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Nucleation of FAU and LTA Zeolites from Heterogeneous Aluminosilicate Precursors Matthew D. Oleksiak, Jennifer A Soltis, Marlon T Conato, R. Lee Penn, and Jeffrey D Rimer Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b01000 • Publication Date (Web): 20 Jun 2016 Downloaded from http://pubs.acs.org on June 29, 2016

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Chemistry of Materials

Nucleation of FAU and LTA Zeolites from Heterogeneous Aluminosilicate Precursors Matthew D. Oleksiak 1, Jennifer A. Soltis 2,4, Marlon T. Conato 1,3, R. Lee Penn 2, and Jeffrey D. Rimer 1* 1

University of Houston, Department of Chemical and Biomolecular Engineering, Houston, TX 77204 USA

2

University of Minnesota, Department of Chemistry, Minneapolis, MN 55455 USA

3

University of the Philippines, Institute of Chemistry, Diliman, Quezon City 1101 Philippines

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Current Address: Pacific Northwest National Lab, Physical and Computational Sciences Directorate, Richland, WA 99354 USA

ABSTRACT: The nucleation of many natural, biogenic, and synthetic crystals involves the initial formation of metastable precursors that provides a kinetic pathway for an amorphous-to-crystalline transformation. This nonclassical mechanism is believed to be the dominant crystallization pathway for microporous zeolites. Despite significant research on zeolite growth mechanisms, molecular level details regarding the assembly, physicochemical properties, and structural evolution of amorphous (alumino)silicate precursors remain elusive. Here we use a combination of diffraction, scattering, and microscopy techniques to characterize the amorphous precursors that assemble and evolve during the synthesis of zeolites FAU and LTA – two materials that are widely used in commercial applications such as catalysis, adsorption, separations, and ion-exchange. Nucleation occurs by a two-step mechanism involving the initial formation of aggregates that serve as heterogeneous sites for nucleation. Using colloidal silica as a reagent, we observe that precursors are comprised of heterogeneous silica and alumina domains due in part to the negligible dissolution of silica during room temperature aging. This indicates substantial Si-O-Si bond breakage must occur during hydrothermal treatment with concomitant exchange of soluble alumina species to achieve a final crystalline product with Si/Al ratio = 1.0 – 2.5. All syntheses were performed with molar compositions of Si/Al ≥ 2.0, which favors the formation of FAU; however, we observe that certain growth conditions are capable of creating a “false” environment (i.e., Al-rich regions) that favors LTA nucleation, followed by intercrystalline transformation to FAU. Time-resolved ex situ transmission electron microscopy of extracted solids during zeolite crystallization indicates that nucleation occurs on the exterior surface of precursors. This observation is consistent with our proposed hypothesis that posits exterior surfaces are more energetically favorable sites for nucleation compared to the particle interior on the basis of confinement effects. Given that numerous zeolite syntheses involve the initial formation of metastable precursors with heterogeneous composition, the pathway for nucleation proposed in this study may prove to be generalizable to other zeolite structures and related materials.

INTRODUCTION There is growing evidence that the nucleation pathways of many geological, biological, and synthetic crystals deviate from classical nucleation theory (CNT). The basic principle of CNT posits that nucleation occurs via the assembly of monomers (i.e., ions or molecules) into a particle of critical radius.1 This theory was originally challenged in studies of lysozyme crystallization that revealed the formation of protein clusters as precursors to nucleation.2 It was shown that the formation of dense liquid droplets of protein leads to a higher localized concentration of solute, which presumably lowers the energetic barrier(s) for nucleation relative to that predicted by CNT. This socalled “two-step" nucleation mechanism is believed to occur or has been directly observed for organic molecules, colloids, polymers, and biominerals.3-7 Precursors formed in these systems tend to be metastable particles that evolve in size and/or structure

via pathways that involve multiple (and potentially indistinguishable) energetic barriers. Only a fraction of precursors become nuclei, while the remainder either dissolve or remain in the synthesis mixture beyond the induction period. In many cases, the remaining precursors serve as growth units via a process generally referred to as crystallization by particle attachment (or CPA).8 Examples include (but are not limited to) the crystallization of metal oxides,9-12 chalcogenides13, calcium minerals,14 and zeolites.15 Zeolites are an exemplary class of materials that form via a combination of nonclassical and classical growth mechanisms. These aluminosilicate microporous materials exhibit tunable acidity, pore size/topology, and shape-selectivity for applications spanning catalysis,16,17 photonics,18,19 drug delivery,20 separations,21 and ion exchange.22 Zeolite synthesis is typically carried out in one of three environments: clear solutions, dispersed low density sols, or separated high density sols (viscous gels).23 Recent studies by Xiao et al.24-26 demonstrated

