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Article Cite This: ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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Surface Nanostructuring of Ti6Al4 V Surfaces for Parylene‑C Coatings with Ultradurable Adhesion Silvan Staufert,*,† Pascal Gutzwiller,† Fajer Mushtaq,‡ and Christofer Hierold† †

Micro and Nanosystems, Department of Mechanical and Process Engineering, ETH Zurich, Tannenstrasse 3, 8092 Zurich, Switzerland ‡ Multi-Scale Robotics Lab, Department of Mechanical and Process Engineering, ETH Zurich, Tannenstrasse 3, 8092 Zurich, Switzerland S Supporting Information *

ABSTRACT: While Parylene-C is known for its excellent mechanical properties and biocompatibility, its applicability for long-term implants is limited by poor adhesion to many commonly used implant alloys. While the commonly adapted route for adhesion promotion offers significant improvement for dry adhesion, it does not ensure long-term stability in permanent contact with fluids. In this work we investigate different synthesis routes for titanium oxide nanostructures on Ti6Al4 V which serve to physically anchor Parylene to the metal surface. The obtained titanium oxide nanostructures have a range of differing morphologies and dimensions and achieve approximately a 550 larger specific surface than an ideally flat surface. The influence of different nanostructures’ characteristics on the wet-adhesion of Parylene-C was investigated by accelerated aging of Parylene at 80 °C for 72 h in Hank’s balanced salt solution (HBSS) and subsequent blisterlike adhesion testing. The aging procedure caused reference samples with polished surfaces and silane A-174 adhesion promoter to experience an 88% reduction in failure load at which the Parylene film delaminates compared to the unaged samples. Nanostructured samples performed equally well before and after aging, corresponding to an 840% failure load increase compared to the reference surfaces. These findings may allow new applications of Parylene-C for long-term applications in liquid environments, which were previously prohibited because of poor adhesion. KEYWORDS: Parylene, adhesion, blister testing, Ti6Al4 V, titanium, nanostructures, hydrothermal synthesis

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INTRODUCTION Metals such as titanium-based alloys and stainless steel have long been used to fulfill the demanding requirements of mechanical durability and corrosion resistance for long-term (>1 year) and permanent medical implants such as stents or ventricular assist devices.1−3 However, these materials cannot always sufficiently meet the specific application’s demands concerning biocompatibility and surface functionality. Thin-film polymeric coatings are sometimes used to augment implant metals in order to fulfill these essential criteria.4−7 Unfortunately, due to the large difference in mechanical properties of metals and polymers, as well as the varying degrees of adhesion strengths, such coatings tend to delaminate and reduce the overall reliability of the implant over time.4 An example of an implant application with extreme requirements in both longterm durability as well as surface functionality is found in ventricular assist devices (VADs). The mechanical complexity and required manufacturing precision of such VADs urge the use of metal as bulk material, typically titanium grade 5 (Ti6Al4V).8 Despite suboptimal hemocompatibility of titanium, VADs make only limited use of organic coatings and modification. The threat of delamination, and thereby potential © XXXX American Chemical Society

fatal failure, of such a coating is one of the prohibiting factors for their application. Parylene-C (poly(chloro-p-xylylene)) is a commonly used polymeric thin-film coating. Due to its excellent moisturebarrier properties, it is often used as thin-film insulator for electronic components, also in biomedical applications.9,10 Further, its excellent biocompatibility allows Parylene-C to be used in direct contact with tissue and bodily fluids, e.g., as a corrosion-inhibiting agent on stainless steel.11 Initially, a Parylene-C surface is bioinert and hydrophobic, rendering it effective against biofouling.12 However, a simple oxygen plasma treatment can be used to make the surface more amenable to cell growth and allow the structuring of 2D cell-scaffolds.13,14 Such an oxygen plasma treatment also allows the binding of drug-loaded, biodegradable polymer layers to its surface, hence making the implant resistant to microbial colonization.7 Beyond Parylene-C, novel types of poly-p-xylylenes promise new opportunities for biomedical applications.15,16 Parylene-C Received: January 16, 2018 Accepted: March 26, 2018

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DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials

Figure 1. (a) Illustration of the fabrication process of the die samples with suspended membrane for adhesion testing. (b) Illustration of the adhesion testing setup. (c) Typical pressure curve for adhesion testing. There is a constant increase in pressure until small defects occur in the membrane. Shortly after, the membrane ruptures.

