Spontaneous Rippling and Subsequent Polymer Molding on Yttria-Stabilized Zirconia (110) Surfaces Haris M. Ansari,†,‡ Zhiyuan Niu,† Chen Ge,† Suliman A. Dregia,† and Sheikh A. Akbar*,† †
Department of Materials Science and Engineering, The Ohio State University, Columbus, Ohio 43210, United States School of Chemical and Materials Engineering, National University of Sciences and Technology (NUST), H-12, Islamabad, Pakistan
‡
ABSTRACT: Spontaneous nanoripple formation on (110) surfaces of yttria-stabilized zirconia, YSZ-(110), is achieved by diffusional surface doping with rare-earth oxides. Periodic arrays of parallel nanobars separated by channels (period ∼100 nm) grow out of the dopant sources, covering relatively wide areas of the surface (∼10 μm). The nanobars mound up on the surface by diffusion, exhibiting morphological uniformity and alignment, with their long axis lying parallel to the [11̅0] direction in the YSZ-(110) surface. The process for forming these nanobar arrays can be as simple as sprinkling of rare-earth oxide powder (dopant source) on YSZ-(110) surface and annealing in a high temperature air furnace. However, higher control on dopant dispersion on the surface is demonstrated with other techniques, including photolithography and inkjet printing. The ripple arrays extend anisotropically on the (110) surface, obeying the parabolic growth law, and showing principal values of the rate constant along [11̅0] (maximum) and [001] (minimum), as expected from the symmetry of the (110) surface. The self-patterned ceramic substrates are well-suited for pattern transfer by replica molding, as illustrated by single-step molding with polydimethylsiloxane (PDMS), which is a widely used biomaterial in cell-culture studies. KEYWORDS: self-assembly, nanobars, rare-earth oxides, surface patterning, nanoimprinting bility.12,13 The physical basis of ATG is that a stressed solid may be unstable with respect to subdividing because the introduction of free surfaces, where tractions must vanish, leads to some relaxation of stress (and elastic strain energy), albeit at the expense of increasing surface area (and surface energy). For the onset of instability on the surface of a uniaxially stressed semi-infinite solid, Asaro and Tiller considered the evolution of periodic infinitesimal ripples by surface diffusion. They predicted that all ripple waves with a wavelength greater than a critical wavelength, λ ≥ λc, would be unstable and become amplified, with a maximum amplification 8π E γ 4 rate for a wavelength λ* = 3 λc = 2 , where E denotes the 3σ Young modulus, γ the surface energy, and σ the stress. In epitaxy, the stress arises from coherency strains in the deposit, in order to achieve lattice matching with the substrate, and it is therefore dependent on the value of misfit between the deposit and the substrate. The ATG theory provides a simple estimate to compare with experimentally observed length scales in selfassembled arrays of epitaxial nanostructures. There have been
O
ne dimensional (1D) inorganic nanostructures have recently attracted significant attention because their remarkable functional properties have enabled a host of promising applications in electronics, ferroics, and sensing.1−4 Several top-down and bottom-up approaches have been developed to grow 1D nanostructures, both discrete and periodic, although it is quite challenging to satisfy the conflicting demands of low cost and high throughput with a high degree of organization. Among the various approaches to fabricate 1D planar nanostructures on surfaces, self-assembly can offer a practical, low-cost route to covering large surface areas with self-organized arrays of nanostructures. Self-assembly can also be combined with lithographic techniques to exercise remarkable control over multiscale patterning, as well as spatial distribution and orientation for applications requiring high complexity and precision.5,6 The stress-induced self-assembly of surface nanostructures has been widely studied and extensively reviewed for epitaxially grown semiconductor systems.7,8 Thin deposits may subdivide spontaneously (during growth or on subsequent annealing) into arrays of islands that exhibit varying order in size, shape, crystal orientation, spacing and spatial alignment.9−11 The phenomenon is widely understood on the basis of the stressinduced Asaro−Tiller−Grinfeld (ATG) morphological insta© 2017 American Chemical Society
Received: January 5, 2017 Accepted: February 6, 2017 Published: February 6, 2017 2257
DOI: 10.1021/acsnano.7b00081 ACS Nano 2017, 11, 2257−2265
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Figure 1. (a) SEM micrograph showing the GDC patch configuration after photoresist removal and annealing at 1200 °C for 5 h. (b) SEM micrograph showing the nanobar array at the edge of one of the GDC patches (inset, FFT). (c) SEM micrograph showing that the nanobars maintain their orientation around the corner of a typical GDC patch.
