Nanotribological Printing: A Nanoscale Additive Manufacturing Method

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Nanotribological Printing: A Nanoscale Additive Manufacturing Method Harmandeep S Khare, Nitya Nand Gosvami, Imène Lahouij, Zachary Milne, John Brandon McClimon, and Robert W Carpick Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b02505 • Publication Date (Web): 16 Oct 2018 Downloaded from http://pubs.acs.org on October 19, 2018

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Nanotribological Printing: A Nanoscale Additive Manufacturing Method H.S. Khare1§, N.N. Gosvami1†, I. Lahouij1#, Z.B. Milne1, J.B. McClimon2, R.W. Carpick1* 1 Department

of Mechanical Engineering & Applied Mechanics of Materials Science & Engineering University of Pennsylvania, Philadelphia, PA 19104 2 Department

Abstract Additive manufacturing methods are transforming the way components and devices are fabricated, which in turn is opening up completely new vistas for conceiving and designing products and engineered systems. Small-scale (sub-micrometer) additive manufacturing methods are largely in their infancy. While a number of methods exist, a particular challenge lies in finding methods that can produce a range of materials while obtaining sufficiently robust mechanical properties. In this paper, we describe a novel nanoscale additive manufacturing technique deemed “Nanotribological Printing” (NTP), which creates structures through tribomechanical and tribochemical surface interactions at the contact between a substrate and an atomic force microscope probe, where material pattern formation is driven by normal and shear contact stresses. The “ink” consists of nanoparticles or molecules dispersed in a carrier fluid surrounding the AFM probe, which are entrained into the contact during sliding. Being stress-driven, patterning only occurs locally within regions which experience contact and sufficiently high stresses. Thus, imaging and measurement to characterize the morphology and properties of the deposited structures can be conducted in-situ during the manufacturing process. Moreover, using local mechanical energy as the kinetic driver activating the solidification process, the method is compact and does not require application of a bias voltage or laser exposure, and can be performed at ambient temperatures. We demonstrate: (1) control of pattern dimensions with sub-100 nm lateral and sub-5 nm thickness control through variations in contact size and applied stress, (2) creation of amorphous, polycrystalline, and nanocomposite structures including sequential multimaterial deposition, and (3) formation of manufactured structures which exhibit mechanical properties approaching those of bulk counterparts. The ability to create nanoscale patterns using standard AFM cantilever probes and operation modes (contact mode scanning in fluid) with commercial AFM instruments, independent of substrate, establishes NTP as a versatile and easily accessible method for nanoscale additive manufacturing. Keywords: Scanning Nanomanufacturing

Probe

Lithography,

AFM,

Additive

* corresponding author Robert W. Carpick, Ph.D. Department of Mechanical Engineering and Applied Mechanics University of Pennsylvania Philadelphia, PA 19104 [email protected] (215) 898-4608 1 ACS Paragon Plus Environment

Manufacturing,

Nanolithography,

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current address Harmandeep S. Khare, Ph.D. Department of Mechanical Engineering Gonzaga University 502 E. Boone Ave., SEAS AD Box 26 Spokane, WA 99258 §

†

current address Nitya N. Gosvami, Ph.D. Block-III, Room No. 235 Department of Applied Mechanics Indian Institute of Technology Delhi Hauz Khas, New Delhi 110 016, India current address Imène Lahouij, Ph.D. MINES ParisTech, PSL – Research University, CEMEF – Centre de mise en forme des matériaux, CNRS UMR 7635, CS 10207 rue Claude Daunesse, 06904, Sophia Antipolis Cedex, France #

Graphical Abstract:

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Advances in nanomanufacturing and nanofabrication have enabled the proliferation of novel nano and microscale systems and devices, which owe their unique physical, chemical, mechanical and electronic properties to the small scale of materials and structures comprising them.1-3 Top-down methods for nanofabrication, including techniques such as photolithography and electron beam lithography are now widely used for patterning structures with a sub-20 nm and sub-5 nm resolution, respectively.4-7 Current photolithography methods provide a viable route for scalable nanomanufacturing. However, methods such as electron-beam lithography and focused-ion beam printing work in a serial writing mode and are expensive. In addition, these top-down nanomanufacturing methods are intrinsically planar, requiring multiple postprocessing steps which limit patterning flexibility and result in a large excess of material waste during manufacturing.7 Additive nanomanufacturing provides design flexibility and allows direct write patterning of metallic, ceramic and insulating nanostructures, which are typically precluded with photolithography.3

