Nanoionic Proton Conductivity Enhancement in Solid-State Reactive

NANOIONIC PROTON CONDUCTIVITY ENHANCEMENT IN SOLID-STATE REACTIVE

SINTERED BaCe0.7Zr0.1Y0.1Yb0.1O3−δ

by Daniel Ryan Clark A thesis submitted to the Faculty and Board of Trustees of the Colorado School of Mines in partial fulﬁllment of the requirements for the degree of Master of Science (Materials Science).

Golden, Colorado Date:

Signed: Daniel Ryan Clark

Signed: Dr. Ryan P. O’Hayre Associate Professor Thesis Advisor

Golden, Colorado Date:

Signed: Dr. Michael J. Kaufman Professor and Head Department of Metallurgical and Materials Engineering

ii ABSTRACT

The "holy grail" in the field of proton-conducting oxides is to decrease their minimum operating temperature so that they may be used in a wider variety of applications. In addition to this challenge, new synthesis techniques are needed so that these materials can be produced with greater ease and lower cost. The possibility of exploiting nano-scale interfacial phenomena in solid-state ionic materials has risen to prominence. However, this approach, dubbed "nanoionics" by the scientific community, has not been widely applied to proton conductors, with only a few studies showing modest improvements. In this study, a novel approach is taken to create nanoionic interfaces in a high-performing proton conductor, BaCe0.7Zr0.1Y0.1Yb0.1O3−δ (BCZYYb). Thin nickel metal films are created at the grain boundaries of the BCZYYb bulk phase, creating space-charge layers at the interfaces. These nanoionic space-charge layers create a dramatic enhancement (up to 32X) in ionic conductivity. Using isotope and concentration cell experiments, which have never before been conducted on a proton conductor/metal nano-composite, it is shown that the enhancement is indeed protonic, suggesting space-charge layer enhancement. In addition to demonstrating a large nanoionic enhancement, the synthesis technique employed in this work used offers commercial viability. Using the recently discovered solid-state reactive sintering (SSRS) technique, these membranes are 1/10th the cost of normal polymeric/sol-gel synthesized materials, and in addition require less processing time due to the removal of the calcination step and the lower required sintering temperatures.

iii TABLE OF CONTENTS

ABSTRACT ...... iii

LIST OF FIGURES ...... vi

LIST OF TABLES ...... ix

ACKNOWLEDGMENTS ...... x

CHAPTER 1 INTRODUCTION ...... 1

1.1 Motivation ...... 1 1.2 Applications ...... 2 1.3 Thesis Organization ...... 3

CHAPTER 2 BACKGROUND ...... 4

2.1 Proton Conducting Oxides ...... 4 2.1.1 Barium Cerates ...... 5 2.1.2 Barium Zirconates ...... 6 2.1.3 Barium Cerate-Zirconates ...... 7 2.2 Solid-State Reactive Sintering ...... 10 2.3 Nanoionics and the Space-Charge Region ...... 13 2.3.1 Homogeneous Systems ...... 14 2.3.2 Heterogeneous Systems ...... 18

CHAPTER 3 EXPERIMENTAL RESULTS AND DISCUSSION ...... 27

3.1 Experimental Design ...... 27 3.1.1 Synthesis and Fabrication ...... 28 3.1.1.1 Solid-State Reactive Sintering ...... 28 3.1.1.2 Polymeric Synthesis ...... 28 3.1.1.3 Nickel Reduction ...... 30 3.1.2 X-ray Diﬀraction ...... 30 3.1.3 Electron Microscopy ...... 30

iv 3.1.3.1 FESEM ...... 30 3.1.3.2 FIB ...... 30 3.1.3.3 TEM ...... 31 3.1.4 Electrochemical Characterization ...... 31 3.1.4.1 Electrochemical Impedance Spectroscopy ...... 31 3.1.4.2 Hydrogen/Deuterium Isotope Experiment ...... 32 3.1.4.3 Concentration Cell ...... 33 3.2 Results and Discussion ...... 36 3.2.1 X-ray Diﬀraction ...... 37 3.2.2 Electron Microscopy ...... 37 3.2.3 Magnetometry ...... 41 3.2.4 Electrochemical Impedance Spectroscopy ...... 43 3.2.5 Hydrogen/Deuterium Isotope Experiment ...... 46 3.2.6 Concentration Cell ...... 47 3.2.7 Discussion ...... 49

CHAPTER 4 CONCLUDING REMARKS ...... 51

4.1 Summary and Conclusions ...... 51 4.2 Future Work ...... 51

REFERENCES CITED ...... 53

v LIST OF FIGURES

2.1 Cubic perovskite crystal structure, ABX3...... 5 2.2 Nyquist plot (imaginary vs. real impedance) of BZY10 and BCY10 single crystals. . . .7 ◦ 2.3 Conductivities of acceptor-doped compositions for the BaCeO3-BaZrO3 system at 100 C and 600 ◦C...... 8

2.4 Conductivities of BaCe0.7Zr0.1Y0.2−xYbxO3−δ where 0.0 ≤ x ≤ 0.2 at different temperatures (◦C)...... 9 2.5 (a) Ionic conductivities of BCZYYb, BCZY, GDC, and YSZ from 400 to 700 ◦C in wet oxygen (∼3 vol. % H2O). (b) Typical current-voltage behavior for Ni-BCZYYb | BCZYYb | BCZY-LSCF and Ni-BCZYYb | SDC | LSCF in H2,H2 with H2S and dry propane at 750 ◦C...... 10 2.6 Secondary electron SEM images of fracture cross-sections of BZY20 with different sintering aids (∼1 wt. %) sintered at 1500 ◦C for 24 hours...... 12 2.7 Secondary electron SEM images of fracture cross-sections of BCY20 with different amounts of NiO sintered at 1500 ◦C for 24 hours...... 13 2.8 Examples of a heterogeneous interface with Phase 1 having a higher work function than Phase 2 (Φ1 > Φ2), giving rise to a space-charge layer...... 14 2.9 Percentage of atoms in grain boundaries as a function of grain size assuming grain boundary widths of 0.5 and 1 nm...... 15 •• 2.10 Oxygen vacancy (VO ) profiles of 8YSZ and 8YZA (8YSZ + 0.4 mol.% Al2O3) in the space-charge region...... 17

2.11 HRTEM image of nanocrystalline ceria and conductivity as a function of pO2...... 17

2.12 Comparison of the electrical conductivity of SrCe0.95Yb0.05O3 (a) thin ﬁlm on Al2O3 substrate (dg = 70 nm) and (b) polycrystal (dg = 3 µm) as a function of hydrogen concentration at 1000 ◦C...... 19 ◦ 2.13 Conductivity of LiI/Al2O3 electrolyte as a function of Al2O3 content at 25 C...... 20 2.14 Conductivity enhancement of AgX (X = Br,I) systems using mesoporous alumina particles (MPA) at 25 ◦C...... 21

2.15 (a) Conductivity of CaF2/BaF2 heterolayers with interfacial densities in the 430 - 16 nm range. (b) The concentration or (parallel) conductivity proﬁles for the semi-inﬁnite space- charge situation (left) and for the mesoscale situation (right) in which the space share regions overlap and the bulk values are exceeded even in the centers of the individual layers...... 22

vi 2.16 Conductivity of STO|YSZ|STO trilayers as a function of YSZ layer thickness and a high resolution dark-field STEM micrograph of the interfacial region and a comparison to thin film and single crystal YSZ...... 23 ◦ 2.17 TEM bright field micrographs of highly reduced (dry 5% H2 for 24 hours at 650 C) BZY15Pd2.0 and Pd nanoparticles...... 24 2.18 Backscatter electron images (composition) of polished cross-section of as-sintered BZY15 ◦ control, as-sintered BZY15Pd2.0, and highly reduced (dry 5% H2 for 24 hours at 650 C) BZY15Pd2.0 and the corresponding Pd-L and Y-L EDS maps...... 25 2.19 Grain boundary, bulk, and total conductivity versus temperature for BZY20Pd2.0 compared with sol-gel control BZY15 (BZY15 Control) for highly reducing (dry 5% H2 for 24 ◦ ◦ hours at 650 C) and moderately reducing (wet 5% H2 for 24 hours at 650 C) conditions. 26 3.1 Block diagram comparing traditional polymeric precursor synthesis routes against solid- state reactive sintering for yttrium-doped barium zirconate...... 29

3.2 Nyquist plot for BCZYYbNiO1 in a wet (pH2O ∼ 0.03 atm) 5% H2 (bal. Ar) environment at 100 ◦C and the equivalent circuit model used to solve for the bulk and grain boundary resistances...... 32 3.3 Schematic of concentration cell reactor design. Outside tube is made of quartz and the inside tubes are made of alumina...... 36 3.4 Schematic of concentration cell experiment gas flow...... 37 3.5 X-ray diffraction comparison (CuKα) between BCZYYbNiO1 and BCZYYb-C from 10 - 120◦ 2θ. Peak intensities have been normalized for comparison...... 38 3.6 Secondary electron images of fractured cross-sections of BCZYYb-C (sintered at 1450 ◦C for 24 h), BCZYYbNiO1 (before reduction - 1350 ◦C for 24 h) and a solid-state BCZYYb (1500 ◦C for 24 h) which had no NiO added to the oxide precursors...... 39 3.7 Electron back-scatter image (compositional) of a polished cross section of BCZYYbNiO1 (sintered at 1350 ◦C for 24 h) before reduction and the corresponding EDS maps for BaL, CeL, NiK, YL, YbM, and ZrL x-rays...... 40 3.8 Electron back-scatter image (compositional) of a polished cross section of BCZYYbNiO1 ◦ ◦ (sintered at 1350 C for 24 h) after reduction (5% H2 bal. Ar for 48 h at 750 C) and the corresponding EDS maps for BaL, CeL, NiK, YL, YbM, and ZrL x-rays...... 40 3.9 Bright-field TEM images of BCZYYbNiO1 (sintered at 1350 ◦C for 24 h) after reduction ◦ (5% H2 bal. Ar for 48 h at 750 C)...... 41 3.10 EDS spectrum from 0 - 20 keV of BCZYYbNiO1 (sintered at 1350 ◦C for 24 h) after ◦ reduction (5% H2 bal. Ar for 48 h at 750 C) for a grain boundary with Ni compared to the bulk. TEM micrograph of the measured region is shown. Unlabeled peak at 8.04 keV corresponds to CuKα which is caused by the Cu grid...... 42 3.11 Schematic demonstrating the calculation of Ni metal concentration by subtracting the diamagnetic response (χdia) from the total M vs. H signal...... 43

3.12 SQUID magnetometry data for BCZYYbNiO1 at 50 K for as-sintered, reduced (5% H2 ◦ bal. Ar for 48 h at 750 C), and reduced conductivity (measured in 5% H2 bal. Ar after reduction for ∼400 h)...... 44

vii 3.13 SQUID magnetometry data for BCZYYbNiO1 at 150 K for as-sintered, reduced (5% ◦ H2 bal. Ar for 48 h at 750 C), and reduced conductivity (measured in 5% H2 bal. Ar after reduction for ∼400 h) This is the data set that was used to calculate Ni metal concentrations...... 45 3.14 Total conductivity as a function of temperature for BCZYYbNiO1 and BCZYYb-C measured via EIS. Corresponding activation energy ﬁt lines for BCZYYbNiO1 (solid line) and BCZYYb-C (dashed line)...... 46 3.15 Speciﬁc grain boundary and bulk conductivity as a function of temperature for BCZYYb- NiO1 and BCZYYb-C measured via EIS and normalized using the brick-layer model. . . 47 3.16 Total conductivity as a function of temperature for BCZYYbNiO1 and BCZYYb-C in 5% H2 (bal. Ar) and 5% D2 (bal. Ar) atmospheres...... 48

t • t 2− t − 3.17 Transport numbers of protons ( OHO ), oxygen-ions ( O ) and electrons ( e ) for diﬀerent temperatures measured via a concentration cell in 5% H2 balance Ar (pH2O ∼ 0.03 atm). 49 3.18 Schematic of nanoionic nickel metal regions observed in BCZYYb and the corresponding change in H+ charge carriers due to a work function mismatch between the Ni and BCZYYb...... 50

viii LIST OF TABLES

t • t 2− t − 3.1 Transport numbers of protons ( OHO ), oxygen-ions ( O ) and electrons ( e ) for diﬀerent temperatures measured via a concentration cell in 5% H2 balance Ar (pH2O ∼ 0.03 atm) for BCZYYbNiO1...... 48

ix ACKNOWLEDGMENTS

I would ﬁrst like to thank my advisor, Dr. Ryan O’Hayre, for his mentorship, constant support and advice, and ski tickets over the past four years. Additionally I would like to thank my committee numbers Dr. Neal Sullivan and Dr. Jianhua Tong for their insights and comments. I owe an extra debt of gratitude to Dr. Tong for showing me most of the skills that I have today with ceramic synthesis/processing and electrochemical characterization, as well as being a great mentor and friend.

Funding for this research was provided by the Petroleum Institute in Abu Dhabi, UAE, the National Science Foundation (NSF) Graduate Student Fellowship, and NSF Grant No. DMR- 0820518. Additional support was provided by the Renewable Energy Materials Science and Engi- neering Center (REMRSEC), Colorado Center for Advanced Ceramics (CCAC), and Colorado Fuel Cell Center (CFCC). Thanks to Michael Sanders for helping greatly with the mass spectrometer for the concentration cell measurements which were imperative to this thesis.

No list could be complete without acknowledging everyone in the Metallurgy and Materials Engineering Department. Whether professor, support staﬀ, or fellow student, it is all of you that have made this such a rewarding and enjoyable chapter in my life. Thank you.

