Systematic design and functionalisation of amorphous zirconium metal–organic frameworks

Abstract

Controlling the structure and functionality of crystalline metal–organic frameworks (MOFs) using molecular building units and post-synthetic functionalisation presents challenges when extending this approach to their amorphous counterparts (aMOFs). Here, we present a new bottom-up approach for synthesising a series of Zr-based aMOFs, which involves linking metal–organic clusters with specific ligands to regulate local connectivity. In addition, we overcome the limitations of post-synthetic modifications in amorphous systems, demonstrating that homogeneous functionalisation is achievable even without regular internal voids. By altering the acidity of the side group, length, and degree of connectivity of the linker, we could control the porosity, proton conductivity, and mechanical properties of the resulting aMOFs.

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Introduction

The directionality of coordination bonding enables predictable and precise positioning of atomic arrangements in three-dimensional space,^1,2 enabling the design of crystalline metal–organic frameworks (MOFs) by selecting appropriate metal ion and multitopic organic ligand building units.^3–9 Additional chemical functionalisation of assembled MOFs is also feasible through post-synthetic modifications.^10,11 The recent emergence of amorphous MOFs (aMOFs), some showing glass transition,^12–14 presents unique or complementary properties, including improved ion mobility, transparency, and processability.^15–20

Typically, aMOFs are synthesised in crystalline form and subsequently amorphised by melt-quenching,^15,21 dehydration or desolvation,^22,23 mechanical milling,²⁴ pressure,^25,26 or radiation.²⁷ Nevertheless, these approaches can only be applied when the MOFs are stable during amorphisation. In practice, the majority of materials are prone to organic linker decomposition. Additionally, controlling their properties and local structure poses a challenge, given their top-down preparation¹⁴ and since post-synthetic modification is typically confined to the outer surface due to the absence of regular pores.²⁸

This article presents a bottom-up strategy for designing a series of zirconium-based aMOFs. The approach uses predefined zirconium oxo clusters (Zr₆) cross-linked with a selection of flexible ligands (Fig. 1). Analysis of the local structures revealed the preservation of the Zr₆ cluster units. Post-synthetic modification is showcased in the gel phase before densification, resulting in homogeneous functionalisation with sulfonic acid groups. The presented strategy shows how the porosity, proton conductivity, and mechanical properties of the formed aMOFs can be systematically controlled via the presence of the side group, ligand length, and degree of connectivity.

Fig. 1 Schematic overview of the synthesis routes and functionalisation of zirconium-based amorphous metal–organic frameworks (aMOFs). The Zr–nL₁, Zr–nL₂, and Zr–nL₃ aMOFs are synthesised by pairing the Zr₆O₄(OH)₄(H₂O)₈(CH₃COO)₈Cl₄ (ZrOAc) cluster with three bidentate ligands, including succinic acid (L₁), adipic acid (L₂), and thiomalic acid (L₃). Post-synthetic modification of a cross-linked Zr–nL₃ in the gel phase provides Zr–nL₄ frameworks with sulfosuccinate (L₄). Zr, O, C, and S atoms are represented by light blue, red, grey, and yellow spheres, respectively. H atoms are omitted for clarity.