Additive

nanomanufacturing

methods

such

as

dip-pen

lithography,

electrohydrodynamic jet printing, and direct laser writing have been used for patterning ordered molecules as well as metallic, polymeric, and ceramic nanostructures.3, 8-16 Scanning probe microscopy (SPM) based nanomanufacturing is a class of nanoscale manufacturing methods which employ a scanning nanoscale probe for direct-write patterning (i.e. patterning in a single step) of nanostructures.17-22 More broadly, SPM methods allow direct-write patterning and in-situ imaging by contact or near-contact interactions of a substrate with a nanoscale probe, typically at the end of an atomic force microscope (AFM) cantilever.23-25 Each particular SPM technique can be classified as additive or a subtractive. However, all SPM methods exhibit direct-write characteristics, unlike photolithography. Subtractive SPM patterning occurs when high contact stresses or temperatures at the AFM probe apex induce wear of the substrate, either mechanically or thermomechanically.26 Additive patterning occurs when interactions of the probe with the substrate result in surface modification or net-material addition, typically due to electrochemical or thermochemical reactions in response to an applied electric field or at elevated temperatures.3, 26-31 Subtractive SPM necessarily requires the starting bulk workpiece to be machinable and finished patterns, created through material removal, retain chemical identity of the starting workpiece. Therefore, subtractive SPM methods can be used to pattern resists and masks that easily wear at suitable temperatures and/or stresses (behavior often exhibited by soft polymers), whereas additive SPM/SPL can in principle be used for patterning thermally stable materials with high hardness and stiffness. Despite its advantages, the ability of current SPM methods to generate such mechanically robust nanoscale patterns is limited. In the present work, we introduce nanotribological printing (NTP), a novel additive manufacturing technique for additively creating nanoscale structures and patterns with tunable microstructure, thickness, 3 ACS Paragon Plus Environment

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and lateral dimensions. While NTP may exhibit characteristics of certain other SPL methods (e.g. thermochemical patterning), it is an additive direct-write nanomanufacturing technique which fundamentally relies on applied tribological stresses for initiating tribomechanical or tribochemical interactions. In NTP, patterning is achieved through tribomechanical and tribochemical interactions driven by contact stress (normal and shear), at the interface of an AFM probe sliding in contact with a substrate. Ink units – molecules or discrete nanoparticles – are dispersed in a suitable fluid within which the AFM probe is submerged. This is typically achieved by using an AFM fluid cell, as illustrated in Fig. 1 (a). The choice of liquid carrier (particularly, whether to use a polar or non-polar liquid) is determined primarily by the solubility and dispersion stability of ink units. During sliding, ink units (molecules or nanoparticles) are entrained into the contact at the leading edge of the AFM probe, as illustrated in Fig. 1 (b). Once within the contact, applied normal and shear stresses result in surface-mediated tribochemical or tribomechanical interactions between multiple ink units, as well as ink units and the substrate, leading to the formation of stable surface-bound solidified structures.

Figure 1 (a) Schematic of the NTP process on a substrate, using an AFM microcantilever submerged in an inkcontaining fluid carrier in the fluid cell, (b) Schematic illustration of the entrainment of dispersed ink units at the leading edge of the sliding AFM probe, resulting in the stress-assisted formation of robust surface-bound solid structures on the substrate.

Here

we

demonstrate

the

versatility

of

NTP

by

printing

with

two

different

“inks”: molecules dissolved in carrier fluid that undergo tribochemical reactions to form a solid phase, and solid nanoparticles suspended in a carrier fluid that undergo tribomechanical interactions to form a solid phase. For both inks, polyalphaolefin (PAO) oil is used as the carrier fluid. Tribochemical Nanotribological Printing using ZDDP Tribochemical NTP refers to the use of tribochemically reactive ink molecules as the source material, where applied stresses kinetically activate chemical reactions which result in the formation of surface-bound structures. This is inspired by the known formation of surface-bound tribofilms derived from 4 ACS Paragon Plus Environment

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zinc dialkyl dithophophate (ZDDP) molecules.32 ZDDP molecules, when subjected to shear and compression at an interface, form solid surface-bound films with chemically and mechanically graded properties along their depth; they are used commercially in engine lubricants as the films they form protect the underlying surfaces from wear.33-35 More recent studies at the micro and nanoscale show that growth of ZDDP-derived surface films is thermally activated, and accelerated in response to applied tribological stresses (normal and shear) at the interface.36, 37 The mechanism of film deposition is understood to follow from transition state theory, where applied stresses perform mechanical work that reduce the activation energy barrier(s) to reach the transition state(s) between the chemical reactants and products. Stressactivated transition state theory has been successfully used to describe the formation of reaction products in a number of tribological systems.38-42 Following from transition state theory, an increase in interfacial temperature or applied stress can increase the reaction kinetics, accelerating the tribochemical reaction rates. In particular, growth rates of ZDDP-derived surface films were previously shown to increase exponentially with an increase in temperature and applied normal (compressive) contact stress,37 although shear stresses are present too and likely play a strong role.36 These ZDDP-derived surface films are known to consist of an amorphous polyphosphate glass structure, with its composition and mechanical properties varying through the depth of the film.32, 35 To demonstrate NTP with ZDDP, ZDDP molecules (chemical formula: [(RO)2PS2]2Zn, Fig. 2(a)) were added at a concentration of 0.8% by weight to a polyalphaolefin (PAO4) synthetic oil. The specific ZDDP used is a secondary ZDDP where the pendant alkyl groups are of the form -CH(CxHy)(CmHn) (LubrizolTM LZ1371). The substrate and oil bath were heated to 100°C. In addition to thermally accelerating the growth rate, it is known that the mechanical properties of macroscopic ZDDP-derived tribofilms are significantly improved (higher modulus and hardness) at higher temperatures, typically above at least 50°C depending on the specific type of ZDDP.32 Two patterning experiments were performed with ZDDP using probes of different apex geometries. In a first series of experiments, a steel microsphere approximately 52 µm in diameter was attached to a stiff (tapping-mode) AFM cantilever (nominal stiffness: 55 N/m), and a piece from a polished silicon (100) wafer was used as the substrate. A custom MATLABTM script with sequential commands for varying scan size, speed and angle was used to automate scanning in a commercial AFM, operated in contact mode. The resulting ZDDP-derived pattern of the word “RESEARCH”, has an average thickness of 8 nm, and is shown in Fig. 2 (b). This was produced by creating a series of line segments at different locations and orientations. Each patterned segment was subject to 200 reciprocations at 11 µm/s, under a nominal contact stress of 410 MPa; the entire pattern being formed in approximately 4 hours. The deposited ZDDP-derived pattern is robust, showing no degradation or measurable wear after subsequent solvent rinse cycles, and multiple tapping and contact mode imaging over the course of > 3 months (Fig. 2 (b)). 5 ACS Paragon Plus Environment