deposition29 or by changing the substrate surface orientation to LAO-(011).22,30 The symmetry-breaking that occurs due to the two aforementioned approaches selectively promotes the growth of nanowires parallel to the preindented grooves or the single [100] direction in the LAO-(011) surface. The previous examples demonstrate that the experimental environment and the choice of the substrate are key parameters in controlling the morphology and alignment in strain-mediated 1D oxide nanostructures. However, the challenges of achieving selforganized patterns with spatial correlations between 1D oxide nanostructures clearly underscore the need for a simpler and scalable process with wide-area applicability. Here, we report a spontaneous nanopatterning process for producing pseudoperiodic arrays of parallel nanobars on the (110) surface of commercially available substrates of yttriastabilized zirconia (YSZ). The simple process is based on solidstate doping of the substrate surface with rare-earth oxides at a low homologous temperature (∼1/2 the absolute melting point) where surface diffusion is expected to be much more active than bulk diffusion. We report results based on doping with gadolinia-doped ceria (GDC) and Eu2O3 (EO). The nanopatterning of advanced ceramics can have an impact on a number of important technological applications. Doped ceria and zirconia are technologically important hightemperature ceramics with broad ranging applications in solidoxide fuel cells (SOFCs),31 superconducting tapes,32 sensing,33 and catalysis.34,35 Moreover, when synthesized in strained nanocrystalline forms, e.g., epitaxial thin films and multilayers, YSZ and other ionic conductors can exhibit superionic conductivities, with great technological advancement promised
extensive developments of the original ATG theory, in order to account for effects of nonlinear evolution dynamics beyond the onset of the instability, as well as the combined effects of elasticity and surface energy anisotropy on the evolution of shapes and coarsening of self-assembled nanoislands.7,8 Although most previous studies have focused on vapordeposited semiconductor materials, it has recently been shown that self-assembly of nanostructures could also be achieved on ceramic surfaces with remarkable order,14−16 Also, some thermodynamic and kinetic models have predicted the possibility of achieving 1D nanostructures,17,18 and this approach has been able to yield nanostructure elongation in two orthogonal directions on atomically flat substrate surfaces,19 or else required miscut substrates to confine the growth along a single direction.20,21 Some success has also been achieved in a few oxide systems via deposition of nano features on single crystal substrates from chemical solutions,22−24 by vapor deposition with molecular beam epitaxy (MBE),25 and by confined crystallization in track-etched nanoporous polymer templates from sol−gel based precursors.26,27 Gibert et al. used chemical solution deposition to grow ceramic nanowires on the (001) surface of LaAlO3 (LAO) substrates,23 and they showed the individual feature morphologies were dependent on the experimental environment. In a reducing atmosphere, elongated nanowires were formed in two mutually orthogonal ⟨100⟩ directions in the surface; whereas discrete nanoislands were observed under an oxidizing atmosphere.23,28 Preferential elongation of the nanowires along one of the two equivalent ⟨100⟩ directions on (001) was achieved by scratching grooves on the substrate surface with a nanoindenter prior to 2258
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Figure 2. SEM micrograph depicting the nanobar morphology after annealing at 1200 °C for (a) 15 h and (b) 25 h. (c) SEM micrograph of the GDC patch corner and the nanobar array emanating from it.
by reductions in their operating temperatures.36−38 The nanoscale roughening of the YSZ surface can impact its chemical reactivity39,40 in low-temperature gas−solid reactions that are important in sensing and fuel cell applications. In addition, the higher doping of the nanobars relative to their substrate provides a means for modifying their state of stress, possibly leading to anisotropic superionic conductivity. Thus, surface patterning by solid-state doping of YSZ may offer a simple alternative route to engineering materials for ionic devices. And owing to its hardness and thermochemical stability, patterned YSZ can also serve as a mold for pattern transfer to other materials. Also, the processing of these nanostructures requires only simple steps and can be easily scaled up to cover large surface areas without the need for lithography and thin film deposition techniques. In this paper we are concerned with the pattern formation, including the processing, characterization, and the transfer of the nanoripple arrays to biofunctional polymers. Studies in tissue engineering show that biological cells respond to grooveridge topographies41−43 that mimic features of the extracellular matrix.44 For example, corneal epithelial membrane has features similar to the groove/ridge topography of the nanobar pattern, with dimensions ranging from 20−200 nm.45 While YSZ ceramic is not a substrate of choice in cell culture studies, proof of concept studies have shown that cells do respond to topography on YSZ surfaces directly.46,47 It has also been demonstrated that self-assembled discrete nanoislands on YSZ(100) surfaces can be used as master patterns for nanoimprinting important biomedical polymers such as polydimethylsiloxane (PDMS), polystyrene (PS) and ethylene glycol dimethacrylate (EGDMA).48 Unlike a pattern of discrete islands on a surface, a ripple pattern is equivalent to its own
negative, and therefore it can be transferred in a single step. We demonstrate pattern transfer with high fidelity molding with PDMS, which can in turn be used as a platform to carry out biological cell attachment and proliferation studies.