On a personal note, I would like to thank all of my fellow graduate student colleagues in the CCAC that I have had the pleasure of working with over the years, Mike, Tim, Josh, Jason, Archana, Kevin, and Aaron, for making my time here extremely enjoyable as well as all the students/post- docs in the Advanced Energy Materials Laboratory group. A special thank you to Amy Morrissey who performed the SQUID magnetometry measurements. I would like to thank my band (Virga) members Clay, Genti, and Sloan for their constant support and sharing my passion outside of school and all my other friends, too numerous to name. A large thank you to my family, whose support got me to where I am today. I owe a debt of gratitude to Vanessa for her love and support and always being there for me. I love you. I can’t forget to mention Hazel, Bixby, and Sookie for your squiggles, monkey chirping, and snuggling that has put a smile on my face and kept me sane.

x For my loving and supporting Mother CHAPTER 1

INTRODUCTION

1.1 Motivation

Since their discovery roughly 30 years ago, proton conducting ceramics have garnered a lot of attention from the scientific community. In this unique class of materials, the perovskite crystal structure (ABX3) has shown particularly great promise in proton-conducting applications. In spite of this scientific attention, proton conducting ceramics are still poorly understood compared to many other ion conductors. A majority of the previous work on proton conductors has been spent synthesizing and testing new perovskites in the hope of finding perovskites with higher conductivity and stability. The barium cerate and zirconate systems were discovered in the late 80s and 90s, respectively, and since then very few major improvements in this field have been made in respect to enhanced proton conductivity. The best performing materials to date are just slight variations in chemistry of the same principle systems that have been explored over the past two decades.

In this work, a completely different approach is taken in order to probe for better performing proton conductors: harnessing interfacial phenomena to increase ionic conductivity. Interfaces have been exploited in the world of electronic conductors for over 50 years, bringing modern day life into the the electronic age with processors, capacitors, photovoltaics, etc. These same interfacial principles, which cause changes to the local electronic structure of metals and semiconductors, can also be applied to ion conductors as well. When these interfaces are engineered at the nanometer length scale - i.e., when the interfacial density is extremely high, dramatic changes in bulk behavior can occur. A relatively new field of research has applied this concept to ion conductors and is dubbed "nanoionics." With the creation of nanoscale interfaces, both homogeneous and heterogeneous, extreme departures from bulk material behaviors have been observed. These departures from the bulk state (typically either enhancement-type or depletion-type departures) are due to the formation of space-charge regions at the interface. So far, most systems that have been probed for their nanoionic effect have have been model systems that use expensive synthesis and fabrication techniques and don’t serve much practical commercial importance.

In this thesis, the novel idea of exploiting nanoionic interfaces is combined with a high performing proton conductor. Previously, very few studies have sought to apply nanoionic principles to proton conductors. The studies that have been completed have only showed marginal gains in conductivity (partially due to doping effects, not nanoionic interfaces), nothing like the orders of magnitude changes seen in other ion conducting systems. Additionally, the systems that were studied were also expensive and had very little commercial relevance. In the present study, nanoionic interfaces are introduced with order of magnitude effects, especially at lower temperatures. In addition to demonstrating true nanoionic effects, a cost-effective synthesis route is used in order to

1 achieve metal/proton-conductor interfaces within the material. The synthesis route used is solid- state reactive sintering, where raw material costs are an order of magnitude lower than conventional synthesis techniques, along with reduced sintering temperatures and times. This work not only presents a nanoionic eﬀect in one of the highest conductivity proton conductors to date, but also demonstrates how it can be achieved using scaleable, commercially relevant and scaleable synthesis techniques.

1.2 Applications

While there are a myriad of applications for proton conducting ceramics, fuel cell technologies are the most prominent application for these materials. Fuel cells offer the advantage of much higher efficiencies for energy conversion compared with traditional internal combustion engines and turbines. Fuel cells are able to run on hydrogen (can be produced via renewable methods, such as electrolysis) as well as hydrocarbon fuels, giving them many potential commercial applications. Despite their obvious thermodynamic advantages, fuel cells are currently plagued by a few main problems. On one end there are the low temperature polymer electrolyte membrane (PEM) fuel cells (20 - 95 ◦C) which offer the advantage of low operating temperatures, allowing for easier system design and less exotic materials. The low temperature also causes one major disadvantage: chemical kinetics are slow at low temperatures. To get around this, exotic expensive catalysts are needed (usually Pt and Ru based), which makes them lose much of their commercial viability (until the catalyst performance is increased by at least an order of magnitude). On the other end of the fuel cell spectrum, there are traditional solid oxide fuel cells (SOFCs). Because of their high operating temperatures (800 - 1200 ◦C), inexpensive catalysts (such as Ni) can be used. However, the high temperature causes difficulties in designing a system to work with such high temperatures and requires extremely expensive interconnects that can handle the high temperatures and extremely reducing and oxidizing environments. Proton conducting fuel cells offer performance in the temperature "sweet spot", between low temperature PEM fuel cells and the high temperature SOFCs. This intermediate temperature range (300 - 600 ◦C), often referred to as the "Norby" temperature gap, offers the advantage of being able to use inexpensive catalysts as well as being able to use less exotic materials (such as stainless steel) to design the fuel cell system. Along with fuel cells, these material systems can also be used in related electrochemical applications such as electrolysis, sensors, ion/gas pumps, and hydrogen/water separation membranes.

The introduction of engineered interfaces further expands the applications of these materials. These ideas can be applied to other ion conductors in order to alter their transport behavior to desired applications. Since the carrier concentration is modiﬁed at these nanoionic interfaces, they can be exploited to make new exciting architectures, such as ionic transistors and switches. With these devices, the applications of these material systems are endless in their possibilities.

2 1.3 Thesis Organization

Chapter 2 provides relevant background information on the past and current research efforts in proton conducting ceramics with emphasis on the barium cerate and barium zirconate systems. It also includes background on the nanoionic effects and past research efforts that have lead the field to where it is today. Chapter 3 begins by describing some brief theory and technical details of the experiments conducted within this thesis. The second half of Chapter 3 describes the experimental results and discussion. Chapter 4 ends the thesis with concluding remarks about the implication of the findings in this thesis along with future work.

3 CHAPTER 2

BACKGROUND

2.1 Proton Conducting Oxides

For more than 100 years, scientists have known that certain oxides exhibit oxygen-ion conductivity [1]. However, it wasn’t until the 1950s that hydrogen defects were found in ZnO single crystals and conductivity changes as a function of hydrogen partial pressure were observed [2,3]. It was initially proposed that hydroxyl groups formed from the hydrogen and oxide ions and that protonic conduction could result. A decade later, the formalism of hydrogen in the oxide being ionized to protons was developed. These interstitial defects with a positive eﬀective charge were observed in Cu2O, CoO, and NiO [4]. Now it is understood that protons (or the eﬀects of them) are found in most binary and higher oxides; in fact, very few oxides exhibit the complete absence of protons [5]. It is out of the scope of this section to review proton conduction in all oxides (readers should refer to Refs. [6–13] for in-depth review articles). This section will instead focus on proton conducting perovskites with special attention paid to the doped barium cerate, zirconate, and cerate-zirconate systems since they relate most directly to the material used in this nanoionic study.

The mineral perovskite (CaTiO3) was discovered and named by Gustav Rose in 1839 (found in the Ural Mountains, Russia), and is named after a Russian mineralogist, Count Lev Aleksevich von Perovski [14]. This crystal structure has gained much notoriety for its plethora of applications such as dielectrics [15–17], ferroelectrics [18], magnetics [19], colossal magnetroresistance [20], catalysis [21], superconductors [22], and solid oxide fuel cells (SOFCs) [23]. The perovskite crystal structure has 2+ a general stoichiometry of ABX3. In the cubic form (Pm3m space group), the A cation has a coordination number (c.n.) of 12, the B4+ cation has a c.n. of 6, and the X2− anion also has a c.n. of 6 (shown in Figure 2.1). The perovskite crystal structure is commonly distorted, with the most common resulting crystal structure being orthorhombic (Pmna) [24].

•• For proton conducting perovskites, oxide-ion vacancies (VO ) are extremely important as they are needed for the formation of protonic defects by the dissociative absorption of water. These vacancies are usually formed extrinsically with a lower-valent cation substituting on the B-site. It is very common for these dopants to be chosen from the rare earth elements (designated Re) and an example of the doping to form oxygen-ion vacancies is shown in Equation 2.1.

2ABO3 0 •• × Re2O3 −−−−→ 2ReB + VO + 3OO (2.1)

In order to form protonic defects, water from the gas phase dissociates into a hydroxide ion and a proton; the hydroxide ion ﬁlls an oxide-ion vacancy, and the proton forms a covalent bond with a

4 A2+

B4+

_ X2

Figure 2.1 Cubic perovskite crystal structure, ABX3 (Pm3m). A-site cations on cell corners, B-site cation in body-centered position and X-site anions on face-centered positions. lattice oxygen, as shown in Equation 2.2.

•• × • H2O + VO + OO → 2OHO (2.2)

The proton conductivity of certain perovskites began to garner attention in the 1980s with acceptor doped LaYO3, LaAlO3, and SrZrO3. It was observed that these materials exhibited slight proton conductivity in hydrogen containing atmospheres [25]. The ﬁrst system to be studied for its high proton conductivity was SrCeO3 doped with Y, Yb, Sc, and Mg [26]. They were conﬁrmed to be proton conducting by demonstrating EMF potential in a wet hydrogen environment that could not be explained by oxygen-ion or electron conductivity. The SrCeO3 demonstrated useful applications in electrolysis [27], humidity sensors [28], and fuel cells [29]. After pioneering proton 2+ conductors with the perovskite SrCeO3, Iwahara et al. shifted their attention to using Ba as the A-site cation, pushing proton conductors to where they are today.

2.1.1 Barium Cerates

In 1988, the proton conductor BaCeO3 (doped with Nd2O3) arrived on the scene, demonstrating conductivity almost an order of magnitude higher than the doped SrCeO3 system [30]. When put in fuel cell conﬁgurations, the BaCeO3 system, with various dopants, demonstrated power densities of up to 0.2 W cm−1 at 800 ◦C [30–35]. Considering that these fuel cells were not optimized for performance, the results showed that proton conducting oxides could reduce the oper-

5 ating temperatures of SOFCs, which had normally been running from 1000 - 1200 ◦C in traditional yttrium-stabilized zirconia (YSZ) systems. While the BaCeO3 system looked promising as an alter- native energy material, it had one critical ﬂaw: stability.

The thermodynamic properties of this material were measured using drop calorimetry and it was shown that BaCeO3 is only slightly stabilized with respect to decomposition into its binary oxides, shown in Equation 2.3 [36].

BaCeO3 + CO2 → BaCO3 + CeO2 (2.3)

In the presence of even small amounts of CO2, it was found that this decomposition reaction would take place [37]. Additionally, BaCeO3 also proved to be unstable in H2O environments (pH2O ∼ ◦ 430 torr) in temperatures below 900 C, decomposing into CeO2 and Ba(OH)2 [38,39]. This meant that hydrocarbon fuels (i.e. methane) were ruled out because they would inevitably lead to the creation of CO2 and H2O and the decomposition of the electrolyte. As a result of these issues, high proton conductivity no longer became the only goal in creating new materials, stability became just as crucial. One answer was found in replacing (either partially or completely) the Ce4+ on the 4+ B-site with Zr , leading to BaZrO3 and BaCe1−xZrxO3 compounds.

2.1.2 Barium Zirconates

While the BaCeO3 system had great proton conductivity, its stability (in H2O and CO2) left something to be desired if these perovskites were to make more of an impact on SOFCs. First attention to doping the more stable Zr4+ came in the form of partially substituting it on the B- site for Ce4+, creating barium cerate-zirconates, which are covered in the next section. In 1993, Iwahara et al. synthesized BaZrO3 and examined Y, Ga, In, Dy, and Nd as dopants and found that Y and Dy provided the highest proton conductivity (tested in pure H2), but still lower (∼50%) than the best doped BaCeO3 materials [40]. It was also seen that these compounds conducted protons in a wet inert atmosphere (N2) [41]. Because the conductivity of Y-doped BaZrO3 (BZYx where x is mol. % of Y) wasn’t as high as the BaCeO3 system, it wasn’t investigated much until 1999, when Kreuer et al. showed that this stable electrolyte perhaps had more potential than was originally believed. It was seen that, while the total conductivity of the material was lower than the cerate-based compounds, the bulk ionic conductivity was actually higher, as demonstrated by electrochemical impedance spectroscopy (EIS) [42]. This is shown in Figure 2.2 where EIS was done on single crystal BZY10 and BCY10 (BaCeO3 with 10 mol. % Y), showing higher resistance (lower conductivity) for BCY10 compared with BZY10. This points to the problem with BZY having poor grain boundary conductivity. There is a high sensitivity of the concentration and mobility of protonic defects against any symmetry reduction, and the inevitable structural distortions in the grain boundary lead to decreased proton mobility along with a depletion of protonic defects [43]. This leads to the grain boundaries having in increased resistance to proton transport. Because

6 4 T ~ 140°C

pH2O = 23 hPa 3 BZY10

2 BCY10

100 kHz 10 kHz 1

1 kHz 1 MHz 100 Hz 0 0 1 2 3 4 5 6 7 8

Figure 2.2 Nyquist plot (imaginary vs. real impedance) of BZY10 and BCY10 single crystals. Adapted from Ref. [42]. of this, much of the focus with BZY has been fabricating highly dense, large-grained samples in order to reduce the grain boundary density. In Section 2.2, we cover a technique called solid-state reactive sintering, which is used to cost-eﬀectively sinter these ceramics and achieve large grain sizes. Naturally with the trade oﬀ between high total conductivity of the BCY system and the high stability of the BZY system, work has evolved to understand if the two can be combined to simultaneously achieve high stability and high conductivity - thereby leading to cerate-zirconate solid solution compounds.