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for select cases that zeolite synthesis can be carried out with minimal water (solid phase) in what is referred to as a solvent-free process. Clear solutions are ideal for mechanistic studies because they are amenable to characterization by scattering or microscopy techniques; however, clear solutions are rarely employed in commercial synthesis due to their low zeolite yield and high cost of reagents. Two reagents commonly employed in zeolite synthesis are colloidal and fumed silica, which produce either turbid sols or viscous gels that pose challenges for in situ characterization. The term “sol gel” is often invoked to describe zeolite synthesis mixtures; however, the physicochemical properties and microstructure of the gel state are not well understood. Gel is a nebulous term that applies to materials where the continuous and dispersed phases may be interchangeable. Indeed, zeolite precursors are comprised of colloidal aggregates and/or a network of (alumino)silicate polymers in addition to entrained solution, which is believed to be an integral component of precursors.27 The spatial distribution of silicon and aluminum in the gel can vary significantly, often leading to a large disparity in the Si/Al ratio (SAR) of the bulk amorphous phase relative to the final zeolite product. As such, the amorphous-tocrystalline transformation of precursors involves an exchange of silica and/or alumina species between the solid and solution. There are many divergent hypotheses regarding the role of the solid phase in zeolite crystallization.28,29 Flanigen and Breck30 were among the first to propose that crystallization occurs in the solid phase through either solid-state transformation or exchange of silica/alumina with the solution. These processes are hypothesized to involve the breakage and formation of Si-O-Si and Si-O-Al bonds.31 The concept of “gel transformation” has been proposed for zeolites such as ZSM-5 (MFI),32-34 zeolite A (LTA),35 and zeolite X/Y (FAU).36,37 For LTA, it was suggested that nucleation occurs within the interior of the gel, followed by crystal growth (from the gel interior to its exterior) by consuming nutrient within the gel particle.38 Alternative mechanisms include reverse crystallization wherein heterogeneous nucleation occurs on the exterior surface of the precursor and the crystal grows by consuming the amorphous material (from the precursor exterior to its interior). The latter mechanism has been proposed for zeolites such as FAU, analcime (ANA),39 sodalite (SOD),40 and more recently EMT.41 Two commercial zeolites used for catalytic cracking and adsorption are FAU16 and LTA,42 respectively. FAU is a large-pore zeolite with 12-membered-ring (MR) pores and a SAR that ranges from 1.0 to 2.5. LTA is a small-pore zeolite with 8-MR pores and a SAR ≈ 1.0. As shown in Scheme 1, both zeolite crystal structures share a common composite building unit (CBU), the sodalite cage (sod). In FAU, sod cages are

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connected by a double-6-MR, while in LTA the sod cages are connected by a double-4-MR. Herein, we systematically examine the effect of synthesis composition and conditions on LTA and FAU nucleation. Using a combination of microscopy and diffraction techniques, we show that precursors (the amorphous bulk phase) are distinctly heterogeneous at the microscopic level. Our findings reveal that the amorphous-to-crystalline transformation involves a twostep nucleation pathway, beginning with the formation of precursors. We postulate that confinement effects within the interior of precursors engender large energetic barriers that favor nucleation on the exterior surface of amorphous precursor particles.

Scheme 1. Illustrations of LTA and FAU crystal structures and their corresponding CMUs: double-4-membered-ring (d4R, LTA), double-6-membered-ring (d6R, FAU), and the sodalite cage (sod, LTA and FAU).

EXPERIMENTAL Zeolite growth mixtures were prepared with molar compositions x SiO2: y Al2O3: 10 NaOH: 173 H2O. All reagents were used as received without purification. First, sodium aluminate (technical grade, Alfa Aesar) and sodium hydroxide (98%, Sigma Aldrich) were added to deionized (DI) water that was purified with an Aqua Solutions Type I RODI filtration system (18.2 MΩ). The resulting mixture was stirred until well-mixed followed by the addition of one of the following silica sources: LUDOX AS-40 colloidal silica (40%, Sigma Aldrich), LUDOX SM-30 (30%, Sigma Aldrich), or tetraethyl orthosilicate (TEOS, 98%, Sigma Aldrich). AS-40 was used as the nominal silica source for zeolite synthesis unless otherwise stated. The dispersion (low density sol) was stirred at room temperature for 24 h in a polypropylene bottle, which was then heated under static conditions in a ThermoFisher Precision oven at 65°C. Bottles were removed at periodic times between 2 and 168 h. Upon removal from the oven, bottles were cooled to ca. 25°C in a water bath. The solid sample was

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recovered by three cycles of centrifugation and washing with DI water. Centrifugation was performed using a Beckman Coulter Avanti J-E at 5°C and 23,220 G for 45 min. The product was dried in air at room temperature overnight prior to analysis. A kinetic phase diagram was constructed for synthesis mixtures at a constant molar ratio of 10 NaOH: 173 H2O. The quantities of silica and alumina used in each synthesis were varied to adjust the silicon-to-aluminum ratio of the growth mixture, which we refer to herein as SAR(liq), and the silicon-to-hydroxide ratio, Si/OH. The crystal structure of each zeolite was determined by powder X-ray diffraction (XRD) using a Siemens D5000 X-ray Diffractometer equipped with a Cu source and Ni filter. The crystal phase(s) were indexed with simulated patterns obtained from the International Zeolite Structural Database.43 Samples for dynamic light scattering (DLS) were prepared in 10 mL amounts and placed in 12-mL centrifuge tubes. These tubes were regulated at the appropriate set point temperature using a Julabo ED water bath. Samples were removed from the bath at specified times and quenched in an ice bath for 30 sec. Aliquots of 10 drops (ca. 0.5 mL) of each growth dispersion were diluted into 12 mL of DI water and then filtered with a 0.45 µm membrane to achieve adequate counts per second (cps) for DLS measurements (e.g. 20 to 150 kcps). Light scattering measurements were conducted using a Brookhaven Instruments BI-200SM machine equipped with a TurboCorr Digital Correlator, a red HeNe laser diode (35mW, 637 nm), and a decahydronapthalene index-matching bath. The liquid sample cell was regulated at 25°C with a Polyscience digital temperature controller. DLS autocorrelation functions were collected in a 2 min period assuming a refractive index of pure water and a viscosity that was measured at room temperature using a Cannon Ubbelohde Viscometer (0.4 – 2.0 mm2/s). Specimens for transmission electron microscopy (TEM) were prepared by adding ca. 1 mL of growth dispersion to 30 mL of DI water and drop-casting onto 200 mesh copper TEM grids coated with a holey carbon film (SPI Supplies). Samples were dialyzed using previously reported methods.44 High resolution (HRTEM) and dark field imaging were performed using a FEI Tecnai G2 F30 field-emission gun microscope equipped with a cryo stage and operated at 300 keV under low-dose conditions. Images were collected with a Gatan UltraScan charge-capture device (CCD) camera using Digital Micrograph v3. All specimens characterized using this microscope were imaged at cryogenic temperatures to prevent damage from the electron beam. The conventionally-prepared TEM grid was placed in a cryo TEM holder at room temperature and inserted into the microscope. The dewar on the holder was then filled with liquid nitrogen and the temperature of the holder monitored until the tip reached

