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Extending the Compositional Range of Nanocasting in the Oxozirconium Cluster-based Metal-Organic Framework NU-1000 – A Comparative Structural Analysis Wenyang Zhao, Zhao Wang, Camille D Malonzo, Thomas E. Webber, Ana E. PlateroPrats, Francisco Sotomayor, Nicolaas A. Vermeulen, Timothy C. Wang, Joseph T. Hupp, Omar K. Farha, R. Lee Penn, Karena W Chapman, Matthias Thommes, and Andreas Stein Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b04893 • Publication Date (Web): 29 Jan 2018 Downloaded from http://pubs.acs.org on January 30, 2018

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Chemistry of Materials

Extending the Compositional Range of Nanocasting in the Oxozirconium Cluster-based Metal-Organic Framework NU-1000 – A Comparative Structural Analysis Wenyang Zhao,†,1 Zhao Wang,†,1 Camille D. Malonzo,1 Thomas E. Webber,1 Ana E. PlateroPrats,2 Francisco Sotomayor,3 Nicolaas A. Vermeulen,4 Timothy C. Wang,4 Joseph T. Hupp,4 Omar K. Farha,4,5 R. Lee Penn,1 Karena W. Chapman,2 Matthias Thommes,3 Andreas Stein*,1 1

Department of Chemistry, University of Minnesota, 207 Pleasant St. SE, Minneapolis, MN

55455, U.S.A. 2

X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, 9700 S. Cass

Ave., Argonne, IL 60439, U.S.A. 3

Quantachrome Instruments, 1900 Corporate Drive, Boynton Beach, FL 33426, U.S.A.

4

Department of Chemistry and Chemical and Biological Engineering, Northwestern University,

2145 Sheridan Road, Evanston, Illinois 60208, U.S.A. 5

Department of Chemistry, Faculty of Science, King Abdulaziz University, Jeddah, Saudi Arabia

†

These authors contributed equally.

*

Corresponding author: [email protected]

1 ACS Paragon Plus Environment

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ABSTRACT The process of nanocasting in metal-organic frameworks (MOFs) is a versatile approach to modify these porous materials by introducing supporting scaffolds. The nanocast scaffolds can stabilize metal-oxo clusters in MOFs at high temperatures and modulate their chemical environments. Here we demonstrate a range of nanocasting approaches in the MOF NU-1000, which contains hexanuclear oxozirconium clusters (denoted as Zr6 clusters) that are suitable for modification with other metals. We developed methods for introducing SiO2, TiO2, polymeric, and carbon scaffolds into the NU-1000 structure. The responses of NU-1000 towards different scaffold precursors were studied, including the effects on morphology, precursor distribution, and porosity after nanocasting. Upon removal of organic linkers in the MOF by calcination/pyrolysis at 500 ˚C or above, the Zr6 clusters remained accessible and maintained their Lewis acidity in SiO2 nanocast samples, whereas additional treatment was necessary for Zr6 clusters to become accessible in carbon nanocast samples. Aggregation of Zr6 clusters was largely prevented with SiO2 or carbon scaffolds even after thermal treatment at 500 ˚C or above. In the case of titania nanocasting, NU1000 crystals underwent a pseudomorphic transformation, in which Zr6 clusters reacted with titania to form small oxaggregates of a Zr/Ti mixed oxide with a local structure resembling that of ZrTi2O6. The ability to maintain high densities of discrete Lewis acidic Zr6 clusters on SiO2 or carbon supports at high temperatures provides a starting point for designing new thermally stable catalysts.


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1. INTRODUCTION As an approach for shaping materials, casting has been used for centuries in manufacturing and materials processing. The casting process involves the formation of a mold that predefines the product structure, and filling the mold with fluid precursors that are then allowed to solidify. Depending on the shape and size of the mold used for casting, various materials with different structures can be obtained by well-defined processes. The same idea can also be applied to the area of nanomaterials synthesis and nanofabrication. If critical features of the mold are scaled down to the nanometer level, the mold acts as a hard template in a process referred to as nanocasting. Similar to conventional casting, nanocasting provides a relatively general approach to synthesize materials that maintain certain features of structure and morphology of the original template. It has been applied to prepare porous three-dimensional materials with complex structures that are sometimes difficult to synthesize using direct methods. Ordered mesoporous silica, carbon, and colloidal crystals are some examples of templates for nanocasting.1-3 Products of nanocasting include a wide range of oxides, metals and other compositions that target applications in catalysis, photonics, and separations.4-10 Typically, in nanocasting, the template consists of a single component (e.g., carbon) that is completely removed after the nanocasting process, with only structural information being utilized and replicated. One approach to utilize a template more efficiently involves multi-component templates, so that it is possible to selectively remove one component at a time, possibly leaving another component intact in the final product to provide a specific function.11 Metal-organic frameworks (MOFs), also known as porous coordination polymers (PCPs), are organic-inorganic hybrid materials that consist of organic linkers as ligands and metal ions or metal-oxo clusters as coordination centers. As one of the important categories of porous materials, MOFs have attracted much attention in the past two decades.12-14 Owing to the large variety of organic linker molecules, metal species, and complex coordination environments between them, MOFs with different three-dimensional structures have been synthesized and characterized. These materials have porous structures with pore sizes ranging from a few angstroms to several nanometers.15 The well-defined pores provide high surface areas, leading to potential applications in gas storage,16-17 separation,18-19 heterogeneous catalysis20-23 and chemical sensors.24 MOFs can be further altered by post-synthetic modifications, including decorations of the metal nodes with other active metals or exchange of the linkers.25-26 For example, atomic layer deposition in MOFs 3 ACS Paragon Plus Environment


(AIM) and solvothermal deposition in MOFs (SIM) have been reported as efficient ways to introduce active species onto the metal nodes, so that novel properties can be created that extend the applications of the original MOFs.27-29 Furthermore, MOFs can also act as sacrificial templates for pseudomorphic transformations, during which the components of the MOFs are converted to other materials while the general morphology is conserved.30 In this way, the porous structure and high surface areas of MOFs are largely maintained in the target materials, but the composition is modified to provide new functionality. Nanocasting MOFs followed by template removal is another promising approach to modify MOFs, which may lead to products with different structures and morphologies. Early nanocasting studies focused on the synthesis of porous carbon materials using a variety of MOFs as the sacrificial template, such as MOF-5, Al-PCP and ZIF-8.31-33 Various carbon sources can be selected for nanocasting as well, including furfuryl alcohol, glycerol, glucose, phenolic resins and resorcinol.31,

33-35

Pyrolysis is generally applied to effect the conversion, given that both the

precursors and the organic linkers from MOFs can be carbonized when treated at a high temperature. Apart from carbon, the metal oxide titania has also been reported as a product from nanocasting in MOFs, where microporous titania was synthesized using HKUST-1 as a template.36 Both an acid and an oxidant were employed to remove the metal nodes from the MOF, with only titania left in the final product as the inverse replica. One important point to address is that the cost of synthesizing MOFs is often much higher than the value of the corresponding templated products, so that the economics of using MOFs merely as sacrificial templates may be unfavorable. A significant component of the MOF that has been overlooked in this sacrificial approach is the metal or metal-oxo node. These nodes are connected by the organic linkers and evenly distributed throughout the entire MOF particles. Together with the well-defined pores and channels in their structures, they provide important functionality, such as high selectivity and activity towards catalytic reactions.18, 21 However, MOFs are not suited for high-temperature gas-phase reactions because they are usually stable only below 350 ˚C, especially in oxidizing environments.37 At higher temperatures, the organic linkers in MOFs decompose and the metal nodes eventually agglomerate into larger metal or metal oxide crystallites. This irreversible process is thermodynamically favored and reduces the amount of accessible metal sites, thus compromising the catalytic activity of MOFs. Nanocasting in MOFs is a promising strategy to stabilize the structure and maintain the activities of MOF-derived catalytic sites at high 4 ACS Paragon Plus Environment

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temperatures (see Figure 1). It has been demonstrated through the formation of a silica skeleton in the mesopores of the MOF NU-1000, which was able to keep the catalytically-active oxozirconium clusters isolated and accessible even after high temperature treatment at 500 ˚C or higher.11, 38 In order to extend the concept of nanocasting in MOFs to other skeletal compositions, an understanding of the precursors, the interactions between the MOFs and the precursors, and the ability of the skeleton to stabilize the structure at high temperature is necessary. This article provides a comprehensive study of these factors.