RESULTS AND DISCUSSION Nanostructures grown by our technique require dispersing solid dopant sources on single crystal substrates with appropriate elastic modulus anisotropy, followed by annealing at a temperature T ∼ 0.5Tm. Initially, dopant sources were arrayed on YSZ-(110) substrates by a combination of RF sputtering and liftoff photolithography of GDC thin films from GDC targets, leaving behind isolated GDC patches on a bare YSZ surface, as illustrated in Figure 1a. Upon annealing at 1200 °C for 5−25 h in air, nanobar arrays appear near the edges of the dopant patches. In Figure 1a, which was acquired after a relatively short anneal (5 h at 1200 °C), a band of uniform width is visible around all 9 patches in the image. Figure 1b is a higher magnification SEM micrograph revealing that the bands consist of self-assembled nanobar arrays. It is noteworthy that the nanobars appear on surfaces that were masked by the photoresist during sputter deposition of the GDC film, hence they did not have prior contact with GDC. As seen in Figure 1b, the nanobars have uniform widths along their lengths, are closely spaced, and run remarkably parallel to one another without any external influences. A typical nanobar has a length of several microns and a width of ∼90 nm. The evidence that this remarkable alignment is not guided by the edge of the GDC patch is presented in Figure 1c, which includes a corner of a GDC patch. It is clear that the nanobars maintain the same orientation around the patch corner, which means the intrinsic YSZ surface structure/symmetry is guiding nanobar alignment 2259
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Figure 3. (a) SEM micrograph showing a 15 μm × 10 μm area covered with self-assembled nanobars. Footprints of remnant source particles are also visible (marked by arrows). (b) Higher magnification SEM micrograph showing the nanobar morphology produced by the powder suspension method (inset, FFT). (c) STEM micrograph showing the nanobars in cross-section. The SAED pattern in the inset has a [11̅0] zone axis, which is also the long axis of the nanobars. The sample was annealed at 1200 °C for 25 h in air. (d) An EDS depth profile from the top to the bottom of a typical nanobar. The intensities are normalized by the maximum intensities of the respective species.
as the experimentally observed periods of the nanobar patterns, although the latter can be altered by coarsening, as discussed further below. Both doping and nanobar formation occur concurrently by surface diffusion during the high temperature anneal resulting in unidirectional ceramic oxide nanobar arrays without the necessity for orientational ordering through artificial means. Since doping of the YSZ surface and subsequent breakup occur by surface diffusion, the extent of the nanobar arrays emanating from the GDC patches can be controlled by longer soak times at the same temperature. By using appropriate annealing conditions, the nanobar arrays can be made to extend wall-to-wall between the GDC patches. This full coverage is illustrated in Figures 2a and 2b where the samples were annealed for 15 and 25 h, respectively, at 1200 °C. The morphology of the nanobar arrays after long anneals is different from the one seen in the early stages, e.g., in Figure 1b, where the bars were straight along their principal axis. In Figure 2a, the bars are not as straight and frequently kink to accommodate other bars growing around them. Also, the bars have an average width of ∼160 nm, which suggests that the bars coarsen at longer soak times. Annealing for longer times causes the average nanobar width to increase further to ∼190 nm, and frequent breaks appear along their lengths (Figure 2b). We believe that the continuity of the bars is a function of the dopant concentration. Bars that are continuous at low dopant concentrations (low stress) break up into shorter segments as their dopant concentration (stress) increases. Figure 2c shows the corner of a GDC patch and the nanobars around it. All the
rather than the shape of dopant patch. The nanobar array forms by a mechanism similar to the one that produces self-assembled nanoisland arrays on YSZ-(001) surfaces.16 The GDC patches predominantly act as sources to dope the YSZ surface in their proximity. At the employed temperatures, surface diffusion is expected to be several orders of magnitude faster than bulk diffusion.16,49 Therefore, the dissolution of dopants (Gd3+ and Ce4+) with larger ionic radii than the host ions (Zr4+)50,51 creates stresses in a thin surface layer, and these conditions lead to the ATG-like morphological instability, where stress relaxation is achieved by introduction of free surfaces at the expense of increased surface energy. For quantitative comparison, we estimate the characteristic length scale λ* predicted by the ATG theory for a thin surface layer of doped YSZ, using the Y2O3−ZrO2 system for illustration because of availability of the requisite information. At a temperature of 1200 °C the YSZ phase can dissolve up to 31 mol % Y2O3,52 and based on the measured composition dependence of lattice parameter in Y2O3−ZrO2 solid solutions,53 the Y2O3-saturated surface layer would have a lattice parameter mismatch ε = 1.3% with the 8-mol % substrate. Thus, the coherency stress can be estimated as σ ∼ E[001]ε where E[001] is the Young modulus of the substrate in the [001] direction (normal to the bar axis), and based on the measured elastic constants of YSZ,54 E[001] = 360 GPa. For the (011) surface of YSZ, the calculated surface energy γ = 1.44 J/m2.55 On the basis of the available parameters, the ATG theory would predict an optimal 8π E γ wavelength λ* = 3σ 2 = 198 nm, which is on the same order 2260
