Self-Assembled Gyroidal Mesoporous Polymer-Derived High

Subscriber access provided by UCL Library Services

Article

Self-Assembled Gyroidal Mesoporous PolymerDerived High Temperature Ceramic Monoliths Ethan M. Susca, Peter A. Beaucage, Margaret A. Hanson, Ulrike WernerZwanziger, Josef Wilson Zwanziger, Lara A. Estroff, and Ulrich Wiesner Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b05011 • Publication Date (Web): 07 Mar 2016 Downloaded from http://pubs.acs.org on March 8, 2016

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

Chemistry of Materials is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

Self-Assembled Gyroidal Mesoporous Polymer-Derived High Temperature Ceramic Monoliths

Ethan M. Susca†, Peter A. Beaucage†, Margaret A. Hanson§, Ulrike Werner-Zwanziger§, Josef W. Zwanziger§, Lara A. Estroff†,ø, Ulrich Wiesner†,*

†

Cornell University, Department of Materials Science and Engineering, 214 Bard Hall Ithaca NY 14853, USA §

Dalhousie University, Department of Chemistry, 6274 Cobourg Rd, Halifax NS B3H 4R2 Canada

ø

Kavli Institute at Cornell for Nanoscale Science, Ithaca, NY 14853, USA

ACS Paragon Plus Environment

1


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 22

Abstract Polymer-derived ceramics (PDCs) have enabled the development of non-oxide ceramic coatings and fibers with exceptional thermo-mechanical stability. Here we report the self-assembly based synthesis of gyroidal (space group Q230, Ia3d) mesoporous silicon oxynitride ceramic monoliths by pyrolysis of blends of commercially available preceramic polysilazane with a structure-directing triblock terpolymer up to temperatures of 1000 ˚C. Monoliths had pore diameters of 9.4 ± 1.1 nm and surface area of 160 m2/g. The three-dimensionally (3D) ordered periodic pore structure of the as-made hybrids acts to relieve stresses by allowing the escape of gases formed during ceramization. This process in turn enables the retention of smooth monoliths during ceramization under ammonia, a process that both adds nitrogen to the material and removes carbon pyrolysis products. The monoliths are appealing for high-temperature applications such as catalyst supports, MEMS devices including gas and pressure sensors, as well as strong, stiff, and creep-resistant scaffolds for ordered interpenetrating phase composites. Introduction High temperature ceramics synthesized by the thermal decomposition of polymeric precursors, or polymer-derived ceramics (PDCs),1 have enabled breakthroughs in advanced ceramics such as high-temperature-stable2 ceramic coatings and fibers with exceptional creep3 and oxidation/corrosion resistence.4 These materials also exhibit high piezoresistivity,5 making them ideal as pressure sensors in high-temperature applications. Furthermore, as liquid precursors to ceramics, they have enabled precisely engineered ceramic nanostructures through templating and self-assembly processes.


2

Page 3 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Numerous studies have focused on tailoring nanostructure and porosity in PDCs for applications requiring large surface area and tailored pore sizes.6–10 The “hardtemplate” route to nanostructured PDCs uses pre-structured mesoporous materials (templates) that are infiltrated by liquid polymeric ceramic precursors. The template and ceramic precursors undergo pyrolysis before the template is usually removed with a strong acid etch post-pyrolysis.7 Hard-templating is limited by the dimensions with which the mesoporous template can be synthesized. This multistep process additionally requires complete polymer preceramic infiltration, the success of which depends strongly on surface chemistry. Finally, pore sizes of the resulting ceramic are limited by the wall thickness of the hard template used. In contrast to hard-templating, the co-assembly of pre-ceramic polymers with structure-directing macromolecules is a more direct and scalable self-assembly route to nanostructured ceramics. We have previously focused on solvent evaporation-induced self-assembly (EISA) of diblock copolymers to form well-ordered mesoporous non-oxide ceramics with a variety of morphologies. Systems of preceramic poly(urea-methyl-vinylsilazane) (Ceraset ®) with linear AB diblock copolymers: poly(isoprene)-bpoly(dimethylamino ethyl methacrylate), PI-b-PDMAEMA (IA), and poly(isoprene)-bpoly(ethylene oxide), PI-b-PEO (IO)11,12 have been thoroughly studied, but failed to realize 3D-bicontinuous morphologies and have been limited to one-dimensional hexagonal or two-dimensional lamellar structures.13 However, three-dimensionally (3D) ordered and interconnected pore structures are most favorable for applications requiring large accessible surface areas and/or specific pore sizes, i.e. production and storage of


