Metal–Organic Nanosheets Formed via Defect-Mediated Transformation of a Hafnium Metal–Organic Framework

We report a hafnium-containing MOF, hcp UiO-67(Hf), which is a ligand-deficient layered analogue of the face-centered cubic fcu UiO-67(Hf). hcp UiO-67 accommodates its lower ligand:metal ratio compared to fcu UiO-67 through a new structural mechanism: the formation of a condensed “double cluster” (Hf12O8(OH)14), analogous to the condensation of coordination polyhedra in oxide frameworks. In oxide frameworks, variable stoichiometry can lead to more complex defect structures, e.g., crystallographic shear planes or modules with differing compositions, which can be the source of further chemical reactivity; likewise, the layered hcp UiO-67 can react further to reversibly form a two-dimensional metal–organic framework, hxl UiO-67. Both three-dimensional hcp UiO-67 and two-dimensional hxl UiO-67 can be delaminated to form metal–organic nanosheets. Delamination of hcp UiO-67 occurs through the cleavage of strong hafnium-carboxylate bonds and is effected under mild conditions, suggesting that defect-ordered MOFs could be a productive route to porous two-dimensional materials.


Measurements
N 2 adsorption isotherms were carried out at 77 K on a Micromeritics 3Flex gas adsorption analyser. Samples were degassed in situ under vacuum at 120 • C for 20 h using the internal turbo pump. Warm and cold free-space correction measurements were performed by using ultrahigh purity He gas (grade 5.0, 99.999% purity). Ultrahigh purity N 2 (i.e. 99.9992%) was provided by Air Products.

Molecular Simulation
The simulated adsorption of N 2 was investigated using grand canonical Monte Carlo (GCMC) simulations S1 performed in the multi-purpose code RASPA S2 . We used an atomistic model of the hcp UiO-67 MOF structure for which the framework atoms were kept fixed at the crystallographic positions. We used the standard Lennard-Jones (LJ) 12-6 potential to model the interactions between the framework and fluid atoms. In addition, a Coulomb potential was used for N 2 /N 2 interactions. The parameters for the framework atoms were obtained from the UFF S3 or Dreiding S4 force field and N 2 was modelled using the TraPPE potential with charges placed on each atom and at the centre of mass S5 (Supporting Table 1). The Lorentz-Berthelot mixing rules were employed to calculate fluid/solid LJ parameters, and LJ interactions beyond 12.8 Å were neglected. The Ewald sum method was used to compute the electrostatic interactions. Up to 80,000 Monte Carlo cycles were performed, the first 50% of which were used for equilibration, and the remaining steps were used to calculate the ensemble averages. Monte Carlo moves consisted of insertions, deletions, displacements, and rotations. In a cycle, N Monte Carlo moves are attempted, where N is defined as the maximum of 20 or the number of adsorbates in the system. To calculate the gasphase fugacity we used the Peng-Robinson equation of state S6 . Geometric surface areas were calculated by rolling a 3.681 Å diameter sphere, which corresponds to a nitrogen molecule, across the surface of the material S7 . Supporting

in situ synthesis analysis
Full pattern analysis of powder patterns was performed using the Pawley method within the Topas Academic software to determine lattice parameters S8 . The peak shape was allowed to vary freely throughout the refinement. The background was modelled using a combination of a broad Gaussian function, to account for the scattering due to the apparatus and solvent, and a freely refining Chebyshev polynomial. An additional broad Gaussian peak was used to model the small angle scattering present early in the reaction. Additional peaks due to the instrumental background and undissolved ligand were modelled as Gaussians, with width and intensity freely refining. Selected patterns were also refined using the Rietveld method in order to quantify phase fractions. Initially the kinetics of the four distinct phases: SAXS scattering, UiO-67, hcp UiO-67 and ligand were modelled using the model proposed by Gualtieri S9 : (1) A is a normalisation factor, N is the dimensionality of crystal growth, k g is the rate constant of crystal growth and a and b are constants related to nucleation. The induction time is dependent on the nucleation process and so is determined by constants a and b. This five parameter fit was found to over-parameterise the time evolution, and so a reduced functional form was used, Where A and k g have equivalent meaning, but the induction time is explicitly accounted for by the constant t i . t f as discussed in the main text is t f = 1 kg . Both functional forms were unable to account for the growth of hcp UiO-67.

Powder diffraction analysis
All refinements were carried out using Topas Academic 4.1 S8 . The background was modelled using a combination of a broad Gaussian function and a freely refining Chebyshev polynomial (nine parameters, except for hxl, for which seven parameters was sufficient). A small quantity of unreacted ligand was modelled as a single Gaussian peak with width and intensity freely refining. In order to account for the anisotropic broadening observed in the hxl UiO-67 due to reduced order in the c direction, an additional Gaussian size broadening term was also included in the refinement with its orientational dependence modelled using fourth order spherical harmonics. A small impurity peak was detected in the hcp phase and fitted separately as a Gaussian. This broadening corresponded to a size of approximately 40(5)nm. Rietveld refinements were carried out from the DFT optimised structural models, with the position and isotropic displacement parameter of Hf atoms allowed to refine. The observed discrepancies in the low Q peak intensities, particularly acute for hxl UiO-67, are likely due to the presence of large quantities of disordered guests, as often observed in porous materials. We were not able to crystallographically identify the guests through structural modelling, but structural envelopes derived from difference Fourier maps showed that residual electron density was primarily present within the pores S10, S11 . Although for fcu UiO phases there is uncertainty over the proton configuration, this uncertainty is eliminated for hcp because the two oxygens which are facing each other cannot be protonated (c.f. Ref. S 12), and so we did not introduce additional disorder into the structure.
Supporting 0.529775 0.182476 0.250000 2 Estimated standard errors are given in parentheses. * Hf displacement parameters constrained to be identical.
Supporting   The experimental powder diffraction data does not allow for discrimination between these models. The random stacking model was generated by sequentially adding layers of either single clusters (i.e. generating condensed clusters) or ligands and clusters at random. As only (hk0) reflections are present, the experimental diffraction data is sensitive primarily to the projected electron density in the (001) plane. This is in large part responsible for the similarity of the diffraction patterns calculated from single layer sheets.