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Figure 4. (a) SEM micrograph showing self-assembled nanobars produced by using Eu2O3 as dopant. (b) A higher magnification SEM micrograph showing the nanobar morphology (inset, FFT). The sample was annealed at 1200 °C for 25 h in air. (c) SEM micrograph depicting the onset of nanobar formation.
another. Impingement along the tips creates a defect that resembles a shearing fault (oval in Figure 3b) and that along the sides creates a defect in the pattern that resembles an edge dislocation (circle in Figure 3b). In a previous publication, we have shown that self-assembled nanoislands on GDC-doped YSZ-(001) surfaces are truncated pyramids with square bases, {111} inclined side facets and (001) top facets.16 The nanobars, however, are smoothly curved as revealed by scanning transmission electron microscopy (STEM) cross-section imaging (Figure 3c). Also, they are ∼27 nm in height which is significantly smaller than that of a typical nanoisland on YSZ-(001) (∼50−75 nm).16 Figure 3c also shows a selected area electron diffraction pattern (SAED) with a [11̅0] zone axis in the inset, which reveals that the bars are epitaxial and have long axes parallel to the [11̅0] in the YSZ-(110) surface. Figure 3d shows a STEM-EDS depth profile from the top to the bottom of a typical nanobar, which reveals that the nanobars are composed predominantly of YSZ and are richer in dopants at the top than at the base. Since a nanobar is formed by mounding up of material, the whole process can be thought of as one in which the transient diffusion process stacks layers of material sequentially at different points in time.16 Since surface diffusion is much faster than bulk diffusion of Ce4+ and Gd3+ at these processing temperatures, lateral compositional homogeneity (across the instantaneous top layer) can be attained. However, slower bulk diffusion kinetics essentially preserve the composition of a layer once it is buried by the mounding of additional material, hence preventing the nanobars from attaining compositional homogeneity in the depth direction. Also, since solute concentration at a point from the source increases with time, the nanobar top
bars around the patch edge have their long axes parallel to a particular direction in the YSZ-(110) surface. Second, a band of ∼5 μm of a connected porous morphology is visible near the patch edges, which is a result of excessive GDC doping due to proximity to the source. To avoid the complexity of liftoff photolithography, we have developed a simplified process of producing nanostructure arrays on YSZ surfaces.16 The solid dopant sources are dispersed on the surface as a droplet of aqueous powder suspension delivered via an eyedropper. The sample is allowed to dry naturally in air, leaving a residue of powder particles randomly distributed on the substrate surface. Subsequent annealing was performed in a high temperature furnace at 1100−1200 °C in air. The result of such a heat treatment at 1200 °C for 25 h is shown in Figure 3. A wide-area scan of the annealed sample is shown in Figure 3a, showing an area of 15 × 10 μm2 fully covered with nanobars. The footprints of dopant powder particles are also visible. In this scan area, the source particles were totally consumed during the process. Figure 3b is a higher magnification SEM image to show the nanobar morphology. The bars are ∼140 nm wide and several microns long with some of them running from top to bottom of the image. Low magnification images reveal nanobars that are in excess of 10 μm long, which results in an aspect ratio (length/ width) of ∼70. Since the dopant sources in and around the scan area of Figure 3a are small and far apart, the bars are expected to have a lower dopant concentration, hence their continuity. Furthermore, as also evident in Figure 2, there are several defects in the pattern that are a result of nonuniform doping, which causes several nanobar colonies to nucleate and grow at the same time, and later different colonies impinge on one 2261
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Figure 5. (a) SEM micrograph showing nanobars produced by using controlled powder method. The sample was annealed at 1200 °C for 25 h in air. (b) SEM micrograph of the nanobars at a different location on the same sample showing the angular dependence of the pattern extent. (c) Polar plots of measured square extent X2 as a function of orientation angle, the solid curves are fits to eq 1. (d) Growth kinetics of maximum and minimum extents.