Table 1. Synthesis Parameters and Characteristics of the Resulting Nanostructures Process none soaking hydrothermal

EG hydrothermal

Vessel

Temperature [°C]

open open autoclave autoclave autoclave autoclave autoclave

60 60 180 180 180 180 200

Solvent

NaOH conc. [M]

Time [h]

Morphology

H2 O H2 O H2 O H2 O H2 O H2 O 1:1 H2O + C2H6O2

15 5 1 1 3 3 3

8 24 3 6 3 6 6

roughness porous porous flowers flowers fibers flowersd fibers

Characteristic dimensiona [nm] 2.2 111.3 174.5 12.7 33.0 17.9 21.7 14.9

± ± ± ± ± ± ±

31.5 31.5 4.1 11.3 3.6 5.7 4.6

Structure thickness [nm] 245 857 550 1582 1086 1503 1197

± ± ± ± ± ± ±

97 177 157 731 266 531 346

SSA increaseb [m2/m2] 185.0 247.1 111.4 256.2 190.4 255.4 564.2

± ± ± ± ± ± ±

37 49 22 51 38 51 112

Acronymc ref SK-15-8 SK-5-24 HT-1-3 HT-1-6 HT-3-3 HT-3-6 EG-3-6

Surface roughness rms, pore diameter, flower petal diameter, fiber diameter for reference, porous, flower, and fiber morphology, respectively. bSSA, specific surface area. Increase of surface as compared to substrate of same weight and dimensions with an ideally flat surface. cConsists of the process method, NaOH concentration, and reaction duration. dMostly flower structures with a few, long fibers. a

further has excellent mechanical properties and is sometimes used as active element in micro electromechanical (MEMS) devices and sensor packaging.17−20 These versatile applications of Parylene-C and other poly-pxylylenes are however only viable for long-term implants, if a sufficiently robust adhesion of the film can be demonstrated. Unfortunately, the adhesion of Parylene depends strongly on the substrate material and morphology.21,22 Furthermore, even though Parylene-C is an excellent moisture barrier, in a permanently wet environment water will eventually diffuse through the polymer layer and start to degrade the interface.11,23,24 This most critical issue for long-term implant applications cannot be solved by the standard adhesion promoter (silane A-174) used for Parylene-C. Huang et al.24 proposed an alternative adhesion promotion route on silicon substrates via macroscale anchoring. Deep trenches are etched using deep reactive ion etching (DRIE) allowing the Parylene to physically interlock with the substrate. While the resilience of the film was greatly increased, this technique is limited to silicon substrates and further only provides locally increased adhesion.

Parylene-C is deposited in a low-pressure chemical vapor deposition (LPCVD) process which allows it to penetrate into pores as small as 5 nm in diameter.25,26 Based on this knowledge, we propose to create nanostructured surfaces to allow a uniform anchoring over a large surface. This is a universal, material-independent approach to provide a strongly adherent Parylene-C layer, provided that a method for nanostructuring of the substrate exists. In this study, we provide such a nanostructuring approach for a widely used biomedical alloy (Ti6Al4 V) and show its effect on the wetadhesion of Parylene-C.

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EXPERIMENTAL SECTION

Two types of samples were prepared: cuboidal samples with a suspended Parylene membrane for adhesion testing and foil samples for more detailed characterization of TiO2 nanostructures. Where not explicitly mentioned, both types of samples were treated identically. An overview of the entire process is shown in Figure 1a. Materials and Substrate Preparation. Titanium grade 5 (Ti6Al4 V) cuboidal samples were prepared from a rolled and annealed titanium slab purchased from Thyssenkrupp Materials Schweiz AG. The 10 × 10 × 6 mm cubes were formed in a waterjet cutting process, and a 1.5 mm diameter, centered through-hole is B

DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials drilled into the 10 × 10 mm front face. The front face was finished on a planetary grinding machine with three grinding papers of 120/220/ 500 grit and subsequently with three differing silica polishing suspensions with particle size of 9/3/1 μm. Ti6Al4 V foils were purchased from E. Wagener GmbH and cut into 4 × 10 × 0.2 mm strips with a guillotine shear. All samples were cleaned by sonication in acetone, 2-propanol, and deionized (DI) water for 30 min, separately, followed by treatment with an oxygen plasma at 100 W power at a gas-flow rate of 5 sccm O2 for 5 min. The polished and cleaned samples were used as such as flat reference surfaces. Preparation of TiO2 Nanostructures. Three different surface modification processes with varying process parameters were utilized in order to obtain nanostructures with a range of morphologies and sizes. A full list of synthesis parameters can be found in Table 1. The first process, hereafter referred to as soaking (SK), was performed according to Kokubo et al.27 The samples were immersed in aqueous NaOH (Pellets ACS grade, VWR International) solution heated to 60 °C under constant stirring. The second process, hereafter referred to as the hydrothermal (HT) route, was performed according to Wang et al.28 The reaction was carried out in a 50 mL Teflon-lined stainless steel autoclave. The cuboidal samples, with their polished faces in a vertical position, were added to 28.5 mL of aqueous NaOH. For the foil samples the amounts were adjusted such that all process conditions remained constant. The autoclaves were placed in a preheated oven at 180 °C for the respective reaction times and subsequently cooled for 2 h at room temperature before removing the samples from the autoclave. The third process, hereafter referred to as ethylene glycol enhanced hydrothermal (EG) route, was performed according to Yin et al.29 Identical to hydrothermal route, here only the solution was changed from pure NaOH to NaOH dissolved in a 1:1 solution of deionized water and ethylene glycol (C2H6O2, >99% purity, Sigma-Aldrich), and the reaction temperature was increased to 200 °C. For all routes, the samples were carefully washed in deionized water after synthesis of the nanostructures. Then, the cubes were immersed in 0.5 mmol/L aqueous HCl solution at 40 °C for 24 h. After drying in vacuum at 40 °C for 2 h the samples were heated to 400 °C and left to calcinate for 1 h. The heating rate was 5 °C/min, and the furnace was naturally cooled to room temperature before removing the samples. Preparation of Suspended Parylene-C Coatings. In order to obtain suspended Parylene membranes for the subsequent adhesion testing, the cavities in the cubes were filled with a sacrificial material prior to Parylene coating. A temporary bonding adhesive (Crystalbond 555-HMP, Structure Probe Inc.) was used as this sacrificial material. The cubes were mounted upside down on a thin PTFE foil and gently pressed down by a fixture. The assembly was then degassed in a vacuum oven at 100 °C and in vacuum. Both the sacrificial material and the cubes were then heated at 65 °C. At this point a drop of the molten sacrificial material was added to the cavity in the cube. In order to remove trapped air in the cavity, a vacuum was applied before letting the adhesive slowly harden at 60 °C for 1 h. Subsequently, the oven was cooled to 40 °C over 2 h. Parylene deposition was performed by APM Technica AG. All samples were cleaned in situ three times with O2 plasma at 250 W power and 200 sccm O2 flow for 1 min. Samples treated with adhesion promoter were coated in situ with vaporized silane A-174 (3(trimethoxysilyl)propyl methacrylate) at a flow rate of 100 μL/min for 3 min with ambient air as a carrier gas. The Parylene-C monomer was vaporized at 120 °C for 3 h and 125 °C for 5 h and subsequently pyrolyzed at 760 °C. The samples in the reaction chambers were held at room temperature, and the initial chamber pressure was 5 Pa. The thickness of the deposited Parylene film was 9.7 ± 0.2 μm as measured on the cross-section of the film. The Parylene membranes were released by removing the temporary adhesive in deionized water at 55 °C for 90 min. Accelerated Aging. Accelerated aging of the as-prepared samples with suspended membrane was performed on a hot plate. In order to simulate the conditions within the human body, Hank’s balanced salt solution (HBSS, with calcium and magnesium, no phenol red, Thermo