Figure 1. Schematic showing the concept of nanocasting in the MOF NU-1000.

Here we chose the MOF NU-1000 as the template for nanocasting and studied products obtained with different precursors, including tetramethyl orthosilicate for SiO2 nanocasting, tetrakis(dimethylamido) titanium and titanium ethoxide for TiO2 nanocasting, and furfuryl alcohol for polymer/carbon nanocasting. NU-1000 consists of oxozirconium clusters ([Zr6(µ3-O)4(µ3OH)4(OH)4(OH2)4]8+, also denoted as Zr6 node) and 1,3,6,8-tetrakis(p-benzoate)pyrene linkers (TBAPy4-). The Zr6 nodes are unsaturated and 8-connected by the pyrene linkers, forming regular arrays of hexagonal and triangular channels along the c-axis of the crystal.39 NU-1000 is suitable for nanocasting because of its relatively large hexagonal channels (diameter ~3.1 nm) that can 5 ACS Paragon Plus Environment


accommodate a variety of precursors during infiltration processes. These mesopores also provide sufficient capillary forces to prevent precursors from leaching out after infiltration. The precursors inside the channels of NU-1000 are further transformed into a network structure either by gelation or polymerization, with the aim of providing secondary supports that can prevent the Zr6 nodes from aggregating or forming into large crystallites at high temperature. Subsequent calcination (SiO2 and TiO2 nanocasts) or pyrolysis (carbon nanocast) converts the precursors to the corresponding oxide or carbon scaffolds and removes the linkers from the original NU-1000. The scaffold composition can alter the chemical environment of the Zr6 nodes. While the silica scaffold is amorphous and relatively inert, a titania scaffold can, in principle, be crystalline and exhibit redox activity and strong metal-support interactions that alter catalytic properties.40-42 Because of these metal-support interactions, titania-supported clusters, for example, Pt/TiO2,43 Pd/TiO244 and Au/TiO2,45-46 exhibit increased catalytic activity compared to the metal clusters themselves. Carbon was chosen as another possible scaffold to introduce electrical conductivity, because a conductive scaffold may enable electrocatalytic processes, such as the conversion from methane to methanol under mild conditions.47 Here we compare the differences in structure, morphology, and composition of the nanocast materials derived from NU-1000 crystallites that were infiltrated with the different scaffold components listed above. We investigated the interactions between the precursors and the NU1000 host during nanocasting, which led to different distributions of the precursors in the pores of NU-1000 and in some cases altered the nature of the Zr6 nodes. We also demonstrate that, by applying appropriate nanocasting methods, SiO2 and carbon scaffold materials can stabilize Zr6 nodes and prevent their aggregation at high temperature, whereas TiO2 reacts with Zr6 and forms ZrxTi1-xO2, a mixed oxide phase of ZrO2-TiO2 that is stable at high temperature.48 The accessibility of the Zr6 nodes in SiO2 and carbon nanocast samples and the presence of Lewis acidic sites in TiO2 nanocast sample was verified by pyridine adsorption experiments. These nanocasting approaches are deemed suitable for the preparation of single-site catalysts with high thermal stability.


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2. EXPERIMENTAL METHODS 2.1 Materials Chemicals used in this study were obtained from the following sources: ethanol (anhydrous, 200 proof) from Pharmco-AAPER; titanium (IV) ethoxide (TEOT, 95%), titanium (IV) tetrakis(dimethylamide) (TDMAT, >99%) from Gelest; N,N-dimethylformamide (DMF, certified ACS, 99.9%), acetone (certified ACS, 99.7%), pentane (certified, >98%), sodium hydroxide solution (50% w/w, certified) from Fisher Chemical; tetramethyl orthosilicate (TMOS, 98%), furfuryl alcohol (98%), methanol (>99.8%), hydrochloric acid (ACS reagent, 37%), benzoic acid (ACS reagent, 99.5%), zirconium (IV) oxychloride octahydrate (reagent grade, 98%), and pyridine (99.8%) from Sigma-Aldrich; dimethyl sulfoxide-d6 (D, 99.9%) from Cambridge Isotope Laboratories, Inc; 1,3,6,8-tetrakis(p-benzoic acid)pyrene (H4TBAPy) was synthesized following a previously reported procedure.39 Deionized water was purified to a minimum resistivity of 18.2 MΩ·cm with a Milli-Q PLUS reagent-grade water system and was used in all experiments. 2.2 Nanocasting NU-1000 with Different Precursors Silica nanocasting. The synthesis of NU-1000 and nanocasting of NU-1000 with SiO2 were carried out following previously reported methods.11 Details are provided in the Supporting Information. Titania nanocasting. Two different organotitanium precursors were used for nanocasting NU1000, TEOT and TDMAT. For TEOT, 15 mg activated NU-1000 was placed in a 2-mL centrifuge tube and 1 mL TEOT was added. The centrifuge tube was then sealed with parafilm and placed on a vortex mixer until all the solid was suspended in the liquid precursor. After settling for 24 h, the mixture was centrifuged to separate solids and excess precursor material. The residue was washed either once or twice with ethanol, each time followed by centrifugation to isolate the solid. After the last centrifugation step, the centrifuge tube was left open to dry in air for 24 h and then placed in a vacuum oven (200 mTorr) at 80 ˚C for 2 h for further drying. A yellow powder was obtained. The resulting material is denoted as TEOT1@NU-1000 for samples washed once or TEOT2@NU1000 for samples washed twice. Nanocasting with the air-sensitive and moisture-sensitive alternate titania precursor, TDMAT, was conducted in a glovebox under a nitrogen atmosphere. 100 mg TDMAT was added to 30 mg activated NU-1000 and allowed to infiltrate the MOF for 24 h. The infiltrated sample was then washed with pentane, separated by vacuum filtration, exposed to air for 24 h, and then dried in a vacuum oven (200 mTorr) at 80 ˚C for 2 h. The resulting material is 7 ACS Paragon Plus Environment


