to MgH2, Mg2NiH4 and MgO phases. It is consistent with the XRD results. The appearance of MgO phase can be ascribed to the reaction between O (provided by the degradation of TMA-Ni MOF) and Mg during (de)hydriding cycles. To show the morphology of Mg2NiH4 phase on MgH2 particles, a dark field (DF) micrograph presented in Fig. 9 (c) is obtained from Mg2NiH4 (220) and MgO (200) diffraction rings. The Mg2NiH4 and MgO nano-sized particles with bright contrasts uniformly disperse on the large-size MgH2. The sizes of Mg2NiH4 nano-particles range from 5 nm to 20 nm. This is in accordance with the EDS mappings of Ni and O in SEM observations. Additionally, High Resolution Electron Microscope (HRTEM) is performed to reveal the morphology of the Mg-MOF-H sample. It is clearly seen that the interplanar spacings of the selected zones in Fig. 9(d) are determined to be about 0.2312 and 0.2291 nm, corresponding to the Mg2NiH4 (220) plane (0.2309 nm).
Hydrogen storage performances of the pure MgH2 and the Mg-MOF-H composite The hydrogen storage performances of the MgH2 and the Mg-MOF-H composite are evaluated by using the PCT facility. Fig. 10 (a) and (c) illustrate the PC isotherms of the pure MgH2 and the Mg-MOF-H composite tested at 598, 623 and 648 K in a H2 pressure ranging from 0.01 to 4 MPa, respectively. The single slant plateaus owing to the redehydrogenation of MgH2 phase are presented under assigned temperatures in Fig. 10 (a). It can also be found that no obvious plateau can be found in the PCT profile during dehydrogenation of the pure MgH2 at each temperature. Furthermore, a sharp diminution of the H2 pressure at the initial stage of dehydriding process can be easily observed in desorption profile at 598 K, indicating the poor dehydriding kinetic property of the pure MgH2. In contrast, the flat plateaus of the Mg-MOF-H composite at each temperature are
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shown in Fig. 10 (c). Each of the H2 desorption curves possesses a distinct plateau and no decline of pressure can be found even the testing temperature decreases to 598 K. It means that the addition of TMA-Ni MOF has prominent catalytic efficiency in promoting the (de)hydriding kinetics of Mg. It should be pointed out that, only one plateau can be observed in each PC isotherms of the Mg-MOF-H sample. In comparison to major phases (Mg/MgH2), neither the quantity of Mg2Ni nor Mg2NiH4 is enough to form another distinct plateau during hydriding/dehydriding cycles [31]. The hydrogen uptake amounts for the Mg-MOF-H powder at 598, 623 and 648 K are 5.09, 5.04 and 5.17 wt %, respectively. The uptake amounts of hydrogen for the composite under series temperatures are all lower than the theoretical capacity of
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MgH2 (7.6 wt %) [3,4], which is ascribed to the addition of TMA-Ni MOF and the formation of MgO during hydrogen de/ absorption cycles. The de-hydrogenation enthalpies of the pure MgH2 and the Mg-MOF-H samples are estimated using the van't Hoff equation. Fig. 10 (b) and (d) reveal the linear fittings (ln P vs. 1000/T) of the van't Hoff plots and subsequently, the enthalpies of the hydriding/dehydriding processes for Mg-MOFH sample are estimated to be 74.3 and 78.7 kJ mol1 H2, respectively. They are fairly close to the dehydrogenated MgH2 (73.8 and 76.1 kJ mol1 H2) as well as reported values of Mg (±74.7 kJ mol1 H2) [32]. With the aim of investigating the hydriding kinetics of the Mg-MOF-H composite, the profiles of hydrogenation for the
Fig. 8 e SEM micrographs of the as milled (a), hydrogenated (e) and dehydrogenated (i) pure MgH2 powders, the as milled (b), hydrogenated (c) and dehydrogenated (d) Mg-MOF-H composite powders as well as the corresponding EDS mappings showing the distributions of Ni (feh) and O (jel).
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dehydrogenated MgH2 and the Mg-MOF-H composite tested at 448, 473, 498, 523, 548 and 573 K under 3 MPa H2 are given in Fig. 11 (a) and (c), respectively. It is worth noting from Fig. 11 (a) and (c) that the disparity of hydrogenation profiles of the dehydrogenated MgH2 and the Mg-MOF-H is getting remarkable with the reduction of the temperature, especially at relative low temperatures including 448, 473 and 498 K. The hydrogen absorption rate of the dehydrogenated pure MgH2 observably decreases at low temperatures and the sample can merely absorb 2.1 wt % of hydrogen at 448 K for 3 h. In contrast, an evident improvement in the H2 uptake rate of the dehydrogenated Mg-MOF-H powder is achieved under low temperatures. The inset of Fig. 11 (c) reveals that the absorption rate of the Mg-MOF-H does not sharply decline with decreasing the temperature, even at 448 K, indicating the much better absorption kinetics with respect to that of dehydrogenated pure MgH2 sample. In addition, Fig. 12 illustrates the changes of the hydriding percentages of the dehydrogenated MgH2 and the Mg-MOF-H samples at low temperatures. The time needed for dehydrogenated pure MgH2 to absorb 80% of its maximum hydrogen capacity is 7430 s at 473 K under 3 MPa H2 pressure, which is much longer than that of the dehydrided Mg-MOF-H sample (2988 s), implying the enhanced hydriding rate of the Mg-MOF-H sample. The beneficial role of Mg2Ni played during hydrogenation process is responsible for the superior H absorption rate of the Mg-MOF-H sample [31,33].
It has been established that the rate controlling factors of the hydrogenation process in metals involve three steps including the hydrogen molecules dissociation, the penetration of H into the bulk metal and the diffusion of H in the metal matrix [34]. An activation energy (Ea) is taken to summarize the energy barrier during hydrogenation. According to the Johnson-Mehl-Avrami-Kolmogorov (JMAK) model, the data can be systematically analyzed and the linear equation of the sorption kinetics can be expressed as: ln½lnð1 aÞ ¼ h ln k þ h ln t
(2)
where the value of a represents the fraction of phase transformation from metal (Mg) to its hydride (MgH2) at time t, k and h donate a parameter of kinetics and the reaction order or Avrami exponent, respectively. With the rate constant k calculated from the equation above, the Ea can be estimated as follows: k ¼ A expð Ea =RTÞ
(3)
where A donates a temperature-independent coefficient, R and T represent the gas constant and the absolutely temperature, respectively. As shown in Fig. 11 (b) and (d), the Ea values in the dehydrogenated MgH2 and Mg-MOF-H powders are evaluated to be 100.7 and 51.2 kJ mol1 H2, respectively. Such an evident reduction in activation energy suggests a
Fig. 9 e A typical bright field TEM micrograph of the Mg-MOF-H powder (a) hydrogenated at 623 K under 3 MPa H2 atmosphere for 2 h, the corresponding SAED pattern (b), the dark field micrograph obtained from Mg2NiH4 (220) and MgO (200) diffraction rings and the corresponding HRTEM image (d) of the Mg-MOF-H composite.
