Mechanical Control of Surface Adsorption by Nanoscale Cracking

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Mechanical Control of Surface Adsorption by Nanoscale Cracking

Qiang Zhuang , Scott C. Warren , Bilge Baytekin , Ahmet F. Demirörs , Pramod P. Pillai , Bartlomiej Kowalczyk , H. Tarik Baytekin , and Bartosz Grzybowski *

DOI: 10.1002/adma.201306335

between adjacent cracks must be tunable. Top-down, serial approaches [ 14 ] for initiating cracks provide a simple way to control crack spacing, but even in these cases the crack width typically varies by more than an order of magnitude. Similar variability was observed when thin coatings were cracked using uniaxial strain. [ 26–28 ] Fourth, the crack should not only allow a material to deposit but also impart some degree of selectivity. Ideally, the crack should control the size, shape, and crystallo-graphic orientation of the material that deposits. Fifth, the pro-duction of cracks in only specifi c locations of a material would be desirable. Although top-down strategies for initiating cracks one-by-one has been presented, [ 14,15 ] for many materials and applications it would be desirable to identify principles that allow the high throughput production of cracks in a spatially controlled manner.

In the present work, we fabricated a composite material that is comprised of a stretchable, [ 29–31 ] polymeric substrate cov-ered with a reactive/charged polymer onto which a monolayer of nanoparticles is adsorbed ( Figure 1 ). The nanoparticle monolayer acts like a mask and, when stretched, fragments controllably as the large, macroscopic deformation divides into the small, nanoscopic movements of adjacent particles. Because the nanoparticles are deposited as a monolayer, the process of cracking unmasks the chemically distinct substrate, thereby enabling selective surface design. Unlike in continuous materials, the cracking process is guided by the discrete nature of nanoparticles and is therefore not limited by a material’s inherent faceting and fracture characteristics. [ 14 ]

We typically used silica nanoparticles to form the mask-like monolayer (Figure 1 ) but also used other metallic (gold, silver, and platinum) and iron oxide nanoparticles either as a mask layer or as the adsorbing species. The silica particles had diameters of 17 ± 1 nm, were dispersed from aqueous solu-tions with a pH between 4 and 11 and had a silica content of 0.7 wt.%. Under these conditions the particles had negative surface charges as evidenced by zeta potential measurements: at a pH of 6.5, the zeta potential was –21.5 mV; at a pH of 11, it was –30 mV. The particles were adsorbed as a monolayer onto a positively charged substrate in a self-limiting process; [ 32 ] the adsorption rate followed a Langmuir isotherm (see Sup-plementary Information, SI). The substrate was fabricated from 80 μm-thick sheets of elastomeric poly(dimethylsiloxane) (PDMS, Sylgard 184) that were cut to a width ( W ) of 1 cm and a length ( L ) of 8 cm. The surface of PDMS was chemically modifi ed to covalently bond branched poly(ethyleneimine) (PEI, 25 kg mol –1 , M w / M n = 2.5). Based on titration data, the PEI began to acquire a positive charge at pHs below 10.5 and was fully protonated by pH 6. The PEI-functionalized PDMS

Nanoparticles located at surfaces within fl exible devices [ 1,2 ] have the potential to impart novel functions [ 3–5 ] and/or device char-acteristics. [ 6–8 ] However, great care must be taken to engineer these surfaces to prevent crack formation, which ultimately leads to device failure. [ 9 ] The desire to prevent crack forma-tion [ 10–13 ] has led to a relatively limited focus on the potential benefi ts of cracks, [ 14,15 ] especially as a useful materials design tool. This proposal—that cracks could facilitate materials design—has at its core the great variety of cracking phenomena that have been observed in granular, [ 16,17 ] polymeric, [ 18,19 ] crys-talline, [ 14,20 ] and glassy materials [ 21–23 ] but have so far not been harnessed towards the controlled functionalization of surfaces.