Fischer Scientific) was used as the aging medium. The samples were fully immersed into HBSS, and any trapped air bubbles preventing the solution from reaching the membrane were removed by gentle agitation. Aging temperatures were 50 and 80 °C for pristine samples and 80 °C for samples treated with silane A-174. All samples were aged for 72 h. In order to prevent the samples from drying out and/or aging further until testing is performed, the samples were stored in HBSS in a refrigerator at 5 °C. Prior to testing, the samples were dried under gentle nitrogen flow. Adhesion Testing. The adhesion of both as-prepared and aged cuboidal samples was performed using a blisterlike adhesion test. As illustrated in Figure 1b the cubes were clamped on to a sealed pressure chamber connected to a high-pressure reservoir. The pressure inside the chamber was controlled using an electromagnetic proportional valve PVQ33-6G (SMC Pneumatics) and measured with an analogue relative pressure sensor XFPM-700 KPGR (Fujikura). Both the valve and the sensor were operated using a Keithley 2400 SourceMeter (Tektronix Inc.) instrument connected to a PC for data recording and control. The pressure in the chamber was increased at a constant rate until failure of the Parylene film occurred. In order to observe the delamination front, the Parylene-C film was observed using a white light interferometer (WLI; Zygo Corporation). The failure mode was further investigated by scanning electron microscopy (SEM). Characterization Methods. A FEI NovaNanoSEM 450 (Thermo Fisher Scientific) instrument was used to analyze the nanostructures’ morphologies, characteristic dimensions, Parylene-C penetration depth into the nanostructures, as well as the blister-test failure mode at an acceleration voltage of 5 kV. Cross-sectional views of both pristine and Parylene-C coated nanostructures were prepared using focused ion beam (FIB) milling (Zeiss FIB-SEM NVision 40, Carl Zeiss AG). The structures were milled at a gallium ion current of 13 nA. SEM micrographs were evaluated using ImageJ (National Institutes of Health NIH, U.S.A.) software; for each characteristic dimension a minimum of 20 measurements was performed. Transmission electron microscopy (TEM) lamella samples were prepared by embedding a stack of nanostructured foils in epoxy resin, followed by subsequent thinning through grinding, dimpling, and ion milling using a GATAN 691 precision ion polishing system (Gatan Inc.). TEM, selected area electron diffraction (SAED), and energy-dispersive X-ray spectroscopy (EDX) measurements were performed using a FEI Talos F200X (Thermo Fisher Scientific) instrument. Brunauer−Emmett−Teller (BET) specific surface analysis of physisorption isotherms were performed using the Nova 4000e measurement bench (Quantachrome instruments). Prior to measurement the samples were outgassed at 200 °C for 1 h. N2 was used as the adsorptive, and the measurement temperature was 77.3 K. BET analysis was performed in the pressure range 0.05 < P/P0 < 0.3, and the adsorption cross-section was assumed to be 0.162 nm2. The obtained BET constants range from 21.3 to 37.7, and we therefore assume an uncertainty in specific surface area of ±20%. Material composition of the prepared nanostructures was analyzed from Raman spectra, collected using a confocal Raman spectroscope (WITec, CRM 200; laser wavelength, 532 nm) at a laser power of 2 mW.

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RESULTS AND DISCUSSION

Synthesis and Characterization of Nanostructures. The morphologies and dimensions of the synthesized nanostructures were assessed using SEM imaging of the foil type samples. Table 1 summarizes the characteristic dimensions, nanostructure thickness, the specific surface area measurements, as well as the chosen acronyms of all obtained structures. The chosen synthesis parameters resulted in three distinct types of morphologies, specifically pore-shaped structures (nanopores), flower-petal-shaped structures (nanoflowers), and fibrous network structures (nanofibers). Differing synthesis parameters generating the same morphology typically resulted in differing characteristic dimensions and structure heights. C

DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials

Figure 2. (a) SEM image of the as-fabricated nanostructures at a tilt angle of 45°. The inset shows a high-magnification detail in a top-down view. (b) SEM images of FIB cuts showing the cross-section of the nanostructures. (1) SK-5-24, (2) HT-3-3, (3) EG-3-6, and (4) HT-1-6.

Figure 2 shows SEM images of the nanostructures’ surfaces and their cross-sections obtained by FIB milling of four representative samples of the respective morphologies. The porous structures seen in Figure 2a1 were obtained by the soaking process, and exhibited a fractal character with larger pores on the surface that enclose smaller pores at lower heights. The flowerlike structures seen in Figure 2a4 were covered with large petal-shaped structures and smaller, upright, needlelike structures in between. Similarly, the fiber nanostructures in Figure 2a2,3 consisted of upright needle structures which were covered by a dense fiber network. The so overgrown fibers formed nestlike superstructures. Figure S1 shows structures synthesized on the cuboidal samples with similar features. The FIB cross-sections reveal that all samples mainly consist of vertical, slitlike pores, with all samples possessing progressively smaller pores from the top surface toward the bulk of the substrate. Figure 3a,b shows the statistical distribution of pore area and density as obtained by visual inspection of the FIB cross-section images with Image-J software. The results revealed that there are significant differences in pore area distribution and the average in-plane pore density. The above-mentioned, smaller pores toward the bulk of the material are of similar size for all morphologies. The porous samples had slightly smaller pores and a lower pore density in comparison to the other morphologies, indicating an overall lower porosity. Fibrous structures had a high number of average-sized pores due to their crisscrossing nature. In addition to medium-sized pores, flowerlike structures possessed a number of large pores possessing an area above 0.02 μm2. Such large pores were created by the overarching petals. Surface characteristics were further investigated using BET analysis of N2 physisorption isotherms.30 A list of the obtained specific surface areas is given in Table 1 and ranged from approximately 0.5 to 1.2 m2/g. The low magnitude of these absolute values stems from measuring a thin, porous layer (≈1 μm) on a thick substrate (≈200 μm). A more useful measure is perhaps the effective surface area increase of nanostructured samples over a theoretical, ideally flat sample of the same macroscopic dimension and weight. The corresponding, calculated normalization value is 0.0024 m2/g, and the resulting effective surface area increase is 247.1 ± 49, 190.4 ± 38, 564.2

