Investigating the Factors Affecting the Ionic Conduction in Nanoconfined NaBH4

Nanoconfinement is an effective strategy to tune the properties of the metal hydrides. It has been extensively employed to modify the ionic conductivity of LiBH₄ as an electrolyte for Li-ion batteries. However, the approach does not seem to be applicable to other borohydrides such as NaBH₄, which is found to reach a limited improvement in ionic conductivity of 10⁻⁷ S cm⁻¹ at 115 °C upon nanoconfinement in Mobil Composition of Matter No. 41 (MCM-41) instead of 10⁻⁸ S cm⁻¹. In comparison, introducing large cage anions in the form of Na₂B₁₂H₁₂ naturally formed upon the nanoconfinement of NaBH₄ was found to be more effective in leading to higher ionic conductivities of 10⁻⁴ S cm⁻¹ at 110 °C.

Keywords: sodium borohydrides; nanoconfined; ionic conductivity; solid-state electrolyte

1. Introduction

The development of solid-state electrolytes is critical for the advancement of all solid-state high-energy density batteries. In this respect, several families of inorganic solid-state electrolytes have been intensively investigated because of their non-flammability and capability in leading to a high Li⁺/Na⁺ transference number and thus safety with extended battery life-time [1]. However, current solid-state electrolytes still suffer from low ionic conductivity [2]. Current materials reported to lead to high ionic conductivity are oxide solid-state electrolytes like La_0.52Li_0.35TiO_2.96 (ionic conductivity of 10⁻³ at room temperature) [3]. However, in this case, the poor contact between the electrode and the ceramic electrolyte causes a high interfacial resistance [4]. Sulfide-based solid-state electrolytes have also been found to lead to high ionic conductivity (e.g., Li_9.54Si_1.74P_1.44S_11.7C_l0.3 has an ionic conductivity of 2.5 × 10⁻² S cm⁻¹ at room temperature) [5], and this is comparable to the ionic conductivity of organic electrolytes [6]; however, sulfide electrolytes have a narrow electrochemical stability window and tend to decompose upon operation and release toxic H₂S gas [7,8].

Recently, metal borohydrides have drawn considerable interest because of their chemical and electrochemical stability, as well as chemically compatibility with Li/Na metal [9], but modifications are needed to enable high ionic conductivity [10] in complex borohydrides at ambient and not high temperatures. To date, the most common strategies to tailor the ionic conductivity of complex borohydrides are based on anionic substitution [11,12,13] or nanoconfinement approaches [10,14]. For example, partial anionic substitution of BH₄⁻ by NH₂ in NaBH₄ leads to an increase in ionic conductivity from 1 × 10⁻⁹ to 2 × 10⁻⁶ S cm⁻¹ at 26 °C [11]. Nanoconfining LiBH₄ (LiBH₄@MCM-41) within the pores of ordered silica scaffolds such as MCM-41 was reported to lead to a high ionic conductivity of 0.2 mS cm⁻¹ at 55 °C instead of 1 × 10⁻⁸ S cm⁻¹ [10] owing to the interface interaction between LiBH₄ and MCM-41 [15,16]. Further testing of such a material showed that nanoconfined LiBH₄ in an Li/S cell could deliver a high capacity of 1220 mAh g⁻¹ at a working voltage of 2 V and at 0.03 C rates for 40 cycles, which is comparable to the sulfide electrolytes [17,18]. To date, such improvements have only been reported for LiBH₄, with no report on the possibility of such an observation to other borohydrides and in particular NaBH₄, which is of interest as a solid electrolyte for Na batteries. Except for borohydrides materials, nanoconfinement has been proven as an effective method to tailor the ionic conductivity of electrolyte materials including solid composite polymer electrolytes and ionic liquid electrolytes confined in Metal Organic Frameworks [19,20,21]. Herein, we report on the successful confinement of NaBH₄ into the MCM-41, and the resulting limited improvement in ionic conductivity. As a reference, LiBH₄ was also confined into the scaffold material MCM-41 and the existence of an oxide phase in LiBH₄@MCM-41 further proved our hypothesis, i.e., the strong oxidation of NaBH₄ upon nanoconfinement in MCM-41 and thus the formation of an extensive oxide phase at the borohydride/MCM-41 interface limits Na and Li conduction. Better improvement in ionic conductivity could be achieved by partially decomposing NaBH₄ into a mixture of Na₂B₁₂H₁₂ and NaBH₄.