directions suggests that the 1D criterion of the problem (assuming alignment along the most compliant direction in the surface) is insufficient as the nanobar strain fields interact in 3D through the substrate, which essentially makes it a 3D problem. Tersoff et al. have suggested that stress anisotropy can drive the alignment of the nanostructures perpendicular to the direction of maximum stress.17 This argument might be applicable since the bars run perpendicular to the direction of highest modulus, [001]. The nanobar morphology reported in the paper is not limited to GDC as a dopant source material. Other dopant sources can also be used to produce similar results.16 This is testament to the robustness of the powder process that allows exploratory studies of this kind by using commercially available single crystal substrates and oxide powders. To demonstrate this robustness, we dispersed pure Eu2O3 (EO) powder on YSZ-(110) surface which upon annealing at 1200 °C resulted in nanobar arrays similar to the ones obtained with GDC as dopant. This behavior is illustrated in the SEM micrographs shown in Figure 4a and b. Figure 4c shows an area on the sample surface far away from a dopant source. The shallow surface ripples capture the onset of nanobar formation at low dopant concentrations. It is clear from the SEM micrograph that the nanostructures start out as bars rather than small islands that evolve into a bar morphology such as those
is richer in solute than the base because it is formed later in the process (when the solute concentration is higher), while the base, which was formed when the solute concentration was lower, is buried and shielded from solute enrichment. Zhang et al. grew epitaxial In2O3 nanorods with (110) tops and {111} side facets on YSZ-(110) substrates.25 They argued that surface energy minimization alone drives self-organization of In2O3 nanorods, as the energy of the {111} surfaces in In2O3 is significantly smaller than that of the {100} and {110} surfaces. We believe that in our case, the driving force for epitaxial nanostructure formation on YSZ surfaces is the reduction in strain energy. Although the hierarchy of surface energies in YSZ is γ(111) < γ(110) < γ(100),25 the smoothly curved nanobar morphology even after a 25 h anneal at 1200 °C suggests that surface energy minimization alone cannot account for the 1D shape and strong alignment along the [11̅0] direction. YSZ has a strong intrinsic capacity to facilitate nanostructure alignment due to its strong elastic modulus anisotropy. The elastic modulus has a maximum value in ⟨100⟩ directions and a minimum in ⟨111⟩ directions.54 On YSZ-(001) surfaces the nanostructures align along the ⟨110⟩ directions,15,16 which are the lowest-modulus directions in that surface. However, in the case of YSZ-(110) surfaces, two ⟨111⟩ directions are available apart from one [11̅0] direction. The alignment of the nanobar long axes along the [11̅0] direction instead of the ⟨111⟩ 2262
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ACS Nano reported in the GDC/LAO system.23 It is pertinent to mention here that the dimensions of the nanobars are a function of time, temperature and proximity to the source for a particular dopant. For example, in the GDC/YSZ system with the powder method, an average nanobar width of ∼140 nm was observed at 1200 °C while that at 1125 °C was ∼60 nm for 25 h anneals. A similar trend was observed for the nanobars in the EO/YSZ system where a 25 h anneal at 1200 °C yielded nanobars with average width of ∼120 nm while that at 1100 °C was ∼80 nm. To perform an in-depth investigation of the kinetics of the morphological evolution, we applied the powder source in a more controlled manner. Instead of using an aqueous suspension in an eyedropper, we delivered a concentrated source ink onto the substrate via a small capillary tube. The dried ink droplet forms a dense patch and serves as a single isolated dopant source that allows measurement of the spreading of the ripple pattern as a function of time and inplane orientation. This capillary delivery method circumvented the manufacturing difficulties of lithography and making a sputtering target of the source material, but it maintained the simplicity and scalability of the powder process. In order to obtain reasonable continuity and sharpness of the patch edge, EO nanopowder was used to make the ink. Figure 5a shows an example of the morphology of the nanobars prepared by this method. In this case, the source edge is clearly delineated. Also, the dopant patch has a radius of curvature (∼500 μm) that is much larger than the observation scale, i.e., within the field of view of a typical SEM image