referred to as TDMAT@NU-1000. All three samples were heated in static air to 500 ˚C in a tube furnace with a heating ramp rate of 2 ˚C/min and maintained at that temperature for 1 h. The resulting white powders are referred to as TEOT1_Zr6@TiO2, TEOT2_Zr6@TiO2 and TDMAT_Zr6@TiO2, respectively. Carbon nanocasting. For nanocasting NU-1000 with carbon, 30 mg of NU-1000 was placed in an uncapped vial and the vial was put into an autoclave containing 10 mL of furfuryl alcohol (FA), taking care to prevent direct contact between the liquid FA and the NU-1000 powder. The autoclave was heated to 90 ˚C to create FA vapor. Different heating times (12 h, 24 h, 48 h, 72 h and 144 h) were examined for the infiltration step because the exposure time to the FA vapor affected the amount of FA loaded inside the pores of NU-1000. The NU-1000 product loaded with FA is denoted as FA@NU-1000. Polymerization of FA inside the pores of NU-1000 was performed in a tube furnace at 250 ˚C for 6 h using a 5 ˚C/min temperature ramp rate and an N2 flow of 1000 mL/min. NU-1000 with polymerized poly(furfuryl alcohol) inside the pores is denoted as PFA@NU-1000-xh (x represents the different infiltration times in hours). Pyrolysis of PFA@NU-1000 was performed in a tube furnace at 600 ˚C for 1 h using a 5 ˚C/min temperature ramp rate and an N2 flow of 1000 mL/min. The product formed from pyrolysis is denoted as Zr6@C-xh. 2.3 Pyridine Adsorption Pyridine was used as the probe molecule to test the presence and accessibility of Lewis and Brønsted acid sites. The Zr6@SiO2 and Zr6@TiO2 samples were first activated at 200 ˚C for 1 h under dynamic vacuum (< 200 mTorr) to remove any residual surface-adsorbed water. Then the samples were cooled down to room temperature and excess pyridine was added to soak the samples for 1 h under static vacuum. These samples were heated again at 200 ˚C for 1 h under dynamic vacuum (< 200 mTorr) to remove physisorbed pyridine and then cooled down to room temperature. Finally, the samples were analyzed by FT-IR spectroscopy. For Zr6@C, prior to the pyridine adsorption experiment, the sample was heated in static air at 300 ˚C for 30 min, with a ramp rate of 14 ˚C/min to increase accessibility of the oxozirconium clusters. 2.4 Characterization FT-IR spectra were obtained in transmission mode on a Nicolet Magna 760 IR spectrometer, using powdered samples compressed between a pair of NaCl windows. Raman spectroscopy was performed on a WITec Alpha300R confocal Raman microscope with 514.5 nm laser radiation, and 8 ACS Paragon Plus Environment

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the scattering was detected using a DV401 CCD thermoelectrically-cooled detector. Powder Xray diffraction patterns were collected using a PANalytical X’Pert Pro diffractometer. The X-rays were generated with a Co anode (Kα radiation, λ = 1.789 Å) which operated at 45 kV accelerating voltage, 40 mA emission current. Scanning electron microscopy (SEM) was performed on a JEOL6700 field emission scanning electron microscope with an accelerating voltage of 5.0 kV. Scanning electron microscopy combined with energy-dispersive X-ray spectroscopy (SEM-EDS) was conducted on a JEOL JXA-8900 electron probe microanalyzer with an Ultradry-SDD detector and an accelerating voltage of 15.0 kV to 20 kV. All samples were coated with a 50 Å platinum film prior to SEM imaging or a 70 Å carbon film for EDS analyses. Gas sorption experiments were performed using ultra-high purity grade argon or nitrogen adsorptive. All samples were degassed under dynamic vacuum (0.003 mTorr) at 120 ˚C for 12 h before analyses. Nitrogen sorption experiments were performed on a Quantachrome Autosorb-iQ2 analyzer. Brunauer-Emmett-Teller (BET) surface areas were evaluated from the adsorption isotherms in the relative pressure range 0.01–0.20. High-resolution argon isotherm measurements were collected using an Autosorb-iQ2 MP analyzer (Quantachrome Instruments) over a relative pressure range of 10-7 to 1. Temperature was controlled using either a liquid argon cryogen bath or a Cryosync (Quantachrome Instruments). Pore size distributions were obtained from argon and nitrogen isotherms by application of proper quenched solid density functional theory (QSDFT) and nonlocal density functional theory (NLDFT) methods assuming argon adsorption at 87 K or nitrogen adsorption at 77 K in cylindrical silaceous pores (all samples) or in cylindrical carbonaceous pores (NU-1000 and PFA@NU-1000 samples). Thermogravimetric analysis (TGA) was performed on a Netzsch STA 409 instrument. Samples were heated in air or nitrogen at a ramp rate of 5 ˚C/min. CHN elemental analysis was carried out at Atlantic Microlab (Norcross, GA). Transmission electron microscopy (TEM) measurements were obtained using a FEI Tecnai T12 transmission electron microscope with an accelerating voltage of 120 kV and a LaB6 filament. Samples were sonicated for 5 min in ethanol and then deposited onto carbon film-coated copper grids. Microtomed samples for TEM were prepared by embedding them in the epoxy resin Polybed 812 and heating at 60 ˚C overnight. The 65 nm sections were cut by using a Leica UC6 microtome and a Diatome diamond knife. Highangle annular dark-field (HADDF) imaging and elemental maps were obtained with an FEI Titan G2 60-300 X-FEG S/TEM operating at an accelerating voltage of 80 kV with an E. A. Fischione annular detector. A probe current of ~0.1 nA was used, and maps were collected with a minimum 9 ACS Paragon Plus Environment


collection time of 5 min. Electrochemical impedance spectroscopy measurements were carried out with a Solartron S1 1287 electrochemical interface and a Solartron 1255B frequency response analyzer. The materials were ground to a powder and pressed into pellets using a hydraulic press at a ram pressure of 10 tons for 5 min. Each pellet contained around 100 mg of sample. The diameter and thickness of the pellet were measured with a vernier caliper. 2.5 Synchrotron X-ray Structure Analysis Powder X-ray diffraction (XRD) data and total scattering data suitable for pair distribution function (PDF) analysis were collected at beamlines 17-BM and 11-ID-B, respectively, at the Advanced Photon Source at Argonne National Laboratory using 27.4 keV (0.45220 Å) and 58.6 keV (0.2114 Å) X-rays, respectively. Data were collected using an amorphous silicon-based area detector. Geometric corrections and reduction to one-dimensional data were carried out using GSAS-II.49 Lattice parameters and peak intensities were extracted from diffraction patterns via Le Bail whole-pattern fitting50-51 based on the reported structural model for NU-1000 (csq topology, P6/mmm, a ~ 40 Å, c ~ 17 Å).52 Lattice and pseudo-Voigt profile parameters were refined over a 0.5-10° 2θ range. Structure envelopes were generated using the intensities of low-index reflections.53-54 Difference electron density (DED) maps53 were then obtained via subtraction of the envelope for pristine NU-1000 from the envelope for infiltrated NU-1000 samples.55 PDFs were obtained from the data within pdfgetX356 to a Qmax=24 Å−1. Differential PDFs were calculated by subtracting the PDF measured for the casting phase from the calcined nanocast samples within Fityk.57 PDFs for structural models were simulated in PDFgui.58

3. RESULTS AND DISCUSSIONS 3.1 Silica Nanocasting The structures and properties of SiO2@NU-1000 and Zr6@SiO2 materials obtained by silica nanocasting in NU-1000 were described in a previous publication.11 Here we briefly review the nanocasting steps and major conclusions for subsequent comparison with the materials obtained by titania and PFA/carbon nanocasting. TMOS was used as the precursor for SiO2, because TMOS molecules are small enough to infiltrate the pores of NU-1000. TMOS condenses slowly at room temperature to afford amorphous SiO2, and this process can be catalyzed by acid. In order to better control the hydrolysis rate of TMOS, HCl vapor treatment was applied instead of using an HCl solution. The resulting material is denoted as SiO2@NU-1000. After linker removal at 500 ˚C, 10 ACS Paragon Plus Environment