2.1.3 Barium Cerate-Zirconates

The doped barium cerates offer the advantage of higher total proton conductivity, but have the disadvantage that they are not stable in CO2 and H2O environments, limiting their potential applications. However, the doped barium zirconates demonstrate great stability in different atmospheres, but have a lower total proton conductivity due to high grain boundary resistance. Naturally, it was hypothesized that partially substituting the Zr4+ for the Ce4+ would lead to an increase in stability, while hopefully maintaining the high proton conductivity. This was first explored in depth by Wienströer et al. in 1997 by examining the Nd-doped barium cerate-zirconates (BaCe0.9−xZrxNd0.1O2.95 where 0.1 ≤ x ≤ 0.9) [44]. It was found that the ionic conductivity for x ≤ 0.6 was nearly as high as the pure neodymium doped barium cerates. Additionally, the material also showed higher chemical stability than the zirconium-free compounds. This trend was also observed with Gd-doped barium cerate-zirconate, which had higher conductivity than the Nd- doped system, but demonstrated lower stability, making it clear that there is a trade off between stability and conductivity [39]. The Y-doped system (BaCe0.9−xZrxNd0.1O2.95) was studied in 2000

7 and was found to have the highest conductivity of all the doped barium cerate-zirconate systems. Not surprisingly, it was found that stability increased with zirconium content while conductivity decreased [45]. Similar results were confirmed by Kreuer, who showed that there is a relative maximum conductivity in the cerate-zirconate system which reflects the antagonistic effects of local ordering and symmetry reduction on one hand versus lattice expansion on the other [43]. Figure 2.3 shows a comparison of the Nd, Ga, and Y-doped barium cerate-zirconates.

bulk total Ryu et al. [39] Katahira et al. [45] Wienstöer et al. [44]

A: BaCeO3 B: BaZrO3

Figure 2.3 Conductivities of acceptor-doped compositions for the BaCeO3-BaZrO3 system at 100 ◦C and 600 ◦C. Adapted from Ref. [43].

Much work is currently being done on the barium cerate-zirconate electrolytes in order to make the materials more viable for fuel cell applications. One problem that most of the fuel cell materials have is they are very susceptible to coking (building up of carbon) and sulfur poisoning (from H2S gas). In order for fuel cells to be more attractive, they must demonstrate that they can

8 operate on many types of fuels, such as hydrocarbons, which typically have H2S and also coke traditional electrolytes. In 2003, Yang et al. developed barium cerate-zirconates that were co-doped with yttrium and ytterbium (BaCe0.7Zr0.1Y0.2−xYbxO3−δ where 0.0 ≤ x ≤ 0.2). This material demonstrated ∼2 times the total ionic conductivity of BaCe0.7Zr0.1Y0.2O3−δ (BCZY), along with stability in 50 ppm H2S for over 100 hours and resistance to coking [46]. In this systematic study, it was found that the BaCe0.7Zr0.1Y0.1Yb0.1O3−δ composition produced the highest conductivity, as shown in Figure 2.4, and is designated as BCZYYb. Figure 2.5 shows the ionic conductivity com-

10-1

700 650 600 550 -2 10 500 450 400

Conductivity (S/cm) Conductivity

10-3 0.00 0.05 0.10 0.15 0.20 0.25 Concentration of Yb

Figure 2.4 Conductivities of BaCe0.7Zr0.1Y0.2−xYbxO3−δ where 0.0 ≤ x ≤ 0.2 at diﬀerent temperatures (◦C). Adapted from Ref. [46]. pared with other common electrolytes along with fuel cell voltage-current behavior in H2 (with and without 20 ppm H2S) and propane. Because of the interest in using the electrolyte in hydrocarbon applications, BCZYYb was chosen for as the material of interest in this nanoionic study.

It is not exactly known why BCZYYb and Ni-BCZYYb composite anodes exhibit such great stability in sulfur and coking environments. When compared to a Ni-YSZ anode, the Ni-BCZYYb shows a critical pH2S/pH2 that is two to three orders of magnitude higher [47]. It is hypothesized that dissociative adsorption of water on the surface of BCZYYb facilitates the oxidation of the adsorbed sulfur to SO2 at or near active sites [48]. This is because sulfur poisoning is observed in dry H2 atmospheres, but not in wet H2 [46]. There are two possible mechanisms for removing the ∗ ∗ adsorbed sulfur (S ) through combination with adsorbed water (H2O ), shown in Equations 2.4 and 2.5. ∗ × ∗ • − S + 4OO + 2H2O → SO2(g) + 4OHO + 4e (2.4)

9 Power Density (W/cm2) 1.4 1.2 1.0 0.8 0.6 0.4 0.2 0.0 2 2.8 S/H 8 2 2 C in (2.5) 750 °C ◦ H H ) S and 3 2 2 2.4 700 Dry C 20 ppm H 2.0 with H Ni-BCZYYb/SDC 2 Ni-BCZYYb/BCZYYb 1.6 ,H 2 1.2 Current Density (A/cm Density Current 0.8 0.4 2 0.0

1.2 1.0 0.8 0.6 0.4 0.2 0.0 Voltage (V) Voltage + 2H (b) 2 SO -5 -1 -2 -3 -4 : 10 10 10 10 10 2 → 10 ∗ O 2 1.5 400 BCZY O). (b) Typical current-voltage behavior for Ni-BCZYYb BCZYYb 2 + 2H 1.4 ∗ O). In a way, the hydrophilic nature of these barium cerate- 2 S C. Adapted from Ref. [46]. GDC YSZ ◦ ) 1.3 -1 500 750 3 vol. % H ∼ 1.2 1000/T (K 600 , for a given pH Temperature (°C) Temperature 1.1 O 2 ∗ H a | BCZYYb | BCZY-LSCFdry and propane Ni-BCZYYb at | SDC | LSCF in H wet oxygen ( 700 1.0 800 0.9 -4 -5 -1 -2 -3 While a lot of work has been done on ﬁnding new electrolytes that exhibit high proton While it has been observed that the coking and sulfur poisoning resistance of BCZYYb is This reaction is expected to be slow due to the large number of reaction species and electrons

10 10 10 10 10 Conductivity (S/cm) Conductivity (a) conductivity, these materials typicallymaterials remain are very expensive generally to synthesizedexpensive metal synthesize. via nitrate sol-gel and This chloride or is saltssteps and similar to because require ensure polymer the one complete or precursor phase more conversion. lengthy methods calcination Yttrium-doped which and barium grinding use zirconate (BZY) is an example much better then thatBCZYYb. of YSZ, It it is isadsorb suspected still and unclear desorb that why water, since the thedue these to forward above the proton direction reactions large of conducting are amount Equation ofwater materials more adsorbed 2.5 activity, favorable have water is for (i.e. a thermodynamically BCZYYb great accelerated yields ability a higher to intrinsic activated surface 2.2 Solid-State Reactive Sintering zirconates allows them to eliminate(and sulfur oxygen-ion) conductivity. and coke build-up, as well as demonstrate great proton involved. Another hypothetical reactionadsorbed that water is and adsorbed much sulfur, more forming simple SO is the direct reaction between Figure 2.5 (a) Ionic conductivities of BCZYYb, BCZY, GDC, and YSZ from 400 to of a material that, because of its refractory nature, presents extreme challenges for commercial implementation. In order to achieve sufficiently dense BZY, typically extreme conditions, such as high sintering temperature (2100 - 1700 ◦C), long sintering times (>24 hours), and nanometer- sized starting powders (obtained from expensive and hard-to-scale sol-gel, polymeric precursor, or combustion synthesis routes) are needed in order to make fully dense pellets. This results in much higher raw material costs as well as barium volatilization (due to the high sintering temperatures required) which is detrimental to proton conductivity [39, 49–51]. Barium volatilization, caused by high temperature sintering and long sintering times, causes the formation of second phases (i.e. Y2O3), which lower the conductivity [52–54]. So not only do traditional synthesis methods for these perovskites (more specifically BZY and BCY systems) cause detrimental second phases, but they also make it difficult to cost-effectively produce the material on a commercial scale.

In order to move these materials closer to commercialization, sintering aids have been explored to reduce sintering temperature and times, as well as to enable the use of less expensive raw materials (e.g. oxides as opposed to nitrate or chloride salts). For the BZY system, Al2O3, MgO, and Y2O3 were initially evaluated as sintering aids and they had some success in improving the grain growth and densification, but still required relatively high sintering temperatures (>1500 ◦C) [56]. ZnO was then used to successfully sinter BZY at temperatures below 1350 ◦C, but it resulted in a lower conductivity material [57]. It is believed that this decrease in conductivity is due to strong proton 00 trapping at ZnZr sites [58]. While these initial studies used sintering aids, they were added to phase pure powders, meaning that expensive sol-gel or polymer procurer routes were still used to generate the perovskite phase prior to adding the sintering aids. Solid-state reactive sintering (SSRS) subsequently was born when the addition of sintering aids was used not only to increase densification and grain growth, but also to promote the phase formation of the desired perovskite from oxide precursors. A detailed study on many sintering aids (LiF, Al2O3, SnO2, and NiO) found that NiO (∼ 1 wt. %) acted as a great sintering aid that enhanced perovskite phase formation and demonstrated no negative effect on conductivity [55]. In this thesis, SSRS systems using NiO as a sintering aid will be designated as BMYxNiOy where M is the B-site cation, x is the mole percent of yttrium dopant and y is the weight percent of NiO (for example BZY20NiO1 is BaZr0.8Y0.2O3−δ with 1 wt. % NiO). Shown in Figure 2.6, it can be seen that large grain sizes (> 5 µm) were achieved using NiO as a sintering aid. Since oxide precursors (BaCO3, ZrO2,Y2O3) are used directly in this process, the raw material cost is about an order of magnitude less than traditional polymer/sol-gel routes, in addition to less energy consumption because of the lower sintering temperatures/times and the removal of the calcination step.

Solid-state reactive sintering is not exclusive to the barium zirconate system. Since NiO was so successful with BZY, naturally it was examined with the BCY system. BCY was successfully fabricated using nickel oxide at lower sintering temperatures than conventional polymeric/sol-gel methods [59]. Figure 2.7 shows how the addition of nickel oxide greatly improved densiﬁcation and grain growth during sintering. Both the BCY and BZY systems showed evidence of a second phase, BaY2NiO5 [55,59]. It turns out that this BaY2NiO5 phase plays an important role in the sintering

11 Control LiF

SnO2 NiO

Figure 2.6 Secondary electron SEM images of fracture cross-sections of BZY20 with different sintering aids (∼1 wt. %) sintered at 1500 ◦C for 24 hours. Adapted from Ref. [55]. mechanism of these perovskites. At sintering temperatures above ∼ 1000 ◦C, the intermediate BaY2NiO5 phase emerges and begins making "centers" which aid in the densification and grain growth [60]. Once the temperature exceeds 1300 ◦C, partial liquid-phase sintering is observed and it is believed that this phase, as it decomposes into the perovskite BZY phase, is preferentially located at the grain boundaries. While the SSRS mechanism has only been explored in detail for the BZY system, it is believed that it also applies for the BCY system. The addition of NiO has also been confirmed to aid in the sinterability of BCZYYb, where the formation of the BaY2NiO5 phase was again observed. The method used for synthesis was not SSRS, however, since NiO was added to the pure BCZYYb phase, and not the oxide/carbonate precursors. It was found that the addition of the NiO had a negligible effect on the conductivity [61]. However, in this study only temperatures >450 ◦C were tested and the nickel was in solid solution. For information on the specifics of the SSRS synthesis used in this thesis, see Section 3.1.1.1.

12 Figure 2.7 Secondary electron SEM images of fracture cross-sections of BCY20 with diﬀerent amounts of NiO sintered at 1500 ◦C for 24 hours. Adapted from Ref. [59].

2.3 Nanoionics and the Space-Charge Region

Interfaces have a paramount role in many of the materials that we use today. For the better half of a century, engineered interfaces have been used in semiconductors and electronics, pushing our civilization into the modern electronic age [62–66]. Just as with electronic conductors, bulk ionic transport properties can be modiﬁed at interfaces within a material. At interfaces, symmetry is bro- ken and a structural discontinuity is created. In the bulk of an ionic material, electroneutrality must be obeyed, which is in stark contrast to the situation at interfaces where small electrically charged regions can develop. The charge accumulation (or depletion) in these narrow regions, commonly

13 referred to as space-charge regions, are not only tolerable, but thermodynamically necessary [67–72]. Ionic material properties, such as charge carrier concentration, deviate from the bulk at interfaces due to factors such as a work function mismatch (heterogenous interfaces) [73], defect segregation along the interface [74], and structural disorder [75]. Figure 2.8 shows an example of a space-charge layer formation at a heterogeneous interface between two phases with diﬀerent work functions (Φ). These space-charge layers are typically on the order of the Debye length, making them approximately 1-100 nm long [76, 77]. In recent years, it has been proposed that these space-charge layers can perhaps be purposefully controlled and exploited in ionic applications with either homogeneous or heterogeneous interfaces, giving birth to the ﬁeld known as "nanoionics." _ _ + + _ + + _ + + _ + + Phase 2 Phase 1 _ + + + _ + _ + _ + + + _ + _ + + _ + + + + _ + + + + _ +

Ф 1 Ф 2

Figure 2.8 Examples of a heterogeneous interface with Phase 1 having a higher work function than Phase 2 (Φ1 > Φ2), giving rise to a space-charge layer.