the patch edge appears straight, and the concentration gradient of diffusing dopant is expected to be locally normal to the patch edge. The bars exhibit a finer scale near the source, and their period and width increase with the distance from the edge, i.e., with decreasing dopant content. Figure 5b shows another region around the dopant patch, where the edge normal (concentration gradient) makes an angle θ with the [11̅0] direction. The extent of the pattern, X, was measured as a function of time and orientation θ. Assuming anisotropic surface diffusion is the rate-limiting step, we analyze the data on the basis of the parabolic growth law, namely X2(θ,t) = K(θ)t, where t denotes the growth time, and K(θ) is a parabolic rate constant that is proportional to anisotropic surface diffusivity across the patterned surface. Therefore, for a given growth time, X2(θ) and K(θ) are expected to vary as symmetric second rank tensors, i.e., X2(θ ) = X12 cos2(θ ) + X 22 sin 2(θ )
X21
Figure 6. SEM micrograph showing the nanobar pattern transferred onto PDMS via single step molding.
pattern upon imprinting.48 Another advantage is related to the method by which these patterns are produced and imprinted on polymers. Previously, such groove/ridge patterns were created on Si by expensive lithographic processes,41,42 which require clean room facilities and skilled operators. The processing steps involved in forming these master patterns can be as simple as sprinkling of rare-earth oxide powder on YSZ-(110) surface followed by annealing in a high temperature furnace. Hence, these nanopatterned substrates or their polymer counterparts have the potential to be employed as inexpensive model systems for exploratory studies on cell response to nanotopography.
CONCLUSION We have produced self-assembled arrays of nanobars with wide coverage on YSZ-(110) substrate surface by solid-state surface doping with rare-earth oxides. The epitaxial nanobars extend along the [11̅0] direction on the surface, reaching several microns in length while maintaining uniform width on the order of 100 nm. Arrays of adjacent nanobars exhibit remarkable order in spacing and orientation, and these selforganized nanopatterns survive even after long exposures above 1000 °C. Cross sectional imaging reveals that the nanobar surfaces are smoothly undulated periodic ripples without faceting. The ripple period is subject to coarsening, but its magnitude is of the same order as estimated on the basis of the Asaro−Tiller−Grinfeld instability. The processing involved in forming these nanobar arrays is as simple as sprinkling of rareearth oxide powder on YSZ-(110) surface and annealing in a high temperature furnace. It is shown that the ripple arrays extend anisotropically by diffusion on the (110) surface, obeying the parabolic growth law, and showing principal values of the rate constant along [11̅0] (maximum) and [001] (minimum), as expected from the symmetry of the (110) surface. The simplicity of the process and the ease by which it can be scaled up to cover large surfaces by appropriately distributing source particles of various compositions on the YSZ-(110) surface is apparent. The hard ceramic rippled surfaces are shown to be effective templates for pattern transfer to an important polymeric biomaterial, and can thus serve as reusable molds for casting nanopatterned materials with widely diverse functionality.
(1)
X22
where and are the principal values. Polar plots of measured X2(θ,t) are shown in Figure 5c along with curve fits to eq 1. It is noted that the measurements are consistent with eq 1, with a maximum extent reached at θ = 0°, along [11̅0], and minimum along [001], for all annealing times. In Figure 5d, the time dependence of maximum and minimum extents confirms the parabolic growth law, with nearly an order of magnitude difference in principal values of the rate constants, Kmax/Kmin ∼ 7. This level of anisotropy implies that the time required for covering a surface of a given area could depend strongly on how the dopant sources are arrayed on it. Finally, we show an example of using EO-doped YSZ nanobar patterns as replicas for single-step molding with polymethyldisiloxane (PDMS). Figure 6 is an SEM micrograph showing the nanobar-patterned PDMS surface. One major advantage of imprinting a ripple-like pattern like this is that the resulting pattern is indistinguishable from the master pattern, unlike the nanoisland counterpart which produces an inverse 2263
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METHODS
ACKNOWLEDGMENTS This work was supported in part by the OSU Institute for Materials Research (IMR) and the Fulbright − HEC Pakistan Scholarship Program.