Results and discussion

Sol–gel synthesis of cross-linked Zr clusters

Water-soluble Zr₆ oxo cluster, Zr₆O₄(OH)₄(H₂O)₈(CH₃COO)₈Cl₄, denoted as ZrOAc, was synthesised as reported, and its formation was confirmed by powder X-ray diffraction (PXRD), Fourier transform infrared (FTIR), and thermal gravimetric analysis (TGA) (Fig. S1–S3†).²⁹ We selected three flexible aliphatic bidentate ligands—succinic acid (L₁), adipic acid (L₂), and thiomalic acid (L₃)—to prevent the formation of crystalline products (Fig. 1). The ZrOAc clusters were cross-linked by adding 2, 4, 8, and 12 molar equivalents of aqueous solutions of ligand (n ratio) to ZrOAc, followed by heating at 80 °C for 72 h. The process resulted in the formation of gels (Fig. S4†).^30–35 Considering that all starting components are highly soluble in water, the distinct change in solubility suggests the cross-linking of ZrOAc clusters.³⁶ For L₁ and L₂, the minimum n ratio required to initiate gelation is 4, while n = 2 is sufficient for L₃. The gels were then washed and dried at 60 °C for 72 h. The resulting aMOFs are referred to as Zr–nL_x, where n represents the molar equivalent of ligands (L_x = L₁, L₂, or L₃) used during cross-linking. The amorphous nature of all samples was confirmed by the board diffuse scattering patterns observed by laboratory (Fig. S5–S7†) and synchrotron PXRD (Fig. 2A). Optical images of Zr–nL₁ samples show transparent, glass-like monolith, while Zr–nL₂ and Zr–nL₃ became more opaque at higher linker ratios (Fig. 2B and S8–S10†). Scanning electron microscopy (SEM) images of as-synthesised Zr–4L₁, Zr–4L₂, and Zr–4L₃ samples revealed smooth surfaces without any trace of distinguishable microparticles (Fig. S11–S13†). Based on TGA data under N₂ flow before and after dehydration at 200 °C (Fig. S14–S23†), the interstitial water contents were determined at 10.0–15.2 wt%. Additionally, all cross-linked compounds were found to be thermally stable up to at least 220 °C.

Fig. 2 Sol–gel synthesis of cross-linked zirconium-based aMOFs. (A) Synchrotron PXRD (λ = 0.109795 Å) of Zr–4L₁, Zr–4L₂, and Zr–4L₃. (B) Optical image of the as-synthesised Zr–4L₁. Scale bar = 0.5 mm. (C) FTIR spectra of Zr–4L₁, Zr–4L₂, and Zr–4L₃. (D) Zr-cluster model representations for (E and F) peak assignments. The structures are determined from single crystal data of (top) [Zr₆O₄(OH)₄(H₂O)₈(Gly)₈]·12Cl·8H₂O³⁷ and (bottom) [Zr₆O₄(OH)₄(CH₃COO)₁₂]·8.5H₂O.³⁸ Zr, C, and O atoms are represented by light blue, grey, and red, respectively. The longer Zr–μ₃-O pair distance in the trigonal zirconium plane is marked with an asterisk. (E) Pair distribution function (PDF) and (F) extended X-ray absorption fine structure (EXAFS) at the Zr-K edge of ZrOAc, Zr–4L₁, Zr–4L₂, and Zr–4L₃.

The incorporation of the ligands was confirmed by FTIR and ¹H NMR (Fig. 2C and S24–S37†), with ethanol residue detected in Zr–nL₂ and Zr–nL₃. FTIR Peaks at 2926–2936, 1300–1306, and 1153–1173 cm⁻¹ were assigned to the ν(CH), τ(CH₂), and τ(CH₂) vibrations of succinate.^39,40 In Zr–nL₂, an additional ν(CH) band appeared around 2864–2871 cm⁻¹, corresponding to longer aliphatic chain units. For Zr–nL₃, a peak associated with the S–H bond was identified at 2556–2561 cm⁻¹.⁴¹ Additionally, the presence of C=O stretching at 1690–1730 cm⁻¹ suggests the coexistence of non-coordinating carboxylate or carboxylic acid arising from monodentate ligands.⁴² When n = 4, the ratios of incorporated bidentate ligands and acetate were comparable for L₁, L₂, and L₃ (Fig. S38†). Consequently, Zr–4L₁, Zr–4L₂, and Zr–4L₃ were selected for further characterization. The TGA analysis under air indicated the presence of linker vacancies, revealing the ratios between bidentate ligands and Zr₆ clusters in Zr–4L₁, Zr–4L₂, and Zr–4L₃ to be 3.5, 4.3, and 3.7 instead of 6 for full cross-linking (Fig. S39†).⁴³