an equilibrium temperature of ca. 95 K before imaging. Low-resolution imaging and energy dispersive X-ray spectroscopy (EDS) measurements were performed using a FEI Tecnai T12 microscope equipped with a LaB6 filament that was operated at 120 kV. Images were collected with a Gatan MSC794 CCD camera. EDS measurements were taken using an Oxford Inca EDS system equipped with an ultrathin window Si(Li) detector. Energy filtered transmission electron microscopy (EFTEM) data was obtained using a FEI Tecnai G2 F20 ST FE-TEM instrument at the Texas A&M University Microscopy and Imaging Center. RESULTS AND DISCUSSION Time-resolved analysis of zeolite crystallization. Zeolite nucleation and growth occur by complex processes8,28 that are generally not well understood. Numerous synthesis parameters influence both the crystalline phase (framework type) and physicochemical properties of the final product.45,46 The composition of a zeolite synthesis mixture is typically reported as molar ratios of SiO2, Al2O3, base (e.g., alkali hydroxide), H2O, and any other chemicals required for synthesis, such as organic structure-directing agents. We previously reported the use of kinetic phase diagrams to visualize (or map) regions of phase purity where the composition of the growth dispersion is plotted on a ternary diagram with axes of Si, Al, and MOH mole fractions (where M+ = alkali metal; Na+ for this study). All other parameters are held constant.47 In a previous study we showed that growth dispersions with molar ratio Si/OH < 1 heated for 168 h at 65°C result in the formation of FAU in Sirich mixtures (Figure 1, blue region), LTA in Al-rich mixtures (Figure 1, red region), and a binary mixture of each phase in a small range of compositions (overlapping regions). Here, we apply the same methodology by varying SiO2 and Al2O3 content in Sirich media to examine nine compositions with a fixed molar ratio of 10 NaOH: 173 H2O. All growth dispersions were prepared using sodium aluminate as the aluminum source. Here, we focus on LUDOX AS40 (25nm colloidal silica) as the silica source; however, later in this study we will examine two alternative silica sources: LUDOX SM-30 (8nm colloidal silica) and tetraethylorthosilicate (TEOS). Our findings confirm prior reports45,48 that the Si/Al ratio of the synthesis mixture and alkalinity (i.e., Si/OH ratio) impact nucleation and the initial crystalline phase. To test the effect of synthesis mixture composition, we prepared samples C1 – C9 with Si/Al ratios of either 2 or 5 and Si/OH ratios spanning 0.10 to 0.75, as listed in Table 1. Solids from each synthesis were extracted at various times between 2 and 168 h, and the crystal structure(s) was determined by powder XRD. The majority of growth dispersions initially produced FAU, which is consistent with the kinetic phase diagram in Figure 1. The onset of crystallization, as identified by

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hydroxide concentration. Growth dispersions at high alkalinity (Si/OH = 0.1) began to crystallize at shorter time periods (i.e., 4 h for C5 and C9), while nucleation of less alkaline synthesis mixtures (compositions C2, C6, and C7) required at least 12 h for crystallization to begin. The lone exception is composition C1, which exhibits an induction period of 4 h, similar to growth dispersions at high alkalinity. Interestingly, the first zeolite phase to form in C1 dispersions is LTA. We previously reported47 that composition C3 within the Sirich region of the diagram also results in the initial formation of LTA, followed by an intercrystalline transformation to FAU at longer heating time. Comparison of the XRD powder patterns (Figure 2) of solids that were extracted from C1 dispersions at periodic times (2 – 168 h) revealed an intercrystalline transformation from LTA to FAU. Solids extracted from growth dispersions during the first 2 h were amorphous. Within 3 h we observed the formation of LTA crystals, and after 8 h of heating, XRD revealed the presence of both LTA and FAU until ca. 24 h. Additional heating time led to the complete transformation of LTA to pure FAU. Indeed, compositions C1 – C9 all yielded pure FAU within 168 h. Composition C3 exhibited a similar LTA-to-FAU transformation, while the first crystalline peaks observed in XRD patterns of composition C5 revealed a binary mixture of LTA and FAU. When comparing all compositions leading to intercrystalline transformation (C1, C3, and C5), the initial formation of LTA does not seemingly occur in a specific region of Si/OH ratios within the phase diagram; rather, the three compositions have widely varying Si/OH ratios. This phenomenon was only observed at Si/Al = 2, in close proximity to the boundary separating the FAU and LTA phases. The hypothesis for the LTA-to-FAU phase

Figure 1. Kinetic ternary phase diagram with axes in mole fractions of Si, Al, and NaOH. Growth dispersions were prepared with an initial molar ratio of x Si: y Al: 10 NaOH: 173 H2O. Compositions were selected at Si/Al = 2 and 5 with Si/OH = 0.10, 0.25, 0.50, and 0.75. Symbols correspond to the first zeolite phase detected during hydrothermal treatment at 65°C (see Table 1). The shaded regions represent areas of FAU (blue) and LTA (red) formation after 168 h of heating (as reported by Maldonado et al.).47 The approximate times associated with the first observation of Bragg peaks in powder XRD samples for each composition are: (C1) 3 h, (C2) 12 h, (C3) 6 h, (C4) 6 h, (C5) 4 h, (C6) 12 h, (C7) 12 h, (C8) 6 h, and (C9) 4 h.

the first appearance of Bragg peaks in XRD patterns, occurred at different times depending on the composition. The symbols in Figure 1 refer to the first zeolite phase detected by XRD. We observed that the induction period generally increases with reduced

Table 1. Growth dispersion compositions and time-resolved phases of extracted solids.