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lower energy barrier of hydrogen absorption, explaining the better hydriding kinetics of the dehydrogenated Mg-MOF-H sample over the dehydrogenated pure MgH2, particularly at relative low temperatures. As shown in Fig. 13, the Mg-MOFH composite can be obtained after ball-milling process and the particle size of the composite powder gradually increases from about 10 mm to 30 mm after re-dehydrogenation cycles at high temperatures (see Fig. 8). Seemingly, the agglomerations of the Mg-MOF-H particles inhibit the hydrogen absorption, which correspondingly results in sluggish sorption kinetics. Nevertheless, Fig. 13 illustrates that the big-sized dehydrogenated Mg-MOF-H particles (average particle sizez30 mm) are comprised of many Mg2Ni-covered Mg particles (average particle sizez10 mm). It has been established that the Mg2Ni on the surface of Mg particles plays a catalytic role on the chemical absorption of H2 and easily dissociate them into H atoms [35]. As is well known, MgH2 layer outside of Mg particles hampers the diffusion of H atoms from surface into Mg bulk. While the Mg2Ni uniformly distributed on the Mg particles can play a “gateway” role as channels to accelerate the diffusion of hydrogen atoms into the bulk of Mg during hydrogenation process [36]. It results in the remarkable improvement of absorption performance (see Fig. 11). Furthermore, Liang et al. [37] reported that Mg2Ni not only played a catalytic role on the chemisorption of hydrogen, but also markedly facilitated the hydrogen
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diffusion rate in Mg. In addition, Liu et al. have pointed out that MgH2 phase preferentially nucleates at the interface between Mg/Mg2Ni during hydriding, through which the absorption kinetics of the Mg-MOF-H composite can be significantly accelerated [31]. In brief, the beneficial effects mentioned above lead to the outstanding absorption kinetic performances of the dehydrogenated Mg-MOF-H composite. To characterize the dehydrogenation behaviors of the MgMOF-H composite, the hydrogen desorption profiles of the MgMOF-H and MgH2 powders at different temperatures are presented in Fig. 14 (a) and (b), respectively. It can be found that the dehydriding rates of the Mg-MOF-H powders are all much higher than those of pure MgH2 powder at various temperatures. The hydrogen desorption capacity of the MgH2 sample sharply declines with the decreasing temperature. The MgH2 sample releases merely 0.3 wt % H2 at 573 K while the MgMOF-H composite still exhibits a desorption capacity of about 3 wt % H2 at the same temperature, implying the faster dehydriding kinetics of the composite than that of the pure MgH2 powder. The DSC curves of MgH2 and Mg-MOF-H composite are given in Fig. 14 (c) and (e), respectively. As found in Fig. 14 (c), a strong endothermic peak can be observed in each DSC curve at different heating rates in the pure MgH2 sample. The endothermic peak is due to the degradation of b-MgH2. However, it is interesting to note that the DSC profiles of the Mg-MOF-H sample (see Fig. 14 (e)) possesses two endothermic
Fig. 10 e PC isotherms and the van't Hoff plots of the pure MgH2 (a), (b) and the Mg-MOF-H composite (c), (d) tested at 598, 623 and 648 K (Ab: absorption; De: desorption).
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peaks at the heating rate of 10 K/min, demonstrating a twostep dehydrogenation process. Similar phenomenon on the DSC results of the Mg-Ni composite has also been reported in the previous work [36]. According to the XRD patterns in Fig. 5 and the researches on the Mg-Ni powders [36e38] and Mg2Ni compound [39], the weak endothermic peak at relatively low temperature owing to the dehydrogenation of Mg2NiH4 phase. The strong broader peak corresponds well to the decomposition of b-MgH2. It should be noticed that, the first endothermic peak mentioned above gets unconspicuous with decreasing the heating rate and even disappears when the heating rate declines to 3 K/min. The interpretation might be that all endothermic peaks get weak with descent of the heating rate from 10 to 3 K/min, leading to a further weakening of the first small endothermic peak. The subdued small peak may be involved into the strong endothermic peak at the heating rate of 3 K/min, making it difficult to identify the weak peak of the decomposition for Mg2NiH4. It can be summarized from Fig. 14 (c) that peak temperatures of DSC profiles of the pure MgH2 powder at the heating rates including 3 K/min, 5 K/min and 10 K/min are 677.1, 688.4 and 701.8 K, respectively. Whereas, the peak hydrogen desorption temperatures of the Mg-MOF-H powder at the same heating rates are prominently reduced to 545.9, 556.3
Fig. 12 e Hydrogen absorption percentages (%) of the dehydrogenated pure MgH2 and the Mg-MOF-H composite measured at relative low temperatures including 448, 473 and 498 K under 3 MPa H2 pressure for 3 h (De: dehydrogenated).
Fig. 11 e Hydrogen absorption profiles and the corresponding ln k-1000/T plots of the dehydrogenated pure MgH2 (a), (b) and the Mg-MOF-H composite (c), (d) measured under 3 MPa H2 pressure for 3 h.
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Fig. 13 e Schematic illustration showing the preparation of the Mg-MOF-H composite and the reactive details during the hydrogenation process.
Fig. 14 e Hydrogen desorption profiles of the Mg-MOF-H composite (a) and the pure MgH2 (b), DSC profiles and the corresponding ln (b/T2p)-1000/Tp plots for the pure MgH2 (c), (d) and the Mg-MOF-H composite (e), (f) at different heating rates.
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Fig. 15 e Cycling curves of the Mg-MOF-H composite at 623 K for 10 times.
