Degradation Mechanisms in Lao.
8
Sro.
2
CoO
3 as Oxygen Electrode Bond Layer in
Solid Oxide Electrolytic Cells (SOECs)
By
Vivek Inder Sharma
MASSACHUSETTS INS
OF TECHNOLOGY
B.Tech. Chemical Engineering (2007)
Indian Institute of Technology, Guwahati, India
MAR 12 2010
SUBMITTED TO THE DEPARTMENT OF NUCLEAR SCIENCE
AND ENGINEERING
IN PARTIAL FULFILLMENT FOR THE REQUIREMENTS FOR THE DEGREE OF
MASTER OF SCIENCE IN NUCLEAR SCIENCE AND ENGINEERING
AT THE
MASSACHUSETTS INSTITUTE OF TECHNOLOGY
September 2009
C 2009 Massachusetts Institute of Technology
All rights reserved
Signature of Author:
Vivek Inder Sharma
Department of Nuclear Science and Engineering
Aug 31, 2009
.
P
/
Certified by:
V Bilge Yildiz, Ph.D.
Assistant Professor of Nuclear Science and Engineering, MIT
Thesis Supervisor
7/ (~ r4
Certified by:
Harry L. Tuller, Ph.D.
Professor of Mater als Science and Engineering, MIT
Thesis Reader cx
'rv
I j A
Accepted by:
Jacquelyn C. Yanch
Professor of Nuclear Science and Engineering, MIT
Chair, Department Committee on Graduate Students
E

0
.
8
2
3
High temperature steam electrolysis is an efficient process and a promising technology to convert electricity and steam or a mixture of steam and C0
2
, into H
2 or syn-gas (H
2
+
CO) respectively. It is carried out in Solid Oxide Electrolytic Cells (SOECs). At the high temperature of operation, above 8000 C, loss in the rate of hydrogen (or syn gas) production by SOECs has been observed. This loss of performance has been a scientific and technological challenge. The goal of this thesis is to identify the mechanisms for the loss in the electrochemical performance of SOECs due to the oxygen electrode and bond layer degradation. Our specific research objectives were focused on two main mechanisms: 1) Cr transport into the oxygen electrode and bond layer, and 2) Long-range segregation of cations in the bond layer. For SOECs provided by Ceramatec Inc. for this analysis, Lao.
8
Sro.
2
CoO
3
(LSC) was the bond layer and Ao.
8
Sro.
2
MnO
3
(ASM*) was the oxygen electrode, both comprised of perovskite structure. The approach in thesis integrated complementary spectroscopy and microscopy techniques in a novel manner to carry out the 'post-mortem' analysis of SOECs from a high level to a high resolution.
Raman spectroscopy was employed to identify secondary phases on the top surface of
LSC near the interconnect interphase. Surface chemistry and microstructure of the air electrode and the bond layer was studied using scanning Auger Electron Spectroscopy
(AES) with nano-probe capability. High-resolution analysis of the cation distribution in the bulk of the LSC bond layer was achieved by employing Energy Dispersive X-ray
Analysis (EDX) coupled with Scanning Transmission Electron Microscopy (STEM).
Electrochemical treatment and characterization was performed to isolate the mechanism(s) governing the long-range segregation of cations, leading to the dissociation of the LSC bond layer.
Less-conducting, secondary phases of Cr
2
O
3
, LaCrO
3
, La
2
CrO
6 and Co
3
0
4 were identified on the top surface of LSC bond layer. The bond layer exhibited: 1) presence of
Cr, with average Cr-fraction of approximately 0.07 at the surface of its grains, and 2) surface composition variation locally, with La/Co ranging widely from 0.67 to 16.37 compared to the stoichiometric La/Co value of 0.8. Sr and Co cations migrated from the bond layer structure to the LSC/interconnect interface, over a distance of 10-20 microns.
Furthermore, STEM/EDX results showed the presence of phase separated regions at the nano-scale rich in Cr and La but lacking Co, and vice-versa. This indicates the dissociation of bond layer bulk structure at nano-scale. Cr fraction in LSC bulk varied
* The constituent A, in ASM, is proprietary information of Ceramatec Inc.

from 10 to 33%, which is higher than the average Cr-content at the surface of LSC grains.
The maximum Sr fraction observed in LSC bulk was 4.16%, confirming the migration of
Sr to LSC/interconnect interface.
We hypothesize that the long-range transport of Sr, Co, and Cr cations can be caused by two primary mechanisms: 1) Driven by Cr-related thermodynamics, where the Crcontaning species (i.e. at the vicinity of the interconnect) could thermodynamically favor the presence of select cations (i.e. Sr and Co) at the region interfacing the interconnect. 2)
Driven by the electronic or oxygen ion current. To test these hypotheses and to isolate the governing mechanism, we simulated controlled electrochemical conditions on reference cells having ASM electrodes coated with LSC, on both sides of SSZ electrolyte, without any Cr-containing layers on the LSC bond layer. The reference cells degraded even in the absence of Cr. AES results showed that the microstructure and surface composition of the reference cells stayed stable and uniform upon the electrochemical treatment, in spite of the degradation. Thus, this thesis concludes that the Cr-related thermodynamics could be the dominant mechanism driving the uneven dissociation and segregation of cations in
LSC as observed in the stack cells.
As a mechanism for Cr-deposition in the LSC bond layer, we suggest that a thermodynamically-favored reaction between the La-enriched phase (at the surface of the
LSC grains) and the volatile Cr-species (Cr0
3 and CrO
2
(OH)) is responsible for the formation of poorly-conducting secondary phases. This interaction is likely to be limited
by the presence of the segregated La-O-species which can serve as a nucleation agent for this reaction.
Keywords: Solid Oxide Electrolysis Cell, SOEC, degradation, Raman, Auger, spectroscopy, Scanning Transmission Electron Microscopy
Thesis Supervisor: Professor Bilge Yildiz
Assistant Professor of Nuclear Science and Engineering
Thesis Reader: Professor Harry L. Tuller
Professor of Ceramics and Electronic Materials

I would like to express my thanks to a number of individuals for their contributions to this thesis.
First, I would like to thank my advisor, Prof. Bilge Yildiz, for her guidance, encouragement and patience during the completion of this thesis. It has been an amazing experience working under her supervision and learning from her. I will be forever thankful to her for this opportunity.
My gratitude extends to Ceramatec Inc., for providing the SOECs for analysis. I am also thankful to DoE for funding the project and to Dr. S. Elangovan, Dr. D. Carter and Dr. S.
Herring for their valuable inputs and discussions throughout the course of my thesis. I am also thankful to Prof. Harry Tuller for kindly consenting to act as the thesis reader. I am
highly grateful to Libby Shaw, Tim McClure and Dr. Yong Zhang at Center for Materials
Science and Engineeing (CMSE) at MIT and Dr. Richard Schalek and Nicholas Antoniou at Center for Nanoscale Systems (CNS) at Harvard University for their help and support while using the different facilities.
I would like to express my thanks to my colleagues, Joe Fricano, Burc Misirlioglu and
Bulat Katsiev, who helped and supported me at different stages of my thesis.
I would also like to thank my friends Sukant Mittal, Dipanjan Sen, Srikanth Patala,
Vignesh Sundaresh and Yu-Chih Ko, who have been my constant support system here at
MIT and with whom I have had some wonderful experiences, and whose friendship I will value for life.

Last but not the least, I sincerely owe my deepest thanks to my parents and my sister, who always encouraged me in all my endeavors and believed in me all these years.
Vivek Inder Sharma

E
R
T
AG Gibb's Free Energy change for a reaction
AGO Gibb's Free Energy change for a reaction at standard conditions (25
0
C and 1 atm) ai Activity of species
Reversible voltage of electrolysis
Universal gas constant
Temperature
Cell overpotential

TABLE OF CONTENTS
1 Introduction............................................................................................................. 13
1.1 Operating Principle of SOECs ...........................................................................
1.2 Cell Thermodynamics and Kinetics..................................................................
14
15
1.3 C ell C om ponents............................................................................................... 17
1.3.1 E lectroly te ............................................................................................... . . 1 7
1.3.2 O xygen E lectrode................ ............................ ......................................... 18
1.3.3 H ydrogen Electrode .................................................................................. 19
1.4 Perovskite structure and mixed conduction mechanism................................... 21
1.5 Organization of the thesis ................................................................................. 23
2 Degradation Mechanisms and Research Objective .......................................... 25
2.1 B ackground ........................................................................................................... 25
2.1.1 C hrom ium poisoning.................................................................................. 27
2.1.1.1 Gas and solid-state diffusion of Cr-containing species from the interconnect into the oxygen electrode .............................................................. 27
2.1.1.2 Cr-deposition: Role of nucleation agents............................................... 28
2.1.1.3 Chromium deposition in bond layer ..................................................... 29
2.1.2 Cation interdiffusion and segregation ....................................................... 30
2.2 R esearch O bjective ............................................................................................ 33
3 Approach and Techniques Employed................................................................
4 Results and Discussion.........................................................................................
35
45
5
4.1 Preliminary identification of secondary phases at the surface of LSC bond layer:
R am an Spectroscopy.................................................................................................. 46
4.1.1 Sum m ary ................................................................................................. . . 50
4.2 Bond layer and oxygen electrode microstructure and surface composition:
Nanoprobe Auger Spectroscopy ............................................................................... 50
4.2.1 A SM oxygen electrode................................................................................... 52
4.2.2 L SC bond layer .......................................................................................... 56
4.2.2.1 Cr presence in LSC and cation segregation at electro-catalyst surface .... 56
4.2.2.2 Cation variation across the LSC cross-section...................................... 59
4.2.2.3 Migration of Sr and Co to LSC/interconnect interface.......................... 62
4.2.2.4 Cation Migration Mechanism: Experiments on reference cells in controlled environment ......................................................................................
4.2.3 Sum m ary .................................................................................................
65
. . 74
4.3 Cation distribution in the bond layer at a nano-scale: TEM/EDX..................... 76
4.3.1 Elemental distribution in the bond layer .............................. .......... 77
4.3.2 Possible secondary phases in the bond layer ......... ............................ 80
4.3.3 Sum m ary ................................................................................................. . . 82
Conclusions.............................................................................................................. 84
6 Appendix A Techniques Used............................................................................. 88

6.1 Ram an Spectroscopy.........................................................................................
6.2 A uger Spectroscopy ...........................................................................................
6.3 Transm ission Electron M icroscopy ..................................................................
6.4 Focused Ion Beam .............................................................................................
7 A ppendix B A ES D ata ......................................................................................
8 R eferences..............................................................................................................
88
89
91
92
95
100

LIST OF FIGURES
Figure 1-1: Representation of a SOEC .......................................................................... 15
Figure 1-2: Structure of perovskite type oxides [7]....................................................... 21
Figure 1-3: Bulk oxygen transport mechanism, involving random hopping of oxygen ion vacancies on the oxygen sub lattice [7] ............................................................... 22
Figure 1-4: The various steps in the electrochemical reaction of oxygen reduction in a
SOFC. Oxygen is adsorbed at the surface and accepts electrons to form oxide ions, which are transported to the electro-catalytically active sites through either the bulk
(a) or surface (b) paths [7] ................................................................................... 23
Figure 2-1: Operation history of the SOEC stack showing loss in performance.......... 25
Figure 2-2: Bright-field TEM image of oxygen electrode of a SOFC showing presence of
Cr-containing nanoparticles. EDX and SAED pattern of nanoparticles corresponds to (Cr, M n)
3
0
4 spinel structure [22] ...................................................................... 29
Figure 2-3: a) SEM image of the cross section of a SOFC showing the steel interconnect, the LSM bond layer and the scale formed at their interface after heat treatment in air at 8000 C for 300 h with elemental maps of Mn (b) and Cr (c) [23].................... 30
Figure 2-4: Chemical map analyses of La, Sr, Mn, Zr, Y and 0 in a small area in SOFC cathode showing clear inter diffusion between the YSZ and LSM grains [26]........ 32
Figure 3-1: A picture of one of the tested cells studied in the thesis ............................ 36
Figure 3-2: (a) Side view of the reference cells REF#1 and REF#2, operated under controlled electrochemical conditions, without Cr-containing layers, with its specific dim ensions in top view shown in (b).................................................................... 37
Figure 3-3: SEM image of sample showing ROI to prepare TEM membrane with protective coating of Pt on it.................................................................................. 40
Figure 3-4: FIB image of sample showing under-cuts on the membrane to free it from 2 of the 3 sides that it is attached to the sample. The probe is brought in touch with the membrane and is attached to it using a deposition of Pt, C or W ......................... 40
Figure 3-5: SEM image of TEM membrane attached to the Cu-grid. It is finally thinned down to electron transparency using FIB milling................................................ 41
Figure 3-6: Copper grids used to prepare TEM samples .................................................. 43
Figure 3-7: Cross-sectional view of the bond layer and oxygen electrode showing where the TEM samples are prepared from. To compare the changes in microstructure and chemical composition across the bond layer cross-section, we need to prepare TEM samples using traditional milling from both the bond layer/interconnect and bond layer/oxygen electrode interfaces. On the other hand, one FIB prepared is enough for the sam e requirem ent ............................................................................................ 43
Figure 4-1: Raman spectrum for cell CER#1 showing ASM peak................ 46
Figure 4-2: Raman spectrum for cell CER#2 showing peaks for LaCoO
3
........
.. . . . . . . . . .
47
Figure 4-3: Raman spectra collected from a point each from the dark and light region on the oxygen side of cell CER#3, showing the presence of new phases formed after decomposition of the perovskite structure ........................................................... 48
Figure 4-4: Raman spectrum collected from region with the exposed electrolyte on the oxygen side of cell CER#3, showing only the zirconia peaks below 1000cm-. Here the prefix 'm ' stands for m onoclinic.................................................................... 49

Figure 4-5: SEM image of the cross section of cell CER#2............................................. 51
Figure 4-6: SEM image of the cross section of cell CER#2, showing densified regions in
A S M .......................................................................................................................... 5 1
Figure 4-7: SEM image of a region in ASM layer of cell CER#1, showing uniform rounded grains...........................................................................................................
Figure 4-8: SEM of ASM region from cell CER#3 showing faceted grains with sharp and A SM layers......................................................................................................
52 boundaries .................................................................................................................
Figure 4-9: SEM image of cross-section of CER#3 showing delamination between LSC
53
54
Figure 4-10: SEM image of the cross section of cell CER#3 (LSC region).......... 56
Figure 4-11: AES data from points 1, 2 and 3 in Figure 4-10 ....................................... 57
Figure 4-12: SEM image of region 2 in cell CER#3 ........................................................ 58
Figure 4-13: AES spectra from the areas in Figure 4-12, showing Cr fraction as high as
0 .3 4 ............................................................................................................................ 5 9
Figure 4-14: Variation of the averaged cation content, normalized with respect to the sum of La, Sr, Co and Cr, in the bond layer as a function of cross-sectional region for
C ell C E R #3 ............................................................................................................... 6 1
Figure 4-15: Variation of the averaged cation content, normalized with respect to the sum of La, Sr, Co and Cr, in the bond layer as a function of cross-sectional region for
C ell C E R # 5 ............................................................................................................... 6 1
Figure 4-16: SEM image of the LSC surface of cell CER#5, showing numerous crystallites. NAES analysis showed these crystallites to be rich in Co ................. 62
Figure 4-17: AES spectrum from one of the Co-rich crystallites on the surface shown in
Figure 4-16 compared with that from the LSC interior cross-section .................. 63
Figure 4-18: AES spectrum from the region free of Co-rich crystallites on the surface shown in Figure 4-16 compared with that from the LSC interior cross-section....... 63
Figure 4-19: Schematic representation of the movement of various species across the cell cross-section ..............................................................................................................
Figure 4-20: Cell performance for REF#1: Under a constant current density of 0.4A/cm
2
64 at 8200C in air, the potential difference across the cell increased with time and stabilized after 108 hours of operation.................................................................. 66
Figure 4-21: Nyquist plots for anode (oxygen evolution in electrolytic mode) of REF# 1, before and after its operation ................................................................................. 67
Figure 4-22: Nyquist plots for cathode (oxygen reduction in fuel cell mode) of REF#1, before and after its operation ................................................................................. 67
Figure 4-23: SEM image of LSC surface of cell REF #1 showing that it does not have secondary crystallite form ation............................................................................. 68
Figure 4-24: SEM image of REF#1 cross-section ............................................................ 68
Figure 4-25: Comparison of AES spectra from LSC cross-section of cell REF#1 against that of cell C E R #5................................................................................................. 69
Figure 4-26: SEM image of LSC surface of cell REF #1 ................................................. 70
Figure 4-27: AES spectra from the 3 areas shown in Figure 4-26. The three areas probed have similar chemical composition, showing that REF# 1 has not experienced nonuniform segregation of cations at surface of LSC grains (unlike CER#3 and CER#5)
................................................................................................................................... 7 0

