Metal Oxide Nanotubes and Photo-Excitation Effects: New Approaches for Low-to-Intermediate Temperature Solid Oxide Fuel Cells

Investigators Paul C. McIntyre, Associate Professor, Materials Science & Engineering; Michael Shandalov, Postdoc, Materials Science & Engineering; Nana Ayensu, Cynthia Ginestra, Irene Goldthorpe and Andy Lin, Graduate Researchers, Stanford University. Abstract

During the second year of this project, we made advances in 1) atomic layer deposition (ALD) of yttria-stabilized zirconia (YSZ) solid state electrolytes with controlled composition; 2) characterization of nanoscale size effects on crystalline phase stability of fluorite structure metal oxides; 3) ALD of La(Sr,Co)O3 SOFC electrodes. In addition, we completed construction of a microfabricated fuel cell testing apparatus which will allow us to determine the through-thickness ionic transport properties of both planar and nanotube ALD-grown YSZ membranes. These new capabilities compliment the more sophisticated electrochemical measurements being done in the laboratory of our collaborator, Prof. Shriram Ramanathan, at Harvard University. Construction of a new ALD reactor specifically for deposition of LSCO electrodes using commercially available ALD precursors that have relatively low volatility was also undertaken this year.

Introduction The goal of this project is to enhance the power density and low-temperature

efficiency of solid oxide fuel cells manufactured by atomic layer deposition. These enhancements will be achieved by engineering the morphology of the electrolyte at the nanoscale, and by studying the effects of photo-excitation on the catalytic properties of the membrane surface and on the concentration of electronic carriers and ionic defects in the thin film SOFC structure.

The efficiency of a fuel cell is limited by the loss mechanisms inherent to its operation. The power density of a fuel cell is limited by the area of its electrolyte membrane. These two operational parameters are related by the fact that thinner electrolytes not only limit the resistive losses within the fuel cell, but they also allow the incorporation of more active area into a given stack volume. Ultra-thin electrolyte membranes for solid oxide fuel cells must meet at least two criteria: they must be free of pinholes through which gases could leak, and they must be of the correct chemical composition and crystalline structure. Atomic layer deposition (ALD) is capable of making a pinhole-free layer of material 2 to 3 orders of magnitude thinner than state-of-the-art electrolyte membranes. Furthermore, this layer is highly conformal to the substrate, even over rough or convolved surfaces. The chemical composition of the membrane is determined by the mix of precursors admitted to the deposition chamber, and it has been shown that the crystalline structure can be controlled through a sequence of deposition steps at different compositions and thicknesses, followed by annealing at a specified temperature.

Electrolyte thickness is not the only parameter that governs efficiency. The nature of the chemical reactions that take place at the surface of the membrane (the "triple phase boundary") is also responsible for lost work in a fuel cell. Absorption by surface chemical species of photons in the ultraviolet range may reduce the activation energy for these reactions or provide alternative reaction pathways, thereby improving cell efficiency. Illumination studies are also of fundamental interest as a means of altering the electronic carrier concentrations in metal oxides and thus, perhaps, the concentration of charged point defects. Our research into novel structures for low-temperature SOFCs involves a thorough investigation of the properties of ALD fabricated yttria stabilized zirconia (YSZ) and related metal oxide membranes, and the associated microstructure-property relations. Working with our collaborators at Harvard, the oxygen ion-conductivity of the ultra-thin films will be compared to that reported for bulk crystals and ceramics. A porous electrode layer must be applied to the surface of the electrolyte membrane for the fuel cell to be operational. The team will investigate the use of ALD to form a perovskite mixed electron-ion conductor layer for this purpose. Once a perovskite deposition process has been developed, the team will conduct experiments to determine both the oxygen diffusivity through the YSZ and the surface oxygen exchange kinetics on bare YSZ and on perovskite-coated YSZ. While an increase in membrane performance can be achieved via reduction of its thickness, the major increase in membrane area proposed by this research will be achieved by forming the electrolyte as a series of YSZ nanotubes instead of planar YSZ sheets. Metal oxide nanotubes are synthesized by growing a forest of germanium nanowires on a silica substrate. This forest acts as a template, and the YSZ layer will be grown atop the germanium. The initial germanium nanowire outer diameter will thus define the inner diameter of the final metal oxide nanotubes. ALD-coated nanowires will then be released from the substrate surface by selective etching of the underlying silica and Ge nanowires. Background Atomic layer deposition continues to gain attention as a method for synthesizing solid oxide fuel cell materials. In the past year, several groups have reported results on ALD of thin layers of electrode materials that would be compatible with micro-fabricated SOFC structures of the type of interest in this project. [1-3] Hoover and Tolmachev have investigated ALD Pt deposition of sub-monolayer thickness onto base metal Ni electrodes. [1] Their work focuses on minimizing the use of expensive noble metal catalyst materials (ultra-thin ALD coatings) while enhancing the catalytic activity of PEM fuel cell electrodes. However, the same approach can be pursued in making composite electrodes for solid oxide fuel cells. The resulting ALD-coated electrodes were to found to exhibit a substantial enhancement in catalytic activity for oxygen reduction, comparable to PtNi alloy cathode materials.

