Fabricating Ceramic Nanostructures with Ductile-Like Compression

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Fabricating Ceramic Nanostructures with DuctileLike Compression Behavior via Rapid Self-Assembly of Block Copolymer and Preceramic Polymer Blends Lisa M. Rueschhoff, Luke A. Baldwin, Robert Wheeler, Matthew J. Dalton, Hilmar Koerner, John Daniel Berrigan, Nicholas M Bedford, Soenke Seifert, Michael K. Cinibulk, and Matthew B. Dickerson ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.8b01820 • Publication Date (Web): 04 Dec 2018 Downloaded from http://pubs.acs.org on December 10, 2018

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ACS Applied Nano Materials

Article type: Full Article Fabricating Ceramic Nanostructures with Ductile-Like Compression Behavior via Rapid Self-Assembly of Block Copolymer and Preceramic Polymer Blends

Lisa M. Rueschhoff1, Luke A. Baldwin1, Robert Wheeler1, Matthew J. Dalton1, Hilmar Koerner1, John D. Berrigan1, Nicholas M. Bedford2, Soenke Seifert3, Michael K. Cinibulk1, and Matthew B. Dickerson1* *Corresponding author: [email protected] 1. Materials and Manufacturing Directorate, Air Force Research Laboratory, WrightPatterson AFB, Ohio, 45433, United States 2. School of Chemical Engineering, University of New South Wales, Sydney, NSW, 2052, Australia 3. X-Ray Sciences Division, Argonne National Laboratory, Argonne, Illinois, 60439, United States Keywords: block copolymers, ceramics, nanostructures, porous materials, self-assembly, in situ observation

ABSTRACT Here, we report a rapid and scalable process for the fabrication of nanostructured silicon carbide (SiC)-based ceramics displaying mechanical metamaterial properties. These novel mesoporous structures are achieved through patterning at the nanoscale via block copolymer (BCP) self-assembly. In this facile process, a blend of preceramic polymer (PCP) and BCP are dissolved in warm solvent, cast, and quickly solidify as an organogel comprised of a 3D micelle network. Significantly, these materials rapidly self-assemble and do not necessitate the annealing steps that are typically required for block copolymer self-assembly. The PCP nanostructure of the films is thermally stable and maintained through pyrolytic soft template removal and conversion to ceramic. Exploration of PCP, BCP, and PMMA homopolymer blends resulted in the discovery of a co-continuous wormlike micelle phase, which after 1 ACS Paragon Plus Environment

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pyrolysis translated into a ceramic nanocoral-like structure with a network of high aspect ratio ceramic struts punctuated with mesopores. In-situ nanomechanical compression testing reveals ductile-like deformation, complete strain recovery up to 14% strain, and enhanced energy absorption over bulk ceramics. The confluence of rapid self-assembly, affordability, and mechanical metamaterial properties offered by this system surmount many of the challenges associated with producing materials nanostructured over large areas. As such, these materials hold considerable promise for a variety of applications including energy storage, filtration, and catalytic materials.

1. INTRODUCTION The ability to control ceramic structure on the nanoscale enables the realization of new and/or improved properties and functions compared to conventional microstructured bulk materials.1–4 This has led to an increase in research during the last decade to template nanoscale ordered ceramics with tailored pore size, high specific surface area, and increased toughness.5– 16

Tailored pore size and high specific surface area are of interest for membrane, battery, and

catalyst applications, while increased toughness is advantageous for ceramic structural applications. Since no local plastic deformation mechanisms exist in traditional ceramics at room temperature, their strength is limited by crack size. Thus, strength can be increased through smaller maximum possible flaw sizes in nanostructured ceramics.17 Using this mentality, ceramic mechanical metamaterials (i.e., possessing mechanical properties such as ductility or strain recovery that are not possible in the bulk state) have been produced through manipulation at the nanoscale.12,17–19 To surmount inherent difficulties with processing ceramics, intricate processing techniques using either hard or soft templates are often utilized to achieve controlled ceramic nanostructures. Hard templates rely on a physical sacrificial template with desired morphology