Figure 4-28: Comparison of AES spectra from LSC surface to cross-section, for cell
REF# 1. The composition at LSC surface is uniform throughout, both at the surface and across the cross-section................................................................................. 72
Figure 4-29: Comparison of AES spectra from LSC surface of cell REF#1 to REF#2 ... 72
Figure 4-30: Comparison of AES spectra from LSC layer as anode vs cathode for REF#1.
Even though anode and cathode degrade unequally, they show similar microstructure and surface composition ............................................................... 74
Figure 4-31: (a) A dark field TEM image of a region (2.5ptm x 1.7pm) of the TEM sample A with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr .............. 78
Figure 4-32: (a) A dark field TEM image of a region (1.3ptm x 0.9pm) of the TEM sample A with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr ............... 79
Figure 4-33: (a) A dark field TEM image of a region (1.3pim x 0.9pm) of the TEM sample B with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr ............... 80
Figure 5-1: Possible mechanism for Cr-reaction in the LSC bond layer microstructure.. 87
Figure 6-1: Pictorial view of the steps in Auger spectroscopy. a) Incident photon ejecting an electron from the shell of an atom b) A secondary 'Auger' electron ejected from an outer shell............................................................................................................. 90
Figure 6-2: a) Schematic illustration of the H-bar FIB technique. b) SEM image showing the top-down view of an H-bar FIB specimen (taken from [35]).......................... 94
Figure 6-3: In-situ lift out of a sample piece. a) A wedge shaped specimen being dug out from the bulk. b) The wedge mounted on the TEM grid (taken from [35]) ...... 94
Figure 7-1: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 1 of CER#3 ......................................... 96
Figure 7-2: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 2 of CER#3 ......................................... 96
Figure 7-3: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 3 of CER#3 ......................................... 97
Figure 7-4: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 1 of CER#5 ......................................... 97
Figure 7-5: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 2 of CER#5......................................... 98
Figure 7-6: Cation composition, normalized with respect to the sum of La, Co and Cr, for different AES acquisition areas in region 3 of CER#5......................................... 98

LIST OF TABLES
Table 3-1: Summary of our characterization approach and techniques used with their respective objectives ............................................................................................ 35
Table 3-2: Summary of FIB conditions used to prepare TEM samples ............. 42
Table 4-1: Nomenclature for the cells studied in this thesis.......................................... 45
Table 4-2: Conductivity data for the chemical species that were found in the Raman spectra of tested SOEC oxygen electrodes. Conductivity values are obtained from literature and from Ceram atec Inc. ....................................................................... 50
Table 4-3: Comparison of ASM compositions normalized with respect to the sum of A,
Sr and Mn cations, before (cell CER#1) and after use (cell CER#3) ................... 53
Table 4-4: Variation of cation percentages, normalized with respect to the sum of A, Sr and Mn cations, over the different regions of ASM for cells CER#3 and CER#5. For as-prepared ASM, the average fractions of A, Sr and Mn were seen to be 0.63, 0.15, and 0.21, respectively. This corresponds to a (A+Sr)/Mn ratio of 3.71, different from the results from the surfaces of the tested cells CER#3 and CER#5. ................... 55
Table 4-5: Variation of cation fractions, normalized with respect to the sum of La, Sr, Co and Cr compositions, over the different regions of LSC for cells CER#3 and CER#5
......................................................... ......................................................... . 60
Table 4-6: Nomenclature for TEM samples prepared by traditional ion milling ...... 77
Table 4-7: Comparison of chemical composition in the LSC bulk and at LSC surface. It is observed that in LSC bulk, Cr-rich regions are deficient in Co, and vice-versa. Crcontent in the LSC bulk is more than at LSC surface............................................ 81

The environmental consequences of energy production have led many nations in the world to impose stricter guidelines on the production and consumption of energy.
Further, the search for new sources of energy and more efficient means of employing energy has accelerated. Engineers and scientists have always looked for forms of energy that could be produced with reduced emissions and with more efficiency [1]. Phenomena such as global warming have led to an ever-increased emphasis on clean sources of energy.
Nuclear fission energy is one of the promising sources for clean energy at large scale. In addition to electricity production, nuclear plants can also be used for the production of alternative transportation fuels such as hydrogen gas. High temperature electrolysis is one of the most efficient electrochemical processes for hydrogen generation from water with no CO
2 emissions using electricity and heat from nuclear plants [2, 3]. It is carried out in devices called Solid Oxide Electrolytic Cells (SOECs) at high temperatures. In SOECs, either steam or a mixture of steam and carbon dioxide can be used as a fuel to produce hydrogen or syn gas (H
2
+ CO), respectively. The syn gas can be used to produce liquid fuels. Since the conversion takes place via an electrochemical process and at high temperature, the process is clean, quiet and efficient. In addition to low or zero emissions, the benefits of electrolytic cells include high efficiency and reliability, multi-fuel capability, siteing flexibility, durability and ease of maintenance, just as in Solid Oxide

Fuel Cells (SOFCs)[4]. Besides, due to solid electrolyte construction, with impermeability to gas flow, the problem of gas crossover from one electrode to the other is obviated. Another benefit of SOECs is the elimination of expensive catalysts such as platinum, rhodium etc. due to the high temperature of operation. Electrolytic cells are also scalable and can be stacked modularly until the desired hydrogen throughput is reached. SOECs also provide the advantage of being compact (they have small volume to weight ratio), especially in planar design.
At the same time, it must be kept in mind that the SOFC and SOEC technologies are accompanied by challenges, the most important of them being their degradation which leads to a loss in their performance over a period of operation. As will be discussed later in this thesis in more detail, at the temperatures of operation, the different components of the SOFCs and SOECs can interact with each other, forming secondary phases that hamper functionality. This degradation due to chemical interactions becomes more severe at higher temperatures.
1.1 Operating Principle of SOECs
A SOEC is an electrochemical device that converts electrical energy into chemical energy, based on the converse of the working mechanism of a Solid Oxide Fuel Cell
(SOFC). It uses electricity from an external source and can be used to convert steam (or a mixture of steam and carbon dioxide) into hydrogen (or syn gas).
An electrolytic cell consists of the hydrogen electrode (cathode) and the oxygen electrode
(anode), with an oxygen ion-conducting electrolyte sandwiched between them. The electrodes are electrically coupled to an electricity source by external lines outside the cell. Figure 1-1 below shows a schematic of one such SOEC. Oxygen passes over the anode, and hydrogen and steam over the cathode. At the cathode, two water molecules accept four electrons to form hydrogen gas and oxide ions. The oxide ions, then, move through the electrolyte and go to the positive electrode (anode) where they lose electrons and form oxygen gas. Cells are stacked together, and metallic interconnects are used

between cells. Cell interconnects provide cell-to-cell electrical connection and separate fuel and oxidant gas atmospheres in a stack.
The reactions involved in a SOEC are:
At cathode: H
2
0+2e- ->0
2 -+ H
2
(1)
At anode: 02- _1 _02+2e-
2
(2)
Overall: H20 ->-02+ H2
2
(3)
Fuel in
O e-| X
0--
4 Air in
Excess
Fuel ando
H
2 out
__|"2
Caffiode Anode
Electrolyte
Figure 1-1: Representation of a SOEC
-.
Air out
1.2 Cell Thermodynamics and Kinetics
The splitting of water into hydrogen and oxygen is an endothermic reaction. The Gibb's free energy of this reaction (H
2
0(g) ->
1
Hz,)
+
-O2,)) can be written as:

where,
AGO + RTln a2a aH20
0
AGO >0
2
(4)
(5)
The reactions in SOEC are normally carried out at high temperature, and heat from a nuclear reactor or concentrated solar plant can be used to provide the thermal and electrical energy. A high temperature is preferred because that leads to a decrease in the
Gibbs' free energy (electrical input) that is required for the reaction to occur.
The reversible voltage of the electrolysis is:
E AG neF
(6) where ne is the number of electrons exchanged and F is the Faraday's constant. Since the
Gibb's Free energy change for the electrolytic reaction is positive (energy is required to split water into hydrogen and oxygen), the voltage obtained by equation 6 above is negative. Thus, substituting Eq. 4 into Eq. 6, we can obtain the reversible voltage in terms of the Nernst Equation: as a Y
E=AGO+RTln
2 2 aH20
(7)
However, in actual electrochemical systems, there are factors that lead to energy dissipation. Hence, the voltage required for a reaction is higher than the theoretical value.
This energy loss (difference between the actual and the theoretical voltage) is called overpotential (rj).
1=E Eth (8)
In electrolytic cells, this energy loss occurs due to irreversibilities arising as a result of three primary mechanisms: a) the 'activation' overpotential the extra energy required to provide the activation energy for the surface reaction (involving the electron transfer) to take place, b) the 'ohmic' overpotential the extra energy required to overcome the resistance to the diffusion of electrons and ions in the system and c) the 'concentration'

overpotential energy required for the mass transfer of the ions and electrons in the system [5]. These are denoted by fa, ino and ie respectively.
1.3 Cell Components
The electrodes for SOECs have to possess properties important for their proper functioning, including high electronic conductivity, high catalytic activity and compatibility with other cell components. Since we will be comparing some of the degradation results to those obtained for SOFCs, to avoid confusion regarding an electrode acting as anode in one mode and as cathode in the other, we will call them
'oxygen electrode' and 'hydrogen electrode' respectively, for both SOEC and SOFC. For example, the oxygen electrode is the anode for SOEC mode and cathode for SOFC mode of operation. Another target of the materials selection is to minimize the overpotential losses (discussed in the previous section) to achieve the maximum hydrogen (or syn gas) production rates, at the temperature of operation.
1.3.1 Electrolyte
The electrolyte for SOFCs and SOECs must stay stable and must have high ionic conductivity and very low electronic conductivity at the operating temperatures. Besides they must be able to be formed into thin and very strong films. Zirconia based electrolytes, possessing the fluorite structure, have been the most popular choice for the electrolyte material. For cells operating at high temperatures (>800
0
C), doped zirconia, especially yttria stabilized (YSZ) is most commonly used because of the long-term chemical stability of YSZ and compatibility with electrode materials. Doped ceria- and lanthanum gallate based electrolytes have been researched for use in low and intermediate temperatures of operation (500 800
0
C) [6]. But ceria based electrolytes are susceptible to an internal short circuit leading to the conduction of electrons along with the ions. More recently, a number of other materials, especially perovskites, brownmillerites and hexagonal structured oxides have also been investigated as potential

electrolyte materials as they possess good ionic conductivity. These include materials like lanthanum gallate and calcium titanate [1].
Oxygen Electrode
The doped ABO
3 perovskite family of compounds (A = Lanthanide group and B
= transition metal group) is the most common set of materials used as oxygen electrodes in
SOECs inspired from SOFCs. Sr-doped lanthanum manganites (LSM) are conventionally put to use as oxygen electrodes. The A-site dopant ratio is usually modified to improve the conductivity and stability of the perovskite. In the early stages of SOFC development,
LaCoO
3 was investigated as the oxygen electrode and its initial performance in the cells was good. But severe degradation was seen to occur over long hours of operation due to its reaction with the electrolyte (YSZ) [1]. Investigations on the oxygen electrode then focused on doped lanthanum manganites. Sr- and Ca- doped lanthanum manganites have conventionally been used as the oxygen electrode at 800 1000 0 C [6]. The degradation in manganites was not seen to be as severe as for cobaltites, but some reactions with
YSZ, especially at higher temperatures, were recognized [1]. Further modifications such as doping of the A- and B-sites can lead to performance improvement and dimensional stability. At lower temperatures (<800 0 C), the decrease in ionic conductivity in manganite electrodes leads to higher polarization. Efforts have also been made to use cobaltite-based electrodes with ceria-based electrolytes because ceria has less reactivity with perovskites, compared to YSZ. It is also possible to use lanthanum strontium cobaltite based electrodes with ceria-based electrolytes at intermediate temperatures [1].
Composite ceria-lanthanum ferrites can be used at intermediate temperatures [6].
However their high temperature phase behavior and long-term stability remain unclear and need further investigation.

1.3.3 Hydrogen Electrode
Like the oxygen electrode, the hydrogen electrode must combine catalytic activity with electrical and ionic conductivity. Ionic conductivity allows the cathode to spread the oxide ions across a broader region of electrode/electrolyte interface and electronic conductivity ensures the transport of electrons required for the electrode reaction from the external circuit. The cathode generally consists of a nickel-zirconia cermet electrode. The conventional cermet microstructure consists of uniform zirconia dispersion in the Ni metal with metallic inter particle contact for electronic conduction. In the initial stages of
SOFC research, precious metals such as platinum and gold and transition metals such as nickel and iron were tested as hydrogen electrodes. But platinum degrades rapidly in a
SOFC and nickel aggregates at high temperatures, inhibiting the fuel access [1]. Mixing
YSZ electrolyte particles with a nickel matrix to form a composite electrode solved nickel aggregation problems [1]. Such composite electrodes were seen to provide adequate performance at the temperature of operation. Ni-ZrO
2 cermet electrodes have demonstrated long-term stability and high electrical conductivity over wide temperature ranges (600 1000
0
C). But they undergo a performance reduction in the presence of sulfur impurities due to the formation of Ni
3
S
2
. Other materials used for the hydrogen electrode are copper based ceria electrode and redox and sulfur tolerant ceramic singlephase electrodes [6]. In the copper-ceria cermet electrode, copper suppresses carbon formation during the reaction while ceria improves the reaction kinetics. But copper based electrodes are susceptible to sulfur poisoning and high temperature sintering.
Electrodes such as gadolinia-doped ceria have been investigated recently to minimize carbon deposition and increase tolerance to sulfur [6].
Throughout the different ranges of temperature and gas environment, it is necessary to maintain the necessary compatibility with other materials with which the electrodes come into contact, specifically the electrolyte, interconnect and any relevant structural components. Physical compatibility requires a match of thermo-mechanical properties and for chemical compatibility, there should be no solid-state contact reaction.