Jiang et al. [2] of the Bent and Prinz groups at Stanford, have reported on ALD using methylcyclopentadienyl)trimethylplatinum precursor of blanket Pt thin films and patterned Pt. Patterned Pt deposition by area-selective ALD, using a microcontact printed self-assembled monolayer to locally block precursor adsorption, was performed to prepare current collector grids for micro-fabricated SOFC’s. These grids resulted in greater than a 10-fold increase in peak power density. Holme et al., [3] also affiliated with the Prinz group, used ALD to prepare (LaxSr1−x)MnO3 (LSM) mixed electronic/ionic conducting cathode layers for SOFCs. Crystalline LSM layer were grown directly on crystalline YSZ substrates. Dense ALD-grown LSM cathode layers of 95 nm thickness were found to limit the performance LSM/YSZ/Pt SOFC membranes because of low oxygen diffusivity in LSM and sluggish reduction kinetics on the LSM surface at the 450°C measurement temperature. In work focusing on graded ALD-electrodes, Brahim et al. from the Cassir group at CNRS (France) and the University of Helsinki (Finland) recently reported a study of composition effects in ALD-grown ZrO2-In2O3 thin films on electronic and ionic conduction [4]. This material is reported to exhibit either predominantly ionic or electronic conductivity depending on the alloy composition. Relatively thick (~ 1.1 µm) films with a graded ZrO2-In2O3 composition, such that the top layers had 90 mol% In2O3 while the layer near the YSZ electrolyte had 25 mol% In2O3, which should produce a gradation in conduction mechanism, were deposited. This graded cathode layer was found to improve the electrochemical performance of the resulting membranes, as measured by impedance spectra, relative to Pt/YSZ/Pt structures. Results YSZ Nanoscale Electrolyte Membranes by ALD In our initial work on atomic layer deposition (ALD) of metal oxide nanolaminates to prepare YSZ fast ion conducting membranes, we observed an apparent interaction between the tris(cyclopentadienyl)yttrium and ZrCl4 chemistries with unusually low Y2O3 growths per cycle. In part because of this limitation, we subsequently changed our zirconia ALD process precursor from zirconium tetrachloride (ZrCl4) to tetrakis(diethylamino)zirconium, or TDEAZr. This new liquid precursor also provides increased species volatility and eliminates the common clogging problem of chloride-based chemistries. We are now investigating whether similar precursor interference effects are present with the new zirconium precursor. This year, we established reproducible ZrO2 growth and verified film chemistry with x-ray photoelectron spectroscopy. Surface roughness studies suggest that higher precursor source temperatures resulted in a rougher ZrO2 surface, as show in Figure 1, which might be attributed to higher species concentration. Furthermore, higher substrate temperatures result in a smoother ZrO2 surface, as shown in Figure 2, which might be attributed to more complete chemical reactions and higher species mobility at the film surface. Our studies suggest that a higher zirconium precursor temperature coupled with higher substrate temperatures yield smoother films with uniform thickness.

Figure 1: Surface roughness of 6nm-thick ZrO2 increases with TDEAZr source temperature, all other processing conditions identical.(Note: vertical scale is 20 nm, scan area is 1µm x 1µm).