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Figure 2 (a) Molecular structure of ZDDP, showing organo-thiophosphate coordination with zinc, with R representing alkyl or aryl groups, (b) NTP performed with ZDDP as the ink dispersed in a synthetic oil as the carrier liquid, using a steel microsphere attached to an AFM microcantilever, and preprogrammed tip motion, (c) NTP performed with ZDDP dispersed in synthetic oil using a sharp DLC-coated AFM probe: 0° and 90° refer to direction of the fast-scan relative to the projection of the cantilever’s longitudinal axis onto the sample plane (parallel and perpendicular, respectively), (d) and (e) Corresponding cross-sections for 0° and 90° scan direction respectively show that the pattern thickness increases with increasing contact stress. Corresponding FWHM values for each patterned line segment are shown in parentheses.

A second experiment was conducted with a pyramidal silicon AFM probe coated with a 15 nm layer of diamond-like carbon (DLC) at the end of a 225 µm long microcantilever (nominal stiffness: 1.5 N/m), also using a silicon substrate in a 0.8 wt.% dispersion of ZDDP in PAO4 at 100°C (Fig. 2 (c)). DLC (specifically, tetrahedral-amorphous carbon, ta-C) coated AFM probes generally offer significantly higher wear resistance than corresponding silicon probes.43 Contact mode scans were performed by sliding the probe in a direction parallel (0°) or perpendicular (90°) to the projection of the cantilever’s long axis onto the surface plane. To generate single-line patterns of minimal transverse width, the AFM slow scan axis was disabled for these measurements. Successive line scans in each direction were performed for increasing values of applied normal load by manually offsetting the AFM probe by 2 µm. Each line segment shown 6 ACS Paragon Plus Environment

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in Fig. 2(c) was subject to 200 reciprocation cycles. As shown in Fig. 2 (c), ZDDP-derived patterns increase in thickness with increasing normal load for both scan angles, varying between 1 nm and 6 nm. The average surface roughness (Ra ) of these patterned lines was less than 1 nm for all applied loads and in both scan directions. Fig. 2 (d) and (e) show the corresponding values of pattern full-width at half maximum (FWHM) for the 0° and 90° scan directions, respectively. Since AFM lateral positioning typically has a resolution < 1 nm, patterning resolution in NTP is limited only by the pattern FWHM, which may vary depending on applied stress and ink units being used. For ZDDP, we estimate the lateral pattern resolution to be of the order of 200-400 nm. Since some thermal and electronic drift is inevitable in the AFM, these values represent an upper bound for the true lateral resolution of the method, which is likely to be somewhat lower. Interestingly, pattern thickness and values for patterns created at 0° are measurably greater than values for 90° for nominally similar normal loads. For example, for scans at 0° and 627 nN, measured FWHM and height are approximately 426 nm and 4.1 nm, respectively, whereas for the same load at 90°, FWHM and height are measured to be approximately 315 nm and 2.8 nm, respectively. It is plausible this occurs due to rotational asymmetry of the AFM probe, which would yield different contact widths when sliding along 0° and 90°. As well, when sliding along 0°, the AFM cantilever experiences an additional normal bending mode due to the frictional moment at the probe apex; such frictional coupling can further increase the apparent flexural stiffness. We hypothesize that the observed differences in pattern quality may occur due to some combination of tip asymmetry and friction coupling in the 0° direction. Patterns in Fig. 2 illustrate how variations in probe shape, size and applied normal load can be used to create nanoscale patterns with varying vertical and lateral dimensions, as well as lateral resolution. Mechanical properties (modulus and hardness) of a similar NTP-generated ZDDP pattern measuring 5 µm  5 µm were measured using instrumented nanoindentation. ZDDP pattern hardness was found to be 3.7±1.8 GPa, whereas the reduced modulus was measured to be 83.8±29.8 GPa. Measured modulus and hardness are within 20% of values reported in the literature for a broad range of ZDDP-derived surface films generated at the macroscale. This illustrates the ability of NTP to generate patterns with mechanical properties that are similar to those of macroscopically-generated structures.35, 44


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Figure 3 A ZDDP-derived pattern (top) was generated on a silicon substrate to determine the effect of stressassisted patterning on wear of a DLC-coated AFM probe. Corresponding TEM images before and after patterning (bottom) show that although the DLC-coating provides some wear protection, some flattening of the probe apex inevitably does occur, as seen in the after-patterning micrograph.