Compatibility must also extend to the behavior of materials towards the ambient gases including corrosion or poisoning by trace impurities.
We should note that an increase in the operating temperature also increases the degradation rate of the electrolytic cells. This degradation is principally due to the chemical interaction between the electrolyte and the electrodes at the electrolyte/electrode interface at the high temperatures of operation, which leads to the formation of secondary phases that block the active sites where electro-catalysis occurs. These secondary phases block the bulk diffusion path (discussed in the following sub-section) in the perovskite materials, leading to a decrease in the electronic and ionic conductivity and hence a loss in performance of the cell. Thus, it is important to maintain an optimum balance between these conflicting effects of the operating temperature. Currently, SOECs are tested at 800
830
0
C. In order to minimize degradation, operating temperatures down to 5000C can be favorable. Besides, high temperature causes thermal expansion mismatches among different cell components. Also, high temperature operation places severe constraints on material selection and fabrication process.
As we have seen in the preceding discussion, the loss in performance of cells due to degradation over a period of time is a commonly seen behavior. As a result of this degradation, the rate of hydrogen production goes down with time. A better understanding of the mechanisms of degradation can help us identify ways and compositions that can counter the loss in cell performance. The primary focus of our research is to investigate the degradation mechanisms, primarily in terms of the microstructural and chemical composition changes coming about in the oxygen electrode and the bond layer coated on the oxygen electrode to provide better contact between the electrode and interconnects both of which have the perovskite structure, it is worthwhile to discuss the perovskite structure and the electronic as well as ionic conduction mechanisms in perovskites. The following sub-section is devoted to the same.

1.4 Perovskite structure and mixed conduction mechanism
The perovskite structure has the generic form of ABO
3 as illustrated in Figure 1-2. The A sites at the corners of the cubic unit cell are generally occupied by the lower valent cations (La, Sr, Ca and Pb). On the other hand, the B sites in the cube centers are occupied by the higher valent cations (Ti, Cr, Ni, Fe, Co and Zn). The B-site is surrounded by six oxygen ions in an octahedral co-ordination. The perovskite structure is very interesting because there are two cation sites on which to substitute lower valence cations, leading to a much wider range of possible oxide-ion conducting materials.
A
Figure 1-2: Structure of perovskite type oxides [7]
Usually, the A and B-sites in the perovskites are doped in order for them to have better structural properties as well as to improve their electronic and oxygen ion conductivity.
The enhanced electronic and ionic conductivity of doped perovskites is a result of the change in stoichiometry. Normally, the A-site cation is doped with divalent alkaline earth elements such as Sr, Ba etc. To maintain charge neutrality, this either changes the valency of some of the B-site cations and/or leads to formation of oxygen vacancies. This change in the valence state and the production of vacancies in the crystal structure is responsible for the rise in the electronic and ionic conductivity of the perovskite. An example illustrating the formation of vacancies is explained here. In the example considered here, LaCoO
3 is the host material in which SrCoO
3 is dissolved. The dissolution of SrCoO3 can be represented as (following the Kroger-Vink notation)
SrCoO
3
(-+ LaCoO
3
) -> Sr + Coc, + 30x (9)

The incorporation of Sr
2
+ thus leads to charge compensation by the formation of Co 4
+ ions and/or formation of oxygen vacancies to maintain electrical neutrality [8].
2og
1
+ O0x
2
+ Co-
(10)
(11)
The oxygen vacancies formed in the bulk are doubly positively ionized. Thus the doping of A and B site cations leads to their valency change and/or formation of oxygen vacancies. These are responsible for the good ionic and electronic conductivity of the perovskites. The oxygen ion vacancies in the perovskite structure allow for bulk transport of oxide ions. Oxide ions can hop from one vacancy to the other this hopping mechanism as shown in Figure 1-3 providing for the bulk transport. The mixed ionicelectronic conductivity leads to a large effective active area.
oxygen vacancy
Figure 1-3: Bulk oxygen transport mechanism, involving random hopping of oxygen ion vacancies on the oxygen sub lattice [7]
In addition to the bulk transport, surface transport is also important in perovskites for use as electrolytic and fuel cell electrodes. The first step in the reaction, i.e. adsorption and/or partial reduction of oxygen on the surface of the electronic phase (for SOFC mode of operation) precedes the subsequent bulk or surface transport of 02- or O' ions to the electrode/electrolyte interface. This is schematically shown in the Figure 1-4 below.
Ultimately, for the bulk transport of the oxide ions to proceed, the oxide ions have to be exchanged between the surface and the bulk of the electro-catalyst. In the SOEC mode of operation, the direction of these reaction steps is reversed. Oxide ions are formed at the interface of the hydrogen electrode/electrolyte (as shown in Figure 1-1) and travel across

the electrolyte to the oxygen electrode. Then they are transported through the electrode bulk and reach the surface of the electro-catalyst where they eventually lose electrons to form oxygen gas, which is desorbed from the surface. Both the SOFC and SOEC modes of operation, thus, require that both the bulk and the surface of the oxygen electrode have oxide ion vacancies so that the bulk transport of the oxide ions and the adsorption/desorption of the oxide ion is possible at the surface.' a) b
102
O,
Figure 1-4: The various steps in the electrochemical reaction of oxygen reduction in a SOFC. Oxygen is adsorbed at the surface and accepts electrons to form oxide ions, which are transported to the electro-catalytically active sites through either the bulk
(a) or surface (b) paths [7]
From this point on in the thesis, we will refer to the cells obtained from Ceramatec Inc.
The oxygen electrode of the cells under analyses in our project is a perovskite material
Ao
8
Sro.
2
MnO
3
, referred from here on as ASM (The element A is not disclosed here because it is proprietary information). Scandia stabilized zirconia (ScSZ) is used as the electrolyte. The cathode consists of a Ni-ScSZ cermet. Also, a lanthanum strontium cobaltite (Lao.Sro.
2
CoO
3
, also known as LSC) bond layer is used to provide better contact between the oxygen electrode and the metallic interconnects. Its higher conductivity is another reason for its use. LSC also has the perovskite structure.
1.5 Organization of the thesis
In this thesis, we are investigating the degradation mechanisms that contribute to the loss in performance of SOECs. We focus on the oxygen electrode and study two main degradation modes: 1) chromium diffusion into the oxygen electrode and the LSC bond

layer and 2) segregation of the cations of the LSC catalyst. The outline of the report is as follows: we discuss the background of the SOEC and SOFC degradation in Chapter 2, followed by the approach used in Chapter 3. The details of the results obtained are discussed in Chapter 4 and we summarize the key conclusions in Chapter 5. At the end of the thesis, we provide an Appendix A, which explains briefly the various experimental techniques employed for the research. Appendix B provides the entire AES data collected.

2.1 Background
A lot of interest in electrochemical devices based on solid electrolytes has been generated over the last few decades. It has been observed that over a long period of operation, the
Solid Oxide Fuel Cells and Solid Oxide Electrolytic Cells show a loss in performance [9].
As summarized in Chapter 1, this loss in performance is attributed to a range of degradation modes. O'Brien et al [10] measured an approximately 18% loss in production of hydrogen over 1000 hours of operation of SOECs. The following Figure 2-1 shows the operation history of the SOEC stack, the cells from which are under analysis in our project (data from Ceramatec Inc.).
2x60 Half ILS Module Load History
10~
500.0
1000.0
Time (hours)
1500.0
2000.0
2500.0

These cells are operated at a constant voltage, and as is seen from Figure 2-1, current passing through them keeps decreasing with time, indicating degradation of the cells.
Similar degradation issues are important in SOFCs as well but research to date has enabled a better understanding of SOFC degradation mechanism and has helped control the degradation to as low as 1.7% per 1000 hours of operation [11].
Various tests were conducted in order to investigate the degradation mechanisms on the oxygen and hydrogen electrodes of SOFCs. Degradation can be characterized in terms of the chemical and micro-structural changes that occur in the active components of the electrolytic cell. There are many possible underlying degradation mechanisms for the cells to develop a high resistance during operation. The reasons that relate to the electrodes include reactions between electrode/electrolyte leading to the formation of
highly resistant oxides, formation of local hot spots leading to local changes in microstructure and material properties, electrode/electrolyte delamination, and poisoning
by external chemical species [12]. For oxygen electrodes, if the electrodes are prepared
by co-firing, they have inferior electrochemical performance than those that are screenprinted. The co-fired electrode shows higher degradation over 8000 C even without polarization. This may be due to imperfect morphology, over-sintering and interfacial delamination [13]. Thus, we see that the procedure in which the electrodes are prepared is also an important factor in the degradation rates. Another reason for the degradation can be attributed to the change in conductivity of the doped zirconia electrolyte due to transformation of the phase into a less symmetrical crystal structure, an increase in the grain boundary resistivity due to grain boundary segregation of glassy phases and precipitation of long-range ordered phases such as Y
2
Zr
2
O
7 or Y
4
Zr
3
O
12 discussed by
Kondoh et al [14].
Chromium poisoning of the oxygen electrode, segregation of select cations to the surface of the catalyst and inter-diffusion of cations between the electrolyte and oxygen electrode grains are considered crucial processes leading to the degradation of SOECs, and are of interest to us in our research. We will devote the subsequent sections in this chapter to discuss each in more detail.

2.1.1 Chromium poisoning
Previous studies on SOFCs and SOECs show that diffusion of chromium from the metallic interconnects into the electrodes leads to loss in cell performance. The so called
"chromium poisoning", especially on the oxygen electrode, was previously recognized
[15, 16]. It occurs due to the Cr volatilization from the steel that is used for the cell interconnects. Cr deposition has been attributed to both chemical and electrochemical mechanisms [15]. For an electrochemical reaction, deposition can occur only where both ions and electrons are available. This can occur only at the three-phase gas-electrolyteelectrode interface for a purely ionic conducting electrolyte and a purely electronic conducting oxygen electrode. However, if the oxygen electrode possesses ionic conductivity, then the deposition can occur away from this three-phase interface and thus alter its effect on the cell performance to a larger spatial extent. In this section, we discuss two modes of Cr-related degradation mechanisms that have been observed in SOFC literature, namely a) the gas- and solid-state diffusion of Cr-containing species from the interconnects into the oxygen electrode and b) the theory regarding the importance of a
"nucleation agent" required for the deposition of Cr-containing species in the oxygen electrode.
2.1.1.1 Gas and solid-state diffusion of Cr-containing species from the interconnect into the oxygen electrode
According to the first hypotheses, chromium poisoning of SOFC cathodes (oxygen electrode) occurs by gas phase transport of chromium from the interconnect material to the oxygen electrode. The chromium transport occurs primarily through the formation of
Cr6*-containing gaseous species such as Cr0
3 or CrO
2
(OH)
2 from oxidation of chromium oxide in the interconnect [17]. The volatile Cr species are reduced at the triple phase points of electrode, electrolyte and air and form solid Cr
2
0
3 and other Cr-rich phases thereby inhibiting the electrochemistry of the cell [18]. This effect is called 'poisoning' of the electrode by gaseous Cr species. An accompanying hypothesis suggests that along with vapor phase, solid-state diffusion of the chromium containing species into the

oxygen electrode is also an underlining mechanism of the chromium deposition [19].
2.1.1.2 Cr-deposition: Role of nucleation agents
Another important and recent hypothesis suggests that the Cr deposition process at the oxygen electrode is kinetically limited by the nucleation reaction between the Cr species being transported and a "nucleation agent" on the electrode [20]. According to this theory, a nucleation agent is necessary for the volatile Cr species to be able to deposit, or react and stabilize, on the active sites in the oxygen electrode. The nucleation reaction can be represented as follows:
Cr
2
O
3
(oxide)
-+
CrO
3
(g)
CrO
3
(g)+N -+ Cr -N O,(nuclei)+0
2
Cr-N-O,+ CrO
3
-+ Cr20
3
(12) where N is the nucleation agent. In the case of the LSM electrode, the nucleation agent was identified to be the manganese species (Mn 2 ), and for the (La,Sr)(Co,Fe)0
3
(LSCF) electrode, it was the SrO species segregated at the electrode surface [20]. In case of
LSCF, Cr species deposited on the electrode and electrolyte surfaces, forming isolated chromium particles [21]. Experiments concluded that the electrodes that lack this nucleation agent show no or much less Cr deposition than electrode materials which had the nucleation agent [20]. Mn-free La(Nio.
6
Feo.
4
)0
3
(LNF) and Sr-free
(Lao.
6
Bao.
4
)(Coo.
2
Feo.
8
)0
3
(LBCF) were studied as Cr-tolerant electrodes. These studies showed that both LNF and LBCF are more stable for oxygen reduction (in SOFCs) in contact with a Fe-Cr interconnect alloy than LSM and LSCF respectively, under the same experimental conditions [20].
Analytical and High Resolution Transmission Microscopy have been employed to study, in depth, the degradation mechanisms related to chromium poisoning. In a recent study on SOFCs with LSM constituting the oxygen electrode [22], changes from the original electrode microstructure, the observation of chromium containing nanoparticles and, the

observation of Cr in solid phases in the electrode were attributed to Cr-poisoning. It was observed that the electrode pores were filled with nanoparticles and decomposition of the electrode material was noticed. The electrolyte grains were covered by Cr containing phases (Cr
2
O
3 or spinel phase i.e, MnxCr
3
.x0
4
) as shown in Figure 2-2 [22].
Figure 2-2: Bright-field TEM image of oxygen electrode of a SOFC showing presence of Cr-containing nanoparticles. EDX and SAED pattern of nanoparticles corresponds to (Cr, Mn)
3
0
4 spinel structure [221
2.1.1.3 Chromium deposition in bond layer
Previous investigations have shown Cr presence, not only in the oxygen electrode but in the bond layer as well. The most commonly used bond layers have been manganites
(LSM) and cobaltites (LSC, same as in this thesis). According to one of the theories, the degradation in the SOFC performance then occurs due to a poorly conducting oxide scale that is formed at the bond layer/interconnect interface [23]. The electronic nature of this scale and its thickness depends upon the composition of the bond layer. This oxide scale has much lower electronic conductivity than the bond layer. For example, there has been evidence of formation of a (Mn,Cr)
3
0
4 spinel at the LSM/interconnect interface.
Furthermore, elemental mapping showed that Cr had diffused into the LSM bond layer as shown in Figure 2-3. Cr content in the LSC bond layer was observed to be much higher

than for LSM, suggesting that LSC might act as a potential 'getter' of chromium. This can lead to drastic dissociation of the LSC bond layer and hence to degradation of the cells.
LSM A
Figure 2-3: a) SEM image of the cross section of a SOFC showing the steel interconnect, the LSM bond layer and the scale formed at their interface after heat treatment in air at 8000 C for 300 h with elemental maps of Mn (b) and Cr (c) [231
Thus we see that Cr related degradation has previously been reported in SOFCs using
LSC bond layers. Though the basic electrochemical principle of operation of SOECs is opposite to that of SOFCs, similar concerns regarding Cr poisoning of the bond layer and/or the oxygen electrode remain. The deposition of Cr-containing species (in the fuel cell mode of operation) in the bond layer as well as the anode occurs by the reduction of
Cr-containing species that have Cr in its +6 vacancy. Even in electrolytic mode of operation, the oxidation of Cr
2
O
3 in the steel interconnects can form species having Cr*
6
.
The bond layer, as well as the oxygen electrode, both have electrons which can be picked up by this Cr+ 6 -containing species leading to its reduction and formation of secondary
(blocking) phases similar to that in the fuel cell mode. Similarly, since the compositions of various components are similar for both SOFCs and SOECs, the same/similar nucleation agents can be important even in electrolytic mode. Investigation of such Crrelated degradation mode is one major goal of our research.
2.1.2 Cation interdiffusion and segregation
In this section, we discuss two key processes, in addition to Cr-poisoning, that can potentially contribute to the degradation of SOFCs and SOECs cation interdiffusion and segregation. Both the interdiffusion of cations across the electrode/electrolyte interface

and the segregation of cations on the oxygen surface, could lead to localized changes in an electrode's electrochemical properties. If the resulting microstructure is inferior, then this amounts to a loss in the performance of the cells.
Cation interdiffusion between the oxygen electrode and the electrolyte has also been the subject of detailed studies [24, 25, 26] due to its importance in electrode activity and stability. The high operating temperature of the SOECs and SOFCs gives rise to mutual interaction between the components of the cells. For example, formation of highly resistant SrZrO
3 and La
2
Zr
2
O
7
by the reaction between LSM electrode and YSZ electrolyte and diffusion of Mn from LSM into the electrolyte (thus affecting the electrical properties of YSZ) has been previously documented [24]. In a recent work,
Grosjean et al. [26], using Analytical and High Resolution Transmission Electron
Microscopy, showed that the electrode material decomposed over a period of time due to diffusion of the cations from the electrode into the electrolyte grains. Figure 2-4 below shows the chemical map analyses of La, Sr, Mn, Zr, Y and 0 in a small area (2x2 pm
2
) in the SOFC oxygen electrode and shows the interdiffusion between the LSM and YSZ grains clearly. This causes instability of the electrode and enhances the reactivity between the electrode and electrolyte, which leads to formation of parasite phases like SrZrO
3
,
La
2
Zr
2
0
7
[26].
Besides interdiffusion, segregation of cations in the electrodes of SOFCs is well documented. It has been seen that certain electrodes such as LSCF, though offer high power densities than LSM electrode, do not possess long-term stability under several conditions. For example, in a SOFC with LSCF electrode and electrode-electrolyte interlayer of samarium doped ceria (SDC), Sr enrichment at the electrode/electrolyte and electrode/interconnect interfaces has been identified as one of the causes of degradation
[27]. Such segregation can alter the local and extended properties of the cell and would not only increase the ohmic resistance but also affect the oxygen reduction reactions and the charge transfer mechanisms at the interfaces of the electrodes.