2. Results and Discussions

2.1. Infiltration of NaBH₄ in MCM-41

To confirm the degree of infiltration of NaBH₄ into MCM-41, Brunauer–Emmett–Teller (BET) measurements were carried out. As expected, melt infiltration of NaBH₄ into MCM-41 resulted in a pore volume reduction from the original MCM-41 value of 1.02 to 0.02 cm³ g⁻¹, and this corresponded to a 78% pore filling. As reflected by BET, the melt infiltration of NaBH₄ into MCM-41 also led to a significant decrease of the surface area of NaBH₄@MCM-41 from 1110.91 m² g⁻¹ for pristine MCM-41 to 3.5 m² g⁻¹ (Table S1 and Figure S1).

Further analysis by transmission electron microscopy (TEM) showed that the ordered porous structure of MCM-41 was filled upon nanoconfinement of NaBH₄ with the disappearance of a clear porous structure (Figure S2a). Additionally, elemental mapping showed a signal of Na and Si overlapping, and this was taken as additional evidence that NaBH₄ was melt infiltrated within the porosity of MCM-41 (Figure S3).

X-ray diffraction (XRD) patterns of the pristine and nanoconfined NaBH₄ are shown in Figure 1a. For NaBH₄@MCM-41, all the peaks are assigned to NaBH₄, which indicates that no detectable additional phase was formed during the melt infiltration process at 560 °C. However, all the diffraction peaks showed a significant boarding and shift to lower diffraction angles as compared with pristine NaBH₄ (Figure 1b), which suggested some confinement of the borohydride [22,23,24]. Indeed, a shift has been reported to occur for confined borohdyrides owing to the lattice strain imposed by the MCM-41 scaffold [22,25], as further revealed by the small angle X-ray diffraction analysis of MCM-41 and NaBH₄@MCM-41. As shown in Figure 1c, the main (100) diffraction peak of MCM-41 broadens. This peak also shifted to higher diffraction angles, further indicating that NaBH₄ was located within the internal pore of MCM-41 [26]. From these results, it can thus be concluded that NaBH₄ is infiltrated within the porosity of MCM-41. Another indirect evidence of the nanoconfinement is the shift in the dehydrogenation peak of NaBH₄ upon infiltration in MCM-41, from 550 to 520 °C (Figure S4).

To further determine any amorphous phases that may have formed during the nanoconfinement of NaBH₄ in MCM-41 at 560 °C, Fourier Transform Infrared Spectrometery (FTIR)and nuclear magnetic resonance (NMR) analyses were carried out. By FTIR (Figure 2), the typical BH stretching and bending vibrations corresponding to the BH₄⁻ anion in NaBH₄ were observed in NaBH₄@MCM-41 in the range from 2400 to 2200 cm⁻¹ and at 1091 cm⁻¹, in agreement with previous reports [27]. The broad peaks between 3800 and 3200 cm⁻¹ were assigned to OH stretching modes corresponding to a partial oxidation of NaBH₄ in contact with the walls of the MCM-41. As a scaffold material, MCM-41 possesses silanol (Si–OH) and hydrogen-bonded terminal hydroxyl (Si–OH–O–Si) groups located within its internal structure [28]. Therefore, during the melt infiltration of NaBH₄, it is not surprising that silanol and/or hydroxyl groups readily react with NaBH₄ to lead to the formation of boron oxide phases in NaBH₄@MCM-41, as evidenced by the peaks at 1626, 883, and 794 cm⁻¹ owing to the B–O vibrational modes (Figure 2) [29,30,31]. To eliminate the concern regarding the formation of other oxide compounds (e.g., NaO₂), we carefully checked the FTIR spectrum of NaBH₄@MCM-41 in the range of 800–400 cm⁻¹ (Figure S8), and the only peak located at 473 cm⁻¹ was assigned to NaBO₄ instead of NaO₂ [32]. A similar oxidation has previously been observed upon the infiltration of LiBH₄ in SBA-15, and in this case, this led to the formation of LiBO₂ [33]. Besides these oxidized phases, the peak observed at 2496 cm⁻¹ was attributed to the formation of Na₂B₁₂H₁₂, which is commonly reported to occur upon a partial decomposition of NaBH₄ [34,35]. These results are also in agreement with the ¹¹B NMR spectrum (Figure 3) showing the typical resonance of the BH₄⁻ anion centered at −41.95 ppm [36]; a single sharp peak at −2 ppm, assigned to boron in tetrahedral BO₄ environments [37]; and at −15.58 ppm, a peak corresponding to the dodecahedral [B₁₂H₁₂]²⁻ anion in Na₂B₁₂H₁₂ [38].

