Journal of Neuroscience Methods 131 (2003) 9–26 Fabrication and characteristics of an implantable, polymer-based, intrafascicular electrode Stephen M. Lawrence, Gurpreet S. Dhillon, Kenneth W. Horch∗ Department of Bioengineering, University of Utah, 50 S. Central Campus Dr., Salt Lake City, UT 84112, USA Received 27 November 2002; received in revised form 8 July 2003; accepted 11 July 2003 Abstract We describe new manufacturing techniques and physical properties of an improved polymer-based longitudinal intrafascicular electrode (polyLIFE). Modifications were made to correct: (1) poor metal film adhesion and fatigue resistance, (2) inconsistent insulation adhesion and control over recording/stimulation zone length, and (3) insufficient tensile strength for clinical use. Metal adhesion was significantly improved by both plasma treatment and fiber rotation (about the long axis) during metal deposition. Fatigue resistance was improved by reduction in sputtering energy (time × power) combined with long axis rotation, resulting in thin metal films that were 250 times more resistant to cyclic bending fatigue. Insulation adhesion was enhanced with the application of an adhesion-promoting silicone (MED2-4013, Nusil), while the recording/stimulation zone length was controlled to 1 ± 0.2 mm (mean ± S.D.). The polyLIFE was made more robust by the inclusion of three individually metallized fibers, improving its tensile strength by a factor of 4 while producing minimal changes to its overall stiffness. However, the metallized fiber redundancy did not significantly affect fatigue resistance. The manufacturing changes described in this study enable the construction of more mechanically robust polyLIFEs, which should provide greater success when chronically implanted in peripheral nerves. © 2003 Elsevier B.V. All rights reserved. Keywords: polyLIFE; Intrafascicular electrode; Polymer 1. Introduction The design and development of peripheral nerve interfaces has been a major goal of the researchers interested in aiding patients with motor deficits such as foot drop, paralysis, and limb amputations. In order to provide more natural and fine control over prosthetics that utilize these neural interfaces, a concerted effort has been made to increase electrode selectivity for both stimulation and recording. Selectivity for stimulation has been accomplished to a certain degree by modifying cuff technology to add more active sites, but recording selectivity has not been reported (Grill and Mortimer, 1996; Rozman et al., 1993; Veraart et al., 1993; Walter et al., 1997). A much larger increase in selectivity has been obtained by the use of interfaces whose recording/stimulation zones (active sites) have either been placed within individual fascicles or rely on neural regeneration at a nerve stump. Several different designs incorporating this intraneural approach have ∗ Corresponding author. Tel.: +1-801-585-1981. E-mail address: k.horch@m.cc.utah.edu (K.W. Horch). 0165-0270/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0165-0270(03)00231-0 been developed including silicon-based (Akin et al., 1994; Branner et al., 2001; Edell, 1986; Kovacs et al., 1992; Rutten et al., 1991; Veltink et al., 1989) and polyimide-based (González and Rodrı́guez, 1997; Stieglitz and Meyer, 1997) microelectrode arrays as well as longitudinally implanted microwires (Bowman and Erickson, 1985; Malagodi et al., 1989). In acute experiments, these intraneural interfaces have allowed activation of axons at low current densities and recording from small subsets of neurons as opposed to the relatively high current densities required with cuffs and mass activity obtained with extraneural recordings. Key design requirements for intrafascicular devices include well defined, low impedance active sites, small dimensions in order to minimize damage during insertion and prevent long-term foreign body reaction, and the ability to penetrate the tough epineurial and perineurial connective tissue. While silicon-based microelectrode arrays for peripheral nerves take advantage of long-proven micromachining techniques to design devices with high densities of active sites, on-chip circuitry, and manufacturing repeatability, they have achieved limited success in long-term recordings. The failure of silicon-based probes to provide long-term peripheral nerve recordings is thought to be due in part 10 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 to their high stiffness compared to the surrounding neural tissue (mechanical mismatch), causing large foreign body reactions and extensive encapsulation (Stensaas and Stensaas, 1978). Long-term intrafascicular recordings (>4 months) from intact nerves have only been reported with the use of longitudinally implanted intrafascicular electrodes constructed from Teflon-coated, 25 ␮m diameter Pt-Ir wire (Lefurge et al., 1991). These longitudinal intrafascicular electrodes (LIFEs) demonstrated selective recordings of afferent activity from the radial nerve of cats for up to 6 months with good biocompatibility. However, we believe that the mechanical stiffness of the Pt-Ir wire caused gradual drift of these electrodes within the implanted fascicle over time, leading to recording instability and dense encapsulation tissue formation around the electrode. Therefore, we feel that further minimization of mechanical mismatches between the implanted electrodes and the surrounding neural tissue will increase the biocompatibility and recording stability of intrafascicular technologies. Towards this end, our lab has developed a polymer-based longitudinal intrafascicular electrode (polyLIFE), constructed by metallizing single 12 ␮m diameter Kevlar® fibers and insulating them with silicone (McNaughton and Horch, 1996). These polyLIFEs have shown similar acute recording capabilities and tensile strength to Pt-Ir LIFEs, but have a 60 times greater flexibility. In addition to their acute recording capabilities, the polyLIFEs have demonstrated long-term biocompatibility in both peripheral nerve (Lawrence et al., 2002) and dorsal rootlets (Malmstrom et al., 1998), but have yet to achieve the reliability required for chronic recording. Previous work has identified problems associated with the conducting and insulating layers, including poor adhesion of the metal to the Kevlar® fiber, limited mechanical fatigue resistance, inconsistent adhesion of the silicone to the metal layer, and inadequate control over active site length (unpublished observations). Clinical experimentation has also demonstrated a need to increase the tensile properties of the original polyLIFE design in order to better survive surgical handling and percutaneous applications (preliminary unpublished data from clinical trials). The aim of the present study was to improve mechanical and electrical properties of polyLIFEs by modification of their fabrication processes. The effectiveness of these modifications was evaluated by several bench tests. The result was a fabrication process that can produce more mechanically robust polyLIFEs with specifically tailored electrical properties. 2. Materials and methods Fig. 1. General design of a polymer-based intrafascicular electrode (polyLIFE). The polyLIFE consists of a 12 ␮m diameter Kevlar® fiber, metallized with sputter-deposited Ti, Au, and Pt and insulated with silicone (metal and silicone layers not drawn to scale, but drawn to relative thickness). The recording/stimulation zone consists of approximately 1 mm non-insulated portion of the metallized fiber (diameter:length ratio of ∼500:1). lated with silicone (Fig. 1). polyLIFE manufacture involves four steps: (1) fiber preparation and mounting, (2) surface treatment, (3) metal deposition, and (4) insulation. In order to produce the most mechanically and electrically robust polyLIFE possible, variations of each manufacturing step were conducted and subsequently evaluated with bench-top testing methods. Certain processes within each step were identical to all constructed test samples and are listed below. Other procedures within each of the four manufacturing steps were conducted on some samples but not others, and are described in detail in the following sections. Step 1 (fiber mounting): Kevlar-119® fibers were obtained from Dupont in yarn form (∼1500 denier). Individual fibers of 1 m length were teased from these bundles while soaking in alcohol (to help prevent fiber-to-fiber abrasion). These fibers were then wrapped under constant tension onto metal looms and secured with cyanoacrylate. Step 2 (surface treatment): Once the fibers were secured to the loom, they were washed by sequential rinsing in four baths for 3 min per bath: acetone, methanol, isopropanol, and DI water. Step 3 (metal deposition): Metal deposition was accomplished by dc sputter deposition (Denton, Discovery 18 sputter system) at vacuum of <0.67 × 10−3 Pa. Three metal layers, Ti, Au, and Pt, were sequentially sputtered onto the loom-bound fibers. This sputter system contains three sputter guns, enabling deposition of up to three different metals onto a substrate without breaking vacuum. Step 4 (insulation): The metallized fibers were then transferred to another metal loom, secured with cyanoacrylate and insulated with a dip-coating process. In this process, a drop of medical-grade silicone, at the end of a 2 mm diameter steel rod, was applied to the metal clad fiber as it was drawn slowly along the fiber’s longitudinal axis. Reproducible control of dip-coating speed was provided by a modified perfusion pump motor. 