Figure 2-4: Chemical map analyses of La, Sr, Mn, Zr, Y and 0 in a small area in
SOFC cathode showing clear inter diffusion between the YSZ and LSM grains [26]
Not just in SOFCs, but also in perovskite type oxide separation membranes used for oxygen production, segregation of the constituents and the formation of new phases on the outer surfaces of the membrane have been reported [28]. For example, while using
LSCF as oxygen separation membrane, the Co/Fe ratio for fresh membrane was 1:4.75
and this ratio was seen to vary from as low as 1:12 to as high as 1.82:1 after long term operation [28]. This indicates that segregation of the membrane cations occurred during oxygen separation under an oxygen gradient, and cobalt is the fastest element moving towards the membrane outer surface. In another investigation of LSCF as a SOFC electrode, severe changes in the cation concentration at the surface of the electro-catalyst occur were shown. Under cathodic polarization, the concentration of La at the surface got depleted while that of Sr and Co increased. Such segregation is expected to change the electro-chemical properties of the material surfaces and is a key component of our research goals [29].

2.2 Research Objective
It is clear from the preceding discussion that there are a multitude of reasons that can contribute to the degradation of the performance of a SOEC. The goal of our research is to identify the important causes for the loss in performance of the oxygen electrode of the
SOEC. Previous investigations on SOFCs suggest that the oxygen electrode governs a large part of the losses [30]. Therefore we focus only on the oxygen electrode for this thesis. As has been discussed in this chapter, the mechanisms of Cr related degradation are considerably well studied for SOFCs. Secondly, cation segregation from the surface of the electro-catalysts has also been reported. Nevertheless, no consistent conclusion has yet been drawn. Furthermore, there has not been much work on the same issue for
SOECs.
Our research objectives are two folds. The first objective focuses on identifying the mechanism in which the Cr and its compounds transport from interconnects and form electrochemically inactive phases in the oxygen electrode and the bond layer of the
SOECs. Particularly important to us is to study the role of a transporting agent or the so called 'nucleation agent', as has been suggested for SOFCs [20, 21], if any. Some of our preliminary results suggest that the Cr content is much larger in certain regions of the bond layer as compared to other regions, which have a very small amount of Cr. This intriguing observation has opened the possibility of Cr deposition being more favorable in the presence of certain chemical species in the perovskite bond layer and electrode, similar to the "nucleation agent" hypothesis discussed above.
The second objective in our research is to investigate the cation segregation at the surface of the catalyst. As discussed in the preceding sub-section, the formation of secondary phases due to cation interdiffusion and the alteration of the electro-chemical properties due to cation segregation especially on the surface of the electro-catalysts are also important contributors to the degradation of the cell. The possibility of this type of elemental segregation in SOECs is also of interest in our research. Our preliminary results using Auger Electron Spectroscopy (AES) show that the ratio of La/Co at the

surface of the LSC catalyst varies from as low a value as 0.6 to as high as 19 in certain regions, thus hinting towards the elemental segregation and dissociation similar to that observed in [28]. Hence a detailed investigation of segregation is a very interesting and challenging problem that we aim to address in our research.

We employ spectroscopic techniques in an integrated manner, which allows us to investigate the degradation mechanisms from a higher level and move into higher resolution analysis with subsequent steps. The techniques used and our particular objectives utilizing each technique are summarized in Table 3-1.
Table 3-1: Summary of our characterization approach and techniques used with their respective objectives ScSZ
Technique Objective
10 cm
Raman Spectroscopy
Preliminary identification of secondary phases formed on the surface of the bond layer
Electrode surface chemistry and
Nanoprobe Auger Electron microstructure and its
Spectroscopy (NAES) variation across the cross section at a small scale
Focused Ion Beam (FIB)
(pm-nm)
Selective choice of the interface of interest to prepare TEM samples from
Energy Dispersive X-Ray
Spectroscopy (EDX)
Transmission Electron
Microscopy (TEM) cheica tion formed and chemical composition and secondary structures

We investigated two batches of samples in this thesis. Batch 1 consisted of cells provided to us by Ceramatec Inc. This batch consists of actual commercial SOECs, both tested and as-prepared. The tested cells were operated at 830
0
C for 2000 hours of continuous operation and the as-prepared cells are analyzed to provide a reference so as to track the changes in chemical composition and cell microstructure over the period of operation.
Figure 3-1 is a picture of one such tested cell from Batch 1.
Peeled off bond layer and electrode
Figure 3-1: A picture of one of the tested cells studied in the thesis
We aim to investigate whether the bond layer dissociation is an affect of the Cr-related degradation, i.e. whether it is thermodynamically favored by the presence of Crcontaining species, or is an affect of migration of cations under the influence of physical impact by electronic current. To do so, we simulate the actual operating conditions on smaller Batch 2 half-cells having the same oxygen electrode and bond layer composition as the cells under analysis in our project. The schematic of these half-cells is shown in
Figure 3-2. We use gold and platinum mesh as current collectors instead of steel interconnects. This ensures that the Cr-containing species are eliminated and can thus enable us isolate the mechanism of bond layer dissociation

Worldng
Electrolyte
I
Reference
-- 'Electrode amm
1cm
Coutr,, --
Electrode
Reference
Electrode mm4m 4 44m
(a) (b)
Figure 3-2: (a) Side view of the reference cells REF#1 and REF#2, operated under controlled electrochemical conditions, without Cr-containing layers, with its specific dimensions in top view shown in (b)
The first objective of the research is the preliminary identification of the phases present on the top of the LSC bond layer near the interconnect. We used Raman Spectroscopy
[31] to achieve this objective. Raman spectroscopy was carried out in a Kaiser Optical
Instruments' Halogram 5000 series Raman spectrometer at the Center of Materials
Science & Engineering (CMSE) at MIT. For data collection, the laser incident on the sample had a wavelength of 785 nm and the incident power was 13-15 mW. We used collection fibers that achieved a spot size of 80 pm. The spectra were collected with 1 accumulation each and an exposure time of 60 s. This part of the research allowed us to investigate if the bond layer had disintegrated into the secondary phases from the original perovskite phase. This technique enables us to confirm the presence of chromium containing species at the bond layer/interconnect interface region, which subsequently drives us to investigate the microstructure of the oxygen electrode and the bond layer at a higher resolution.
The following objective is to identify the changes in the microstructure and chemical composition of the bond layer and the oxygen electrode as well as the microstructure dependent variations in the presence of the cations and chromium. This step allows for the study of segregation of the cations at the electro catalyst surface. We employ
Nanoprobe Auger Electron Spectroscopy (NAES) [32] to accomplish this. We used the
Physical Electronics Model 700 Scanning Auger Nanoprobe, available at the CMSE at

MIT. Incident electrons on the samples, for the analysis, were accelerated through a potential of 10 kV and electronic current of 10 nA was used. The range of kinetic energy of Auger electrons collected was 20 2000 eV. The smoothing and differentiation of the
AES spectra collected was carried out using the Savitsky-Golay algorithm. AES is a surface sensitive technique and provides us information about only the top few atomic layers of the surface (refer to the Appendix A for more details about the technique).
Using AES, we identify the chemical composition at the surface of the electrocatalyst at its prepared state and after the electrolytic operation. The technique has the capability to enable us to probe the surface secondary composition with high resolution (1 pm) to probe the non-uniform decomposition on the surface.
To systematically investigate the structural and chemical evolution at a high resolution in the grains and at the grain boundaries, we employ Transmission Electron Microscopy
(TEM) [33] coupled with Energy Dispersive X-Ray Spectroscopy (EDX). The information obtained from this technique complements the findings obtained from the
NAES results: TEM/EDX provides information from the bulk of the microstructure while the data in NAES is essentially from the surface of the electrocatalyst grains. TEM was carried out in a JEOL 2010F with a field emission electron source, available at the CMSE at MIT. Electrons incident on the samples were accelerated through a potential of 200 kV. The EDX microanalysis system Oxford Instruments INCA was used for collecting energy dispersive X-ray spectra. The point EDX spectra were acquired with an acquisition time of 300 s and a spot size of 2.4 nm. The quantification of the EDX was obtained by standard less analysis using the Cliff-Lorimer correction. For STEM elemental maps, a JEOL detector was used. Probe size of 0.2 nm with a camera length of
15 cm was employed. The elemental maps were collected within 60 s to 120 s each.
To prepare TEM samples, we employ two methods:
1. The Focused Ion Beam (FIB) [34, 35]
2. Traditional Ion Milling.

FIB allows us to choose a particular area of interest to investigate (e.g. electrode/bond layer or bond layer/interconnect interfaces), as the sample preparation is done in situ during SEM-imaging of the area of interest. It provides us with the flexibility of seeing the region of interest to lift out; enabling the selective choice of regions, which have undergone a change from the initial microstructure. The provision of this flexibility is our reason behind pursuing FIB over the conventional ion milling to prepare TEM samples.
Furthermore, the use of FIB for TEM allows one to repeatably investigate the comparable regions to form correlations between interface structures and interface chemistry in the
SOECs subjected to long term electrochemical conditions.
To prepare TEM samples using FIB, we used the Zeiss NVision 40 Dual-Beam focused ion beam and scanning electron microscope at the Center for Nanoscale Systems (CNS) at Harvard University. We first identify the region of interest (ROI) and deposit a protective layer of Pt on this region. This initial deposition of Pt measures 2x20 jm
2
, is around 2 pm in thickness and is shown in Figure 3-3.
After this deposition, the lift-out technique was employed to mount the membrane on
Omiprobe Cu-grids. In this technique, two trenches are cut out on either side of the ROI using FIB milling. We used an accelerating voltage of 30 kV for the incident beam of Ga ions. The trenches measure 20x8 m 2 and are dug out to a depth of 7-10 jim. The membrane is around 1 jim in thickness after this stage. Then, undercuts are made on two of the three sides that the membrane is attached to the sample and it is attached to a probe using a deposition of Pt, C or W. Once the membrane is attached to the probe as shown in
Figure 3-4, the remaining side, where it is connected to the sample, is FIB-milled to free the membrane. It is then lifted out and mounted on the Omniprobe Cu-grids using deposition of Pt, C or W (shown in Figure 3-5). Final milling is done using lower FIB currents, until the membrane is transparent to electrons. The exact conditions used for milling and depositions are summarized in Table 3-2.

Figure 3-3: SEM image of sample showing ROI to prepare TEM membrane with protective coating of Pt on it
Figure 3-4: FIB image of sample showing under-cuts on the membrane to free it from 2 of the 3 sides that it is attached to the sample. The probe is brought in touch with the membrane and is attached to it using a deposition of Pt, C or W.

Figure 3-5: SEM image of TEM membrane attached to the Cu-grid. It is finally thinned down to electron transparency using FIB milling
We faced several challenges in preparing TEM samples using FIB. The membranes that were successfully thinned down to electron transparency after being mounted on Cu-grids showed evidences of considerable FIB-related damage. To counter this, we did the final thinning currents as low as IpA. The brittle and fragile nature of the sample makes it difficult to lift it out using the probe without breaking the membrane in the process.
Thirdly, once the membrane is attached to the probe and freed from the sample by milling the last region where it is attached to the sample, the membrane repeatedly got detached from the probe in the process of lifting it out. This detachment of the once-attached membrane from the probe was a frequent occurrence and leads us to believe that there could be a development of tensile forces in the membrane when the probe is pressed against it before deposition. The tensile forces could repel the membrane from the probe when the membrane is completely freed from the sample, though this is not a conclusive theory.

Table 3-2: Summary of FIB conditions used to prepare TEM samples
Step
Membrane
Thickness
Cut Size Current Duration
Stage 1:
Preparation of thick membrane for lift-out
Coarse trenches
Deepenig trenches
Membrane thinning
____
3 pm
2.5
___ ___ pm
1 pm
8x20 pm
8x20 pm
__________
1x20 pm
2
2
2
6.5 nA
3 nA
_ ___ ___
700 pA
3 minutes on both side
3 minutes on both side
1 minute on both side
Stage 2:
Removing membrane from sample and attaching to TEM grid
Step
Under cut
Cut Size
1x20 pm
2
Current
1.5 nA
Duration
30 sec
Side cut I gm x sample 1.5 nA height
30 sec
Bring probe to sample
Attach probe to sample
-
3x2 gm 2
-
40 pA 5 minutes
by deposition
Side cut: To free membrane from sample
1.5 pm x sample
1.5 nA height
30 sec
Move membrane to grid
-
Attach membrane to grid 3 pm x sample 40 pA 5 minutes
by deposition height
Cut attached membrane to free it from probe
3 pm x sample height
300 pA 2 minutes
Stage 3: Final thinning to electron transparency
Step
Thinning
Thinning
Fine thinning
Fine thinning
Fine Thinning
Final Polish
Membrane Thickness Current Duration
500 nm
400 nm
300 nm
250 nm
200 nm
<-200 nm
150 pA 5 minutes
80 pA 3 minutes
40 pA 5 minutes
40 pA 5 minutes
10 pA 5 minutes
1 pA 5 minutes

Considering the challenges accompanying the FIB technique, we also employed the more traditionally used ion milling to prepare TEM samples. Since the bond layer is loosely bound on the oxygen electrode and is about 20 microns thick, it was physically lifted off the oxygen electrode using very sharp tweezers and was glued onto 3 mm outer diameter copper grids with a 1mm x 2mm slot in the middle. The grids used are shown in Figure
3-6. Once LSC material was glued to the grid, it was thinned down to electron transparency using a Gatan ion miller. The Ga ions accelerated through a potential of 5
kV were used to achieve the milling of the sample.
0
Figure 3-6: Copper grids used to prepare TEM samples
TEM s
TEM sample B etm
FIBprepared sample
Figure 3-7: Cross-sectional view of the bond layer and oxygen electrode showing where the TEM samples are prepared from. To compare the changes in microstructure and chemical composition across the bond layer cross-section, we need to prepare TEM samples using traditional milling from both the bond layer/interconnect and bond layer/oxygen electrode interfaces. On the other hand, one FIB prepared is enough for the same requirement

One FIB sample, if prepared from across the cross section of the bond layer, can make it possible to compare the TEM results across the depth of the bond layer. To be able to carry out a similar comparison of the variation in microstructure and chemical composition of the bond layer across its cross section, we used traditional milling to prepare two classes of samples one from near the bond layer/interconnect interface
(TEM sample A) and the second from near the bond layer/oxygen electrode interface
(TEM sample B). This is schematically depicted in Figure 3-7.