Since ZDDP tribofilm formation is stress-driven, reducing applied normal load (and therefore, normal and shear stresses) reduces the rate of pattern growth, which increases patterning time. In addition to being undesirable from a process standpoint, longer patterning times (e.g., on the order of several hours) also increase the likelihood of electronic and thermal drift in the AFM, which can result in poorer patterning fidelity. An increase in stress during patterning reduces patterning time, but also increases the likelihood of probe wear, which results in an increase in probe radius over time, reducing lateral resolution. An optimal patterning load therefore balances patterning timescales with probe wear and the resulting changes in probe radius. To estimate probe wear during patterning, a ZDDP-derived pattern was generated using a DLCcoated AFM probe on silicon, at 140 °C. A custom MATLABTM script was used to pattern the letters ‘PENN’, as shown in Fig. 3. Each linear segment in the pattern was subject to 300 reciprocations at 60 µm/s. Small probe apex radii (typically ca. 7 nm) of as-fabricated DLC probes can result in extremely high initial contact stresses when the probe first makes contact with a substrate, increasing the likelihood of probe wear and fracture. To distinguish probe wear due to patterning from wear due to the initial contact with a substrate, DLC probes were pre-worn by sliding 100 reciprocations at 90 nN and 25°C in base PAO4, and subsequently imaged in the TEM, before patterning (referred to as ‘before patterning’ in Fig. 3). From a number of DLC-coated AFM probes otherwise analyzed in the TEM after patterning with ZDDP, there is 8 ACS Paragon Plus Environment

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typically little evidence of a ZDDP-derived tribofilm forming on the probe apex. However, to minimize confounding effects of adsorbed ZDDP thermal decomposition products that may still form on the apex, the DLC probe used for patterning in Fig. 3 was rinsed for 10 min. in an ethylenediaminetetraacetic acid (EDTA) solution (NaOH buffered to pH 9), post-patterning. TEM imaging was again performed to determine the extent of probe wear due to patterning. TEM micrographs of the DLC probe before and after patterning are shown in Fig. 3. The after-patterning probe represents cumulative sliding equivalent to 4200 reciprocations at a nominal load of 90 nN, with a nominal speed of 60 µm/s. A residual surface layer (presumably a combination of remnant thermally degraded reaction products and oil residue) and the 15 nm thick DLC coating are clearly seen in both micrographs. Although the probe apex is somewhat flattened in the post-pattern image, some DLC, approximately 15 nm thick, remains and still separates the probe apex (i.e. the region of contact) from the underlying silicon. The images in Fig. 3 illustrate that although AFM probe wear occurs during sliding, there is potential for it to be regulated by suitably reducing contact stress, in addition to using wear-resistant AFM probes (such as DLC, ultra-nanocrystalline diamond, and single-crystal diamond).

Tribomechanical Nanotribological Printing using ZrO2 nanoparticles To explore the ability of tribological stresses to pattern nanostructures, ZrO2 nanoparticles (Pixclear PC14-10-L01, Pixelligent LLC, Baltimore, MD) were dispersed in PAO4 and evaluated as a metal-oxide ink system for creating polycrystalline nanoscale patterns of zirconia. These commercially available nanoparticles are highly monodisperse with an average diameter of 5 nm45. Metal-oxide nanoparticles, including ZrO2 nanoparticles have previously been shown to form surface films due to tribological action at the macroscale, believed to be through a process of in-situ sintering.46, 47 These nanoparticles are well dispersed in PAO with no evidence of settling out after extended periods of time (> 4 years). To ensure adequate entrainment of sufficiently large number of nanoparticles for patterning, patterning with ZrO2 nanoparticles was performed with steel microspheres attached to AFM cantilevers; attempts to form patterns with regular AFM tips resulted in slow deposition rates of incomplete films. The microspheres were attached to tapping mode cantilevers whose higher force constants (typically 40 N/m compared to 0.01 - 3 N/m for contact mode cantilevers) were needed to apply sufficiently high contact stresses to produce patterns at appreciable rates. Larger microspheres ensure the relative size of the nominal Hertzian contact diameter with a microsphere (100 nm) was larger than the measured diameter of a single ZrO2 nanoparticle (5 nm). ZrO2 nanoparticles were added to PAO4 in a concentration of 10 wt.%. This relatively high concentration was chosen to expedite patterning. The AFM setup for patterning with ZrO2 was otherwise


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nominally identical to that used for patterning with ZDDP, with the exception that ZrO2 patterning was performed at 25°C. Fig. 4 (a) shows ZrO2 patterns generated on a polished 52100 stainless steel substrate using a 40 µm diameter steel probe at an applied normal load of 50 µN (equivalent to a nominal contact stress of 457 MPa according to the Hertz contact model). A central rectangular pattern was generated by raster scanning the probe for 1000 scan cycles at 244 µm/s, in an area measuring 10 µm × 3 µm, and then two separate line patterns 10 µm long were generated by disabling displacement along the AFM slow-scan axis. Fig 4 (b) and (c) show representative cross-sectional profiles of the three patterns along two axes: longitudinally through the central pattern (B-B’) and transversely through the three patterns (A-A’). From section A-A’, The FWHM of the single-track, shoulder pattern is measured as approximately 254 nm (the width at the base of the shoulder pattern is measured to be 524 nm). As with patterns created with ZDDP ink particles, there is a strong likelihood that thermal and electronic drift during ZrO2 patterning will yield FWHM values which are somewhat higher than the true lateral resolution of the method. In contrast, for the applied load and probe diameter, the Hertzian contact diameter is estimated to be 452 nm. The Hertzian estimate assumes a perfectly spherical probe apex with no local surface roughness; actual probes are neither perfectly round nor smooth. Together with the fact that stress distribution within a Hertzian contact results in stresses near the center of contact appreciably higher than at the contact edges, we postulate that an estimate of the Hertzian contact diameter can be used as an approximate measure for the expected single-line pattern widths. Section B-B’, shown in Fig. 4 (c) reveals that the mean height of the central ZrO2 pattern is 81 nm, with an average roughness (Ra) of 9 nm. Interestingly, for other ZrO2 patterns with varying thicknesses (varied by modulating contact stress), the measured average surface roughness is always measured to be less than 12% of the mean pattern thickness. For this central ZrO2 pattern, the sidewall average roughness was evaluated for the two edges along the AFM fast scan axis, as shown in Fig. 4 (d), and were measured to be 57.2 nm and 63.8 nm. It is worth noting that sliding parameters (speed, stress, contact radius, etc.) were not optimized for minimizing pattern roughness.