2.2. Ionic Conductivity of Nanoconfined NaBH₄

The ionic conductivity of the pristine NaBH₄ and NaBH₄@MCM-41 was determined by electrochemical impedance spectroscopy (EIS) (Figure 3). The ionic conductivity of NaBH₄@MCM-41 is 10 times higher (i.e., 7.4 × 10⁻¹⁰ S cm⁻¹ at 20 °C) than that of pristine NaBH₄ in the temperature range of 20–70 °C. However, above 70 °C, the ionic conductivity of NaBH₄ and NaBH₄@MCM-41 was found to be of the same magnitude. To our surprise, encapsulating NaBH₄ into MCM-41 did not significantly enhance the ionic conductivity. This may be because of the extensive reaction of nanoconfined NaBH₄ with the walls of MCM-41 during melt infiltration, which resulted in the formation of an insulating oxide phase [39]. Indeed, boron oxide phases (e.g., NaBO₄) have been reported to be poor ionic conductors [40]. In previous reports, it has been suggested that the ionic conductivity of MCM-41 nanoconfined LiBH₄ was the result of the LiBH₄/MCM-41 interface promoting the reorientation of BH₄⁻ [10,25]. To verify this hypothesis, we reproduced the nanoconfinement of LiBH₄ by melt infiltration in MCM-41 [10]. Successful nanoconfinement of LiBH₄ was verified by BET and thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) analysis (Figures S5 and S6). In this case, the ionic conductivity of LiBH₄@MCM-41 was found to be 1.34 × 10⁻⁷ S cm⁻¹ at 110 °C (Figure S9), which is four times lower in magnitude as compared with the reported value of 10⁻³ S cm⁻¹ [10]. Through careful analysis by FTIR (Figure S7), strong B–O vibrations were observed in LiBH₄@MCM-41 as compared with pristine LiBH₄. These were assigned to lithium borates (Li₂B₄O₇) [41], and indicated a strong oxidation of the borohydride phase during nanoconfinement. The discrepancy between the ionic conductivity of the current LiBH₄@MCM-41 and reported data may thus be attributed to the extensive oxidation of the complex borohydride within the scaffold material. Here, we emphasize the inevitable reaction between the pore wall in MCM-41 and complex borohydrides and the resulting negative effect on ionic conduction.

One factor that may explain the increase in ionic conductivity observed for NaBH₄@MCM-41 below 70 °C is the presence of Na₂B₁₂H₁₂, which was reported to lead to improved ionic conductivity [42]. Through integrating the chemical shifts of B₁₂H₁₂ located at −15.6 ppm in NaBH₄@MCM-41(Figure 4), 18% of amorphous Na₂B₁₂H₁₂ appeared in NaBH₄@MCM-41 owing to decomposition of NaBH₄ (Figure 4). We believed that the appearance of the Na₂B₁₂H₁₂ in the nanoconfined material is due to the interaction between the NaBH₄ and scaffold materials, triggering a partial decomposition of NaBH₄ upon its oxidation and the release of B₂H₆, further reacting with NaBH₄ to lead to Na₂B₁₂H₁₂ [43,44].