Figure 2: Surface roughness of 6nm-thick ZrO2 decreases with substrate temperature, all other processing conditions identical. (Note: vertical scale is 20nm, scan area is 500nm x 500nm). The ionic conductivity of bulk YSZ is closely related to its microstructure. Bulk YSZ is fully stabilized into the cubic phase for compositions near and above 8 mol% (8YSZ). At the threshold composition of 8 mol% Y2O3, bulk YSZ exhibits a peak in ionic conductivity, which is key for SOFC electrolytic performance. In our study of nanoscale YSZ, we are investigating phase changes related to composition in order to determine whether (and how) the nanostructure of ultrathin YSZ affects its ionic conductivity—and ultimately, its potential as an SOFC electrolyte. We are currently producing zirconia films containing 0 to 25 mole percent yttria in the same nanolaminate approach described in our previous work [5]. We will send these films as well as blanket ZrO2 films to the Ramanathan group for in-plane total conductivity measurements. Through-plane conductivity measurements will be performed in our new fuel cell test station. Of particular interest is correlating in-plane and through-plane measurements for more rapid electrolyte electrochemical characterization. As deposited, nanoscale YSZ contains mixed amorphous/polycrystalline material that becomes fully crystalline (see Figure 3 below) upon annealing (950°C, 2.5hrs).

Columnar, hexagonal grains (diameters 5-25nm) span the entire thickness of the nanoscale YSZ films, as indicated by distinct boundaries between crystallites. A tetragonal-to-cubic phase transition occurs for a composition between 6 and 10 mol% Y2O3.

Figure 3: Electron diffraction analysis, showing YSZ phase determination, and plan view TEM images of annealed YSZ nanolaminates. The cubic, fast-ion-conducting, phase is apparently stabilized at a Y2O3 composition of 10 mol%, which is in the bulk composition range for the cubic phase at annealing temperatures in the vicinity of 900°C. After an additional annealing treatment (900°C, 6 hrs) to ensure equilibration of point defects, detailed in-plane electrical measurements were performed on nanoscale YSZ films using impedance spectroscopy (see Figure 4). Nanoscale 6YSZ grown on MgO showed a total conductivity comparable to the conductivity of single crystal 10YSZ. Even with lower yttria content, nanoscale YSZ films exhibits high total conductivity, comparable to the conductivity of bulk 8YSZ at higher temperatures.

Figure 4: Temperature-dependent total electrical conductivities of various ALD-YSZ compositions and bulk YSZ samples. The ALD films were measured in an in-plane electrode geometry. Both highly insulating MgO substrates and Si substrates with various surface insulator layers were investigated. The Si substrates are of interest because they are used in microfabrication of fuel cells for through-thickness electrolytic measurements. The results on MgO are most reliable because this substrate avoids current shunting at high temperatures, given its very large band gap. The 6 mol% Y2O3 ALD-YSZ films on MgO have absolute electrical conductivities similar to or greater than those of single crystal 10YSZ, but the apparent activation enthalpy of conduction is larger for the ALD films. Hafnia nanotubes – Phase Stability at the Nanoscale

While an increase in SOFC membrane performance can be achieved via reduction of its thickness, the major increase in membrane area proposed by this research will be achieved by forming the electrolyte as an array of metal oxide nanotubes instead of planar films. In this part of our work, we investigated new membrane synthesis methods for increased SOFC power density by using ALD-HfO2 oxide nanotubes and conducted fundamental studies on HfO2 phase stability at the nanoscale. Single component hafnium oxide films were investigated in these initial studies because HfO2 exhibits a similar set of polymorphic phase transitions as does ZrO2 in the bulk. Moreover, the ALD-HfO2 process is well-developed in our laboratory and is less complicated than the YSZ laminate approach.

For fabrication of well defined, vertical HfO2 nanotubes with conformal wall

thickness, pre-grown vertical Ge nanowires were used. Ge(111) nanowires on Si(111)

substrates were grown using vapor-liquid-solid (VLS) technique by Irene Goldthorpe from McIntyre’s research group and were used as template for HfO2 nanotubes fabrication. The initial Ge nanowires outer diameter defined the inner diameter of the final metal oxide nanotubes. HfO2 nanotubes process sequence includes steps illustrated by flow-chart in Figure 5.

Figure 5: Flow-chart illustrating HfO2 nanotubes process sequence.

After HfO2 deposition, the samples were annealed at 800oC in Ar atmosphere for 3

hrs in order to crystallize as-deposited amorphous HfO2. Ge is highly susceptible to aqueous corrosion as result of solubility of GeO2 in water. We showed that Ge nanowires were dissolved after few minutes in 30% H2O2 aqueous solution at 40oC in ultrasonic bath. These conditions did not produce any etching of HfO2 thin films. TEM characterization of the etched sample encapsulated in SiO2 (5th step in Fig. 5) showed hollow HfO2 nanotubes (Fig. 6). For preparation of the samples used in phase stability study, SiO2 encapsulation was not used. SiO2 encapsulation and further SiO2 etching is much more useful and convenient for HfO2 nanotubes with thicker walls for future large area SOFC membranes. Instead, wet etching of Ge out of ultrathin ALD HfO2 shell was preformed in an ultrasonic bath directly after HfO2 deposition or annealing procedure.