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that must be removed after processing after coating or infiltrating with a ceramic material or precursor. On the other hand, soft templates utilize molecular templates such as structuredirecting polymers or surfactants in order to directly manipulate nanostructure. The technique used has historically depended on the desired resulting structure and/or properties of the ceramic nanostructure. In an exemplary instance of a hard templating approach, polymer lattices produced via 3D laser lithography were coated with a ceramic material via atomic layer deposition, and the polymer template subsequently removed to leave behind a hollow walled ceramic nanolattice.10,12,18–20 Careful manipulation of the lattice design and dimensions resulted in a ceramic metamaterial that was ultralight, strong, and energy absorbing.10,12,19 The high strength of these materials was attributed to the low probability of preexisting flaws in the nanosized structure, while the elastic shell buckling of the truss walls resulted in the ductile-like deformation and strain recovery.10,12,19 Elastic shell buckling acts as the preferred failure mechanism, rather than more detrimental brittle fracture or yielding, due to the geometry of the truss walls.10,12,19 Additionally, the stretching-dominated lattice structure has stiffness and strength values that scale proportionally with solid volume fraction.17 This allows for even extremely lightweight structures to maintain high enough relative strength for certain scaffold or bending applications. These studies provided the basic science understanding of ceramic nanostructure and property relationship needed to unlock the field of ceramic mechanical metamaterials, yet remain limited in large-scale applications by existing processing methodologies. The use of soft templating techniques have historically been successful in manufacturing ordered mesoporous ceramic materials with high specific surface area for functional applications.1,3,4,6,7,13 Block copolymers (BCP) are an attractive template material due to their ability to self-assembly on the nanoscale into well-defined and ordered nanostructures.21 Self-



assembly of the distinct polymer blocks occurs due to unfavorable monomer-monomer interactions and commonly requires a thermal and/or solvent annealing step. Generally, preceramic polymers (PCPs) are used with the BCP self-assembly process as a ceramic precursor due to their chemical modularity and ease of processing.7,22–29 The blending of PCPs and BCPs, as well as the custom synthesis of BCPs containing PCP blocks, have been used as a soft templating techniques for generating nanostructured ceramics.4,7,9 Structures are limited to the phases accessible by the BCPs, with only a few studies achieving three dimensional and interconnected nanostructures using this method.7,30 The templating of ceramic nanostructures with block-copolymer self-assembly is well established, but the mechanical properties have yet to be explored. Though elegant in approach and successful in achieving nanostructured materials, prior studies have inherent limitations in scalability and template removal or self-assembly kinetics. Therefore, there is a demand for the production of high-temperature ceramic materials with controlled nanostructure using an economical and scalable process. Moreover, the remarkable meta-mechanical properties achieved in prior ceramic nanostructure studies have only been realized using hard templating approaches and well-ordered ceramic nanolattices.10,12,18 To our knowledge, no studies have explored the use of soft template approaches to explore nanostructuring ceramics for increased mechanical behavior. In this work, we report the preparation of ceramic films with a novel coral-like nanostructure through the soft templating of a PCP with a BCP soft template. The rapid kinetics of our soft templating process, along with the low cost and commercial availability of the polymer blend, offers the potential for scalable processing techniques such as tape casting or roll-to-roll manufacturing. The ductile-like compression and strain recovery of our random interconnected nanocoral structure offers potential as a lightweight and energy-absorbing structural or ballistic protection material. The aforementioned capabilities and advantages