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Zr6@SiO2 was obtained. As confirmed by XRD and PDF analyses of synchrotron diffraction data, Zr6 clusters did not aggregate to a significant extent, indicating that SiO2 acted as an inert scaffold to stabilize the clusters. Elemental mapping results showed that SiO2 was homogeneously distributed throughout NU-1000 crystallites with minimal SiO2 on the exterior. Pyridine sorption experiments further confirmed the accessibility of the Lewis acidic Zr6 clusters. 3.2 Titania Nanocasting The organotitanium compounds TEOT and TDMAT were adopted as TiO2 precursors. We estimated the mean molecular diameters for TEOT (1.2 nm) and TDMAT (0.8 nm) from the densities and molecular weights of the corresponding liquids.59 These precursors should be able to penetrate at least the larger mesopores in NU-1000 (diameters: 1.1 nm for the triangular micropores and 3.1 nm for the hexagonal mesopores).60 After infiltration, the high reactivity of the precursors allows them to react with moisture in the air at ambient conditions, making the transformation of precursor to TiO2 fast and efficient. Single or multiple washing steps were employed to remove excess precursor and to adjust the TiO2 content in the sample. Without sufficient washing prior to hydrolysis, TiO2 particles formed on the outside of the NU-1000 crystallites. After the washing step, the precursor was allowed to hydrolyze to afford amorphous TiO2. On the basis of low-angle XRD patterns, the periodicity of the NU-1000 host was maintained after infiltration with TEOT, washing, and condensation (Figure 2). Interactions of the titania precursor with the NU-1000 host resulted in broadening of the low-angle XRD peaks, whose width increased with the extent of titania infiltration into the pores of NU-1000 (see later discussion), with extreme broadening in the case of TDMAT@NU-1000.



Figure 2. XRD patterns of NU-1000 and SiO2, TiO2 or PFA infiltrated NU-1000 as indicated in the figure.

The amount of TiO2 introduced into the NU-1000 host depended on both the organotitanium precursor and the number of washing steps. Analyses of SEM-EDS data (Table 1) indicate that the highest Ti:Zr loading (2.7:1.0) was obtained for TEOT1@NU-1000. However, Raman spectroscopy (characteristic anatase vibrational peaks at 147 cm-1, 398 cm-1, 517 cm-1, and 636 cm-1, Figure S1), XRD (Figure S2), TEM and SAED data (Figure S3) all revealed that after calcination to produce TEOT1_Zr6@TiO2, crystallites in this sample contained a thin external crust of anatase.61 More extensive washing to remove excess precursor (TEOT2@NU-1000) reduced the Ti:Zr ratio to 1.6:1.0 and eliminated this crystalline phase in the calcined materials. However, TEOT2_Zr6@TiO2 showed evidence of reaction between the organotitanium compound and the Zr6 clusters. Although no crystalline component was observed in this material (Figures S2, S3), the Raman spectrum (Figure S1) showed the typical absorptions of ZrxTi1-xO2 with characteristic peaks at 157 cm-1, 268 cm-1, 400 cm-1, 595 cm-1, and 790 cm-1.62 By switching to the TDMAT precursor, a relatively high Ti:Zr atom ratio of 2.0:1.0 was achieved in TDMAT@NU1000 without any evidence for an external TiO2 crust (Figure S3). However, closer examination by pair distribution function (PDF) analysis using synchrotron X-ray scattering revealed that the dominant phase (~67%) in TDMAT_Zr6@TiO2 consists of small domains of ZrTi2O6, whose 12 ACS Paragon Plus Environment

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short-range structure resembles that of srilankite with random Zr and Ti substitutions and some distortion compared to bulk srilankite (Figure 3).63 This observation is evidence for the reaction between TiO2 and Zr6 clusters during calcination at relatively low temperature.64 The estimated size of the ZrTi2O6 domains is ~ 1.5 nm on the basis of the PDF data. In addition, cubic ZrO2 clusters slightly larger than the Zr6 clusters are present (~15 wt%), which are attributed to the aggregation of Zr6 clusters.65 Some monoclinic ZrO2 (~ 16 wt%) and a small fraction (~ 2 wt%) of distorted anatase-like titania are responsible for the PDF features above 1.5 nm. STEM-EDS elemental maps of all TiO2 nanocast samples showed homogeneous distributions of Ti throughout the TiO2@NU-1000 crystallites (Figure S4). Neither X-ray diffraction data nor TEM images revealed any evidence for the formation of more extended ZrO2 phases.



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Table 1. Textural and compositional data of NU-1000, infiltrated NU-1000 materials, and the corresponding calcined (500 ˚C) or pyrolyzed materials (600 ˚C).

Sample

BET area 2

(m /g)

a

Pore

Mesopore a

volume

(cm3/g)

(cm3/g)

volume

a

Composition X:Zrb

NU-1000

2127

1.42

1.19

NU-1000_Cal

57

0.12

0.11

SiO2@NU-1000

1313

0.68

0.61

1.9:1.0c

Zr6@SiO2

444

0.46

0.46

2.2:1.0c

PFA@NU-1000-12

1707

1.10

0.99

15.8:1.0d

Zr6@C-12

287

0.17

0.07

11.1:1.0d

PFA@NU-1000-24

1558

0.97

0.86

16.7:1.0d

Zr6@C-24

295

0.17

0.07

13.1:1.0d

PFA@NU-1000-48

1035

0.76

0.70

24.1:1.0d

Zr6@C-48

350

0.22

0.11

15.5:1.0d

PFA@NU-1000-72

805

0.57

0.56

25.7:1.0d

Zr6@C-72

266

0.16

0.08

16.4:1.0d

PFA@NU-1000-144

131

0.20

0.19

nde

Zr6@C-144

425

0.46

0.40

nde

TDMAT@NU-1000

569

0.29

0.22

2.0:1.0c

TDMAT_Zr6@TiO2

40

0.04

0.03

2.0:1.0c

TEOT1@NU-1000

1252

0.75

0.61

2.7:1.0c

TEOT1_Zr6@TiO2

111

0.12

0.11

2.2:1.0c

TEOT2@NU-1000

1608

0.91

0.76

1.6:1.0c

TEOT2_Zr6@TiO2

144

0.15

0.13

1.8:1.0c

Pore volumes were calculated using the NLDFT, N2 at 77 K, cylindrical model. To allow comparison, a siliceous

kernel was used for all samples in this table and a cutoff of 2 nm was used to determine the mesopore volume. b

X refers to Si, Ti, or C.

c

Determined by SEM-EDS analysis.

d

Calculated through the combination of TGA data and CHN analysis data. For TGA measurements, samples were

heated in air (Figure S6 and Figure S7). The mass did not change above 550 °C, and the residual component was considered to be ZrO2. e

nd = Not determined, because of the large excess of carbon in these samples.


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Figure 3. PDFs for TDMAT_Zr6@TiO2 and simulated PDF patterns for ZrTi2O6 (srilankite, powder-diffraction file 00-046-1265) and a Zr6 cluster, showing Zr/Ti–O, Zr/Ti...Zr/Ti, Zr–O or Zr...Zr correlations. The fitting results indicate that this Zr6@TiO2 sample mainly consists of 1.5 nm ZrTi2O6 (~67 wt%) and cubic ZrO2 clusters (0.9 nm, cubic, a = 4.88 Å, ~15 wt%). Features beyond 1.5 nm correspond mainly to monoclinic ZrO2 (1.8 nm, monoclinic, a = 5.41 Å, b = 5.51 Å, c = 4.30 Å, β = 100º, ~16 wt%) and a small fraction of anatase TiO2 (2.6 nm, tetragonal, a, b = 3.8 Å; c = 9.58 Å, ~2 wt%). Residuals at short range are attributed to a trace amount of carbonaceous material left in the sample (~1 wt%) and possible distortion of the idealized structural components in this material.