If cracking is to emerge as a useful design tool, we must learn how to control the formation of cracks with a high degree of reproducibility, spatial precision, and versatility. In this work, we have identifi ed and then implemented fi ve principles that enable these objectives. First, cracks must be chemically dis-tinct from the non-cracked regions. This chemical patterning could then enable spatially controlled surface design: one region is chemically or physically altered while leaving the other untouched. Second, the method should provide a means to control crack orientation. Although cracks often form in a direction perpendicular to the strain fi eld, [ 15 ] this is not uni-versally true. [ 24 ] In general, it remains unclear how to guide crack orientation in many classes of materials and a material’s microstructure often presents an obstacle to guiding crack ori-entation (e.g., when a crack follows a cleavage plane in a crystal or defect sites at grain boundaries rather than propagating in a desired direction. [ 25 ] Third, the width of a crack and spacing

Q. Zhuang,[+] Prof. S. C. Warren,[+] Dr. B. Baytekin, Dr. A. F. Demirörs, Dr. P. P. Pillai, Dr. B. Kowalczyk, Dr. H. T. Baytekin, Prof. B. Grzybowski Department of Chemistry Northwestern University Evanston , IL 60208 , USA E-mail: [email protected] Q. Zhuang, Prof. S. C. Warren, Dr. B. Baytekin, Dr. A. F. Demirörs, Dr. P. P. Pillai, Dr. B. Kowalczyk, Dr. H. T. Baytekin, Prof. B. Grzybowski Department of Chemical and Biological Engineering Northwestern University Evanston , IL 60208 , USA Q. Zhuang School of Materials Science and Engneering Northwestern Polytechnical University Xi’an 710072 , PR China [+]Equal contributions

Adv. Mater. 2014, DOI: 10.1002/adma.201306335

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was mounted onto a homemade sample holder that imparted strains (Δ L / L ) of between 0% and 45% (±3%). Initially, the sample was mounted with 0% strain and immersed in a solu-tion of silica nanoparticles for 2 hours to form a complete nan-oparticle monolayer (Figure 1 c). The sample was subsequently rinsed with water of varying pH (between 2 and 11) and was stretched in the long direction of the PDMS sheet, achieving strains of 15%, 30%, or 45% (Figure 1 e). Because the density of PDMS is conserved upon extension (the Poisson coeffi cient is 0.5), stretching simultaneously reduced the width, W , by ca. 7%, 12%, and 17%, respectively.

As described above, the fi rst and perhaps most crucial test of whether cracks can be usefully employed in material design

is to examine whether cracking exposes a chemically distinct surface. In order to perform this test, we analyzed the chem-ical composition of the top few nanometers of a series of our stretchable fi lms (Figure 1 a,c,e) using secondary ion mass spectroscopy (SIMS). The fi rst system, which was composed of an unstrained PDMS sheet with grafted PEI, had two major peaks in the SIMS spectrum at mass/charge ratios of 28 and 40 (Figure 1 b). These peaks are consistent with the presence of both PDMS and PEI. Next, adsorbing a monolayer of nega-tively charged silica nanoparticles onto the positively charged PEI caused the mass/charge peak at 40 to disappear entirely (Figure 1 d), which is consistent with a surface containing [ 28 ] Si but no PEI. Finally, upon stretching the PDMS (Δ L / L = 45% ) ,


Figure 1. Exposing “masked” reactive sites via strain-induced cracking. a,b) A stretchable substrate, PDMS, is functionalized with PEI. A SIMS analysis in (b) shows peaks that are characteristic of PDMS and PEI. c,d) Subsequently, silica nanoparticles are adsorbed as a monolayer onto the PEI. SIMS reveals that the PDMS and PEI are hidden, showing only peaks that are characteristic of silica (mass/charge = 28). e,f) Upon stretching, cracks form in the silica nanoparticle layer. SIMS shows that the PDMS and PEI are again exposed. Also see the SI for XPS data. g) The chemical structure of the branched PEI used in these experiments. h,i) Analysis of cracking immediately after applying a strain of 45% (the samples were strained horizontally). h) Bruker Inphase image of a stretched PDMS-PEI-silica surface with vertically oriented cracks. Vertical arrows identify the yellow, PEI-rich cracks; red corresponds to silica. Inphase scale = 40 mV. i) SEM image of the same sample. Vertical arrows identify the dark PEI cracks; silica nanoparticles are light.

