Synthetic nanocomposite MgH2/5 wt. % TiMn2 powders for solid-hydrogen storage tank integrated with PEM fuel cell

Storing hydrogen gas into cylinders under high pressure of 350 bar is not safe and still needs many intensive studies dedic ated for tank’s manufacturing. Liquid hydrogen faces also severe practical difficulties due to its very low density, leading to larger fuel tanks three times larger than traditional gasoline tank. Moreover, converting hydrogen gas into liquid phase is not an economic process since it consumes high energy needed to cool down the gas temperature to −252.8 °C. One practical solution is storing hydrogen gas in metal lattice such as Mg powder and its nanocomposites in the form of MgH2. There are two major issues should be solved first. One related to MgH2 in which its inherent poor hydrogenation/dehydrogenation kinetics and high thermal stability must be improved. Secondly, related to providing a safe tank. Here we have succeeded to prepare a new binary system of MgH2/5 wt. % TiMn2 nanocomposite powder that show excellent hydrogenation/dehydrogenation behavior at relatively low temperature (250 °C) with long cycle-life-time (1400 h). Moreover, a simple hydrogen storage tank filled with our synthetic nanocomposite powders was designed and tested in electrical charging a battery of a cell phone device at 180 °C through a commercial fuel cell.

Hydrogen is an energy carrier, which holds tremendous promise as a new clean energy option 1,2 . It is a convenient, safe, versatile fuel source that can be easily converted to a desired form of energy without releasing harmful emissions 3,4 . A key advantage of hydrogen is that when burned, carbon dioxide (CO 2 ) is not produced. Hydrogen storage is one of the most crucial difficulties restricting utilization of hydrogen energy for real applications. However, storage hydrogen gas into gas cylinders under high pressure reached to 350 bar is a well-known technology, using such pressurized hydrogen gas tanks as a source of fuel in vehicles is not safe at present and still needs many intensive studies dedicated for improving the structural and mechanical properties of the materials used in tank's manufacturing. Extreme conditions for on-road vehicle service should be defined, to demonstrate performance of storage systems demonstrated both under the stresses of normal vehicle operation and under externally imposed stresses 5 . Likewise pressurized hydrogen gas, liquid nitrogen possesses many difficulties related to its very low density. Accordingly, the size of liquid hydrogen requires larger tanks reaches to about three times larger than the traditional gasoline tank 6 . Practically, converting hydrogen gas into liquid hydrogen is not an economic process since it consumes large amount of energy required to cool down the gas temperature to −252.8 °C. For instance, liquefying 1 kg of hydrogen gas in medium-size plant requires 10 to 13 kWh of electrical energy 7 . Moreover, liquid nitrogen is not safe since it has a high flammability range. Boil-off losses associated with the storage, transportation and handling of liquid nitrogen can consume up to 40% of its available combustion energy 6,8 .

Solid Hydrogen
Apart from gaseous and liquidus phases of hydrogen, solid hydrogen has been considered as the most reliable and safe practical solution for providing clean energy required for different applications, using proper fuel cells such as proton-exchange fuel cells membrane (PEM) 9 . Hydrogen can be simply stored in nanocrystalline metal

Solid Hydrogen Tanks
Merits of using solid hydrogen tanks are obvious for those who live in isolated communities having lack in connection to the public power grid. However, renewable energy sources that generate electricity directly through for example photovoltaics (PVs) and windmills are environmentally friendly and can provide cost-effective solutions, their energy availability varies drastically from time to day per one day. Accordingly, energy storage is necessary to meet the electricity demand with the required reliability 34 . Thus, the combination of stored solid hydrogen generated from excess electricity and a fuel cell is a promising solution 35 .
In general solid hydrogen storage tanks differ in their design and materials used when compared with those used to store pressurized hydrogen gas and liquid hydrogen. This because diffusion of hydrogen atoms into metal lattice (e.g. Mg) does not require the application of very high pressure since the gas-solid exothermic reaction between the two phases takes place simultaneously at relatively low pressure (below 15 bar/300 °C). Moreover, the need of using expensive cryotanks with controlled pressure (similar to those used for liquid hydrogen) is not required. Accordingly, the basic process for design, materials selections, and manufacturing of vessels or tanks contain metal hydride, as a source of hydrogen becomes inexpensive in materials science point of view. In fact, the success of using solid hydrogen tank for providing hydrogen to operate an electrical system with constant rate of hydrogen flow through a suitable fuel cell can be firstly attributed to the hydrogenation/dehydrogenation kinetics behavior and cyclability of the metal hydride materials stored in the tank.
