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Article

Liquid Channels Built-In Solid Magnesium Hydrides for Boosting Hydrogen Sorption

by
Zhi-Kang Qin
1,
Li-Qing He
2,
Xiao-Li Ding
1,*,
Ting-Zhi Si
1,
Ping Cui
1,*,
Hai-Wen Li
2 and
Yong-Tao Li
1,3
1
School of Materials Science and Engineering & Low-Carbon New Materials Research Center, Anhui University of Technology, Maanshan 243002, China
2
Hefei General Machinery Research Institute, Hefei 230031, China
3
Key Laboratory of Green Fabrication and Surface Technology of Advanced Metal Materials of Ministry of Education, Anhui University of Technology, Maanshan 243002, China
*
Authors to whom correspondence should be addressed.
Inorganics 2023, 11(5), 216; https://doi.org/10.3390/inorganics11050216
Submission received: 20 March 2023 / Revised: 25 April 2023 / Accepted: 2 May 2023 / Published: 17 May 2023
(This article belongs to the Special Issue State-of-the-Art and Progress in Metal-Hydrogen Systems)
Browse Figures
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Abstract

:
Realizing rapid and stable hydrogen sorption at low temperature is critical for magnesium-based hydrogen storage materials. Herein, liquid channels are built in magnesium hydride by introducing lithium borohydride ion conductors as an efficient route for improving its hydrogen sorption. For instance, the 5 wt% LiBH4-doped MgH2 can release about 7.1 wt.% H2 within 40 min at 300 °C but pure MgH2 only desorbs less than 0.7 wt.% H2, and more importantly it delivers faster desorption kinetics with more than 10 times enhancement to pure MgH2. The hydrogen absorption capacity of LiBH4-doped MgH2 can still be well kept at approximately 7.2 wt.% without obvious capacity degradation even after six absorption and desorption cycles. This approach is not only through building ion transfer channels as a hydrogen carrier for kinetic enhancement but also by inhibiting the agglomeration of MgH2 particles to obtain stable cyclic performance, which brings further insights to promoting the hydrogen ab-/desorption of other metal hydrides.