Ultrasonic treatment caused ultrathin wall HfO2 nanotubes to fracture and assisted the selective etching of the Ge nanowire cores.

Figure 6: TEM cross-section image showing a HfO2 nanotube; high-resolution images of the HfO2 nanotube wall and of the HfO2 film deposited on Si(111) between the nanotubes.

At atmospheric pressure and the annealing temperature of 800°C employed in our experiments, the bulk crystalline polymorph with the lowest Gibbs free energy is the monoclinic (m-HfO2) phase. In annealed samples, we observed reversed phase stability, where in structures with smaller dimensions (thinner nanotube walls and films) the monoclinic phase was either not observed or present in a smaller amount than the tetragonal (t-HfO2) phase. Amorphous (a-HfO2) was observed in the thinnest 1.8 and 2.0 nm-thick samples, despite the long high-temperature post deposition annealing of 3 hrs. Figure 7 illustrates three pairs of similar-thickness samples, having different preparation sequences. Figures 7a-7c show SAED patterns from samples which were annealed first, then went through Ge etching; Figs. 7d-7f are SAED patterns from samples which went through Ge etching first, then were annealed. Due to the small fraction of monoclinic phase inside the large area sampled by plan-view SAED, the resulting diffraction patterns were dominated by t-HfO2 reflections for 3.1 and 3.5 nm-thick samples; however, the more detailed nano-beam diffraction studies did detect a minor component of m-HfO2. Plan-view SAED confirmed monoclinic structure for the 7.6 and 7.0 nm samples. Note that the relatively weak diffraction rings in Figs. 7b, 7c, 7e and 7f are probably due to the small thicknesses (and thus the small total volume) of the examined films and nanotubes.

Differences between electron diffraction patterns of t-HfO2 and the cubic phase are

small. Interplanar spacings of similar planes in both phases are very similar [6,7];

however, their intensities will differ in a randomly-oriented sample. Because of the similarity in interplanar spacings, most SAED diffraction rings in the 3.1 and 3.5 nm-thick plan-view sample can be indexed to either the tetragonal or cubic phases. However, in these two samples, the {311} ring had a high intensity of 52% and 67%, respectively, relative to the {111} ring, which is reported to have the greatest diffracted intensity for random samples of both phases. This is similar to the powder diffraction file data for t-HfO2, which shows 70% for tetragonal {311}[6] but only 1% for cubic phase [7] relative to 100% of {111} ring. Thus, we tentatively identify these films as t-HfO2, rather than the cubic phase. In the 7.6 and 7.0 nm-thick samples, the measured SAED rings correspond to the monoclinic phase only.

Figure 7: SAED patterns taken from plan-view samples: (a) 1.8 nm-thick sample showing an absence of diffraction rings (amorphous structure), (b) 3.1 nm-thick sample showing tetragonal structure, and (c) 7.6 nm-thick sample showing monoclinic structure, (d) 2.0 nm-thick sample showing an absence of diffraction rings (amorphous structure), (e) 3.5 nm-thick sample showing tetragonal structure, and (f) 7.0 nm-thick sample showing monoclinic structure. Insets in (b) and (c) are TEM nanobeam diffraction patterns of disks taken from cross-section samples having 3.1 nm and 7.6 nm thickness accordingly.

Average crystallite sizes were calculated for each sample directly from measurements of TEM images and are summarized in Table I. A clear trend of crystal size increase with nanotube wall or film thickness is seen.

Table I: Thickness and crystallite size in annealed ALD-HfO2 nanotube walls and thin films. a) Samples which were annealed first, then went through Ge etching. b) Samples which went through Ge etching first, then were annealed.