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enables considerable follow-on research and potential applications in the areas of energy storage, filtration, and catalytic materials. 2. RESULTS AND DISCUSSION Structure and Morphology Here, we report the preparation of ceramic films with nanostructures arising through the soft templating of a preceramic polymer (PCP) with the triblock copolymer poly(methyl methacrylate)-b-poly(n-butyl acrylate)-b-poly(methyl methacrylate), hereby denoted pMBM. When dissolved in a midblock selective solvent (e.g., 2-ethylhexanol), the pMBM endblocks form micelle cores that remain connected via the solubilized midblock, as illustrated in Fig. 1.31–34 The pMBM/PCP polymer blend creates spherical micelles, while the addition of supplemental endblock identical homopolymer poly(methyl methacrylate) (PMMA) leads to wormlike micelles (Fig. 1b). This macromolecular self-assembly acts to template a soluble polycarbosilane PCP. Thermal cross-linking is used to solidify the polycarbosilane (PCS), followed by high temperature pyrolysis for conversion to an amorphous polymer-derived silicon carbide (SiC) based ceramic nanostructure (Fig. 1c). There are several distinct differences between this pMBM system utilized in our study compared to those previously utilized in ceramic templating that bear noting here. In previous studies, self-assembly processing most often begins with a low viscosity solution containing diluted template and preceramic components.35–38 In these prior studies, structure development is then accomplished through lengthy annealing processes.35–38 This prior work contrasts with the temperature dependent micelle formation used here that creates a thermally reversible organogel with well characterized viscoelastic behavior.31–34 The instantaneous cooling of the organogel causes self-assembly of the pMBM micelles, without the need for additional thermal or solvent annealing steps. While the organogel micelle structure and rheology is well understood, it has yet to be explored as a removable soft template for preceramic polymers or



other soluble polymer systems. The commercially available PCS utilized in this work (SMP10) offers enhanced processing and rheology control over the use of crystalline ceramic powders and converts to SiC based ceramic at a conversion yield of ~80 wt.% upon high temperature (≥ 800 oC) pyrolysis.26,39,40 Films were fabricated for ease of characterization via flow coating polymer organogels (polymer concentration [Φp] = 0.14) onto a heated substrate. When using pMBM alone to template PCS, as-deposited films contained spherical micelles with a PMMA core with diameter of nominally 13 nm (as measured from the AFM image in Fig. 2a). Micelles were removed upon pyrolysis while maintaining PCS structure to create an amorphous SiC based ceramic with nominally spherical mesopores (SEM Fig. S1, ESI). A higher concentration or larger molecular weight of the micelle forming phase (PMMA) is known to create more three dimensional, bicontinuous micelle morphologies such as rod and wormlike micelles.41 Rather than modify the structure via chemical synthesis, a PMMA homopolymer was added to the pMBM/PCS polymer blend system (this blend hereby denoted as pMBMM). The use of all commercially available polymer systems maintains the modularity and accessibility of these materials for potential scalable processing. The as-deposited films made from a blend of pMBMM/PCS exhibit wormlike micelles (AFM phase Fig. 2b), a morphology that, to the best of our knowledge, has not been reported for the pMBMM blend system. These wormlike cylindrical micelles are approximately 13 nm in width (as measured from Fig 2b) with features such as y-junctions, loops, and long threadlike micelles. The nonuniformity of the micelle cylinders is likely due to the high molecular weight of the added homopolymer (115,000 g/mol) relative to the cumulative weight of the PMMA blocks in pMBM (32,000 g/mol). Wormlike micelles form due to interphase-curvature restraints within a small critical packing parameter window (1/3 < P < 1/2), and have been previously reported for a variety of block copolymer and homopolymer blends.41–44 With just 6 ACS Paragon Plus Environment

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the pMBMM blend (no PCS present), films exhibited interconnected micelles with spherical morphology rather than wormlike structures (Fig. S2). In fact, wormlike micelles were only observed with the presence on PCS. Analysis of small-angle x-ray scattering (SAXS) profiles of both pMBM + PCS and pMBMM + PCS bulk samples lead to subtle differences in the observed morphology between the two that is consistent with primary scatterers observed in AFM (Fig. S3). While the pMBM+PCS system was best fit with a simple sphere form factor, the pMBMM + PCS was fit with a cylinder form factor with high aspect ratio (wormlike). However, while this approach lead to the lowest residuals in the fitting, the low q upturn in both systems reflects clustering/agglomeration and lower q data is required to substantiate this hypothesis. Both scattering curves exhibit a peak (Q = 0.02 Å-1) superimposing the form factor, which is fitted with a hard sphere structure factor of the same particle population. This structure factor is similar in both and agrees with the order seen in AFM images. The similarities between SAXS (structure factor in bulk samples) and AFM (films) enforces the existence of order across all sample scales investigated.