Graphical Abstract

1. Introduction

Green and renewable energy development is the key to reducing carbon dioxide emissions and fossil fuel overuse [1,2,3,4,5]. Of these, hydrogen energy has garnered the most interest because of its abundant sources, high combustion heat value, and pollution-free combustion products [6]. However, achieving safe and efficient hydrogen storage is a key challenge. Solid-state hydrogen storage has relatively high storage volume density and transport safety, which have become the focus of hydrogen storage research in recent years [7,8,9,10,11]. Metal hydrides are widely used as solid hydrogen storage materials because they can store large quantities of hydrogen under milder conditions in a reversible manner. MgH2 is a candidate with sufficient reserves, a broad application, and high efficiency and safety. The Department of Energy considers its high reversible hydrogen storage capacity (7.6 wt.%) and volume hydrogen storage density (106 kg·m−3) to be among the most promising of the solid-state materials for meeting the technical requirements for onboard hydrogen storage [12,13,14,15]. However, high thermodynamic stability, high oxidation reactivity, and slow hydrogen sorption kinetics have become the primary obstacles to the practical application of hydrogen storage in vehicles. In the last two decades, numerous techniques for modifying the kinetic and thermodynamic properties of MgH2 have been developed to circumvent these issues: (i) alloying Mg with single transition metals and other metallic elements such as Ni [16,17], V [18], and Ti [19]; and (ii) nanoscale adjustment by confinement into single-walled carbon nanotubes [20] or graphene nanosheets (GNS) [21]. Unfortunately, these solutions typically have several drawbacks, including (i) low hydrogen storage capacity due to the addition of metals without hydrogen affinity, and (ii) irrepressible nanostructure agglomeration and instability [22,23,24,25]. Consequently, further investigation of a novel strategy to improve the hydrogen sorption performance of Mg-based materials at lower temperatures with faster kinetics is necessary.
In recent years, complex hydrides have been introduced into magnesium-based systems, proving to be a promising strategy for enhancing the kinetics of hydrogen storage [26,27,28,29,30]. For example, Liu et al. [31] reported a favorable desorption capacity of 4.5 wt.% at a relatively low temperature of 250 °C for the MgH2+ Li3AlH6 mixture. Li et al. [32] further discovered that the LiNH2−MgH2 system began desorbing hydrogen at 150 °C and exhibited improved reversibility, but this combined system released quantities of undesirable ammonia (NH3) gas. In light of these findings, introducing borohydrides into the MgH2 system will also improve the kinetic properties, albeit at the expense of inflexible thermodynamic properties and the formation of byproducts. Kato et al. [33] discovered that the altered hydrogen desorption in the NaBH4 and MgH2 systems could be attributed to the migration of metallic Mg into the surface of NaBH4. The intrinsic mechanisms of these combined systems are still unknown, but these results suggest that the enhanced properties are primarily a result of the rapid migration of ionic hydrogen in MgH2 [34].
In light of these findings, we present a novel method for improving the hydrogen sorption of MgH2 by introducing lithium borohydrides, namely the construction of an ion transfer channel in MgH2. Complex borohydrides such as LiBH4 and Li2B12H12 are known to be fast ion conductors composed of Li+, [BH4]−, and [B12H12]2− [35], which can serve as intermediates for high ionic conduction and activity in the diffusion of H- from MgH2. This new method involves the construction of ion transfer channels as hydrogen carriers for kinetic enhancement and inhibiting the aggregation of MgH2 particles from achieving stable cyclic performance. The desorption process of the pseudo-eutectic LiBH4–MgH2 system involves liquid borohydride phases in particular. In addition, the science underlying the remarkable kinetic enhancements of MgH2 brought about by the introduction of liquid borohydride channels was elucidated.

2. Results and Discussion

2.1. Hydrogen Storage Properties of LiBH4-Doped MgH2

We first introduce LiBH4 as an ionic conductor to enhance the milling kinetic performance of MgH2 and then compare it with pure and Li2B12H12-doped MgH2 systems. Figure 1 compares the hydrogen absorption and desorption kinetics of pure, Li2B12H12-doped, and LiBH4-doped MgH2 at 300 °C, which are detected by their sixth hydrogen cycle. As depicted in Figure 1a, the desorption hydrogen kinetic properties of MgH2 are sluggish, and less than 0.7 wt.% hydrogen is desorbed within 40 min. In contrast, the addition of complex borohydrides improves the desorption kinetics of MgH2. The sixth dehydrogenation can be completed rapidly and releases ~7.1 wt.% hydrogen within 40 min, representing enhancement by ten-fold compared to pure MgH2. Similar desorption enhancements were observed in the L2B12H12-doped MgH2 system. Figure 1b further indicates that LiBH4-doped MgH2 has superior hydrogen absorption kinetic properties. As compared to the capacity of ~1.6 wt.% for pure MgH2 within 10 min, the absorption capacities of L2B12H12-doped MgH2 and LiBH4-doped MgH2 are significantly increased by three and four times, respectively.
In order to explore and compare the initial desorption behaviors and cyclic performances of the pure MgH2 and LiBH4-doped MgH2, they were tested for six cycles at 300 °C and the curve of hydrogen ab-/desorption at a constant temperature was shown in Figure 2. It can be seen that the desorbed capacity of pure MgH2 is low, with only ~1.5 wt.% H2, and the absorbed capacity was ~1.8 wt.% H2 after the sixth cycle. More excitingly, that of the LiBH4-doped MgH2 can directly increase by releasing ~7.1 wt.%, and the hydrogen absorption capacity can still be kept at ∼7.2 wt.% even after six cycles. These results indicate that the kinetics and cyclic performance of MgH2 can be significantly enhanced by introducing ionic borohydrides.