and 568.8 K, respectively. Furthermore, the onset temperature for dehydriding is significantly declined to 497.2 K, which is 167.8 K lower than that of pure MgH2 sample (665 K), as found in Fig. 14 (e). The dehydrogenation activation energy (Ed) can be evaluated to ulteriorly investigate the excellent dehydriding performance of the Mg-MOF-H composite. The activation energy of dehydriding process can be calculated by the Kissinger equation: . ln b Tp 2 ¼ A Ed RTp
(4)
where b donates the heating rate, Tp is the peak dehydriding temperature, A is a linear constant and R is the gas constant. The Ed of the pure MgH2 is calculated to be 181.4 kJ mol1 H2 with the linear fitting the plots of ln (b/T2p) vs. 1000/Tp in Fig. 14 (d) and (f). The Ed of MgH2 can be effectively declined to 126.7 kJ mol1 H2 with the addition of TMA-Ni MOF and consequently, leading to an outstanding improvement on desorption kinetic performance of the Mg-MOF-H. During the dehydriding of the Mg-MOF-H sample, Mg2NiH4 nanoparticles disperse on the surface of MgH2 firstly decomposes to Mg2Ni and H2, followed by the dehydrogenation of MgH2. This has been demonstrated by the DSC measurement. Zalusk et al. reported a synergetic effect from the mixture of Mg2NiH4 and MgH2 to enhance the desorption kinetics during dehydriding. Results in the literature revealed that the addition of Mg2NiH4 by ball-milling destabilized the magnesium hydride and consequently, led to the superior desorption kinetic property [39]. Also, those Mg2NiH4 coated homogeneously on MgH2 particles act as channels for hydrogen diffusion from magnesium hydride matrix [36,40]. It is necessary to point out that the magnesium oxide possesses a beneficial effect of grain boundary pinning, preventing the growth of Mg grains at elevated temperatures [41]. Meanwhile, the catalytically active sites which is attributed to the special vacancies on the MgO, make it easier to recombine atomic hydrogen during hydrogen desorption of MgH2 [42]. In short, the nanoscale Mg2NiH4 on the surface of MgH2 particles as well as the existence of magnesium
oxide, contribute to the improved dehydrogenation performance of the Mg-MOF-H composite. To investigate the effect of the TMA-Ni MOF additive on the cycling performance of MgH2, continuous hydriding/dehydriding tests of the Mg-MOF-H composite are performed at 623 K under 3 MPa H2 for hydrogenation and in vacuum for dehydrogenation. Fig. 15 illustrates the hydrogen absorption and desorption curves of the composite for 10 cycles. As seen in Fig. 15, no evident hydrogen storage capacity loss is detected after 10 cycles, indicating the superior cycling stability of the Mg-MOF-H composite. In addition, the hydrogen absorption/desorption kinetics of the composite remains unchanged and the composite possesses a constant reversible capacity of 3.4 wt % at 623 K. The results gathered in this study have shown the efficient catalytic effect of TMA-Ni MOF on hydrogen sorption properties of Mg/MgH2. As compared to additives of Mg2NiH4 or Ni nanoparticles [43e45], TMA-Ni MOF can be introduced into MgH2 through simple ball milling with fairly short duration and homogeneous distribution on the MgH2 particles. In addition, SEM and TEM observations (see Figs. 8 and 9) confirm the formation of nano-sized Mg2NiH4 coated on the surface of MgH2 particles without severe agglomeration after hydrogen sorption cycles. Furthermore, a relatively low amount of TMANi MOF addition can significantly improve the hydrogenation/ dehydrogenation performances of Mg/MgH2, revealing its high catalytic effectiveness.
Conclusions In this work, the TMA-Ni MOF was successfully synthesized for the first time in an environmentally friendly media, which was introduced into MgH2 to form the Mg-MOF-H composite through ball milling. The main conclusions can be drawn as follows: 1) XRD results show the formation of Mg2NiH4 phase when the dehydrogenated Mg-MOF-H composite is rehydrogenated and the Mg2NiH4 phase completely decomposes to Mg2Ni when the sample is dehydrided, indicating that Ni in the MOF is attracted by Mg to form Mg-Ni compound during de/re-hydrogenations. 2) SEM and TEM observations reveal that the nano-sized Mg2NiH4 is uniformly coated on the surface of large MgH2 particles in the Mg-MOF-H composite. 3) The hydriding and dehydriding enthalpies of the Mg-MOFH composite are determined to be 74.3 and 78.7 kJ mol1 H2, respectively. These values are close to the standard values of Mg (±74.7 kJ mol1 H2). 4) The dehydrogenated Mg-MOF-H composite exhibits enhanced hydrogen absorption kinetics as compared to the dehydrogenated pure MgH2. The Ea of the dehydrogenated Mg-MOF-H powder is estimated to be 51.2 kJ mol1 H2. 5) The hydrogen desorption temperature of the Mg-MOF-H composite is notably reduced in comparison to that of the pure MgH2. The onset dehydriding temperature at the heating rate of 10 K/min is 167.8 K lower than that of the pure MgH2. The corresponding Ed for the composite is calculated to be 126.7 kJ mol1 H2.
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6) The nanoscale Mg2Ni/Mg2NiH4 coated on the Mg/MgH2 powders dramatically enhances the hydrogen storage kinetics of the Mg-MOF-H composite by providing “gateway” for hydrogen diffusion during hydriding/dehydriding processes.
Acknowledgment Prof. Zou would like to thank the support from the National Science Foundation (No. 51771112), the Shanghai Science and Technology Commission under No. 14JC1491600 and Shanghai Education Commission “Shuguang” scholar project (16SG08).
references
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Available online at www.sciencedirect.com
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Effects of trimesic acid-Ni based metal organic framework on the hydrogen sorption performances of MgH2 Zhewen Ma a, Jianxin Zou a,b,*, Chuanzhu Hu a, Wen Zhu a, Darvaish Khan a, Xiaoqin Zeng a,b, Wenjiang Ding a,b a
National Engineering Research Center of Light Alloys Net Forming & State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, PR China b Shanghai Engineering Research Center of Mg Materials and Applications & School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, PR China
article info
abstract
Article history:
A metal-organic framework based on Ni (II) as metal ion and trimasic acid (TMA) as organic
Received 6 December 2018
linker was synthesized and introduced into MgH2 to prepare a Mg-(TMA-Ni MOF)-H com-
Received in revised form
posite through ball-milling. The microstructures, phase changes and hydrogen storage
26 January 2019
behaviors of the composite were systematically studied. It can be found that Ni ion in TMA-
Accepted 29 January 2019
Ni MOF is attracted by Mg to form nano-sized Mg2Ni/Mg2NiH4 after de/rehydrogenation.
Available online 23 March 2019
The hydriding and dehydriding enthalpies of the Mg-MOF-H composite are evaluated to be
Keywords:
remains unchanged. The absorption kinetics of the Mg-MOF-H composite is improved by
Hydrogen storage
showing an activation energy of 51.2 kJ mol1 H2. The onset desorption temperature of the
74.3 and 78.7 kJ mol1 H2, respectively, which means that the thermodynamics of Mg
Mg hydride
composite is 167.8 K lower than that of the pure MgH2 at the heating rate of 10 K/min. Such
TMA-Ni MOF
a significant enhancement on the sorption kinetic properties of the composite is attributed
Kinetics
to the catalytic effects of the nanoscale Mg2Ni/Mg2NiH4 derived from TMA-Ni MOF by providing gateways for hydrogen diffusion during re/dehydrogenation processes. © 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Increasing demand of conventional energy and the increasing environmental contamination caused by the consumption of traditional fuels have hampered the worldwide sustainable development [1,2]. “Hydrogen energy”, as one kind of sustainable and clean energy with zero carbon emission, has attracted numerous of attention to replace long-standing traditional energy. Nevertheless, the lack of reliable and
efficient hydrogen storage medium is now inhibiting the wide application of hydrogen energy. As a solid state hydrogen carrier, MgH2 has been regarded one of prospective candidates for hydrogen storage, which is attributed to its high H2 uptake capacity (7.6 wt %), environmental friendly nature, low cost as well as the high abundance of Mg on earth [3,4]. Nevertheless, the hydrogenation enthalpy value (74.7 kJ mol1 H2) of MgH2 indicates the high thermodynamic stability of magnesium hydride, leading to a high operation temperature (>623 K) [3].