The thermal behaviour of dehydrated Zr–4L₁, Zr–4L₂, and Zr–4L₃ was further investigated using differential scanning calorimetry (DSC, Fig. S40†).⁴⁴ All samples exhibit an endothermic baseline shift between 104.4 and 112.7 °C at a heating rate of 10 °C min⁻¹. However, these baseline shifts are smaller than previously reported values. For example, the measured change in heat capacity (ΔC_p) of Zr–4L₁, using sapphire reference, is 0.013 J g⁻¹ °C⁻¹ (Fig. S41†), which is significantly lower than the ΔC_p associated with the glass transition temperature in other MOF and coordination polymer glasses, such as 0.16 J g⁻¹ °C⁻¹ in ZIF-4 or 0.34 J g⁻¹ °C⁻¹ in [CsHSO₄]_0.5[ZnPIm]_0.5 glasses.^45,46

Chemical structures and short-range ordering

Using synchrotron X-ray total scattering, we characterised the short-range ordering within Zr–4L₁, Zr–4L₂, and Zr–4L₃. The Fourier transformation and the Lorch modification of the Faber–Ziman total structure factor, S(Q), were employed to obtain atomic pair information represented as pair distribution functions (PDFs, Fig. 2D and E).^47–49 The peak features of Zr–4L₁, Zr–4L₂, and Zr–4L₃ closely resemble those of the pristine ZrOAc cluster up to 5.0 Å. The range includes the retention of Zr–O bonds (2.2 Å), Zr⋯Zr correlations of the adjacent site (Zr–Zr_adj, 3.5 Å), and Zr⋯Zr correlations of the opposite site (Zr–Zr_opp, 5.0 Å). These observations confirm that the Zr₆ building unit is maintained in all aMOFs.

The peak between 6.5 and 6.8 Å (Fig. 2E, labelled with I) in Zr–4L₁, Zr–4L₂, and Zr–4L₃ corresponds to the distance between Zr⁴⁺ and the fourth neighbour C atoms of the stretched bridging ligand and opposite carboxylate O atoms that link the nearby cluster. The pair distance closely aligns with the Zr⋯C (6.4–7.0 Å) and Zr⋯O of the carboxylate (6.9–7.1 Å) observed in Zr-MOFs comprising fumarate or aspartate ligands.^50,51 To support the peak assignments, we conducted additional X-ray total scattering and PDF analysis on two crystalline MOFs with Zr₆ units, MOF-801 and UiO-66 (Fig. S43†), together with a partial PDF simulation for MOF-801 (Fig. S44†). Above 7 Å, the PDF profiles of Zr–4L₁, Zr–4L₂, and Zr–4L₃ show less intense and broader peaks, reflecting the diminishing periodicity characteristic of their amorphous nature.⁵² The peak labelled II in Fig. 2E is the expected position for inter-cluster Zr⋯Zr correlations for Zr–4L₁ and Zr–4L₃. For Zr–4L₂, the same correlation is expected around 9–12 Å due to a wide conformational arrangement of the adipate linker.^40,53

The local coordination environments of the Zr₆ clusters before and after cross-linking were examined using Fourier-transform (FT) Zr K-edge extended X-ray absorption fine structure (EXAFS, Fig. 2F).³⁸ The main characteristic peaks for Zr₆ units at ca. 1.58, 3.10, and 4.64 Å, corresponding to the radial distances of Zr–O, Zr–Zr_adj, and Zr–Zr_opp pairs (Fig. 2D), are observed in all samples: ZrOAc, Zr–4L₁, Zr–4L₂, and Zr–4L₃. The peak assignments were confirmed by EXAFS fitting for Zr–4L₁, Zr–4L₂, and Zr–4L₃ using the Zr₆(O)₄(OH)₄(CH₃COO)₁₂·8.5H₂O³⁸ crystal structure as a model (Fig. S47 and Table S1†). The fitted Zr–Zr distances of 3.54–3.55 Å align well with the pair distances observed in PDF analysis. Apart from the ligand types, the local environments of Zr–4L₁, Zr–4L₂, and Zr–4L₃ are indistinguishable. When comparing the ZrOAc building unit with the cross-linked frameworks, an additional peak at ca. 1.80 Å only observed for ZrOAc is attributed to longer Zr–μ₃-O linkages in the trigonal zirconium plane. The identical elongation of the Zr–O bond is also observed in other Zr₆ clusters with 8 bidentate bridging carboxylate coordination, such as [Zr₆(O)₄(OH)₄(H₂O)₈(Gly)₈]·12Cl·8H₂O and [Zr₆(OH)₈(H₂O)₈(HGly)₄(Gly)₄]·(SO₄)₆·14H₂O (HGly = protonated glycine), where the equatorial edges of both clusters are not coordinated by glycine ligand but are instead capped with 8 water molecules (Fig. S48†).^37,54 This Zr–O bond elongation, however, is absent in the 12-coordinated Zr₆ cluster, such as Zr₆(O)₄(OH)₄(CH₃COO)₁₂·8.5H₂O,³⁸ or in defective Zr₆ clusters capped with water or hydroxide molecules in irregular positions.⁵⁵ These results suggest that cross-linking occurs in an alternate position and differs from the regular 8-coordination in the starting Zr₆ cluster (ZrOAc).