Most the authors published fine articles related to the Mg-based hydrogen storage materials, focused mainly on materials preparations, characterization and investigation the hydrogenation/dehydrogenation behaviors of the metal hydride systems. However, there are some studies reported interesting and promising results related to designing and utilizing of solid hydrogen tank. Of these, Zu et al. 36 proposed an advanced design that can be used for real manufacturing of hydrogen storage tank. However, their design may improve the geometric flexibility and structural performance of composite toroidal hydrogen storage tanks; they have not examined the system with any types of metal hydride to ensure the validity of their tank model for hydrogenation/dehydrogenation processes. A numerical work proposed by Gkanas et al. 37 showed a three-dimensional computational model regarding coupled heat and mass transfer during both the hydrogenation and dehydrogenation process. In this study, two different types of hydrides, LaNi 5 and an AB2-type (Ti-Zr-Mn intermetallic system) were selected as possible candidate materials. Their results showed the possibility of hydrogen gas uptake and release with moderate kinetics at 20 °C and 15 bar of H 2 . A very interesting and promising model was recently introduced by Gattia et al. 38 . They used compacted powders prepared by ball milling MgH 2 with Nb 2 O 5 with mixed with 5 wt.% ENG. The pellets were coated by different type of metal coating, using sputtering and thermal evaporation techniques. The material were tested by inserting the pellets into a cylindrical tank. Unfortunately, the kinetics of the system did not show attractive characteristics since the temperature required for cyclic procedure taken place at 310 °C.
On the other hand, significant number of patents related to designing hydrogen storage tanks are available. One of the earliest invention related to this regard dates back to 1969, when Lyon et al. proposed a process and an apparatus used to utilize MgH 2 as a source of hydrogen for useful fuel cell applications 39 . In their prototype, Mg reacted with hydrogen gas to form MgH 2 and then decomposed at high temperature (277 °C to 649 °C) under pressure ranging from 1 bar to 207 bar 39 . Applications of such high temperatures and pressure are consider as drawback of their system. In 2014, Ornath 40 , introduced an interesting idea for manufacturing a hydrogen storage tank that can lead to lower the temperature of hydrogen uptake and release. Unluckily his patent did not contain any experimental results to prove his idea. Details and more information on the employing different types of metal hydride materials on solid-state hydrogen storage are recently reviewed and published by Rusman and Dahari et al. 41 .
Apart from the interesting and useful solid hydrogen storage models invented by many authors, our present has been addressed to satisfy two objectives; the first is focused on synthesizing and characterizing a new nanocomposite MgH 2 /5 wt. % TiMn 2 system with advanced kinetics behaviors and cyclability, where the second objective is focused on utilizing the as-synthesized nanocomposite powders in a self-manufactured hydrogen storage tank interfaced to 40 W/4.5 A proton exchange membrane (PEM) fuel cell. This system feeds the PEM-fuel cell with the released hydrogen gas form the tank. The converted electrical energy was utilized to charge the battery of a cell phone device through 5 V voltage regulator. As far as the authors know, this is the first time to examine the validity of a new Mg-based nanocomposite in real applications.

Results and Discussion
Structure. The general and local structure beyond the nano level for the nanocomposite MgH 2 /5 wt. % TiMn 2 powders were investigated by means of X-ray diffraction (XRD) and field emission-high resolution transmission electron microscope (FE-HRTEM) techniques, respectively. The XRD pattern of elemental Mg powders (precursor) obtained after 200 h of RBM under 50 bar of a hydrogen gas atmosphere is shown in Fig. 1(a). The powders consisted of fine nanocrystalline mixture of β-MgH 2 (PDF file# 00-012-0697) and γ-MgH 2 (PDF file# 00-035-1184) phases, implied by the broad Bragg peaks shown in Fig. 1(a). Low intensity Bragg peak related to fcc-MgO (PDF file# 01-079-9866) was detected at scanning angle of about 42° due to the oxidation of the sample during XRD sample preparation outside of the glove box.
Based on the purpose of the present study, the end-product (200 h) of as-synthesized MgH 2 powders was manually doped with 5 wt. % TiMn 2 powders inside the glove box and then charged into the milling vial together with the milling media. The vial was then pressurized with 50 bar of hydrogen gas and mountained onto a high-energy ball mill where the RBM process taking place for different milling time. After 3 h of milling, the powders revealed Bragg peaks related to MgH 2 phase, where the Bragg peaks corresponding to TiMn 2 powders were hardly seen, as shown in Fig. 1(b). This may be attributed to the effect of cold welding generated by the milling process on TiMn 2 powders, leading to sticking the powders onto the milling media (vial's internal wall and balls). Further ball milling time led to fragmentation of TiMn 2 agglomerate powders to fall into the vial's milling zone and then milled together with MgH 2 powders. After 12.5 h, the milled powders consisted of ultrafine particles of MgH 2 coexisted with TiMn 2 (PDF file# 01-071-9712) powders, as implied by the broad Bragg peaks related to both phases ( Fig. 1(c)). This broadening manifested in the Bragg peaks raised from both refinement of the MgH 2 and TiMn 2 crystallites and accumulated macrostrain during the RBM process. It should be notified that neither reacted powders related to the formation MgTiH x , and TiH 2 42 phases nor elemental phases corresponding to metallic Mg and Ti could be detected even after longer RBM time, ranging between 37.5 h ( Fig. 1(d)) and 50 h ( Fig. 1(e)). Moreover, the Bragg-peaks of metallic hcp-TiMn 2 maintained their peak positions after 50 h of RBM, implying the absence of solubility into the MgH 2 lattice, as elucidated in Fig. 1(e).