To investigate if the presence of Na₂B₁₂H₁₂ within NaBH₄ can effectively lead to improved ionic conductivities, we synthesised NaBH₄ with higher amounts of Na₂B₁₂H₁₂. After the reaction between B₂H₆ and NaBH₄, 43% of the starting NaBH₄ was converted to Na₂B₁₂H₁₂ (Figure S10). Such a mixture of NaBH₄ and Na₂B₁₂H₁₂ is stable up to 300 °C (Figure S11) and exhibited a drastic enhancement in ionic conductivity, with an ionic conductivity of 10⁻⁶ S cm⁻¹ at 110 °C (Figure 5). This is two magnitudes higher than pristine NaBH₄ [45]. Thus, the enhancement in ionic conductivity observed in nanoconfined NaBH₄@MCM-41 can be assigned to the existence of the B₁₂H₁₂²⁻ anion. Anion rotation is believed to significantly enhance cation hopping [46,47], and B₁₂H₁₂⁻ exhibits an intrinsically high dynamic motion [48]. Accordingly, the enhanced ionic conductivity in NaBH₄@MCM-41 was attributed to the co-existence of B₁₂H₁₂²⁻ and BH₄⁻. We thus believe that fine tuning the ratio between Na₂B₁₂H₁₂ and NaBH₄ will further increase the ionic conductivity of NaBH₄ and other complex borohydrides. Further investigations along this path are underway.

3. Materials and Methods

3.1. Synthesis

All the operations were carried out under an inert atmosphere in an argon-filled LC-Technology glove box (<1 ppm O₂ and H₂O, Salisbury, MA, USA). Sodium borohydride (NaBH₄, 99%) was purchased from Sigma-Aldrich (Sydney, NSW, Australia) and further purified. ZnCl₂ (≥98%) was purchased from Ajax Finechemn (Sydney, NSW, Australia), and dried at 120 °C overnight on a Schlenk line under vacuum (0.01 MPa) before use. Lithium borohydride (LiBH₄, 95%) was purchased from Acros (Sydney, NSW, Australia). Prior to use, LiBH₄ was purified following the reported procedures [49]. MCM-41 was purchased from ACS materials and dried under vacuum at 200 °C for 2 h before use to remove any water traces.

3.1.1. Synthesis of the Nanoconfined Complex Borohydrides

For the nanoconfinement of NaBH₄ into MCM-41 (noted NaBH₄@MCM-41), a mixture of 1.5 g NaBH₄ grinded with 1.27 g MCM-41 was heated at 5 °C min⁻¹ and kept at 560 °C for 3.5 h under 8 MPa H₂ pressure. This temperature was chosen because NaBH₄ melts at 500 °C (Figure S4)

LiBH₄@MCM-41 was synthesized as per previous report [10]. LiBH₄ infiltration was carried out to fill 100% of the pore volume of MCM-41. This was achieved by mixing 0.85 g of LiBH₄ with 0.127 g of MCM-41 in a mortar and pestle. The mixture was then placed in a stainless-steel sample holder and heated to 295 °C at 5 °C min⁻¹ under an H₂ pressure of 10 MPa. Infiltration was done at this temperature for 30 min.

3.1.2. Synthesis of the Mixture of NaBH₄ and Na₂B₁₂H₁₂ Was via Solid–Gas Reaction

NaZn₂(BH₄)₅ was synthesised by ball milling pristine NaBH₄ and ZnCl₂ with a molar ratio of 2:1 with a Retsch MM301 mill operated at a frequency of 20 Hz. The mixture of NaBH₄ and ZnCl₂ was milled in a stainless-steel vial (25 mL) containing a single stainless-steel ball (15 mm diameter) for 10 min. Sodium borohydride (120 mg) was ball-milled in a similar manner under high purity argon in a 15 mL stainless steel vial filled with a single stainless-steel ball (1.5 g and 15 mm diameter).