Thermal Evolution; Template Removal and Ceramic Conversion The as-deposited film structure of the pMBMM +PCS polymer blend, as characterized used atomic force microscopy (AFM), is shown in Fig. 3a and Fig. S5. An omnipresent challenge in using soft templates is maintaining template fidelity during high temperature PCP to ceramic conversion. However, the gel soft template used here provides structural support to preserve the nanostructure. The as-deposited films were freeze-dried overnight to remove the solvent before a thermal cure step at 230 oC. Further optimization of the solvent and removal process would negate the need of freeze-drying, allowing for more continuous processing methods to be used. This relatively low temperature curing process results in crosslinking reactions that proceed through hydrosilylation, dehydrocoupling and oxidation pathways.45,46



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Minor changes are observed in the FTIR after curing due to the high percentage of silane and alkene groups present in the polymer, whereas the spectra dramatically changes upon further heating to 800 °C. FTIR spectra (Fig. S6) of the material through thermal conversion highlights the disappearance of the characteristic polymer OH, C-H, Si-H, and C=C bonds in the high temperature material and the appearance of prominent Si-O, C-O, and Si-C peaks, demonstrating conversion to an amorphous Si-C based ceramic. Fig. 3b shows an AFM height image of the cross-linked, solid-state PCS shaped in the inverse micelle template. The film was then heated to 800 oC for conversion to ceramic material and pMBMM template removal (Thermogravimetric analysis (TGA) shown in Fig. S4). The resulting ceramic nanostructure is shown in Fig. 3c (AFM height). Further analysis of the ceramic nanocoral structure was accomplished via scanning and transmission electron microscopy (SEM and TEM, respectively), shown in Fig. 4a-b. Removal of individual micelles left behind discreet wormlike mesopores surrounded by an interconnected amorphous ceramic strut structure. These images were used to estimate ceramic struts average thickness (t) of 20 nm. The length (L) of the struts was much more variable and difficult to accurately measure. An average length 200 nm was calculated from 10 separate ligaments ranging from nominally 100 – 500 nm in length. Elemental maps obtained via scanning

transmission

electron

microscopy

(STEM)

energy

dispersive

X-ray

spectroscopy (EDS) are shown in Fig. S7. Silicon, carbon, and oxygen are all detected, as expected for an amorphous SiC ceramic derived from PCS.5,40

Mechanical Behavior The nanocoral ceramic structures generated in this study has mechanical advantages compared to a bulk material due to the 1) low possible flaw size within the nanostructure, and 2) elastic buckling mechanism possible of the high aspect ratio ceramic struts.47,48 To analyze 8 ACS Paragon Plus Environment

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the compressive behavior of these materials, in-situ SEM testing was done on micropillars with radii of 1.5 and 3 μm fabricated via focused ion beam (FIB) milling. A full compression test was carried out on a micropillar in order to estimate compression behavior as well as the maximum attainable energy absorption. Subsequent cyclic compression testing was used to further understand the maximum strain recovery of the material. The resulting stress-strain curve and annotated micropillar from a single compression test are shown in Fig. 5 and Fig. S8. The stress values are calculated using the overall area of the micropillar and are therefore lower than the actual stress on the mesoporous material (see discussion in ESI). The three regions typically seen in bending-dominated foams are seen in the curve: I) linear elasticity (bending), II) plateau stress (inelastic buckling, yielding, or fracture), and III) densification.47,48 SEM images taken during these three regions are shown in Fig. 5c-e and are marked on the stress-strain curve with larger red markers. The general compression behavior shows a much more gradual failure compared to bulk ceramics which fail via catastrophic fast fracture, with greater elastic behavior than that seen for micro and macroporous ceramics.49 The elastic region (I) has a Young’s Modulus of 1.4 GPa (the slope of this region) and extends to 15% engineering strain, with potential recovery up to this point. In region II, the structure continues to collapse starting at a stress of 224 MPa (the ‘plateau stress’, σplateau) until the struts impinge at a densification strain of 57% (εd) when the stress rises steeply. Once the struts fully collapse within the pore structure, densification occurs (region III). The compressed film visual geometry (Fig. 5e) indicates a failure mechanism of inelastic buckling rather than brittle fracture. The energy absorbed in this structure (the shaded area under the curve in regions I and II) is 207 MPa per unit volume (or 0.9 nJ) and is larger than typically observed for dense ceramic materials. This energy absorption behavior is similar to that seen in micropillars of