The bright field image (BFI) of nanocomposite MgH 2 /5 wt.% TiMn 2 powders obtained after 50 h of RBM time is displayed in Fig. 1(f). The nanopowers obtained after this stage of milling were aggregated to form larger particles due to van der Waals forces 43 containing fine nano-dimensional dark-grey lenses related to TiMn 2 phase embedded into the light-grey MgH 2 matrix, as shown in Fig. 1(f). The FE-HRTEM image of zone I indexed in Fig. 1(f) is displayed in Fig. 1(g). The nanocomposite powders revealed Moiré-like fringes with nanocrystalline-structure morphology (5 nm to 17 nm). Moreover, the lattice fringes of MgH 2 (β and γ phases) and hcp-TiMn 2 were regularly separated with an interplanar spacing (d) matching well with the reported PDF files cited in Fig. 1(a). Based on careful analysis performed of at least 30 examined zones for three individual samples, we could not detect the existence of any other phase(s). This implies the formation of binary nanocomposite MgH 2 /5 wt.% TiMn 2 powders. A different examined zone located at the top edge of the particle, zone (II) was analyzed by HRTEM ( Fig. 1(h)). Two selected zones (III, and IV) in Fig. 1(h) were selected to get their corresponding nano beam diffraction patterns (NBDPs). The grains located in each zone contained ultrafine nano-grains ranging in size between 5 to 8 nm, as displayed in Fig. 1(h). It can be notified that the grains did not reveal specific orientation ( Fig. 1(h)) and the development of dislocation tangles into sub-boundaries can be seen clearly in zones III and IV ( Fig. 1(h)). The dislocations presented inside the individual grains suggesting the nanostructure development by severe plastic deformation created by the ball milling media.
The NBDPs taken from the zone III and IV elucidated in Fig. 1(h) are shown in Fig. 1(i) and (j), respectively. Both electron diffraction revealed Debye continuous rings corresponding to βand γ-MgH 2 phases, plus hcp-TiMn 2 . The present of these diffracted rings indicates the random orientation of the grains in the nanocomposite powders ( Fig. 1(i)). In addition, the circular-like NBDPs of MgH 2 ( Fig. 1(j)) were overlapped with numerous closely spaced spots, each from diffraction of a single crystallite for TiMn 2 .
Morphology and elemental analysis. In order to understand the distribution effect of TiMn 2 powders into MgH 2 matrix on the kinetics behaviors and thermal stability of the hydride phase, carful energy-dispersive X-ray spectroscopy (EDS) elemental mapping experiments were conducted for the powders obtained after the early (1-3 (Fig. 2(f)). The EDS analysis of the powders obtained after the early stage of milling indicated that the concentration of the kinetic-modifier phase (TiMn 2 ) was widely varied from particle to particle and even within the particle itself. The local chemical analysis of TiMn 2 measured with an average area of 0.3 μm was in the range between 0.7 and 64 wt. %.
Increasing the RBM time (25 h) led the large TiMn 2 flaky particles to be disintegrated into smaller particles, ranging in size between 10 to 150 nm in diameter, as shown in Fig. 2(i-l). However, these metallic particles are heterogeneously embedded into the MgH 2 matrix (Fig. 2(j)). The TiMn 2 concentration in the host MgH 2 matrix was significantly improved, as indicated by the near concertation values varied from 3.6 to 7.8 wt. %. Toward the end of the RBM time, the TiMn 2 particles (dark lenses presented in Fig. 2(m)) revealed significant changes in their shapes and possessed spherical-like morphology with fair distributions (Fig. 2(o), and (p)) into MgH 2 matrix ( Fig. 2(n)). This morphological improvement of TiMn 2 particles was followed by a dramatic deduction in their sizes laid within a narrow size distribution range between 5 nm to 27 nm ( Fig. 2(m,o,p)). Accordingly, and based on EDS local analysis of different locations in several particles, the concentration of TiMn 2 in the MgH 2 matrix was outstandingly improved to be in the range between 4.8 to 5.3 wt. %. Thermal stability. Differential scanning calorimetry (DSC) performed at a constant heating rate of 20 °C/ min under helium gas flow of 100 ml/min was employed to investigate the effect of RBM time and TiMn 2 additive on the decomposition temperature (dehydrogenation temperature at normal pressure) of as-synthesized MgH 2 powders. The DSC trace of as-synthesized MgH 2 powders obtained after 200 h of RBM revealed a broad endothermic event laid at a peak temperature of 728 K, as shown in Fig. 3(a). This endothermic peak notably shifted to the lower temperature side (697 K) upon milling for short time (3 h), as displayed in Fig. 3(b). The peak decomposition temperature for those samples obtained after 6 h ( Fig. 3(c)) and 12.5 h ( Fig. 3(d)) were 671 K and 628, respectively. The tendency of MgH 2 powders to be decomposed at lower temperature (581 K) with increasing the RBM time was continuously occurred during the intermediate stage of RBM (25 h), as displayed in Fig. 3(e). During the final stage of RBM (37.5 h) the decomposition temperature was retreated to 546 K, as elucidated in Fig. 3(f). This value did not remarkably changed (541 K) for the nanocomposite powders obtained after 50 h of RBM time (Fig. 3(g).