2.2. Manufacturing process modifications and LIFE construction 2.1. General design The basic structure of the polyLIFE consists of a Kevlar® mono-filament core, clad with three metal layers and insu- 2.2.1. Fiber mounting Due to the cylindrical geometry of the fibers, it is challenging to obtain uniform coating thickness, not only around S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 the fiber’s circumference, but also along its longitudinal axis. Gas-phase scattering makes it possible to coat surfaces that do not view the source; however, these “shadowed” surfaces receive metal at a much reduced sputtering flux, causing non-uniformities in thickness (Fancey and Matthews, 1991). Furthermore, sputtering flux is inversely proportional to the distance from the sputtering source. An additional geometrical concern involves the angle of incidence of impinging atoms. Sputter deposition at normal angles of incidence and higher substrate temperatures tends to produce more dense films with lower surface roughness (according to simulations; Dong et al., 1996), whereas sputter deposition at oblique angles of incidence tends to cause columnar development (a common microstructural feature of thin-films deposited at low pressures and temperatures) (Nieuwenhuizen and Haanstra, 1966). This columnar morphology can significantly affect thin films in terms of mechanical integrity (crack propagation is frequently seen along weak, low-density intercolumnar regions), electrical properties, and surface roughness (Ohring, 2002). For geometrically complex surfaces (such as miniature diameter fibers), control of various thin-film properties can be 11 exercised by both optimization of the source-metal to substrate distance, reduction of the shadowing effect, and minimization of oblique angles of incident atoms. This optimization might be accomplished by continuous rotation and short sample lengths. However, in our sputter system, rotation was possible only around one axis, and maximization of fiber length was a critical goal. Thus, positioning of the mounted fibers inside the sputter system could significantly affect the mechanical performance of the deposited metal film. In order to ascertain the optimal mounting strategy, fibers were mounted on two different types of looms: (1) a flat, rectangular-shaped, metal frame (referred to as the “flat loom” shown in Fig. 2A) and (2) a cylindrical-shaped metal frame (referred to as the “cylindrical loom” shown in Fig. 2B). 2.2.2. Surface treatment Plasma treatment of the surfaces has long been known to enhance the adhesion of metal to polymers. At high pressures (>9.3 Pa) and with plasmas created from reactive gases such as O2 , dry chemical etching (commonly known as reactive-ion etching) of the substrate can enhance Fig. 2. Loom orientation during sputtering. (A) The geometrical condition during sputtering of fibers mounted on a flat loom. In this case, rotation of the loom is accomplished by rotation of the entire platform. (B) The geometrical condition during sputtering of fibers mounted on a cylindrical loom. In this case, rotation of the loom is accomplished via a feed-through to a motor external to the sputter chamber. The loom axle is connected to the feed-through by a flexible rotation coupler. 12 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 adhesion by chemically modifying, removing, or roughening the surface (Lamendola et al., 1999; Sugawara and Stansfield, 1998). For instance, reactive-ion etching of Kevlar-149® fibers with 100 W of O2 plasma demonstrated significant increases in oxygen content of the surface within 15 s and visible surface roughening after 10 min (Sheu and Shyu, 1994). At low pressures (<0.93 Pa) and with chemically inert plasmas, physical sputtering (commonly known as sputter etching) of the substrate can enhance adhesion by removing or dispersing contaminants, creating dangling bonds or free radicals, and roughening the surface (Baglin, 1995; Ohring, 2002; Sugawara and Stansfield, 1998). In this study, four different surface treatment processes were evaluated: (1) no treatment, (2) reactive-ion etch (30 s of 100 W O2 plasma at 27 Pa with an Oxford, Plasmalab 80 Plus), (3) sputter etch (6 min of 250 W rf power of an N2 plasma at <1.4 Pa with a Denton, Discovery 18 sputter system), and (4) a reactive-ion etch and a subsequent sputter etch. Each surface treatment process will be referred to as “S#”, where S stands for surface treatment and # stands for the specific process (e.g. S2 stands for surface treatment process #2, Table 1). 2.2.3. Metal deposition The mechanical properties of vacuum deposited thin films have been shown to be quite different from bulk properties. For instance, depending on the deposited film microstructure or morphology and particular testing methods, different researchers have found widely different Young’s modulus for Au films (53–55 GPa, Espinosa and Prorok, 2001 or 30–78 GPa, Nix, 1989) as compared to bulk gold (78 GPa, Nix, 1989). Other mechanical properties such as yield stress have been found to be strongly related to film thickness (170 MPa for 0.5 ␮m thick Au films versus 50–55 MPa for 1.0 ␮m thick Au films) (Espinosa and Prorok, 2001) and crystallographic texture (Thompson, 1993). Furthermore, the electrical properties can be related to the film density, where films with greater density can exhibit bulk-like resistivity (Maissell, 1973). The mechanical and electrical properties of thin metal films can be tailored by controlling sputtering parameters Table 1 Surface treatment processes Process Etch S1 S2 S3 S4 None SP RIE RIE SP Pressure (Pa) – <1.4 27 <1.4 27 Power (W) Gas Time (min) – 250 100 250 100 – N2 O2 N2 O2 – 6.0 0.5 0.5 6.0 Reactive-ion etching (RIE) was conducted in an Oxford Plasmalab 80 Plus. Sputter etching (SP) was conducted in a Denton Discovery 18 sputter system immediately prior to metal deposition. In the case where both sputter and reactive-ion etching were performed, the reactive-ion etching was performed prior to sputter etching. such as deposition rate and film thickness. Both deposition rate and film thickness affect microstructural properties such as density and grain size. For instance, higher deposition rates tend to result in greater film density and larger grain size (Ohring, 2002). Each metal has its own power/rate relationship that is defined by the number of sputtered metal ions per bombarding argon ion (or sputtering yield) (Stuart, 1983). The substrate, platform, and source-metal target heat up during sputtering in relation to power and time (Lamont, 1979). Density and grain size have been shown to increase with greater film thickness and substrate temperature during metal deposition (Ohring, 2002). Higher substrate temperatures, directly related to higher deposition rates (Maissell, 1973), can also directly affect adhesion, nucleation, and microstructure of growing thin films. Metal deposition for this study involves several features: (1) setup/positioning of the loom inside the sputter chamber, (2) application of continuous rotation, and (3) deposition parameter setup (source-metal, power, session time, and number of repetitions). Due to geometrical constraints inside the Denton sputter system, flat and cylindrical looms were positioned inside the chamber with different methods. Flat looms were positioned such that the centroid of the loom was placed over the center of the substrate platform (Fig. 2A). For the cylindrical loom, a special holder was fashioned that would allow rotation about the loom’s radial axis without contacting the substrate platform. The cylindrical loom was placed in the holder and centered on the substrate platform. The cylindrical loom system was oriented such that one end was positioned equidistant from the Ti and Au sputter guns and the other end was below the Pt sputter gun (Fig. 2B). The three sputter guns were positioned equidistant from each other and oriented towards the center of the substrate platform at an angle 30◦ from normal to the platform surface. Rotation was supplied to the flat loom via rotation of the entire substrate platform at a speed of 60◦ /s (Fig. 2A). For the cylindrical loom, the substrate platform remained fixed; only the loom was rotated. In this case, a