The change in oxidation state and chemical environment was probed by X-ray photoelectron spectroscopy (XPS). The presence of Cl⁻ in ZrOAc (267 and 196 eV) further emphasises that the Zr₆ cluster comprises 8 coordinated water molecules, as counter anions are necessary to compensate for the overall charge balance requirements (Fig. S49†).^37,54 This is not the case for Zr–4L₁, Zr–4L₂, and Zr–4L₃, since Cl⁻ are eliminated after the cross-linking process and the charge neutrality should be maintained by hydroxide anions (Fig. S50–S52†).⁵⁵ XPS further confirms the presence of the S atom in Zr–4L₃ (Fig. S52†), and XPS at the Zr 3d region remains identical to ZrOAc, indicating the retention of the Zr⁴⁺ oxidation state for all samples (Fig. S53–S56†).

Post-synthetic modification

The absence of regular internal voids in aMOF usually impedes the diffusion of reactants, resulting in inhomogeneous functionalisation. To circumvent this limitation, we conducted post-synthetic modifications in the gel phase (Fig. 3A). Oxidation of thiols (–SH) to sulfonic acid groups (–SO₃H) was achieved by soaking Zr–nL₃ gels in a 35% H₂O₂ solution for 2 h.^41,57 The resulting material was washed with deionised water, and protonation was completed with 0.02 M H₂SO₄ (see ESI†). Subsequently, the protonated gels underwent further washing cycles, solvent exchange with ethanol, and densification to obtain Zr–nL₄ (n = 2, 4, 8, and 12; Fig. S57 and S58†). PXRD patterns measured for the resulting material confirmed the retention of its amorphous nature after the modification without forming any crystalline oxide byproduct (Fig. 3B and S59†). Additionally, TGA results revealed that Zr–nL₄ samples are thermally stable to at least 250 °C and contain 15–17 wt% of interstitial water (Fig. S60–S63†).

Fig. 3 Post-synthetic modification of cross-linked zirconium-based aMOF. (A) Schematic overview of the oxidation of thiols (–SH) to sulfonic acid groups (–SO₃H) in gel phase before densification. (B) Synchrotron powder X-ray diffraction (PXRD, λ = 0.109795 Å) of Zr–4L₃ and Zr–4L₄. (C) FTIR spectra of Zr–4L₃ and Zr–4L₄. XPS spectra at (D) Zr 3d and (E) S 2p of Zr–4L₃ and Zr–4L₄. (F) PDF and (G) EXAFS at Zr K-edge of Zr–4L₃ and Zr–4L₄. (H) Model for EXAFS fitting derived from the crystal structure of [Zr₆O₄(OH)₄(OH₂)₈(HCOO)₄(SO₄)₄]·2HCl·3H₂O,⁵⁶ with Zr, C, O, and S atoms represented by light blue, grey, red, and yellow, respectively. (I) Fourier transforms (magnitude and real component) of EXAFS at the Zr K-edge with fitting curves for Zr–4L₄. The fitted results are presented in Table S2.†