Thickness (nm) a) 1.8 b) 2.0

a) 3.1 b) 3.5

a) 4.3 b) 4.8

a) 5.5 b) 5.8

a) 7.6 b) 7.0

Average crystal size measured

from plan-view TEM (nm)

a) NA b) NA

a) 4.5 ± 0.6 b) 4.7 ± 0.7

a) 5.3 ± 1.1 b) 5.4 ± 1.0

a) 10.7 ± 1.7 b) 11.2 ± 1.9

a) 16.2 ± 2.8 b) 14.5 ± 2.4

Transitions in phase stability from a-HfO2 to t-HfO2 and from t-HfO2 to m-HfO2 with increasing film and nanotube wall thickness were observed in ALD-grown samples, after high temperature annealing. In this part of the project, our studies confirmed the existence of a microstructure-related crossover in HfO2 polymorph stability at the nanoscale, as they were predicted from thermodynamic evaluations [8]. Phase stability studies of SOFC membranes at the nanoscale have a great importance for explanation of high-temperature behavior and conductivity of the membranes. Microstructural Characterization and High-Temperature Conductivity of ALD-Grown HfO2 Films on MgO (100)

Despite substantial scientific and technological interest in ultrathin HfO2 films, to the best of our knowledge, there have, to date, been no reports on the electrical conductivity of nanocrystalline hafnia films as a function of temperature and oxygen partial pressure. In this part of study, we report on the electrical conductivity of high quality HfO2 films grown by ALD and annealed at 800°C, and on crystallographic and structural details of these films observed by TEM. Analysis of the HfO2 film conductivity suggested that hole conduction via charged oxygen interstitial formation can be a dominant conduction mechanism in undoped ALD-HfO2 films. Two sets of HfO2 films with different thickness of ~ 35 nm and ~63 nm, respectively, were annealed for 3 hr at 800 °C in flowing Ar gas. Structural information on films, such as the crystal structure, film thickness, and crystallite morphology were obtained by plan view and cross-sectional TEM characterization. The electrical conductivity of the films was studied by electrochemical measurements in a home-built furnace-type system in Harvard University by Prof. Shriram Ramanathan’s group.

Figure 8: (a) Cross-section TEM image of the 35 nm HfO2 thin film grown on the MgO(100) substrate after annealing at 800°C in Ar for 3 hrs. The inset is the plan view TEM image. (b) SAED pattern image of the identical sample after electrochemical measurement. (c) Lattice image by high resolution cross-section TEM of the annealed 35 nm film prior to conductivity measurement. The inset is the FFT analysis of the marked area.

Figure 8a displays the cross-section TEM image of a representative HfO2/MgO sample. Inset shows plan view TEM image obtained from the identical sample after electrode stabilization and electrochemical measurements conducted for ~3 days above 800°C. It was observed that the film thickness is ~35 nm on average and the film surface is rough. The lateral size of columnar crystallites varies from ~18 nm to 40 nm in the annealed film and the comparison between Fig. 8a and its inset confirms that the grains were grown up to a few hundreds of nanometers during high temperature experiment. As shown in Fig. 8b, SAED patterns were successfully indexed to reported diffraction patterns of monoclinic HfO2. Dark field TEM imaging at various tilting conditions confirmed that single-phase monoclinic HfO2 was fully-crystalline, without presence of amorphous matrix. The SAED patterns obtained before and after the conductivity measurement were the same, confirming that grain growth did not affect them and no structural transformation occurred during the measurement. To investigate the interface structure in detail, fast Fourier transformation (FFT) analysis was conducted. In Fig. 8c, FFT analysis of HfO2 and MgO lattice fringes near the interface (see inset) showed that HfO2 (022) planes are almost parallel to MgO (022) planes, suggesting existence of local epitaxy between the crystals and the substrate. Angle between these planes can be addressed to differences in phase structure of two materials: monoclinic structure of HfO2 and cubic structure of MgO. The weak HfO2 (100) SAED ring in Fig. 8b suggested existence of <100> texture in the films. Furthermore, the strong HfO2<100> texture was confirmed by XRD spectra (not shown).

Figure 9: Conductivity σ vs. 1000 /T plots measured on both 35 nm and 63 nm HfO2 films in air and 5% FG accompanied by activation energy acE∆ calculations (Changhyun Ko from Ramanathan group, Harvard University). Figure 9 shows the conductivity of 35 and 63 nm HfO2 films in air and forming gas obtained as a function of temperature from 700ºC to 950°C. The oxygen partial pressure is roughly 10-17 atm in for the forming gas annealing condition. acE∆ s are calculated as 1.06 eV and 1.17 eV in 35 nm and 67 nm films in air respectively and well consistent with the acE∆ value of bulk HfO2 (1.18 eV). Moreover, it was found that the activation energy for conduction is higher in the reducing environment than in air and this implies different dominant conduction mechanisms in low and high