both carbon and epoxy with porous nanostructures (1.0 – 1.5 nJ).50 A more detailed and annotated stress-strain curve is shown in Fig. S9. During cyclic compression, the micropillar was compressed and unloaded to increased maximum stress values each cycle, up to 1000 MPa. A full stress-strain graph of the cyclic test is shown in Fig 6a, with each compression cycle shown with a different colored line. The first cycle consisted of compression to 100 MPa followed by unloading and immediate compression to 200 MPa (cycle 2). This was continued by cycles to 300, 400, 600, and 1000 MPa. An enlarged view of the cycles to 100, 200 and 300 MPa are shown in Fig. 6b, with all three cycles showing primarily elastic behavior. The red data points in Fig 6b corresponding to images of the pillar from cycle 2 at full compression to 200 MPa/14% engineering strain (Fig. 6c) and unloaded to 0.05% strain (Fig. 6d), displaying 99.5% strain recovery of the material. This recovery could be attributed to elastic buckling of the individual ceramic struts. The Young’s modulus of the first three cycles (average of 1.3 GPa) closely matches that of the full compression test (overlaid graph shown in Fig. S10 with values in Table S4), but increases in successive tests due to strut collapse and structure densification. Even with some of these failure mechanisms, 65% of the strain was recovered in the last compression cycle (1000 MPa, 48% strain). This amount of strain recovery is extraordinary compared to bulk ceramic materials that experience catastrophic, non-recoverable failure at strains less than 1%.51 However, nanoceramic structures, such as octet-truss hollow nanolattices, have been found to experience such ductile deformation and strain recovery.10 These structures were found to first compress in a linear elastic regime followed by ductile-like controlled deformation with a plateau in stress after yielding.10 In this previous study, structures with the smallest wall thickness (10 nm) recovered up to ~98% of height after compression to 50% strain.10 Increasing the thickness to 20 nm resulted in marginal recovery after compression to 50% strain, and no recovery was


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observed for thick wall lattices (60 nm).10 The random cellular architecture of our ceramic nanocoral structure (and of ceramic foams and aerogels) results in bending dominated deformation of the struts.17,52 Since the random architecture is difficult to control, the greatest benefit would come from altering the aspect ratio (i.e. slenderness) of the nano ceramic struts. A decrease in ceramic strut thickness (t), and therefore the slenderness ratio (t/L), would favor elastic buckling to unlock greater strain recovery and absorbed energy. We are currently exploring this line of research by increasing the micelle width through altering the PMMA concentration and/or molecular weight.

3. CONCLUSIONS A novel polymer-derived SiC-based ceramic nanostructure was generated through templating with a block copolymer/homopolymer. Block copolymer self-assembly was driven by PMMA endblock micelle formation, with wormlike micelles achieved through the addition of a supplemental PMMA homopolymer. The templated PCS shape was retained through thermal treatment up to 800 oC for ceramic conversion. The resulting SiC-based ceramic material had discreet mesopores from the PMMA wormlike micelle removal that were interconnected with ceramic struts. This novel porous ceramic nanostructure afforded ductile compression behavior and higher energy absorbance during deformation than typical ceramic materials due to buckling of the ceramic struts. Cyclic compression testing revealed full strain recovery is possible up to 14% engineering strain. Realization of multiple micelle, and therefore pore, morphologies were achieved through varying polymer compositions. Significantly, the process detailed here utilizes inexpensive pMBMM templates and a commercial PCP without the use of additional annealing steps for BCP self-assembly. These factors make our process highly tractable for high throughput manufacturing techniques such as roll-to-roll deposition, making it possible to affordably produce nanostructured materials