Lithographically Defined GDC Patches. Single crystal 8 mol % YSZ-(110) substrates were purchased from MTI Corporation (Richmond, CA). The substrates were 5.0 mm × 5.0 mm × 0.5 mm in dimensions and as supplied they were chemically polished to reduce the surface roughness below 5 Å. Liftoff photolithography was used to pattern RF-sputtered GDC patches on YSZ-(110) substrates (60W, 4 h, 5 mTorr Argon atmosphere) using a Discovery-18 DC/RF magnetron sputter deposition system (Denton Vacuum, Moorestown, NJ)). The resulting samples consisted of arrays of 50 μm × 50 μm GDC patches with 50 μm wide channels in between that had been shielded by photoresist during deposition. The sputtering target was made from 5 mol % GDC powder (purchased from Nextech Materials (now Nexceris), Lewis Center, OH) by Sputtering Target Manufacturing Co., LLC (Westerville, OH), which has since been purchased by Kurt J. Lesker Company. After dipping in acetone for photoresist removal, the samples were air-dried and then annealed at 1200 °C for 5−25 h in a high temperature furnace (box or tube) in air. Powder Suspension Method. For the powder suspension method, we used 0.1−0.3 g/L suspensions of 5 mol % GDC powders (Nextech Materials (now Nexceris), Lewis Center, OH) and similarly concentrated suspensions of Eu2O3 (Alfa Aesar, Ward Hill, MA), in distilled, deionized water. The suspensions were ultrasonicated for 5 min to yield a milky suspension then applied on the YSZ substrate surface via an eyedropper. Prior to application on the YSZ substrate, it was cut into pieces which were roughly 2.5 × 2.5 mm2. The droplet size was controlled such that the substrate was immersed completely in it and was allowed to dry in air in order to leave powder particles dispersed on the YSZ surface. Controlled Powder Method. For the controlled powder method, a concentrated Eu2O3 ink was made by adding 1 wt % Bicine (SigmaAldrich, St. Louis, MO) and 3 wt % Eu2O3 nanopowder (Alfa Aesar, Ward Hill, MA) into 30:70 wt % water:ethylene glycol solution followed by 15 min sonication. A 200 μm glass capillary was used to deliver the ink onto the substrate and a PV830 Pneumatic Picopump (World Precision Instruments, Sarasota, FL) was used to supply a small positive pressure for the ink ejection. The samples were heat treated at 1100−1200 °C for 5−25 h in a high temperature furnace in air. Scanning electron microscopy was performed on FEI Sirion FEG− SEM (Hillsboro OR, U.S.A.). Transmission electron microscopy (TEM) samples were prepared by focused ion beam (FIB, FEI Helios, FEI Company, Hillsboro OR) and imaging was performed on FEI Tecnai F-20 STEM (FEI Company, Hillsboro OR). Nanoimprinting on PDMS. A YSZ-(110) substrate patterned with Eu2O3 doped YSZ nanobars was sonicated for 5 min and air-dried afterward to remove the powder residue. For pattern transfer, Sylgard 184 polydimethylsiloxane (PDMS, Dow Corning, Midland, MI) in a weight ratio of 10:1 PDMS to curing agent was used. After thorough mixing of the PDMS and curing agent, the mixture was put in a vacuum chamber for 10 min to remove air bubbles. The YSZ substrate was placed face-up in a Petri dish and PDMS mixture was poured into it until the entire substrate was submerged. For better contact between PDMS and YSZ, the Petri dish was put in the vacuum chamber again for 10 min to remove trapped air at the interface. The Petri dish was later placed in an oven at 60 °C for 3 h to cure the PDMS. After curing, the PDMS was peeled off from the substrate and gold coated for SEM characterization.
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AUTHOR INFORMATION Corresponding Author
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[email protected]. Tel: +1-614-292-6725. Fax: +1-614292-1537. ORCID
Haris M. Ansari: 0000-0002-7321-5358 Sheikh A. Akbar: 0000-0003-3567-274X Notes
The authors declare no competing financial interest. 2264
DOI: 10.1021/acsnano.7b00081 ACS Nano 2017, 11, 2257−2265
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ACS Nano
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