2.3.1 Homogeneous Systems

Many homogeneous systems have been shown to have nanoionic effects due to space-charge layers at grain boundaries. Grain boundary effects on a material’s properties are nothing new and in the field of electroceramics the effect of grain boundaries on electrical, defect, and transport properties has been known and exploited for quite a while. One example of this is the use of ZnO as a varistor (variable resistor) material, where the grain boundaries are modified in order to give

14 the desired electrical properties of the material. These grain boundaries, which have also been mapped using a Scanning Tunneling Microscope (STM) [78], are depleted of electrons causing an increase in the overall resistance of a ZnO device by a factor of 1010 [79]. There are several ionic material systems in which the effect of the grain boundaries on bulk behavior have been observed. As the grain size approaches the nanometer regime, these interfacial effects, in many cases, begin to dominate overall material behavior. Figure 2.9 shows the importance of grain boundary interfaces in nanomaterials (< 10 nm) [75]. In this section, nanoionic and space-charge effects will be examined in the homogeneous yttria stabilized zirconia (YSZ), CeO2, SrTiO3, and SrCe0.95Yb0.05O3 systems.

100

1 nm 80 0.5 nm

Percentage of atoms in boundaries Percentage 0 1 10 100 Grain size (nm)

Figure 2.9 Percentage of atoms in grain boundaries as a function of grain size assuming grain boundary widths of 0.5 and 1 nm. Adapted from Ref. [75].

YSZ (xYSZ refers to doping amount of Y2O3 in ZrO2 where x is mole % of Y2O3) is a well- known ionic conductor that has been studied for years and is used in many applications, such as fuel cells and oxygen sensors [80–84]. Nernst observed the ionic conductivity of ZrO2 +9% Y2O3 (9YSZ) as early as 1899 [1, 85]. Doping ZrO2 with Y2O3 stabilizes the cubic ﬂuorite structure and creates extrinsic oxygen vacancies responsible for ion conduction [80]. Equation 2.6 shows the Kröger-Vink •• notation of the doping of ZrO2 with Y2O3, creating oxygen vacancies (VO ).

ZrO2 0 •• × Y2O3 −−−→ 2YZr + VO + 3OO (2.6)

15 Recently, much attention has been turned to modifying this staple ion conductor by creating nanocrystalline YSZ in an attempt to reduce the operating temperature [86]. It has been predicted that space-charge regions would form along surfaces and interfaces, leading to a change in ionic and electrical conductivity due to carrier concentration changes [87,88]. Nanocrystalline 16YSZ showed band-gap energy changes of 0.25 eV as the grain size decreased from 100 nm to 1 nm [89]. This change was explained by the formation of space-charge layers which increased in predominance with decreasing grain boundary size, showing that the microstructure can play an important role in the ionic and electronic conductivity of this material. To further explore the effect of grain size on YSZ, thin polycrystalline films of 9.5YSZ were made and it was observed that a 17 nm thick film had an electrical conductivity enhancement of nearly one order of magnitude vs. that of a 210 nm single crystal [90]. This was also attributed to space-charge effects at the interface and grain boundaries, which was shown by a higher activation energy observed in a conductivity relaxation measurement, also explaining the lower ionic conductivity.

While these space-charge regions dramatically increased the electrical conductivity, it is these same space-charge regions which are responsible for a blocking effect of ionic transport across the grain boundaries in YSZ. This blocking effect was first noted in 1975 by Ioffe et al. when it was observed that 5.7YSZ grain boundary conductivity increased linearly with grain size, however the sample with 0.2 µm did not fit the linear trend [91]. The grain boundaries were found to have an activation energy 7 kJ mol−1 higher then that of the bulk [92]. At first, this blocking effect was attributed to an intergranular amorphous SiO2 phase [93–98]. A problem arose in this explanation when the same blocking effects were noted even when the siliceous phase was not present [100]. Therefore, it meant that much of the contribution came from ZrO2/ZrO2 interfaces and that the increase in ionic grain boundary resistance was due to a space-charge layer [99, 101–106]. The 0 •• space-charge potential of YSZ is negative, which corresponds to YZr segregation and VO depletion •• [101,106]. Figure 2.10 shows predicted defect concentration of oxygen vacancies (VO ) in the space- charge region along ZrO2/ZrO2 grains.

Another ion conductor that has demonstrated nanoionic effects is cerium oxide. CeO2 is a well − 2− known mixed conductor (e and O ) which has the same fluorite structure as ZrO2 (doesn’t need to be stabilized for cubic structure) but is different in that pure CeO2 undergoes large departures from stoichiometry (CeO2−x) at elevated temperatures in reducing atmospheres, causing electronic conduction [107,108]. Equation 2.7 shows the defect formation upon oxygen leaving the lattice.

× •• 0 1 O → V + 2e + O2(g) (2.7) O O 2

Chiang et al. synthesized nanocrystalline CeO2, shown in Figure 2.11a, with an average grain size of 10 nm and compared its conductivity behavior against coarse-grain (∼5 µm) polycrystalline CeO2. The nanocrystalline CeO2−x demonstrated a shift from ionic to electronic conductivity along with a 104 enhancement in electronic conductivity (extrapolated from the coarse-grain material), as shown in Figure 2.11b [109]. The grain boundary impedance was found to be greatly reduced (> 103 lower

16 0 500 °C

-1 ](bulk))

. O -2 )/[V x ](

. O -3 . 8YSZ 8YZA

log([V -4

0 1 2 3 Distance from Grain Boundary Core (nm)

•• Figure 2.10 Oxygen vacancy (VO ) proﬁles of 8YSZ and 8YZA (8YSZ + 0.4 mol.% Al2O3) in the space-charge region. Adapted from Ref. [99].

0 (a) (b) T=600°C -1

-2 dg=10 nm -3

-4 d ~ 5 μm ] (S/cm) reduced g (electronic)

bulk extrinsic (ionic)

σ -5 ~10 4 log[ -6 extrapolated electronic -7 -8 -22 -18 -14 -10 -6 -2 2 6 8

log[pO2] (Pa)

Figure 2.11 (a) High resolution TEM image of nanocrystalline CeO2−x of 10 nm average grain size prepared from chemically processed powder. (b) The grain conductivity of a coarsened polycrystal exhibits electronic and ionic regimes as a function of pO2 . In contrast, the nanocrystal conductivity follows electronic behavior characteristic of a 4 reduced oxide even at high pO2, and is ∼10 greater than the extrapolated electronic conductivity of the coarsened polycrystal at 600 ◦C. Adapted from Ref. [109].

17 resistance) which corresponded to the activation energy being half that of the coarse-grained material (more than 2.4 eV lower per oxygen vacancy compared to the coarse-grained counterpart) [110]. The loss of ionic conductivity for electronic conductivity is assumed to be correlated with a depletion of acceptor dopants in the bulk as a consequence of their segregation at grain boundaries [111].

Kim et al. showed with a Mott-Schottky model (both for CeO2−x and 0.15 mol % Gd-doped CeO2), that the grain boundaries form oxygen-ion depletion space-charge regions along with an accumulation of electrons. Space-charge potential was found to be 0.3 eV and weakly dependent on temperature and pO2 [112].

Fluorite crystal structure ceramics aren’t the only ion conductors to have demonstrated homogeneous space-charge effects, perovskites have also been studied. As a prototypical example, strontium cerate doped with ytterbium is another mixed conductor (proton, oxygen-ion, and electron) that has been widely studied over the last two decades. This perovskite (perovskites covered in further detail in Section 2.1) has attractive potential applications in gas sensors due to its sensitivity to ambient gas concentrations [114,115]. SrCe0.95Yb0.05O3, similar to CeO2, has shown space-charge effects as interfacial area is increased with decreasing grain size. Thin films, deposited on Al2O3, with grain size ∼ 70 nm have shown dramatic enhancements in conductivity compared with coarse-grain polycrystalline samples (dg =3 µm) [116]. This result was attributed to the interfacial effects becoming dominant and was related to the size-dependent grain boundary impurity segregation, which has been confirmed by other studies [117]. Later, nanostructured SrCe0.95Yb0.05O3 showed greatly enhanced proton conductivity and faster reaction kinetics in hydrogen-containing atmospheres due to grain boundary and surface controlled diffusion rates in comparison to microcrystalline speci- mens. This meant that the sensitivity and equilibrium time characteristics relating to a change in a hydrogen atmosphere of a hydrogen sensor based on this nanostructured film are entirely controlled by the nanostructured material [113]. Figure 2.12 shows the enhancement in conductivity obtained with the nanostructured material vs. the coarse-grain polycrystalline sample.

These are just some spotlight homogeneous systems demonstrating space-charge eﬀects. While homogeneous systems exploit defect segregation and structural disorder in the creation of a space- charge region, heterogeneous interfaces bring about new consequences of material work function mismatch, making their behavior intriguing and extremely useful.

2.3.2 Heterogeneous Systems

The mixing of two materials to create behavior diﬀerent than both bulk phases is nothing new to material science. In 1929, Jander showed that the ionic conductivity of a composite can be greater than each material’s individual conductivity [118]. As new material processing techniques such as pulse laser deposition (PLD) and molecular beam epitaxy (MBE) continue to advance, new exciting possibilities arise for fabricating heterostructures that allow for the exploitation of these eﬀects with greater precision and at smaller and smaller length-scales. In this section, nanoionic

18 -0.5 T = 1000°C ¼ (a) -1.0

-1.5

-2.0

¼ -2.5 (b) log[conductivity] (S/ cm) -3.0 ½ σ ¼ ~ P H2

-3.5 -1.5 -1.0 -0.5 0

log[P H2 ]

Figure 2.12 Comparison of the electrical conductivity of SrCe0.95Yb0.05O3 (a) thin ﬁlm on Al2O3 substrate (dg = 70 nm) and (b) polycrystal (dg = 3 µm) as a function of hydrogen concentration at 1000 ◦C. Adapted from Ref. [113]. heterostructures will be explored using four examples: Al2O3 (insulator) particles in diﬀerent ion conductors, BaF2/CaF2 heterolayers, ZrO2/SrTiO3 heterolayers, and Pd metal in BaZrO3.

The ﬁrst systematic studies looking at the use of a second phase to increase ion transport in a solid-state ion conductor were conducted 1973. Liang explored the lithium iodide (LiI) system which is a Li+ ion conductor that is attractive for solid-state batteries because of its relatively high ionic conductivity, low electronic conductivity, and compatibility with active anode materials such as lithium [119–121]. Remarkably, however, as shown in Figure 2.13 Liang was able to obtain + large increases in the Li ion conductivity by adding small quantities of non-conducting Al2O3 (OCV measurements conﬁrmed negligible electronic conductivity) [122]. It was clear that the large enhancement could not be explained by doping mechanisms. The cause of the enhancement was instead attributed to the formation of interfacial space-charge regions where the ion conductor and insulator were in contact [123, 124]. This work led to the exploration of the heterogeneous compositing of other conducting systems with insulators in order to increase the overall conductivity. Similar enhancements were seen for AgCl:Al2O3, AgCl:SiO2 [125], AgI:Al2O3 [126], AgBr:Al2O3 [127], and TiCl:Al2O3 [128].

19 utr)i re oudrtn h ffcso ukbhvo u otesaecag ein[134]. CaF region alternating space-charge by the to nm due 16 behavior - bulk 430 on of effects densities the Interfacial understand to order in ductors) [131]. sequences layer preparing and for geometry, choice Sata spacings, of interfacial method periodicity, the defined been well has with MBE heterosystems [125]. required was thermodynamically is density layers charge interfacial space-charge the of role the conductor:ion and ion exemplified. explored interfaces, further also conductor:insulator conductors ion were ion to of interfaces enhancement addition conductor conductivity demon- In of the [130], mechanism. advantage in system space-charge Taking role LiI a the crucial via [129]. with a plays confirmed 2.14 area also Figure was interfacial in particles that strating mesoporous shown of are area and surface achieved large were the magnitudes Al of mesoporous order Using four significant more obtained. much however, be Al systems could the conductor:insulator enhancements of these size in large density relatively interfacial the the creasing to due enhancement magnitude of order LiI/Al of Conductivity 2.13 Figure tal. et hntomxdinccnutr r ncnat eitiuino hreoe ohspace- both over charge of redistribution a contact, in are conductors ionic mixed two When one than more no showed typically they enhancement, demonstrated systems those While sdMEgot fCaF of growth MBE used dpe rmRf [122]. Ref. from Adapted

Conductivity (S/cm) × 1010 12 0 2 4 6 8 0 10 2 /BaF 2 O 3 Al 2 20 lcrlt safnto fAl of function a as electrolyte seRf.[3,3]frifraino hs F these on information for [132,133] Refs. (see 2 O 3 content (mol.%) 20 30 2 2 O /BaF 3 40 odciiyehneet falmost of enhancements conductivity , 2 eeoaeswr rprd(same prepared were heterolayers 50 2 O 3 2 60 O atce sd yin- By used. particles 3 otn at content − o con- ion 25 ◦ C. 10 4

AgX

σ AgI:MPA

σ 10 3

AgBr:MPA 10 2

AgI:Al 2O3 (60 nm) 10 1

AgBr:Al 2O3 (60 nm) 0 Conductivity enhancement enhancement Conductivity 10 0.0 0.1 0.2 0.3 0.4 0.5 0.6

Volume fraction of Al 2O3, φ

Figure 2.14 Conductivity enhancement of AgX (X = Br,I) systems using mesoporous alumina ◦ particles (MPA) at 25 C. The conductivities for AgI:Al2O3 [126] and AgBr:Al2O3 [127] also depicted (60 nm dense particles). Adapted from Ref. [129]. overall thickness of ∼ 500 nm) and the ionic conductivity was measured. Figure 2.15a shows the conductivity enhancement due to the increasing of interfacial densities between the two ion conductors. The change of conductivity with temperature is the same as expected for composite electrolytes where at lower temperatures there is conductivity enhancement characterized by an activation energy close to the migration enthalpy of the relevant carriers [135]. In the composite system, this enhancement was attributed to a space-charge layer with an enhanced carrier concentration and almost temperature-independent interfacial concentration. In the high temperature regime, the activation energy increases, indicating that the the CaF2 space-charge conductivity determined the behavior in this regime. The activation energy of 0.95 eV is almost identical to the migration en- − − ergy of F interstitials in CaF2 [133] which suggests suggests a transfer of F ions from BaF2 to − CaF2, rather than a bilateral segregation of the F to the interfacial core. Figure 2.15b shows the space-charge effects which qualitatively account for the conductivity effects observed, showing the semi-infinite space-charge situation as well as the mesoscopic situation in which there is no longer an electro-neutral bulk. In the mesoscale situation, the layers have lost their individuality and an ensemble has formed with qualitatively new conductivity properties [135]. This model system clearly demonstrates the importance of interfaces within these ion conductors and shows the possibility of creating dramatic improvements in ionic/electrical properties of materials by engineering heterogeneous interfaces.