Apparent activation energy of dehydrogenation.
In order to realize the effect of doping MgH 2 with 5 wt. % TiMn 2 powders on the apparent activation energy (E a ) of the metal hydride phase, individual DSC experiments were conducted with different heating rates (5, 10, 20, 30 and 40 °C/min). In this study the effect of RBM on E a were investigated for all samples obtained after different milling stages. Figure 4(a) to (c) displayed selected DSC curves conducted at different heating rated (k) for 3 samples obtained after selected RBM times (3 h, 37.5 h, and 50 h). All the scans revealed single endothermic events related to the decomposition of MgH 2 into metallic Mg and hydrogen gas, as confirmed by XRD technique. While the peak height increased proportionally with the increasing the k from 5 °C/min to 40 °C/min, the peak temperatures (T p ) were significantly shifted to the higher temperature side, as presented in Fig. 4(a-c)). The E a of dehydrogenation related to each sample was calculated according to the Arrhenius equation 44 : where k is a temperature-dependent reaction rate constant, R is the gas constant, and T is the absolute temperature. The E a values were determined by measuring the T p corresponded to the different k and then plotting ln(k) versus 1/T p . The E a values were then obtained from the slope of line (−E/R, where R is the gas constant). Based on these measurements, the nanocomposite powders obtained after 3 h of RBM time showed a high E a value (138.03 kJ/mol), as shown in Fig. 4(d). This indicates a high thermal stability of the powders against decomposition. In contrast, E a of nanocomposite MgH 2 /5 wt.% TiMn 2 powders obtained after 37.5 h (Fig. 4(e)) of RBM showing a lower value (117.76 kJ/mol), indicating a significant destabilization of the MgH 2 upon high-energy ball milling with TiMn 2 phase. The E a value did not show notable improvement (116.99 kJ/mol) upon increasing the RBM time to 50 h, as elucidated in Fig. 4(f). The apparent E a of our system is closed to the reported one for MgH 2 /5 wt. % Zr 70 Ni 30 Pd 10 powders (92 kJ/mol) 4 . However, our system showed lower E a value when compared with pure MgH 2 45 (164 kJ/mol) 40 51 , and MgH 2 -Ta 2 O 5 (74 kJ/mol) 52 systems. Figure 5 summarizes the DSC experiments by presenting the effect of RBM time on T p ( Fig. 5(a)). It is well established that high-energy ball milling process is a typical destabilization process 11,17,18,53 generating several lattice imperfections (e.g. dislocations, staking faults, point defects, etc.) in the milled powders 10 . In thermodynamics point of view, the existence of these defects are favorable for destabilize the stable MgH 2 phase into a metastable phase [45][46][47] , where the hydrogen gas can be released at lower temperature values. As the milling time increases from 0 h to 12.5 h (early stage of RBM), the grain size of MgH 2 was drastically decreased from 231 nm to 84 nm, as displayed in Fig. 5(b). The T p was not affected by such deduction on the MgH 2 grain size only, but in fact it was greatly affected by the continuous morphological improvement taken place with increasing RBM time, as discussed in Fig. 2. The morphological effect including grain size refining on improving the hydrogenation/dehydrogenation kinetics of MgH 2 binary system was demonstrated by many authors [48][49][50][51] . During this early stage of RBM time, the T p was dropped from 725 K to 628 K, as shown in Fig. 5(a). Moreover, the continuous size reduction happened in this early stage of milling led to a notable improvement in E a that tended to decrease from 162 (0 h) to 130 kJ/mol (12.5 h), as elucidated in Fig. 5(b). The influence of grain size reduction on E a and kinetics is well established and attributed to the short-diffusion bathes of hydrogen atoms 27 . We should emphases that intermetallic TiMn 2 hard particles did not only played a heterogeneous catalytic role for improving the poor kinetics of MgH 2 , but also considered as micro-milling media leading to accelerate the grain size refinement of the metal hydride powders.