Analysis of N2 sorption isotherms confirmed the partial filling of pores by the TiO2 precursors (Figure S5), leading to decreases in BET areas and pore volumes with increasing titania loading (Table 1). The sorption data will be discussed in more detail in Section 3.4.

3.3 Carbon Nanocasting Furfuryl alcohol (FA) has been widely used as a carbon precursor for nanocasting, because the relatively small dimensions of an FA molecule (0.84 Å × 0.64 Å × 0.43 Å) allow it to penetrate a wide variety of nanoporous templates.32,

66-69

For infiltrating NU-1000, we applied a vapor



treatment to introduce FA into its porous structure. During the heating process, FA gradually diffused into the pores of NU-1000 to form FA@NU-1000. To polymerize the FA inside the pores of NU-1000, FA@NU-1000 was heated at 250 ˚C for 6 h to form PFA@NU-1000. The PFA@NU1000 was then pyrolyzed at 600 ˚C for 1 h under N2 to form Zr6@C. The amount of FA introduced into the pores of NU-1000 could be tuned by controlling the infiltration time. With increasing infiltration time from 12 h, 24 h, 48 h, 72 h, to 144 h, the BET area and pore volume progressively decreased, as pores were filled with more PFA (Figure S8 and Table 1). Pore size distribution curves of the PFA@NU-1000 samples showed a gradual shift of the peak associated with mesopores to smaller diameter as the infiltration time was increased. With an infiltration time of 144 h, most pore space was filled with PFA or access to any remaining pores was extensively blocked (Figure S9), which resulted in a dramatically decreased surface area and pore volume (Table 1) compared to shorter infiltration times. The pore filling could also be confirmed with the help of XRD patterns of PFA@NU-1000 samples (Figure S10). The d-spacing for the lowest-angle XRD peak in NU-1000 (010) near 3° 2q corresponds to the spacing of the hexagonal channels (~3 nm). With increasing infiltration time, the intensity of this peak decreased, as a result of contrast matching between the MOF and PFA in the large mesopores, similar to the contrast matching behavior found when silica filled the pores of NU-1000 in SiO2@NU-1000 (Figure 2).11 XRD patterns of the pyrolyzed Zr6@C samples were relatively featureless and did not show any distinct X-ray peaks (Figure S11). In particular, no crystalline ZrO2 phase was observed, which suggests that the PFA/carbon matrix was able to support the Zr6 clusters and keep them from aggregating during the pyrolysis process. This contrasts with pyrolyzed NU-1000, where broad reflections of nanosized tetragonal zirconia are observed in the XRD pattern (particle size estimated to be around 1.5 nm). The presence of mainly isolated Zr6 clusters in the Zr6@C samples was confirmed by PDF analysis (Figure 4, Zr6@C-48h was used for the analysis). After subtracting the pattern of a reference sample, which was derived from pyrolysis of a mixture of PFA and H4TBAPy linker from that of the Zr6@C sample, two pronounced peaks centered at 2.15 and 3.41 Å could be observed, which are associated with the Zr–O and Zr…Zr atom-atom distances, respectively. The fitted simulated pattern corresponds to 0.6 nm (69 wt%), 0.8 nm (27 wt%), 2.5 nm (0.4 wt%) ZrO2 particles as well as 1.7 nm (3.6%) graphitic particles, which indicates that most of the Zr6 clusters were still intact. After the nanocasting and pyrolysis processes, Zr6@C samples 16 ACS Paragon Plus Environment

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retained the rod shapes of the original NU-1000 (Figure S12). On the basis of Raman spectra of the Zr6@C samples, a combination of disordered and graphitic carbon structures was present in these materials (Figure S13), consistent with previous reports on pyrolyzed PFA that during pyrolysis up to 500 ˚C, the furan rings in the polymer chains fracture and are gradually converted into graphitic structures.70-72 The linker molecules from the NU-1000 template consist of aromatic pyrene units that also contribute significantly to the carbon content of the pyrolyzed product. In summary, on the basis of pore size distributions and XRD patterns of the multiple PFA@NU-1000 samples, it could be concluded that the FA molecules penetrate the pores of NU1000, gradually filling more pore space with increasing infiltration time. The optimal infiltration time appears to be between 48 and 72 h to provide sufficient robustness for the carbon skeleton but avoid overfilling, which may compromise the accessibility of the Zr6 clusters.



Figure 4. PDFs for Zr6@C-48h and a reference carbon sample (the reference sample was prepared by pyrolyzing a mixture of PFA and the organic linker H4TBAPy), dPDF, for Zr6@C sample showing Zr–O and Zr…Zr correlations after subtracting the signals from the carbon phase and a simulated PDF pattern for a mixture of oxozirconium clusters (0.6 nm cubic, a = 4.88 Å (~69 wt%) and larger cubic (0.8 nm, cubic, a = 4.88 Å (27 wt%) and 2.5 nm cubic, a = 4.88 Å (0.4%) ZrO2 and 1.7 nm graphitic particles (3.6 wt%)). Residuals at short range are attributed to distortion of the idealized Zr6 clusters64 and the fact that the carbon matrix still contained some hydrogen and oxygen groups that could not be perfectly described by the graphitic phase used in the simulation.

3.4 Comparative Structural Analysis of Infiltrated Materials As a general principle for nanocasting, the introduction of the additional scaffold should not destroy the original structure of the host itself. In the case of nanocasting NU-1000, this can be verified by the XRD patterns of the precursor-infiltrated products. The XRD patterns of the infiltrated and condensed materials matched reasonably well with that of NU-1000 at low angles except for TDMAT@NU-1000 (Figure 2), indicating that the long-range structure of NU-1000 was maintained after precursors condensation. For TEOT1@NU-1000, broadening of the peaks was observed as evidence of strain from TiO2 inside the pores,73 which indicates a moderate loss 18 ACS Paragon Plus Environment