2.2. Structural Features of LiBH4-Doped MgH2

The high-resolution transmission electron microscopy (HRTEM) technique is used to reveal the microstructure characteristics of the as-prepared LiBH4-doped MgH2 to clarify the reason for these improvements. As shown in Figure 3a, the light gray massive structures are embedded homogeneously in the dark gray zonal distribution, forming a transmission channel corresponding to the selected regions in Figure 3b,d. From the selected regions, we obtain lattice stripes by transposing the selection using the Fourier transform, in which all the d-spacings of approximately 0.3801, 0.322 and 0.252 nm can be easily indexed to the (011), (111), and (112) planes of the LiBH4 phase, respectively. In addition, an amorphous layer is visible at the edge of the associated structures for LiBH4-doped MgH2 (yellow dotted line in Figure 3a); LiBH4 and MgH2 have a relatively well-balanced distribution within this structure, as depicted in Figure 3f–i. This novel structure strongly indicated that the LiBH4-doped MgH2 system had successfully constructed the zonal channel for hydrogen transfer. We also found the presence of the element O from the EDS spectrum in Figure 3g, indicating that the passivation layer on the surface of MgH2 is unavoidable even for commercial MgH2. Moreover, the existence of associated structures of LiBH4-doped MgH2 and the formation of an amorphous layer collectively inhibit the Mg grain particle agglomeration during kinetic cycling, thereby facilitating the diffusion of hydrogen for improved kinetic and cyclic properties.
In order to determine the current state and distribution of the LiBH4-doped MgH2 system upon desorption or absorption, we analyzed the morphological and structural characteristics of the as-prepared LiBH4-doped MgH2 materials before and after six cycles. Figure 4a shows that, except for the MgH2/Mg phases, no LiBH4 peaks were detected in the XRD patterns. Intriguingly, Figure 4b of the FTIR results reveals that the characteristic bands of B-H vibration and stretching bonding of LiBH4 at ~2359, 2293, 2225 and 1128 cm−1 can always be detected before and after cycling, although their intensities diminish slightly. Apparently, LiBH4 indeed exists in amorphous and/or nanocrystal states during both ball milling and cycling, rather than decomposing or reacting to form a new phase, as further demonstrated by the XPS analysis in Figure 4c, where the electronic binding energy of B1s at approximately 188 eV demonstrates the stable existence of LiBH4 in the cyclic LiBH4-doped MgH2. Moreover, Figure 4d–i shows the morphological evolution and elemental distribution of LiBH4-doped MgH2 during cycling. Before and after cycling, a particle morphology with an average size of 1~2 μm was observed, along with sintering-induced connection phenomena upon heating for sorption. Furthermore, Mg and B elemental distributions were well dispersed. These results indicate that the good dispersion of LiBH4 significantly inhibits the growth of Mg grains, which explains why superior kinetic and cyclic properties were obtained in Figure 1.

