ACTAUNIVERSITATIS
UPSALIENSISUPPSALA
2019
Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 1835
Anion redox processes in novelbattery cathode materialsinvestigated by soft X-rayspectroscopy
FELIX MASSEL
ISSN 1651-6214ISBN 978-91-513-0714-5urn:nbn:se:uu:diva-390623
Dissertation presented at Uppsala University to be publicly examined in Room 80101,Ångström Laboratory, Lägerhyddsvägen 1, Uppsala, Friday, 27 September 2019 at 09:15 forthe degree of Doctor of Philosophy. The examination will be conducted in English. Facultyexaminer: Professor Kevin Smith (Department of Physics, Boston University).
AbstractMassel, F. 2019. Anion redox processes in novel battery cathode materials investigated bysoft X-ray spectroscopy. Digital Comprehensive Summaries of Uppsala Dissertations fromthe Faculty of Science and Technology 1835. 73 pp. Uppsala: Acta Universitatis Upsaliensis.ISBN 978-91-513-0714-5.
This thesis presents experimental investigations of the electronic structure of emerging andnovel cathode materials used in lithium- and sodium-ion batteries. The investigated materialsinclude a range of oxide materials containing the elements nickel and manganese. Central goalsare to find fundamental explanations for favorable, respectively, unfavorable electrochemicalcycling behavior and to arrive at a better understanding of the roles that the different elementalconstituents of the compounds play. The experiments are based on the application of X-rayAbsorption Spectroscopy (XAS) and Resonant Inelastic X-ray Scattering (RIXS) in the soft X-ray region and have been performed at synchrotron radiation facilities such as The AdvancedLight Source (USA), The Swiss Light Source (Switzerland) and SPring-8 (Japan).
XAS and RIXS of spinel LiNi0.44Mn1.56O4 at the O K-edge as well as the Ni and Mn L-edges were measured for two different crystal structures, namely, transition-metal-ordered and-disordered, respectively. The results show that both Ni and O contribute strongly as redoxcenters for the charge compensation during electrochemical cycling. The Ni L-RIXS spectrashow evidence of a more stable Ni--O bond in the disordered material.
In the layered manganese oxide materials Li[Li0.2Ni0.2Mn0.6]O2, Na0.67[Mg0.28Mn0.72]O2, andNa0.78[Li0.25Mn0.75]O2, as well as the disordered Li1.9Mn0.95O2.05F0.95 one observes that reversible Oredox leads to two distinct features in O K-RIXS. Both features resonate in a narrow incidentenergy range suggesting that localized O hole states are formed, one close to the elastic peakand the other as a strong emission peak at an energy loss of about 8 eV. These features appearreversibly on the voltage plateau of the charge-discharge curve and can be used to identify acertain type of O redox reactions.
The work also includes investigations that compare two different compositions of thestructurally related material Li2MnO3 grown epitaxially as thin films. Evidence is found foranionic activity during the initial cycle that is of a different kind than the above as no evidencefor localized O holes is found. Instead, excess Li in the transition metal layer is shown to leadto a more rapid loss of covalency in the Mn--O bonds.
In short, this work presents some of the first explorations into the role of different types ofanionic redox centers in cathodes, by means of XAS and RIXS thereby also demonstrating theutility and power of synchrotron based techniques for gaining atomic-level understanding ofbattery electrode materials.
Keywords: soft X-ray spectroscopy, X-ray absorption spectroscopy (XAS), resonant inelasticX-ray scattering (RIXS), lithium-ion battery (LIB), sodium-ion battery (SIB), anionic redox,cathode materials, layered manganese oxide, spinel LNMO
Felix Massel, Department of Physics and Astronomy, Molecular and Condensed MatterPhysics, Box 516, Uppsala University, SE-751 20 Uppsala, Sweden.
© Felix Massel 2019
ISSN 1651-6214ISBN 978-91-513-0714-5urn:nbn:se:uu:diva-390623 (http://urn.kb.se/resolve?urn=urn:nbn:se:uu:diva-390623)
Dedicated to my family.Past, present, and future generations.
List of papers
This thesis is based on the following papers, which are referred to in the textby their Roman numerals.