3


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 22

energy and catalysis.14 To the best of our knowledge such 3D-ordered structures have not hitherto been realized for non-oxide ceramics using self-assembly. Linear ABC triblock terpolymers, often employing poly(styrene) as a middle block, have been shown to have significantly more area of their phase diagrams containing 3D-ordered morphologies when compared with diblock systems.15,16 In subsequent years, ABC triblock terpolymers have been used to structure direct inorganic materials.17–19 Furthermore, triblock terpolymers have been employed in conjunction with various additives to assemble into cubic mesoporous gyroidal materials that retain their monolithic identity post-pyrolysis.20,21 With this understanding, we rationally designed the triblock terpolymer, poly(isoprene)-b-poly(styrene)-b-poly(N,N-dimethylamino ethyl methacrylate), PI-b-PS-b-PDMAEMA (ISA), to self-assemble with a preceramic poly(methyl-vinyl-silazane) (PMVS) into the majority matrix phase of the double gyroid, space group Ia3d. This chemistry and the gyroidal structure are depicted in Scheme 1. The double gyroid is based on a minimal surface with zero mean curvature at every point on the surface, the so-called Schoen G minimal surface. It divides space into two equal volumes that are 3D continuous, interpenetrating, chiral, and related to each other by an inversion at the center of the unit cell (see red G minimal surface shown in Scheme 1a).22 Block copolymer self-assembly based double gyroids are not minimal surface structures, but are derived from the G minimal surface via the construction of surfaces on both sides of the minimal surface and then filling the volume in between so as to reach the correct volume fraction of the majority phase.23 The triblock terpolymer directed double gyroid hybrid structure discussed below is a core-shell structure, in which the two minority networks are constituted by a PI core surrounded by a PS shell (green


4

Page 5 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


and blue in Scheme 1, respectively). These two core-shell networks are in turn separated by the PDMAEMA + PMVS majority domain which upon pyrolysis constitutes the ceramic phase separating two interpenetrating pore networks. The triblock terpolymer (ISA) and preceramic polymer (PMVS) system used in this study retains a monolithic identity through high-temperature treatment. The highly symmetric and fully interconnected pore architecture provides a relief network for stresses caused by thermal gradients and gas formation that can otherwise result in uncontrolled pore formation, cracking, and fracture. In fact, this self-assembled monolith survives heat treatment under ammonia, a process that for monolithic preceramics obtained without a structure-directing polymer typically results in macroscopic pores and cracking (Figure S1).24 Here, adventitious processing is achieved under an ammonia atmosphere that simultaneously removes free carbon from the preceramic while condensing the network into a nitrogen-rich (oxy)nitride ceramic.


5


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 6 of 22

a

b

c n m O

H l O

i) 50 ºC ii)130 ºC EISA X-link N

H H Si N Si

N x H

y

450 ºC

1000 ºC

N2/H2

NH3

Scheme 1. (a-c) Structures, chemical compounds, and processing used in the present study. Colors correspond to the location of chemical components: PI (green), PS (blue), PDMAEMA + PMVS (red). In (a) the left side shows the core-shell double gyroid structure with the PMVS + PDMAEMA phase reduced to the gyroidal (G) minimal surface with zero mean curvature everywhere. On the right, the G minimal surface is shown by itself, splitting space into two equal-volume halves with interpenetrating architecture. Chemical structures (b) of the macromolecules and indication of hydrogen bonding promoting selfassembly of poly(methyl-vinyl-silazane), PMVS, with PI-b-PS-b-PDMAEMA (ISA with n = 140, m = 160, l = 190 and PDI = 1.11). Note diphenylethylene is employed after PI-b-PS diblock synthesis to prevent living chain ends from attacking the carbonyl group in the methacrylate block. PMVS composition is approximately x = 0.2 and y = 0.8, and includes oligomeric cyclic units with a wide polydispersity. (c) Schematic illustrating the co-assembly of PMVS with ISA supported by hydrogen bonding (dotted lines) and subsequent pyrolysis steps to yield a carbon-free silicon oxynitride.

Experimental Methods General. Every effort was taken to prevent the hydrolysis of poly(methyl-vinylsilazane) (PMVS) used in this work. Samples were kept in inert atmospheres and had minimal exposure to moisture during and between processing steps.