During the intermediate stage of RBM (25 h) the MgH 2 grain size was reduced to 38 nm (Fig. 5(b)), where T p and E a reached to lower values of 582 K (Fig. 5(a)) and 124 kJ/mol (Fig. 5(b)), respectively. Toward the end of the RBM processing time (37.5 h-50 h) the MgH 2 grain continued their tendency to be reduced in sizes (14 nm-7 nm) according to the double successive effects of balls-and TiMn 2 -micro milling media. Further destabilization of MgH 2 taken place during this final stage of RBM, referred by the reduction occurred in T p values (546 K-540 K), as shown in Fig. 5(a). Most importantly and as a result of MgH 2 grain refinement, Ea tended to drop into a lower value (116-117 kJ/mol), as elucidated in Fig. 5(b). Hydrogenation/dehydrogenation kinetics behavior. The improvement of kinetics related to hydrogen absorption and desorption processes attained upon mechanical doping MgH 2 with 5 wt. % TiMn 2 powders were monitored after different RBM time, using Sievert's method 45 .
Kinetics of hydrogenation. The hydrogenation kinetics behavior investigated at 250 °C under 10 bar H 2 gas pressure for nanocomposite MgH 2 /5 wt.% TiMn 2 powders obtained after selected RBM time are shown in Fig. 6(a). The shaded zone indexed in Fig. 6(a) is elucidated in Fig. 6(b) with a different time scale. After 25 h of RBM, the sample absorbed 3.5 and 4 wt. % H 2 after 1.25 and 4 min, respectively as shown in Fig. 6(b). As previously discussed in Fig. 5(b), increasing the RBM time resulting a significant decreasing in the MgH 2 grain size, allowing an easier hydrogenation process. Accordingly, the sample obtained after 37.5 h of RBM time showed a significant development on absorbing 5.1 wt. % H 2 after only 1.25 min (Fig. 6(b)). Marginal improvement (5.4 wt. % H 2 ) was notified with increasing the absorption time to 3 min ( Fig. 6(a)). Further increasing of the absorption time (~7.75 min) did not cause any improvement on hydrogen uptake and the sample saturated at 5.4 wt. % H 2 , as elucidated in Fig. 6(a). However, longer RBM time (50 h) did not show a major beneficial effect on the hydrogen uptake (5.4 wt. % H 2 ) and hydrogenation kinetics, the sample possessed a notable ability to absorb higher H 2 value with shorter time (~4.7 wt. %/0.6 min), as displayed in Fig. 6(b). This value is higher than the one shown by  56 .
The XRD pattern of 50 h sample taken after hydrogenation process at 250 °C is shown in Fig. 6(c). The sample revealed a domain structure of γ-MgH 2 coexisted with small molecular fractions of γ-MgH 2 . Moreover, the XRD analysis did not confirm the formation of any intermediate or reacted phases (MgTi, TiH 2 , MgTiH) where the TiMn 2 additive powder maintained its original hcp-structure (PDF file #01-071-9712), as displayed in Fig. 6(c). This implies that TiMn 2 additive can be considered as heterogeneous catalysts, in which it enhanced the hydrogenation/dehydrogenation kinetics of MgH 2 without entering into a chemical reaction with MgH 2 matrix.
Kinetics of dehydrogenation. In order to get the complete picture of TiMn 2 effect on enhancing the behavior of MgH 2 powders, The kinetic of hydrogen releasing for the samples obtained after 25 h, 37.5 h, and 50 h of RBM time were carefully investigated at 250 °C at 200 mbar H 2 gas pressure. Figure 6 Fig. 6(e). Significant improving on the hydrogenation kinetics was attained for 50 h sample that desorbed −1.9 wt. % H 2 within 2 min (Fig. 6(e)). In contrast, a moderate improvement could be notified for the sample obtained after 37.5 h and 25 h that succeed within 2 min to discharge −1.  The crystal structure related to 50 h sample obtained after the dehydrogenation test was examined by XRD. The sample revealed sharp Bragg peaks corresponding to hcp-Mg (PDF file #00-004-0770) coexisted with fine hcp-TiMn 2 (PDF file #01-071-9712) particles having broad Bragg peaks patterns, as shown in Fig. 6(f). Neither reacted nor intermediate phases could be detected, implying the absent of any undesirable reactions during the dehydrogenation process.

Cyclability of hydrogen absorption/desorption.
Measuring the cycle-life-time, which reflects the capability and performance of synthesized nanocomposite MgH 2 /5 wt.% TiMn 2 powders for achieving continuous hydrogenation/dehydrogenation cycles was investigated. For comparison, the cycle-life-time test of pure MgH 2 powders obtained after 200 h of RBM time (before doping with TiMn 2 powders) was conducted and the results are displayed in Fig. 7(a). To ensure the cyclic continuity and to improve the of hydrogenation/dehydrogenation kinetics of pure MgH 2 , the test was achieved at high temperature (300 °C). The pressures used for hydrogenation/ dehydrogenation processes were 8/0.2 bar, respectively.