* Corresponding author. National Engineering Research Center of Light Alloys Net Forming & State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, PR China. E-mail address: [email protected] (J. Zou). https://doi.org/10.1016/j.ijhydene.2019.01.288 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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Additionally, the sluggish hydrogen sorption kinetics is another obstacle for the practical application of Mg based hydrogen storage materials [4,5]. Despite the diminution of the particle size of Mg to nano-scale [6e9], addition of catalytic elements e.g. Ti [10,11], Fe [10e13], Co [12e14], V [15,16], Cu [17,18], Mn [10,12] and Ni [10,13,14,17e20] etc., metal oxides including CeO2 [21], Cr2O3 [22], Nb2O5 [23], TiO2 [24] or other kinds of metal hydrides e.g. LiBH4 [25] into Mg to form Mg based composites is another approach to solve the issues mentioned above. At the same time, like other kinds of solid state carriers including metallic compounds, pure metals, carbon nanostructures and chemical hydride, metal organic frameworks (MOFs) have been applied as hydrogen storage materials [1,26]. For instance, Zn4O(BDC)3 named MOF-5 and Zn4O(BTB)2 (MOF177) can uptake 10 wt % and 7.0 wt % of H2 at 77 K, respectively [27,28]. Nonetheless, numerous works related to MOFs mainly focus on the physical absorption performances of hydrogen at 77 K, while the beneficial effects of MOFs on the hydrogen storage behaviors in Mg/MgH2 have rarely been reported yet. In this study, trimesic acid (TMA) and Ni (II) based MOF (TMA-Ni MOF) with no toxic solvents was firstly synthesized in deionized water (DI water) at room temperature. Considering the catalytic role of the TMA-Ni MOF played in MgH2, the overall hydrogen absorption/desorption performances and corresponding thermodynamic and kinetic parameters of pure MgH2 and Mg-(TMA-Ni MOF)-H composite were carefully studied.
Experimental Sample preparation Synthesis of TMA-Ni MOF in DI water NiSO4 $6H2O (99%) was chosen as the metal ion source, and trimesic acid (TMA, 99%) was used as the organic linker for the preparation of TMA-Ni MOF. Sodium hydroxide (99%) was used to adjust the pH value of the TMA solution. A NiSO4 solution was obtained by dissolving 0.009 mol NiSO4 $6H2O into 60 mL DI water. Meanwhile, 0.0045 mol TMA was dissolved into same amount of DI water and the pH of the solution was adjusted to 7 by dropping wise addition of 0.5 M NaOH. Then the prepared metal ion solution was added into the TMA solution under a constant stirring rate of 1200 rpm. Seconds later, green-colored precipitate (TMA-Ni MOF) appeared in the solution. Ten hours of mixing time was set for the complete formation of MOF. The synthesized precipitate was washed with excess DI water to remove residual reactants. Finally, the prepared TMA-Ni MOF was dried at 333 K for 8 h, followed with another 24 h at 373 K. Then the dried MOF was stored in glove box for further measurements.
container (material: SUS304, capacity ¼ 100 mL) under Argon atmosphere (1 bar) and ball milled at 150 rpm for 4, 8 and 16 h. For the preparation of the as milled TMA-Ni MOF, about 2.0 g pure TMA-Ni MOF powder was added into the container and ball milled with same parameters. Same amount (2.0 g) of pure MgH2 was ball milled in another container at 150 rpm for 8 h for comparison. The rotation speed of 150 rpm for ball milling was selected for the homogenization of TMA-Ni MOF powder with MgH2 in the Mg-MOF-H composite without severe destruction of the MOF structure.
Characterization The pore parameters of the TMA-Ni MOF were determined by N2 sorption on a 3H-2000PS2 (Beishide Co. Ltd). The structural analysis of the TMA-Ni MOF was performed via Fourier transform infrared spectroscopy (FTIR) by a Nicolet iS5 (Thermo Fisher Scientific Inc.) equipped with a horizontal ATR accessory. The thermal stability of the MOF was studied using thermolgravimetry analysis (TG, Netzsch STA449F3 Jupiter) by heating samples from 293 K to 823 K with the heating rate of 5 K/min under 1 bar of Ar atmosphere. X-ray diffraction (XRD, D/max 2550VL/PCX) measurements were carried out for the phase identification of the TMA-Ni MOF, Mg-MOF-H and the pure MgH2 powders at different states. A scotch tape was used on the XRD holder to prevent the samples from oxygen and H2O contamination. The collected data were analyzed using Jade 6.5 software. The microstructures and morphology of the Mg-MOF-H sample were identified by a transmission electron microscope (TEM, JEM-2100F) and a scanning electron microscope (SEM, Phenom-XL) equipped with an energy dispersive spectrometer (EDS). The hydriding/dehydriding performances of the Mg-MOF-H composite (ball milled for 8 h) and the pure MgH2 (ball milled for 8 h) were evaluated using a Sievert type pressure-composition-temperature (PCT) apparatus. The auto-PCT sorption tests at various temperatures including 598, 623 and 648 K were performed with minimum and maximum hydrogen pressures of 0.01 and 4 MPa, respectively. Before the PCT measurements, the Mg-MOF-H powder and the as milled pure MgH2 were both activated at 693 K for 2 h by operating a dehydrogenation process in vacuum. Dehydrogenation properties of the samples were examined by PCT at different temperatures and differential scanning calorimetry (DSC, Netzsch STA449F3 Jupiter) under 1 bar argon atmosphere at different heating rates from 298 K to 773 K. The charging/discharging performance of the Mg-MOF-H composite was measured at 623 K for 10 cycles. Each cycle was performed under 3 MPa H2 for absorption and in vacuum for desorption within 60 min. The sample powders were both hydrogenated at 623 K under 3 MPa hydrogen pressure for 2 h and dehydrogenated at 673 K in vacuum for 2 h in the PCT apparatus for XRD, SEM, TEM as well as DSC measurements.
Preparation of Mg-(TMA-Ni MOF)-H composites and as milled TMA-Ni MOF The Mg-(TMA-Ni MOF)-H (Mg-MOF-H) composites and the as milled TMA-Ni MOF with different milling durations were all prepared in a planetary ball mill with a ball-to-powder weight ratio of 20:1. About 1.8 g MgH2 (produced by Shanghai Mg Power Technology Co. Ltd., purity >98%) [29] and 0.2 g TMA-Ni MOF powder were added into a sealed stainless steel
Results and discussions Characterization of the Ni-MOF Despite synthetic method, metal salt and solvent used in this work, the interaction of Ni2þ and trimesic acid could be
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Fig. 1 e Schematic diagram of the preparation of TMA-Ni MOF.
assumed to be electrostatic in essence. During preparation process, green-colored precipitates that indicated the formation of TMA-Ni MOF rapidly formed with the addition of NiSO4 into the TMA solution. The schematic diagram of the prepared TMA-Ni MOF is shown in Fig. 1. It can be seen that
every Ni2þ bonding with two carboxylic functional groups and every two TMA necessitates the use of three Ni ions under ideal conditions. Fig. 2(a) and (b) present the micrographs of the as prepared TMA-Ni MOF at different magnifications. It can be found that the diameter and length of
Fig. 2 e SEM micrographs of the TMA-Ni MOF (a, b) under different magnifications and the raw material (pure MH2) (c) used in this work.