SiO2 C1 C2 C3 C4 C5 C6 C7 C8 C9

7.5 5.0 3.9 2.5 1.0 7.5 5.0 2.5 1.0

Molar Composition Al2O3 NaOH H2O 1.88 1.25 0.90 0.63 0.25 0.75 0.50 0.25 0.10

10 10 10 10 10 10 10 10 10

173 173 173 173 173 173 173 173 173

Products with Heating Time (h) at 65°C ‡ 2

4

6

12

24

168

am am am am† am† am† am† am† am

LTA am† Am am† FAU (LTA) am† am† am† FAU

LTA am† LTA am† (FAU) FAU (LTA) am† am† FAU FAU

FAU (LTA) FAU FAU(LTA) FAU FAU(LTA) am† (FAU) FAU FAU FAU

FAU (LTA) FAU FAU FAU FAU(LTA) FAU FAU FAU FAU

FAU FAU FAU FAU FAU FAU FAU FAU FAU

‡

am = amorphous, LTA = Linde type A (zeolite A), FAU = faujasite (zeolite X or Y); † Sodium sesquicarbonate peaks detected in powder XRD (crystalline impurity); Structures in parentheses indicate the minor crystalline phase; Composition C3 is identical to the one previously reported by Rimer and coworkers.47

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transformation, which was initially postulated in our earlier work,47 will be explored in more detail herein. Heterogeneous growth dispersions. Here we use TEM, EFTEM, and EDS to examine the evolution of precursor size, shape, and spatial distribution of elements (Si and Al) during crystallization. The initial growth dispersion for zeolite synthesis was comprised of silica particles (i.e., LUDOX AS-40 colloidal silica source) suspended in an Al-rich aqueous solution. Focusing on composition C3, we examined a dispersion that was aged for 48 h at room temperature using EFTEM to determine the spatial distribution of Al and Si. Solids extracted from the aged C3 dispersion revealed the presence of monodisperse spherical particles with an average diameter nearly identical to that of the colloidal silica source (refer to Figure 4A).

Figure 2. Powder XRD patterns of the solids extracted from a C1 growth dispersion after heating at 65°C for the following times: 2, 3, 8, 24, and 168 h. After 3 h of heating, LTA is the first structure to nucleate, and within 8 h there is a binary mixture of LTA and FAU. After 24 h of heating, LTA is a minor phase, and by 168 h the LTA-to-FAU intercrystalline transformation is complete. The driving force for this transformation is not well understood. The thermodynamic stability of zeolite structures tends to increase with higher framework density and increased Al content.49 Here there are two competing effects: FAU is denser while LTA is more Al-rich.

EFTEM images precursors indicate that the majority of Al (Figure 3A, red) is present in solution surrounding the colloidal silica particles. Images of Si for the same sample (Figure 3B, green) show the majority of silica is contained within the colloidal particles, with very little Si observed in the surrounding regions of the sample. An EDS line scan across a single particle (Figure 3C) confirms the presence of core-shell particles where Al is observed at the particle exterior (shell) and Si is observed within the particle interior (core). These measurements are evidence of the precursor’s core-shell structure, although the coverage and average thickness

of Al on silica surfaces is unknown (i.e., the depiction in Figure 3C is merely illustrative). For instance, Iler50 reports that the introduction of alumina to an initially saturated amorphous silica sol leads to coprecipitation of alumina and soluble silica on solid colloidal silicate surfaces. This seems to suggest that the core-shell illustration should contain an aluminosilicate shell wherein the boundary separating the shell from the core is less defined than the one depicted in Figure 3C schematic. This is consistent with the EDS line scan that reveals the presence of silica within the alumina shell. Transmission electron micrographs of samples extracted from a C3 dispersion after 48 h aging at room temperature show that colloidal silica dissolution does not lead to an apparent reduction in particle diameter (Figure 4A). Notably, precursor particles are similar in size (ca. 27 nm, see Figure S13) to the colloidal silica source (ca. 25 nm). The addition of alumina resulted in the immediate aggregation of silica particles, which can be inferred from the visual clear-to-opaque transformation of the growth dispersion (refer to Figure 6). The presence of aggregates was quantitatively verified using DLS to detect an appreciable increase in the average hydrodynamic diameter and polydispersity during aging. The appearance of aggregates in TEM images is partially attributed to sample preparation (i.e., drying); however, the spherical domains of each

Figure 3. TEM analysis of samples prepared with composition C3 reveals a precursor structure comprised of a Si-rich core and an Al-rich shell. EFTEM images were captured to show the aluminum (A, red) and silicon (B, green) content in particles extracted from a dispersion prepared at room temperature without heating. Elemental mapping reveals that particles are predominantly SiO2 and exhibit sizes comparable to the original colloidal silica reagent. There is a thin layer of Al2O3 on the exterior surface. (C) EDS cross-sectional scan of a single particle shows the increase in Al content (dashed red line) on the exterior of the particle, and an interior comprised solely of silica (solid green line). The corresponding STEM image is provided in Figure S20.