Figure 3. Box charts of the (a) pore area distribution and (b) pore density distribution as determined by ImageJ evaluation of the crosssectional SEM images. Pore area was determined by manually fitting polygons. Pore density denominates the number of pores along the width of the cross-section. (c) N2 adsorption/desorption isotherms. Each measurement was performed on a batch of at least 12 foil samples.

± 112, and 256.2 ± 51 for the SK-5-24, HT-3-3, EG-3-6, and HT-1-6 samples, respectively. It is obvious that the sample synthesized using the ethylene glycol enhanced hydrothermal route demonstrated a twice higher effective surface area in comparison to samples synthesized from other routes. The comparison of pore count, size, and layer thickness between the samples by SEM images did not vary significantly enough to explain this large difference in surface area. D

DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials The corresponding adsorption and desorption isotherms are shown in Figure 3c. According to IUPAC recommendations, the isotherms were of type IIb (type II isotherm with type H3 hysteresis), indicating the presence of structure aggregates forming slit-shaped pores. This is in good agreement with the observed morphologies seen in Figure 2b, where the open space between the needlelike structures resulted in the slitlike pores. The extension of the hysteresis into the low-pressure regime (P/P0 < 0.4) is indicative of very small pores of similar size as the adsorptive.30 This phenomenon likely corresponds to very small pores observed toward the bulk of the substrate. The isotherm for ethylene glycol enhanced sample appeared as an intermediate form of the above explained isotherm and an additional type H1 hysteresis at 0.4 < P/P0 < 0.8. This possibly indicates that, instead of nanofibers, this sample consisted of nanotubes, and the observed hysteresis corresponds to the capillary filling of the nanotube’s cavity. The doubling of the effective surface area increase could also be explained through this observation. Figure 4 shows the Raman spectra of the as-synthesized nanostructures after calcination. The peaks at 149, 399, 516,

Figure 5. (a) DF TEM image of the flowerlike nanostructures of sample HT-1-6. The inset shows the SAED diffraction pattern of the nanostructures. (b) Higher magnification of a cross-section of a fiber structure, showing that it is hollow. (c) Higher-magnification, bright field image of the porous oxide layer underneath the nanostructures. (d) High-angle annular dark field (HAADF) image and EDX map of a zoomed-out section, showing that Al does not migrate into the nanostructures. (e) HAADF, SAED, and EDX map of nanofibers scratched off from an EG-3-6 sample.

(200) and (101) diffractions of anatase TiO2. The EDX analysis presented in Figure 5d and Figures S2 and S3, as expected, show titanium, aluminum, and vanadium in the bulk metal. The nanoporous interface layer appears to be an oxide of titanium, as seen by the large increase in oxygen in that region. Further, there is no evidence of aluminum or vanadium migrating into the nanostructures, which predominantly consist of titanium and oxygen. Figure 5e shows TEM analysis performed on nanofibers from sample EG-3-6, which were detached from the bulk substrate. Similar to sample HT-1-6, the nanofibers are tubular and resemble the structures obtained by Yin et al.29 From the SAED pattern obtained from a single nanofiber, we can observe the presence of twinning defects in the material, as previously observed by Kolen’ko et al.34 The diffraction pattern matches closely to sodium trititanate (Na2Ti3O7) with the measured d-values = 1.88, 3.11, 3.52, and 6.744 Å corresponding to the (020), (111), (110), and (−101) diffractions, respectively. EDX maps and spectra shown in Figure 5e and Figure S3, respectively, further confirm the presence of sodium, titanium, and oxygen for sample EG-3-6. Scheme 1 shows the proposed growth mechanism of the respective nanostructures. For the soaking synthesis route, the structures are formed through dissolution of bulk titanium through corrosive attack of the hydroxyl groups. In presence of the alkali ions, the dissolved TiO2 forms a sodium titanate hydrogel layer which is dehydrated and densified in the subsequent processing steps.35 Due to the isotropic nature of the dissolution process, the resulting structures are porous in nature. Under hydrothermal conditions, precipitation effects of the sodium titanate begin to dominate the dissolution, thus encouraging a more anisotropic structure growth.36 The further