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Table 1 e Pore structural parameters of the TMA-Ni MOF obtained using the N2 absorption-desorption isotherms. Metal salt NiSO4
BET surface area (m2/g)
Pore size (nm)
Pore volume (cm3/g)
346.22
7.40
0.1736
individual needle-shaped TMA-Ni MOF are about 0.5 and 5 mm, respectively. Meanwhile, as shown in Fig. 2 (c), the MgH2 powder (raw material) used in this work is comprised of spherical particles with average particle size of 10 mm. The surface area with other pore specifications of the TMA-Ni MOF is listed in Table 1. As seen in Table 1, the BET specific surface area of the TMA-Ni MOF is 346.22 m2 g1 and its pore size as well as pore volume are estimated to be 7.40 nm and 0.1736 cm3 g1, respectively. The FTIR spectra of the TMA and the TMA-Ni MOF are given in Fig. 3 for structural determination. It is obvious that the most crucial functional groups in the MOF synthesized here are CeO single stretching bonds at about 1437.3 and
1373.6 cm1 and C]O bonds at about 1700 cm1. It can be found that the carbonyl groups peak divided into two peaks between 1500 and 1650 cm1 for the as prepared TMA-Ni MOF, which can be put down to the interaction of nickel ion and carboxylate groups [30]. Compared with the peaks in FTIR spectra of TMA, the disappeared peak at about 930 cm1 in that of the TMA-Ni MOF is ascribed to the OeH bond breaking during preparation process. This can be explained by the linking of Ni ion and trimesic acid in Fig. 1. The thermogravimetry (TG) curve of the TMA-Ni MOF is given in Fig. 4 to evaluate the thermostability of the MOF. According to Fig. 4, the degradation profile of the TMA-Ni MOF possesses two steps at about the temperature ranges of 310e500 and 670e740 K, indicating the solvent (DI water) loss and decomposition of MOF, respectively. Moreover, the as prepared MOF shows approximately 54 wt % of weight loss upon heating to 773 K.
Phase compositions and microstructural characterization of the Mg-MOF-H composite and the pure MgH2 Fig. 5 (a) shows the XRD patterns of the as prepared TMA-Ni MOF and the as milled pure MOFs after different ball milling durations including 4, 8 and 16 h. The characteristic peaks of the TMA-Ni MOF are located at relatively low angles (2q), such as 13.92 , 15.34 and 18.74 . It can be found that peaks of MOF remain unchanged after different ball milling durations, demonstrating that neither crystal structure change nor phase transformation takes place during ball milling at 150 rpm for different durations. The corresponding SEM
Fig. 3 e FTIR spectrum of the trimesic acid (TMA) and the TMA-Ni MOF synthesized in DI water.
Fig. 4 e Thermogravimetric curve of the TMA-Ni MOF.
Fig. 5 e XRD patterns of the pure TMA-Ni MOF powders (a) and Mg-MOF-H composites (b) ball milled for different time.
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micrographs of the as milled pure TMA-Ni MOFs after different ball milling durations (4, 8 and 16 h) are given in Fig. 6 (a)-(c), respectively. It can be clearly observed from Fig. 6 (a) that, some of the TMA-Ni MOF needles crush into small particles after ball milling (4 h), while rest of the individual TMANi MOF remains to be needle-shaped with length of about 1.5 mm. According to Fig. 6 (b) and (c), all of the TMA-Ni MOF crushes into small particles with particle sizes range from 0.1 to 0.5 mm after ball milling for 8 h and the particle size of the as milled MOF cannot be further lowered down with increasing the ball milling duration (16 h). Fig. 5 (b) presents the XRD results of the as milled Mg-MOFH composites for different durations (4, 8 and 16 h). Despite the weak peaks from Mg, the patterns of composites are all indexed as MgH2 without MgO or other phases. It means that there is no reaction or phase transformation during ball milling process with the increase of ball milling duration from 4 h to 16 h at 150 rpm. The SEM images of the ball milled MgMOF-H composites for different ball milling times are shown in Fig. 6(def) and Fig. 6(gei), (j-l) and (m-o) display the EDS
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mapping of Mg, Ni and O, respectively. The distribution of Ni and O on the surface of MgH2 particles are more homogeneous in the Mg-MOF-H powders after ball milling for 8 and 16 h than those of the powder ball milled for only 4 h, implying that the ball milling duration of 8 h at 150 rpm is an optimum ball milling parameter for the preparation of the Mg-MOF-H composite. Fig. 7 (a)-(f) illustrate the XRD patterns of the TMA-Ni MOF, as milled Mg-MOF-H composite (8 h), as milled pure MgH2 (8 h) and the powders after hydriding/dehydriding, respectively. Despite a weak peak from Mg, the patterns of the as milled Mg-MOF composite (Fig. 7 (b)) and MgH2 powders (Fig. 7 (c)) can be identified as pure MgH2 without any other phases. Comparatively, same phases can also be identified in the pattern of the pure MgH2 powder after hydrogenation at 623 K for 2 h (Fig. 7 (d)). According to Fig. 7 (e), the re-hydrogenated Mg-MOF-H powder composed of b-MgH2 as major phase with small quantity of Mg2NiH4, Mg and MgO phases. The presence of remnant Mg demonstrates that some large-sized Mg cannot be completely hydrogenated at 623 K for 2 h.
Fig. 6 e SEM micrographs of the pure TMA-Ni MOF powders (aec) and Mg-MOF-H composites (def) ball milled for different time as well as the corresponding EDS mappings of Mg-MOF-H powders showing the distributions of Mg (gei), Ni (jel) and O (meo).
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Fig. 7 e XRD patterns of the TMA-Ni MOF (a), Mg-MOF-H composite and pure MgH2 in varying states: as milled (b), (c), re-hydrogenated at 623 K (d), (e), and dehydrogenated at 673 K (f), respectively.