The synthesis of NaBH₄/Na₂B₁₂H₁₂ mixed compounds was undertaken in an in-house built Sievert apparatus with a customized sample holder, which had two compartments separated by a stainless-steel mesh (20 μm porosity). Then, 100 mg of ball milled NaBH₄ was placed on the top of the mesh and 400 mg NaZn₂(BH₄)₅ was placed at the bottom of the sample holder. The reaction was carried out at 150 °C for 4 h under 10 MPa H₂. The materials synthesized by this solid–gas reaction were then ball milled in a 15 mL stainless steel vial filled with a single stainless steel ball (1.5 g and 15 mm diameter) at a frequency of 20 Hz for 10 min, and the previous synthesis route was repeated to maximise the Na₂B₁₂H₁₂ yield.

3.2. Characterization

The crystalline nature of the materials was determined by X-ray diffraction (XRD) using a Philips X’pert Multipurpose XRD system (Bruker, Preston, Victoria, Australia)operated at 40 mA and 45 kV with a monochromatic Cu Kα radiation (λ = 1.541 Å)-step size = 0.01°, time per step = 10 s·step⁻¹. The materials were protected against oxidation from the air by a Kapton foil.

Infrared analysis was carried out on a Bruker Vertex 70 V (Bruker, Preston, Victoria, Australia) equipped with a Harrick diffuse reflectance Praying Mantis accessory. The materials were loaded in an air-tight chamber in the glovebox and the chamber was fitted on the Praying Mantis. Spectra were acquired with a 1 cm⁻¹ resolution with an MCT-detector.

Solid-state ¹¹B magic angle spinning (MAS) nuclear magnetic resonance (NMR) experiments were carried out on a narrow-bore Bruker Biospin Avance III solids-700 MHz spectrometer (Bruker, Preston, Victoria, Australia) with a 16.4 Tesla superconducting magnet operating at a frequency of 224.7 MHz ¹¹B nucleus. Approximately 3–10 mg of material was packed into 4 mm zirconia rotors fitted with Kel-f^® caps or vespel caps, respectively. The 4 mm rotors were spun in a double resonance H-X probe head at 14 kHz at the magic angle. The ¹¹B spectra were acquired with hard 1 to 3 μs radio frequency pulses corresponding to a 30° tip angle. The recycle delays of up to 10 s were used to ensure full relaxation of the signals of all nuclei, and up to 512 transients were co-added to ensure sufficient signal to noise. The spectra were obtained at room temperature and chemical shifts were referenced using a 1 M NaCl_(aq) for ²³Na, NaBH_4(s) for ¹¹B, and 1 M LiCl_(aq) for ⁷Li. The spectral deconvolution was carried out using the Dmfit software [50].

Thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) in conjunction with mass spectrometry (MS) were conducted at 10 °C min⁻¹ under an argon flow of 20 mL min⁻¹ using a Mettler Toledo TGA/DSC 3 (Mettler Toledo, Melbourne, Victoria, Australia) coupled with an Omnistar (Pfeiffer) MS. Masses between m/z = 2 and 100 were followed and 40 μL alumina crucibles were used.

The MCM-41 and nanoconfined NaBH₄ morphology and elemental mapping were determined by transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) using a Philips CM200 (Philips, Eindhoven, Netherlands) operated at 200 kV. For TEM analysis, the materials were ground and dispersed in cyclohexane, sonicated for a few seconds, and then dropped onto a carbon coated copper grid. Brunauer–Emmett–Teller (BET) was performed using a Micromeritics TriStar 3000 Analyzer from Micrometrics Instrument Corporation (Norcross, GA, USA).

The ionic conductivity of the materials was determined by electrochemical impedance spectroscopy (EIS) (Biologic, Sydney, NSW, Australia). Here, 30 mg of materials was placed in a 10 mm diameter die and uniaxially cold pressed at 9 MPa with a hydraulic press to make a pellet of 0.03 cm thickness and 0.712 cm². Two polished sheets of stainless steel were used as electrodes. The pellet was then placed into a controlled environment sample holder (CESH) from BioLogic (Biologic, Sydney, NSW, Australia). The cell was assembled in the glove box under argon atmosphere. EIS was conducted using an alternating current impedance spectroscopy method with a VMP3 potentiostat from BioLogic. The AC impedance measurement was set from 100 mHz to 1 MHz. The measurement was conducted in the temperature range from 25 to 135 °C with an interval of 10 °C. Before each measurement, the sample dwelled for 20 min for temperature equilibration.