Initial evidence of complete thiol oxidation was obtained from FTIR spectra (Fig. 3C and S64–S67†). The S–H band (2556–2561 cm⁻¹) was replaced by newly formed bands centred at 1205, 1125, and 1038 cm⁻¹, corresponding to v_as(O=S=O), vs.(O=S=O), and S–O stretching, respectively.^41,58 Additionally, the post-synthetic modification eliminated non-coordinating carboxylate or carboxylic acid, as C=O stretching at 1690–1730 cm⁻¹ disappeared in all samples. Following this confirmation, Zr–4L₄ was chosen for further characterisation. XPS confirmed no changes in the oxidation state of Zr⁴⁺ (Fig. 3D), while the S 2p peaks shifted to a higher binding energy after the oxidation treatment (Fig. 3E), with S 2p_3/2 shifting from 163.4 eV of –SH to 168.8 of –SO₃H in Zr–4L₃ and Zr–4L₄, respectively.^41,59 The ¹H NMR of the digested Zr–4L₄ confirmed a complete conversion, as evidenced by the shift of the CH peak adjacent to the functional group from 3.50–3.54 ppm in Zr–4L₃ (Fig. S35†) to 3.93–3.97 ppm in Zr–4L₄ (Fig. S69†). TGA analysis under air indicated a partial linker elimination after the treatment, as the ratio between bidentate ligands and Zr₆ clusters decreased to 2.8 in Zr–4L₄ (Fig. S70†).

The local structure after post-synthetic modification (Zr–4L₄) was analysed through X-ray total scattering (Fig. 3F) and EXAFS (Fig. 3G). PDF data revealed a distortion in the Zr₆ node, as evidenced by a shift in the Zr–Zr_adj peak from 3.53 Å in Zr–4L₃ to 3.49 Å in Zr–4L₄, along with the emergence of a peak at 4.0 Å. A minor contraction in the Zr–O distance was also observed, decreasing from 2.21 to 2.18 Å. While the Zr–O, Zr–Zr_adj, and Zr–Zr_opp peaks are retained, their relative intensities are altered. In the FT-EXAFS (Fig. 3G), we observed a decrease in the peak intensities for Zr–Zr_opp pairs and a splitting of the Zr–Zr_adj peak in Zr–4L₄ compared to Zr–4L₃. The relative intensity of the Zr–O peak was increased in both PDF and EXAFS results. This is explained by incorporating capping water or hydroxide ions around Zr⁴⁺ to complete the coordination sphere following partial ligand elimination.⁵⁵ Similar node distortions and alterations in PDF peaks have been observed in a series of crystalline NU-1000 MOFs, where varying node capping ligands influenced the extent of ligand elimination during thermal treatment.^60,61 No amorphous ZrO₂ or Zr(OH)₄ byproducts were observed (Fig. S71†).

Another cause of the node distortion is the coordination of sulfate anions to the outer surface of Zr₆ clusters during the protonation step. EXAFS fitting confirmed the incorporation of sulfate anions in Zr–4L₄, with the data fitted well to the crystal structure of the [Zr₆O₄(OH)₄(OH₂)₈(HCOO)₄(SO₄)₄]·2HCl·3H₂O⁵⁶ (Fig. 3H, I and Table S2†), showing a reduced Zr–Zr distance of 3.47–3.50 Å. Similar behaviour has been observed in MOF-808 (Zr₆O₅(OH)₃(BTC)₂(HCOO)₅(OH₂), BTC = benzenetricarboxylate), where sulfate anions replaced coordinated formate groups on Zr₆ clusters after washing with dilute sulfuric acid.^62,63 These sulfate anions increase node distortion by coordinating with the Zr₆ clusters in multiple binding modes.

Porosity, proton conductivity, and mechanical properties

The influence of the ligand on gas-accessible porosity was investigated via N₂ physisorption. All samples were activated at 130 °C for 12 h under a dynamic vacuum. At 77 K, Zr–4L₁ and Zr–4L₂ with aliphatic ligands exhibited negligible N₂ adsorption (Fig. S72†). Positron annihilation lifetime spectroscopy (PALS) on Zr–4L₁ and Zr–4L₂ was used to further probe the free spaces in both samples that are quasi-inaccessible to probe molecules (Fig. S73 and S74†).^35,64 PALS detects free space within materials by measuring the lifetime of ortho-positronium (o-Ps) when exposed to a ²²Na positron source. The average pore diameter of each sample is calculated from the lifetime of o-Ps, with a longer lifetime corresponding to a larger average pore diameter. The o-Ps lifetime of Zr–4L₁ is 2.90 ± 0.04 ns, corresponding to an average pore diameter of 3.57 ± 0.03 Å. The longer ligand in Zr–4L₂ results in a longer o-Ps lifetime of 3.71 ± 0.09 ns, translating to a pore diameter of 4.08 ± 0.05 Å. These results confirm the presence of free spaces in both samples that are large enough to accommodate N₂ (3.6 Å) but remain inaccessible.