Figure 4. Raman spectra of the synthesized nanostructures. As indicated, all but the HT-1-6 and reference spectra were scaled up 10fold to make the peaks more visible.

and 637 cm−1 were exhibited very clearly for the HT-1-6 sample and less distinctively for the EG-3-6 sample, corresponding well with reported values of the Raman modes of titanium anatase.28,31 Samples SK-5-24, HT-3-3, and EG-3-6 further show Raman modes of sodium titanate at 190, 278, and 702 cm−1.32,33 The peak at 430 cm−1 is possibly related to the presence of sodium titanate as well; however, there is a clear downshift with respect to the reported values of 446 cm−1. TEM analysis of sample HT-1-6, was performed on a TEM lamella and is presented in Figure 5a−c. The cross-section, seen in Figure 5a reveals that the nanostructures are firmly attached to the Ti6Al4 V substrate. Figure 5b shows that the needlelike structures between the larger sheets are tubular in nature. Interestingly, it further reveals the presence of an interface layer between the bulk metal and the nanostructures. The magnified image in Figure 5c shows that this layer is nanoporous. SAED analysis, performed on multiple nanostructures, shows rings with d-values at 1.88 and 3.56 Å corresponding well to the E

DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials Scheme 1. Illustration of the Influence of Processing Parameters on the Nanostructure Morphology

formation of flowerlike and fibrous structures strongly depends on the NaOH concentration during the course of the reaction. The formation of titanate nanotubes has been attributed to the following process: First, nanosheets become detached from the bulk of the substrate through corrosive attack. Due to a difference in surface energy on both sides of the sheet, they start rolling up and form nanotubes.28,37 The highly alkaline conditions, in the initial phases of the reaction, stabilized the tubes in this state. As the reaction progresses and the sodium in the solution is consumed, further formation of nanotubes is inhibited, and the sheets start to unroll and form flowerlike structures. Due to an oriented attachment crystal growth, the flowers grow larger with longer reaction times. These mechanisms are supported by the fact that we obtained highly crystalline, flowerlike structures for a process (HT-1-6) with low concentration and long duration, whereas the reaction (HT-3-3) at higher concentration and shorter time led to a fibrous structure. Yin et al.29 hypothesized that the presence of ethylene glycol inhibits the initial formation of nanosheets and, instead, forms nanowires through oriented attachment crystal growth. Since the directionality of this mechanism is not dependent on the alkaline concentration, growth of 1D structures is supported for longer reaction times, leading to a denser fiber network. Suspended Parylene-C Films. Figure 6 shows SEM images of the nanostructure−Parylene interface obtained by FIB milling. From these figures we can observe that the penetration of all nanostructures with Parylene was complete, with the Parylene reaching the pores closest to the bulk of the substrate. As can be expected from the morphologies of the nanostructure pore networks, and the highly conformal nature of Parylene deposition, a significant number of voids were present in the deposited polymer. While voids were present for all samples, fibrous nanostructures appeared to have fewer and smaller voids than the flowerlike morphologies. For the porous SK-5-24 sample the top layer of the nanostructures had a high amount of large voids, rendering the resulting structure almost entirely hollow. Such voids form when a narrow pore entrance is obstructed before the main cavity of the pore is completely filled. Since morphology and contact-area of the interface are not changed by their presence, no significant effect on the microscale adhesion performance is expected. However, the macroscopic mechanical properties such as elasticity and yield strength are likely to be impacted. Due to high thickness of the polymer layer (10 μm) as compared to the nanostructure layer (0.5−1.5 μm), the morphology of the Parylene’s top surface was not altered by the presence of such nanostructures.

Figure 6. SEM images of the FIB cuts showing a cross-section of the (a) porous, (b, c) fibrous, and (d) flowerlike nanostructures after depositing Parylene-C. The inset in part b shows the formation of a Parylene void within a titanate structure.