4. Conclusions

In conclusion, we encapsulated NaBH₄ into MCM-41 and the resulting NaBH₄@MCM-41 material was expected to lead to higher Na⁺ ionic conductivities. However, any enhancement in ionic conductivity was found to be hindered by the formation of oxide phases in NaBH₄@MCM-41, and similar observations made with LiBH₄@MCM-41 tend to prove this hypothesis. We thus propose that the inevitable oxidation between complex borohydrides and MCM-41 may have a negative effect on the ionic conductivity. However, the presence of the large cage-like anions (B₁₂H₁₂)²⁻ formed upon the partial decomposition of the borohydrides during their nanoconfinement may lead to a more effective path to improve their ionic conductivity.

Supplementary Materials

The following are available online at https://0-www-mdpi-com.brum.beds.ac.uk/2304-6740/9/1/2/s1, Figure S1: BET analysis for MCM-41 and NaBH₄@MCM-41: (a) N₂ physisorption and (b) pore size distribution; Figure S2: Typical TEM image of (a) the empty scaffold MCM-41 and (b) NaBH₄@MCM-41; Figure S3: STEM elemental mapping of NaBH₄@MCM-41; Figure S4: TGA-DSC-MS profiles of (a) pristine NaBH₄ and (b) NaBH₄@MCM-41; Figure S5: BET analysis for MCM-41 and LiBH₄@MCM-41: (a) N₂-physisorption and (b) pore size distribution; Figure S6: TGA-DSC-MS profiles of (a) pristine LiBH₄ and (b) LiBH₄@MCM-41; Figure S7: FTIR analysis of LiBH₄@MCM-41 and pristine LiBH₄; Figure S8: FTIR analysis of NaBH₄@MCM-41 and pristine NaBH₄; Figure S9: Arrhenius plot of LiBH₄@MCM-41 and pristine LiBH₄; Figure S10: ¹¹B NMR of the NaBH₄ + Na₂B₁₂H₁₂ composite synthesised by exposing NaBH₄ to B₂H₆; Figure S11: Arrhenius plot of NaBH₄ and the NaBH₄+Na₂B₁₂H₁₂ composite synthesised by exposing NaBH₄ to B₂H₆; Table S1: Summary of BET analysis for MCM-41, NaBH₄@MCM-41, and [email protected]

Author Contributions

X.L. performed the research and analyzed the data. A.R. and K.-F.A.-Z. conceived and supervised the work, and analyzed the data. All authors have read and agreed to the published version of the manuscript.

Funding

We acknowledge support from the UNSW Digital Grid Futures Institute, UNSW Sydney, under a cross-disciplinary fund scheme and the ARC Research Hub on Integrated Energy Storage solutions.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. (a) X-ray diffraction (XRD) of nanoconfined NaBH₄ along with that of pristine NaBH₄ and (b) magnification of (a) for 2θ = 23–42°; (c) small angle XRD of MCM-41 and NaBH₄@MCM-41.

Figure 2. FTIR analysis of NaBH₄@MCM-41 and pristine NaBH₄.

Figure 3. Arrhenius plot of NaBH₄@MCM-41 and pristine NaBH₄. The cause for the small “jump” in ionic conductivity for NaBH₄ is unknown as it does not correspond to any known phase transition.

Figure 4. ¹¹B nuclear magnetic resonance (NMR) of pristine NaBH₄@MCM-41 and NaBH₄.

Figure 5. Arrhenius plot of NaBH₄ and the NaBH₄+Na₂B₁₂H₁₂ composite synthesised by exposing NaBH₄ to B₂H₆.

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Investigating the Factors Affecting the Ionic Conduction in Nanoconfined NaBH₄

Abstract

1. Introduction