6

Page 7 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


ISA Polymer Synthesis. Standard Schlenk line techniques were used throughout the synthesis. Poly(isoprene)-b-poly(styrene)-b-poly(dimethylamino ethylmethacrylate) (ISA) was synthesized via sequential anionic polymerization as reported elsewhere.25,26 Briefly, sec-butyl-lithium (Aldrich, 1.4 M in cyclohexane) initiated poly(isoprene)-bpoly(styrene) active chain ends were synthesized in benzene and end-functionalized with diphenylethylene before a solvent swap to twice distilled tetrahydrofuran (Anhydrous, inhibitor free, ≥99%, Aldrich). Dimethylamino ethylmethacrylate (DMAEAMA) (98%, Aldrich) was distilled over triethylaluminum and added to the reactor at -40˚C. The reaction continued for 2.25 hours before degassed methanol was used to terminate the living polymerization. The polymer was thoroughly dried and used without further purification. Solvent EISA and thermal crosslinking. The processing route is shown in Scheme 1c. In a typical synthesis, 0.654 g of ISA, 0.359 g of PMVS, and 3.6 mg of radical initiator (dicumyl peroxide) were rapidly stirred for three hours at 5 wt% in anhydrous toluene. This mixture was poured into a 50 mL PTFE evaporating dish at 50 ˚C and evaporated slowly under a hemispherical dome in a continuously nitrogen-purged dry box for 48 hours. Samples were then placed under vacuum for 24 hours at 50 ˚C, and 3 hours at 130 ˚C to initiate crosslinking of the vinyl side groups on PMVS. The dry samples contained 34 wt% PMVS, 65.6 wt% ISA. Densities of 1.00 g/cm3 and 1.18 g/cm3 were used for PMVS and PDMAEMA (A) block, respectively, to estimate that the volume fraction of the hydrophilic phase composing the double gyroid matrix was 66 vol% (A + PMVS phase).


7


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 22

Thermal Processing and Ceramization. Samples were placed in an alumina boat in an alumina tube furnace and ramped at 1 ˚C/min to a dwell temperature of 450˚C under continuously purged nitrogen forming gas (5% H2 in N2; flow rate: 10 L/hr) and dwelled for 3 hours at 450 ˚C before slowly cooling to room temperature at 1 ˚C/min. Samples were placed in an alumina boat in a quartz tube furnace under anhydrous ammonia (8 L/hr flow rate) and ramped from room temperature to 600 ˚C in the course of 5.5 hours (1.75 ˚C/min) and dwelled at 600 ˚C for 3 hours. Samples were then ramped to 1000 ˚C in 4.5 hours (1.5 ˚C/min) and dwelled at 1000 ˚C for 3 hours. Samples were finally allowed 10 hours to cool to room temperature (cooling rate of 1.6 ˚C/min) before being removed from furnace. Characterization methods are described in the Supporting Information. Results and Discussion Experimental Design. Our route to polymer-derived ceramic monoliths results in a 3D-ordered, fully percolating, periodic pore structure, and can be macroscopically shaped by the mold in which solvent EISA is performed. Our procedure uses an ISA terpolymer containing 16 wt% PI (I), 30 wt% PS (S), and 54 wt% PDMAEMA (A) (Mn(ISA) = 57 kDa, PDI = 1.11) as a structure-directing agent for the oligomeric ceramic precursor, PMVS (tradename: Durazane 1800, AZ Electronic Materials), as depicted in Scheme 1b. PMVS (density 1.00 g/cm3) is incorporated into the PDMAEMA block (density 1.18 g/cm3) domain through selective hydrogen bonding and microphase separates from both the PI- and PS-blocks into the majority matrix phase (at 66 vol % PMVS + PDMAEMA) of the double gyroid.


8

Page 9 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


This work employs two thermal processes, in series, for the ceramization of composite films exposed to 130 ˚C for initial crosslinking. The first process at 450 ˚C further crosslinks the polysilazane network27 while simultaneously converting the I and S blocks into a disordered carbonaceous phase most likely lining the pores left behind by decomposition of these blocks (see black color, Figure 1a middle photo). Subsequent pyrolysis under ammonia to 1000 ˚C removes the carbonaceous phase from the monoliths as observed by the disappearance of black color (Figure 1a right photo).28–30 The carbonaceous phase helps preserve the macroscopic shape while processing under ammonia to 1000 ˚C (Figure 1a, S1). Samples synthesized without this intermediate pyrolysis step deformed, by curling into a rolled-carpet shape, during treatment in ammonia to 1000 ˚C (Figure S1b). In contrast, samples thermally treated from room temperature to 1000 ˚C under reducing nitrogen atmosphere only (5% H2 in N2; flow rate: 10 L/hr) stayed nanostructured and black, but did not curl up (data not shown). Attenuated total reflectance Fourier transform infrared spectroscopy (ATR-FTIR) performed after the intermediate (450 ˚C) processing step shows signatures for methyl and vinyl groups, which disappear upon final processing to 1000 ˚C (Figure S2). These results agree well with FTIR spectra previously established for these materials.30,31 The use of an in-situ formed carbonaceous phase to preserve nanostructure has previously been referred to as the combined assembly by soft and hard chemistries (CASH).32 It is interesting to note, however, that the primary effect of the carbonaceous phase here is not prevention of nanostructure collapse but rather the preservation of monolithic shape (no curling) during heat treatment.