I The role of anionic processes in Li1−xNi0.44Mn1.56O4 studied by
resonant inelastic X-ray scattering
F. Massel, B. Aktekin, Y.-S. Liu, J. Guo, D. Brandell, R. Younesi, M.Hahlin, and L.-C. DudaIn manuscript.
II How Mn/Ni ordering controls electrochemical performance in
high-voltage spinel LiNi0.44Mn1.56O4 (LNMO) with fixed oxygen
content
B. Aktekin, F. Massel, M. Ahmadi, M. Valvo, M. Hahlin, W. Zipprich,F. Marzano, L.-C. Duda, and R. Younesi, K. Edström, D. BrandellIn manuscript.
III Excess lithium in transition metal layers of epitaxially grown thin
film cathodes of Li2MnO3 leads to rapid loss of covalency during
first battery cycle
F. Massel, K. Hikima, H. Rensmo, K. Suzuki, M. Hirayama, C. Xu, R.Younesi, Y.-S. Liu, J. Guo, R. Kanno, M. Hahlin, and L.-C. DudaSubmitted to J. Phys. Chem. C.
IV Anion Redox Chemistry in the Cobalt Free 3d Transition Metal
Oxide Intercalation Electrode Li[Li0.2Ni0.2Mn0.6]O2K. Luo, M. R. Roberts, N. Guerrini, N. Tapia-Ruiz, R. Hao, F. Massel,D. M. Pickup, S. Ramos, Y.-S. Liu, J. Guo, A. V. Chadwick, L.-C.Duda, and P. G. BruceJ. Am. Chem. Soc. 138 (2016)
V Lithium manganese oxyfluoride as a new cathode material
exhibiting oxygen redox
R. A. House, L. Jin, U. Maitra, K. Tsuruta, J. W. Somerville, D. P.Förstermann, F. Massel, L.-C. Duda, M. R. Roberts, and P. G. BruceEnergy Environ. Sci. 11 (4 2018), pp. 926–932
Continued on next page.
VI Oxygen redox chemistry without excess alkali-metal ions in
Na2/3[Mg0.28Mn0.72]O2U. Maitra, R. A. House, J. W. Somerville, N. Tapia-Ruiz, J. Lozano, N.Guerrini, R. Hao, K. Luo, L. Jin, M. A. Pérez-Osorio, F. Massel, D. M.Pickup, S. Ramos, X. Lu, D. E. McNally, A. V. Chadwick, F. Giustino,T. Schmitt, L.-C. Duda, M. R. Roberts, and P. G. BruceNat. Chem. 10.3 (2018), pp. 288–295
VII What Triggers Oxygen Loss in Oxygen Redox Cathode Materials?
R. A. House, U. Maitra, L. Jin, J. G. Lozano, J. W. Somerville, N. H.Rees, A. J. Naylor, L.-C. Duda, F. Massel, A. V. Chadwick, S. Ramos,D. M. Pickup, D. E. McNally, X. Lu, T. Schmitt, M. R. Roberts, and P.G. BruceChem. Mat. 31 (2019), pp. 3293–3300
VIII Understanding charge compensation mechanisms in
Na0.56Mg0.04Ni0.19Mn0.70O2L. A. Ma, F. Massel, A. J. Naylor, L.-C. Duda, and R. YounesiSubmitted to Comms. Chem.
Reprints were made with permission from the publishers.
Extended bibliographyThe following paper has been omitted from this thesis due to the character ofthe investigated material.
i Transition metal doping effects in Co-phosphate catalysts for water
splitting studied with XAS
F. Massel, S. Ahmadi, M. Hahlin, Y.-S. Liu, J.-H. Guo, T. Edvinsson,H. Rensmo, and L.-C. DudaJ. Electron Spectros. Relat. Phenomena 224 (2018), pp. 3–7
Comments on my own participationThe level of my personal contribution to the overall products is approximatelyreflected by my position in the respective author list. In all projects I took partin or led the conducting of experiments at the synchrotron facilities and theanalysis of the X-ray spectroscopy data while much of the electrochemistrywas performed by my collaborators. I also took part in the discussion of theoverall results. I wrote the majority of the manuscripts for Papers I and III aswell as the XAS and RIXS portions of Papers II and VIII.