However, the as-prepared nanocrystalline MgH 2 powders possessed excellent hydrogen storage capacity of about 6.9 wt. % ( Fig. 7(a)), the powders failed to maintain their capacity. This is indicated by a monotonic degradation on the storage capacity upon increasing the cycle-life-time, reaching a lower value (6.3 wt. % H 2 ) after only 75 h, as shown in Figs 7(a) and 8(a). When the powders were subjected to longer cycle-life-time in the range between 75 −275 h (Figs 7(a) and 8(a)), the hydrogen storage capacity was severely degregated to reach a lower value of 5.4 wt. % H 2 , after 200 h (Fig. 8(a)). Toward the end of the test (275 h), the MgH 2 powders were disable to store more than 5.24 wt. % H 2 , as elucidated in Fig. 7(a). The tendency of nanocomposite MgH 2 /5 wt.% TiMn 2 powders to manifest hydrogen storage degradation was almost absent even after 1400 h of continuous hydrogen uptake/discharge, as presented in Figs 7(a) and 8(b).
In order to realize the reasons responsible for such severe degradation, the morphology of the material powders obtained after 275 h of cycle-life-time (Mg powders) were examined by FE-SEM technique. Noticeably, the metallic powders after this test tended to agglomerate and composite larger particles with average size of 13 μm in diameter, as presented in Fig. 8(c). This agglomeration behavior was a result of the continuous applications of rather high hydrogen pressure on the powders at a high temperature (300 °C). Moreover, the topology of the powder was solid, where the pores facilitating the hydrogen diffusion into/out the powder were absent (Fig. 8(c)). The absent of favorable pores and cavities degrade the performance of MgH 2 powders to maintain their hydrogen storage capacity. Unluckily, and according to the application of high temperature and pleasure the metallic Mg grains revealed severe grain growth with an average diameter of 80 nm, as shown in the STEM-dark field image (DFI) in Fig. 8(e).
The beneficial effect of doping nanocrystalline MgH 2 with 5 wt. % TiMn 2 powders on the hydrogenation/ dehydrogenation cyclability conducted at 250 °C under hydrogenation/dehydrogenation pressure of 8/0.2 bar is clearly realized in Fig. 7(b). The nanocomposite powders possessed very high performance for achieving 1400 cycle-life-time (~1000 continuous cycles) without severe degradation, as displayed in Fig. 7(b). During the first 400 h, the sample maintained its original hydrogen storage capacity (5.4 wt.%), that was slightly decreased to about 5 wt. % H 2 after 900 h, as shown in Fig. 7(b). Further unserious decreasing can be seen after 1200 h when the hydrogen storage capacity dropped a little to the level of 4.95 wt. % H 2 , as presented in Fig. 7(b). The last 200 h of the cycle-life-time test is individually displayed in Fig. 8(b). During this last stage of the test, the nanocomposite powders maintained their hydrogen storage capacity at almost a constant value of 4.95 wt. %, as elucidated in Fig. 8(b). Moreover, the kinetics of hydrogenation/dehydrogenation processes remaining constant with no obvious failure or decay.
Since milling MgH 2 with TiMn 2 powders did not lead to the formation of any reacted intermediate phase (Fig. 6(c) and (f)) responsible for enhancing the kinetics and performance of charging/discharging hydrogen cyclability, the justification beyond such obvious improvements can be related to a morphological reason.  ranging in sizes between 0.22 μm to less than 0.5 μm in diameter, as shown in Fig. 8(d). In addition, TiMn 2 fine particles (~5 nm in diameter) maintained their tendency for adhering onto the MgH 2 surface of powders, even after 1400 h of cycle-life-time, as shown in the STEM-DFI presented in Fig. 8(f). Comparing the STEM-DFI presented in Fig. 8(d) with that one for nanocomposite powders before conducting the cycle life-time test (Fig. 2(mp)) led us to confirm the absence of any undesired grain growth for both MgH 2 grain matrix and TiMn 2 particles. The limited grain growth seen in MgH 2 grains is probably attributed to the distribution of TiMn 2 particles (grain-growth inhibitors) into the MgH 2 matrix. It was pointed out by Yao et al. 57 that hydrogen diffusion is much faster when Mg grains were in the nanoscale level.