2.3. Electrochemical Analysis of LiBH4-Doped MgH2

Electrochemical impedance spectroscopy and in situ morphological analysis were utilized to elucidate the intrinsic role of LiBH4 as an ionic conductor in enhancing hydrogen sorption on MgH2. The semicircle Nyquist plots of LiBH4-doped MgH2 are significantly smaller than those of pure MgH2 (see Figure 5a), indicating lower electrolyte resistance and exceptionally fast electron conductivity in a simulation of an all-solid-state battery. Figure 5b depicts the Nyquist plots of LiBH4-doped MgH2 at various temperatures, where the characteristic impedance semicircle gradually decreases with increasing temperatures, corresponding to the resistance value decreasing from approximately 4 × 106 Ω at 55 °C to approximately 1.05 × 105 Ω at 125 °C. Figure 5c compares the Arrhenius curves of pure MgH2, LiBH4, and LiBH4-doped MgH2. The enhanced ionic conductivity apparently indicates that MgH2 shifts from an insulator to a conductor with ionic conductivity of approximately 3.2 × 10−7 S cm−1 at 125 °C by introducing LiBH4. In addition, the in situ optical images of the LiBH4-doped MgH2 sample during heating shown in Figure 5d reveal that above 275 °C (i.e., the melting point of LiBH4 [36]), liquid droplets form on the surface of the matrix. Intriguingly, hydrogen desorption from MgH2 was captured by continuous and rapid bubbling from the liquid LiBH4 phase (Figures S6 and S7, ESI).
Figure 5e proposes a model of hydrogen desorption enhanced by liquid ion channels incorporated into solid magnesium hydrides based on previous results. The liquid LiBH4 droplets uniformly embedded in the MgH2 matrix serve as channels for the rapid carrier of hydrogen, which can be understood from three perspectives: (i) the H2 bubbles can move more rapidly in the liquid phase compared to the solid phase due to lower migration energy; (ii) as we know, the LiBH4 would transfer from a low-temperature phase into high-temperature phase at ~120 °C, corresponding to the significant increase in ion conduction as shown in Figure 5a–c, and its enhanced ion conduction of composites would also be kept even at higher temperature. This further forces us to deduce that the enhanced desorption of MgH2 is related to the faster ion conduction of LiBH4; and (iii) the introduction of fast Li+ ions from LiBH4 can easily carry H− from ionic MgH2 via the ionic interactions and can also alter the ionization degree and ionic density of the Mg-H bond to improve the migration of H atoms, as further confirmed by in situ spectroscopy. Along with enhancing the Li-ion migration, the H2 sorption is also significantly promoted above phase-transition temperature. Thus, we deduce that introducing LiBH4 into MgH2 matrix would enhance the migration of H− by interacting with fast Li+ transfer based on positive and negative ionic attraction to each other. In this regard, more experiments with advanced technology are needed to directly detect the important correlation that will provide insight into property improvement.

3. Experimental Procedure

3.1. Material Preparation

MgH2 (purity > 95% from anabai Medicine Co., Ltd., Wuhan, China) and LiBH4 (purity > 95% from China Aladdin) were combined in a weight ratio of 19:1 and loaded into a 300 mL stainless-steel ball milling tank (SUS304). The milling tank filled with 1 g of mixed material, and different masses and diameters of stainless steel (DECO-304-B) balls (5 mm—18.6 g; 6 mm—9.2 g; 8 mm—8 g; 10 mm—4.2 g), with a ball-to-sample mass ratio of 40:1 was mechanically milled for 10 h at 400 rpm with a planetary ball mill (QM-3SP2). All prepared samples were milled for twenty periods of 30 min, with a 2 min interval between each period. The same experimental procedures and parameters were used on the control group of pure MgH2 and that with 5 wt% Li2B12H12 doping (purity > 95% from Aladdin). All mechano-chemical treatments were performed in an Argon-filled glovebox (ρ(O2) < 0.1 ppm, ρ (H2O) < 0.1 ppm).