The following Table 4-1 defines the nomenclature for the cells, which are presented in this thesis.
Cell Number
CER#l
CER#2
Table 4-1: Nomenclature for the cells studied in this thesis
Air electrode description
As-prepared cell with no bond layer on top of the Ao.
8
Sro.
2
MnO3
(ASM*) electrode
As-prepared cell with Lao.Sro.
2
CoO
3
(LSC) bond layer on top of the
ASM electrode
CER#3 bond layer, operated at 830
0
C for 2000 hrs
CER#5
830
0
C for 2000 hrs, with bond layer removal after operation
'/2
ILS right side stack, cell# 48 - ASM with LSC bond layer, operated at 830
0
C for 2000 hrs
REF# As-prepared reference cell, ASM electrodes with LSC bond layer, operated at 8200 C for 322 hrs: Pt and Au mesh instead of stainless steel
REF#2 As-prepared reference cell, ASM electrodes with LSC bond layer, sintered at 830
0
C for 108 hrs
*The component A is proprietary information belonging to Ceramatec Inc.

4.1 Preliminary identification of secondary phases at the surface of LSC bond layer: Raman Spectroscopy
Raman spectroscopy was performed on cells CER# 1-5 on the top surface of LSC, which was in contact with the stainless steel interconnects. The results for the as-prepared cells
(CER#1 and CER#2) were used as a reference for Raman peaks while analyzing the
Raman data for cells CER#3 and CER#4. Publications from literature [36-42] were used as references for Raman peaks of the phases of interest found in our data, namely Cr
2
O
3
[36], C [37], ZrO
2
[38], LaCrO
3
[39], La
2
CrO
6
[40], LaCoO
3
[41] and Co
3
0
4
[42]. Figure
4-1 and Figure 4-2 below show the Raman Spectra for cells CER#1 and CER#2 respectively. The data in Figure 4-2 is relatively noisy and broad, however consistent with literature showing that the Raman spectroscopy on the LSC generally gives a noisy spectrum with broadening [41]. Our Raman spectrum on LSC is of better data quality than the published data in 41.
0 500 1000
Raman Shift (cm')
1500
Figure 4-1: Raman spectrum for cell CER#1 showing ASM peak

0 500 1000
Raman Shift(cm')
1500
Figure 4-2: Raman spectrum for cell CER#2 showing peaks for LaCoO
3
Figure 3-1 shows the air side of the cell CER#3 after disassembly. We see dark and light lines that appear on the surface. This distinction occurs due to the design of the corrugated flow channels to enable air flow. The LSC regions in contact with the flow channels were comparatively lighter.
We carried out Raman spectroscopy on the dark and light colored regions on the oxygen electrode side of this cell. Besides, we found that the bond layer and the oxygen electrode were peeled off at certain regions, possibly exposing the zirconia electrolyte. Raman spectroscopy at these regions clearly showed that the perovskite structure of the LSC has in fact decomposed and formed secondary phases. Presence of chromium is manifest in the form of different Cr-containing compounds. As discussed in the background of SOEC degradation, chromium comes from the stainless steel interconnects. More elaboration on the nature of Cr in LSC is provided in the discussion of results obtained using Auger
Electron Spectroscopy (AES) and Scanning Transmission Electron Microscopy (STEM).

CI
3
04 dark
light
0 500 1000
Raman Shift (cm')
1500
Figure 4-3: Raman spectra collected from a point each from the dark and light region on the oxygen side of cell CER#3, showing the presence of new phases formed after decomposition of the perovskite structure
Figure 4-3 shows Raman data from two points from the LSC surface of cell CER#3 one each from the dark and light regions. This figure makes it clear that there is no structural difference between the light and dark regions. Both have the bond layer structurally intact, even though chemically decomposed. The color contrast is not distinguishable in terms of the phases present in these regions. As marked in the Figure 4-3, LSC perovskite structure has decomposed into the Co
3
0
4
, Cr
2
0
3
, LaCrO
3 and La
2
CrO
6 phases. Previous work done on the thermodynamics of La-Cr-O system [43] has observed that LaCrO
3 and
La
2
CrO
6 are formed in 02 atmosphere at temperatures similar to the operating conditions for the SOECs under analysis in our research. Specifically, LaCrO
3 is stable till 1000
0 C in atmosphere ranging from pure 02 to p02 = 10-16.1 Pa. La
2
CrO
6 forms in oxygen atmosphere at T > 700 0
C and is stable within a wide temperature range and Po
2
> 10
5
Pa.
While the SOEC stacks are operated at atmospheric pressure, local P
0 2 in anode and bond layer of the cells can be greater than 105 Pa due to the generation and evolution of 02
[44]. Such high Po
2 favors the stability of La
2
CrO
6 as shown in Ref. 43. (The exact value of local Po
2 in the anode and LSC bond layer is not determined for this thesis). Thus, the

findings reported in Ref. 43 support that the secondary phases of lanthanum chromite and chromate formed were stable at the operating conditions of SOECs. This decomposition of LSC perovskite structure leads to an adverse change in the electronic properties of the bond layer. Thus, the decomposition of the LSC perovskite into secondary, low conducting and less active phases can be a chief cause for the degradation.
m - monoclinic m-ZrOm m-ZrO, m-ZrO, unidentified peaks
0 500 1000
Raman Shift (cm'')
1500
Figure 4-4: Raman spectrum collected from region with the exposed electrolyte on the oxygen side of cell CER#3, showing only the zirconia peaks below 1000cm
1
.
Here the prefix 'im' stands for monoclinic
Figure 4-4 shows Raman results for a point where the electrode and the bond layer had
peeled off. Clearly, the electrolyte has been exposed here. The first few peaks correspond to the zirconia electrolyte. The peaks above 1000 cm
1 were not found in the literature as reference. The new secondary phases formed are all in the LSC but we do not see them here because those layers have peeled off.
Compared to the original compositions, the secondary phases that were identified have a much lower conductivity and hence lead to a loss in performance of the cells. Table 4-2 below compares the conductivity values for the various phases identified here.

Table 4-2: Conductivity data for the chemical species that were found in the Raman spectra of tested SOEC oxygen electrodes. Conductivity values are obtained from literature and from Ceramatec Inc.
Constituent
Lao.6Sro.
8 mol % Sc
2
4
CoO
LaCoO
O
3
3
-
3 at 8000 C at 900) C
ZrO
2 at 8000 C
Conductivity (S/cm)
1.6 x10
3 [45]
~ 7.6x10
2
[46]
~ 4.78 x10-
2
[47]
~ 3.9x10
1
[48] C0
3
0
4 at 800( C
Cr
2
O
3 at 10000 C
LaCrO
3 at 800u C
1.0 x
[49]
3.4 xlO-' [50]
4.1.1 Summary
The important results obtained from Raman Spectroscopy are summarized below:
* The LSC bond layer clearly decomposed and shows the presence of chromium containing Co
3
0
4
, Cr
2
O
3
, LaCrO
3 and La
2
CrO phases. All the secondary phases have much lower electronic conductivities than the initial phase (as seen in Table
4-2), and thus, hamper the cell performance.
" The presence of Cr in the bond layer is a clear indication of Cr transport from the steel interconnects. This is likely a major cause for the dissociation of the active bond layer and, hence, the degradation of the cell, as is discussed in more details in sections 4.2 and 4.3
" Monoclinic zirconia was observed at electrolyte surface.
4.2 Bond layer and oxygen electrode microstructure and surface composition:
Nanoprobe Auger Spectroscopy
Upon confirming the presence of the less conducting secondary phases by Raman spectroscopy, nanoscale scanning Auger Electron Spectroscopy (NAES) was used for identifying the variation in the cation distributions over the cross section of the bond layer and the oxygen electrode particularly on their surfaces.

Several pieces were broken off from cell CER#2 and were subjected to cross sectional
Auger spectroscopy. Different spots in the ASM and the LSC region were scanned.
Figure 4-5 shows an SEM image of the cross section, including the oxygen electrode and bond layer. Since the LSC layer was not sintered (only as-applied and dried), it looks very dense due to the presence of the binder and is not conducting.
Oxygen Oxygen
Bond electrode: electrode: layer layer 1 layer 2 Electrolyte
LSC ASMI ASMI/ScSZ
2000 X 00P
Figure 4-5: SEM image of the cross section of cell CER#2
Oxygen Oxygen
Bond electrode: electrode: layer layer 1 layer 2 Electrolyte
LISC
ASM
Figure 4-6: SEM image of the cross section of cell CER#2, showing densified regions in ASM

We present the AES results in terms of the material under investigation, i.e. oxygen electrode (ASM region) and bond layer (LSC region). The following sub-sections discuss the key results for both in detail.
4.2.1 ASM oxygen electrode
A fresh as-prepared ASM sample (without any LSC bond layer), i.e. cell CER#1, was broken and subjected to cross sectional Auger spectroscopy. The AES results obtained from CER# 1 serve as reference ASM AES spectra. Figure 4-7 below shows an SEM image of the region analyzed in the cross-section of the electrode. The microstructure has rounded grains of 0.5-1 pm size.
20000 X
Figure 4-7: SEM image of a region in ASM layer of cell CER#1, showing uniform rounded grains
Figure 4-8 is a SEM image of the ASM region of cell CER#3, the tested cell. Clearly, the
ASM microstructure has changed from rounded grains in CER#1 to faceted grains in
CER#3 after the long-term high-temperature electrolytic conditions.

10000 X 10.0 keV0
Figure 4-8: SEM of ASM region from cell CER#3 showing faceted grains with sharp boundaries
AES spectrum for the ASM region of this sample was compared with that of cell CER# 1.
Both CER# 1 and CER#3 surface compositions are similar at the measurement conditions of UHV and room temperature, thus ASM seems has stayed stable. Due to the confidentiality agreement with Ceramatec Inc., we cannot present the AES spectrum from the ASM region here. Table 4-5 below compares the composition of cell CER#1 and cell CER#3, normalized with respect to the sum of compositions of A, Sr and Mn.
From Table 4-5, it is clear that the composition of the ASM region stays practically unchanged during operation and it does not disintegrate. Auger spectroscopy did not show the presence of chromium in the ASM layer.
Table 4-3: Comparison of ASM compositions normalized with respect to the sum of
A, Sr and Mn cations, before (cell CER#1) and after use (cell CER#3)
Element
Cell CER#I
Cell CER#3
A
0.63
0.60
Sr
0.15
0.19
Mn
0.21
0.21
(A+Sr)/Mn
3.71
3.76

Oxygen Oxygen
Bond electrodeelectrode: layer layer layer 2 Electrol e
Delamination.
ASM;
LSC
20pm
Figure 4-9: SEM image of cross-section of CER#3 showing delamination between
LSC and ASM layers
The ASM layer in cell CER#5 was also studied by AES. It had stayed stable and showed no presence of Cr. We noticed delamination between the ASM electrode and the LSC bond layer, shown in Figure 4-9. The gap between ASM and LSC was roughly 1-2 microns and was present throughout the length of the cells CER#3 as well as CER#5. The delamination can be attributed to high current densities of operation: High current density leads to a higher evolution rate of gaseous oxygen in ASM electrodes, which can exert pressure on LSC. This can lead to LSC delaminating from ASM. Whether the delamination occurs over the course of cell operation is, yet, unknown. It is likely that the delamination between electrode and bond layers prevented Cr-containing species to diffuse into the electrode and limited them to the LSC region. The exact role that the delamination plays, if any, in the stability of ASM electrode is still to be understood clearly and needs further investigations.
To facilitate our analysis of comparing the variations in chemical composition across the
ASM cross section, ASM was divided into three regions of investigation a region close to the LSC/ASM interface (region 1), a middle region (region 2) and a region closest to

the electrolyte (region 3). Table 4-4 summarizes the compositions of areas scanned by
AES, normalized with respect to the sum of compositions of A, Sr and Mn, in these different regions for both cells CER#3 and CER#5. As seen here, the local surface compositions probed by the nanoscale-scanning Auger spectroscopy shows that the ASM surface composition did not vary significantly, and stayed comparably similar upon longterm testing.
Thus, we see that the A/Mn ratio on the surface of both of the cells in different regions stays fairly constant, though it is different from the bulk average of 0.8. Also, the ratio of
(A+Sr)/Mn for cells CER#3 and CER#5 is different, and on average larger than the corresponding ratio on the surfaces of as-prepared cells (CER#1). This implies a further segregation of the A-site cations to the surface upon electrolytic conditions.
Table 4-4: Variation of cation percentages, normalized with respect to the sum of A,
Sr and Mn cations, over the different regions of ASM for cells CER#3 and CER#5.
For as-prepared ASM, the average fractions of A, Sr and Mn were seen to be 0.63,
0.15, and 0.21, respectively. This corresponds to a (A+Sr)/Mn ratio of 3.71, different from the results from the surfaces of the tested cells CER#3 and CER#5.
Region 1
Region 2
Region 3
Cell CER#3 Cell CER#5
A Sr Mn (A+Sr)/Mn A Sr Mn (A+Sr)/Mn
0.53- 0.26- 0.09- 6.07- 0.48- 0.27- 0.20- 3.80-
0.59 0.37 0.15 9.74 0.55 0.33 0.22 4.03
0.53- 0.26- 0.13-
0.57 0.33 0.17
4.73- 0.48- 0.30- 0.17- 4.24-
6.08 0.51 0.34 0.19 4.87
0.54- 0.23- 0.07- 5.22- 0.54- 0.14- 0.18- 2.97-
0.60 0.36 0.14 12.66 0.58 0.44 0.26 5.22

4.2.2 LSC bond layer
After the detailed investigation of ASM layer using NAES, the next objective in our research was to identify surface compositions on LSC via a thorough AES examination.
This section is devoted to the discussion of key AES results obtained from LSC.
4.2.2.1 Cr presence in LSC and cation segregation at electro-catalyst surface
We probed the LSC region of cell CER#3. Figure 4-10, a SEM image, shows the crosssectional microstructure of the LSC region for cell CER#3. Figure 4-11 is the AES spectra of the different areas in this region.
10000 X 10.0 keV 2
Figure 4-10: SEM image of the cross section of cell CER#3 (LSC region)

t
Area I
-- Area 2
Area 3
U
400
04
0
-
Cr Element 0
La o C'-Co
Cr La Co Crl(La+Co+Cr)
Area 1 0.70 0.08 0.08 0.13
Area 2 0.57 0.01 0.38 0.04
Area 3 0.61 0.02 0.29 0.07
0.27
0.02
0.05
500 600
Kinetic Energy (eV)
700 800
Figure 4-11: AES data from points 1, 2 and 3 in Figure 4-10
The presence of La, Co, 0 and Cr is clearly manifest on the surfaces of this region. From
Figure 4-11, three important observations are drawn:
1) It is clear that the three different points have significant differences in their local compositions and chemical signature even though they are separated by only a few of microns from each other. The ratio of La to Co was found to vary roughly from a mere 0.67 to 16.37. Ideally, this ratio of the as prepared LSC should be 0.8.
2) 'Crystallite' like structures, differing from the LSC microstructure, were found in the bond layer. One such crystallite is visible in the SEM figure above (area 1). A noteworthy observation is that these crystallites show a significantly higher Cr content than the rest of the neighboring microstructure, as is visible from the table inset in the spectra in Figure 4-11.
3) Sr-signal, intriguingly, lacked in the AES spectra from the LSC cross section. Since
Auger spectroscopy is a surface sensitive technique (top 1-2 monolayers on the surface), these observations lead us to hypothesize that the surface composition of the bond layer changed such that:
.................