vacuum-sealed, translation feed-through was used to supply rotation at a speed of 24◦ /s (Fig. 2B). The deposition parameters were subdivided into six categories: (1) source-metal, (2) session power (where one session is defined as period between creation and extinction of the plasma), (3) session time, (4) number of session repetitions, (5) session energy (session time × session power), and (6) total sputter energy (sputter energy for entire process). For this study, each layer of metal was typically completed in one session; however, in order to minimize substrate temperatures, some processes incorporated multiple sessions per metal layer, where a cooling period of 3 min would follow each session. Process energy is reported as the sum of all session energies per process and is related to the amount of metal deposited. Table 2 summarizes the parameters involved for each of the S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 13 Table 2 Metal deposition processes Process Loom Metal Power (W) Time (min) Energy (kJ) M1 Flat Ti Au Pt 750 500 380 2 6 4 90 180 91.2 1 1 1 361 M2 Flat Ti Au Pt 750 250 250 2 4 2 90 60 30 1 2a 2a 270 M3 Flat Ti Au Pt 750 250 250 2.5 16 4 113 240 60 1 1 1 413 M4 Cylindrical Ti Au Pt 50 50 50 4 3b 6 12 9 18 1 20a 1 210 M5 Cylindrical Ti Au Pt 50 50 50 4 3b 6 12 9 18 1 18 1 192 M6 Cylindrical Ti Au Pt 250 100 – 0.5 15 – M7c Cylindrical Ti Au Pt 50 50 50 7.5 90 – 4 50d , 3b 3–6 12 150d , 9 9–18 Reps (#) Process energy (kJ) 1 1 – 97.5 1 1d , 18–60a 1 129–309 A “session” is defined as the period between plasma ignition and extinction while sputtering from a single metal source. The “Reps” category indicates how many sessions consisting of the same power/time parameters were completed. Session energy was calculated as powermetal × timemetal and process energy was calculated as the sum of energy from all sessions in a single process. a A break in vacuum was made to flip or reorient the loom. b A 3 min stoppage or “cooling” period was made immediately following each session. c M7 is a collection of several processes that were mainly used to help demonstrate the relationship between longitudinal resistance and process energy for metallized fibers. d A single process identical to that of M6 with the exception of session power and times. metal deposition processes which are referred to as “M#” where M stands for metal deposition and # stands for the specific process (e.g. M2 stands for metal deposition process #2). Even though rotation of the flat loom was intended to minimize the shadowing effect, a small portion of the fiber circumference on the bottom side of the loom was continuously shadowed. Therefore, the flat loom was manually flipped over midway through deposition of the Au and Pt layers for process M2. Due to the geometry of the samples, standard methods of measuring film thickness were not possible. Hence, we were not able to estimate deposition rate or film coating uniformity in a quantitative way. Based on several longitudinal resistance measurements taken at various levels along samples from cylinder looms, a non-uniformity in coating thickness was assumed. In an attempt to minimize this thickness non-uniformity, the cylindrical loom was reoriented end-for-end by breaking vacuum midway through deposition of the Au layer in the process M4. After deposition, several metallized fibers constructed with the process M4 were annealed at 45, 100, 130, 150, and 170 ◦ C for the purpose of determining the effect of temperature on longitudinal and interfacial impedance of the deposited metal film. 2.2.4. Insulation All medical-grade silicones used in this study were provided by Nusil Silicone Technology© . Three insulation processes were employed: (1) several coats of MED-6015 (an optically clear potting and encapsulating silicone elastomer) cured at 150 ◦ C for 2 h, (2) an adhesion coating of MED2-4013 (a fast cure silicone adhesive) with an additional three coats of MED-4210 (silicone elastomer) cured at 45 ◦ C for 12 h, and (3) an adhesion coating of MED2-4013 with an additional three coats of MED-4210 cured at 150 ◦ C for >30 min (Table 3). Silicone coatings were applied with a coating speed less than 5.5 cm/h. Vacuum deaeration was not conducted prior to curing, and may Table 3 Insulation of the various medical-grade silicones consisted of a dip-coating procedure at a controlled rate <5.5 cm/h Process Base layer silicone (type) Top layer silicone (type) Coats (#) Cure temperature (◦ C) Cure time (h) I1 I2 I3 MED-6015 MED2-4013 MED2-4013 MED-6015 MED-4210 MED-4210 2 3 3 100 45 150 2.0 12.0 1.5 14 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 account for observed pinholes. Each insulation process is referred to as “I#”, where I stands for insulation and # stands for the specific process (e.g. I2 stands for insulation process #2). The insulation processes have been ordered in terms of improving the adhesion performance. 2.2.5. polyLIFE construction Two groups of polyLIFEs were constructed to evaluate the effect of insulation and the specific insulation process on the mechanical fatigue properties of complete electrodes. The first group was constructed with surface treatment, metal deposition and insulation processes S4·M4·I2, and the second group was constructed with processes S4·M4·I3. In order to increase the tensile strength and electrical redundancy of the polyLIFE, several electrodes were assembled with three individually metallized fibers. These three metallized fibers were wound around each other and subsequently insulated with process I3. These polyLIFEs will be referred to as (S4·M4)R ·I3, where R indicates the metallized fiber redundancy. 2.2.6. Miniature wire LIFE construction As an alternative to the polyLIFE and the multi-stranded polyLIFE, two groups of solid conductor LIFE electrodes were constructed. One group of LIFEs was constructed from Teflon insulated, 25 ␮m diameter Pt-Ir wire (Lefurge et al., 1991). A second group of LIFEs was constructed from 14 ␮m diameter Au wire buttressed with two strands of bare Kevlar® and subsequently insulated with process I3. Throughout the following text, the LIFEs described in this section will be referred to as Pt-Ir LIFEs and Au LIFEs. approximate 1.5 cm long section of the metallized fiber to electrical tape that had been laid out on a flat surface, (3) pulling the fiber off the tape normal to the tape’s surface, and (4) measuring the final resistance of the metallized fiber. The fiber was adhered to the tape using a pressure of 206 ± 20 kPa (n = 20, pooled data) and pulled off the tape with 23 ± 7 mg of force (n = 20, pooled data). The adhering pressure and pull-off force was measured with a mass balance on two separate days (n = 10 for each day). Failure for the tape test was defined in the same manner as the wipe test. Samples from all groups were pooled together and subsequently tested in random order. Samples subjected to the tape test were also inspected for specific adhesion failure modes using a scanning electron microscope. 2.3.2. Metallized fiber longitudinal resistance and interfacial impedance The longitudinal resistance (/cm) of metallized fibers was tested using an ohm-meter. Interfacial impedance (k/mm) of the metallized fibers was estimated by immersing the fiber in a 4 mm diameter drop of saline and measuring the resulting impedance to a 1 kHz application of a 10 nA constant-current sine wave. The return electrode consisted of a stainless steel needle whose surface impedance was below the sensitivity of the measurement. Statistics used for these tests were two-tailed Student’s t-tests in which unequal variance was assumed. 