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of crystallinity perpendicular to the channel direction (c-axis) of NU-1000, and is believed to result from strong interactions of TiO2 with the Zr6 clusters in NU-1000 (see below). In the case of TDMAT@NU-1000, the characteristic diffraction peaks were broadened to an extreme extent, which may be caused by the distortion of the unit cell, resulting from the high reactivity of the TDMAT precursor and strong interactions between the precursor and NU-1000. Broadening of the XRD peaks was not observed in the TMOS and PFA infiltrated materials, which suggests weaker interactions between these precursors and NU-1000. The identity and loading of the precursor used for infiltrating NU-1000 had a significant effect on the distribution of the nanocast components in the pore system of the MOF and some effect on particle morphology. The precursor loading during nanocasting influences skeletal integrity, porosity, aggregation of clusters, and accessibility of clusters in the nanocast products. One synthetic target was to achieve compositional uniformity in the products and minimize the formation of SiO2, TiO2, and carbon on the exterior of NU-1000, because an excess crust is likely to block pores and lower the accessibility of Lewis acidic Zr sites. In the optimized samples, SiO2@NU-1000 and TDMAT@NU-1000, the Si:Zr and Ti:Zr atom ratios were 1.9:1.0 and 2.0:1.0, respectively, whereas the C:Zr ratio reached significantly higher values in PFA@NU-1000 (see Table 1). In order to elucidate the distribution of the infiltrated and condensed skeletal components in NU-1000, difference envelope density (DED) analysis of synchrotron X-ray diffraction data was performed on selected composite samples. By calculating the difference between observed and calculated electron densities, DED maps were generated, which provide an average distribution and occupancy of the added components in NU-1000,55 where the extra electron density originates from SiO2, TiO2, or PFA in the pores of NU-1000, respectively. As shown in Figure 5, the DED maps indicate that the silica precursor primarily enters the large hexagonal mesopores of NU-1000 (Figure 5a, d), with the highest density near the hexagonal walls. The carbon precursor in a highly loaded sample (PFA@NU-1000-72h) occupies both hexagonal and triangular pores, as well as the pockets that are formed by four organic linker molecules between Zr6 clusters (Figure 5b, e). The filler produces continuous excess electron density in the mesopore channels. When less PFA is loaded into the structure (PFA@NU-1000-48h), the electron density of infiltrated PFA becomes more localized with less filling of the large pores (Figures S14). In contrast to the silica or PFAloaded NU-1000 samples, the titanium precursor TEOT was found mainly in the pockets between 19 ACS Paragon Plus Environment


adjacent Zr6 nodes. This placement may have been inducive to further reaction of the precursor with aqua and oxo groups on the Zr6 clusters, which resulted in the ZrxTi1-xO2 phase observed for some of the TEOT-infiltrated materials. A DED map for TDMAT@NU-1000 could not be obtained owing to the poor crystallinity of this sample, but it is conceivable that most TDMAT ended up between Zr6 clusters and caused more distortion to NU-1000 because of its higher reactivity compared to TEOT.

Figure 5. Difference envelope density maps of SiO2@NU-1000 (a, d), PFA@NU-1000-72h (b, e), and TEOT2@NU1000 (c, f), viewed from different projections. The purple, pink or yellow colored regions correspond to the electron density differences between each sample and NU-1000. They are projected onto the NU-1000 structure.

Analysis of gas sorption isotherms (Figures 6 and S15) confirmed the introduction of SiO2, TiO2, and PFA into the pore space of NU-1000 and provided additional insight into the component distributions. After nanocasting, all three types of sample showed decreases in BET areas and pore volumes as expected (Table 1). The DFT pore size distribution of NU-1000 shows two main peaks, corresponding to the larger hexagonal mesopores and the smaller triangular micropores. In all cases, a reduction in pore diameter and pore volume of the mesopores is observed. These changes are apparent in pore size distributions derived from both N2 and Ar sorption. For the micropore region, Ar sorption is the preferred method of analysis. Unlike N2, Ar does not possess a quadrupole moment and hence does not exhibit specific interactions with surface functional groups, 20 ACS Paragon Plus Environment

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which leads to a straightforward correlation between pore size/shape and pore filling pressure. Ar adsorption at 87 K fills narrow micropores at significantly higher relative pressures in comparison with nitrogen at 77 K, which shortens equilibration times and permits the measurement of high resolution adsorption isotherms.74 The DFT pore size analysis reveals a decrease in microporosity with the most reduction in micropore volume for PFA@NU-1000-72h, the sample for which the DED maps revealed the highest increase in electron density for the micropore channels. It is interesting to note that in the analysis of Ar sorption isotherms of PFA@NU-1000 samples, application of a siliceous model (Fig. 6b) in the DFT calculations resulted in better fits of the isotherms than a carbon model (Fig. S16), both assuming cylindrical pores. This may be an indication of significant Ar–oxide interactions, implying that Ar could access the Zr6 clusters in the infiltrated MOF. The choice of model (siliceous vs. carbonaceous) affects mainly the interpretation of micropore filling, because here pore filling is strongly influenced by the type of adsorptive-adsorbent surface interactions as well as pore geometry. In the case of mesopores, the adsorption sites/pore walls are already occupied with a mono- or multilayer of adsorbate prior to capillary condensation, and hence the effect of adsorptive-adsorbent surface interaction on the isotherm is reduced.



Figure 6. (a) Ar sorption isotherms, (b) DFT pore size distributions and (c) cumulative pore volumes of NU-1000 and NU-1000 infiltrated with different precursors as indicated in the legend. Pore size distributions were calculated by QSDFT using a siliceous model.75-76


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SEM images for SiO2, TiO2 and carbon nanocast NU-1000 before infiltration, after infiltration and condensation, and after calcination or pyrolysis illustrate the effects of the different precursors on morphology during the nanocasting process (Figure 7). After infiltration and condensation of TMOS, the particle size remained approximately unchanged for SiO2@NU-1000 compared to the original NU-1000 host (Figure 7b, before linker removal), indicating a weak interaction of TMOS with the components of the NU-1000 framework through van der Waals forces, with little preference over linkers or nodes. A similar trend is also observed for carbon nanocasting (Figure 7n, q), where the size of PFA@NU-1000 particles is similar to that of the original NU-1000. In contrast, NU-1000 particles shrank along their channel directions after infiltration with TDMAT@NU-1000 (Figure 7e), but this effect was not observed in TEOT1@NU-1000 or TEOT2@NU-1000 (Figure 7h, k and Figure S17). The TDMAT precursor is more ionic due to the polarized Ti-N bonds and has a higher tendency to bind to oxozirconium clusters, causing the shrinkage of NU-1000 particle (Figure 7e), whereas with less ionic TEOT the particle sizes remained unaffected after infiltration and condensation. The particle size for all samples shrank after the final heat treatment owing to the loss of organic linkers, mostly along the directions perpendicular to the mesopores for SiO2 and carbon, and along all three directions for TiO2. With the protection of the carbon generated from the decomposition of the PFA molecules and the linkers under an inert atmosphere, the size of carbon-nanocast NU-1000 particles changed the least (Figure 7o, r). All TiO2@NU-1000 samples shrank drastically after linker removal, which further confirmed the strong interactions between TiO2 and Zr6 clusters that led to the formation of ZrxTi1-xO2 clusters (Figures 7f, i, l and S17).



Figure 7. SEM images of the NU-1000 host materials used for each of the corresponding nanocasting procedures (a, d, g, j, m, p), nanocast SiO2@NU-1000 (b), TDMAT@NU-1000 (e), TEOT1@NU-1000 (h), TEOT2@NU-1000 (k), PFA@NU-1000-48h (n), and PFA@NU-1000-72h (q), and the corresponding calcined samples Zr6@SiO2 (c), TDMAT_Zr6@TiO2 (f), TEOT1_Zr6@TiO2 (i), and TEOT2_Zr6@TiO2 (l), or pyrolyzed samples Zr6@C-48h (o) and Zr6@C-72h (r). Note that different NU-1000 batches were used as starting materials for silica nanocasting, titania nanocasting, or PFA/C nanocasting.