over large areas. Such nanostructured SiC-based ceramics may be attractive for applications requiring materials that are high-temperature resistant, porous, and energy absorbing. 4. EXPERIMENTAL METHODS Materials and Sample Preparation Poly(methyl methacrylate)-b-poly(n-butyl acrylate)-b-poly(methyl methacrylate) block copolymer [Mw total = 64,000 g/mol, PnBa = 50 wt.%, PDI = 1.1953 ; denoted as pMBM] was synthesized through living anionic polymerization and provided by Kuarary Co., Ltd.. Poly(methyl methacrylate) (PMMA) homopolymer was purchased from Sigma-Aldrich Chemical Company [Mw = 120, 000 g/mol, PDI = 2 – 2.4 (reported from Sigma Aldrich)]. A commercially available PCS polymer [Trade name SMP-10, Starfire Systems, Mw =7,500 g/mol, PDI = 8.60] was used as a SiC PCP for the structural templating process. All chemicals were used as received without modification or purification. Blends of block copolymer (7 wt.% of the total solution) and homopolymer (2 wt.%) were prepared by dissolving in 2-ethylhexanol solvent [purchased from Alfa-Aesar] (84 wt.%) at 90 °C into a capped 20 mL glass vial on a stir plate (~30 min). After the polymers were fully dissolved, polycarbosilane (7 wt.%) was mixed in at 70 °C (~5 min). Film samples (~1.0 μm thickness) were prepared by flow coating onto a silicon wafer substrate (washed with ethanol, methanol, and UV-ozone cleaned) on a peltier plate at ~ 65 °C. The glass slide used as a blade was set to 500 μm and the stage velocity was 10 mm/s. As-deposited films were left in a fume hood overnight for limited solvent evaporation, followed by freeze drying for full solvent removal. The films were thermally cured (to cross-link the polycarbosilane) in a vacuum oven where they were heated at 1 oC/min and held at 160 oC and 230 oC for 1 h each. Pyrolysis was carried out up to 800 oC (1 h hold) in flowing argon gas in an alumina tube furnace equipped with a graphite sleeve at a heating rate of 1 oC/min. Characterization


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A Bruker Dimension AFM was used in tapping mode with a ~ 5 nm tip on a silicon cantilever (SSS-NCHR). The cantilevers resonance frequency was roughly 330 kHz. SEM images were taken in a Zeiss Gemini 500 field emission SEM operated at 0.5 keV and TEM images were taken in a FEI Talos operating in STEM mode with all samples sputter coated with ~10 nm of iridium. Phase dimensions calculated from AFM, SEM and TEM images were done so using ImageJ and were averaged from 25 manual measurements. In-situ mechanical testing and SEM images during testing were acquired in a FEI Quanta SEM operating at 5.0 keV. Ceramic film micropillars (~0.8 μm film thickness, pillar radius 3- 6 μm) were prepared through focused ion beam (FIB) milling in a Tescan Lyra-3 instrument using a gallium ion beam operating at 30 kV with a current of 66 pA. In-situ compression tests were completed using a micromechanical test frame manufactured by MicroTesting Solutions LLC with sample positioning (three degrees of freedom) enabled by piezoelectric stepper motion. Tip displacement was controlled through direct piezoelectric actuation.54 A flat tungsten tip (prepared via FIB milling to a diameter of approximately 6 μm) was used as a compression platen at a stroke rate of 100 nm/s, corresponding to a tip displacement of ~ 10 nm/min and a strain rate of 2 x 10-4 1/s. SEM images were acquired at 6 μs/pixel scan rate, with image resolution of 2000-3000 pixels/line. These images were used for strain measurements by application of digital image correlation methods. The planar surface area of the film micropillar was used for the calculation of stress. The porosity was not accounted for in this area and is therefore and underestimation of applied load on the individual struts. SAXS was performed at the 12-ID-B station of the Advanced Photon Source, Argonne National Laboratory. Samples (14 wt.% pMBM in 2-ethylhexanol) were loaded into 1 mm OD thin-walled quartz tubes and irradiated with 14 keV X-rays. SAXS patterns were collected on a Pilatus 2M detector at a sample to detector distance of 2 m.



CONFLICTS OF INTEREST There are no conflicts to declare. ACKNOWLEDGEMENTS The authors acknowledge support from the Air Force Research Laboratory Materials, Manufacturing Directorate Lab Director’s Fund, and the Air Force Office of Scientific Research under the Low Density Materials Portfolio. This research was performed while the author (LMR) held an NRC Research Associateship award at the Air Force Research Laboratory. This research used the 12-ID-B beamline of the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. The authors are grateful to Dr. Kathleen Cissel for SEM training, Kristi Singh for AFM assistance, and to Prof. Chelsea Davis (Purdue University) for helpful discussion.