2.3. Material properties characterization 2.3.3. Metal film fatigue resistance Fatigue resistance of the metallized fibers was tested by repeated cycling around a 500 ␮m diameter hypodermic needle (Fig. 3). Fatigue resistance was defined as the number of 2.3.1. Metal film adhesion Adhesion performance of the sputter-deposited metal layer was evaluated using methods that would challenge the interface with forces comparable to forces that would be expected during surgical handling and long-term implantation. Towards the end, we subjected the metallized fibers to a wipe test and a tape test. The wipe test consisted of approximately 5 cm length of the metallized fiber drawn between two dry fingers five times and measuring the resistance before and after the test. The applied pressure during the wipe tests was held as consistent from test to test as possible. Applied pressure was measured with a manometer on two separate days (n = 10 for each day). According to the manometer, the pressure used for the wipe tests was 22.7±1.0 kPa (pooled data). Failure on this test was defined as a measured longitudinal resistance greater than 5 k. Electrode impedances greater than 30 k result in poor signal to noise ratios. As a typical recording zone impedance ranged from 15 to 25 k, a maximum lead resistance of 5 k was chosen to maximize the potential signal to noise ratio. The tape test consisted of four steps: (1) measuring the fiber’s initial longitudinal resistance, (2) adhering an Fig. 3. Fatigue test setup. The fatigue test consisted of continuously wrapping and unwrapping the test electrode around an approximate 500 ␮m diameter hypodermic needle until electrode failure (where failure meant that the electrode was “non-conductive at rest”). Fatigue resistance was defined as the number of cycles until the test electrode failed. Each cycle was comprised of a 180◦ counter-clockwise rotation and a 180◦ clockwise rotation at 240◦ /s. Longitudinal resistance was evaluated by measuring the voltage drop (Ve ) across a test electrode placed in series with a 70 k resistor and a 25 V supply. During cycling a constant axial load of ∼3 g was applied to the test electrode with a micro-clip dangling from the electrode’s tip end. This micro-clip was also used to complete the current path. S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 cycles until the electrode’s longitudinal resistance became greater than 5 k (tests were terminated at 200 kilocycles). In order to measure the metallized fiber’s resistance, the fiber was placed in series with a 70 k resistor and a 25 V dc power supply. The resistance of the metallized fiber was measured at the end of each cycle when the cycled portion was not subjected to dynamic or static bending. The axial load consisted of approximately 3 g of micro-clip connected to the tip-end of the fiber. This micro-clip, in combination with a flexible seven-stranded stainless steel cable, was also used to complete the current path for the resistance measurements. Each cycle consisted of a +180◦ rotation followed by a −180◦ rotation at 360◦ /s. Rotation speed was chosen such that a 180◦ rotation occurred at least once every second, a value within the physiological rate of movement for humans. The diameter of the hypodermic needle was chosen so that the metallization layer was subjected to bending well in the plastic range (strain >2%) of each metal’s stress–strain curve. Due to the non-normal distribution of the data, the Mann–Whitney U-test was used to compare fatigue resistance between the tested groups. Variability within tested groups was evaluated with a two-tailed F-test. 2.3.4. Insulation adhesion and leakage impedance Adhesion performance of the insulation layer and leakage impedance were measured. Adhesion of the insulation layer was tested by the wipe and tape tests described above. The wipe test consisted of wiping saline wetted fingers along the insulated fibers and confirming adhesion with a light microscope and surface impedance measurements before and after the test. Insulation leakage impedance (M/cm) was measured by immersing a 4 mm length of the insulated fiber in saline and measuring the resulting impedance at 1 kHz. Failure on the wipe and tape tests were defined as a resulting leakage impedance of <0.1 M/cm (a value greater than two orders of magnitude above uninsulated surface impedances capable of recording neural activity and greater than three orders of magnitude greater than the longitudinal impedance of the metallized fibers). 2.3.5. Improved mechanical properties with insulation and redundancy Electrode toughness for this study was based on polyLIFE fatigue resistance, tensile strength, and long-term leakage impedance stability. The fatigue test was used to evaluate changes to fatigue resistance of single and redundant metallized Kevlar® fibers after application of the insulation layer. For comparison, Pt-Ir LIFEs and Au LIFEs were also subjected to the fatigue test. A uniaxial tensile test (described elsewhere) (McNaughton and Horch, 1996) was used to evaluate changes to tensile strength. The flexibility of the multi-stranded polyLIFE was measured as described elsewhere (McNaughton and Horch, 1996) and compared to the flexibility of the Pt-Ir LIFEs in order to determine the relative change in stiffness due to redundancy. Long-term leakage 15 impedance (M/cm) was tested on several multi-stranded polyLIFEs by soaking them in saline for a 6-month period. Leakage impedance measurements were conducted on day 0, 10, 90, 150, and 180 of the soak test. Leakage impedance was estimated by immersing the fiber in a 2 mm diameter drop of saline and measuring the resulting impedance (at 1 kHz). Fatigue resistance differences were tested for statistical significance (α = 0.05) with the Mann–Whitney U-test. Ultimate tensile strength differences were tested for statistical significance (α = 0.05) with a two-tailed, Student’s t-test. 3. Results 3.1. Modified polyLIFE material properties 3.1.1. Metal film adhesion It has been found that adhesion performance is directly affected by certain mechanical properties of the metal film such as internal residual stress, thickness, and Young’s modulus. Barenblatt (1962) found that adhesion performance is negatively affected by thin films with larger residual stresses, greater thickness, and lower values of Young’s modulus. Wipe tests and tape tests were conducted on six groups of metallized fibers (n = 7 for each group), manufactured with processes S1·M1, S3·M2, S3·M3, S4·M4, S2·M5, and S3·M6. Except for one sample from group S1·M1, all groups passed the wipe test. However, groups S1·M1, S3·M2, and S3·M3, failed the tape test 100, 43, and 29% of the time, respectively. Among the groups that passed the tape test, the cylindrical loom groups experienced a significant lower mean resistance change (percentage ± S.E.) due to the tape test than the flat loom groups (P < 0.05 for pair-wise comparison of fibers in groups S3·M2, S3·M3, S4·M4, S2·M5, and S3·M6 whose longitudinal resistance had changed 14 ± 10%, 66 ± 27%, 1 ± 0.6%, 2 ± 0.7%, and 0 ± 0.1%, respectively, due to the tape test). In general, metallized fibers from flat looms faired worse on the wipe and tape tests than those from cylindrical looms. Likewise, metallized fibers that had been fabricated with higher process energies also tended to perform worse on the wipe and tape tests. Fig. 4 shows scanning electron micrographs of an exemplary fiber from each group subjected to the tape test. These micrographs reveal that the tape test caused micro-cracks in the metal layer. These cracks appear to have propagated along the metal layer’s grain boundaries. Fibers from flat looms showed tendencies to crack not only circumferentially, but longitudinally, leading to the failure and peeling of large (>50 ␮m long) sections of the metal film. However, fibers mounted on flat looms were always metallized with high process energies (>270 kJ), which could have resulted in thicker films and higher substrate temperatures during the metal deposition process. Damage due to the wipe or tape tests on fibers from groups S4·M4, S2·M5, and S3·M6, appeared as circumferential or spiral fractures that did not result in a significant increase in longitudinal resistance. 16 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 Fig. 4. Scanning electron micrographs of metallized fibers subjected to the tape test. Metallized fibers from flat looms where delamination adhesion failures are visible: (A) S1·M1, (B) S3·M2, (C) S3·M3. Metallized fibers from cylindrical looms where no cracking or adhesion failure is visible: (D) S4·M4, (E) S2·M5, and (F) S3·M6 (gray bars: 10 ␮m). From the micrographs, it was difficult to determine in some cases whether complete delamination was prevented by better adhesion of the metal to the Kevlar® or by better cohesion of a non-adhered metal tube surrounding the Kevlar® fiber. 3.1.2. Metallized fiber longitudinal resistance and interfacial impedance Longitudinal resistance and interfacial impedance were measured on metallized fibers from groups manufactured with processes S1·M1, S3·M2, S3·M3, S4·M4, S2·M5, and S3·M6 (n = 7 for each group). Longitudinal resistance was also measured from a single metallized fiber from each of the metal deposition processes S2·M7. Longitudinal resistance was found to be related to process energy used to metallize these fibers (Fig. 5A). The interfacial impedance was found to be logarithmically related (r 2 = 0.997) to the session time and session energy for Au deposition (Fig. 5B). 3.1.3. Annealing effects on longitudinal resistance and interfacial impedance While no relationship was found between the annealing temperature and the longitudinal resistance (Fig. 6A), annealing caused both positive and negative effects on S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 Fig. 5. (A) Process energy vs. longitudinal resistance. For details on processes M1–M7, see Table 2. M7∗ is comprised of several separate processes conducted solely to help demonstrate the process energy/longitudinal resistance relationship. (B) Session time × session energy vs. interfacial impedance. A logarithmic relationship exists between (session time × session energy) and interfacial impedance. This graph suggests that interfacial impedance is related to the time taken to deposit a certain amount of metal. interfacial impedance (Fig. 6B). For instance, the mean surface impedance was reduced by factors of 33% (P < 0.01) and 47% (P < 0.01) for annealing temperatures of 45 and 100 ◦ C, respectively, as compared to the control values. However, the mean surface impedance was increased by 350% (P ≪ 0.01) and 930% (P ≪ 0.01) for annealing temperatures of 150 and 170 ◦ C, respectively, as compared to control values. 