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colloidal silica particle within the aggregate are clearly visible in electron micrographs, which indicates that particles do not appreciably fuse together or coalesce. Once the dispersion is heated for a period of time (e.g., 2.5 h), the particles continue to aggregate (Figure 4B) and the disappearance of distinct 25-nm spherical domains suggests that particles fuse and/or ripen to generate denser clusters that are less fractal than samples aged at room temperature. EFTEM analysis of samples heated for 0.5, 1.5, and 2.5 h reveal that there is still a significant heterogeneous distribution of Si and Al within the aggregates (Figure S11), which seems to suggest that the core-shell structure is preserved during the early stages of hydrothermal treatment. The addition of alumina to synthesis mixtures clearly facilitates colloidal silica aggregation, although the exact mechanism for this process is unclear from our studies. Notably, we cannot differentiate if the driving force for aggregation is interparticle interactions between core-shells or the bridging of two or more silica particles through covalently bound alumina shells (i.e., similar to polymer-induced flocculation).51 Given the high alkalinity of growth mixtures employed in this study (i.e., pH > 13), silica dissolution is expected to occur; however, colloidal silica maintains a constant diameter during room temperature aging, indicating very little, if any, particle dissolution. This is consistent with Iler’s work showing that alumina deposition on silica particles effectively inhibits the rate of silica dissolution.50 This phenomenon can also be observed in other zeolite systems where undissolved silicates selfassemble (or dynamically evolve) into so-called wormlike particles,52-57 which are comprised of high-Si domains that appear to be aggregates of colloidal silica.58 Even in the absence of alumina, we observed only marginal dissolution of colloidal silica. Notably, we

Figure 4. TEM images of solids extracted from a dispersion with molar composition 3.6 SiO2: 0.9 Al2O3: 10 NaOH: 173 H2O. (A) Particles after room temperature aging have a spherical shape, a monodisperse size distribution, and an average diameter of 27 ± 3 nm, similar to that of the colloidal silica source (25 nm). (B) Particles after 2.5 h of heating at 65°C reveal a more polydisperse distribution of particle size and denser aggregates.

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monitored by DLS the temporal reduction in the hydrodynamic diameter of silica particles in an Al-free synthesis mixture (Figure 5). Room temperature aging resulted in a slow rate of dissolution (ca. 0.02 nm/min), while an increase in temperature resulted in a concomitant increase in the rate of dissolution. The corresponding Arrhenius plot (Figure S12) yields an apparent activation energy of dissolution, EA = 87.2 kJ/mol, that is similar to previously reported values for amorphous silica.59,60 Dissolution assays revealed that colloidal silica particles reach a lower limit beyond which particle size remains constant with heating time, thus indicating the dispersion has reached thermodynamic equilibrium (i.e., solubility of amorphous silica). This general trend is depicted in Figure 5 for the 45°C data (dashed line) where dissolution ceases once the hydrodynamic diameter of silica reaches ca. 70% of its original size. This indicates that, irrespective of the alumina concentration, the synthesis mixture cannot attain a homogeneous composition of soluble silica. The majority of the silica is “sequestered” within the solid state, rendering this fraction of silica initially unavailable for zeolite growth. The amorphous-tocrystalline transformation of precursors must therefore occur through one of two possible routes: (i) temporal dissolution of the solid phase to replenish soluble species that are consumed during zeolite crystallization; or (ii) solid-state rearrangements that originate from crystal nucleation on the surface of core-shell particles

Figure 5. Ex situ DLS measurements of an Al-free synthesis mixture with molar composition 3.6 SiO2: 10 NaOH: 173 H2O that was heated at different temperatures: 25.0, 35.0, 40.0, 42.5, and 45.0°C. The high alkalinity of the dispersion (pH = 13.6) results in the dissolution of colloidal silica particles. Here we plot the temporal change in the relative hydrodynamic diameter of silica particles, DH/DH(t=0), where DH(t=0) is the initial size measured at t = 0 min. Solid lines are linear regression revealing that the kinetic rate of dissolution (slope of the line) increases with increasing temperature. Measurements of DH account for the viscosity of DLS samples, η = 0.924 cP. The dashed line corresponds to the plateau in particle diameter. The apparent activation energy is similar to reported values for silica dissolution (see Figure S12).

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(i.e., solid-liquid interface). It should be noted that the latter mechanism cannot occur without solutionmediated Si–O–Si bond breakage to facilitate the transition from amorphous silica (Si/Al = ∞) to zeolite crystals (Si/Al = 1 – 3). Despite crystallization occurring in an overall Si-rich suspension, most of the silica is retained within the original colloidal silica particles. As such, nucleation on the exterior surfaces of core-shell precursors occurs in a local environment enriched with Al, taking into account both the alumina shell and the Al-rich solution. This inhomogeneity likely explains the observed LTA-toFAU transformation. A similar scenario may not occur in situations where an appreciable silica concentration is present near the site of nucleation, i.e., if a sufficient quantity of silica is able to dissolve from the core or if nucleation occurs at Al-depleted regions on the shell surface. All compositions in Table 1 yield FAU after sufficient heating time; therefore, the initial formation of LTA is seemingly attributed to localized Al-rich zones within the synthesis mixture. In order to provide additional evidence for this hypothesis, we performed similar studies using alternative silica sources to modify the initial state of precursors in the growth dispersions, as described in the following section.

Figure 6. Photographs and idealized schematics of zeolite growth dispersions prepared with two different silica sources: LUDOX AS-40 (left) and TEOS (right). These images correspond to synthesis mixtures that have been stirred for 24 h at room temperature. Both samples contain distinct heterogeneous domains comprised of Si-rich and Al-rich particles and/or solution. The majority of silica is inaccessible throughout much of the induction period owing to either undissolved cores (i.e., LUDOX samples) or immiscible liquid layers (i.e., non-hydrolyzed TEOS).