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Figure 4 (a) A central ZrO2 pattern measuring 10 µm × 3 µm, and two single-line ZrO2 patterns were generated using a colloidal steel probe on a 52100 steel substrate, (b) A cross-sectional profile across the three patterns (section A-A’) shows the single-line patterns have an approximate FWHM of 254 nm, (c) Cross-sectional profile along the central pattern (section B-B’) reveals a nominal pattern thickness of 81 nm, with surface averaged roughness (Ra) of 9 nm, (d) Sidewall average roughness (Ra) of the central pattern was measured to be 57.2 nm and 63.8 nm along the two fast-scan edges.

To examine the microstructure of ZrO2 patterns, printed ZrO2 on steel was sectioned using a standard focused ion beam (FIB) milling procedure to generate an approximately 100 nm thin crosssectional lamella, which was subsequently imaged in a transmission electron microscopy (TEM). Fig. 5 shows TEM micrographs of representative ZrO2 tribofilm cross-sections, with varying magnifications. These micrographs show that NTP-patterned ZrO2 has a highly dense, polycrystalline microstructure. Subsequent electron dispersive spectroscopy (EDS, not shown) and micrograph Fourier transform (Fig. 5, FFT inset) confirm that patterns comprised nearly entirely of ZrO2 and are polycrystalline. TEM micrographs in Fig. 5 reveal that the transition from steel substrate to ZrO2 film is relatively abrupt, and any mixing which may have occurred between ZrO2 and steel is limited to an interlayer which is thinner than 2-4 nm. The nearly abrupt transition from steel to ZrO2 illustrates that the patterned substrate undergoes negligible roughening, wear, or damage during patterning, and demonstrates a high level of precision and process control with NTP. Interestingly, several grains in the TEM micrographs are found to be larger than 5 nm, the measured diameter of discrete nanoparticles.45 Separate TEM-based EDS studies show that the tribofilm is nearly free of carbon, demonstrating that the organic capping agents have likely been removed; these results will be presented in an upcoming manuscript. We hypothesize that a high specific surface energy of these 5 nm diameter nanoparticles and the reactive nature of the ligand-free ZrO2 surface, together with applied compressive and shear stresses, results in grain coarsening through grain boundary migration 11 ACS Paragon Plus Environment

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and surface diffusion of Zr and O. Prior studies support both the size scale-driven reduction in temperatures and stresses required for nanoparticle sintering, as well as surface-diffusion driven coalescence and grain coarsening for ZrO248, 49. Although porosity of the ZrO2 pattern was not quantified, TEM micrographs from several FIB sections from several tribofilm growth experiments did not reveal the existence of any pores or voids. To probe the structural integrity of patterned ZrO2, mechanical properties were measured using nanoindentation. For ZrO2 patterns generated from ZrO2 nanoparticles, which are known to be a combination of tetragonal and cubic phases, the Young’s modulus and nanoindentation hardness were measured to be 154 ± 5 GPa and 7.3 ± 0.7 GPa, respectively, assuming a Poisson’s ratio of 0.3 for zirconia.50 For comparison, Young’s modulus of nanocrystalline monoclinic ZrO2, sintered at 1100°C with a theoretical density of 92% was measured to have a Young’s modulus of 199 ± 2 GPa51 and hardness values ranging from 4.1 – 9.2 GPa52-54; single crystal monoclinic ZrO2 was measured to have an orientationdependent Young’s modulus varying between 200 - 350 GPa and a hardness of 6.6 GPa.55 It should be noted that in bulk form, ZrO2 is known to exist only in its monoclinic form at temperatures below 1170°C and at pressures below 3 GPa.56

Figure 5 TEM micrographs of a 100 nm thick lamella of ZrO2 nanoparticle-derived tribofilm, extracted from an AFM-generated ZrO2 pattern using FIB milling. The image reveals a highly dense, polycrystalline microstructure, free of voids. A clearly defined transition is also observed at the steel-zirconia interface.

Multimaterial Nanotribological Printing The ability to create ZrO2-derived and ZDDP-derived nanoscale patterns can be combined and applied to create more complex, multi-component patterns. These include layered heterostructures, where each layer has a distinct composition, as well as composite nanostructures where two or more material chemistries are mixed in controllable loading fractions.