9


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 22

a 130° C

450° C

1000° C

b

Figure 1. (a) Monolithic pieces exposed to different processing temperatures as indicated pictured against a quarter-inch grid. (b) Small angle X-ray scattering patterns of materials exposed to final temperatures of 130 ˚C (black d211 = 37 nm), 450 ˚C (teal, d211 = 28 nm), and 1000 ˚C (red, d211 = 19 nm). The d(211)-spacing was calculated using the primary peak (q* = 2π/d(211)). Tick marks indicate the first 16 expected reflections for a Q230 (Ia3d) symmetry with relative peak ratios of √3, 2, √7, √8, √10, √11, 2√3, √13, √15 … √25.33

Structural Materials Characterization. The ordered mesostructure was probed with small-angle X-ray scattering (SAXS), transmission electron microscopy (TEM) (Figure S3), and scanning electron microscopy (SEM). The structure evolution for the monoliths exposed to different processing temperatures of 130 ˚C, 450 ˚C, 1000 ˚C as observed by azimuthally-integrated plots of 2D SAXS patterns is shown in Figure 1b in conjunction with photographs of the macroscopic monoliths in Figure 1a. SAXS patterns are consistent with the symmetry of the double-gyroid space group, Ia3d. Five peaks


10

Page 11 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


matching the Ia3d symmetry were observed in all samples, and include the characteristic (211) and (220) reflections at relative locations of √6 to √8 inverse nm. Remarkably, the ordered mesostructure survives processing to 1000 ˚C. During processing, monoliths shrink nearly 50% in linear dimension to accommodate the increase in density accompanied by the polymer-to-ceramic conversion. The primary peak (q*) in each SAXS pattern corresponds to 2π/d(211) and can be used to follow the percent linear contraction during the evolution from polymer to ceramic materials. In Figure 1, q* clearly moves to larger values upon thermal processing indicating shrinkage of the structure. The d(211)-spacing decreases from 37 nm to 28 nm after 450 ˚C heat treatment, and shrinks further to 19 nm for samples processed to 1000 ˚C. These d211spacings suggest a 49% decrease in linear dimension starting from the sample treated to 130 ˚C which agrees well with observed reduction in macroscopic monolith size, diminishing by 48% laterally (Figure 1a). SEM was used to confirm the continuous percolating gyroidal network throughout the monoliths (Figure 2a–d). Scanning electron micrographs of samples processed to 1000 ˚C show characteristic facets of the double gyroid matrix, specifically the (211) and (110) planes, as depicted in Figure 2c and d, respectively. Two-dimensional fast-Fouriertransform (FFT) patterns of electron micrographs with the characteristic (211) “double wavy” pattern were used for samples processed to temperatures of 130 ˚C, 450 ˚C, and 1000 ˚C, respectively, to determine structural periodicities of 39, 25, and 20 nm (Figure S3) which agrees well with the corresponding SAXS-determined d-spacings described above. A summary of the contractions as observed from monolithic dimensions, SAXS dspacings, and TEM/SEM FFTs is provided in Table 1.


11


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 22

Table 1. Linear contractiona as determined by measuring macroscopic monolith dimensions, as well as SAXS and TEM/SEM image analysis. TEM/SEM FFT(211) Processing Monolith SAXS d(211) (nm) (nm) Temperature Diameter (cm) 130 ˚C 1.69 37 39 450 ˚C 1.35 (20%) 28 (24%) 25 (36%) 1000 ˚C 0.88 (48%) 19 (49%) 20 (49%) a

Parentheticals denote percent contraction relative to the dimensions measured for the 130 ˚C treated sample.

We observe large grains of gyroidal mesostructure several microns in diameter by SEM (Figure S4). The interface between grains is coherent, indicating that pores are connected between the grains. We suspect the isotropy of the cubic gyroidal mesostructure may aid in the retention of monolithic materials with coherent grain boundaries during the pyrolysis steps that result in large volume contractions.

Figure 2. SEM micrographs of fractured surfaces of gyroidal porous samples processed to 1000 ˚C (porous silicon oxynitride by 29Si-NMR), sputtered with Au-Pd for increased contrast. The colored rectangles in (a) are shown at higher magnifications in (b–d). (c) and (d) depict two commonly observed facets of the double gyroid structure: (211) and (110) planes, respectively. The mesoscale grain boundary between these facets can be observed in (a) and (b) and stretches from the upper left to lower right of each image.


12

Page 13 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Solid-state 29Si cross-polarization and magic angle spinning (CP/MAS) NMR spectroscopy was used to determine local structure of samples processed to 1000 ˚C. The 29