Alternative silica sources. Here we switched from LUDOX AS-40 to TEOS as the silica source. The latter generates a two-phase solution comprised of an initially immiscible Si-rich TEOS layer (Figure 6) that progressively hydrolyzes to release silica into an Al-rich aqueous solution. This scenario is analogous in some respects to synthesis mixtures prepared with colloidal silica wherein the dispersion of core-shell particles leads to spatially segregated silica and alumina regions. Our

findings reveal that a C3 composition prepared with TEOS exhibits a LTA-to-FAU intercrystalline transformation, identical to the LUDOX AS-40 dispersion shown in Figure 2. As illustrated in Figure 6, colloidal silica and TEOS synthesis mixtures are subdivided into Si- and Al-rich domains. In both scenarios, silica is slowly released into the Al-rich solution. The initial formation of LTA is seemingly dependent upon the heterogeneity of the growth dispersion wherein spatio-temporal variations in elemental composition can engender localized regions that promote the nucleation of the more metastable structure, LTA, in a pseudo Al-rich environment. The immiscible layer of TEOS persists during room temperature aging and throughout the induction period and early stage of crystallization (Figure S21) at elevated temperature. Slow hydrolysis of TEOS leads to a progressive accumulation of silica within the aqueous solution. Silica in the aqueous solution does not remain solubilized as monomers or small oligomers, but forms a bulk amorphous phase. TEM images of samples extracted from the aqueous phase after room temperature aging (Figure 7A) reveal aggregates containing 5 – 10 nm particles, analogous to the findings reported by Tsapatsis and Penn for silicalite-1.44,61,62 The aggregates appear to be loosely packed, suggesting that both solvent and dissolved (alumino)silicate species can diffuse into the interior of the aggregate. EFTEM of a single aggregate (Figure 7C – E) shows that the Si and Al content is more uniformly dispersed in the TEOS sample compared to the core-shell precursors derived from colloidal silica (see Figure 3); however, the resolution of EFTEM is not sufficient to discern the spatial distribution of elements within the 5-nm subdomains of the aggregates. It is feasible that alumina forms a shell surrounding the silica particles, but it is also possible that alumina condenses with silica to yield a more homogeneous bulk amorphous phase. Nevertheless, the majority of silica resides in the nonhydrolyzed TEOS layer while the aqueous solution remains Al-rich throughout much of the induction period (owing to the slow hydrolysis of TEOS). This presumably facilitates the nucleation of LTA crystals prior to the onset of intercrystalline transformation once the silica concentration has reached an appreciable level in the aqueous solution. The kinetics of zeolite crystallization in TEOS synthesis mixtures is markedly different than the equivalent mixture prepared with 25-nm colloidal silica (LUDOX AS-40 source). The induction time of the former is approximately three times longer. Moreover, the transformation from LTA to FAU is substantially slower for TEOS-containing growth dispersions. As shown in powder XRD patterns (Figure 7B), a LTA– FAU binary mixture is present in the product after 168 h of heating, whereas the AS-40 dispersion yields pure FAU within 24 h (see Figure S3).

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One question posed in this study was whether the LTA-to-FAU phase transformation would transpire if TEOS hydrolysis occurred more rapidly to produce a population of small siliceous particles. To investigate this further, we prepared a C3 composition using 8-nm colloidal silica (LUDOX SM-30), which is similar in size to the ca. 5-nm particles observed in TEOS dispersions. TEM images of solids extracted from C3 mixture prepared with SM-30 after aging at room temperature for 24 h reveal aggregates of small particles (ca. 10 – 15 nm, Figure S15). Electron micrographs show that the degree of aggregation in the SM-30 sample is much higher than that in AS-40. The first crystalline phase that nucleates in SM-30 samples is FAU (Figure S9). Despite the similarity in silica particle size between TEOS and SM-30 growth dispersions, there are two factors that likely contribute to the nucleation of FAU: (i) the total surface area of silica particles in SM-30 samples is larger, which depletes alumina from the aqueous solution during the formation of core-shell particles; and (ii) the aqueous solution for SM-30 samples contains a larger quantity of silica in the aqueous phase during the onset of heating. This hypothesis is consistent with the switch in phase

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behavior observed when replacing the 25-nm colloidal silica source, which has a specific surface area of ca. 135 m2/g, with the 8-nm source that has a specific surface area of 320 – 400 m2/g. A larger surface area increases the probability that silica particles have submonolayer coverage of alumina. Iler reports that complete coverage of alumina on silica surfaces is not required to suppress dissolution.50 For our system, crude estimates of alumina surface coverage on colloidal silica particles using 3.95 g/cm3 as an estimated Al2O3 density, and assuming 100% adsorption of all available alumina, indicate that there is sufficient quantity of alumina to fully coat the 25-nm silica particles, but not enough to form a complete coverage on the 8-nm particles. The study by Iler examined amorphous silica dissolution in undersaturated aqueous solutions at pH 8 and showed that 1.4 Al atoms per 1.0 nm2 surface area is sufficient to reduce silica solubility by an order of magnitude. A direct comparison of Iler’s threshold coverage to our system is not feasible due to the markedly different composition and alkalinity (pH 13.6) of zeolite growth dispersions; however, a corollary of Iler’s study is that sub-monolayer coverage of alumina is not uniquely indicative of silica dissolution.

Figure 7. Analysis of a C3 synthesis mixture prepared with TEOS. (A) TEM image of solids extracted after 48 h room temperature aging. The aqueous layer of the two-phase solution (bottom layer in Figure 6) was extracted for sample preparation. Inset: high magnification image of particles with sizes 5 – 10 nm (scale bar = 20 nm). (B) Powder XRD patterns of solids extracted from dispersions heated at 65°C for the following times: 12, 20, 24, and 168 h. The LTA-to-FAU intercrystalline transformation is similar to that reported by Maldonado et al.47 for a C3 composition prepared with LUDOX AS-40 as the silica source. (C) TEM image of another C3 sample aged for 48 h at room temperature and the corresponding EFTEM elemental mappings of (D) Si (green) and (E) Al (red).