* The ratio of La/Co at the surface deviated from 0.8 and assumed drastically different values at different grains due to local decompositions.
" Sr depleted on the surface, and segregated to the inner layers of the grains, its signal thus getting blocked, or had migrated elsewhere in the cell.
In region 2 of cell CER#3, we found Cr content as high as 34%. This sample was selected from the air-steam inlet region at which a higher current density is expected. However, this observation was not possible to reproduce for the same specimen, using AES, due to the deterioration of the sample upon wait-time, and hence remained inconclusive. Figure
4-12 shows the SEM image of this specific region and it is followed by the corresponding
AES spectra. AES spectra clearly indicate the different chemical signature from the base cations of La and Co, indicating the surface composition variations within small distances on the bond layer surface. As discussed in results from STEM/EDX, this region may be from the bulk composition of a grain exposed during breaking of the sample, rather than the surface.
Figure 4-12: SEM image of region 2 in cell CER#3

-
Region I
Region 2
Region 3
0
0
Coj Co
0 Co
CrL
Element 0 Cr La Co
Ame 1 0.40 0.32 0.25 0.02
Area 2 0.43 0.34 0.15 0.04
Area 3 0.39 0.31 0.27 0.07
Cr/(La+Co+Cr)
0.53
0.59
0.51
0
400 500 600
Kinetic Energy (eV)
700 800
Figure 4-13: AES spectra from the areas in Figure 4-12, showing Cr fraction as high as 0.34
Cell CER#5 was also subject to Auger spectroscopy. It showed similar results as cell
CER#3. Presence of chromium was seen in the various regions of LSC. Formation of crystallite-like structures with higher chromium content and the local variation in the fraction of cations was observed, similar to those in CER#3
4.2.2.2 Cation variation across the LSC cross-section
To facilitate our investigation of the variation of composition of different cations across the LSC cross-section, we divided the analysis of LSC into three different regions a region close to the surface of the bond layer (region 1), a middle region (region 2) and a region close to the ASM/LSC interface (region 3). Table 4-5 summarizes the compositions found by AES in these different regions, normalized with respect to the sum of La, Sr, Co and Cr cations, for both CER#3 and CER#5. The variation of normalized average* fraction of cations in each region is also shown in Figure 4-14 and
Figure 4-15. For both cells, CER#3 and CER#5, as one traverses into the bond layer from
* This section discusses the variations in cation fractions averaged over each region. The exact cation fractions for all data points in the three regions are discussed in Appendix B.

the LSC/air interface, Cr content monotonically decreased. Co-content monotonically increased for CER#3 while it decreased monotonically for CER#5. La-content first increases and then decreases for CER#3 and monotonically increases for CER#5. While there are inconsistencies between the two cells in terms of La- and Co-distribution, and a large range of measured values, the highest Cr-content was seen near the Cr-source, i.e.
interconnects for both cells. The variations in Co- and La-fractions across LSC crosssection, on the other hand, do not follow a particular trend, as more explicitly shown in
Appendix B.
Table 4-5: Variation of cation fractions, normalized with respect to the sum of La,
Sr, Co and Cr compositions, over the different regions of LSC for cells CER#3 and
CER#5
Cell CER#3 Cell CER#5
Cr La Co La/Co Cr La Co La/Co
Region 0.04- 0.22- 0.05- 1.17- 0.03- 0.72- 0.09- 2.99-
AES data, 1 0.50 0.91 0.42 16.37 0.09 0.87 0.24 9.58
except for Region 0.02- 0.33- .05- 0.67- 0.02- 0.74- 0.06- 3.45region 2 of
CER#3
2 0.18 0.91 .50 14.57 0.04 0.92 0.21 14.26
Region 0.01- 0.28- 0.06- 0.62- 0.01- 0.84- 0.08- 6.19-
3 0.27 0.90 0.56 12.42 0.02 0.90 0.14 10.24
Region with very high Cr cen (inh CrRegion 0.75- 0.38- 0.06- 1.97-
2 0.85 0.68 0.18 11.21
region 2 of
CER#3)

..... .....
40
30
20
10
0
0
100
90
80
70
60
50
1
Region
2 3
Figure 4-14: Variation of the averaged cation content, normalized with respect to the sum of La, Sr, Co and Cr, in the bond layer as a function of cross-sectional region for Cell CER#3
100
90
80
70
60
-
50
-
40 -
30
-
20
-
10
-
0
+0Cr
-I
T
-I
Region
Figure 4-15: Variation of the averaged cation content, normalized with respect to the sum of La, Sr, Co and Cr, in the bond layer as a function of cross-sectional region for Cell CER#5

4.2.2.3 Migration of Sr and Co to LSC/interconnect interface
Any signal belonging to Sr in the AES spectra was not found in LSC. Intrigued by the lack of Sr on the surface of the cross section of the bond layer, and recalling that Sr was previously found to migrate from the bond layer to near the interconnect [51], we carried out AES analysis at the surface of the LSC bond layer, i.e. at the LSC/interconnect interface of cell CER#5. We obtained two interesting results by this analysis.
1. SEM imaging at the LSC surface showed the presence of finely dispersed
'crystallites', about 3-5 microns in size (as can be clearly seen in Figure 4-16 below).
These crystallites were loosely bound to the LSC surface. These crystallites were much richer in Co than in the LSC cross section. The AES spectrum from one of these crystallites is shown in Figure 4-17 below. In addition, these crystallites were found to lack La, whereas the LSC cross-section had approximately 70 at. % La
(normalized with respect to the sum of La, Co and Cr compositions).
Co-rich crystallites
Sr-rich region, free of Co-rich crystallites
Figure 4-16: SEM image of the LSC surface of cell CER#5, showing numerous crystallites. NAES analysis showed these crystallites to be rich in Co

-
surface crystallite
LSC interior surface
Cr La -Co
400
SI
500
Element 0
Surface 0.67
In*Mio. A5Q
1
600
Kinetic Energy (eV)
La
-
032
1
700
Co
-*Co
Sr Co C
0.02 0.29 0.01
l04j
800
Figure 4-17: AES spectrum from one of the Co-rich crystallites on the surface shown in Figure 4-16 compared with that from the LSC interior cross-section
-
-- surface crystallite-free region
LSC interior surface
Sr
1600
Ii
1620
I
1640
Sr
Element 0 La Sr Co Cr
Surface 0.48 0.07 0.41 0.01 0.02
Iterior 5
I I
1660 1680
Kinetic Energy (eV)
032nA
I
1700
-
I
1720
0
.
M I
1740
Figure 4-18: AES spectrum from the region free of Co-rich crystallites on the surface shown in Figure 4-16 compared with that from the LSC interior crosssection
2. On the surface regions free of Co-rich crystallites shown in Figure 4-16, we measured a considerably stronger Sr content in the AES spectra as compared to that from the

LSC interior cross section. This comparison is presented in Figure 4-18. While some regions on this Co-crystallite-free surface were extremely rich in Sr, containing as high as 41 at. % Sr, other regions had comparatively lower Sr-content (~12 at. %). In comparison, no Sr-content was observed on the LSC interior surface.
These findings, coupled with the lack of Sr-signal from the LSC cross-section suggest that Sr and Co cations have migrated from the LSC interior, towards the
LSC/interconnect surface, over a distance of tens of microns. This is an interesting observation and further hints at the drastic dissociation of LSC, which can no longer be assumed to remain in its initial perovskite phase. The direction of transport of all ions and electrons in the cell are summarized in the Figure 4-19 below. This migration could be due to two primary reasons: a. The presence of Cr-containing species near the interconnect could thermodynamically drive Sr and Co cations to the top of bond layer.
b. Electronic current from the oxygen electrode towards the interconnects, as shown in
Figure 4-19, could kinetically demix and migrate (physically transport) Sr and Co towards the LSC/interconnect interface.
Sr
0
Co e Cr
LSC
Figure 4-19: Schematic representation of the movement of various species across the cell cross-section

4.2.2.4 Cation Migration Mechanism: Experiments on reference cells in controlled environment
It is important to isolate the mechanism for the long-range migration of Sr and Co as shown schematically in Figure 4-19. Our hypotheses are that the long-range cation migration is driven by: 1) Cr-related thermodynamics or 2) Electronic current, discussed in the previous section. To test these hypotheses, we carried out experiments under controlled electrochemical conditions, on half-cells without any Cr-containing layers on the electrodes. The diagram of the cells is shown in Figure 3-2. These cells are comprised of ASM as the electrode and LSC as the bond layer on both sides of the ScSZ electrolyte.
To investigate the cell performance in the absence of Cr, we replaced steel interconnects used in cells CER#3-5 by Au and Pt meshes in our reference cell test set-up. In order to replicate the actual operating conditions, REF#1 was tested in air, under a constant current density of 0.40A/cm
2
(same as the starting current for CER#3-5) at 816
0
C for
13.5 continuous days. To compare the effects of electrochemistry to the effect of high temperature alone on the reference cells, an additional reference sample was annealed,
REF#2, at 8200C in air for 4.5 days continuously. In the electrochemical set-up, the working electrode ran in electrolytic mode (anode) while the counter electrode ran in fuel cell mode (cathode).
Figure 4-20 shows the potential difference across REF#1, from working electrode to counter electrode, under constant current density of 0.40 A/cm 2 , as a function of time.
The potential difference increased from an initial value of 0.23V and stabilizes at 0.42V
after 4.5 days of operation. This increase indicates that REF#1 degraded, even though there is no Cr in the system.

0.45
0.4 -
"7
0.35
0.3 -
0.25 -
0 .2
-
0 50
I
100
I I
150 200
Time (hours)
I
250
I
300
I
350
Figure 4-20: Cell performance for REF#1: Under a constant current density of
0.4A/cm 2 at 820
0
C in air, the potential difference across the cell increased with time and stabilized after 108 hours of operation
Individual anode and cathode resistances in oxygen reactions of REF#1 were identified, using Electrochemical Impedance Spectroscopy (EIS), before and after the long-term test.
Figure 4-21 and Figure 4-22 are Nyquist plots for the working electrode (electrolytic mode, anodic) and counter electrode (fuel cell mode, cathodic), respectively. The impedance of the anode increased by 0.76 9, from 0.07 i to 0.83 f, while that of cathode increased by 0.42 L, from 0.07 9 to 0.49 K. Thus, the anode, operating in the electrolytic mode degraded more than the cathode, even though their compositions are the same.

+
Before Operation * After Operation
0.3
-
100Hz .
.
0.2 -
0.1
a
100.) Hz
*m 10Hz in
10 Hz7
1U
Hz
0.2
0.4
Z_real(ohm)
0.6
0.8
Figure 4-21: Nyquist plots for anode (oxygen evolution in electrolytic mode) of
REF#1, before and after its operation
* Before Operation 0 After Operation
0.3 i
S0.2 -
0.1
N
100 Hz
10 Hz
+
0 z
1O 0
0.2
0.4
Z_real (ohm)
0.6
0.8
Figure 4-22: Nyquist plots for cathode (oxygen reduction in fuel cell mode) of
REF#1, before and after its operation

Figure 4-23: SEM image of LSC surface of cell REF #1 showing that it does not have secondary crystallite formation
Oxygen- Oxygen|
Bondelectrodgelectrodd: layer: layer 1: layer 2 :Electrolyte
Figure 4-24: SEM image of REF#1 cross-section
Comparing the changes in microstructure and microchemistry of the reference half-cells to those in the full-scale cells obtained from Ceramatec Inc. can shed light on the reasons for cation migration observed in CER#3 and CER#5. We analyzed both the LSC surface as well as the cross-section. There were no crystallites, containing Co or Sr, found at the surface of the REF#l LSC as shown in Figure 4-23, unlike that in CER#5 (shown in
Figure 4-16). Figure 4-24 shows the cross-section of REF#1, there is no delamination between LSC and ASM layers, which was seen in CER#5, even though both REF#1 and

.......
CER#5 operate under same current density. Thus, we conclude that high current density of operation is not the sole mechanism that can cause delamination of LSC from ASM electrode. The precise cause for delamination remains unclear. We divided the analysis of reference half-cells into five sub-objectives:
1. Comparison of the chemical compositions of LSC in REF#l and CER#5 can help in judging whether they are different in terms of dissociation. Figure 4-25 compares the
AES spectra from the cross section of REF#1 to that of the actual full-scale cell tested at Ceramatec Inc. (CER#5). The peak corresponding to Cr is evident in the spectrum from CER#5 and is missing in that from REF#1, which is expected as CER#5 had steel interconnects. On the other hand, REF#1 had Au and Pt meshes as the current collectors, which removed the role of Cr and its species in the cell.
----
-
REF#
CER#5
Cr 'La Co Co Co
0
500
O0
Element 0 L Sr Co Cr (La+Sr)/Co
REF#1 0.48 0.31 0.16 0.07 6.71
CER#5 0.52 10.38 0.04 0.04 9.50
400 600
Kinetic Energy (eV)
700 800
Figure 4-25: Comparison of AES spectra from LSC cross-section of cell REF#1 against that of cell CER#5
.

Figure 4-26: SEM image of LSC surface of cell REF #1
-
Area I
Area 2
Area 3
O
4'
La COI CO CO
0
Element O La Sr CO (La+Sr)/Co
Area 1 0.52 0.26 0.16 0.07 6.00
Area 2 0.53 0.27 0.13 0.07
Area 3 0.51 0.27 0.15 0.07
5.71
6,00
400 500 600
Kinetic Energy (eV)
700 800
Figure 4-27: AES spectra from the 3 areas shown in Figure 4-26. The three areas probed have similar chemical composition, showing that REF#1 has not experienced non-uniform segregation of cations at surface of LSC grains (unlike CER#3 and
CER#5)
2. We investigated whether the LSC/interconnect interface of REF#1 was subject to Sr and Co enrichment upon electrochemical treatment with Au and Pt interconnects. For this purpose, we performed AES analysis on the LSC surface. Figure 4-26 is a SEM image of the region probed by AES followed by the AES spectra in Figure 4-27 from

the three areas shown. All three areas showed uniformity in composition 53 at.% 0,
26 at.% La, ~15 at.% Sr and 7 at.% Co.
3. The microstructure and chemical composition of LSC cross-section in CER#3 and
CER#5 were different from the LSC/interconnect interface. Investigation of the microstructure and composition of reference half-cells, with Au and Pt meshes, can help in understanding whether variations in chemical composition across LSC crosssection are related to the presence of Cr. Figure 4-28 compares the AES spectra for an area on the LSC surface to the LSC cross-section, for REF#1. Results showed that
REF#1 had a uniform microstructure and chemical composition throughout, both at its surface as well as across the cross-section, upon the long-term test. As seen from the table insets in Figure 4-25, Figure 4-27 and Figure 4-29, the ratio of A-site to Bsite cations for REF#1 varied from 5.71 to 6.71. This variation is small in comparison to the drastic variations from 0.67 to 16.37, observed for CER#5. Thus, Cr-driven thermodynamics had a more significant impact on non-uniform dissociation of LSC bond layer and the long-range segregation of cations at the surface of LSC. Electronic current on the other hand was not found to be as important a factor as Cr-deposition.
4. In order to compare the effects of just the heat treatment without electrochemistry
(REF#2) to that of heating with electrochemistry (REF#1), we assess the surface compositions of REF#1 and REF#2. This evaluation also helped us to identify whether any changes occur in the microstructure of reference half-cells over the course of their operation. Figure 4-29 shows AES spectra from these two cells. It is seen that both REF#1 and REF#2 have similar chemical compositions, with ratio of
A-site to B-site cations varying from 5.71 to 6.