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Fig. 6 (a) illustrates patterning of layers sequentially comprised of a ZrO2 and a ZDDP-derived film. Steel colloidal probes of diameter 52 µm were initially run against polished 52100 stainless steel substrates at 25°C in order to create the ZrO2 patterns, again using a 10 wt.% dispersion of ZrO2 in PAO4. After ZrO2 patterning, the AFM fluid cell was flushed and rinsed with hexanes, and subsequently filled with PAO4 containing 0.8 wt.% ZDDP. ZDDP patterns were created at 100°C, with the same colloidal probe, on top of the ZrO2, as shown. ZDDP patterning is readily supported on a base layer of patterned ZrO2 since the modulus of ZrO2 is measured to be higher than typical values associated with ZDDP-derived films and is therefore able to allow sufficiently high contact stresses to be generated. In contrast, for the conditions we tested, a ZDDP base layer, due to its lower modulus and hardness (as discussed above), was unable to support the contact stress required to generate ZrO2 patterns and was found to wear readily at higher contact stresses before ZrO2 nucleation could occur. This highlights a present constraint of NTP: as a stress-activated process, patterning materials on low stiffness and low hardness substrates is inhibited as the underlying material will limit the stresses that can be applied by the probe due to either elastic or plastic deformation. This could potentially be ameliorated by working at higher temperatures, where thermal energy enables the physical and/or chemical mechanisms required for film formation to occur. Fig. 6 (b) similarly shows a pattern generated with 9 wt.% ZrO2 co-mixed with 0.8 wt.% ZDDP in the same liquid dispersion. Co-mixed ZrO2 and ZDDP in varying relative concentrations have the potential to generate patterns with nanocomposite-like microstructures. Possible synergistic interactions between other pairs of ink molecules can similarly be exploited in future work.

Figure 6 (a) Layered patterns with a ZrO2 base pattern, and a ZDDP-derived overlay were generated by sequentially patterning with ZrO2-containing and ZDDP-containing PAO4, at 25°C and 100°C, respectively. Growth of ZDDPderived patterns is just as readily supported on a ZrO2 pattern as it is on a 52100 stainless steel substrate, (b) AFM topographic image of pattern generated in a dispersion containing both ZrO2 and ZDDP, dispersed in PAO4 in 9 wt.% and 0.8 wt.% fractions, respectively.


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In summary, we present nanotribological printing as a novel method for nanolithography using standard contact-mode AFM. With NTP, patterning at the nanoscale is achieved primarily due to applied normal and shear contact stresses and is accelerated for inks such as ZDDP by an increase in temperature. Standard AFM probes such as sharp pyramidal probes, as well as larger colloidal microspheres, were shown to induce patterning of amorphous and crystalline patterns on a range of substrates. Polycrystalline ZrO2 patterns exhibit stress-induced densification and grain coarsening even at ambient temperature and were found to have mechanical properties reaching at least 75% of literature values for bulk sintered zirconia. Finally, multi-component patterns were generated using ZDDP and ZrO2, through sequential and nanocomposite patterning. Future work will explore nanoscale patterning with NTP and other model ink systems, including polymer-based inks and dispersions of metallic nanoparticles (such as Cu and Au), as well as patterning within a suitably large liquid meniscus surrounding the tip-sample contact, rather than a special fluid cell. Its ability to pattern with a range of ink materials and on a variety of substrates, using only contact-mode scanning in an AFM make NTP a highly versatile and accessible technique for additive nanolithography. The results here were used with ZDDP and ZrO2 solutions available as-provided from industrial collaborators, without optimization for enhancing the rate of pattern formation. Strategies for improving patterning rates, including optimizing the tip shape, contact stress, the pendant alcohol group of the ZDDP molecules (which can strongly affect reactivity), the size of the ZrO2 nanoparticles, and the molecular structure of the capping agent on the nanoparticles, among other factors, can all potentially significantly enhance the printing rate of the structures. Explorations of these phenomena are being pursued in ongoing work. MATERIALS AND METHODS ZDDP and ZrO2 formulations ZrO2 and ZDDP inks were each individually dispersed in a polyalphaolefin-based carrier fluid (SpectraSyn PAO4, ExxonMobil Chemical Co., Spring TX), in concentrations of 10 wt.% and 0.8 wt.%, respectively. For patterns generated through co-mixed ZrO2 and ZDDP, each component was added to PAO4 in concentrations of 9 wt.% and 0.8 wt.%, respectively. Each formulation in PAO4 was treated with an additional 5 wt.% of alkylated naphthalene as co-solvent. A secondary ZDDP with pendant alkyl groups of the form -CH(CxHy)(CmHn) (LubrizolTM LZ1371, Wickliffe OH) was dispersed in PAO4 for patterning with ZDDP ink. Commercially available ZrO2 nanoparticles (Pixclear PC14-10-L01, Pixelligent LLC, Baltimore, MD) were synthesized using a solvothermal technique57 and subsequently treated and capped with an organic ligand to improve their dispersion stability in oil.45 . The size of ZrO2 nanoparticles was evaluated using transmission electron microscopy (TEM) as well as dynamic light scattering (DLS). In the 14 ACS Paragon Plus Environment