Contents
1 Introduction and background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111.1 The scope of this thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12
2 Spectroscopy and matter . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 132.1 Microscopic structure of matter . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13
2.1.1 Basics of atom physics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 132.1.2 Quantum mechanical principles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 142.1.3 The electron cloud atom model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 162.1.4 From atoms to molecules to solids . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
2.2 Light matter interaction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 192.3 X-ray spectroscopy techniques . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22
2.3.1 Synchrotron radiation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 222.3.2 X-ray absorption spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 232.3.3 Resonant inelastic X-ray scattering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26
3 Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293.1 Working principle . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29
3.1.1 Cationic versus anionic redox . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 313.2 Emerging cathode materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
3.2.1 From layered to Li-rich layered transition metal oxidecathodes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
3.2.2 Spinel LiNi0.5Mn1.5O4 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 353.2.3 Sodium-ion battery cathodes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 363.2.4 Single crystal thin film cathodes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
4 Summary and discussion of results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 384.1 Lithium-ion battery investigations: spinel materials . . . . . . . . . . . . . . . . . . . 384.2 Lithium-ion battery investigations: Li-rich materials . . . . . . . . . . . . . . . . . . 45
4.2.1 The effect of excess Li on covalency in Li2MnO3 thinfilms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45
4.2.2 The role of oxygen in the lithiation ofLi[Li0.2Ni0.2Mn0.6]O2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49
4.2.3 Localized oxygen hole states in novel, oxygen redoxactive, disordered Li1.9Mn0.95O2.05F0.95 . . . . . . . . . . . . . . . . . . . . . . . . 50
4.3 Sodium-ion battery investigations: layered materials . . . . . . . . . . . . . . . . . . 534.3.1 No oxygen loss in the highly anionic redox active
Na0.67[Mg0.28Mn0.72]O2 cathode . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53
4.3.2 The reason for oxygen loss in anionic redox activecathodes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55
4.3.3 The charge compensation mechanism of theNa0.56Mg0.04Ni0.19Mn0.70O2 cathode . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
5 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61
6 Populärvetenskaplig sammanfattning . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62
7 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65
Abbreviations
AM Alkali metalAO Atomic orbitalEY Electron yieldGS Ground stateICP-OES Inductively coupled plasma optical emission spectroscopyIPFY Inverse partial fluorescence yieldLCAO Linear combination of atomic orbitalsLCO LiCoO2LIB Lithium-ion batteryLMO LiMn2O4LMOF Li1.9Mn0.95O2.05F0.95LNMO LiNi0.5Mn1.5O4LUMO Lowest unoccupied molecular orbitalLrMO Lithium-rich manganese oxide Li[Li0.33Mn0.67]O2LrNMC Lithium-rich NMC, Li[Li1– x– y– zNixMnyCoz]O2MO Molecular orbitalNLMO Na0.78[Li0.25Mn0.75]O2NMC Li[Ni1– y– zMnyCoz]O2NMMO Na0.67[Mg0.28Mn0.72]O2NMNMO Na0.56Mg0.04Ni0.19Mn0.70O2OEMS Operando electrochemical mass spectrometryPFY Partial fluorescence yieldRIXS Resonant inelastic X-ray scatteringSEI solid electrolyte interfaceSHE Standard hydrogen electrodeSIB Sodium-ion batterySoC State of chargeTEM Transmission electron microscopyTEY Total electron yieldTFY Total fluorescence yieldTM Transition metalXAS X-ray absorption spectroscopyXRD X-ray diffraction
1. Introduction and background
Batteries are an important part of energy storage technology and are currentlyrising in relevance due increasing demand for electric vehicles, portable con-sumer electronics, and intermittent energy sources. In order to meet the highand varying demands these applications pose on various aspects of batteryperformance, including gravimetric and volumetric energy density, long-termstability, operational safety, or costs of constituents, immense scientific effortis invested to push the boundaries of current materials and to discover com-pletely new and better performing materials.