21 (a) 10 1 (b) 0.95 eV

10 0 16.2 nm CaF 20 nm 2

-1 BaF 2 10 CaF 2

K) 50 nm -1 CaF2

cm -2 103 nm BaF 2 -1 10 Ω CaF 250 nm 2

T (

σ BaF BaF 2 10 -3 430 nm 2 CaF2 0.72 eV BaF 2 10 -4

BaF 2 CaF 10 -5 2 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 10 3 T -1 (K -1 )

Figure 2.15 (a) Conductivity of CaF2/BaF2 heterolayers with interfacial densities in the 430 - 16 nm range. (b) The concentration or (parallel) conductivity proﬁles for the semi- inﬁnite space-charge situation (left) and for the mesoscale situation (right) in which the space share regions overlap and the bulk values are exceeded even in the centers of the individual layers. Adapted from Ref. [134].

Heterlayers have also been explored for structurally dissimilar ion conductors, which also yielded greatly enhanced conductivities. In attempt to use more practical materials that can be utilized in fuel cells, YSZ was explored and epitaxial heterolayers were sandwiched between two 10 nm thick SrTiO3 (STO) layers; varying the thickness of the YSZ from 1 to 62 nm. Figure 2.16 shows that decreasing the YSZ layer thickness drastically increased the overall conductivity by more than eight orders of magnitude [136]. The enhancement was claimed to be ionic due to the fact that the DC conductance in the sample was three to four orders of magnitude lower than the values obtained from the AC measurements, indicating that the electronic contribution is negligible. It should be noted that this claim was disputed, arguing that the electronic contribution from the STO may not be negligible and that partial pressure dependance of conductivity (Brouwer diagram) and Hebb-Wagner polarization was needed in order to claim the enhancement was indeed ionic [138]. Regardless, this demonstrated that dissimilar phases and real-world ion conductors can also be used to take advantage of the nanoionic effect. While the CaF2/BaF2 and YSZ/STO demonstrate clearly the effects of nanoionics, these systems are merely models to probe scientific understanding. Growing nanometer epitaxial films has very little practical significance for commercial electrochemical devices due to the extreme cost and poor scalability. In order to make practical commercial systems, less expensive routes for preparing these interfaces have been initially explored with the BaZrO3 material system.

Introducing these heterogeneous phases in proton conductors has only gained recent attention. Tong et al. introduced Pd metal, chosen for its high work function and aﬃnity towards hydrogen

22 2

0 YSZ trilayers -2 ) -1 -4

/Scm bulk YSZ dc

σ -6

log( 1 nm -8 5 nm 20 nm 30 nm -10 62 nm thin !lm YSZ single x-tal YSZ -12 1.0 1.5 2.0 2.5 3.0 3.5 4.0 1000/T (K-1 )

Figure 2.16 Conductivity of STO|YSZ|STO trilayers as a function of YSZ layer thickness and a high resolution dark-ﬁeld STEM micrograph of the interfacial region and a comparison to thin ﬁlm and single crystal YSZ (data taken from Ref. [136]). Adapted from Ref. [137].

[139–141], into BZY15 via liquid impregnation (BZY15Pd2.0) [142]. It was found, upon reduction ◦ in dry 5% H2 for 24 hours at 650 C, that the Pd reduced into both large Pd particles (∼1 µm) and small Pd particles (∼20 nm) in the grain boundaries, dubbed "highly reducing" conditions. ◦ Under wet 5% H2 for 24 hours at 650 C (dubbed "moderately reducing" conductions) the Pd only reduced into nanoparticles (∼5-10 nm), as shown in Figure 2.17 and Figure 2.18. Figure 2.19 shows that these treatments result in conductivity increases from 1.2 - 2.7 times, with the highest enhancement seen in the modestly reduced samples, which only had the Pd nanoparticles. It is speculated that the Pd2+ ions can substitute on the B-site cation in the perovskite, creating more oxygen vacancies and thus increasing the bulk conductivity. Meanwhile, the Pd nanoparticles are hypothesized to increase the charge carrier concentration at their interfaces with the BZY15 due to space-charge layer effects, causing an increase in the measured grain boundary conductivity. However, this study did not quantitatively study the effect of the Pd nanoparticles on the different charge carriers, and only examined conductivity at temperatures above 200 ◦C. Additionally, the study employed polymeric sol-gel synthesis along with liquid Pd impregnation, which makes this material impractical and too expensive for commercial applications.

23 ◦ Figure 2.17 TEM bright ﬁeld micrographs of highly reduced (dry 5% H2 for 24 hours at 650 C) BZY15Pd2.0 and Pd nanoparticles. Adapted from Ref. [142].

In this thesis, heterogeneous interfaces are explored in BCZYYb, where nanoionic eﬀects are observed with special attention paid to the lower temperature behavior in which the space- charge layer becomes more dominant. In addition, the transport numbers for the individual charge carriers are found in order to correlate the nanoionic eﬀect to protonic vs. oxygen-ion vs. electronic enhancements.

24 BSE (Composition) Image Pd-L EDS Map Y-L EDS Map

Figure 2.18 Backscatter electron images (composition) of polished cross-section of as-sintered BZY15 control, as-sintered BZY15Pd2.0, and highly reduced (dry 5% H2 for 24 hours at 650 ◦C) BZY15Pd2.0 and the corresponding Pd-L and Y-L EDS maps. Adapted from Ref. [142].

25 Highly Reduced Moderatly Reduced Temperature (°C) Temperature (°C) 700 600 500 400 300 200 700 600 500 400 300 200 -2.4 BZY15Pd2.0 BZY15Pd2.0 -2.6 BZY15 Control -2.5 BZY15 Control

-2.8

] (S/cm) -3.0 -3.0 bulk σ -3.2 log[ -3.5 -3.4

-3.6 -4.0 -3.0 BZY15Pd2.0 BZY15Pd2.0 -3.0 BZY15 Control BZY15 Control -3.5

] (S/cm) -4.0 -4.0 GB σ -4.5 -5.0 -5.0

-5.5 -6.0 BZY15Pd2.0 -3.0 -3.0 BZY15Pd2.0 BZY15 Control BZY15 Control

-4.0

] (S/cm)-4.0 log[ total σ -5.0 log[ -5.0

-6.0 1.00 1.25 1.50 1.75 2.00 2.25 1.00 1.25 1.50 1.75 2.00 2.25 1000/T (K -1 ) 1000/T (K -1 )

Figure 2.19 Grain boundary, bulk, and total conductivity for BZY15Pd2.0 compared with sol-gel ◦ control BZY20 for highly reducing (dry 5% H2 for 24 hours at 650 C) and moderately ◦ reducing (wet 5% H2 for 24 hours at 650 C) conditions. Adapted from Ref. [142].

26 CHAPTER 3

EXPERIMENTAL RESULTS AND DISCUSSION

Many experiments examining nanoionic effects in ion conductors have been undertaken, but usually without practical implications for commercial technologies. The systems are usually made to probe fundamental effects, but would be too costly for applications in fuel cells, sensors, membranes, etc. due to the costly fabrication routes (i.e. MBE or PLD) and synthesis techniques employed. Many of the investigations have been conducted on model systems which garner little commercial interest. On the flip side, most research efforts in proton conductors have been looking at new materials, many of which turn out not to be entirely useful. Solid-state reactive sintering took one step foreword, allowing for these proton conducting perovskites to be synthesized and processed at costs one order of magnitude less than traditional methods. However, in these systems not much attention has been paid to the fate of the sintering aid (Ni in most cases) - i.e., seeking to determine whether it can be exploited for more than just phase formation, grain growth, and densification. In this thesis, SSRS is explored with BCZYYb, a high-performing, coking and sulfur resistant proton conductor, where the NiO is reduced out of solution to form Ni metal at the grain boundaries. These metal:ion conductor interfaces are believed to give rise to a nanoionic space-charge effect causing enhanced charge carrier concentration and consequentially higher proton conductivity. In this Chapter, experiments are done to examine the microstructural and electrochemical behavior of the BCZYYbNiO1 which has been reduced in order to form interfacial regions within the material.

3.1 Experimental Design

It is important to be able to not only characterize the mircostructural behavior in the BCZYYb system when NiO is added then reduced out, but it is imperative that these results be correlated to the electrochemical behavior. In this Section, the experimental details of the material synthesis, microstructural characterization, and electrochemical characterization are given. Additionally, the purpose of this section is to also briefly explain the fundamentals behind some less than common experiments: hydrogen/deuterium isotope and concentration cell experiments. Both the hydrogen/deuterium isotope experiment and the concentration cell measurements serve the purpose of exploring majority charge carriers. The hydrogen/deuterium isotope experiment offers a qualitative way to identify proton conduction in a material. The concentration cell measurement provides a way to determine transport numbers (tx, where x is the charge species), offering a quantitative insight in to what defects are responsible in different atmospheres or temperature regimes. These experiments prove crucial for mixed ionic/electronic conductors where the majority carrier species isn’t always obvious (as is the case with BaCe0.7Zr0.1Y0.1Yb0.1O3−δ).

27 3.1.1 Synthesis and Fabrication

In order to understand the eﬀects of NiO on the system, control BCZYYb was synthesized via a polymeric precursor method in order to properly compare its behavior to the SSRS BCZYYbNiO1.

3.1.1.1 Solid-State Reactive Sintering

The mixed ion conducting ceramic pellets of BaCe0.7Zr0.1Y0.1Yb0.1O3−δ containing ∼1 wt.% NiO (BCZYYbNiO1) were synthesized by the solid-state reactive sintering (SSRS) method from BaCO3 (Alfa Aesar, 99.8%), CeO2 (Alfa Aesar, 99.9%), ZrO2 (Alfa Aesar, 99.7%), Y2O3 (Alfa Aesar, 99.9%), Yb2O3 (Alfa Aesar, 99.9%), and NiO (Alfa Aesar, Ni 78.5%) precursor powders. Briefly, stoichiometric amounts of BaCO3, CeO2, ZrO2,Y2O3, and Yb2O3 were weighed and ball milled with isopropyl alcohol solvent (Sigma-Aldrich, 99.5%) with 3 mm diameter spherical yttria- stabilized zirconia (YSZ) grinding media (Union Process). 1.0 wt.% sintering aid of NiO based on the total weight of BaCO3, CeO2, ZrO2,Y2O3, and Yb2O3 was added into the above powder mixture (approx. 5 atomic % NiO). The powders were ball-milled for 24 h, and then dried at 80 ◦C for 48 h. Before pressing, 1 drop of 10 wt. % polyvinyl alcohol (Alfa Aesar, 98 - 99%, high molecular weight) in DI water was added to 2 g of BCZYYbNiO1 powder. To examine the effect of the NiO, similar solid-state samples were prepared except without NiO added (the rest of the synthesis steps are the same as above) and was sintered at 1500 ◦C for 24 h in order to get a sample dense enough to conduct electron microscopy. Green pellets (2 g) were then fabricated by hydraulic compaction under 375 MPa pressure for 2 min in a 19 mm diameter circular carbon-aided steel dry- press die using 1 wt.% steric acid (Alfa Aesar, 98%) in isopropyl alcohol as a release agent on the sides of the die. These green pellets were sintered in ambient air at 1350 ◦C for 24 h to achieve >96% relative density (theoretical density = 6.211 g cm−3 [46]). Density measurements were made using geometrical measurements of pellets ground flat then weighed due to the fact the pellets are extremely hydrophilic and degrade in water if the the Archimedes technique is used. Small pores were found for all materials synthesized by this SSRS process and they are attributed to a second phase formation with the Yb, possibly BaYb2NiO5. The cause of the pores was not investigated due to the overall sintered relative density being >96% and is discussed further in the Future Work section.

3.1.1.2 Polymeric Synthesis

As an additional comparison, BCZYYb powder was synthesized by the polymeric sol-gel method (BCZYYb-C) from metal nitrate precursors: Ba(NO3)2 (Alfa Aesar, 99+%), Ce(NO3)3·6H2O (Alfa Aesar, 99.5%), ZrO(NO3)2 solution (Sigma-Aldrich, 35wt.% in dilute nitric acid), Y(NO3)3·6H2O (Alfa Aesar, 99.9%), and Yb(NO3)3 · xH2O (Alfa Aesar, 99.9%, x was determined to be 10.48 by thermogravimetric analysis). Stoichiometric quantities of these metal nitrates were dissolved

28 into de-ionized water to form a clear solution. At the same time, as combined chelating agents, ethylenediaminetetraacetic (EDTA) acid (Alfa Aesar, 99.4%) and citric acid monohydrate (Alfa Ae- sar, 99+%) were dissolved into an ammonium hydroxide (Alfa Aesar, 28-30% NH3) aqueous solution while keeping the molar ratio of EDTA acid : citric acid : total ions to be 1 : 1.5 : 1. The metal nitrate solution was slowly added to the chelating solution and the resultant transparent solution was adjusted to pH = 10.0 using ammonium hydroxide or nitric acid. A typical solution of 2000 mL has a cation concentration of 0.10 M of Ba2+ giving rise to about 64 g powder. The polymerization was fulﬁlled by vaporizing excess water at 80 ◦C with magnetic stirring. The sticky polymerized gel was dried in oven in ambient air at 150 ◦C for 48 h to form a black charcoal powder. This black charcoal powder was calcined at 1000 ◦C for 10 h to produce the BCZYYb phase pure powder. Green pellets were then fabricated by hydraulic compaction under 375 MPa pressure for 2 min in a 19 mm diameter circular carbon-aided steel dry-press die. The optimized sintering conditions of 1450 ◦C for 24 h were used for BCZYYb-C. Figure 3.1 shows a block diagram comparing traditional polymeric precursor synthesis methods against the SSRS method.