Figure S4 shows a WLI height profile of a released, suspended Parylene-C membrane. The membrane was slightly receded from the titanium surface. The average magnitude of this deformation over all samples is 134 ± 50 μm. This was a result of an imperfect shaping of the sacrificial material during the formation process. As a consequence, a section of the membrane was attached to the interior sidewall of the cavity over which the membrane is suspended. This is an important factor to keep in mind for the pressure testing of the membranes, as it influences the stresses on Parylene−substrate interface. This effect is more pronounced at lower pressures. Comparing surface profiles from before and after release, the membrane does not show any signs of deformation during the release process. This indicates that there is no prestress present in the membrane. Figure S5 shows micrographs of both reference and modified cuboidal samples after accelerated aging at 80 °C in HBSS. Since the cuboidal samples were fully immersed into the liquid, the interface was expected to be degraded by an undercut attack starting from the circumference of the suspended membrane. The lack of a clear, visual presence of such an attack front can either indicate a uniform and complete disintegration of the interface or that the interface was not attacked at all. The nanostructured samples aged at 50 or 80 °C both featured small blisters in the Parylene layer. Such blisters can be attributed to diffusion of water vapor through the Parylene layer and subsequent accumulation at the interface.38 Based on the water vapor permeation rates observed by Menon et al.,39 a water vapor transmission rate of 9 × 10−4 g/(mm2 day) was expected for the aging process at 80 °C. This high transmission rate readily justified the assumption of a diffusion-based attack of the Parylene−titanium interface. The number and size of blisters varied widely between all nanostructured samples. As Figure S6 shows, the majority of the blisters are small in area, with a few larger blisters appearing after aging at 80 °C. The variation of blister sizes between the samples and the relative similar distribution at the different aging temperatures lead us to assume these blisters either appear at defects in the F

DOI: 10.1021/acsanm.8b00081 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials nanostructure film or due to inflation of voids in the Parylene film due to outgassing of the substrate. Further investigations are required to determine the definitive cause of these blisters. Parylene-C Adhesion. Figure 7 shows the results of the blisterlike adhesion testing of the cuboidal samples after

that there is an upper limit of structure density at which the adhesion performance starts to reduce. As shown in Figure 8b, the failure for the nanostructured samples occurred along the circumference of the cavity, while

Figure 7. Bar chart of the failure loads obtained by the adhesion testing. The red line and area depict the average of failure load and of an unaged, dry-tested sample. The * and *** labels indicate distribution differences at a significance level of 0.05 and 0.001, respectively (t-test).

Figure 8. (a, b) Micrographs of a fractured reference and nanostructured sample, respectively, after accelerated aging at 80 °C for 72 h and blisterlike testing. The reference sample delaminated over large areas, while for the nanostructured sample only the membrane burst. (c, d) SEM images of a fractured reference and nanostructured sample. The SEM images were taken with samples at a tilt of 60°. The magnified insets clearly show the difference in Parylene residuals, indicating a cohesive failure mode for the nanostructured sample. The scale bars in the insets are 10 and 5 μm for parts c and d, respectively.

accelerated aging. As indicated by the red line in the figure, the average failure load of unaged, dry samples serves as a benchmark value. There was no appreciable performance difference for unaged samples with and without nanostructuring. The loss of durability for the reference samples is very evident. Already at an aging temperature of 50 °C, the failure load was reduced from 5.9 × 105 to 1.6 × 105 Pa. Aging at 80 °C further reduced the failure load to 0.7 × 105 Pa, a reduction of 88% as compared to the unaged samples. The durability reduction for samples with a silane A-174 pretreatment was almost identical. The WLI measurements of the blister expansion show that this low-pressure failure was a consequence of an advancing delamination front. As soon as the delaminated area reached a critical size, catastrophic failure occurred, and the Parylene film ruptured. The porous nanostructure samples had a significantly improved performance over the reference but were falling short of the benchmark by 14%. This slight performance reduction is possibly related to the presence of larger voids in the Parylene, as they might result in an interface with an overall reduced yield strength. For all samples with fiber- and flowerlike morphology, the performance matches the benchmark of the unaged samples. Furthermore, there was no significant change in adhesion when increasing the aging temperature from 50 to 80 °C. This indicates that the degradation of the interface was fully suppressed. No significant performance difference between the fibrous and flowerlike can be observe. EG-3-6 with its thick and dense fiber layer, however, exhibited a slight increase in spread between the measured failure loads. This might indicate