3.2. Material Preparation

The XRD analysis was performed on a MiniFlex 600 XRD unit (Rigaku, Japan), Cu Kα radiation (λ = 0.154056 nm) utilized at 40 kV and 15 mA. The 2θ angle ranged from 10° to 90° with increments of 0.02°. The powder samples were placed in custom-made molds and sealed with polyimide thin-film tape, ensuring that they were under an argon atmosphere during the measurement process. The morphologies of the samples were observed by scanning electron microscopy (SEM, Zeiss Sigma 300) and transmission electron microscopy coupled with an EDS (TEM, FEI Talos F200X). The FTIR analyses were conducted on a TENSOR27 with a wavelength range from 400 to 3000 cm−1. A Thermo Fisher Scientific Spectrometer K-Alpha was used to conduct XPS analyses. The powder sample was contained in an argon-filled glove box before being mounted on a sample holder and transferred to the XPS facility using a special container to prevent air exposure.
Electrochemical impedance spectroscopy (EIS) was used to measure the ionic conductivities of pure MgH2 and 5 wt.% LiBH4-doped MgH2 by using a Solartron impedance analyzer. The material was weighed as 120 mg by electronic balance in a high-purity argon glove box, and then poured into a metal sheet mold with a press at 7 MPa for 5 min. The pressing tablet was the positive material, while the negative was a lithium tablet that was wrapped in a plastic bag and then rolled with a metal rod into a uniform sheet. Before electrochemical testing, all experimental and control groups were mixed with ~12 mg of acetylene black to increase their electrical conductivity.
The de-/absorption kinetics of samples were measured by using an automated Sieverts-type apparatus that allowed for accurate determination of the evolved hydrogen amount. Approximately 2 g of sample was loaded into an evacuated stainless-steel autoclave that connected with automatic PCT measurements. After rapid heating of the sample to the desired temperatures, the autoclave was immersed into the heating furnace. After activation, the hydrogen de-/absorption curves at 300 °C were successively performed by a back pressure of 0.01 and 4 MPa, respectively.

4. Conclusions

In conclusion, we successfully incorporated 5 wt.% LiBH4 into the MgH2 hydrogen storage system to significantly improve the hydrogen kinetic and stable cyclic properties. Our experimental findings support the design of a hydrogen-optimized ion transfer channel, which is facilitated by forming associated structures. For superior cyclic performance, an amorphous layer and the uniform dispersion of LiBH4 are the two most important factors inhibiting the growth of Mg grains. During the heating process, the LiBH4 droplets uniformly embedded in the MgH2 matrix serve as ion migration channels for rapid transport of hydrogen. These findings suggest a new strategy for enhancing the hydrogen sorption of ionic hydrides and other hydride systems, as well as for accelerating the search for candidate materials that are suitable for hydrogen storage.

Supplementary Materials

The following supporting information can be downloaded at: https://www.mdpi.com/article/10.3390/inorganics11050216/s1, Figure S1: Temperature-programmed kinetics of desorption; Figure S2: Isothermal absorption of pure MgH2; Figure S3: Isothermal absorption of LiBH4-doped MgH2; Figure S4: SEM images of LiBH4-doped MgH2 before kinetic cycles; Figure S5: SEM images of LiBH4-doped MgH2 after six kinetic cycles; Figure S6: High temperature laser confocal image for liquid phase transition during hydrogen desorption at 276 °C; Figure S7: High temperature laser confocal image for liquid phase transition during hydrogen desorption at 281 °C.

Author Contributions

Conceptualization, Z.-K.Q. and X.-L.D.; methodology, L.-Q.H.; validation, P.C. and T.-Z.S.; formal analysis, Z.-K.Q.; investigation, H.-W.L.; writing—original draft preparation, Z.-K.Q.; writing—review and editing, X.-L.D.; visualization, P.C.; supervision, H.-W.L.; project administration, Y.-T.L. All authors have read and agreed to the published version of the manuscript.

Funding

This work was financially supported by the Key Program for International S&T Cooperation Projects of China (No. 2017YFE0124300), National Natural Science Foundation of China (Nos. 51971002, 52171205, 52101249 and 52171197), Scientific Research Foundation of Anhui Provincial Education Department (Nos. KJ2020ZD26, KJ2021A0360), Anhui Provincial Natural Science Foundation for Excellent Youth Scholars (No. 2108085Y16), and the Provincial University Outstanding Youth Research Project (No. 2022AH020033).

Data Availability Statement

Data is contained within the article.

Conflicts of Interest

The authors declare no conflict of interest.