Meanwhile, the formation of Mg2NiH4 phase in the hydrogenated Mg-MOF-H composite reveals a reaction of H2, Mg and Ni to generate Mg2NiH4 phase during hydrogenation process. The reaction can be expressed as follows: 2 Mg þ Ni þ 2H2 ¼ Mg2NiH4
(1)
Furthermore, when the re-hydrogenated Mg-MOF-H composite powder were dehydrogenated at 673 K, the XRD pattern shows the appearance of Mg and Mg2Ni phases, which indicates that all of the MgH2 and Mg2NiH4 transform to Mg and Mg2Ni, respectively (see Fig. 7 (f)). Additionally, small amount of MgO is present in re/dehydrogenated Mg-MOF-H composites. The existence of oxygen in MgO and nickel in Mg2NiH4/ Mg2Ni phases suggests the decomposition of the TMA-Ni MOF additive at high temperatures. It can be deduced from Figs. 4 and 7 that the TMA-Ni MOF in Mg-MOF-H composite is getting unstable with the increment of temperature. The structure of the TMA-Ni MOF collapses at 693 K during activation process. The chemical bonds among component elements in MOF structure were broken down and consequently, nickel, oxygen as well as amorphous carbon remain in the composite. The contact area between the MOF additive and MgH2 particles increases during the ball-milling process and part of MOF needles can penetrate the extremely thin MgO layer covered on MgH2 particles. In the following hydrogenation process at 623 K, the large contact area between the MOF and MgH2 makes it easy for nickel and oxygen from the TMA-Ni MOF to react with surrounding Mg under 3 MPa H2 atmosphere to produce Mg2NiH4 and MgO, respectively. The typical SEM images of the as milled, hydrogenated and dehydrogenated MgH2 powders are shown in Fig. 8 (a), (e) and (i), respectively. It can be observed that the average particle size of the pure MgH2 remains at 10 mm, which indicates that the ball-milling process on the given condition (150 rpm, 8 h) cannot prominently reduce the average particle size of the MgH2 powder. According to Fig. 8 (e) and (i), the particle sizes
of MgH2 and Mg both increase from about 10 mm to 30 mm, demonstrating a significant agglomeration of MgH2/Mg particles during re/dehydrogenation processes. Same phenomenon can be found on the Mg-MOF-H composite in varying states (see Fig. 8(bed)). The particle size of the composite powder goes up from about 12 mm (as milled state) to 30e40 mm (re/dehydrogenated states), suggesting that the addition of TMA-Ni MOF additive cannot reduce the average particle size of the Mg-MOF-H composite in varying states. Besides, Fig. 8(feh) and (j-l) display the EDS mapping of Ni and O, respectively. It can be found that, Ni and O are homogeneously distributed on MgH2 (as milled and hydrogenated) and Mg powders (dehydrogenated), indicating the uniform distribution of Mg2NiH4, Mg2Ni and MgO phases on the surface of Mg particles in varying states. It is mainly attributed to the appropriate ball-milling process during the preparation of the composite. The TEM observations were conducted in Fig. 9 to display the micro-morphology of the hydrogenated Mg-MOF-H composite. Fig. 9 (a) presents a bright field image of the Mg-MOF-H sample. It can be observed from Fig. 9(a) that two irregularshaped Mg-MOF-H particles aggregating together and the size of the composite is measured to be 500 nm. Based on the SAED pattern of the Mg-MOF-H composite in Fig. 9 (b), all the diffraction rings corresponding well to MgH2, Mg2NiH4 and MgO phases. It is consistent with the XRD results. The appearance of MgO phase can be ascribed to the reaction between O (provided by the degradation of TMA-Ni MOF) and Mg during (de)hydriding cycles. To show the morphology of Mg2NiH4 phase on MgH2 particles, a dark field (DF) micrograph presented in Fig. 9 (c) is obtained from Mg2NiH4 (220) and MgO (200) diffraction rings. The Mg2NiH4 and MgO nano-sized particles with bright contrasts uniformly disperse on the large-size MgH2. The sizes of Mg2NiH4 nano-particles range from 5 nm to 20 nm. This is in accordance with the EDS mappings of Ni and O in SEM observations. Additionally, High Resolution Electron Microscope (HRTEM) is performed to reveal the morphology of the Mg-MOF-H sample. It is clearly seen that the interplanar spacings of the selected zones in Fig. 9(d) are determined to be about 0.2312 and 0.2291 nm, corresponding to the Mg2NiH4 (220) plane (0.2309 nm).
Hydrogen storage performances of the pure MgH2 and the Mg-MOF-H composite The hydrogen storage performances of the MgH2 and the Mg-MOF-H composite are evaluated by using the PCT facility. Fig. 10 (a) and (c) illustrate the PC isotherms of the pure MgH2 and the Mg-MOF-H composite tested at 598, 623 and 648 K in a H2 pressure ranging from 0.01 to 4 MPa, respectively. The single slant plateaus owing to the redehydrogenation of MgH2 phase are presented under assigned temperatures in Fig. 10 (a). It can also be found that no obvious plateau can be found in the PCT profile during dehydrogenation of the pure MgH2 at each temperature. Furthermore, a sharp diminution of the H2 pressure at the initial stage of dehydriding process can be easily observed in desorption profile at 598 K, indicating the poor dehydriding kinetic property of the pure MgH2. In contrast, the flat plateaus of the Mg-MOF-H composite at each temperature are
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shown in Fig. 10 (c). Each of the H2 desorption curves possesses a distinct plateau and no decline of pressure can be found even the testing temperature decreases to 598 K. It means that the addition of TMA-Ni MOF has prominent catalytic efficiency in promoting the (de)hydriding kinetics of Mg. It should be pointed out that, only one plateau can be observed in each PC isotherms of the Mg-MOF-H sample. In comparison to major phases (Mg/MgH2), neither the quantity of Mg2Ni nor Mg2NiH4 is enough to form another distinct plateau during hydriding/dehydriding cycles [31]. The hydrogen uptake amounts for the Mg-MOF-H powder at 598, 623 and 648 K are 5.09, 5.04 and 5.17 wt %, respectively. The uptake amounts of hydrogen for the composite under series temperatures are all lower than the theoretical capacity of
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MgH2 (7.6 wt %) [3,4], which is ascribed to the addition of TMA-Ni MOF and the formation of MgO during hydrogen de/ absorption cycles. The de-hydrogenation enthalpies of the pure MgH2 and the Mg-MOF-H samples are estimated using the van't Hoff equation. Fig. 10 (b) and (d) reveal the linear fittings (ln P vs. 1000/T) of the van't Hoff plots and subsequently, the enthalpies of the hydriding/dehydriding processes for Mg-MOFH sample are estimated to be 74.3 and 78.7 kJ mol1 H2, respectively. They are fairly close to the dehydrogenated MgH2 (73.8 and 76.1 kJ mol1 H2) as well as reported values of Mg (±74.7 kJ mol1 H2) [32]. With the aim of investigating the hydriding kinetics of the Mg-MOF-H composite, the profiles of hydrogenation for the
Fig. 8 e SEM micrographs of the as milled (a), hydrogenated (e) and dehydrogenated (i) pure MgH2 powders, the as milled (b), hydrogenated (c) and dehydrogenated (d) Mg-MOF-H composite powders as well as the corresponding EDS mappings showing the distributions of Ni (feh) and O (jel).