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3.5 Cluster Stabilization at High Temperatures One of the goals of nanocasting is to utilize scaffolds as supports to stabilize the Zr6 clusters provided by NU-1000 at high temperatures by preventing their aggregation. The clusters must remain isolated from each other even after organic linkers are lost, yet remain accessible and functional. In order to test their limits of stability, the SiO2, TiO2, and PFA infiltrated NU-1000 materials were heated in air (SiO2, TiO2) or nitrogen (PFA/C) at different target temperatures, and structural changes and the development of new phases was monitored by XRD, N2 sorption, and Raman spectroscopy. As stated in Section 3.2, TiO2 nanocast samples formed ZrxTi1-xO2 after linker removal, and the transformation from TiO2@NU-1000 to ZrxTi1-xO2 at calcination temperatures will be discussed. For SiO2 and C nanocast samples, directly after linker removal, no ZrO2 reflection peaks were observed in any of the XRD patterns, only a broad peak at around 35º appeared (Figure 8). By comparison with simulated XRD patterns of ultrasmall particles of tetragonal ZrO2 (Figure S18), the pattern for a single unit cell most closely resembled the experimental patterns of the nanocast samples, supporting the notion that the Zr6 clusters remained mainly non-aggregated in these samples. For comparison, when NU-1000 was calcined at 500 ˚C (Figure S2), distinct reflections for crystalline tetragonal ZrO2 were observed in the XRD pattern as a result of cluster aggregation by thermal condensation after loss of the organic linkers. Thermal stability limits of Zr6@SiO2 and Zr6@C differed when these materials were heated to progressively higher temperatures. As shown in Figure 8a, the scaffold in Zr6@SiO2 remained effective up to at least 600 ˚C. Only after heat treatment at 700 ˚C, small, more distinct diffraction peaks for tetragonal ZrO2 appeared in the XRD pattern of Zr6@SiO2, indicating that Zr6 clusters had begun to aggregate. Based on the broad nature of the diffraction peaks, the crystalline ZrO2 was still only a minor phase. For Zr6@C samples, the thermal stability depended on carbon loading. The PFA@NU-100048h sample with a lower carbon content showed several broad peaks corresponding to the ZrO2 after pyrolysis at 800 ˚C (Figure 8f). However, the more extensively infiltrated PFA@NU-100072h sample remained thermally stable even after heating to 800 ˚C; no peaks for ZrO2 could be observed, indicating the effective isolation of the Zr6 clusters (Figure 8e). As suggested in the introduction, the carbon scaffold in the nanocast sample after pyrolysis may act as a conductive support. Electrochemical impedance spectroscopy (EIS) was carried out on the Zr6@C-72h 25 ACS Paragon Plus Environment


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samples to measure their conductivities (Table 2). The conductivity of Zr6@C-72h_600 was too low to measure. However, when the sample was heated further to 800 ˚C, the conductivity increased to 0.41 S/m. Interestingly, this value is lower than the conductivity of NU-1000 pyrolyzed at 800 ˚C, possibly due to the presence of additional oxygen introduced by the nanocasting precursor. Raman spectra of the Zr6@C and pyrolyzed NU-1000 samples (Figure S19) showed an decrease in the peak ratio ID/IG with increasing pyrolysis temperature as the carbon matrix became more graphitic, consistent with the increase in conductivity observed by EIS. The reasonably high BET area (270 m2/g) and pore volume (0.3 cm3/g) of the Zr6@C-72h_800 sample indicate that the material remained porous after the high temperature treatment.

Table 2. Conductivities of the Zr6@C and control samples sample

conductivity (S/m)*

pyrolyzed NU-1000 (600 ˚C)

too low to measure

pyrolyzed NU-1000 (800 ˚C)

2.13

pyrolyzed mixture of PFA and linker

2.47

Zr6@C_600

too low to measure

Zr6@C_800

0.41

* The conductivity s was calculated using the equation 𝑅 =

$ %&

. Here, R is the resistance, derived from the impedance

spectrum, L the thickness and A the cross-sectional area of the pellet.

In the case of the TiO2-infiltrated materials, no crystalline phases formed after calcination at 500 ˚C (Figure 8b, c, d), except for the external anatase crust observed for TEOT1_Zr6@TiO2. The ZrxTi1-xO2 phase detected by PDF analysis remained amorphous at this temperature. Depending on the precursor and synthesis conditions, a crystalline ZrTiO4 phase appeared at different calcination temperatures. This was first detected by Raman spectra that showed vibrations characteristic for ZrTiO4 at 158 cm-1, 275 cm-1, 405 cm-1, 637 cm-1 and 790 cm-1 (Figure 9). These 26 ACS Paragon Plus Environment

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features were found for TEOT2_Zr6@TiO2 at temperatures as low as 500 ˚C, but for TDMAT_Zr6@TiO2 only at 700 ˚C, at which point the corresponding X-ray diffraction peaks clearly revealed an extended crystalline ZrTiO4 phase (Figure 8b, c, d), which is thermodynamically stable at high temperatures.77



Figure 8. XRD patterns of (a) Zr6@SiO2, (b) TDMAT_Zr6@TiO2, (c) TEOT1_Zr6@TiO2, (d) TEOT2_Zr6@TiO2 and Zr6@C (derived from PFA@NU-1000-72h (e) and PFA@NU-1000-72h (f)) after treatment at elevated temperatures to study thermal stability. The numbers at the end of the sample names indicate the final calcination (a, b, c, d) or pyrolysis (e, f) temperatures.


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Figure 9. Raman spectra of (a) TDMAT_Zr6@TiO2 and (b) TEOT_Zr6@TiO2 calcined at the indicated temperatures. The pronounced slope of the spectra of TDMAT_Zr6@TiO2 is due to the strong fluorescence of these samples under irradiation with the 514.5 nm Raman laser.

N2 sorption isotherms demonstrated that the calcined or pyrolyzed nanocast materials remained porous (Figure S20, Table 1). The BET area was highest for Zr6@SiO2 (~444 m2/g) due to the amorphous nature of SiO2, with a pore volume of ~0.5 cm3/g that was completely due to mesopores. For Zr6@C samples, the BET areas and pore volumes were similar for different samples, averaging ~300 m2/g and ~0.2 cm3/g, respectively. For Zr6@TiO2 materials, the BET areas and pore volumes were lower, likely due to the reaction of the titania precursor with Zr6 nodes to form ZrxTi1-xO2, which resulted in extensive particle shrinkage. It is conceivable that thermal treatment of NU-1000 materials filled with silica or titania precursors in nitrogen rather than in air could result in the formation of carbon–silica or carbon– titania composites with oxozirconium clusters. This possibility was confirmed for SiO2@NU-1000. 29 ACS Paragon Plus Environment


XRD data of samples pyrolyzed at 600 or 800 ˚C in nitrogen showed that cluster aggregation was prevented (Figure S21). The presence of carbon was confirmed by elemental analysis, Raman spectroscopy (Figure S22) and the black appearance of these samples. However, in the case of TiO2@NU-1000 prepared from TDMAT or TEOT precursors, cluster aggregation was prevented only up to a pyrolysis temperature of 600 ˚C. At 800 ˚C, the carbon phase could not keep clusters from aggregating in these systems, resulting in the formation of a more extended zirconia phase. Interestingly, the presence of the carbon phase in the pyrolyzed samples prevented the formation of ZrTiO4.