EDTA

C6H8O7

NH4OH

Ba(NO3)2 Evaporate Water Gel/Char Calcination Final Powder ~90 °C 150 °C 800 - 1000 °C ZrO(NO3)2 6 - 24 hrs 24 hrs 6 - 24 hrs Press Pellet Y(NO3)2 Sinter at ‡ Polymeric Synthesis Polymeric 1550 - 1700 °C H O Raw Material Cost 2 Polymeric Synthesis = ~$1600/kg SSRS = ~$170/kg † BaCO3

ZrO2

Y2O3 Ball Milling Drying Final Powder Press Pellet

SSRS 24 - 48 hrs 90 °C Sinter at NiO 24 hrs 1300 - 1450 °C Isopropanol

Figure 3.1 Block diagram comparing traditional polymeric precursor synthesis routes against † solid-state reactive sintering for yttrium-doped barium zirconate. BaCO3 is some- ‡ times replaced with BaO or BaSO4. Raw material costs based on August 2012 prices.

29 3.1.1.3 Nickel Reduction

Both the BCZYYbNiO1 and BCZYYb-C pellets were reduced in 5% H2 (balanced with Ar) with a ﬂow rate of 70 standard cubic centimeters per minute (SCCM) at 750 ◦C for 48 h before physical characterization and electrochemical testing.

3.1.2 X-ray Diﬀraction

The crystal structure of the BCZYYb powder and ceramic pellets were characterized by x-ray diﬀraction (XRD) using a Philips diﬀractometer (X’Pert Pro) with CuKα radiation, tube voltage of 45 kV, tube current of 40 mA, and a slit size of 1◦. Intensities were collected at room temperature in the 2θ range of 20 - 120 ◦ with a step size of 0.008◦ and a measuring time of 5 s. All XRD data was processed using Phillips Highscore software which includes the following: the background was removed, the spectra were smoothed (polynomial quintic with a convolution range of 7), and the KαII peaks were stripped (Rachinger method with an intensity ratio of 0.5).

3.1.3 Electron Microscopy

In order to examine the microstructural characteristics of our system, scanning and transmission electron microscopy were employed along with sample preparation via a focused ion beam. In this section the operating details of these instruments will be covered.

3.1.3.1 FESEM

The microstructure was characterized using a JEOL JSM-7000F ﬁeld emission scanning electron microscope (FESEM) equipped with an energy-dispersive x-ray spectrometer (EDS). FESEM EDS samples were prepared by mounting samples in epoxy and polishing to 0.05 µm with colloidal silica and coated with carbon. Grain size measurements were done with multiple FESEM images to achieve statistical reliability. EDS mapping was done with EDAX Genisis software at 15 kV with a resolution of 256 x 200 pixels and each pixel with a dwell time of 200 µs with drift correction enabled.

3.1.3.2 FIB

TEM preparation was accomplished with a FEI Helios NanoLab 600i DualBeam focused ion beam (FIB) with an isotopically pure gallium beam and using standard lift-out method onto a copper half TEM grid and cleaned with 2 kV Ga beam.

30 3.1.3.3 TEM

Microstructure characterization was done using a 200 kV Philips CM200 transmission electron microscope (TEM) with Princeton Gamma-Tech Prism Energy Dispersive X-Ray Spectrometer. All images were collected on an 1024 x 1024 pixel CCD detector. EDS measurements were taken by tilting the sample to 18◦ and Cu peaks were ignored as they originated from the holder.

3.1.4 Electrochemical Characterization

It is important to be able to relate the microstructural phenomena observed by electron microscopy to the material’s performance. Measurements of conductivity and charge carrier type in the system are extremely important to gain insight on the eﬀect of the nanoionic nickel regions. In this section, the experimental setup for the following electrochemical characterization techniques will be explained: electrochemical impedance spectroscopy (EIS), the hydrogen/deuterium isotope experiment, and the concentration cell measurements.

3.1.4.1 Electrochemical Impedance Spectroscopy

For conductivity studies, silver paste was applied to each side of the pellets along with a current collector made out of silver mesh and gold wire. After application, they were fired at 800 ◦C for 5 h to form porous silver electrodes. The conductivities were studied in a wet 5% H2 (balanced with Ar) flown at 70 SCCM at different temperatures. All gas flow measurements were done with Aera FC-PA7800 series mass flow controllers that have an accuracy of ± 1% and repeatability of ± ◦ 0.2%. The wet gas was prepared by passing through a water bubbler at 25 C to make pH2O ∼ 0.03 atm. Final EIS spectra were taken at steady-state when the resistance values changed < 2% over 2 hours. All measurements were done starting at higher temperatures (750 ◦C) and moved down in 50 ◦C increments to a final temperature of 50 ◦C. Note that some data is not available for 50 ◦C due to the resistance being too high and causing too much noise, triggering over-potential alarms. Temperature measurements were taken with a K-type thermocouple (standard error of 2.2 ◦C or 0.75%, whichever is larger) that was placed ∼0.5 cm away from the pellet. The measurements were taken using a Gamry Reference 3000, sweeping from 1 MHz to 0.1 Hz. Figure 3.2 shows the equivalent circuit model that was used to fit the Nyquist plot data (sample data from BCZYYbNiO1 ◦ in 5% H2 bal. Ar at 100 C) in order to calculate the separate bulk and grain boundary resistances.

In order to normalize the grain boundary conductivity to take into account diﬀerent grain sizes, a speciﬁc grain boundary conductivity (σgb,sp) is calculated using the "brick layer" model [143–145] shown in Equation 3.1: g g+G L σgb,sp = (3.1) RgbA

31 50 T = 100 °C Rbulk RG.B.

40 CPE CPE

3 30

(Ω) x10 Grain Bulk Electrode

real Boundary Z 20

0 0 10 20 30 40 50 3 -Zimaginary (Ω) x10

Figure 3.2 Sample Nyquist plot for BCZYYbNiO1 in a wet (pH2O ∼0.03 atm) 5% H2 (bal. Ar) environment at 100 ◦C and the equivalent circuit model used to solve for the bulk and grain boundary resistances. where g is the grain boundary distance (∼2 nm, measured via TEM), G is the average grain size

(diameter), L is the width of the pellet, Rgb is the resistance due to the grain boundary and A is the surface area of the pellet.

3.1.4.2 Hydrogen/Deuterium Isotope Experiment

The hydrogen/dueterium (1H/2D) isotope experiment allows for the qualitative measurement of the contribution of protons to the total conductivity (which may include oxygen-ion, hole, and electron conductivity as well). It has been shown that for the proton transfer reaction [146] and for proton conductivity, the barrier for deuteron transfer is higher than that for proton transfer [147]. This difference is characterized by the difference of the ground-state energies of the O-H and O-2D oscillating states, which is about 1/2h∆v ∼ 60 meV [148]. This difference in energy causes the diffusion of the deuteron to be slower than the proton, causing a marked conductivity difference

32 in the material if proton conduction is indeed the primary charge carrier. While proton tunneling makes it hard to calculate what the exact diﬀerence should be [149], typically it is estimated to be √ approximately a diﬀerence of 2, as shown in Equation 3.2

1 σ ∝ D ∝ √ (3.2) m where σ is the conductivity, D is the diffusivity, and m is the mass of the diffusing species. Since the 2D has a mass of 2 amu, vs H which has a mass of 1 amu, the difference is expected to be √ 2 = 1.4, however ratios greater than 1.4 are often observed [150–153]. The reasons that the difference isn’t always 1.4 is outside the scope of this thesis. For more information on comparing classical, semi-classical, tunneling, and escape theory for isotopic effects, refer to Ref. [151].

Pellets for the isotope measurements were prepared by applying silver paste to each side of the pellets along with a current collector made out of silver mesh and gold wire. After application, ◦ they were fired at 800 C for 5 h to form porous silver electrodes. 70 SCCM of wet 5% H2 (balanced with Ar, pH2O ∼ 0.03 atm) was used to measure the first H2 conductivity data point for each specific temperature. The measurements were taken using a Gamry Reference 3000, sweeping from 1 MHz to 0.1 Hz. Final EIS spectra were taken at steady state when the resistance values changed < 2% over 2 hours. At the same temperature, the gas mixture was switched to 70 SCCM of wet ◦ (D2O bubbler at 25 C, pD2O ∼ 0.03 atm) 5% D2 (balanced with Ar) and the corresponding D2 isotope was measured the same as above. The gas was then switched back to 70 SCCM of wet 5% H2 and the temperature was changed to the next set point. All measurements were done starting at higher temperatures (750 ◦C) and moved down in 100 ◦C increments to a final temperature of 50 ◦C. Note that some data is not available for 50 ◦C due to the resistance being too high causing too much noise and over-potential alarms. Temperature measurements were taken with a K-type thermocouple that was placed ∼0.5 cm away from the pellet.

3.1.4.3 Concentration Cell

Concentration cell measurements (commonly referred to as EMF measurements) allow for the quantitative measurements of the transport numbers (sometimes referred to as transference number) for up to three diﬀerent species [154]. Since BCZYYb is a mixed ionic and electronic conductor • 2− − (conducts OHO, O , and e ) [46], it is important to separate the conductivity contributions from the individual species via their respective transport numbers (ti where i is the transport species). • 2− − For the purpose of this thesis, only OHO, O , and e will be considered, however using diﬀerent atmospheres, transport numbers can also be determined for h•. First it is important to illustrate that the transport numbers for the system must add up to unity, as shown in Equation 3.3.

ionic electronic z }| { z}|{ t • + t 2− + t − = 1 OHO O e (3.3)

33 t t • + t 2− Note that the ionic transport number ( ionic) is OHO O and the electronic transport number (telectronic) is te− + th• , however th• is considered negligible in our reducing environments of interest [61, 155]. Many people distinguish ionic from electronic conduction by performing simple open cell voltage (OCV) measurements and taking the ratio of the measured OCV to the calculated OCV from the Nernst potential, shown in Equation 3.4.

OCVmeasured = tOH• +O2− = tionic = 1 − telectronic (3.4) OCVNernst O

While this approach works well for ionic conductors that conduct only one ionic species (such as YSZ), it has limitations when the material shows mixed ionic conduction (i.e., many of the barium cerates and barium zirconates). In order to separate the individual contributions in the ionic transport number, one more degree of freedom is needed. To do this we must ﬁrst consider four possible reactions and correlate them to their chemical potentials (µx). The four equations that are most commonly included are shown in Equation 3.5 [156, 157].

− 2− O2 + 4e ↔ 2O + − H2 ↔ 2H + 2e − − (3.5) H2 + O2 + 2e ↔ 2OH + − 2H2 + O2 ↔ 2H3O + 2e

If one applies Equation 3.6, the electrochemical potential gradient of the electrons (∇µe− ) in the oxide can be related to the other charge-carrying species (∇µi) [158].

X ti ∇µ − = (∇µi + ∇µ − ) (3.6) e z e i i where zi is the amount of electrons related to the corresponding ionic species. Combining the chemical equilibria in Equation 3.5 with Equation 3.6 and relating ∇µO2− with ∇µOH− and ∇µH+

+ with ∇µH3O via Equation 3.5 yields Equation 3.7 [154].

1 ∇µ − = − t 2− (∇µ 2− + ∇µ − ) + t + (∇µ + + ∇µ − ) e 2 O O e H H e − − − + + − + tOH (∇µOH + ∇µe ) + tH3O ∇µH3O + ∇µe 1 (3.7) = − t 2− + 2t − − 2t + ∇µ 2− 4 O OH H3O O 1 + t + − t − + 3t + ∇µ + 2 H OH H3O H

This chemical potential term can be related to the Nernst potential (more detail found in Ref. [154]) via Equation 3.8 RT X ti ∇µi = − ∆lnci (3.8) F z i i

34 where R is the gas constant, T is the temperature, F is the Faraday constant, and ci is the concentration of species i. This is then applied to Equation 3.7 to yield Equation 3.9, which accounts for the partial pressure of hydrogen (pH2 ) and water (pH2O). ∇µe− is converted to a measurable quantitity, EII−I, which is the voltage diﬀerence across the cell with no current (OCV).

pII RT H2O E = t 2− + 2t − − 2t + ln II−I O OH H3O I 2F pH O 2 (3.9) pII RT H2 − t 2− + t + + t − + t + ln O H OH H3O 2F pI H2

Since proton transport in oxide solids is usually represented by a H+ associating around an oxygen

t + t − t + = t • site [43], the H3O and OH are ignored. Note that H OHO . This yields Equation 3.10, which 2− + − is used to calculate the transport number of O , H , and e by varying the pH2 and pH2O across the concentration cell.

pII pII RT H2O RT H2 EII−I = [t 2− ] ln − [t 2− + t + ] ln (3.10) O 2F pI O H 2F pI H2O H2

OCV measurements are often troubled by thermovoltages and other oﬀsets, so a common practice is swapping the gradient across the cell under "forward" and "reverse" conditions and calculating the gradient via Equation 3.11.