3.1.4. Metal film fatigue resistance Fatigue tests were conducted on groups of metallized fibers manufactured with processes S1·M1 (n = 8), S3·M2 (n = 10), S3·M3 (n = 6), S4·M4 (n = 7), S2·M5 (n = 4), and S3·M6 (n = 4). Fig. 7 shows median fatigue resistance (±S.E.) from each group and the median fatigue resistance from Pt-Ir LIFEs (n = 12). Regardless of process energy used for metal deposition, fibers mounted on flat looms did not demonstrate significantly better fatigue resistance (P > 0.05) than Pt-Ir wire. However, fibers mounted on cylindrical looms demonstrated a fatigue resistance of greater than 20 (P ≪ 0.01), 250 (P < 0.01), and 490 (P < 0.01) times 17 Fig. 6. (A) Annealing temperature vs. longitudinal resistance for metallized fibers. (B) Annealing temperature vs. mean interfacial impedance (±S.E.) for metallized fibers. Metallized fibers made with process M4 were annealed at 45, 100, 130, 150, and 170 ◦ C post-deposition, and subsequently evaluated in terms of longitudinal and interfacial impedance. that of the Pt-Ir wire for groups S4·M4, S2·M5, and S3·M6, respectively. In terms of flat loom mounted fibers, fatigue resistance was not correlated with either process energy (P > 0.34 for S1·M1 versus S3·M3 and P > 0.63 for S3·M2 versus S3·M3) or attempts to improve circumferential thickness uniformity (P > 0.57 for S1·M1 versus S3·M2). For fibers mounted on cylindrical looms, fatigue resistance was not correlated with process energy (P > 0.2 for group S2·M5 versus S3·M6), but was negatively correlated with attempts to improve the longitudinal thickness uniformity (P < 0.01 for group S4·M4 versus S2·M5). However, fibers mounted on cylindrical looms not only improved the metal film’s fatigue resistance, by a factor of 140 (S2·M5, P < 0.05) and 280 (S3·M6, P < 0.05), compared to fibers mounted on flat looms, but dramatically decreased the within-group variability for groups S2·M5 (P < 0.05), S3·M6 (P < 0.05), and S4·M4 (P < 0.01), respectively versus pooled flat loom data. Fig. 8 shows scanning electron micrographs of the representative metallized fibers that had been subjected to fatigue tests. The micrographs from flat looms demonstrate both longitudinal and circumferential cracking, as well as large delaminations. Fig. 8A shows an extreme example from group S1·M1 where the metal film failed after two order of magnitude more cycles than would have been expected according to the group’s median value. The question here is why this 18 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 Fig. 7. Fatigue resistance of metallized fibers from groups S1·M1, S3·M2, S3·M3, S4·M4, S2·M5, and S3·M6 compared to fatigue resistance of Pt-Ir wire. fiber remained conductive when it is obvious from inspection of the micrograph that no continuous metal layer exists. In fact, this fiber was conductive even though only islands existed next to large patches of delaminations. Presumably, electrical conduction was possible through contact between the metal islands when the fiber was in the straightened condition (i.e. at the time the resistance measurement was taken during the fatigue test). It is believed that this behavior seen with metallized fibers from flat looms causes the extreme variability in fatigue tests within groups as compared to the Pt-Ir LIFEs (P ≪ 0.01 two-tailed F-test). Compared to fatigued fibers from cylindrical looms, the length of uncracked metal was greater with fibers from flat looms. For example, fatigued fibers from group S4·M4 showed circumferential cracks that were spaced at substantially closer levels than the circumferential cracks displayed in fatigued fibers from Fig. 8. Scanning electron micrographs of metallized fibers subjected to the fatigue test. Metallized fibers from flat looms where both longitudinal and circumferential cracks are visible: (A) S1·M1 failed after 153 kilocycles, (B) S3·M2 failed after 7.7 kilocycles, (C) S3·M3 failed after 2.2 kilocycles. Metallized fibers from cylindrical looms where only circumferential cracks are visible: (D) S4·M4 failed after 9.3 kilocycles, (E) S2·M5 failed after 25 kilocycles, and (F) S3·M6 test stopped before failure at 200 kilocycles (gray bars: 10 ␮m). S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 19 Fig. 9. Fatigue resistance of polyLIFEs. Results are from two groups of miniature wire LIFE electrodes (Pt-Ir LIFE and Au LIFE) and four groups of polyLIFEs (S4·M4, S4·M4·I2, S4·M4·I3, and (S4·M4)R ·I3). group S3·M2 where the process energy used for each group was similar (210 and 270 kJ, respectively). 3.1.5. Insulation adhesion performance and leakage impedance In order to choose an effective combination of silicones for insulation of the polyLIFE electrode, adhesion performance and leakage impedance were tested on metallized fibers insulated with processes I1, I2, and I3 (Table 2). Results from the wipe test indicated that insulation process I1 provided poor adhesion and process I2 demonstrated variable adhesion. This variable adhesion could be in part due to an incomplete cure of the MED2-4013 layer. Nusil had suggested a cure temperature greater than 100 ◦ C, but a cure temperature of 45 ◦ C was used to avoid interfacial impedance increases of the metal film. Metallized fibers coated with insulation process I3 showed no decrease after the wipe test (n = 6) or tape test (n = 6). Since only process I3 passed the wipe and tape tests, the result suggest that a complete cure (curing the adhesion silicone MED2-4013 above 100 ◦ C) is necessary for improved adhesion. According to the leakage impedance tests, insulation processes I2 and I3 were sufficient to produce leakage impedance values greater than 2 M/cm. For insulation process I3, pinholes were discovered upon inspection with the scanning electron microscope when fewer than three coatings of MED-4210 were applied. When applying process I3, the insulation layer was estimated to be 3–5 ␮m thick as determined from scanning electron micrographs. 3.1.6. Improved mechanical properties of polyLIFEs due to insulation and redundancy Two groups of miniature wire LIFE electrodes and four groups of polyLIFEs were subjected to the fatigue test: (1) Pt-Ir LIFEs (n = 12), (2) Au LIFEs (n = 7), (3) non-insulated polyLIFEs made with processes S4·M4 (n = 7), (4) insulated polyLIFEs made with processes S4·M4·I2 (n = 7), (5) insulated polyLIFEs made with processes S4·M4·I3 (n = 13), and (6) insulated, redundant polyLIFEs made with processes (S4·M4)R ·I2 (n = 7). Fig. 9 shows median fatigue resistance (±S.E.) of the two groups of miniature wire LIFEs and all tested polyLIFEs. All insulated polyLIFE groups demonstrated a fatigue resistance of greater than 40 times than that of Pt-Ir LIFEs (P ≪ 0.01) or Au LIFEs (P < 0.01). The fatigue resistance of non-insulated polyLIFEs was improved by 6.8 times (P < 0.05) and 3.7 times (P < 0.01) when insulated with processes I2 and I3, respectively. No significant relation was found between cure temperature (group S4·M4·I2 versus S4·M4·I3) and fatigue resistance (P > 0.36). While insulated, redundant polyLIFEs improved the median fatigue resistance of non-insulated polyLIFEs by greater than 2.5 times (P < 0.05 for group (S4·M4)R ·I2 versus group S4·M4), their fatigue resistance was not significantly different than that of insulated polyLIFEs (P > 0.73 for group (S4·M4)R ·I2 versus S4·M4·I3). The insulated, redundant polyLIFEs’ ultimate tensile strength (1050 ± 187 mN, mean ± S.D.) was a factor of approximately four times greater than insulated polyLIFEs (P < 0.0001 group (S4·M4)R ·I2 versus S4·M4·I3) and approximately three times greater than Pt-Ir wire (P < 0.0001). Furthermore, the insulated, redundant polyLIFE remained more flexible (60 ± 19 nN m) than the Teflon-coated Pt-Ir wire (330 ± 80 nN m) by a factor of 5 (P < 0.01). Fig. 10 demonstrates that long-term leakage impedance remained stable over the 6-month soak test, with all samples (n = 10) exhibiting values >5 M/cm. Mean active zone length (±S.D.) of the insulated, redundant polyLIFEs was 1.01 ± 0.2 mm, as measured from scanning electron micrographs. Some longitudinal wicking of the silicone between metallized fibers (perhaps due to capillary action) was also observed, which could account for some of the measured variability in active zone length. 