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Putative mechanisms of zeolite nucleation. Prior studies have proposed different, but not necessarily mutually exclusive, hypotheses regarding the role of zeolite precursors as sites for heterogeneous nucleation. These theories are divided based on the spatial origin of nucleation occurring either within the interior35 of the gel phase or on the exterior surface63 of precursors. In 1999 there were two seminal papers published by Bein and Mintova35,63 presenting potential pathways for nucleation in the interior and on the surface of LTA and FAU precursors. Valtchev and coworkers64 expanded upon these theories by proposing that solvent entrained within the interior of FAU precursors can facilitate the dissolution or condensation of aluminosilicate species during zeolite nucleation. In literature, the semantics of referring to zeolite precursors as gels could give the impression that particles are a homogeneous mixture of silica and alumina, which overlooks the inherent complexity of the gel particle. As illustrated in Scheme 2, a hierarchical view from macroscopic to microscopic length scales (i.e., left to right in the schematic) provides an idealized depiction of precursors as agglomerates of heterogeneous silica and alumina domains. Precursor particles may appear to be a uniform gel at the macroscopic level, but our findings suggest that the sol gel is comprised of core-shell particles and entrained solvent. This gives rise to microscopic domains with significantly different spatial distribution of Si and Al compared to the crystalline product. The continual presence of precursors throughout zeolite crystallization is qualitatively consistent with the two-step mechanism of nucleation wherein the first step relates to the self-assembly of metastable (amorphous) aggregates that lower the energetic barrier(s) for nucleation. The large degree of inhomogeneity within precursors (as illustrated in Scheme 2) suggests that there must be substantial rearrangement of silica and alumina in order for the core-shell particles to transform into crystalline FAU or LTA, which exhibit more uniformly-distributed Si and Al within their framework structures. The second step of nucleation (i.e., the amorphous-to-crystalline transformation of precursors) constitutes a series of energetic barriers attributed to molecule dissolution, diffusion, rearrangement, and eventual reordering. One topic that is highly debated in the field of zeolite nucleation is whether these dynamic processes occur within the interior of precursors or on their exterior surfaces. To address this subject, we obtained time-resolved bright field and dark field TEM images of the solids extracted from a C3 synthesis mixture at periodic heating times after observing the onset of Bragg peaks in powder XRD patterns (Figure S2) in order to ensure the presence of partially-crystalline domains. A bright field image (Figure 8A) confirms that the sample is comprised of spheroidal precursor aggregates. Figure 8B is a dark field image of the highlighted area in

Scheme 2. Hypothesized (alumino)silicate precursors.

microstructure

of

Figure 8A where the bright spots correspond to nuclei. This is consistent with a selected area electron diffraction (SAED) pattern showing the presence of rings (see Figure S18). The presence of residual amorphous material is consistent with the broad peak at 2θ = 25° in the powder XRD pattern. The location of nuclei appears to be at the exterior surface of precursor particles. Expanding this analysis to composition C1, we again observe the onset of nucleation on the exterior surfaces of precursor aggregates (Figure 8C and D). High resolution TEM (HRTEM) analysis of individual particles in this sample (Figure 8E and F) reveals the presence of lattice fringes near the exterior of the particle, while many particle interiors of the partiallycrystalline sample did not contain fringes. For a subset of zeolite compositions explored in this TEM study, it appears as though nucleation originates on the exterior of precursors. A similar finding was recently reported by Valtchev and coworkers for zeolite EMT.65 EDS measurements reveal a temporal change in the SAR of precursors during hydrothermal treatment. The nucleation and growth of zeolite crystals from an initially heterogeneous growth dispersion invariably involves considerable silica bond breakage within coreshell particles to generate a more uniform distribution of silica and alumina in the final product. This is supported by comparing the composition of Si-rich amorphous precursors to that of LTA crystals (Si/Al = 1.0) or FAU crystals (Si/Al = 1.0 – 2.5). For instance, a C3 composition heated for 1 h contains aggregates of precursors with a spatially-averaged SAR of 2.2; however, EDS analysis at different locations within a single aggregate after 2.5 h of heating reveals a broad distribution of SAR ranging from 2 to 70 (see Figure S14 in the SI). EFTEM images of a 2.5 h sample indicate that silica and alumina domains are predominantly segregated (Figure S14). In general, the most aluminous regions of these samples correspond to the crystalline domains.

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Figure 8. TEM images of samples prepared shortly after the onset of nucleation. (A and B) Micrographs of a C3 composition heated at 65°C for 9 h in (A) bright field mode and (B) dark field mode. The latter was analyzed in the region outlined by the dashed box in A. The dark field image shows crystal nuclei (bright spots) on the exterior of precursors. (C – F) Micrographs of a C1 composition heated at 65°C for 3 h. The bright field image (C) and corresponding dark field image (D) focus on a large aggregate of precursors where nuclei appear to be located on the exterior of particles within the aggregate. (E) TEM image of a different particle within the same sample. (F) HRTEM image of a region within the dashed box in E showing lattice fringes (indicated by the black line).

In Figure 9A we present a bright field TEM image of solids extracted from a C1 composition that was heated for 3 h at 65°C. The aggregates in this image are subdivided into four regions for EDS area scanning, which provides a spatially-averaged Si/Al ratio. The HRTEM image in Figure 9B clearly shows the presence of crystalline and amorphous domains in agreement with the powder XRD pattern in Figure 2. In Figure 9C we provide the average atomic percentage of Si and Al in each area of analysis as well as the molar Si/Al ratio. Areas 1 and 2 assess the exterior and interior of a single aggregate, respectively. HRTEM indicates that the exterior region (area 1) is partially crystalline whereas the interior region (area 2) is completely amorphous. The partially crystalline region has Si/Al ≈ 1.4 and the amorphous region has Si/Al ≈ 1.7, whereas the entire aggregate (area 3) has an intermediate

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Figure 9. EDS analysis of particulates in a C1 composition heated at 65°C for 3 h. (A and B) TEM images showing (A) a representative image of the sample with four dashed circles marking the approximate areas of EDS analysis, and (B) a higher magnification image of area 1 showing the crystalline region in which lattice fringes are observed. Inset: HRTEM image showing lattice fringes in the direction indicated by the white line. (C) Results of EDS showing Si and Al atomic percentages (left axis) as well as the Si/Al ratio (right axis) for each area.