2OP regimes. Comparison of the in-plane electrical conductivity of ALD-YSZ films (Fig. 4) and ALD-HfO2 films (Fig. 9), we find that the ALD-HfO2 films exhibit lower conductivity at a given temperature, and that their conductivity is rather strongly dependent on oxygen partial pressure. The latter observation suggests a major electronic contribution to the observed conductivity and, indeed, more detailed electrochemical analysis suggests that the conductivity at the higher oxygen partial pressures is dominated by holes rather than oxygen-related point defects. Lanthanum Strontium Cobalt Oxide Thin Films via ALD

Attempts were made at depositing a layered film of lanthanum strontium cobalt oxide. Several successful depositions with each component film present were made, but control over stoichiometry was not achieved due to a variety of factors associated with our custom ALD chamber. Figure 10 shows a thickness map of a typical sample on a single crystal silicon wafer.

Figure 10: thickness map (in Å) for a 3” single crystal silicon wafer with native oxide, with a film of lanthanum strontium cobalt oxide deposited via ALD. Twenty layers each of lanthanum oxide, strontium oxide, and cobalt oxide were deposited in that order.

Figure 11: XPS results the sample shown in the previous figure, with a film of lanthanum strontium cobalt oxide deposited via ALD. Twenty layers each of lanthanum oxide,

strontium oxide, and cobalt oxide were deposited in that order. The ratio of La to Sr to Co was approximately 12 to 2 to 4.5.

Due to the unknown but likely catalytic effect of each layer on the layer deposited on top of it, deposition rates were calculated and found to be different than for the individual oxide depositions. However, problems with clogging and thus gas delivery and vapor pressures have made it difficult to accurately determine the real deposition rates of each film. Depositions repeated with the same parameters often yielded different results, usually in the form of a gradual decrease in concentration of a particular element, which was often solved after the removal of a clog, only to have another element exhibit the same behavior. Thus, the depositions were constantly and negatively affected by clogging problems in the long lines leading from the bubbler sources to the chamber.

Given the history of problems with clogs and the difficulty in creating an environment where all lines are cleared and properly delivering precursor to the chamber, we made the decision to build a new ALD chamber, with significantly smaller volume and much shorter line lengths (e.g. 15 feet to 2 feet). At the time of writing this report, the chamber is nearly complete. The design of the chamber was hampered by budget constraints. The biggest compromise was the size of the chamber. We would have preferred to use a chamber with an internal volume barely larger than that required to hold a silicon wafer, but instead opted to use a spare chamber found in storage, which consists of eight 6” flanges. While the chamber is larger than optimal, it is quite smaller than the existing ALD chamber, and the flanges make it easier to add attachments like pressure gauges, delivery lines, and windows. The chamber itself sits on a rack with wheels, allowing space for control panels and making it easy to roll the entire ensemble out of its nook for maintenance access.

Because there is no loadlock attached to the chamber, access to the heater stage is direct through the door, meaning that a translational stage is not required. To hold the heater stage at an accessible height within the chamber, we designed a rail to support the stage, while minimizing contact and thus reducing heat sinks through direct contact. The heater stage design itself was largely taken from the heater stage in our existing chamber, which incorporate a boronitride heater sandwiched between two stainless steel pieces, with a thermocouple inserted directly through the top piece and measuring the temperature at the center. Because the power leads are exposed, there continues to be a risk of shorts due to the deposition of a conductive film if precursor vapor pressures are not carefully controlled. However, this limitation is unavoidable unless the heater and stage is directly incorporated into the chamber walls. We expect the heater stage to be able to reach temperatures around or above 800°C, making it suitable for in situ anneals.

Of the eight large flanges, two are blanked, two are used for windows (top and side), two are used for precursor delivery (side), one is for the door (side), and the bottom is used for the pump. The pump is a large rotary pump with a chemical scrubber, and connects to a right-angle valve under the chamber through a roughly six foot hose. A mesh on the bottom of the chamber protects the right angle valve from any debris.

The precursor delivery lines are grouped by chemistry. One large flange will hold two feedthroughs for yttrium and zirconium precursors, and the other will hold feedthroughs for lanthanum, strontium, and cobalt precursors. For now, because the lines are relatively short, we plan to rely on the vapor pressure of the precursors for delivery to the chamber, rather than use nitrogen carrier gas as in the existing chamber. Oxidant and purge gases are provided via feedthroughs connected to smaller flanges that are part of the chamber.