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3.6 Cluster Accessibility Upon dehydration of NU-1000, the water molecules bonded to Zr6 clusters are removed, which enhances the Lewis acidity of the clusters.22 This process leads to potential catalytic activity of the materials for acid-catalyzed reactions. Similar to the dehydration of NU-1000, linker removal at high temperatures will also remove the water molecules associated with the Zr6 clusters, affording Lewis acid catalysts that are stable at high temperatures. In order to test cluster accessibility in the nanocast materials via the Lewis acidity of the Zr6 clusters, we used pyridine adsorption experiments. Pyridine is a commonly used probe molecule for the detection of acidic sites on metal oxide surfaces. Due to the lone electron pair on the nitrogen atom, pyridine can either coordinate to Lewis acidic sites or be protonated by Brønsted acidic sites. As a result, the vibrational modes of pyridine are altered, causing shifts in the vibrational frequencies observed in the FT-IR spectrum.78 We previously reported that for Zr6@SiO2, upon pyridine adsorption, the ν8a and ν19b modes of the νCCN vibrations are observed at 1607 cm-1 and 1445 cm-1, respectively, associated with vibrations of pyridine bonded to a Lewis-acid site (Figure 10).11 Even though Zr6 clusters no longer exist in the TiO2 nanocast samples, the Lewis acidic mixed Ti/Zr sites are also of interest as catalytic centers.79-81 Pyridine sorption experiments on TDMAT_Zr6@TiO2 yielded similar observations as for Zr6@SiO2. After pyridine adsorption, several characteristic vibrational peaks of Lewis-acid site bonded pyridine appeared between 1750 cm-1 and 1400 cm-1, including the ν8a and ν19b peaks at 1604 cm-1 and 1445 cm-1. In contrast, the pyridine adsorption spectrum of TEOT2_Zr6@TiO2 resembled more closely that of pyridine on bulk ZrTiO4 (Figure S23), consistent with the observation of ZrTiO4 in the Raman spectrum of TEOT2_Zr6@TiO2. Although in the case of Zr6@SiO2, the spectral patterns are distinct from that of pyridine adsorbed in bulk SiO2, features in the FT-IR spectrum of Zr6@TiO2 exposed to pyridine coincide with those in the spectrum of TiO2 exposed to pyridine. Therefore, the observed Lewis acid sites cannot be uniquely ascribed to Zr and likely contains contributions from Ti.



Figure 10. FT-IR spectra of Zr6@SiO2 and TDMAT_Zr6@TiO2 after adsorption of pyridine. Corresponding spectra for SiO2 and TiO2 after pyridine adsorption are also shown.

For carbon nanocasting, the FT-IR spectrum of Zr6@C that had been exposed to pyridine shows a pronounced peak at 1436 cm-1, associated with the vibration of the free pyridine ring (Figure S24), but no evidence for pyridine bound to Lewis acidic sites offered by the Zr6 clusters. It appears that the clusters are shielded by the carbon network. Although they are accessible to small Ar atoms, access is blocked to larger pyridine molecules in the as-made material. When ammonia was used instead as a smaller probe molecule in temperature-programmed desorption (TPD) experiments, a TPD signal associated with Lewis acidic sites was observed (Figure S25), confirming accessibility of the clusters to ammonia. To expose the clusters and make them more accessible to pyridine, the Zr6@C sample (Zr6@C-48h) was heated in static air at 300 ˚C for 30 min to remove some of the carbon. For the resulting material, two IR peaks associated with the vibration of the pyridine molecules bonded to the Lewis acid sites could be observed, confirming that the post-treatment had made the Zr6 clusters more accessible to pyridine by introducing defects into the carbon skeleton (Figure 11). The Raman spectrum of the post-treated sample shows an 32 ACS Paragon Plus Environment

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increase in the relative intensity ratio of the D-band to the G-band (Figure S26), resulting from partial oxidation of the carbon network. The rod-like particle morphology was preserved after this treatment to expose the Zr6 clusters, but average particle dimensions were slightly smaller (Figure S27). In order to visually determine the presence of Zr6 clusters, TEM images of microtomed Zr6@C-48h, Zr6@C-48h heated at 300 ˚C for 30 min, and Zr6@SiO2 samples were obtained (Figure S28). No extended aggregation of clusters was noticed. However, it is difficult to image individual clusters for these materials.

Figure 11. FT-IR spectra of the Zr6@C-48h sample heated in air at 300 ˚C for 30 min before (black) and after exposure to pyridine (red) and a difference spectrum of the two spectra in (a) (blue).

4. CONCLUSIONS We developed methods for nanocasting NU-1000 with SiO2, TiO2, and carbon. Using appropriate infiltration methods, nanocast porous materials with moderate BET areas could be prepared for all of these systems. Zr6 clusters appear to remain isolated after nanocasting in SiO2 and carbon samples, making them suitable for the design of single-site catalysts. Zr6@SiO2 remains stable up to 600 ˚C in air and Zr6@C samples remain stable at 800 ˚C in N2 before excessive zirconia agglomeration occurs, while TiO2@NU-1000 transformed to ZrxTi1-xO2 during calcination. In all cases, the processes were pseudomorphic, that is the shape of the original NU-



1000 crystals was maintained throughout, although shrinkage occurred after various processing stages. Several differences were identified among the materials and the corresponding nanocasting processes. During the infiltration processes, the spatial distribution of precursors in micro- and mesopores of the NU-1000 differs, resulting from the distinct molecular sizes and interactions of the precursors examined here. TMOS molecules enter mainly the larger mesopores; FA molecules occupy both the mesopore and micropore channels and the pockets between Zr6 clusters; the organotitanium molecules investigated here, because of their stronger interactions with the Zr6 clusters prefer infiltrating the pockets instead of the pores. These stronger interactions lead to significant shrinkage of the TiO2@NU-1000 particles after infiltration, which is amplified after calcination and formation of the mixed ZrxTi1-xO2 phase. On the basis of these studies, we identified several requirements for successful nanocasting in MOF materials, in particular NU-1000. First of all, one needs to find an efficient way to infiltrate the pores of MOFs with the precursor. In order to fulfill this requirement, various precursors containing the target elements can be considered, but molecule size must obviously be compatible with pore accessibility. Moreover, it is important to consider the reactivity of the MOF components with scaffold precursors. An appropriate infiltration method must be applied to avoid damaging MOF structures during the infiltration. During the subsequent condensation and linker removal processes, it is crucial to maintain the accessibility and isolation of the node clusters if they are needed in the product application. Therefore, the amount of precursor infiltrated into the pores of MOF materials becomes important. Additionally, the precursors should not have strong interactions with the node clusters, otherwise the precursors and the clusters may react with each other during the infiltration or linker removal procedure. These nanocasting methods can, in principle, also be adapted to other MOF materials with sufficiently large pores. Different catalytic properties may be achieved by nanocasting MOFs containing different metals as part of the nodes. In addition, by varying the nanocasting conditions (e.g. infiltration time), different precursor loadings can be obtained, which may allow for tuning the properties of the catalysts by adjusting the ratios of components in the nanocast products. ASSOCIATED CONTENT


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Supporting Information Synthetic methods, detailed characterization of the titania and carbon nanocast materials. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author [email protected] Notes The authors declare the following competing financial interest(s): J.T.H. and O.K.F. have a financial interest in the start-up company NuMat Technologies, which is seeking to commercialize metal-organic frameworks. ACKNOWLEDGMENTS This research was supported by the U.S. Department of Energy, Office of Basic Energy Sciences, Division of Chemical Sciences under Award DE-SC-0012702, except for the parts listed below. Parts of this work were carried out in the University of Minnesota Characterization Facility, which receives partial support from the NSF through the MRSEC, ERC, MRI, and NNIN programs. Work done at Argonne National Laboratory was performed using the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. A.E.P.P. acknowledges a Beatriu de Pinós fellowship (BP-DGR 2014) from the Ministry of Economy and Knowledge (Catalan Government). We thank Professor Paul J. Dauenhauer, Dr. Omar A. Abdelrahman, and Ms. Kristeen E. Joseph for carrying out TPD measurements.



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