composition (Si/Al ≈ 1.6). The EDS scan of area 4 encompasses an entire aggregate, which appears to be largely amorphous with a Si/Al ratio that is comparable to that of area 2. Interestingly, precursors at this relatively late stage of crystallization are more homogeneous in composition with an Al content that is markedly closer to the zeolite product compared to the initial core-shell particles. The effect of confinement on zeolite nucleation. Here we discuss the feasibility of nucleation within the interior of precursors (i.e., occluded liquid pockets) by considering the potential energetic barriers imposed by confined geometries. The role of confinement in crystallization has been examined in the context of its effect on crystal polymorphism66 and habit.67 Significantly fewer studies examine the influence of confinement on nucleation (the majority of these studies focus on unrelated materials, such as polymers).68 One notable exception is the work by Veesler and coworkers69-71 who examined nucleation and growth of inorganic salts (e.g., NaCl) in microdroplets. Veesler observed that nucleation is virtually unaffected until liquid droplets reach picoliter (pL) volumes, corresponding to a droplet diameter of ca. 12.5 µm, where the confined geometry led to noticeable

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differences in crystallization. At these conditions, it was observed that the metastability limit (i.e., the minimum concentration required for nucleation) dramatically increased for NaCl nucleation. For instance, NaCl nucleation will not occur under these conditions unless the supersaturation ratio is increased by more than a factor of two.69 To the best of our knowledge, confinement effects in zeolite synthesis have not been reported, let alone considered; however, we anticipate that energetic barriers for nucleation exist in small liquid volumes occluded within zeolite precursors or the interstitial spaces of aggregates formed thereof, analogous to other inorganic crystals. For instance, consider the aggregate of precursors in Figure 9A with an approximate diameter of 1.5 µm. The total volume of this aggregate is 10–3 pL; and only a fraction of this space is occupied by solvent. The significant confinement within the aggregate likely increases the metastability limit for zeolite nucleation, which in turn would lead to a higher energetic barrier for nuclei to form in the confined environment of a single nanoparticle precursor and/or the interstitial space within precursor aggregates. Conversely, nucleation on the exterior surface of precursors, where confinement effects are nonexistent, would presumably have a lower energetic barrier and is more thermodynamically favorable. The exact metastability limit for zeolite nucleation is unknown, but we speculate that the supersaturation must be substantially higher in extremely-confined geometries relative to bulk solution for LTA and FAU nucleation. There are examples of zeolite crystallization in highly-confined porous templates, such as the work by Tsapatsis and coworkers who prepared threedimensionally ordered mesoporous (3Dom) zeolites in 10 – 40 nm diameter pores of carbon using a hierarchical templating scheme.72 In their procedure, the crystallization of (alumino)silicate precursors in narrow pores occurred in the presence of saturated steam. This created high supersaturation regions that were sufficient to promote the nucleation and growth of several zeolite framework types.73,74

CONCLUSIONS In summary, we have shown that sol gels in zeolite synthesis are comprised of heterogeneous silica and alumina domains. Comparison of three different silica sources suggests that the induction period and the initial crystal phase is highly dependent upon the degree of heterogeneity, or more specifically, the rate at which silica is released into Al-rich solutions. The formation of core-shell precursors and aggregates thereof creates a broad distribution of Si/Al ratios where the amorphousto-crystalline transformation requires substantial bond breakage to yield crystalline products with more uniformly-distributed Si and Al composition.

Here we hypothesize that confinement effects are relevant to zeolite nucleation. Future studies to confirm this hypothesis should recognize that different silica and alumina sources in zeolite synthesis can give rise to precursors of varying physicochemical properties (notably chemical heterogeneity and microstructure).75,76 This, in turn, can lead to potentially disparate findings among studies that use different synthesis conditions. In cases where precursors assemble with a more uniform distribution of silica and alumina than those in our study, the barriers for nucleation may be less stringent. Irrespective of the procedure selected for zeolite synthesis, the potential role of confinement should be taken into consideration when elucidating the putative pathway(s) of crystallization. Our findings indicate that heterogeneous nucleation of zeolites involves a nonclassical pathway wherein the origin of nucleation occurs on the exterior surface of amorphous precursors. The collective observations presented in this study do not, in of themselves, prove the existence of confinement effects, nor do they negate the possibility that nucleation can occur within the interior of bulk amorphous precursor phases. Our findings are qualitatively consistent with the premise that higher energetic barriers reduce the probability of nucleation within the interior of sol gels. The viewpoint that nucleation occurs on exterior surfaces of precursors is gaining traction in literature. Given that many zeolites are synthesized by conditions similar to those used in this study, these findings may prove relevant for other framework types.

ASSOCIATED CONTENT Supporting information The Supporting Information is available free of charge on the ACS Publications website at DOI: Additional X-ray diffraction patterns, DLS data, and electron microscopy measurements are available.

AUTHOR INFORMATION Corresponding author *E-mail: [email protected] Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS We are grateful to Dr. Siva Chinta for useful discussions. This work was supported by funding from the National Science Foundation (Award 1151098), The Welch Foundation (Award E-1794), and U of MN Nanostructural Materials and Processes Program; National Science Foundation (Award 0957696). Parts of this work were carried out in the Characterization Facility, U of MN, a member of the NSF-funded Materials Research Facilities Network (www.mrfn.org) via the MRSEC program. Electron microscopy was partially supported by the US Department of Energy, Office of Basic Energy Sciences, Division of Chemical Sciences, Geosciences &

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Biosciences. Pacific Northwest National Laboratory (PNNL) is a multiprogram national laboratory operated for DOE by Battelle.

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