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the mass/charge peak at 40 re-emerged, indicating a change in chemical composition (Figure 1 f). The qualitative similari-ties between the fi rst and fi nal SIMS spectra suggest that the PEI is exposed upon stretching although a quantitative analysis is not possible because SIMS intensities exhibit only a modest correlation with surface composition. The changes in surface composition were further confi rmed by x-ray photoelectron spectroscopy (XPS, see SI), which showed that the nitrogen signal from the PEI appeared only after stretching the PDMS/PEI/silica NP composite. Together, these observations provide strong evidence that a surface's composition can be signifi -cantly and easily controlled by a macroscopically applied strain.

Next, we examined the effect of the mechanical deformation on the material's surface microstructure. Because stretching increases the area of the polymeric substrate but the area occu-pied by the nanoparticles is nearly constant, we expected that nanoscale cracks would arise upon straining the material. This prediction was verifi ed by atomic force microscopy (AFM) and scanning electron microscopy (SEM) on the stretched materials. Immediately after stretching, both AFM and SEM revealed the presence of cracks that were oriented predominantly perpendicular to the stretching direction (Figure 1 h,i). Crack width increased mono-tonically with strain, reaching a crack width of ∼30 nm and a spacing between cracks of ≈160 nm for a strain (Δ L / L ) of 45%. In addi-tion to cracks, AFM revealed the presence of bucking along the direction of the stretching, which is consistent with the Poisson effect leading to a contraction in the material dimensions perpendicular to the stretching direction (see Supporting Information Figure S5 for a height image). However, comparison with SEM of the same sample (in which PEI regions appear darker, Figure 1 i) reveals that only cracks were devoid of the masking silica NPs, while the buckles remained covered with such particles. This conclusion is fur-ther confi rmed by in-phase AFM mechanical imaging whereby the harder silica regions appear red while the softer PEI regions—cor-responding to vertically running cracks in Figure 1 h—appear yellow.

With the above data in place, we began to explore whether the controlled cracking of nanoparticle monolayers could be used to control the adsorption characteristics of these surfaces. In these experiments, we used neg-atively charged, citrate-stabilized gold nano-particles of different average sizes ( Figure 2 a) that interacted with the positively charged PEI found in the cracks.

Figures 2 b and 2 c illustrate the outcome of experiments in which substrates covered with 17 nm silica NPs and strained by 0, 15, 30, and 45% were exposed to a pH = 6.5 solu-tion of 16 nm Au NPs for 18 hours. As evi-

denced by the darkening color of the substrates (Figure 2 b) as well as the increasing 525 nm SPR peak of AuNPs (Figure 2 c), the degree of adsorption was negligible for the unstrained substrate but increased approximately linearly with increased strain (note: the extinction cross section scales linearly with the density of adsorbed AuNPs for discrete, non- or weakly-inter-acting NPs). It is important to note that Au NPs covered with alkane thiols terminated in the positively charged N(CH 3 ) 3 + groups [ 33 ] showed no adsorption even for the highest strain—these control experiments confi rm that absorption is mediated by the electrostatic interactions between the AuNPs and the unmasked PEI regions and that the area of PEI available for these interactions increases with strain.

In order to better clarify the relationship between strain, crack size, and the existence of sites that enable nanoparticle adsorption, we performed a series of experiments in which we


Figure 2. Strain, size, and charge-selective modifi cation of strained surfaces. a) Size distribu-tions (by TEM measurement) of the Au NPs used in adsorption experiments. b) Photograph of PDMS-PEI-silica nanoparticle fi lms with strain-induced deposition of gold nanoparticles. The silica particles had an average diameter of 17 nm and the gold particles were 16 nm in size. Strains ranged from 0% to 45%. c) UV-Vis extinction spectra of the surfaces correspond to the samples shown in (b). Peaks at 525 nm are due to the SPR of AuNPs and the increasing intensity evidences that at larger strains, more AuNPs have deposited. d,e) To explore the size selectivity of adsorption, gold particles of four different sizes ( d gold / d silica = 0.94, 1.2, 1.5, and 2.0) were adsorbed at four different strains (0%, 15%, 30%, and 45%) for a total of 16 data points in each of the two 3D plots in (d) and (e). The experiments in d and e were identical except for the pH of nanoparticle deposition: d) pH 6.5; e) pH 8.5.