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dehydrogenated MgH2 and the Mg-MOF-H composite tested at 448, 473, 498, 523, 548 and 573 K under 3 MPa H2 are given in Fig. 11 (a) and (c), respectively. It is worth noting from Fig. 11 (a) and (c) that the disparity of hydrogenation profiles of the dehydrogenated MgH2 and the Mg-MOF-H is getting remarkable with the reduction of the temperature, especially at relative low temperatures including 448, 473 and 498 K. The hydrogen absorption rate of the dehydrogenated pure MgH2 observably decreases at low temperatures and the sample can merely absorb 2.1 wt % of hydrogen at 448 K for 3 h. In contrast, an evident improvement in the H2 uptake rate of the dehydrogenated Mg-MOF-H powder is achieved under low temperatures. The inset of Fig. 11 (c) reveals that the absorption rate of the Mg-MOF-H does not sharply decline with decreasing the temperature, even at 448 K, indicating the much better absorption kinetics with respect to that of dehydrogenated pure MgH2 sample. In addition, Fig. 12 illustrates the changes of the hydriding percentages of the dehydrogenated MgH2 and the Mg-MOF-H samples at low temperatures. The time needed for dehydrogenated pure MgH2 to absorb 80% of its maximum hydrogen capacity is 7430 s at 473 K under 3 MPa H2 pressure, which is much longer than that of the dehydrided Mg-MOF-H sample (2988 s), implying the enhanced hydriding rate of the Mg-MOF-H sample. The beneficial role of Mg2Ni played during hydrogenation process is responsible for the superior H absorption rate of the Mg-MOF-H sample [31,33].
It has been established that the rate controlling factors of the hydrogenation process in metals involve three steps including the hydrogen molecules dissociation, the penetration of H into the bulk metal and the diffusion of H in the metal matrix [34]. An activation energy (Ea) is taken to summarize the energy barrier during hydrogenation. According to the Johnson-Mehl-Avrami-Kolmogorov (JMAK) model, the data can be systematically analyzed and the linear equation of the sorption kinetics can be expressed as: ln½lnð1 aÞ ¼ h ln k þ h ln t
(2)
where the value of a represents the fraction of phase transformation from metal (Mg) to its hydride (MgH2) at time t, k and h donate a parameter of kinetics and the reaction order or Avrami exponent, respectively. With the rate constant k calculated from the equation above, the Ea can be estimated as follows: k ¼ A expð Ea =RTÞ
(3)
where A donates a temperature-independent coefficient, R and T represent the gas constant and the absolutely temperature, respectively. As shown in Fig. 11 (b) and (d), the Ea values in the dehydrogenated MgH2 and Mg-MOF-H powders are evaluated to be 100.7 and 51.2 kJ mol1 H2, respectively. Such an evident reduction in activation energy suggests a
Fig. 9 e A typical bright field TEM micrograph of the Mg-MOF-H powder (a) hydrogenated at 623 K under 3 MPa H2 atmosphere for 2 h, the corresponding SAED pattern (b), the dark field micrograph obtained from Mg2NiH4 (220) and MgO (200) diffraction rings and the corresponding HRTEM image (d) of the Mg-MOF-H composite.
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lower energy barrier of hydrogen absorption, explaining the better hydriding kinetics of the dehydrogenated Mg-MOF-H sample over the dehydrogenated pure MgH2, particularly at relative low temperatures. As shown in Fig. 13, the Mg-MOFH composite can be obtained after ball-milling process and the particle size of the composite powder gradually increases from about 10 mm to 30 mm after re-dehydrogenation cycles at high temperatures (see Fig. 8). Seemingly, the agglomerations of the Mg-MOF-H particles inhibit the hydrogen absorption, which correspondingly results in sluggish sorption kinetics. Nevertheless, Fig. 13 illustrates that the big-sized dehydrogenated Mg-MOF-H particles (average particle sizez30 mm) are comprised of many Mg2Ni-covered Mg particles (average particle sizez10 mm). It has been established that the Mg2Ni on the surface of Mg particles plays a catalytic role on the chemical absorption of H2 and easily dissociate them into H atoms [35]. As is well known, MgH2 layer outside of Mg particles hampers the diffusion of H atoms from surface into Mg bulk. While the Mg2Ni uniformly distributed on the Mg particles can play a “gateway” role as channels to accelerate the diffusion of hydrogen atoms into the bulk of Mg during hydrogenation process [36]. It results in the remarkable improvement of absorption performance (see Fig. 11). Furthermore, Liang et al. [37] reported that Mg2Ni not only played a catalytic role on the chemisorption of hydrogen, but also markedly facilitated the hydrogen
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diffusion rate in Mg. In addition, Liu et al. have pointed out that MgH2 phase preferentially nucleates at the interface between Mg/Mg2Ni during hydriding, through which the absorption kinetics of the Mg-MOF-H composite can be significantly accelerated [31]. In brief, the beneficial effects mentioned above lead to the outstanding absorption kinetic performances of the dehydrogenated Mg-MOF-H composite. To characterize the dehydrogenation behaviors of the MgMOF-H composite, the hydrogen desorption profiles of the MgMOF-H and MgH2 powders at different temperatures are presented in Fig. 14 (a) and (b), respectively. It can be found that the dehydriding rates of the Mg-MOF-H powders are all much higher than those of pure MgH2 powder at various temperatures. The hydrogen desorption capacity of the MgH2 sample sharply declines with the decreasing temperature. The MgH2 sample releases merely 0.3 wt % H2 at 573 K while the MgMOF-H composite still exhibits a desorption capacity of about 3 wt % H2 at the same temperature, implying the faster dehydriding kinetics of the composite than that of the pure MgH2 powder. The DSC curves of MgH2 and Mg-MOF-H composite are given in Fig. 14 (c) and (e), respectively. As found in Fig. 14 (c), a strong endothermic peak can be observed in each DSC curve at different heating rates in the pure MgH2 sample. The endothermic peak is due to the degradation of b-MgH2. However, it is interesting to note that the DSC profiles of the Mg-MOF-H sample (see Fig. 14 (e)) possesses two endothermic
Fig. 10 e PC isotherms and the van't Hoff plots of the pure MgH2 (a), (b) and the Mg-MOF-H composite (c), (d) tested at 598, 623 and 648 K (Ab: absorption; De: desorption).
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peaks at the heating rate of 10 K/min, demonstrating a twostep dehydrogenation process. Similar phenomenon on the DSC results of the Mg-Ni composite has also been reported in the previous work [36]. According to the XRD patterns in Fig. 5 and the researches on the Mg-Ni powders [36e38] and Mg2Ni compound [39], the weak endothermic peak at relatively low temperature owing to the dehydrogenation of Mg2NiH4 phase. The strong broader peak corresponds well to the decomposition of b-MgH2. It should be noticed that, the first endothermic peak mentioned above gets unconspicuous with decreasing the heating rate and even disappears when the heating rate declines to 3 K/min. The interpretation might be that all endothermic peaks get weak with descent of the heating rate from 10 to 3 K/min, leading to a further weakening of the first small endothermic peak. The subdued small peak may be involved into the strong endothermic peak at the heating rate of 3 K/min, making it difficult to identify the weak peak of the decomposition for Mg2NiH4. It can be summarized from Fig. 14 (c) that peak temperatures of DSC profiles of the pure MgH2 powder at the heating rates including 3 K/min, 5 K/min and 10 K/min are 677.1, 688.4 and 701.8 K, respectively. Whereas, the peak hydrogen desorption temperatures of the Mg-MOF-H powder at the same heating rates are prominently reduced to 545.9, 556.3
Fig. 12 e Hydrogen absorption percentages (%) of the dehydrogenated pure MgH2 and the Mg-MOF-H composite measured at relative low temperatures including 448, 473 and 498 K under 3 MPa H2 pressure for 3 h (De: dehydrogenated).