20 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 Fig. 10. Mean leakage impedance (±S.D.) during a 6-month saline soak test of polyLIFEs (n = 10). The polyLIFEs were manufactured with processes (S4·M4)R ·I3 (i.e. three metallized fibers made with processes S4·M4 were first wound around each other and subsequently insulated with process I3). 4. Discussion 4.1. Metal–polymer adhesion Adhesion performance of the metal sputtered onto Kevlar® fibers mounted on flat looms was significantly improved by reactive-ion etching with O2 , as demonstrated by the wipe and tape tests. Other studies have shown similar metal–polymer adhesion performance improvements not only by plasma treatments, but also by substrate heating (during and after deposition, Baglin et al., 1991; Faupel et al., 1999) and/or inclusion of a “glue” metal layer such as Ti or Cr (Lee, 1991). Covalent bonding of Ti to polyimide substrates (Ti-C) with evaporation deposition has shown to be promoted at increased coverage (Ohuchi and Fellich, 1986) or after high energy (2 keV) Ar+ ion bombardment of the substrate prior to metal deposition (Bodö and Sundgren, 1988). While we did not attempt direct adhesion of the Au layer to Kevlar® , our application of the Ti (“glue”) layer without additional plasma treatment was shown to be insufficient. Specifically, tenuous adhesion may have been achieved between the Ti and Kevlar® during sputtering, but compressive residual stresses achieved in the film during cooling may have led to adhesion failure (as evidenced in scanning electron micrographs of flat-loom mounted fibers). This residual internal stress was most likely caused by thermal expansion coefficient differences; the transverse thermal expansion coefficient of Kevlar® (60 × 10−6 ⑀/◦ C) is nearly six times that of any of the deposited metals (8.8, 14.1, and 9.1 × 10−6 ⑀/◦ C for Ti, Au, and Pt, respectively). Differential thermal expansion of the metals and polymer during metal deposition may have caused significant compressive internal residual stresses within the metal film after cooling, thus contributing to the poor adhesive and mechanical performance observed from fibers from the flat loom. In addition to thermal residual stresses, Barenblatt (1962) suggests that adhesion failure of thin film metal is affected by other factors such as Young’s modulus and film thickness. In our case, it is not likely that Young’s modulus had any effect on adhesion performance as it has been shown to be independent of film thickness (Espinosa et al., 2003). Likewise, it is unlikely that increased thickness is the cause of the adhesion failure exhibited in our system. For instance, group S3·M2 demonstrated worse adhesion performance than group S3·M3 (a group for which higher process energies presumably resulted in thicker metal films). Furthermore, fibers from group S3·M3 likely experienced a higher temperature and thermal expansion during sputtering than fibers from S3·M2. Therefore, we have attributed the poor adhesion performance of fibers in group S3·M2 to increased thermal residual stresses in the metal film due to the multiple heating and cooling cycles (caused by breaking vacuum and flipping the loom to promote circumferential thickness uniformity). We believe that the improved adhesion performance for metallized fibers from flat looms can be attributed to the minimization of thermal residual stress in combination with plasma surface treatments. Adhesion performance of fibers mounted on cylindrical looms was excellent, as evidenced by minimal changes in resistance in the tape test. However, although no delamination was visible in the scanning electron micrographs from the tested fibers, we cannot conclude that adhesion of the metal to the polymer was responsible for the improved performance. Although no correlation was found between adhesion and any of the surface treatment processes (S2, S3, or S4) when used with cylindrically mounted fibers, this result could have been due more to an improvement in the mechanical properties of the deposited metal film. For instance, rotation of the cylindrically mounted fibers would have minimized the oblique angle of incident atoms around the circumference, perhaps accounting for the absence of S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 longitudinally directed cracking due to the tape test. One might propose the thinner Ti layer was responsible for the improved adhesion performance; however, this trend was not seen with fibers from the flat loom. The lower session energies experienced by the cylindrically bound fibers should have contributed to the lower substrate temperatures, which in turn may have resulted in lower residual stresses and improved adhesion performance. Various “glue” layers, such as Ti, adhere best to polymers when their surfaces have been oxidized (Ohring, 2002). Reactive-ion etching and sputter etching worked equally well in improving the adhesion performance for fibers bound to cylindrical looms; however, we suggested using reactive-ion etching with an O2 plasma over sputter etching as the latter process affects surfaces in a non-specific way and requires high process power, which may lead to excessive substrate heating. 21 The lower interfacial impedances observed in fibers made with higher process energies could be attributed to increased surface area; caused by increased surface roughness or larger dimensions. The larger diameters observed with flat loom-bound fibers could partially account for the lower interfacial impedance. However, the group S3·M6 demonstrated a much lower interfacial impedance than either group S4·M4 or S2·M5 even though they had comparable fiber diameters. Likewise, increased surface roughness must also be excluded as a primary cause of increased interfacial impedance as group S3·M6 is observed to be much smoother than any of the flat loom groups. Metallized fibers from group S3·M6 exhibited a similar surface roughness to the two other cylindrical loom groups (by visual inspection of scanning electron micrographs), but demonstrated the second lowest interfacial impedance. These results suggest that the surface impedance of these sputter-deposited metal layers depends more on film density than on surface roughness. 4.2. Electrical properties of sputter-deposited metal 4.3. Annealing effects on electrical properties Previous work has shown that low film densities, high dislocation content, greater amount of trapped gases, voids, and defects lead to poor electrical properties (Ohring, 2002). In films containing a predominantly columnar microstructure (typically observed in films deposited at low pressures, low substrate temperatures, and oblique angles of incident atoms), large voids and higher concentrations of precipitates can also be found in the intercolumnar regions, which in turn contribute to poor electrical properties. As our films are deposited at low gas pressures, relatively low substrate temperatures (TS /TM < 0.3, where TS is the substrate temperature and TM is the melting temperature of the sputtered metal) and highly oblique angles of incidence, it can be assumed that our metal deposition processes result in low density films with columnar microstructure (qualitatively observed in scanning electron micrographs of our metal films). Only process energy had any significant effect on longitudinal resistance, presumably due to the increased amount of metal deposited with higher process energies, with no changes evident due to post-deposition annealing. However, interfacial impedance was logarithmically related to the session energy, indicating that it is important to control the time taken to sputter a given amount of metal. Longer times may have allowed the substrate temperature to rise thereby increasing the incident atoms lateral mobility, which resulted in greater film density and lower interfacial impedance. For the purpose of recording neural activity, previous studies have suggested that interfacial impedances of the recording/stimulating zone should be lower than 30 k/mm (Malagodi et al., 1989). Therefore, we suggest avoiding short session times combined with low session energies (<9 kJ). The lowest values for interfacial impedance were found with session times >10 min in combination with higher session energies. It would appear that time was a more important factor than session energy, indicating that perhaps substrate temperatures were related more to time than to power. Interfacial impedance was affected by annealing temperatures. In the case where high interfacial impedance was observed after metal deposition, annealing resulted in even higher interfacial impedances. Since only one group was tested for annealing effects, it is not known if a similar trend of increasing interfacial impedance with increasing annealing temperature would be found with the other groups. The exact cause of the increasing interfacial impedance due to higher