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strains applied to the fl exible substrate. The results are sum-marized in Figure 2 d and 2 e and show that the surface density of adsorbed AuNPs increases as their size decreases. Simulta-neously, adsorption increases with increasing strain, though it is limited by the relative dimensions of cracks and particles − for example, even when the sample is strained by 45%, there is almost no deposition of the largest, 34 nm ( d gold / d silica = 1.95), gold nanoparticles onto the substrate (with cracks being ≈30 nm wide, cf. above). We emphasize that this size selec-tivity cannot be explained based on the total exposed PEI area alone but is in line with the widths of individual cracks con-trolling adsorption. In this way, the cracks act as nanoscopic “gates” that selectively permit entry of some objects but not others. We also note that the effect can be further amplifi ed by increasing the pH of the AuNP solution (Figure 2 e). When pH changes, so do the interactions between and the arrangement of the silica particles on the PDMS/PEI surface—specifi cally, at higher pHs these particles tend to aggregate [ 34 ] increasing the crack size and yielding even more sites that support par-ticle adsorption, while still providing the same type of crack-gated surface reactivity (for further discussion of pH effects, see SI, Section 4).

With nanoobjects of non-spherical shapes, the cracks can control not only the degree of adsorption but also the orienta-tion of the adsorbing particles. In a proof-of-concept experiment ( Figure 3 ), we prepared a layer of 80 nm silica nanoparticles, strained the PDMS substrate by 45%, and exposed that strained material to a solution of nanorods having an average width of 20 nm and length of 60 nm. Unlike earlier samples, the silica nanoparticles adsorbed as a submonolayer and as a result, cracks of many different sizes (between 0 and 120 nm wide) and orientations were produced; additionally, even in the

strained sample there was no apparent crack orientation. The large distribution of cracks proved useful, however, because the numerous crack geometries allowed us to perform a large-scale statistical analysis of the relationship between crack orientation and nanorod orientation within a single sample (Figure 3 b,c). The AFM images revealed that nanorod deposition was guided by the cracks, leading to a common situation in which nanorods were oriented along the direction of a crack when the crack width was smaller than the nanorod length but larger than the nanorod’s width (Figure 3 c). A complete statistical anal-ysis showed that nanorods were oriented with the crack direc-tion (angles < ≈30°) only when the cracks were narrower than 60 nm. When the cracks were wider than 60 nm, any orienta-tion was possible although there was still residual preference for nanorods to be oriented along the direction of cracks. Such a preference can be understood by considering that van der Waals interactions are maximized [ 35 ] when the long edge of a nanorod touches silica particles, thereby orienting the nanorod along the crack.

Having shown that cracks can guide deposition based on particle surface chemistry, size and orientation, we sought to achieve our ultimate goal of localizing the formation of cracks to specifi c regions on a material's surface ( Figure 4 ). The prin-ciple that we have relied on is a familiar one: thin regions of a material stretch more than thick regions. By using photolithog-raphy or other fabrication tools to build polymeric materials with complex shapes, we have strategically sought to use this principle to localize crack formation to the regions where mac-roscopic strains becomes focused. As we show, this enables the functionalization of surfaces with patterns that cannot be easily achieved by other methods.

As a fi rst example, we fabricated a PDMS strip that had regions that alternated in thickness of 40 and 80 μm

(Figure 4 a–d). As before, the PDMS was func-tionalized with PEI and 17 nm silica nano-particles. When the ribbon was stretched, the strain was localized more on the thin regions (Δ L / L = 50%) than the thick regions (Δ L / L = 28%). As predicted by the data in Figure 2 d,e, this difference in strain was suffi cient to lead to a large difference in crack size and AuNP adsorption. When the stretched ribbon was exposed to 16 nm gold nanoparticles, signifi -cantly more particles were deposited over the thinner regions, as evidenced by the photo-graph in Figure 4 c and the quantifi cation of the red and green channels of the image in Figure 4 d.