Fig. 11 e Hydrogen absorption profiles and the corresponding ln k-1000/T plots of the dehydrogenated pure MgH2 (a), (b) and the Mg-MOF-H composite (c), (d) measured under 3 MPa H2 pressure for 3 h.
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Fig. 13 e Schematic illustration showing the preparation of the Mg-MOF-H composite and the reactive details during the hydrogenation process.
Fig. 14 e Hydrogen desorption profiles of the Mg-MOF-H composite (a) and the pure MgH2 (b), DSC profiles and the corresponding ln (b/T2p)-1000/Tp plots for the pure MgH2 (c), (d) and the Mg-MOF-H composite (e), (f) at different heating rates.
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Fig. 15 e Cycling curves of the Mg-MOF-H composite at 623 K for 10 times.
and 568.8 K, respectively. Furthermore, the onset temperature for dehydriding is significantly declined to 497.2 K, which is 167.8 K lower than that of pure MgH2 sample (665 K), as found in Fig. 14 (e). The dehydrogenation activation energy (Ed) can be evaluated to ulteriorly investigate the excellent dehydriding performance of the Mg-MOF-H composite. The activation energy of dehydriding process can be calculated by the Kissinger equation: . ln b Tp 2 ¼ A Ed RTp
(4)
where b donates the heating rate, Tp is the peak dehydriding temperature, A is a linear constant and R is the gas constant. The Ed of the pure MgH2 is calculated to be 181.4 kJ mol1 H2 with the linear fitting the plots of ln (b/T2p) vs. 1000/Tp in Fig. 14 (d) and (f). The Ed of MgH2 can be effectively declined to 126.7 kJ mol1 H2 with the addition of TMA-Ni MOF and consequently, leading to an outstanding improvement on desorption kinetic performance of the Mg-MOF-H. During the dehydriding of the Mg-MOF-H sample, Mg2NiH4 nanoparticles disperse on the surface of MgH2 firstly decomposes to Mg2Ni and H2, followed by the dehydrogenation of MgH2. This has been demonstrated by the DSC measurement. Zalusk et al. reported a synergetic effect from the mixture of Mg2NiH4 and MgH2 to enhance the desorption kinetics during dehydriding. Results in the literature revealed that the addition of Mg2NiH4 by ball-milling destabilized the magnesium hydride and consequently, led to the superior desorption kinetic property [39]. Also, those Mg2NiH4 coated homogeneously on MgH2 particles act as channels for hydrogen diffusion from magnesium hydride matrix [36,40]. It is necessary to point out that the magnesium oxide possesses a beneficial effect of grain boundary pinning, preventing the growth of Mg grains at elevated temperatures [41]. Meanwhile, the catalytically active sites which is attributed to the special vacancies on the MgO, make it easier to recombine atomic hydrogen during hydrogen desorption of MgH2 [42]. In short, the nanoscale Mg2NiH4 on the surface of MgH2 particles as well as the existence of magnesium
oxide, contribute to the improved dehydrogenation performance of the Mg-MOF-H composite. To investigate the effect of the TMA-Ni MOF additive on the cycling performance of MgH2, continuous hydriding/dehydriding tests of the Mg-MOF-H composite are performed at 623 K under 3 MPa H2 for hydrogenation and in vacuum for dehydrogenation. Fig. 15 illustrates the hydrogen absorption and desorption curves of the composite for 10 cycles. As seen in Fig. 15, no evident hydrogen storage capacity loss is detected after 10 cycles, indicating the superior cycling stability of the Mg-MOF-H composite. In addition, the hydrogen absorption/desorption kinetics of the composite remains unchanged and the composite possesses a constant reversible capacity of 3.4 wt % at 623 K. The results gathered in this study have shown the efficient catalytic effect of TMA-Ni MOF on hydrogen sorption properties of Mg/MgH2. As compared to additives of Mg2NiH4 or Ni nanoparticles [43e45], TMA-Ni MOF can be introduced into MgH2 through simple ball milling with fairly short duration and homogeneous distribution on the MgH2 particles. In addition, SEM and TEM observations (see Figs. 8 and 9) confirm the formation of nano-sized Mg2NiH4 coated on the surface of MgH2 particles without severe agglomeration after hydrogen sorption cycles. Furthermore, a relatively low amount of TMANi MOF addition can significantly improve the hydrogenation/ dehydrogenation performances of Mg/MgH2, revealing its high catalytic effectiveness.
Conclusions In this work, the TMA-Ni MOF was successfully synthesized for the first time in an environmentally friendly media, which was introduced into MgH2 to form the Mg-MOF-H composite through ball milling. The main conclusions can be drawn as follows: 1) XRD results show the formation of Mg2NiH4 phase when the dehydrogenated Mg-MOF-H composite is rehydrogenated and the Mg2NiH4 phase completely decomposes to Mg2Ni when the sample is dehydrided, indicating that Ni in the MOF is attracted by Mg to form Mg-Ni compound during de/re-hydrogenations. 2) SEM and TEM observations reveal that the nano-sized Mg2NiH4 is uniformly coated on the surface of large MgH2 particles in the Mg-MOF-H composite. 3) The hydriding and dehydriding enthalpies of the Mg-MOFH composite are determined to be 74.3 and 78.7 kJ mol1 H2, respectively. These values are close to the standard values of Mg (±74.7 kJ mol1 H2). 4) The dehydrogenated Mg-MOF-H composite exhibits enhanced hydrogen absorption kinetics as compared to the dehydrogenated pure MgH2. The Ea of the dehydrogenated Mg-MOF-H powder is estimated to be 51.2 kJ mol1 H2. 5) The hydrogen desorption temperature of the Mg-MOF-H composite is notably reduced in comparison to that of the pure MgH2. The onset dehydriding temperature at the heating rate of 10 K/min is 167.8 K lower than that of the pure MgH2. The corresponding Ed for the composite is calculated to be 126.7 kJ mol1 H2.
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6) The nanoscale Mg2Ni/Mg2NiH4 coated on the Mg/MgH2 powders dramatically enhances the hydrogen storage kinetics of the Mg-MOF-H composite by providing “gateway” for hydrogen diffusion during hydriding/dehydriding processes.
Acknowledgment Prof. Zou would like to thank the support from the National Science Foundation (No. 51771112), the Shanghai Science and Technology Commission under No. 14JC1491600 and Shanghai Education Commission “Shuguang” scholar project (16SG08).
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