Based on these results, it can be concluded that one advantage of using hard intermetallic compounds such as TiMn 2 , refractory metal carbides and amorphous alloys for improving the kinetics behavior and cycle-life-time of MgH 2 powders is related to the tendency of these hard particles to surround MgH 2 grains leading to block their attitude of growing at high temperature and pressure. Another advantage of TiMn 2 intermetallic compound additives with their nano-dimensional particle size is the improvement of the cyclability performance for MgH 2 system. This present system can be considered as one of the most reliable performance system among the most well known system, exemplified by MgH 2 -5 wt. % of Ti 0.4 Cr 0.15 Mn 0.15 V 0.3 (73 cycles, 290 °C) 56 , TiMn 2 (100 cycles, 300 °C) 55 , VTiCr (100 cycles, 300 °C) 55 , FeTi (500 cycles, 200 °C) 58 , metallic glassy of Zr 70 Ni 20 Pd 10 (100 cycles, 200 °C) 4 , 10 wt.% of big-cube Zr 2 Ni (2546 cycles, 250 °C) 21 , ZrNi 5 (600 cycles, 275 °C) 45 , 5Ni/5Nb 2 O 5 (180 cycles, 250 °C) 59 , and 5TiC/5Fe-12Cr (530 cycles, 275 °C) 25 . Integrated hydrogen storage system for fuel cell Applications. Our present work have two objectives; the first is focused on synthesizing and characterizing a nanocomposite MgH 2 /5 wt. % TiMn 2 system, where the second objective is focused on utilizing the as-synthesized nanocomposite powders manufacturing of a complete hydrogen storage system for fuel cell applications. To attain this purpose, a simple hydrogen storage tank was manufactured (Supplementary Materials, Fig. S1) composited of high-pressure hollow vessel made of pure titanium (Ti) metal where a hollow-graphite mould with an inner diameter of 10 mm was inserted into the vial. The powders were then charged into hollow-graphite mould and the system was sealed with copper metal gasket inside the glove box. A high pressure ball valve was perfectly installed into the tank cap's to allow hydrogen gas releasing and charging. Figure 9 presents photographs related to our experimental set up of the integrated hydrogen storage system including the Ti-vial placed into a gas-temperature control heater ( Fig. 9(a) and (b)) with a jacket temperature insulator ( Fig. 9(c)). The tank was then mounted on a temperature controlled hotplate (Fig. 9(c)). The hydrogen storage tank was connected 40 W/4.5 A proton exchange membrane (PEM) fuel cell ( Fig. 9(a)) through a pipeline allowing the released hydrogen gas form the tank to be passed to the PEM-fuel cell, as shown in Fig. 9(a). The converted electrical energy required to charge the battery of a cell phone device through 5 V voltage regulator ( Fig. 9(a)). The PEM-fuel cell system is controlled and operated with a software where the data output corresponding to hydrogen flow rate, voltage and current were obtained and stored.
Before starting the fuel-cell testing, the powders were obeyed to pressure-composition-temperature (PCT) analysis ensure the possibility of dehydrogenation process at a low temperature (180 °C) and to investigate the pressure required to achieve the decomposition process, MgH 2 /5TiMn 2 powders obtained after 50 h of milling. In the PCT experiment, the sample was firstly activated at high temperature (350 °C) that corresponding to about 35 bar of H 2 for 12 h. The powders were then charged with hydrogen at 180 °C, where the corresponding pressure reached to 0.4 bar, as shown in Fig. 10(a). Under this low temperature (180 °C) the fully charged powders tended to discharge their hydrogen storage capacity (5.43 wt.%) at 0.3 bar, as displayed in Fig. 10a. The PCT experiments was repeated 5 times under the same conditions to ensure the reprodcutability of the results. It is worth to be mentioned that the presence of minimal pressure-hysteresis gap between the pressure of absorption and desorption (~0.1 bar) indicates the powders' ability to achieve long term of cycle-life-time without failure, as was previously presented (Fig. 7b). Moreover, and in part to ensure our results, new set of kinetic measurements were conducted at 180 °C for both hydrogenation and dehydrogenation reactions taking place under 8-10 bar/0.2 bar (Fig. 10(b)). The possibility of achieving both reactions at 180 °C was approved, however, both reactions showed rather slow kinetics when compared with the results obtained at 250 °C ( Fig. 6(a) and (d)).
Prior to fuel cell experiments processing start, the nanocomposite MgH 2 /5 wt.% TiMn 2 powders were activated at 350 °C by charging/discharging hydrogen gas atmosphere at 30 bar/1 bar, respectively for 10 continuous cycles to obtain MgO-surface free nanocomposite powders. A continuous hydrogen gas flow pressurized at 10 bar was introduced at 180 °C to the Ti-tank containing the powders. The Ti-tank was then kept under this temperature (180 °C), where the readings shown by the hydrogen pressure gauge connected to the tank (Fig. 10(c)) was monitored and recorded every 1 min. These readings were used to construct a relationship between the time required for the powders placed to be charged with hydrogen, as shown by the open red symbols presented in Fig. 10(c). Completion of the absorption process was realized when the hydrogen was completely absorbed by the powders and the tank's pressure dropped to the atmospheric level (Fig. 10(c)). The hydrogen gas cylinder was then removed and disconnected from the tank, where the tank was kept under this 180 °C for almost 40 min. In order to construct a relation between the time required to release hydrogen gas stored in the powder in Ti tank, the readings of the pressure gauge was recorded every 1 min. The monotonical increase of hydrogen pressure inside the closed Ti-tank implies a gradual hydrogen release from the powders, as elucidated by the closed blue symbols shown in Fig. 10(c). The time required to release hydrogen from MgH 2 powders at 180 °C was about 25 min, as shown in Fig. 10(c). At this temperature, we started to open the valve of the tank shown connected to the PEM-FC system ( Fig. 9(a)) with nearly constant hydrogen flow of 175 ml/min ( Fig. 10(d)). Figure 10(d) displays the hydrogen gas flow transported into the PEM-fuel cell. The hydrogen released from the nanocomposite powders was used to feed the PEM-fuel cell through a pipeline, as shown in Fig. 9(a). During the battery charging time (28 min) the hydrogen gas released from the hydrogen storage tank at 180 °C passed delivered into the PEM-fuel cell at constant rate of 175 ml/min, as shown in Fig. 10(d). Providing the fuel cell with such a constant rate of hydrogen gas flow leads to generate constant values of voltage (7 V) and current (1 A), as displayed in Fig. 10(e) and (f), respectively.