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TEM, ZrO2 nanoparticles were found to be crystalline, with an average diameter of 5 nm before capping.45 The size of the capped nanoparticles as determined by DLS was 6.6 nm with a Dv(99.99) value of 17.8 nm (Fig. 7). A DLS Dv(99.99) value represents the diameter where 99.99% of the particles in the suspension are smaller than the given value. The value obtained using DLS is larger than the corresponding diameter from TEM since DLS captures the hydrodynamic diameter of a capped nanoparticle (i.e. a crystalline core with a shell of dispersing ligands), whereas TEM captures radius of only the crystalline core. Atomic force microscopy and substrate materials AFM patterning experiments were performed on a Keysight 5500 AFM, equipped with a fluid cell and a variable temperature substrate holder (Keysight Technologies, Santa Rosa, CA). The fluid cell was filled with an oil containing the desired ink material (ZDDP molecules, ZrO2 nanoparticles, or combinations thereof). Pattern growth was conducted either using a silicon AFM probes with a tip-side 15 nm DLC coating (Multi 75 DLC, Budget Sensors, Bulgaria), or using custom-prepared steel colloidal probes. For the latter, steel microspheres of diameters ranging between 10 µm and 70 µm (NanoSteel Co., Providence, RI) were mounted on tapping mode cantilevers (TAP300-G, Budget Sensors, Bulgaria) with a two-part epoxy (JB Weld, Sulphur Springs, TX), using a custom-built micromanipulator. Normal force calibration of AFM cantilevers was performed using the Sader Method58, 59 prior to attachment of steel microspheres. An off-end flexural length, determined from the actual location of mounted probe, was used to determine the true bending stiffness for the colloid-mounted cantilever.58 This corrected bending stiffness was used to determine the applied normal load, and subsequently to estimate the contact diameter and stress using Hertzian mechanics (the Hertz model was deemed appropriate since probe-sample adhesion was measured to be negligible in the oil solution). The sliding countersurface consisted either of a disc-shaped 52100 steel coupon, hardened to Rockwell C hardness greater than 60 (Heckel Tool and Mfg. Corp., Eagle, WI) or a highly polished doped silicon (100) wafer. As-received steel coupons were polished to a nominal RMS roughness of 1.3 ± 0.1 nm using a polishing slurry. Tribofilm mechanical characterization The mechanical properties of ZrO2 and ZDDP patterns were evaluated using an AgilentTM Nano Indenter G200 (Agilent Technologies, Santa Clara, CA) in the continuous stiffness measurement (CSM) method60 and a Hysitron TI 950 TriboIndenter (Hysitron Corp., Minneapolis, MN) in depth-controlled indentation mode. In both cases, a Berkovich diamond tip was used.


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ACKNOWLEDGEMENTS J.B.M. acknowledges support from the National Science Foundation under Grant No. DMR-1107642. Use of University of Pennsylvania Nano/Bio Interface Center Instrumentation is acknowledged. The authors gratefully acknowledge K.T. Turner at the University of Pennsylvania and G. Feng at Villanova University for use of instrumentation, G.D. Cooper at Pixelligent Technologies Inc. (Baltimore, MD) for providing ZrO2 formulations, and S.L. Cravens for preparing the EDTA solution.