Si CP/MAS NMR spectrum (Figure 3) shows one broad peak with a maximum near a

chemical shift of -50 ppm relative to tetramethyl silane (TMS). Crystalline beta-Si3N4 has a chemical shift of -48.7 ppm,34,35 suggesting that the processed sample contains many SiN4-type tetrahedral sites; while beta-silicon carbide and alpha-quartz are reported at 17.5 and -107.4 ppm, respectively,36,37 regions of the spectrum in Figure 3 where only noise levels are observed. Mixed oxide and nitride resonances are found in the intermittent region between these extremes. The breadth of the single peak indicates the sample is amorphous and silicon bonding varies in composition from the pure nitride to mixed oxygen/nitrogen sites. We expect a bias towards silicon-bonding species near the surface of the mesoporous structure since proton induced cross-polarization was used to amplify the signal. The spectrum is nonetheless consistent with locally amorphous silicon oxynitride reported elsewhere.38 To support the 29Si-NMR characterization, elemental analysis was performed using X-ray photoelectron spectroscopy yielding atomic percentages of 36% Si, 43% N, 17% O, 4% C, suggesting a stoichiometric composition of SiN1.17O0.47C0.11. The amorphous nature of the monolith is further confirmed by the lack of peaks in powder x-ray diffraction (XRD) (Figure S5) and is expected for PDCs processed to 1000 ˚C.39


13

DppmI$int

ß-SiC

0e+00 0e+00

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Intensity (a. u.) 2e+08


Page 14 of 22

α-SiO2 ß-Si3N4

0 0

−50 −50

−100 −100

−150 −150

DppmI$ppm

Chemical Shift (ppm) 29

Figure 3. Solid-state Si CP/MAS NMR spectrum of mesoporous gyroidal sample processed to 1000 ˚C consistent with silicon oxynitride amorphous ceramic. Vertical lines indicate expected chemical shifts for β-Si3N4, β-SiC, and α-SiO2 provided as reference points.34–37

In order to characterize the porosity of the silicon oxynitride monoliths, nitrogen sorption isotherms (Figure 4a) were acquired and analyzed following the BrunauerEmmett-Teller (BET) method.40 A type IV-H1 hysteresis was observed,41 as expected for well ordered mesoporous materials, and a BET surface area of 158 m2/g calculated; only 14 m2/g of this total surface area is from micropores with sizes smaller than 2 nm, suggesting that the silicon oxynitride matrix domains are quite dense. The Barrett-JoynerHalenda (BJH) analysis (Figure 4b) suggests a narrow pore size distribution around 9.4 nm with a standard deviation of 1.1 nm.


14

Page 15 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 4. (a) Nitrogen sorption isotherms, and (b) pore size distribution calculated using the BJH model for cylindrical pores on silicon oxynitride samples processed to 1000˚ C.

A specific pore volume of 0.416 cm3/g was measured at a relative pressure of 0.99. The specific pore volume and density of the oxynitride phase is used to calculate the porosity of the monolith. A safe assumption for amorphous oxynitride density is between 2.8 g/cm3 and 3.1 g/cm3 42 corresponding to a porosity (fractional void space) of 54% to 56% (Supporting Info, S6). These porosity values suggest that the gyroidal matrix wall thickness must shrink relative to d-spacing reduction observed by SAXS in order to obtain a porosity larger than the initial volume percentage occupied by the PI and PS blocks (about 34 vol %). Conclusions This work explores a versatile self-assembly route for the synthesis of 3D bicontinuous gyroidal high temperature ceramic monoliths. We suspect the co-continuous cubic gyroid structure with space group Ia3d enables isotropic volume contraction during


15


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 22

pyrolysis, while a percolated mesopore network enables the release of gases that might otherwise cause cracking or microporosities. This mesostructuring enables formation of smooth gyroidal monoliths not possible from unstructured (ISA-free) PDC materials processed identically. The small wall thickness on the order of 10 nm enables rapid diffusion of gases through the walls during high-temperature solid-state processing and conversion and results in a dense ceramic wall phase with 91% of the measurable surface area of the porous monolith arising from mesopores. The high temperature ceramics described here are appealing for high temperature applications, particularly because they retain their monolithic integrity throughout processing, making them ideal platforms, e.g. for catalyst supports or MEMS sensing devices.3,6 Finally, these ceramics can be used as templates to backfill the 3D periodic interpenetrating pores with high-conductivity metals or tough polymeric materials which may result in materials with unusual properties like negative refractive index43 or high specific energy absorption.44 Associated Content Supporting Information Available: light micrographs of monoliths postprocessing, infrared spectroscopy, electron microscopy, x-ray diffraction, and porosity calculation. Additional details of methods include: gel permeation chromatography (GPC), proton NMR, solid state NMR, scanning and transmission electron microscopy (SEM, TEM), small angle x-ray scattering (SAXS), FT-IR spectroscopy, x-ray photoelectron spectroscopy (XPS), and nitrogen sorption measurements. Author Information *Email: [email protected]


16

Page 17 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Notes The authors declare no competing financial interest. Acknowledgements This work was supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Award No. DE-SC0010560. We thank the Cornell Center for Materials Research (CCMR) and the Cornell High Energy Synchrotron Source (CHESS) for use of their facilities, in particular the electron microscopy and polymer characterization facilities. CCMR Facilities are supported by the National Science Foundation under Award Number DMR-1120296, CHESS is supported by the NSF & NIH/NIGMS via NSF award DMR-0936384. The ssNMR work was performed at the Canadian Nuclear Magnetic Resonance Research Resource (NMR-3). PAB was supported by the National Science Foundation (NSF) Graduate Research Fellowship Program (NSF DGE-1144153). We would like to acknowledge Gregg McCraw for his generous donation of polysilazanes on behalf of AZ Electronic Materials. The authors would additionally like to thank Jon Shu, Francis. J. Disalvo, Ryo H. Wakabayshi, Jörg G. Werner, and Katherine P. Barteau for their help with this research.