annealing temperature is unknown, but potential causes could be microstructural or chemical in nature. For instance, cracking due to differential thermal expansion between the metals and polymer (Owusu-Boahen and King, 1998) has been shown to result in increased surface roughness and presumably lower interfacial impedance. Likewise, elevated temperatures known to cause grain growth, and hence removal of grain boundaries, would have resulted in densification of the film. However, grain growth at substrate temperatures below 0.3 (TS /TM ) is unlikely (Hentzell et al., 1984). Therefore, surface cracking is the most likely cause of decreased interfacial impedance between 22 and 130 ◦ C. At temperature higher than 130 ◦ C, increased interfacial impedance could have been caused by chemical modification of the metal surface. Chemical modification of the surface could be directly related to diffusion characteristics of the metal film. For thin metal films, the diffusion coefficient can be 1015 times greater than in bulk due to the high density of defects, grain boundaries and vacancies, making diffusion times through the small dimension films characteristically short. In thin gold films with ∼500 nm grains and at low temperatures (<0.3 TS /TM ), diffusion tends to occur mainly through the grain boundaries (Gupta and Asai, 1974). Any precipitates or oxide compounds formed at the grain boundaries can out-diffuse, thus chemically modifying the surface (Ohring, 2002). This situation is amplified 22 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 by the larger surface-to-bulk ratios involved in thin films. Although no proof exists that would directly point to any of the specific mechanisms mentioned, it is our belief that diffusion is in some way related to the increased interfacial impedance observed after heating above 130 ◦ C. Surface analysis tools such as electron spectroscopy for chemical analysis (ESCA) or secondary ion mass spectrometry (SIMS) should help reveal the exact compositional changes caused by post-deposition heating. 4.4. Metal film fatigue resistance The greatest fatigue resistance was found with fibers that had been bound to cylindrical looms and metallized with lower process energies. Fibers bound to flat looms and metallized with higher process energies, not withstanding some extreme cases, tended to show worse performance and demonstrated higher variability. Additionally, attempts to create circumferential or longitudinal thickness uniformity either made no difference (flipping of flat loom) or were detrimental (end-for-end reorientation of cylindrical loom) to the metal film’s fatigue resistance. The improved fatigue resistance of metallized fibers bound to cylindrical looms could be explained by their presumably lower film thickness, higher density, or crack propagation morphology. As mentioned earlier, thinner films exhibit higher yield stresses, at the cost of a more brittle behavior. However, lower film thickness, as well as lower incident atom mobility (assumed to be the case for metal deposition with lower session and process energies of the cylindrical bound fibers) has a tendency to negatively affect density. Therefore, it is not likely that density had a significant effect on fatigue resistance, at least for group S2·M5. In terms of crack morphology, crack initiation is promoted by rough surfaces, corrosive environments, and mechanical failure along grain facets or boundaries (Suresh, 1991). For instance, it has been shown that the microscopic valleys or notches of irreversibly roughened metal surfaces cause stress concentrations and additional crack nucleation (Wood, 1958). Likewise, corrosive environments such as oxygen (Gough and Sopwith, 1932; Thompson et al., 1956) or seawater (Suresh, 1991) tend to promote surface roughening, by reacting with freshly revealed material or metal protrusions at the surface, caused by plastic deformation. Once the protruding metal has reacted or oxidized, reverse movement back within the bulk is impeded, resulting in surfaces with greater roughness. Grain boundaries tend to separate highly misoriented grains, and crack initiation is promoted when the deformation within a single grain is directed at the intersection of two neighboring grains located at the film surface (Porter and Levy, 1960). Additionally, preferential oxidation at grain boundaries (Duquette, 1979) and void formation around precipitates found in grain boundaries (Vasudévan and Doherty, 1987) can cause microscopic stress concentrations due to notch development or intergranular cavitation. These stress concentrations can in turn initiate cracks and promote crack propagation along the weak grain boundaries. Perhaps the most likely cause for the improved fatigue resistance of groups S2·M5 and S3·M6 are the relatively smoother surfaces observed (as compared to flat loom-bound fibers) and more densely packed grains. In addition to the inhibited crack initiation and propagation morphology, improved fatigue resistance could be attributed to lower film thickness, improved adhesion performance, and lower residual stresses in the cylindrically mounted fibers. A qualitative comparison of the effects of thickness on metal film morphology due to externally imposed bending is provided in Appendix A. 4.5. Insulation adhesion performance and leakage impedance Insulation process I3 provided the best adhesion consistency and leakage impedance. The improved adhesion was due to the use of MED2-4013 as a base or “glue” layer as it contained adhesion promoters. However, use of multiple cycles of >100 ◦ C is not recommended due to the detrimental effect on the interfacial impedance, mechanical integrity, and adhesion of the metal film. Since a complete cure of MED2-4013 is only possible at or slightly above 100 ◦ C, one might attempt to obtain a silicone that includes adhesion promoters, but with a low temperature cure cycle. We have obtained a low temperature cure version of MED2-4013 in the past, but the work time of this silicone was insufficient for our needs. It was also found that at least three layers of MED-4210 were necessary to prevent pinholes. 4.6. Improved mechanical properties of the polyLIFE due to insulation and redundancy Our study demonstrates improved fatigue resistance by the application of silicone or fiber redundancy. The silicone layer may inhibit crack formation at the metal film surface by the mechanism of either smoothing the surface (a phenomena known to increase the fatigue life of metals) (Basinski et al., 1983) or creating a passivation layer that inhibits continual surface oxidation, thus allowing reversibility of plastic deformations. Alternatively the silicone might hold islands of metal in tight physical contact. The redundancy technique most likely improves fatigue resistance by the mechanism of probability. For instance, the greater number of metallized fibers used for a single electrode improves the chances that one will last to the measured fatigue limit. Observations from scanning electron micrographs of three metallized fibers that had been wound around each other, but not insulated, demonstrate fiber-to-fiber abrasion, resulting in significant damage to the metal film. Fatigue test results demonstrate that the damage resulting from the act of winding the metallized fibers did not negatively affect fatigue resistance of insulated, redundant polyLIFEs. However, no significant positive effect on fatigue resistance was observed either. The most significant effect of redundancy was an increase to the electrode’s tensile strength without S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 a substantial decrease in flexibility, the original polyLIFEs’ primary design advantage over the Pt-Ir LIFEs. 