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29. Wang, D. B.; Kodali, V. K.; Underwood, W. D.; Jarvholm, J. E.; Okada, T.; Jones, S. C.; Rumi, M.; Dai, Z. T.; King, W. P.; Marder, S. R.; Curtis, J. E.; Riedo, E. Adv. Funct. Mater. 2009, 19, 36963702. 30. Fenwick, O.; Bozec, L.; Credgington, D.; Hammiche, A.; Lazzerini, G. M.; Silberberg, Y. R.; Cacialli, F. Nat. Nanotechnol. 2009, 4, 664-668. 31. Wei, Z. Q.; Wang, D. B.; Kim, S.; Kim, S. Y.; Hu, Y. K.; Yakes, M. K.; Laracuente, A. R.; Dai, Z. T.; Marder, S. R.; Berger, C.; King, W. P.; de Heer, W. A.; Sheehan, P. E.; Riedo, E. Science 2010, 328, 1373-1376. 32. Spikes, H. Tribol. Lett. 2004, 17, 469-489. 33. Fujita, H.; Spikes, H. A. Proc. Inst. Mech. Eng., Part J 2004, 218, 265-278. 34. Aktary, M.; McDermott, M. T.; McAlpine, G. A. Tribol. Lett. 2002, 12, 155-162. 35. Bec, S.; Tonck, A.; Georges, J. M.; Coy, R. C.; Bell, J. C.; Roper, G. W. P. Roy. Soc. Lond. A Mat. 1999, 455, 4181-4203. 36. Zhang, J.; Spikes, H. Tribol. Lett. 2016, 63. 37. Gosvami, N. N.; Bares, J. A.; Mangolini, F.; Konicek, A. R.; Yablon, D. G.; Carpick, R. W. Science 2015, 348, 102-106. 38. Makarov, D. E. J. Chem. Phys. 2016, 144, 030901. 39. Adams, H. L.; Garvey, M. T.; Ramasamy, U. S.; Ye, Z.; Martini, A.; Tysoe, W. T. J. Phys. Chem. C 2015, 119, 7115-7123. 40. Liu, J.; Jiang, Y.; Grierson, D. S.; Sridharan, K.; Shao, Y.; Jacobs, T. D. B.; Falk, M. L.; Carpick, R. W.; Turner, K. T. ACS Appl. Mater. Interfaces 2017, 9, 35341-35348. 41. Jacobs, T. D. B.; Carpick, R. W. Nat. Nanotechnol. 2013, 8, 108-112. 42. Yeon, J.; He, X.; Martini, A.; Kim, S. H. ACS Appl. Mater. Interfaces 2017, 9, 3142-3148. 43. Vahdat, V.; Ryan, K. E.; Keating, P. L.; Jiang, Y.; Adiga, S. P.; Schall, J. D.; Turner, K. T.; Harrison, J. A.; Carpick, R. W. ACS Nano 2014, 8, 7027-7040. 44. Bec, S.; Tonck, A., Nanometer Scale Mechanical Properties of Tribochemical Films. In Tribology Series, Dowson, D.; Taylor, C. M.; Childs, T. H. C.; Dalmaz, G.; Berthier, Y.; Flamand, L.; Georges, J. M.; Lubrecht, A. A., Eds. Elsevier: 1996; Vol. 31, pp 173-184. 45. Williams, Z. S. G.; Wang, Y.; Wiacek, R. J.; Bai, X.; Gou, L.; Thomas, S. I.; Xu, W.; Xu, J. Synthesis, capping and dispersion of nanocrystals. WO/2011/133228, 2011. 46. Kato, H.; Komai, K. Wear 2007, 262, 36-41. 47. Battez, A. H.; Gonzalez, R.; Viesca, J. L.; Fernandez, J. E.; Fernandez, J. M. D.; Machado, A.; Chou, R.; Riba, J. Wear 2008, 265, 422-428. 48. Herring, C. J. Appl. Phys. 1950, 21, 301-303. 49. Srdic, V. V.; Winterer, M.; Hahn, H. J. Am. Ceram. Soc. 2000, 83, 729-736. 50. Khare, H. S.; Lahouij, I.; Jackson, A.; Feng, G.; Chen, Z.; Cooper, G. D.; Carpick, R. W. 51. Eichler, J.; Eisele, U.; Rödel, J. J. Am. Ceram. Soc. 2004, 87, 1401-1403. 52. Bravo-Leon, A.; Morikawa, Y.; Kawahara, M.; Mayo, M. J. Acta Mater. 2002, 50, 4555-4562. 53. Cutler, R. A.; Reynolds, J. R.; Jones, A. J. Am. Ceram. Soc. 1992, 75, 2173-2183. 54. Graeve, O. A., Zirconia. In Ceramic and Glass Materials: Structure, Properties and Processing, Shackelford, J. F.; Doremus, R. H., Eds. Springer US: Boston, MA, 2008; pp 169-197. 55. Ingel, R.; Lewis, D.; Bender, B. A.; Rice, R. W., Physical, Microstructural and Thermomechanical Properties of ZrO2 Single Crystals. 1984; Vol. 12, p 408-414. 56. Block, S.; Da Jornada, J. A. H.; Piermarini, G. J. J. Am. Ceram. Soc. 1985, 68, 497-499. 57. Wang, X. M.; Xiao, P. J. Mater. Res. 2007, 22, 46-55. 58. Sader, J. E.; Larson, I.; Mulvaney, P.; White, L. R. Rev. Sci. Instrum. 1995, 66, 3789-3798. 59. Green, C. P.; Lioe, H.; Cleveland, J. P.; Proksch, R.; Mulvaney, P.; Sader, J. E. Rev. Sci. Instrum. 2004, 75, 1988-1996. 60. Li, X. D.; Bhushan, B. Mater. Charact. 2002, 48, 11-36.


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350.0 nm 175.0 0.0

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cantilever

carrier liquid

laser

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photodetector fluid cell

(b)

dispersed molecules or particles sliding direction AFM probe

substrate

shear stress

printed pattern ACS Paragon Plus Environment

entrainment normal stress + temperature

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(a)

RO RO

S P

S

S Zn

S

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OR P

Si substrate

OR increasing contact stress

(b)

8.0 nm 4.0

Si substrate

5 μm

878nN

(d)

6.0

0°

height [nm]

5.0

502nN (374nm)

4.0 3.0 2.0

(419nm)

0.0

12.5

4 μm

-10.0

0°

(435nm)

753nN 627nN (442nm) (426nm)

1004nN

(e)

6.0

2.0

0.0

0.0

position [µm]

9.0

12.0

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(318nm)

753nN (298nm) (217nm)

502nN (314nm)

376nN (286nm)

3.0 1.0

6.0

878nN

4.0

1.0 3.0

90°

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627nN (315nm)

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25.0 nm

ZDDP pattern

12.5 0.0

5 μm Si substrate before patterning

after patterning Si probe

Si probe

15 nm DLC coating

15 nm DLC coating

oil residue

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oil residue


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printed ZrO2

A’

B

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A

50.0

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60.0

ave. roughness = 9nm mean height 81nm

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average sidewall roughness ~ 57.2nm 75.0 nm

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average sidewall roughness ~ 63.8nm

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–

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Platinum

Zir con ia

50 nm

Steel Steel

10 nm


5 nm

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(a) steel substrate

zinc dithiophosphate on zirconia on steel

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nanoprinted zinc dithiophosphate nanoprinted zinc dithiophosphate and zirconia hybrid composite

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nanoprinted zirconia

175.0

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nanoprinted zinc dithiophosphate

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nanoprinted zirconia

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Nanotribological Printing: A Nanoscale Additive Manufacturing Method

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