Eforward − Ereverse EOCV = (3.11) 2

With the concentration cell measurements, the nanoionic eﬀect of the nickel on each individual charge carrier in the BCZYYb system can be examined, giving greater insight into what mechanisms are taking place and whether the enhancement is indeed mostly protonic.

Figure 3.3 shows a schematic of the concentration cell reactor design where gas concentrations are supplied independently on the inside and outside tubes in order to create two different atmospheres across the pellet. The outside tube for the reactor was made of transparent quartz in order to place the thermocouple close to the sample and the inside tubes were made of alumina. The inside alumina tubes along with the pellet were minded and sealed using a glass sealant. Tempera- ture measurements were taken with a K-type thermocouple that was placed ∼0.5 cm away from the pellet. Pellets were prepared by applying silver paste to each side of the pellets along with a current collector made out of silver mesh and gold wire. After application, they were fired at 800 ◦C for 5 h to form porous silver electrodes. Both gas streams were prepared by flowing pure H2 (99.999%) and pure He (99.999%) through individual mass flow controllers, mixed with a total flow rate of 100 SCCM, then run through insulated H2O bubblers with independent temperature control. After being humidified, the gas ran through heated lines to the reactor. Gas composition samples were taken just before entering the reactor and sent through heated lines to an atmospheric sampling mass

35 Outside Quartz Tube

Outside Tube Gas Outlets

Inside Tube Gas Outlets Outside Tube Gas Inlet Inside Tube Gas Inlet

Machined Grooves

Glass Sealant

Electrolyte Pellet

Figure 3.3 Schematic of concentration cell reactor design. Outside tube is made of quartz and the inside tubes are made of alumina. spectrometer (MKS Cirrus Benchtop Atmospheric RGA System). Figure 3.4 shows a schematic of the gas flow for the concentration cell experiment. OCV measurements were taken using a Gamry Reference 3000 for the 750 and 550 ◦C temperature points. It was noted that the error in the OCV measurements (via oscillations in the voltages) increased with decreasing temperatures. Final OCV measurements were calculated by taking the average of the points over 2 h and error in these mean values was calculated at a 90% confidence interval. In order to verify the accuracy of the Gamry 3000, a Keithley 2182A nanovoltmeter was used. While there was error in the voltage, temperature, and gas concentrations, the largest found error in the transport number calculations was the linear regression preformed in order to extract the slope which contains the transport number. Regression error analysis was done in order to extract the error at a 90% confidence interval.

3.2 Results and Discussion

In this section, the results of the diﬀerent characterization methods will be presented and discussed.

36 Inlet He MFC Heated Line Non-heated Line

To Outside

Bubbler Tube Spectrometer

Inlet H2 MFC Mass

Inlet He MFC

To Inside Bubbler Tube

Inlet H2 MFC

Figure 3.4 Schematic of concentration cell experiment gas ﬂow.

3.2.1 X-ray Diﬀraction

X-ray diﬀraction experiments were conducted in order to compare the crystal structures between the BCZYYbNiO1 and BCZYYb-C. It can be seen in Figure 3.5 that the crystal structure of the BCZYYbNiO1 and the BCZYYb-C are only slightly diﬀerent. The BCZYYb-C is a fully cubic perovskite (Pm3m), where the BCZYYbNiO1 demonstrates some slight peak splitting indicating that it is orthorhombic (Pnma). The BCZYYbNiO1 synthesized is consistent with literature and is often characterized as "pseudo cubic" (a ≈ b ≈ c) [61]. No second phases, such as Y2O3, NiO, or BaY2NiO5, were found in the XRD spectrum.

3.2.2 Electron Microscopy

Fully dense samples (relative density > 96%) are needed in order to accurately test the performance of each material. Figure 3.6 shows a comparison between BCZYYb-C (sintered at 1450 ◦C for 24 h), BCZYYbNiO1 (before reduction - 1350 ◦C for 24 h) and a solid-state BCZYYb (1500 ◦C for 24 h) which had no NiO added to the oxide precursors. It can be seen that the BCZYYbNiO1 has a small amount of pores which are attributed to the formation of an intermediate second phase during sintering. Because the overall density of the pellet was above 96%, the cause of the pores was not examined and is outside the scope of this thesis. The dramatic eﬀect of phase

37 irsrcue u eentproaigars h ebae diinly tcnb enthat seen be can it Additionally, the with throughout membrane. along distributed the images were backscatter across regions SEM percolating these the not that on were observed seen was but be It microstructure, can maps. regions EDS These corresponding the 3.8. Figure in shown aries, NiO [61]. with expected BCZYYb be of not reports would Previous sintering maps. during EDS reactions BaY the same the in of regions finding depleted the Zr) show BaYb (or a Ce are deficient and regions and Ba phase Yb second and and these Nickel that possibility in BaY showing a rich the is reduction, are to It that before Zr. regions maps and phase Ce, EDS second Ba, small shows in some 3.7 maps of EDS Figure indication reduction, after the and cross-sections. only before polished Ni of on distribution the collected see to were order In boundary. grain the along the with BCZYYb solid-state the comparing by seen be BCZYYbNiO1. can SSRS growth grain and sintering, formation, (CuK comparison diffraction X-ray 3.5 Figure pnrdcini %H 5% in reduction Upon Ni the of some reduction, Upon 2 NiO

0- 10 Intensity (a.u.) 5 eotdi SSBYadBY[96]wihdmntaeN n ihareas rich Y and Ni demonstrate which [59,60] BZY and BCY SSRS in reported 120 ◦ 2 θ 20 ekitniishv ennraie o comparison. for normalized been have intensities Peak . 2 NiO 2 bl r at Ar) (bal. 5 hs,hwvrNOwsaddt hs ueBZY n the and BCZYYb pure phase to added was NiO however phase, 40 2+ oscm u fsldslto o omnce ea regions metal nickel form for solution solid of out came ions 750 2 θ 60 38 ◦ ⁄CuK for C α ewe CYbi1adBZY- from BCZYYb-C and BCZYYbNiO1 between ) 48 α ,N ea sfuda h ri bound- grain the at found is metal Ni h, 80 BCZYYbNiO1 BCZYYb-C 100 2 NiO 120 5 hs,similar phase, BCZYYb-C

BCZYYbNiO1

Solid-state BCZYYb without NiO

10 μm

Figure 3.6 Secondary electron images of fractured cross-sections of BCZYYb-C (sintered at 1450 ◦C for 24 h), BCZYYbNiO1 (before reduction - 1350 ◦C for 24 h) and a solid- state BCZYYb (1500 ◦C for 24 h) which had no NiO added to the oxide precursors.

39 Figure 3.7 Electron back-scatter image (compositional) of a polished cross section of BCZYYb- NiO1 (sintered at 1350 ◦C for 24 h) before reduction and the corresponding EDS maps for BaL, CeL, NiK, YL, YbM, and ZrL x-rays.

Figure 3.8 Electron back-scatter image (compositional) of a polished cross section of BCZYYb- ◦ ◦ NiO1 (sintered at 1350 C for 24 h) after reduction (5% H2 bal. Ar for 48 h at 750 C) and the corresponding EDS maps for BaL, CeL, NiK, YL, YbM, and ZrL x-rays. there is some grain boundary cracking, caused by the volume expansion of the Ni metal at the grain boundaries, which only occupy a fraction of the grain boundaries. The grain boundary cracking and pores are believed to be the cause of the lower grain boundary conductivity at higher temperatures until the nanoionic eﬀect becomes dominating at lower temperatures (below 350 ◦C) and

40 can explain why the enhancement is only ∼30X as opposed to 104 in other nanoionic systems. Figure 3.9 shows bright-field TEM micrographs of the nickel metal at the grain boundaries. It is quickly apparent that the nickel metal does not form nickel metal nanoparticles, but rather thin metal films at the grain boundaries. These regions are 60 - 100 nm in width and do not cover all of the grain boundaries. Pores were also observed at some of the BCZYYb:Ni interfaces, which may be caused by different diffusion rates between the two materials (Kirkendall voids) [159]. EDS, shown in Figure 3.10, reveals the dominating presence of Ni at the grain boundary, while none (within the ∼1 at. % resolution of the EDS) was detected in the bulk.

Figure 3.9 Bright-ﬁeld TEM images of BCZYYbNiO1 (sintered at 1350 ◦C for 24 h) after re- ◦ duction (5% H2 bal. Ar for 48 h at 750 C).

3.2.3 Magnetometry

While the electron microscopy data shows that nickel reduced out of solid-solution in the BCZYYb matrix, it does not tell the amount of Ni2+ that remains in solid-solution. The parti- tioning of the Ni between Ni2+ and Ni metal can, however, be determined via magnetometry. The magnetization versus ﬁeld (M vs. H) measurements were collected at 150 K using a Magnetic Prop- erty Measurement System 7 (MPMS 7) Superconducting Quantum Interference Device (SQUID)

41 Grain Boundary with Nickel

Bulk

Ni GB

500 nm Intensity (a.u.) Intensity

0 5 10 15 20 Energy (keV) Bulk Grain Intensity (a.u.) Intensity

0 5 10 15 20 Energy (keV)

Figure 3.10 EDS spectrum from 0 - 20 keV of BCZYYbNiO1 (sintered at 1350 ◦C for 24 h) after ◦ reduction (5% H2 bal. Ar for 48 h at 750 C) for a grain boundary with Ni compared to the bulk. TEM micrograph of the measured region is shown. Unlabeled peak at 8.04 keV corresponds to CuKα which is caused by the Cu grid.

Magnetometer in San Diego, CA, USA. Measurements and concentration calculations were completed by Amy Morrissey (Colorado School of Mines). The M vs. H curve collected at 150 K 0 is a combination of Ni and cBCZYYb (diamagnetic response of BCZYYb), shown in Figure 3.11. The largest error in the concentration calculation is in uncertainty in the amount of NiO added to the samples which is approx. ± 0.1 wt. %, causing an approximate change in the calculated concentrations of ± 7%. Magnetometry data for the BCZYYbNiO1 conﬁrms that before reduction (as-sintered) there is only the presence of Ni2+ ions, while after reduction 71 ± 7% of the nickel 0 reduced into Ni metal. The sample in which the conductivity was measured (∼400 h in 5% H2 bal. Ar) was also tested and it was found that 65 ± 7% of the nickel reduced into nickel metal, showing

42 Figure 3.11 Schematic demonstrating the calculation of Ni metal concentration by subtracting the diamagnetic response (χdia) from the total M vs. H signal. that there isn’t a significant change (within error) in reduction over time. This indicates that the electrochemical characterization should not be changing in time due to more Ni being reduced out since it appears that the maximum amount of Ni metal has came out of solution. Figures 3.12 and 3.13 show the specific magnetization as a function of magnetic field strength, where the distinct behavior differences between the reduced and non-reduced samples can be seen. After reduction there is still a large amount of Ni2+ ions in solid solution which is predicted to cause a change in bulk conductivity.

3.2.4 Electrochemical Impedance Spectroscopy

Here we report our findings on the ionic/electric behavior of BCZYYb and the effects of the Ni metal regions for a wide temperature range, focusing on the low to intermediate temperatures (350 - 50 ◦C). Figure 3.14 shows a comparison of the total conductivity of the BCZYYbNiO1 nanoionic composite and the BCZYYb control (BCZYYb-C) where it can be seen that there is an enhancement in conductivity throughout the entire temperature range (750 - 50 ◦C). For temperatures over 450 ◦C, an enhancement of 15 - 25% was observed. As the temperature decreased, the amount of enhancement increased dramatically, resulting in >20X enhancement for temperatures under 100 ◦C (enhancement of 32X at 50 ◦C). More insight into what is happening can be gained from examining the activation energies (derived from the slope of the conductivity curve). For BCZYYb-C, there is the appearance of two distinct slope regions. From 750 - 550 ◦C, there is an activation energy of 0.27 eV. This low activation energy is attributed to dehyration of the membrane at these high temperatures resulting in lower conductivities which is consistent with the literature reports of BZY and BCY [42,43]. From 550 - 50 ◦C there is a second slope of 0.67 eV which corresponds to expected activation energies of proton transport [42, 43, 45, 49]. When looking at the different slope regimes, differences can be seen from BCZYYb-C and BCZYYbNiO1. Unlike BCZYYb-C which has two regions, BCZYYbNiO1 has three distinct slope regions. From 750 - 550 ◦C, there is an activation energy of 0.21 eV. This region is approximately the same as the high-temperature, low-activation energy region for BCZYYb-C and is again attributed to dehyradtion of the membrane at these

43 70 T = 50 K

60 -3 ) x10

-1 50 kg 2 40

30 As-sintered 20 Reduced Reduced-Conductivity

Specific Magnetization (A m 0

-10 0.0 0.2 0.4 0.6 0.8 1.0 µ0H (T)

Figure 3.12 SQUID magnetometry data for BCZYYbNiO1 at 50 K for as-sintered, reduced (5% ◦ H2 bal. Ar for 48 h at 750 C), and reduced conductivity (measured in 5% H2 bal. Ar after reduction for ∼400 h). high temperatures. Since BCZYYbNiO1 has a higher conductivity and a 25% lower activation than BCZYYb-C, it is hypothesized that this is a result of the Ni2+ left in solid solution. From 550 - 250 ◦C there is a second slope of 0.64 eV which corresponds to expected activation energies of proton transport [42, 43, 45, 49]. Intriguingly, there is a third region from 250 - 50 ◦C which has an activation energy of 0.37 eV, approximately half the expected value of typical proton-transport activation energies. This suggests that there an entirely diﬀerent transport mechanism happening at lower temperatures.