At 77 K, Zr–4L₃ and Zr–4L₄ show N₂ uptakes of 22.2 and 17.2 cm³ g⁻¹, respectively (Fig. S72†). The presence of side groups likely introduces steric hindrance, preventing pore collapse and preserving space large enough for N₂.⁶⁵ While Zr–4L₃ displays a type I isotherm characteristic of microporous materials, the type IV isotherm observed in Zr–4L₄ further confirms the partial elimination of bidentate linkers and the formation of mesopores following post-synthesis modification.⁶⁶

Promoting proton (H⁺) conductivity in crystalline MOFs is achievable by increasing the H⁺ carrier concentration or establishing continuous H-bonding networks.^6,71 However, no example of systematic H⁺ conductivity control by gradually modifying the structure in aMOF has been demonstrated.^6,13,71,72 We chose Zr–4L₁, Zr–4L₃, and Zr–4L₄, with comparable ligand lengths, to investigate the effects of tethered functional groups in aMOFs on H⁺ conductivity (Fig. 4A). We measured the H⁺ conductivity of the aMOFs by variable-temperature alternating current (AC) impedance spectroscopy under a humidified atmosphere (95% RH). The H⁺ conductivity of Zr–4L₁ is 6.7 × 10⁻³ mS cm⁻¹ at 30 °C and increases to 8.5 × 10⁻² mS cm⁻¹ at 70 °C. The activation energy for H⁺ conduction (E_a, Fig. S75†) is 0.63 eV. In this case, H⁺ mainly diffuses through the H-bonding network between interstitial water guest molecules and non-coordinating or protonated carboxylates since the main alkyl chain of the L₁ ligand contains no tethered functional groups to act as a hopping site.⁶ Note that Zr–4L₁ contains ca. 10.7 wt% water molecules (Fig. S14†). Zr–4L₃ contains weakly acidic thiol (–SH, pK_a ≈ 7.8) side groups, contributing as an additional hopping site. The H⁺ conductivity values of 1.7 × 10⁻² mS cm⁻¹ and 9.9 × 10⁻² mS cm⁻¹ were achieved at 30 °C and 70 °C, respectively. Replacing thiols with sulfonic acid groups (–SO₃H, pK_a ≈ −1.7), a strong Brønsted acid site, in Zr–4L₄ together with the incorporation of sulfate anion on the surface of the Zr₆ clusters further increases the conductivity to 3.1 × 10⁻² mS cm⁻¹ and 0.16 mS cm⁻¹ at 30 °C and 70 °C. Sulfonic acid groups also contribute to the hydrophilic character of Zr–4L₄, with an interstitial water content of 17 wt%. Incorporating functional groups in Zr–4L₃ and Zr–4L₄ reduced the E_a to 0.44 and 0.40 eV, respectively. The order of H⁺ conductivity and inverse E_a for proton conduction in this series is Zr–4L₁ < Zr–4L₃ < Zr–4L₄. Thus, the H⁺ conductivity is controlled via the acidity of the side group, even in amorphous forms.

Fig. 4 (A) Variable-temperature H⁺ conductivity measurements at 95% RH of Zr–4L₁, Zr–4L₃, and Zr–4L₄. The activation energy for H⁺ conduction was calculated from Arrhenius plots (Fig. S75†). (B) Hardness values obtained with the triangular pyramid nano-indenter with a 115° tip angle of Zr–4L₁, Zr–4L₂, Zr–4L₃, and Zr–4L₄. Data points represent mean hardness with 95% confidence intervals. Reference hardness data from other CP/MOF glasses are taken from ref. 40, 45 and 67–70.