The new chamber is almost online. We need to perform a temperature calibration for the stage and connect the nitrogen purge line to the house nitrogen, as well as a few other minor details. Because of the smaller chamber volume and much shorter line lengths, we believe that this new chamber will dramatically improve the productivity of our ALD depositions.

Figure 12: Images of the mostly built new ALD chamber. The top row shows temperature controllers and mouse valves, the middle row shows the chamber with partially connected lines, and the bottom image shows the heater stage inside the chamber. Testing of ALD- YSZ Structures – SOFC Testing Chamber

In this part of the project, we completed SOFC testing chamber construction and installation. For this purpose, a stainless steel chamber with controlled high-temperature heater was built and was connected to a gas source with controlled gas-flow rate. For impedance measurements of SOFCs, a galvanostat was purchased and was connected to a SOFC testing chamber. A PC was purchased and connected to the galvanostat in order to control the galvanostat and record the measured impedance data. In turn, to accomplish measurements of separate microfabricated fuel cells from an array of cells fabricated on 25x25mm2 Si wafers, a high-precision sub-micron micromanipulator with electric probe was purchased and connected to the SOFC testing system. The micromanipulator was equipped with high-temperature probe and high-temperature probe holder. Heating system of the chamber consisted of high-power (500 Watt) halogen-tungsten lamp housed in massive stainless steel light reflector and ceramic socket under the testing chamber. The bottom part of the chamber has a quartz window sealed with custom-ordered high-temperature vacuum gaskets made from mica. The purpose of the quartz window is to transmit heat from the lamp to the tested SOFC sample by radiative heat transfer. Temperature control of the SOFC sample was accomplished using power and temperature controllers connected to rigid thermocouple placed directly on SOFC sample. In order to prevent heat escape from bare stainless steel testing chamber, 100% pure fiber-glass insulation was wrapped around the chamber and the light reflector. The SOFC testing chamber is shown in Fig. 13.

Figure 13. (a) Schematic sketch illustrating SOFC test chamber. (b) Overview photograph of the SOFC test chamber and support structure with micromanipulator. The constructed SOFC testing chamber was tested for leaks at moderate inner pressure, simulating the real conditions where the gas will flow through the chamber at the rate of ~80 cc per minute. No leaks from the chamber and the surrounding tube connectors were detected. The SOFC testing chamber was also successfully tested for high-temperature heating to 600oC, while planned SOFC testing temperature is intended to be 500oC. At this moment, the SOFC testing chamber is completely ready to use. The first SOFC impedance measurements are planned for the beginning of May 2009. References

1. R.R. Hoover and Y.V. Tolmachev, “Electrochemical Properties of Pt Coatings on Ni Prepared by Atomic Layer Deposition,” J. Electrochem. Soc. 156, A37-43 (2009).

2. X. Jiang, H. Huang, F.B. Prinz, and S.F. Bent, “Application of Atomic Layer Deposition of Platinum to Solid Oxide Fuel Cells,” Chem. Mater. 20, 3897-95 (2008).

3. T.P. Holme, C. Lee, and F.B. Prinz, “Atomic Layer Deposition of LSM Cathodes for Solid Oxide Fuel Cells,” Solid State Ionics 179, 1540-44 (2008).

4. C. Brahim, F. Chauveau, A. Ringuede, M. Cassir, M. Putkonen, and L. Niinisto, J. Mater. Chem. 19, 760-66 (2009).

5. C.N. Ginestra, R. Sreenivasan, A. Karthikeyan, S. Ramanathan, and P.C. McIntyre, “Atomic Layer Deposition of Y2O3/ZrO2 Nanolaminates: A Route to Ultrathin Solid-State Electrolyte Membranes,” Electrochem. Solid-State Lett. 10, B161-65 (2007).

6. ICDD 2007, PDF card #00-008-0342, tetragonal HfO2. 7. ICDD 2007, PDF card #00-053-0560, cubic HfO2. 8. A. Navrotsky, “Energetics of Nanoparticle Oxides: Interplay between Surface Energy and

Polymorphism,” Geochem. Trans. 4, 34-37 (2003). Contacts Paul C. McIntyre: pcm1@stanford.edu (PI) Michael Shandalov: michash@stanford.edu (Postdoc) Sybille Katz: sybkatz@stanford.edu (PI Admin)