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Inorganics 11 00216 g001 550
Figure 1. (a) Hydrogen desorption and (b) absorption kinetic curves of pure, Li2B12H12−doped and LiBH4−doped MgH2 at 300 °C; all sample data are detected by the sixth cycle.
Figure 1. (a) Hydrogen desorption and (b) absorption kinetic curves of pure, Li2B12H12−doped and LiBH4−doped MgH2 at 300 °C; all sample data are detected by the sixth cycle.
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Figure 2. Kinetics and cycling performance of hydrogen desorption and absorption of (a,b) MgH2 and (c,d) LiBH4—doped MgH2 upon six cycles at 300 °C.
Figure 2. Kinetics and cycling performance of hydrogen desorption and absorption of (a,b) MgH2 and (c,d) LiBH4—doped MgH2 upon six cycles at 300 °C.
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Figure 3. Microstructural features of as-milled LiBH4-doped MgH2: (a) TEM images; (b,d) fast Fourier transform (FFT) and (c,e) their lattice images; (f) STEM-HAADF image; (g) EDS spectrum and its corresponding (h) Mg and (i) B elemental mapping.
Figure 3. Microstructural features of as-milled LiBH4-doped MgH2: (a) TEM images; (b,d) fast Fourier transform (FFT) and (c,e) their lattice images; (f) STEM-HAADF image; (g) EDS spectrum and its corresponding (h) Mg and (i) B elemental mapping.
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Figure 4. (a) XRD patterns, (b) FTIR and (c) XPS spectra of the pure MgH2, as-milled and cyclic LiBH4-doped MgH2 samples, as well as the FESEM images of LiBH4-doped MgH2 sample (d) before and (g) after cycling and their corresponding (e,h) Mg and (f,i) B elemental mapping.
Figure 4. (a) XRD patterns, (b) FTIR and (c) XPS spectra of the pure MgH2, as-milled and cyclic LiBH4-doped MgH2 samples, as well as the FESEM images of LiBH4-doped MgH2 sample (d) before and (g) after cycling and their corresponding (e,h) Mg and (f,i) B elemental mapping.
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Figure 5. (a) Nyquist plots of pure and LiBH4−doped MgH2 as well as their (b) Nyquist plots at different temperatures and (c) corresponding Arrhenius plots; (d) In situ optical images of LiBH4−doped MgH2 sample detected by the high temperature laser confocal microscope and (e) its proposed model for formation of fast ions migration channels in MgH2 induced by LiBH4 melting.
Figure 5. (a) Nyquist plots of pure and LiBH4−doped MgH2 as well as their (b) Nyquist plots at different temperatures and (c) corresponding Arrhenius plots; (d) In situ optical images of LiBH4−doped MgH2 sample detected by the high temperature laser confocal microscope and (e) its proposed model for formation of fast ions migration channels in MgH2 induced by LiBH4 melting.
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MDPI and ACS Style

Qin, Z.-K.; He, L.-Q.; Ding, X.-L.; Si, T.-Z.; Cui, P.; Li, H.-W.; Li, Y.-T. Liquid Channels Built-In Solid Magnesium Hydrides for Boosting Hydrogen Sorption. Inorganics 2023, 11, 216. https://doi.org/10.3390/inorganics11050216

AMA Style

Qin Z-K, He L-Q, Ding X-L, Si T-Z, Cui P, Li H-W, Li Y-T. Liquid Channels Built-In Solid Magnesium Hydrides for Boosting Hydrogen Sorption. Inorganics. 2023; 11(5):216. https://doi.org/10.3390/inorganics11050216

Chicago/Turabian Style

Qin, Zhi-Kang, Li-Qing He, Xiao-Li Ding, Ting-Zhi Si, Ping Cui, Hai-Wen Li, and Yong-Tao Li. 2023. "Liquid Channels Built-In Solid Magnesium Hydrides for Boosting Hydrogen Sorption" Inorganics 11, no. 5: 216. https://doi.org/10.3390/inorganics11050216

APA Style

Qin, Z. -K., He, L. -Q., Ding, X. -L., Si, T. -Z., Cui, P., Li, H. -W., & Li, Y. -T. (2023). Liquid Channels Built-In Solid Magnesium Hydrides for Boosting Hydrogen Sorption. Inorganics, 11(5), 216. https://doi.org/10.3390/inorganics11050216

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