In summary, we have synthesized nanocomposite MgH2/5 wt. % TiMn 2 system with advanced storage capacity (~5.3 wt. % H 2 ), hydrogenation/dehydrogenation kinetics and long hydrogen charging/discharging cyclability shown at rather low temperature (250 °C). The new nanocomposite powders were charged into a self-made Ti-tank and employed as a source of hydrogen required to charge a battery of a cell-phone device through a commercial proton exchange membrane (PEM) fuel cell.

Methods
Preparation of MgH 2 powders. Elemental Mg metal powders (~80 μm, 99.8% provided by Alfa Aesar -USA), and hydrogen gas (99.999%) were used as starting materials. An amount of 5 g Mg was balanced inside a He gas atmosphere (99.99%) -glove box (UNILAB Pro Glove Box Workstation, mBRAUN, Germany) and sealed together with fifty FeCr balls into a hardened steel vial (150 ml in volume), using a gas-temperature-monitoring system (GST; supplied by evico magnetic, Germany). The ball-to-powder weight ratio was maintained at 40:1. The vial was then evacuated to the level of 10 −3 bar before introducing H 2 gas to fill the vial with a pressure of 50 bar. The reactive ball milling (RBM) process was carried out at room temperature, using a high-energy ball mill (Planetary Ball Mill PM 400 provided by Retsch, Germany). The RBM process was interrupted after selected milling time (3, 6, 12.5, 25, 50, 100, and 200 h) where the vial was opened inside the glove box to take a small amount (~300 mg) of the milled powders for different analysis. Then, the RBM process was resumed, using the same operational conditions shown above. The as-synthesized MgH 2 powders were then mixed in the glove with the 5 wt. % of TiMn 2 shots, using an agate mortar and pestle. Five gram of the mixed powders for each composite system were charged together with fifty Cr-steel balls into the hardened steel vial and sealed under He gas atmosphere. The vial was then filled with 50 bar of hydrogen gas atmosphere and mounted on the high-energy ball mill. The milling process was interrupted after selected time (3, 6, 12.5, 25, 37.5 and 50 h) and the powders obtained after an individual milling time were completely discharged into 8 Pyrex vails for different analysis. Then, new MgH 2 /5 wt% TiMn 2 mixed powders were charged again and ball milled under the same milling conditions. The contamination contents of Fe and Cr of the powders obtained after 50 h of ball milling were 1.19 and 0.32 wt. %, respectively.
Powder characterizations. XRD AND HRTEM. The crystal structure of all samples was investigated by XRD with CuKα radiation, using 9 kW Intelligent X-ray diffraction system, provided by SmartLab-Rigaku, Japan. The local structure of the synthesized material powders was studied by 200 kV-field emission high resolution transmission electron microscopy/scanning transmission electron microscopy (HRTEM/STEM, supplied by JEOL-2100F, Japan), which is equipped with Energy-dispersive X-ray spectroscopy (EDS) supplied by Oxford Instruments, UK. In addition to the elemental analysis achieved by EDS approach, we employed ICP technique to get the elemental analysis by a chemical analytical approach.
Thermal stability. Shimadzu Thermal Analysis System/TA-60WS-Japan, using differential scanning calorimeter (DSC) was employed to investigate the decomposition temperatures of MgH 2 powders with a heating rate of 20 °C/min. The activation energy for of the powders obtained after different RBM time were investigated, using Arrhenius approach with different heating rates (5, 10, 20, 30, 40 °C/min).
The hydrogenation/dehydrogenation behaviors. The hydrogen absorption/desorption kinetics were investigated via Sievert's method [60][61][62] , using PCTPro-2000, provided by Setaram Instrumentation, France, under hydrogen gas pressure in the range between 200 mbar (for dehydrogenation) to 10 bar (for hydrogenation). The samples were examined at different temperatures of 50, 100, 250, and 275 °C. In the PCT measurements, the dosed pressure in absorption/desorption was gradually increased/decreased by 1000 mbar until equilibrium pressure reached to 13000 and 50 mbar, respectively. The PCT absorption/desorption kinetics were fitted in real-time by the HTPSwin software, to determine the sufficient equilibration time (the next point would start when the uptake had relaxed just 99% to asymptote). A minimum time of 30 minutes per equilibrium point and a maximum timeout of 300 minutes were set for each kinetic step in both the absorption and desorption isotherms.