References: (1)

Colombo, P.; Mera, G.; Riedel, R.; Sorarù, G. D. Polymer-Derived Ceramics: 40 Years of Research and Innovation in Advanced Ceramics. J. Am. Ceram. Soc. 2010, 93, 1805–1837.

(2)

Riedel, R.; Kienzle, A.; Dressler, W.; Ruwisch, L.; Bill, J.; Aldinger, F. A Silicoboron Carbonitride Ceramic Stable to 2,000 °C. Nature 1996, 382, 796–798.

(3)

An, L.; Riedel, R.; Konetschny, C.; Kleebe, H.-J.; Raj, R. Newtonian Viscosity of


17


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 22

Amorphous Silicon Carbonitride at High Temperature. J. Am. Ceram. Soc. 2005, 81, 1349–1352. (4)

Wang, Y.; Fei, W.; Fan, Y.; Zhang, L.; Zhang, W.; An, L. Silicoaluminum Carbonitride Ceramic Resist to Oxidation/corrosion in Water Vapor. J. Mater. Res. 2006, 21, 1625–1628.

(5)

Zhang, L.; Wang, Y.; Wei, Y.; Xu, W.; Fang, D.; Zhai, L.; Lin, K.-C.; An, L. A Silicon Carbonitride Ceramic with Anomalously High Piezoresistivity. J. Am. Ceram. Soc. 2008, 91, 1346–1349.

(6)

Bernard, S.; Miele, P. Nanostructured and Architectured Boron Nitride from Boron, Nitrogen and Hydrogen-Containing Molecular and Polymeric Precursors. Mater. Today 2014, 17, 443–450.

(7)

Shi, Y. F.; Meng, Y.; Chen, D. H.; Cheng, S. J.; Chen, P.; Yang, H. F.; Wan, Y.; Zhao, D. Y. Highly Ordered Mesoporous Silicon Carbide Ceramics with Large Surface Areas and High Stability. Adv. Funct. Mater. 2006, 16, 561–567.

(8)

Schitco, C.; Bazarjani, M. S.; Riedel, R.; Gurlo, A. NH 3 -Assisted Synthesis of Microporous Silicon Oxycarbonitride Ceramics from Preceramic Polymers: A Combined N 2 and CO 2 Adsorption and Small Angle X-Ray Scattering Study. J. Mater. Chem. A 2015, 3, 805–818.

(9)

Malenfant, P. R. L.; Wan, J.; Taylor, S. T.; Manoharan, M. Self-Assembly of an Organic-Inorganic Block Copolymer for Nano-Ordered Ceramics. Nat. Nanotechnol. 2007, 2, 43–46.

(10) Wan, J.; Malefant, P. R. L.; Taylor, S. T.; Loureiro S. M.; Manoharan, M. Microstructure of block copolymer/precursor assembly for Si-C-N based nanoordered ceramics. Mater. Sci. Eng. A. 2007, 463, 78–88. (11)

Kamperman, M.; Garcia, C. B. W.; Du, P.; Ow, H.; Wiesner, U. Ordered Mesoporous Ceramics Stable up to 1500 Degrees C from Diblock Copolymer Mesophases. J. Am. Chem. Soc. 2004, 126, 14708–14709.

(12)

Kamperman, M.; Du, P.; Scarlat, R. O.; Herz, E.; Werner-Zwanziger, U.; Graf, R.; Zwanziger, J. W.; Spiess, H. W.; Wiesner, U. Composition and Morphology Control in Ordered Mesostructured High-Temperature Ceramics from Block Copolymer Mesophases. Macromol. Chem. Phys. 2007, 208, 2096–2108.


18

Page 19 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(13)

Kamperman, M.; Fierke, M. A.; Garcia, C. B. W.; Wiesner, U. Morphology Control in Block Copolymer/Polymer Derived Ceramic Precursor Nanocomposites. Macromolecules 2008, 41, 8745–8752.

(14)

Kamperman, M.; Burns, A.; Weissgraeber, R.; van Vegten, N.; Warren, S. C.; Gruner, S. M.; Baiker, A.; Wiesner, U. Integrating Structure Control over Multiple Length Scales in Porous High Temperature Ceramics with Functional Platinum Nanoparticles. Nano Lett. 2009, 9, 2756–2762.