---
-
Surface
Cross-section
0
I
La CO CO CO
Elmnt 0 La Sr CO (La+Sr)ICO
Surface 0.52 0.26 0.16 0.07 6.00
Cross- 0.48 0.31 0.14 0.07 6.42
0section
I
400 500 600
Kinetic Energy (eV)
700 800
Figure 4-28: Comparison of AES spectra from LSC surface to cross-section, for cell REF#1. The composition at LSC surface is uniform throughout, both at the surface and across the cross-section
---- REF#1
-- REF#2
La CO Co CO
0
O0
Element 0 La Sr Co (La+SryCo REF 0.52 0.26 0.16 0.07 6.00 n 4 n 'r. A IA 571 neum"
400 500 600
Kinetic Energy (eV)
700 800
Figure 4-29: Comparison of AES spectra from LSC surface of cell REF#1 to REF#2

5. The working electrode (anode) and counter electrode (cathode) of REF#1 show unequal increase in resistance over the period of operation. We compared their surface chemical compositions to identify whether the difference in the extent of their degradation can be attributed to dissimilar microstructures. Figure 4-30 shows the comparison of AES plots from LSC surface of both electrodes. The anode degraded almost twice as much as the cathode, both electrodes show little difference in the A- to B-site cation ratio.
These results suggest that there was no significant difference in the microstructure and chemical composition of the LSC region of REF#1 and REF#2 upon electrochemical and thermal treatments, respectively. The cross-sectional (interior) surface of REF#1 had similar composition as its top surface. Thus, REF#1 and REF#2 were stable throughout the LSC region. There was no significant difference in the chemical compositions of the just heat treated cells (REF#2) and the cells that were operated under a constant current density (REF#1). The A-site to B-site cation ratio for the anode of REF#1 and REF#2 varied from 5.71 to 6.71. On the other hand, for CER#5, the dissociation of LSC across its cross-section was drastically different from that at its surface, with Sr and Co cations migrating to the surface. The A-site to B-site cation ratio for CER#5 varied from 0.67 to
16.37. Thus, cation segregation in CER#5 is more severe than in REF#1 and REF#2.
While the same current density for REF#I was used as the initial current density for
CER#5, an uneven segregation of cations at the surface of the grains in REF#1 was not observed. No secondary phases rich in Sr or Co were present at the surface of REF#l.
The resistance of the reference cell was seen to increase with time (before stabilizing), indicating degradation even in the absence of Cr. The fact that REF#l and CER#5 had the same electrode and bond layer compositions, except for the absence of Cr-containing species in REF#1, indicates that the deposition of Cr-containing species in LSC plays a crucial role in the non-uniform dissociation of the LSC phase and the migration of select cations (Sr and Co in this case) to the LSC surface. It must be pointed out here that the exact mechanism by which Cr and its species cause this segregation and the dissociation of LSC remains yet to be clearly understood.

---- Anode
Cathode o
IO ement 0 La Sr Co (La+Sr)/Co
6.42
Anode 0.48 0.31 0.14 0.07
8.00
Cathode 0.46 0.33 0.15 0.06
400 500 600
Kinetic Energy (eV)
700 800
Figure 4-30: Comparison of AES spectra from LSC layer as anode vs cathode for
REF#1. Even though anode and cathode degrade unequally, they show similar microstructure and surface composition.
4.2.3 Summary
Major results obtained from Auger Spectroscopy in this work are summarized below: e
The ASM electrode stayed relatively stable throughout the operation. There was no Cr found in the ASM region. This could be attributed to the observed delamination between ASM and LSC, which could prevent the diffusion of Crcontaining species into the electrode, keeping it stable. The results for tested ASM were compared to those for the as-prepared ASM. While the microstructure varied, the surface composition was very similar, with slight increase in the A-site segregation in the tested cells. For both tested and as-prepared ASM samples, the
A/Mn ratio varied approximately between 2.5 and 3.5 compared to the expected bulk average of 0.8. Thus, A-site segregation of cations (A and Sr) on the ASM grains existed. However, the A/Mn ratio stays fairly stable throughout the ASM region and only small local variations were observed.

* There was clear presence of chromium in LSC bond layer. The Cr content
(fraction) varied along the depth of the bond layer. The average chromium fraction observed was approximately 0.07 (normalized with respect to sum of La,
Co and Cr cations). A general decreasing trend was observed from the air/bond layer interface towards the ASM electrode. In one particular case, normalized Cr fraction as high as 0.53 existed, midway in the LSC cross-section. Cr-containing crystallite-like structures were found and had the highest amount of Cr content in their respective regions. Surface composition locally varied in the LSC grains.
The ratio of La/Co differed largely even in the neighboring grains within a few microns. This ratio, which should ideally be 0.8, varied from 0.67 to 16.37. In conclusion, the bond layer dissociated drastically and it is highly likely that the
LSC no longer remained in its perovskite phase and disintegrated into secondary phases.
* AES signal did not show Sr on the surface of grains in the LSC cross-section.
This suggests depletion of Sr on the LSC grains occurring at the operating conditions. AES on LSC, at the interface between LSC/interconnects, showed that the surface had numerous crystallites that were much richer in Co compared to the
LSC cross section. The Co-rich crystallites were observed to be deficient in La.
Furthermore, this interface was significantly richer in Sr compared to the crosssectional surface of LSC. Sr was found unevenly distributed at the
LSC/interconnect surface, with Sr-content ranging from 12 at. % to 41 at. % .
On the other hand, the LSC cross section did not show presence of Sr. These results suggest that Sr and Co cations migrated to the surface of the LSC bond layer from the cross section. This observation also confirms the severe dissociation of the bond layer and suggests that LSC no longer remained in the initial perovskite phase.
* To isolate the mechanism underlying the dissociation of LSC, accompanied with the migration of Sr and Co to the LSC surface from its interiors, we replicated the electrochemical operating conditions on reference half-cells with ASM electrodes

coated with LSC layer on both side of ScSZ electrolytes, in the absence of Cr.
The current collectors in this experiment were Pt and Au meshes instead of steel interconnects. The reference cell degraded with time, even in the absence of Cr.
The anode, working in electrolytic mode, degraded almost twice as much as the cathode. Both electrodes showed similar microstructures and compositions at the end of the test, indicating that degradation occurred even in the absence of severe cation segregation. The analysis of the reference cells after the test showed that the microstructure and chemical composition of LSC bond layer stayed the same before and after the operation, with 53 at.% 0, 26 at.% La, ~15 at.% Sr and 7 at.%
Co at its cross-section and its surface. On the other hand, the LSC microstructure of cells CER#3-5, operated at Ceramatec Inc, varied drastically across its crosssection. No delamination between ASM and LSC layers was observed for REF# 1.
These observations indicate that the electronic or ionic current alone is not the dominating factor in Sr and Co migration. Presence of Cr in the bond layer is likely a major cause for the non-uniform dissociation of LSC bond layer, although the cells degraded even in the absence of Cr.
4.3 Cation distribution in the bond layer at a nano-scale: TEMIEDX
To identify the length scale at which the structural and chemical evolution took place in the grains and at the grain boundaries of LSC, TEM analysis was carried out on CER#5.
Since AES is a surface sensitive technique, TEM complements the AES results by providing information about the bulk chemistry of samples. AES results showed that the
ASM oxygen electrode surface stayed stable. Hence, we limited the TEM/EDX studies to only the LSC bond layer. In this section, we present results from two types of samples prepared by traditional ion milling method: one represents the region in LSC near the interconnect interface, and one represents the region in LSC near the ASM anode interface. The nomenclature for the TEM samples is presented in Table 4-6. The schematic of the regions that TEM samples are prepared from is shown in Figure 3-7.

Table 4-6: Nomenclature for TEM samples prepared by traditional ion milling
Sample Name
TEM sample A
TEM sample B
LSC region
LSC/Interconnect interface
LSC/ASM interface
4.3.1 Elemental distribution in the bond layer
The chemical dissociation of the LSC bond layer at a scale of a few micrometers to nanometers was studied by carrying out STEM elemental mapping measurements on different regions in both TEM samples A and B. Figure 4-31 and Figure 4-32 show dark field TEM images of different regions of sample A, and the elemental maps for La, Sr,
Co and Cr, at different length scales. Figure 4-33 shows the elemental maps collected from a region of sample B.

(b) (c)
(d) (e)
Figure 4-31: (a) A dark field TEM image of a region (2.5p~m x 1.7jpm) of the TEM sample A with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr

600nm
(b) (c)
(d) (e)
Figure 4-32: (a) A dark field TEM image of a region (1.3pm x 0.9pm) of the TEM sample A with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr

(b)
(c)
(d) (e)
Figure 4-33: (a) A dark field TEM image of a region (1.3pm x 0.9pm) of the TEM sample B with the elemental maps for (b) La, (c) Sr, (d) Co and (e) Cr
4.3.2 Possible secondary phases in the bond layer
It is evident from Figure 4-31, Figure 4-32 and Figure 4-33 that La, Sr, Co and Cr cations distributed non-uniformly in the bulk of LSC bond layer, indicating dissociation of bond layer bulk structure at the nano-scale. There are three noteworthy observations here.

1. Phases which have high Cr- and La-content and a lower Co-content were present, as indicated in Figure 4-31 and Figure 4-32. Similarly, the region marked in Figure 4-33 had high Co- and lower Cr- and La-content. The La- and Cr-rich phases measured approximately 1.5 4m x 1.5 4m. The Co-rich phases were smaller and were approximately 1 pm in length. Table 4-7 summarizes the composition in bulk and at the surface of LSC grains. The average Cr content at the surface of LSC grains, obtained by AES was 7%, whereas in the LSC bulk, it varied from a minimum of 10 at. % to a maximum of 33 at. %. Thus, the average Cr-content at the surface of LSC grains was lower than in LSC bulk. This contrast implies that the mechanism for the change in bulk structure of LSC grains reflected differently on their surface.
Table 4-7: Comparison of chemical composition in the LSC bulk and at LSC surface. It is observed that in LSC bulk, Cr-rich regions are deficient in Co, and vice-versa. Cr-content in the LSC bulk is more than at LSC surface
Atomic %
LSC Surface (AES) LSC Bulk (TEM)
TEM Sample A
Element (Near the interconnect)
TEM Sample B
(Near the anode) Cr-rich Co-rich region region
La
Sr
Co
Cr
Co/Cr
La/Cr
Cr-rich Co-rich Cr-rich Co-rich region region region region
54.92
2.81
22.45 50.00 23.08
15.49
27.03
0.57
2.03
2.85
63.63
10.20
6.24
2.20
4.16
12.50
33.33
0.38
1.50
3.84
57.60
15.38
3.74
1.50
28.01
--
44.62
27.37
1.63
1.02
37.03
--
56.53
6.44
8.78
5.75
2. The La/Cr ratio in the LSC bulk was -2 for sample A, indicating the formation of
La
2
CrO
6
. For sample B, La/Cr ratio was 1.50. The drop in La/Cr ratio from near the interconnect to near the anode suggests formation of different chemical phases. An

increase in the Cr content (relative to La) in sample B suggests the presence of
LaCrO
3 and/or Cr
2
O
3 near the anode. Co-containing phase is likely to be C0
3
0
4
. This is consistent with the Raman Spectroscopy results indicating their presence at the top surface of the bond layer upon long-term operation.
3. The Sr signal in STEM measurements on TEM samples was weak. Thus, we conclude that Sr was missing in the bulk of LSC. From Table 4-7, the highest Srcontent was only 4.16% in the bulk of LSC. This observation conforms to the surface
AES results indicating that Sr completely left the bulk LSC structure and segregated in another phase on top of the LSC at the interconnect interface.
4.3.3 Summary
We prepared TEM samples using traditional ion milling, from both the LSC/interconnect interface region and the LSC/ASM interface region. We carried out STEM elemental mapping measurement and point EDX analysis in both batches of samples. The following points summarize the key results obtained from STEM/EDX on various samples.
" Cr-content in the bulk of LSC varied from 10 33 at. %. At the surface of the
LSC, the average Cr-content was 7 at. %. Thus, the mechanism of reaction with
Cr in the LSC bulk is differently reflected at its surface. Phases that were rich in
Cr and La but low in Co content, and vice-versa, were found in the bulk.
* Cr-rich phases measured approximately 1.5 pm x 1.5 pm while the Co-rich phases were smaller in size (approximately 1pm in length).
* La/Cr ratio in the LSC bulk varied from near the interconnect to near the anode, suggesting the formation of different chemical phases in the LSC bulk. La/Cr and
Co/Cr ratios suggest formation of LaCrO
3
, La
2
CrO
6
, Cr
2
O
3 and Co
3
0
4
. Hence the observations by Raman spectroscopy and TEM concur with each other.

* A maximum of 4.16 at. % Sr was found in TEM samples from the bulk of LSC layer. This result is similar to Sr-depletion at surface of LSC grains, discussed in
Section 4.2.2.1. This suggests that Sr has completely migrated from the LSC structure to the LSC/interconnect interface, which confirms LSC dissociation.

5
High temperature steam electrolysis is a promising technology, to convert electricity and steam or a mixture of steam and CO
2
, into H
2 or syn-gas (H
2
+ CO) respectively. It is carried out, at high temperatures above 800 0
C, in devices called Solid Oxide Electrolytic
Cells (SOECs). SOECs experience a loss in performance during the time of their operation at high temperatures. The chemical interaction of the constituent layers of
SOECs at the temperatures of operation has been suggested as a prevalent cause of their degradation, as in their Solid Oxide Fuel CELL (SOFC) counterparts. Specifically, diffusion of Cr-containing species from steel interconnects into the oxygen electrode and the bond layer in SOECs, and the long-range cation segregation in the bond layer have been of importance. This thesis was targeted at understanding the mechanism of Crmigration into the oxygen electrode and bond layer, and the cation segregation in the bond layer. The oxygen electrode and the bond layer of SOECs under analysis in this thesis had the perovskite structure. Lao
8 o.r
2
CoO
3
(LSC) was the bond layer and
Ao.
8
Sro.
2
MnO
3
(ASM*) was the oxygen electrode. We employed a novel methodology, integrating different spectroscopy and microscopy techniques to carry out the investigations. We began from a higher-level understanding, investigating the top surface of the LSC bond layer using Raman Spectroscopy. Subsequently, we studied the bond layer microstructure and microchemistry at the surface of LSC grains, and its variation across the LSC cross-section, using AES. Identification of the length-scale of the cation segregation and phase dissociation in the bulk of LSC grains at high resolution was done
by utilizing TEM/EDX techniques. The bulk-sensitive results from TEM/EDX
* The constituent A is proprietary information belonging to Ceramatec Inc.