the Parylene stays firmly attached on the entire surface of the sample. Figure 8c,d shows a higher-magnification SEM image of the fracture site. It is clearly visible that the Parylene completely detached from the smooth reference surface, leaving little residues behind. The nanostructured surfaces, however, were still covered in Parylene, even at the microscale. Residual Parylene filaments show that fracture occurred in the bulk of the polymer. This indicates that the introduced nanostructures change the failure mode from adhesive to cohesive. Thus, the failure of the film now depends on its favorable bulk properties. Hence, on the basis of these results, we successfully demonstrated that the anchoring of Parylene on nanostructured Ti6Al4 V surfaces is feasible and dramatically improves its adhesion in wet environments. The most commonly applied adhesion promotion route for Parylene-C is a pretreatment of the substrate with silane A-174. To offer a comparison of silanized reference samples and the nanostructured samples as well as the influence of the silane on the nanostructures, we additionally tested a batch of silanetreated and HBSS aged samples. From Figure 7 we observe no significant adhesion improvement with the addition of the silane; the measured failure loads were ∼88% lower than for the nonsilanized nanostructured samples. This is supported by the studies of Youssefian et al.,23 who observed a sharp decline in adhesion forces for silanized interfaces at a water content of >60%. For all nanostructured samples, there was a drastic decrease in adhesion performance when they were pretreated G

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with silane A-174. The reduction was most pronounced for fibrous samples where we observed a failure load reduction of 72%. Figure S7 shows a comparison of a fiber nanostructure interface with and without added silane. While the Parylene interpenetration of the silanized sample appears to be similar, the fiber structure looked denser and reduced in thickness. We therefore assume that, during its vapor phase deposition, the silane condensed between the fibers and caused the structures to collapse and stick to each other. As a result, a much less interconnected pore network was present for anchoring of the Parylene. This effect is less pronounced for the flowerlike and porous samples as they possessed much larger features, which were less prone to collapse. From these observations we can state that the standard adhesion promoter silane A-174 reduced adhesion promotion in every case and, hence, would not recommend its use on nanostructured surfaces.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support by the Stavros Niarchos foundation. This work is part of the Zurich Heart project under the umbrella of University Medicine Zurich/Hochschulmedizin Zürich. The authors acknowledge support of ScopeM/Swiss Federal Institute of Technology ETHZ, in particular from Dr. Alla Sologubenko, Dr. Eszter Barthazy, and Dr. Joachim Reuteler. We further gratefully acknowledge Sung Min Kim (Laboratory of Energy Science and Engineering, ETH Zurich) for the support with BET analysis and Dr. Mirsoslav Haluska (Micro- and Nanosystems, ETH Zurich) for help with Raman spectroscopy. We gratefully acknowledge APM Technica AG for their advice and support for the Parylene-C deposition.

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CONCLUSIONS In this study we have successfully shown that the delamination of Parylene-C on titanium surfaces in wet environment can be stopped by anchoring the Parylene film at the nanoscale. Titanium oxide and sodium titanate nanostructures with various morphologies were synthesized on bulk Ti6Al4 V by a facile, hydrothermal route. The so-obtained nanostructure layers had thicknesses from approximately 250 to 1500 nm and an increase in specific surface area ranging from 200- to 550fold. FIB cuts through the polymer−nanostructure interface show that the Parylene completely interpenetrates the nanostructures, regardless of their morphology. The adhesion of the Parylene was tested by inflating a suspended portion of the polymer film and was quantified by using the pressure load at which failure occurs. After aging samples at 80 °C for 72 h in a physiological salt solution, the Parylene layer on smooth, unmodified surfaces lost most of its adhesion, leaving behind a clean metal surface after fracture. However, nanostructured surfaces led to a complete retention of adhesion forces, showing no difference to samples prior to aging. The failure load on the nanostructured samples is increased by 840% even in comparison to reference samples using the standard adhesion promoter silane A-174. While small blisters appear in locations away from the suspended membrane, we believe that these shortcomings may be alleviated by further process optimization. The findings of this study might allow the use of Parylene for long-term implant applications with direct contact of bodily fluids without loss of adhesion.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsanm.8b00081. Comparison of Nanostructures on cuboidal and foil samples, additional EDX maps and spectra, WLI surface profile, micrographs, blister size analysis, and FIB crosssection (PDF)

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REFERENCES

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: +41 44 632 71 15. ORCID

Silvan Staufert: 0000-0002-5203-3934 H

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