(15)

Epps, T. H.; Cochran, E. W.; Bailey, T. S.; Waletzko, R. S.; Hardy, C. M.; Bates, F. S. Ordered Network Phases in Linear Poly(isoprene- B -Styrene- B -Ethylene Oxide) Triblock Copolymers. Macromolecules 2004, 37, 8325–8341.

(16)

Shefelbine, T. a.; Vigild, M. E.; Matsen, M. W.; Hajduk, D. a.; Hillmyer, M. A.; Cussler, E. L.; Bates, F. S. Core−Shell Gyroid Morphology in a Poly(isopreneBlock -Styrene- Block -Dimethylsiloxane) Triblock Copolymer. J. Am. Chem. Soc. 1999, 121, 8457–8465.

(17)

Toombes, G. E. S.; Mahajan, O. S.; Weyland, M.; Jain, A.; Du, P.; Kamperman, M.; Gruner, S. M.; Muller, D. A.; Wiesner, U.; V, C. U.; York, N.; High, C.; Synchrotron, E.; Chess, S.; May, R. V; Re, V.; Recei, M.; September, V. SelfAssembly of Four-Layer Woodpile Structure from Zigzag ABC Copolymer / Aluminosilicate Concertinas. Macromolecules 2008, 41, 852–859.

(18)

Stefik, M.; Mahajan, S.; Sai, H.; Epps, T. H.; Bates, F. S.; Gruner, S. M.; Disalvo, F. J.; Wiesner, U. Ordered Three- and Five-Ply Nanocomposites from ABC Block Terpolymer Microphase Separation with Niobia and Aluminosilicate Sols. Chem. Mater. 2009, 21, 5466–5473.

(19)

Arora, H.; Du, P.; Tan, K. W.; Hyun, J. K.; Grazul, J.; Xin, H. L.; Muller, D. A; Thompson, M. O.; Wiesner, U. Block Copolymer Self-Assembly-Directed SingleCrystal Homo- and Heteroepitaxial Nanostructures. Science 2010, 330, 214–219.

(20)

Werner, J. G.; Hoheisel, T. N.; Wiesner, U. Synthesis and Characterization of Gyroidal Mesoporous Carbons and Carbon Monoliths with Tunable Ultralarge Pore Size. ACS Nano 2014, 8, 731–743.

(21)

Robbins, S. W.; Sai, H.; Disalvo, F. J.; Gruner, S. M.; Wiesner, U. Monolithic Gyroidal Mesoporous Mixed Titanium and Niobium Nitrides. ACS Nano 2015, 8,


19


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 22

8217–8223. (22)

Schoen, A. H. Infinite periodic minimal surfaces without self-intersections. NASA Technical Note 1970, No. D-5541.

(23)

Hajduk, D.A., Harper, P.E., Gruner, S.M., Honeker, C.C., Kim, G., Thomas, E.L. and Fetters, L.J. The gyroid: A new equilibrium morphology in weakly segregated diblock copolymers. Macromolecules 1994, 27, 4063-4075.

(24)

Schmidt, W. R.; Marchetti, P. S.; Interrante, L. V; Hurley, W. J.; Lewis, R. H.; Doremus, R. H.; Macielt, G. E. Ammonia-Induced Pyrolytic Conversion of a Vinylic Polysilane to Silicon Nitride. Chem. Mater. 1992, 7, 937–947.

(25)

Creutz, S.; Teyssie, P.; Je, R. Living Anionic Homopolymerization and Block Copolymerization of (Dimethylamino) Ethyl Methacrylate. Macromolecules 1997, 30, 6–9.

(26)

Li, Z.; Hur, K.; Sai, H.; Higuchi, T.; Takahara, A.; Jinnai, H.; Gruner, S. M.; Wiesner, U. Linking Experiment and Theory for Three-Dimensional Networked Binary Metal Nanoparticle-Triblock Terpolymer Superstructures. Nat. Commun. 2014, 5, 3247.

(27)

Chen, Y.; Yang, X.; Cao, Y.; Gan, Z.; An, L. Quantitative Study on Structural Evolutions and Associated Energetics in Polysilazane-Derived Amorphous Silicon Carbonitride Ceramics. Acta Mater. 2014, 72, 22–31.

(28)

Burns, G. T.; Chandra, G. Pyrolysis of Preceramic Polymers in Ammonia: Preparation of Silicon Nitride Powders. J. Am. Ceram. Soc. 1989, 72, 333–337.

(29)

Choong Kwet Yive, N. S.; Corriu, R. J. P.; Leclercq, D.; Mutin, P. H.; Vioux, A. Thermogravimetric Analysis/Mass Spectrometry Investigation of the Thermal Conversion of Organosilicon Precursors into Ceramics under Argon and Ammonia. 2. Poly(silazanes). Chem. Mater. 1992, 4, 1263–1271.