In order to understand the enhancement phenomena and the changes in activation energies, the conductivities were separated and normalized into grain boundary and bulk contributions, shown in Figure 3.15. It can be seen that there is an enhancement (∼2 - 5X) in the bulk conductivity for the BCZYYbNiO1 material throughout the entire temperature range. The bulk enhancement is attributed to the creation of additional oxygen vacancies resulting from the Ni2+ cations substituting for the Ce4+ and/or Zr4+ B-site cations, illustrated in Equation 3.12 (B-site assumed to be 4+). This hypothesis is also consistent with the lower observed activation energy in BCZYYbNiO1 for

44 25 T = 150 K

-3 20 ) x10 -1 kg 2 15

As-sintered 10 Reduced Reduced-Conductivity

5 Specific Magnetization (A m

0.0 0.2 0.4 0.6 0.8 1.0

µ0H (T)

Figure 3.13 SQUID magnetometry data for BCZYYbNiO1 at 150 K for as-sintered, reduced (5% ◦ H2 bal. Ar for 48 h at 750 C), and reduced conductivity (measured in 5% H2 bal. Ar after reduction for ∼400 h) This is the data set that was used to calculate Ni metal concentrations. both the 750 - 550 ◦C and 550 - 250 ◦C temperature range.

2ABO3 00 •• × NiO −−−−→ NiB + VO + 2OO (3.12)

No evidence of nickel metal was found within the bulk BCZYYbNiO1 via electron microscopy. For the grain boundary conductivity, however, the nickel caused a decrease in conductivity for temperatures of 200 ◦C and above, but a dramatic enhancement in conductivity for 150 ◦C and below. These dramatic increases in conductivity at the lower temperatures (8.9X at 100 ◦C and 23X at 50 ◦C) are attributed to the metal:proton conductor interfaces observed at the grain boundaries. At the higher temperatures, these nickel metal regions, which also cause grain boundary cracking, hindering the proton conductivity due to nickel metal’s lack of proton conductivity and the blocking eﬀects of the grain boundary cracks. However, as the temperature decreases, the space-charge layers at these interfaces should become exponentially larger [88, 160, 161], which we hypothesize causes a nanoionic space-charge eﬀect to dominate over the blocking nature of the electronically conducting Ni. It is hypothesized that these regions have a much lower activation energy due to the increase in charge-carrier density. In the following sections, this hypothesis is examined in

45 Temperature (ºC) 750650 550 450 350 300 250 200 150 100 50 -1 10 0.21 eV BCZYYbNiO1 -2 10 0.64 eV BCZYYb-C ) -1 -3 10 0.27 eV

-4 0.37 eV 10

-5 10 0.67 eV -6

Total Conductivity (S cm Conductivity Total 10

-7 10

-8 10 1.0 1.5 2.0 2.5 3.0 -1 1000/T (K )

Figure 3.14 Total conductivity as a function of temperature for BCZYYbNiO1 and BCZYYb-C measured via EIS. Corresponding activation energy ﬁt lines for BCZYYbNiO1 (solid line) and BCZYYb-C (dashed line). the context of additional hydrogen isotope and concentration cell measurements which enable the speciﬁc contribution from protonic conductivity on the observed behavior to be isolated.

3.2.5 Hydrogen/Deuterium Isotope Experiment

In order to identify whether the observed conductivity enhancements are due to proton, oxygen-ion, or electron conductivity, isotope experiments were done in order to further investigate the dominant charge carriers. In particular, we sought to verify that the low-temperature enhancement was not simply due to electronic conductivity caused by the nickel metal. Figure 3.16 shows isotopic conductivity data for BCZYYbNiO1 and BCZYYb-C in order to qualitatively examine the contribution of protons vs. electrons and oxygen-ions. The results show a clear H/D isotopic eﬀect with an average ratio (H+ conductivity/D+ conductivity) of 1.78 ± 0.43 for BCZYYbNiO1 and 1.40 ± 0.31 for the BCZYYb-C, demonstrating both materials are proton conducting throughout the entire temperate range (750 - 50 ◦C), and that the observed enhancement is therefore likely to be protonic in origin. However, this method only qualitatively shows that protons are the dominating

46 Temperature (ºC) 450 350 250 200 150 100 50 -1 10 -2 BCZYYbNiO1 G.B.

) 10 BCZYYb-C G.B. -1 -3 10 BCZYYbNiO1 Bulk BCZYYb-C Bulk -4 10 -5 10 -6 10 -7 10 -8 10 -9 10 Specific Conductivity (S cm -10 10 -11 10 1.5 2.0 2.5 3.0 -1 1000/T (K )

Figure 3.15 Speciﬁc grain boundary and bulk conductivity as a function of temperature for BCZYYbNiO1 and BCZYYb-C measured via EIS and normalized using the brick- layer model. transport species in the system. In order to further separate the ionic (protons and oxygen-ions) and electronic transport, concentration cell measurements were conducted.

3.2.6 Concentration Cell

Concentration cell measurements were done in order to conﬁrm that protons are indeed the dominating transport species as was suggested by the isotope experiments, particularly at the lower temperatures where order-of-magnitude conductivity enhancements are observed. Table 3.1 shows the transport numbers determined for protons, oxygen-ions, and electrons. These values are also displayed in Figure 3.17.

BCZYYb has been claimed to be a mixed-ion conductor and this is conﬁrmed in the concentration cell measurements [46]. As the temperature decreases, the contribution of protons increases while the contribution of oxygen-ions decreases. At 750 ◦C, the oxygen-ion contribution is signiﬁ- cant, and this is attributed to proton dehydration. Since the mobility of oxygen-ions is much lower than that of protons, it causes a decrease in conductivity at the higher temperatures. This explains the conductivity plateau observed at temperatures above 750 ◦C which gives the appearance of a lower activation energy, but is instead attributed to oxygen-ions becoming an increasingly larger

47 Temperature (ºC) 750 550 450 350 250 150 50 -1 10

BCZYYbNiO1 5% H2 -2 BCZYYbNiO1 5% D 10 2 ) BCZYYb-C 5% H

-1 2 BCZYYb-C 5% D2 -3 10

-4 10

-5 10

Total Conductivity (S cm -6 10

-7 10 1.0 1.5 2.0 2.5 3.0 -1 1000/T (K )

Figure 3.16 Total conductivity as a function of temperature for BCZYYbNiO1 and BCZYYb-C in 5% H2 (bal. Ar) and 5% D2 (bal. Ar) atmospheres.

t • t 2− t − Table 3.1 Transport numbers of protons ( OHO ), oxygen-ions ( O ) and electrons ( e ) for diﬀerent temperatures measured via a concentration cell in 5% H2 balance Ar (pH2O ∼ 0.03 atm) for BCZYYbNiO1.

Transport Species 750 ◦C 550 ◦C 350 ◦C 150 ◦C 50 ◦C

t • ± ± ± ± ± OHO 0.76 0.06 0.89 0.05 0.95 0.07 0.96 0.08 0.95 0.09 tO2− 0.24 ± 0.06 0.09 ± 0.05 0.02 ± 0.07 0.02 ± 0.08 0.02 ± 0.09 te− 0.00 ± 0.04 0.02 ± 0.04 0.03 ± 0.05 0.02 ± 0.06 0.03 ± 0.07 fraction of the overall transport. For the BCZYYbNiO1, if the proton-conducting slope of 550 - 250 ◦C was extended to the higher temperatures, it is seen that the measured values are approximately 30% lower, which can quantitatively be seen in the concentration cell data where tO2− = 0.24 ± 0.06. As the temperature decreases, our hypothesis is that the space-charge regions created by the Ni metal interfaces becomes exponentially larger, causing protons to become the dominant charge carrier due to their increased local carrier concentration at the nanoionic interfaces. Equation 3.13 shows the exponential relation of charge carrier concentration with temperature [71,161,162]: c+ (x) c− (x) ∆GF 0 0 = exp − (3.13) c+ − c+ (x) c− − c− (x) kbT

48 1.0

0.8

t • 0.6 OHO 2- tO t 0.4 elec.

Transport Number 0.2

0.0 700 600 500 400 300 200 Temperature (°C)

t • t 2− t − Figure 3.17 Transport numbers of protons ( OHO ), oxygen-ions ( O ) and electrons ( e ) for diﬀerent temperatures measured via a concentration cell in 5% H2 balance Ar (pH2O ∼ 0.03 atm).

0 where cj (x) is the concentration of charge defects, j, as a function of distance, x, and cj is the reference charge defect concentration, ∆GF is the free formation enthalpy of heterojunction eﬀect pairs. As observed via SEM and TEM, these nickel regions do not percolate throughout the entire material, and therefore they provide negligible electronic conductivity. Additionally, as the activation energy for proton transport is lower than that of the oxygen-ion, it is therefore expected that the proton transport number should also increase with decreasing temperature (this same behavior has been observed for BaCe0.2Zr0.7Y0.1O3−δ with 1-2 wt. % NiO added [155]). These measurements are also consistent with the isotope experiment, which also demonstrated that protons are the dominating charge carrier throughout the entire temperature region.

3.2.7 Discussion

In this study, the novel nanoionic concept is applied to construct a nanoionic composite by using the proton conducting ceramic BCZYYb as the matrix and Ni metal nano-regions as the second phase (1 wt. % NiO, BCZYYbNiO1). The conductivity, isotope, and concentration cell data suggest that these Ni-containing samples have a higher proton concentration in the grain boundary regime [67–72,76,77], resulting in a marked enhancement of total proton conductivity. As shown in Figure 3.18, these nickel regions preferentially locate along the grain boundaries and are believed

49 2 + Ф + H + H H + + + + H H H H + + + + + + + H H H H H H H 1 ______Ф + + + + + + + H H H H H H H + + + + H H H H second phase nanoregion + + H H + 2 H Ф 50 charge carriers due to a work function mismatch between + region conductor second phase bulk proton the Ni and BCZYYb. sponding change in H grain boundary to create heterogonous interfaces,demonstration causing protonic of space-charge a regions large tothat nanoionic proton form. effect transport in This is is indeed a being the proton-conducting modified. first oxide that has explicitly shown Figure 3.18 Schematic of nanoionic nickel metal regions observed in BCZYYb and the corre- CHAPTER 4

CONCLUDING REMARKS

4.1 Summary and Conclusions

With proton-conducting perovskites gaining more scientific interest, approaches to increase their conductivity by orders of magnitude are needed in order to make them more commercially relevant. The field of nanoionics offers a potential solution to increase conductivity by exploiting nano-scale interfaces within materials, greatly altering their transport properties. To reduce cost, solid-state reactive sintering has been shown to lower the cost to 1/10th of normal polymeric synthesis routes of proton conductors, creating a major step towards inexpensive commercial-scale synthesis of these materials. In this thesis, these two major advancements are combined in order to create an inexpensive nanoionic composite, which demonstrates dramatic enhancements of over an order of magnitude in conductivity. The SSRS BCZYYbNiO1 was reduced in order to create thin metal films at the grain boundaries, which are responsible for creating protonic space-charge layers (due to the mismatch in work function between the nickel and BCZYYb). While the enhancements in the intermediate temperature range (750 - 350 ◦C) are only 20 - 30%, the lower temperature conductivity is dramatically increased by up to 32X.

To characterize what was occurring in the microstructure upon reduction of the nickel, SEM (EDS), TEM, and magnetometry were all used in conjunction. The SEM and TEM showed the formation of nickel metal regions in some of the grain boundaries, forming the nanoionic interfaces. Magnetometry conﬁrmed that there was no nickel metal present before reduction, but after reduction ∼ 70% of the Ni2+ ions were reduced into nickel metal.

In order to understand the eﬀects of the nanoionic interfaces at the grain boundaries, isotope and concentration cell measurements were completed. They revealed that the enhancement was protonic throughout the entire temperature range, demonstrating that the nanoionic space-charge mechanism is primarily responsible for the remarkable improvements in conductivity. Additionally, they showed the never-before-seen eﬀect of the space-charge region on protonic transport in this mixed-ion conductor.

4.2 Future Work

There are a number of avenues that need to be explored in greater detail on both the theoretical and experimental side. Coupling these results to space-charge models would allow for the further understanding of the role that the space-charge plays within these composites. While this thesis shows the eﬀects of a space-charge region, the region itself was never observed. With great strides in ﬁeld-ion microscopy and the development of atom probe tomography, these interfaces can be

51 chemically measured in order to examine the length and concentration proﬁle of the space charge region. This would allow for great insight on the size and distribution of these space-charge regions which could be fed back into theoretical models.

Experimentally, more investigations are needed to understand the behavior of this nanoionic composite system. First, it is not understood why nickel thin films are formed at the grain boundary instead of nickel nanoparticles. Previous work on introducing metal into these perovskites has shown the formation of small nanoparticles upon reduction. Perhaps nickel satisfies wetting conditions within this material. In order to probe this hypothesis, careful wetting and contact angle experiments could be done on BCZYYb and Ni in order to understand why nickel thin films are formed at the grain boundaries and this behavior could be compared against other similar perovskites.

In this work, small pores were also observed in the BCZYYbNiO1, but not in the BCZYYb- C, indicating that there is a second-phase reaction with the nickel. Investigating the formation of these pores and the role of the second phases would allow for further deconvolution of the changes in conductivity. Additionally, it is not understood if the nickel’s reaction with the Yb is creating these second phases and it is unknown if they enhance/deter the sintering and densiﬁcation of the SSRS BCZYYbNiO1. Additionally, more work is needed in order to understand the kinetics of reduction on this material. Careful reduction experiments, coupled with magnetometry, would identify how fast the material reduces and if it is able to get to full reduction. Such experiments could also probe the eﬀect of reduction conditions on the size and distribution of the nickel metal regions.

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