In addition to porosity and H⁺ conductivity, we investigated the correlation between the choice of ligands and bulk mechanical properties of the aMOFs using nano-indentation equipped with a standard triangular pyramid indenter with a 115° tip angle (Fig. 4B). From load-depth curves, we determined the mean hardness (HT115) for Zr–4L₁, Zr–4L₂, Zr–4L₃, and Zr–4L₄ as 0.82, 0.39, 0.57, and 0.43 GPa, respectively. A notable decrease in hardness was observed in Zr–4L₂ compared to Zr–4L₁, attributed to an increase in ligand length. A similar trend in mechanical properties was also noted in Zr₆-based crystalline MOFs when comparing the theoretical average bulk modulus of 40 GPa in UiO-66 with 17.4 GPa in UiO-67 [Zr₆O₄(OH)₄(bpdc)₆] (bpdc: 4,4′-biphenyl dicarboxylate), an extended counterpart.^73–75 Beyond ligand length, functional side groups and interstitial molecules also decreased the mechanical hardness, likely due to steric hindrance.^70,75 This hypothesis is supported by increased N₂ uptake in Zr–4L₃ and Zr–4L₄ compared to Zr–4L₁ and Zr–4L₂. However, a substantial decrease in mechanical modulus is evident when framework connectivity is lowered,⁷⁶ elucidating the behaviour observed in Zr–4L₃ and node-distorted Zr–4L₄.

Conclusions

We demonstrated a method for systematically designing and synthesising a series of aMOFs by crosslinking a well-defined metal cluster with various flexible carboxylate ligands. After gelation and densification, glass-like monolithic products were formed. Synchrotron X-ray total scattering and EXAFS revealed the preservation of Zr₆ cluster units after forming extended networks in all samples, even without long-range ordering. To demonstrate the versatility of our approach, we conducted post-synthetic thiol oxidation to obtain sulfonic and functionalised aMOFs. Finally, we established a correlation between the choice of functional groups, ligand length, porosity, proton conductivity, and mechanical properties. Our results indicate that the properties of aMOFs can be modified in a way comparable to their crystalline counterparts, emphasising the impact of precursor selection. This work thus establishes a foundation for the systematic design of novel amorphous metal–organic hybrid systems.

Data availability

The data supporting this article are available in the main text and ESI.†

Author contributions

N. M. and R. A. designed the project. N. M. synthesised, analysed, and characterised samples. S. K., S. H., and N. M. collected and analysed XAFS and PDF data. J. T. carried out NMR measurements. N. G. and T. H. conducted XPS measurements. K. O. conducted gas sorption experiments. The manuscript was written with contributions from all authors. All authors have approved the final version of the manuscript.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors acknowledge the Flemish Research Foundation (FWO Vlaanderen) for support in the EOS research project G0H0522N, KU Leuven for research project C14/20/085, and VLAIO for research project HBC.2021.0580. N. M. acknowledges the support from the Flemish Research Foundation (FWO Vlaanderen) for a Junior Postdoctoral Fellowship (1280924N) and a travel grant for a short stay abroad (K1DE523N), from ICYS for a research fellowship, and from the JSPS KAKENHI Grant Number JP24K23109. S. K. acknowledges the JSPS postdoctoral fellowship. J. T. acknowledges the support from FWO Vlaanderen and F. R. S.-FNRS under the Excellence of Science (EOS) program of a doctoral fellowship (G0H0522N). We acknowledge BL04B2 beamlines at SPring-8 for the synchrotron X-ray total scattering experiments with the approval of JASRI (Proposal No. 2023B1721). We acknowledge BL14B2 beamlines at SPring-8 for the XAFS measurements (Proposal No. 2023B1750). We acknowledge the NMR core at Hasselt University. We acknowledge the European Synchrotron Radiation Facility (ESRF) for the provision of synchrotron radiation facilities under proposal number CH-6893 (DOI 10.15151/ESRF-ES-1442319021), and we would like to thank Pierre-Olivier Autran for assistance and support in using beamline ID11. We acknowledge Dr Takashi Nakanishi for providing access to PALS measurements. We acknowledge Dr Renzhi Ma for providing access to TGA measurement under air.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4sc05053c

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