(30)

Wan, J.; Gasch, M. J.; Mukherjee, A. K. Effect of Ammonia Treatment on the Crystallization of Amorphous Silicon – Carbon – Nitrogen Ceramics Derived from Polymer Precursor Pyrolysis. J. Am. Ceram. Soc. 2002, 85, 554–564.

(31)

Scarlete, M.; McCourt, N.; Butler, I. S.; Harrod, J. F. Spectroscopic Investigation of the Synthesis of Thin Silicon Nitride Films on Silicon Single-Crystal Wafers via


20

Page 21 of 22

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Ammonia-Assisted Pyrolysis of Organosilicon Polymers. Chem. Mater. 2001, 13, 655–661. (32)

Lee, J.; Orilall, M. C.; Warren, S. C.; Kamperman, M.; DiSalvo, F. J.; Wiesner, U. Direct Access to Thermally Stable and Highly Crystalline Mesoporous TransitionMetal Oxides with Uniform Pores. Nat. Mater. 2008, 7, 222–228.

(33)

Foerster, S.; Khandpur, A. K.; Zhao, J.; Bates, F. S.; Hamley, I. W.; Ryan, A. J.; Bras, W. Complex Phase Behavior of Polyisoprene-Polystyrene Diblock Copolymers Near the Order-Disorder Transition. Macromolecules. 1994, 23, 6922–6935.

(34)

Carduner, K. R.; Carter III, R. O.; Milberg, M. E.; Crosbie, G. M. Phase Composition of Silicon Nitride Powders by Silicon-29 Magic Angle Spinning Nuclear Magnetic Resonance Spectroscopy. Anal. Chem. 1987, 59, 2794–2797.

(35)

Carduner, K. R.; Carter III, R. O.; Rokosz, M. J.; Peters, C.; Crosbie, G. M.; Stiles, E. D. Silicon-29 Magic -Angle Spinnning Nuclear Magnetic Resonancy Spectroscopy of Sintered Silicon Nitride Ceramics. Chem. Mater. 1989, 1, 302– 307.

(36)

Carduner, K. R.; Shinozaki, S. S.; Rokosz, M. J.; Peters, C. R.; Whalen, T. J. Characterization of β–Silicon Carbide by Silicon–29 Solid-State NMR, Transmission Electron Microscopy, and Powder X-ray Diffraction. J. Am. Ceram. Soc. 1990, 73, 2281–2286.

(37)

Lippmaa, E.; Mägi, M.; Samoson, A.; Engelhardt, G.; Grimmer, A. R. Structural Studies of Silicates by Solid-State High-Resolution 29Si NMR. J. Am. Chem. Soc. 1980, 102, 4889–4893.

(38)

Chollon, G.; Hany, R.; Vogt, U.; Berroth, K. A Silicon-29 MAS–NMR Study of Silicon Nitride and Amorphous Silicon Oxynitride Fibres. J. Eur. Ceram. Soc. 1998, 18, 535–541.

(39)

Mera, G.; Ishikawa, R.; Ionescu, E.; Ikuhara, Y.; Riedel, R. Atomic-Scale Assessment of the Crystallization Onset in Silicon Carbonitride. J. Eur. Ceram. Soc. 2015, 35, 3355–3362.

(40)

Sing, K. S. W.; Everett, D. H.; Haul, R. A. W.; Moscou, L.; Pierotti, R. A.; Rouquérol, J.; Siemieniewska, T. Reporting physisorption data for gas/solid


21


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 22

systems with special reference to the determination of surface area and porosity (Recommendations 1984). Pure Appl. Chem. 1985, 57, 603–619. (41)

Thommes, M.; Kaneko, K.; Neimark, A. V.; Olivier, J. P.; Rodriguez-Reinoso, F.; Rouquerol, J.; Sing, K. S. W. Physisorption of gases, with special reference to the evaluation of surface area and pore size distribution (IUPAC Technical Report). Pure and Appl. Chem. 2015, 87, 1051-1069.

(42)

Grahn, H. T.; Maris, H. J.; Tauc, J.; Hatton, K. S. Elastic Properties of Silicon Oxynitride Films Determined by Picosecond Acoustics. Appl. Phys. Lett. 1988, 53, 2281–2283.

(43)

Hur, K.; Francescato, Y.; Giannini, V.; Maier, S. A.; Hennig, R. G.; Wiesner, U. Three-Dimensionally Isotropic Negative Refractive Index Materials from Block Copolymer Self-Assembled Chiral Gyroid Networks. Angew. Chem. Int. Ed. 2011, 50, 11985–11989.

(44)

Lee, J.; Wang, L.; Boyce, M. C.; Thomas, E. L. Periodic Bicontinuous Composites for High Specific Energy Absorption. Nano Lett. 2012, 12, 4392–4396.

Table of Contents Image:

Polymer/Preceramic Blend

SiN1.17O0.47C0.11 Ceramic


22

Self-Assembled Gyroidal Mesoporous Polymer-Derived High

Recommend Documents