With this initial experiment showing that strain can be localized to enable site-specifi c surface design, we fi nally attempted to selectively pattern a polymeric substrate presenting an array of small wells on its surface. Could only certain surfaces of the wells be decorated with nanoparticles when the substrate is fully immersed in a nanopar-ticle solution? The affi rmative answer to this question is illustrated in Figure 4 g,h, where a PDMS slab with 200 μm diameter × 200 μm


Figure 3. Guiding nano-object orientation by nanoscale cracking . a) Illustration of nanorods depositing along the direction of a crack, corresponding to a nanorod orientation close to zero degrees. b) A height AFM image of 20 × 60 nm nanorods deposited into the cracks formed by stretching a fi lm prepared from 80 nm SiO 2 nanoparticles. Circles draw attention to regions in which nanorods have oriented along the crack (orientations close to zero degrees); elsewhere in the image, nanorods adopt a wider range of orientations. The inset enlarges a region of a single nanorod within a crack. Height scale = 100 nm for both images. c) Analysis of nanorod orientation as a function of crack size. Nanorods orient along the crack direction (nanorod orientation of 0°) when the crack is narrower than the nanorod’s length but wider that the nanorod’s width. Bars sum to 100% within a single crack width/nanorod length row (120 nanorods were measured per row).

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deep wells on its surface and functionalized with PEI/17 nm silica particles was strained by 35% and soaked in a solution of 16 nm gold nanoparticles. As expected, gold deposited on the

bottoms of the wells, which were the thinnest part of the mate-rial (ca. 10 μm) and therefore most strained. More interestingly, only those walls of the well that were elongated along the strain

Figure 4. Strain-selective surface design . a,b) When a substrate of varying thickness is stretched a) the strain is localized primarily in its thinnest regions. b) Larger cracks develop in the more strained regions, and more nanoparticles can selectively deposit therein. c) A PDMS material with the geometry shown in (a) was coated with PEI and 17 nm silica particles, strained, and exposed to a solution of 16 nm gold nanoparticles. The photograph clearly evidences gold deposition in the thin regions. d) This was quantifi ed by analyzing the intensity difference between the red and green channels of the red-green-blue image; redder portions correspond to the presence of gold. This profi le is compared to a control (blue, “no deposition”) which was not exposed to gold. e–h) A 500 μm-thick PDMS sheet was patterned with cylindrical wells. Cracks formed only on those walls that were perpendicular to the strain and on the thin (≈10 μm) well bottom, which has its entire surface perpendicular to the strain. g) A top-down optical microscope image of the well structure in which gold deposition on the well bottom is observed. h) The difference between red and green channels is quantifi ed, as in d. In addition to the bottoms of the wells, the NPs adsorb on the parts of the side walls that are most strained. i–k) This is evidenced by secondary ion mass spectroscopy (SIMS) maps − over the boxed region marked in (f): of i) the total ion distribution, j) the gold ion distribution, and k) the gold ion distribution plotted as a fraction of the total ion distribution (i.e., gold ions/total ions). The scale bar in image (i) applies to all SIMS maps. In the SIMS images, low/medium/high counts are represented by black/red/yellow, respectively.

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direction resulted in additional deposition there (Figure 4 e,f), as confi rmed by spatially resolved compositional analysis using secondary ion mass spectroscopy imaging (SIMS, Figure 4 i–k). In this way, our strain-directed approach provides a conceptu-ally novel yet technologically straightforward avenue to angle-selective surface functionalization along pores or pore-like structures within stretchable materials. This last result has particular signifi cance for asymmetrically functionalizing pores, which has been an outstanding problem in the design of porous materials.

Looking forward, we believe that the simplicity of using strain to break a material's symmetry to functionalize spe-cifi c regions of a surface will allow easy access to many useful patterns such as spatially varying composition gradients. For example, a polymeric sheet fabricated in the form of a triangle could be strained in the direction of each vertex, enabling the sequential deposition of three distinct materials, each of which has a different composition gradient. Such composi-tion spreads could lead to combinatorial searches for specifi c properties of mixed nanomaterials, whether optical, catalytic, electronic, or magnetic. For now, one of the most potentially useful concepts that we have shown here is that cracks can act as controllable, chemically distinct nanoscopic “gates” that control the quantities, sizes, orientations, and chemistries of depositing nanomaterials and, ultimately, perhaps also of macromolecules.

Supporting Information Supporting Information is available from the Wiley Online Library or from the author.

Received: December 31, 2013 Revised: January 30, 2014

Published online:

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Mechanical Control of Surface Adsorption by Nanoscale Cracking