5. Conclusions We have demonstrated improvements to the metal film adhesion, fatigue resistance, insulation adhesion consistency, leakage impedance, and mechanical properties of finished polyLIFE designs. Metal adhesion was improved both by reactive-ion etching of the bare Kevlar® fibers, adjustments to the geometry, and rotation of the fibers during sputtering. The mechanical quality of the deposited films was better for fibers that were bound to cylindrical looms that were axially rotated. It is presumed that this type of rotation and mounting helped minimize the abundance of weak intergranular regions. We feel that crack propagation along weak grain boundaries is the main failure mechanism in loss of adhesion and fatigue failure. Cylindrically bound fibers also performed better on fatigue testing due to thinner films, and minimization of residual internal stresses caused by differential thermal expansion. Electrical properties such as longitudinal resistance were related to the process energy for metal deposition, while the interfacial impedance was found to be logarithmically related to the session energy. Although high temperatures did not affect longitudinal resistance, they did cause significantly elevated interfacial impedances; a situation we would like to avoid in order to ensure good recording properties. The insulation adhesion consistency and leakage impedance was improved with process I3, however the high temperatures necessary to cure the adhesion layer may be detrimental to the surface impedance of the metal film. Application of the insulation layer caused greater fatigue resistance for the finished polyLIFEs, as might have been expected; surprisingly, redundancy did not improve fatigue resistance. However, we feel that the improved tensile strength of the multi-fiber polyLIFE design provides sufficient rationale to recommend redundancy for chronic applications over the single fiber design. Alternative flexible intrafascicular electrodes could involve silicon- or polyimide-based substrates, or the use of intrinsically conductive fibers. Since the silicon- and polyimide-based interfaces are founded on time-tested micromachining techniques they have an advantage over hand made electrodes in that one has the ability to quickly modify designs (fast prototyping) and can assert precise control over active zone size and repeatability. For silicon to obtain any degree of flexibility, the dimensions must be reduced to such a degree that the final product is mechanically fragile. A silicon-based ribbon cable developed for high flexibility demonstrated successful function after chronic implantation of up to 1 year (Hetke et al., 1994), but this interconnect cable was not designed with the highly mobile peripheral nerve in mind and is thought by the authors to be too stiff and brittle to be practical for intrafascicular application 23 (Najafi and Hetke, 1990; Najafi et al., 1990; Petersen, 1982). Several highly flexible polyimide-based devices have been developed for cortex (Boppart et al., 1992; Rousche et al., 2001) or peripheral nerves (González and Rodrı́guez, 1997; Rodrı́guez et al., 2000; Stieglitz and Beutei, 1997), but are not directly applicable to intrafascicular peripheral nerve recording/stimulation, or rely on the regeneration of nerve stumps. Mechanically, the metal layers of these devices have demonstrated resistance to fatigue, but seem to have been tested only in the elastic strain range. Therefore, in addition to the geometrical differences, a direct comparison of fatigue resistance between metal layers deposited on planar polyimide substrates versus Kevlar® fibers was not possible. To the authors’ knowledge, only one polyimide-based longitudinal intrafascicular device has been developed, but it has somewhat larger dimensions and higher interfacial impedance (Yoshida et al., 2000). As a substrate material, polyimide possesses some advantages over Kevlar® in that its thermal expansion coefficients more closely match those of the deposited metals. However, polyimide’s tensile strength and Young’s modulus (55.8 MPa and 3.2 GPa, respectively) are substantially lower than Kevlar’s® (3500–3600 MPa and 59–124 GPa, respectively) (Wen, 1996), and consequently, devices with similar dimensions to our 12 ␮m diameter fibers would not likely be strong enough in tension to be inserted within fascicles. Metal layers on polyimide substrates have been applied with evaporation techniques, generally known to give greater film thickness variability, and worse mechanical and adhesion performance. This thickness variability and worse mechanical performance is primarily due to “spitting” and lower energy of impinging atoms as compared to sputtering (Ohring, 2002). Intrafascicular electrodes made from intrinsically conductive fibers could potentially combine high flexibility with near infinite fatigue resistance. Carbon fibers have demonstrated sufficient intraneural recording capabilities, but their brittle nature prohibits chronic implantation (McNaughton and Horch, 1994). On the other hand, intrinsically conductive polymers (polyaniline, polythiophene, polypyrrole) with sufficiently low longitudinal resistance, have been found to lack sufficient tensile strength for our application (Andreatta et al., 1989; Hsu et al., 1993; Moulton and Smith, 1992; Nemoto and Marks, 1991). Future development of metal clad Kevlar® fibers would involve further adjustment to the surface treatment or metal deposition processes as well as modification of redundancy methods. Our reactive-ion etch time could be increased in order to more extensively roughen and oxidize the surface to promote better adhesion performance. However, extremely long etching times could detrimentally affect the fiber’s tensile strength and its integrity at the interfacial region. This weakening effect has previously been seen with wet chemical etching processes of long duration (Sheu et al., 1994). In order to increase the density of the deposited metal film, a negative substrate bias greater than −200 V could be applied. The negative bias changes the electric field near the 24 S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 substrate, thus enhancing the flux and energy of charged incident atoms (Ohring, 2002). The increased energy of incident atoms is thought to improve adhering atoms lateral mobility and promote resputtering that can remove adsorbed gas atoms and eliminate the phenomena of void formation common at low pressures and substrate temperatures. With the use of bias sputtering, a myriad of metal film properties can be tailored such as increased density and step coverage, decreased resistivity, modification of hardness and residual stresses, modification of columnar microstructures, and improved adhesion performance. Longitudinal coating thickness uniformity could be improved by combining phasic linear translation with the long axis rotation already achieved with the cylindrical loom. This phasic longitudinal motion may also minimize the shadowing effect due to oblique angles of incident atoms experience at either end of the fiber. The interfacial impedance of the films might be minimized by adjusting the session time and energy. Finally, one might enjoy better fatigue resistance with multi-fiber polyLIFE designs if the fibers were not wound around each other, thus avoiding excessive metal-to-metal abrasion. Fig. 11. Effects of metal film thickness on deformation behavior. Due to externally imposed bending, (a) thinner films show brittle fracture, while (b) thicker films show ductile deformation. The greater film thickness in part b is also evidenced by the larger grains. Fig. 12. Effects of thermal residual stresses on metal film integrity. Extremely long session times result in high substrate temperatures and subsequent thermal stress. Due to externally imposed bending of similar thickness films, (a) thermally stressed films demonstrate catastrophic failure, and (b) thermally unstressed films demonstrate ductile deformation. For parts a and b, the session times for Au deposition were 50 and 3 min, respectively. S.M. Lawrence et al. / Journal of Neuroscience Methods 131 (2003) 9–26 Acknowledgements This study was supported by NINDS of the US National Institutes of Health. Appendix A Metal film mechanical properties are directly affected by both metal film thickness and session time, as observed from scanning electron micrographs (Figs. 11 and 12). Even though gold is commonly thought to be a highly ductile material, brittle behavior is seen in thinner films (Fig. 11a). As metal film thickness increases, plastic or ductile deformation precedes ultimate failure (Fig. 11b). During microbeam load-deflection experiments, Espinosa et al., found this transition from brittle to ductile deformation behavior occurred in beams between 0.5 and 1.0 ␮m thick (Espinosa et al., 2003). Fig. 12a shows the effects of residual thermal stress on film integrity. Presumably, an excessively long session duration resulted in a high substrate temperature. The differential expansion and contraction of the involved materials caused excessive thermal stress that, upon bending, caused catastrophic failure in the metal film. With much shorter session times, substrate heating is minimized during sputtering and this failure mode can be avoided (Fig. 12b). With the aid of scanning electron microscopic analysis, we can tailor the metal deposition process to achieve either ductile or brittle failure (according to our specific needs) and prevent residual thermal stresses. References Akin T, Najafi K, Smoke RH, Bradley RM. A micromachined silicon sieve electrode for nerve regeneration applications. IEEE Trans Biomed Eng 1994;41:305–13. Andreatta A, et al. Electrically conductive polyblend fibers of polyaniline and p-phenylene terephthalamide. Polym Commun 1989;31:275–8. Baglin JEE. Interface design for thin film adhesion. In: Lee L-H, editor. Fundamentals of adhesion. New York: Plenum Press; 1991. p. 363–82. Baglin JEE. 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