complement the surface-sensitive results obtained by AES. We also conducted electrochemical experiments in a controlled environment to identify the role of electronic current on the mechanism of cation segregation and migration.
Raman spectroscopy identified secondary phases on the top surface of the LSC bond layer (at the LSC/interconnect interface). The results clearly showed that the LSC bond layer had, at least, partially dissociated and the poorly conducting secondary phases of
C030
4
, Cr
2
O
3
, LaCrO
3 and La
2
CrO
6 were formed. These observations indicated Cr migration from the steel interconnects into LSC.
AES results showed an average chromium fraction of 0.07 at the surface of LSC grains.
Furthermore, the LSC layer exhibited local variation in surface composition, with the
La/Co ratio varying from 0.67 to 16.37. This indicates inhomogeneous segregation of the
A and B site cations in the LSC layer and dissociation of LSC over the period of operation. While there was no Sr and a small fraction of Co was identified as remaining in the LSC cross sectional surface, AES spectra at the top of LSC (LSC/interconnect interface) showed the presence of Co-rich crystallites, lacking in La, and a Sr-rich surface layer, with Sr-content varying from 12 % 41% of all constituent elements. Due to the diffusion of Cr and its species into the bond layer microstructure, and drastic local variations of chemical composition due to cation segregation and migration of Sr and Co species to the LSC/interconnect interface, LSC did not remain in its initial perovskite phase by dissociating into secondary phases.
In order to isolate the exact mechanism of cation segregation, we simulated the operating conditions, in the absence of Cr, on reference cells having ASM electrodes coated with
LSC, on both sides of the ScSZ electrolyte. The cell operated in air at 8200C, under a constant current density of 0.40A/cm 2
, for 13.5 days. Cell resistance increased over the first 4.5 days and then stabilized. Significant degradation observed in this experiment suggests that there exists at least one other mechanism that contributes to the degradation even in the absence of Cr-poisoning. AES analysis on the reference cells showed that they stayed stable upon the half-cell operation. The chemical composition of LSC

remained uniform throughout. Thus, in the absence of Cr, the cells did not show uneven dissociation or segregation of cations in LSC. Secondary phases rich in either Co or Sr were absent at the LSC top surface of reference half-cell. Hence, the electronic or ionic current alone cannot be the dominating factor in Sr and Co long-range migration.
Presence of Cr in the bond layer and the associated thermodynamic driving forces are likely the major causes for the non-uniform dissociation of LSC bond layer structure.
We found that nano-scale Cr-rich regions were associated with enrichment of La and depletion of Co, indicating a phase between Cr and La. The La/Cr ratio in the bulk of
LSC confirmed this phase to be either La
2
CrO
6 or LaCrO
3
, depending on the location where the TEM sample was prepared from: close to the interconnect or close to the anode. Cr-rich phases measured approximately 1.5 jim x 1.5 pm, while Co-rich phases were smaller (approximately 1 pim in length). Furthermore, the Cr-content in the LSC bulk was more than at the LSC surface, on average 20% and 7%, respectively. Results for
Sr were consistent with those from AES: only 4.16% was found in the bulk of LSC and none on the surface. The results from TEM/EDX confirm the severe dissociation of the
LSC, and shows that the secondary phase formation was on the nano-scale.
Thus, we conclude that formation of poorly-conducting secondary phases and dissociation of the LSC bond layer leading to severe degradation of the SOEC performance. Our results indicate that the dominant cause for this is the migration of Crcontaining species from the steel interconnects into the bond layer microstructure, and the corresponding Cr-related thermodynamics under electrolytic polarization conditions.
Suggested Mechanism for Cr-related degradation
Based on the results from the AES, TEM, and electrochemical measurements in this thesis, we suggest the following mechanism as a likely path for the Cr-related degradation of the LSC bond-layer. The secondary phases can be formed by the electrochemical reduction of Cr** ions, present in the volatile Cr0
3 or CrO
2
(OH) species, at the oxygen electrode a process already observed for SOFCs. A non-electrochemical process kinetically limited by nucleation agents, between Cr0
3 or CrO
2
(OH) species and

..........
nucleation agents in the electrode can also form phases with low conductivity. The La-0and Sr-O-segregates, at the surface of LSC grains, are the likely nucleation agents in the latter theory. In our samples, a clear La-enrichment on the surface of both the as-prepared and the tested LSC layer grains were found by AES. Thus, it is likely that the La-enriched
La-0-phase, on the surface of LSC grains, potentially reacts with the Cr0
3 species, leading to the formation of LaCrO
3 or La
2
CrO phases observed. The proposed scheme governing the aforementioned reaction is shown below and its schematic is depicted in
Figure 5-1.
Cr20
3 ->
CrO (g)
CrO
3
(g)+La
-+
CrO
3
+ La
2
0
3
->La
2
(13)
Cr0
3
Figure 5-1: Possible mechanism for Cr-reaction in the LSC bond layer microstructure
The exact mechanism by which Cr causes such long-range transport of cations leading to
La-Cr phase formations, and the relation of this process to the electrochemical potential and gas pressure conditions in SOEC anode should be further investigated in terms of the thermodynamics involved in these reactions. We propose that an extension of this study is to perform controlled-environment experiments probing the phase and composition evolution, ideally under in situ conditions, coupled with material thermodynamics studies. This work then can more deeply address the stability related challenges in SOEC anode and bond layer materials to enable the successful development of the SOEC technology.

-
6.1 Raman Spectroscopy
Raman spectroscopy is a spectroscopic technique that is often employed in chemistry and particle physics to study vibrational and rotational modes in a molecule. It can be looked on as an inelastic collision between the incident photon and the molecule where, as a result of the collision, the vibrational or rotational energy of the molecule is changed
[31]. Monochromatic light, in the ultraviolet or near visible ultraviolet wavelength range, from a laser source is used to illuminate the sample to be analyzed. Generally, this radiation frequency is much higher than the vibrational frequencies but is lower than the electronic frequencies. The incident light interacts with the electron cloud of the bonds of the molecule. The incident energy of the photons excites the electron from the electron cloud into a virtual excited state lower in energy than an electronic transition. This electron in the excited energy state eventually de-excites to a lower energy vibrational state, thus emitting another photon. In order for the energy to be conserved, the energy of the scattered photon must be different from that of the incident photon. Numerically, the
Raman shift can be expressed as
-
V=
1 1
-
AZincident
Ascattered

If the molecule gains energy, then Acaered is larger than Aient
.
At low temperatures, the population of the vibrational excited states is very low. So most of the transitions are from an initial lower energy ground state to a relatively excited vibrational level. Thus the scattered photon has a lower energy and longer wavelength than the incident photon.
This is called Stokes Raman Scattering. At higher temperatures, the molecules are in the higher excited state and Raman scattering leaves the molecules into the low energy ground state. This is called anti-Stokes Raman Scattering. Different molecules have different vibrational and rotational modes and hence their interaction with the incident light is different from one another. So effectively, different molecules have different values for the Raman shifts, which can be thought of as their 'fingerprint' and can be used to identify them.
6.2 Auger Spectroscopy
It is another spectroscopic technique employed in studying surfaces. It is based on the
'Auger effect' named after its discoverer Pierre Auger. It is an electronic process in which low energy (in the range 2 50 keV) electrons or photons are incident on the sample surface. These electrons/photons interact with the electron cloud of an atom and eject an electron from the inner shell of the atom, leaving behind a hole. Hence, an electron from the outer electronic shell makes a transition into the hole left behind, thus emitting a photon with energy equal to the energy difference between the ground state
(with the hole) and the state which the electron transits from. This emitted photon can in turn excite another electron from an outer state if it is more energetic than the binding energy of the outer shell electrons. This emitted electron is called the 'secondary' electron. The secondary electrons are what are collected in Auger Spectroscopy and their energy dependence is chemical specific due to the characteristic orbital energies. These energies act as a 'fingerprint' for the elements and hence can help identify the compositions present. The energy of an Auger electron is small, typically 10 1000 eV and its mean free path in a condensed phase is 0.4 - 4 nm. Consequently, Auger electrons can be examined only when they originate from the surface, or near surface, of a material.
This surface sensitivity leads to the necessity of avoiding or removing any superficial

contamination. For the same reason, Auger is mostly carried out under ultra high vacuum else the electrons can easily be deflected by thermal motions of molecules in surrounding air [32].
Since the exciting electron beam can be ramped across a range of energies and Auger data obtained in seconds, it is possible to produce simultaneous analyses for a number of
N
EAdjgem.
*7
/Z
2S is
2 p
(a) (b)
Figure 6-1: Pictorial view of the steps in Auger spectroscopy. a) Incident photon ejecting an electron from the shell of an atom b) A secondary 'Auger' electron ejected from an outer shell elements on a surface. It is important to note that there are two principal difficulties associated with the Auger method of analysis the surface being analyzed may be modified by the incident electron beam and if the sample is an insulator, there may be a build up of charge upon it. To counter the possibility of sample damage due to the energy of electrons, the incident electrons can be accelerated through a lower potential so as to minimize the potential damage that the incident electrons can cause to the sample. In order to counter the charging on the sample, the sample should be properly grounded. If proper grounding cannot be achieved, one can use an ion gun that shoots positive ions on the sample to neutralize the build up of negative charge of the electrons. This process considerable improves the imaging and Auger spectrum collection.

6.3 Transmission Electron Microscopy
Transmission Electron Microscopy (TEM) [33] is a microscopic technique in which a beam of electrons is focused on a very thin (thickness < 100 nm) sample. During the process, the electron beam interacts with the crystalline structure of the sample and passes through it. As an electron travels through the sample, it is either scattered or it remains unaffected by the specimen, thus giving a non-uniform distribution of electrons.
The transmitted electrons are magnified and focused on an imaging screen or a photographic film and images of the sample can be captured. The image can also be detected on a CCD camera. The different angular distributions of the transmitted electron can lead to formation of diffraction patterns and the spatial distribution of electrons can be used in forming images and obtaining and image contrast. These diffraction patterns are the basis of all crystallographic analysis and image formation in TEM and can be used to obtain valuable information such as the crystallinity of the sample, the crystal lattice parameters and symmetry, phases present, orientation of the sample with respect to the electron beam, grain size etc. All the above-mentioned advantages are accompanied
by inherent drawbacks. First of all, under the TEM, we look at only a small part of the specimen. The higher the resolution, the worse the sampling ability of the instrument.
Thus, before looking at a sample in TEM, we must examine it with other techniques that provide lower resolution but better sampling, for example the visible light microscope or the Scanning Electron Microscope. Another problem with the TEM is that it provides us with a two-dimensional image of a 3-dimensional sample. Thus, mostly all the TEM information that we get is averaged through the thickness of the sample. A single TEM image has no depth sensitivity. Thus other techniques such as AES, which are surfacesensitive or depth-sensitive, need to complement the TEM. Also, since the entire technique is based on transmission of electrons through the sample, the sample needs to be electron transparent and, thus, has to be very thin. The requirement for thin samples is a major difficulty in carrying out TEM studies. Traditional mechanical thinning is a commonly employed method to obtain appropriate samples. In the following section we will also discuss the use of the Focused Ion Beams (FIB) as an alternative. But, as a general rule, the thinning processes can affect the specimen, changing both its structure

and chemistry. So one must be cautious to recognize the artifacts introduced by various sample preparation methods.
6.4 Focused Ion Beam
Focused Ion Beam (FIB) is a technique used for site-specific analysis, deposition and milling of materials. It is a dual beam instrument and has a Scanning Electron
Microscope (SEM) as well as a source of focused beam of metal ions. Gallium is the most commonly used metal in a FIB. In the common set up, a tungsten needle is heated and Ga metal is placed in contact with it and wets the surface of the tungsten needle. A large electric field is applied which ionizes the Ga metal and causes field emission of Ga atoms. These ions are then accelerated to energies normally ranging between 5-50 keV and focused on the sample. These ions can be focused on the samples into diameters smaller than 10 nm with current densities of several A/cm 2 [34]. The two most important features of the FIB tool are the capability to remove material from the sample by sputtering and to add materials to the sample by ion-induced reactions (deposition). Due to these features, FIB has found a variety of applications. One of them is Transmission
Electron Microscope (TEM) sample preparation, where the FIB is used to form a thin slice of material by sputtering trenches on either side of the slice. The prepared slice can be examined in a TEM while still attached to the sample or can be removed from the sample and attached to 'TEM grids' which can then be analyzed under TEM. While we have the traditional mechanical polishing method to prepare samples for TEM, FIB offers certain advantages over those methods [35]. a. No other technique can select the target area as precisely as FIB lamellae can be prepared with a spatial accuracy of within -20 nm.
b. FIB preparation techniques are virtually independent of the nature of the material.
The milling procedures for different materials require only minor adjustments, if at all, depending on the bulk properties of the material.
c. The FIB system that we used is a dual-beam set up and has a SEM in conjunction with the FIB. So we are able to image the sample under the SEM and hence can

choose the particular region of interest for TEM. This capability to selectively prepare TEM samples is our major motivation behind using the FIB instead of traditional mechanical polishing.
It is important to note that FIB has certain disadvantages as well. First, the ion collisions initiating sputter removal can also lead to ion implantation and cause severe damage to the remaining bulk of material [35]. It may lead to complete amorphization or formation of agglomerates in certain metals. It has also been reported in literature that FIB milling of fine grained fcc metals can change the orientation and size of the surface grains and form Ga intermetallic compounds [35].
There are various techniques that can be employed to use the FIB to prepare TEM samples. The techniques keep on evolving with time. The initial methods were based on a
'H-bar technique' in which samples were mechanically polished down to an approximately 50 micron lamella and then FIB was used to cut two trenches, one from each side, leaving behind a thin electron transparent lamella supported by bulk material on two opposite sides (Figure 6-2). In parallel, techniques were developed that make it possible to directly remove an electron transparent lamella from the bulk of the specimen without mechanical polishing. These 'lift out' techniques can be categorized into the ex-
situ and in-situ methods. In the ex-situ lift out (EXLO) technique, the lamella is thinned to electron transparency and is then removed from the trench and mounted on carboncoated TEM grids, which can then be analyzed in the TEM. EXLO is very fast but its difficulty is a major disadvantage. In the in-situ lift out (INLO), a small wedge shaped or parallel-sided piece is dug out from the sample and is attached to the TEM grid. The final thinning of the sample is done when the sample is on the TEM grid. FIB milling then progresses in a manner as for an H-bar specimen (Figure 6-3). For our use, we have employed the INLO technique for TEM sample preparation.

5n 4I
TEM
Figure 6-2: a) Schematic illustration of the H-bar FIB technique. b) SEM image showing the top-down view of an H-bar FIB specimen (taken from [351)
(a) (b)
Figure 6-3: In-situ lift out of a sample piece. a) A wedge shaped specimen being dug out from the bulk. b) The wedge mounted on the TEM grid (taken from [351)

7 Appendix B AES Data
We collected AES data from different regions (classification of regions described in
Section 4.2) in LSC cross-section of cells CER#3 and CER#5. Section 4.2 discussed the observations and results from the averaged AES data for each region of cells CER#3 and
CER#5. In this appendix, we provide the detailed AES data from each acquisition. We collected five different sets of data for each region for CER#3 and one set of data for each region of CER#5. Each data set consisted of three acquisition areas. Thus for each of the three regions of CER#3, we acquired AES spectra from 15 different areas.
Similarly, for each region of CER#5, we acquired AES spectra from 3 different areas.
The following plots show the chemical compositions of La, Cr and Co from each acquisition area in the three regions of CER#3 and CER#5.

CER#3, Region 1
0 1
J
=3=
2 3 4 5 6 7 8 9 1 ) 11 12 13 14 15 16
Data points
Figure 7-1: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 1 of CER#3
CER#3, Region 2
100
90
80
-70 m 60
50
A
40
Qs 30
0
20
10
A
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15
Data Points
Figure 7-2: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 2 of CER#3

CER#3, Region 3
60
50
40
30
20
10
0
100
90
80
70
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16
Data Point
-0-Cr
-+-Co
Figure 7-3: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 3 of CER#3
CER#5, Region 1
-0-C"
-U-Co
Data Points
Figure 7-4: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 1 of CER#5

CER#5, Region 2
-U*-Co
-*e-La
Data Points
Figure 7-5: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 2 of CER#5
CER#5, Region 3
-E'-Cr
-U-Co
-*-L a
1
Data Points
Figure 7-6: Cation composition, normalized with respect to the sum of La, Co and
Cr, for different AES acquisition areas in region 3 of CER#5

It is evident from the figures above that the chemical fractions of La, Cr and Co vary randomly in the different regions of CER#3 and CER#5 and do not follow a particular pattern.

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