Supporting Information Available: Additional AFM of block copolymer film morphologies, SAXS analysis, thermogravimetric analysis of polymer constituents, dynamic light scattering size analysis, FTIR of constituents through thermal treatment and ceramic conversion, STEM energy dispersive X-ray spectroscopy elemental analysis, detailed mechanical analysis calculation, and videos of in-situ mechanical testing.

REFERENCES (1) Wan, J.; Alizadeh, A.; Taylor, S. T.; Malenfant, P. R. L.; Manoharan, M.; Loureiro, S. M. Nanostructured Non-Oxide Ceramics Templated via Block Copolymer SelfAssembly. Chem. Mater. 2005, 17, 5613–5617. (2) Shi, Y.; Wan, Y.; Zhao, D. Ordered Mesoporous Non-Oxide Materials. Chem. Soc. Rev. 2011, 40, 3854. (3) Chan, V. Z. H.; Hoffman, J.; Lee, V. Y.; Iatrou, H.; Avgeropoulos, A.; Hadjichristidis, N.; Miller, R. D.; Thomas, E. L. Ordered Bicontinuous Nanoporous and Nanorelief Ceramic Films from Self Assembling Polymer Precursors. Science (80). 1999, 286, 1716–1719. (4) 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. 14 ACS Paragon Plus Environment

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Fig. 1 Schematic representation of nanostructured porous ceramic film production via micelle self-assembly. a) The starting physical polymer blend containing PMBM triblock copolymer (PMMA (blue) endblocks and PnBA (red) center block) and polycarbosilane (PCS) preceramic polymer (grey). b) Self-assembly of as-deposited polymer films into PMMA spherical or wormlike micelle cores. c) Thermal treatment (800 °C) converts the PCS to ceramic and removes the pMBM template, leaving behind mesopores (white).


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Fig. 2 AFM phase view of as-deposited films of a) polymer blend (pMBM + PCS) with spherical PMMA micelle cores (light phase) in a PCS matrix (dark) and b) pMBM + PMMA homopolymer (pMBMM+PCS) with PMMA wormlike micelle cores (light).



Fig. 3 Thermal evolution of self-assembled pMBMM + PCS film. a) AFM phase view of asdeposited polymer film showing wormlike micelles with PMMA core (light) in PnBA/PCS matrix (dark), b) AFM height view after thermal cure (230 oC) showing cross-linked PCS (light) and pMBMM template (dark). c) AFM height view after pyrolysis (800 oC) containing only nanostructured silicon carbide based ceramic (light). The z-axis height is 63 nm in b) and 68 nm in c).


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Fig. 4 a) SEM and b) TEM images of templated nanocoral ceramic films (post pyrolysis).

Fig. 5 a) Stress-strain curve for a micropillar tested in compression (red data points correspond to images c-e), highlighting the typical elastic (I), plateau (II), and densification (III) regions commonly seen in foam structures. b) SEM image of the micropillar testing set-up. c-e) SEM images of the sample at the following applied loads and corresponding engineering strain, respectively (the same scale is used for all images); c) 150 MPa, 11%, d) 382 MPa, 33%, e) 1200 MPa, 64%.



Fig. 6 Stress-strain curves for micropillar cyclic compression test of sequential full loading and unloading. a) All curves showing compression cycles from up to max values of 100, 200, 300, 600, and 1000 MPa. b) Enlarged view of the cycles to 100, 200, and 300 MPa. The line is shown to draw the eye between data points. The larger red data points in b) correspond to images of the sample c) loaded to 200 MPa (14% engineering strain) and subsequently d) unloaded to 0 MPa (0.05% engineering strain). The scale bar in c) corresponds to 1 µm, with the same scale used for both c) and d).

For Table of Contents Only Full recovery of ceramic compressive strain

Polymer

200

ΔT Ceramic

Stress (MPa)

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

0 0

Engineering Strain

0.15


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Fabricating Ceramic Nanostructures with Ductile-Like Compression

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