When aiming to optimize battery performance it is imperative to understandthe underlying mechanisms at all length scales. The processes responsible forthe charging and discharging mechanism for example take place on atomiclength scales and thus require highly sophisticated experimental techniquesfor their investigation. X-ray core hole spectroscopy utilizes fundamental pro-cesses on this length scale and thus delivers element specific information aboutatomic states and molecular bonds. This makes this technique advantageousfor the investigating of compounds consisting of various chemical elementswhich is the case for batteries. In particular, X-ray absorption spectroscopy(XAS) and resonant inelastic X-ray scattering (RIXS) provide the possibilityto probe specific chemical sites and thus allow the investigation of their rolein the chemical reactions of the respective battery material.
The specifics of energy storage mechanisms vary greatly for different bat-tery types. In the current field of battery technology intercalation materialsrepresent an ever-growing market share especially in the area of mobility andhave been dominating the marked of consumer electronics for some years al-ready. The understanding of the underlying charge and discharge mechanismsof many cathode materials is traditionally geared towards a cationic redox oftheir transition metal components. However, some materials exhibit a capacitythat far exceeds the theoretical limits assuming cationic redox only. The expla-nation of the source of this extra capacity has been an ongoing debate for manyyears and the questions at the center of this debate can only be answered withan understanding of the mechanisms at atomic length scales. The experimen-tal approach used in this thesis delivers key contributions to this understandingand specifically resulted in some of the first XAS and RIXS studies provingthe involvement of anionic redox processes in certain cathode materials.
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1.1 The scope of this thesisIn this thesis I investigate the electronic structure of both various well-knownas well as a number of emerging, novel Li/Na-ion battery cathode materials.My foremost experimental tools have been X-ray absorption spectroscopy andresonant inelastic X-ray scattering. Thus while electrochemical preparationand other chemistry-related tools are essential parts of my published papers(and are the expertise of some of my collaborators), the focus of this thesisis on demonstrating and utilizing the power of these X-ray spectroscopies forelucidating the electronic structure of battery materials. It is shown that usingthese tools and analyzing the spectra adds qualitatively new insight into theworking of such batteries, in particular, by dissecting the roles of anionic andcationic reaction centers.
This thesis is structured in a way to deliver an overview of concepts andtechniques necessary for a proper interpretation of the experimental results.Chapter 2 presents the essential building blocks of the theoretical frame worknecessary to understand the experimental techniques used in this thesis. Chap-ter 3 presents short reviews to the state of the art of the respective cathode ma-terials investigated in this work. And finally Chapter 4 summarizes the mostimportant results from the publications featured in this thesis with a focus onthe work that the author contributed to them.
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2. Spectroscopy and matter
The aim of this chapter is to give a short overview of the most important con-cepts that are the phenomenological and theoretical foundation for the researchpresented in this thesis. These concepts are based on an enormous body ofscientific work performed by a series of many generations of minds infinitelygreater than my own one. Before presenting these concepts, I want to remindthe reader of the famous quote [1] of ISAAC NEWTON:
If I have seen a little farther than others it is because I have stood on the shoul-ders of giants.
Most, if not all observations made throughout the experimental work in thisthesis are virtually meaningless without the proper context set by these giants.
2.1 Microscopic structure of matterBesides presenting fundamental concepts and terms important for the interpre-tation of this work, this section is aiming to show the close historical interac-tion between the development of spectroscopy and our understanding of whatordinary matter truly is. Based on this understanding the following sectionswill introduce different types of X-ray spectroscopy central to the thesis.
2.1.1 Basics of atom physicsMost of the physical world surrounding us is made up of the small buildingblocks known as atoms. While by no means “uncuttable” into smaller subunitsas the Greek origin of the word implies, atoms are a set of the unique entitiesthat separate the elements of the periodic table and therefore determine a largeportion of the properties of ordinary matter. The size of atoms is in the orderof a few Ångströms (1A = 10−10m) and while over 99.9% of the mass iscentered in the positively charged core or nucleus, a yet even larger portionof the volume is taken up by the much lighter, negatively charged electronsfluctuating around the nucleus. The electric charge of a nucleus is an integermultiple of an elementary charge. This integer is called atomic number, whichis the number of protons in the nucleus and defines which element it is, e. g.oxygen always has 8 protons, while nickel has 28. Neutral atoms have an equalnumber of protons and electrons and when a number of electrons are added or
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removed (ionized ), the atom is referred to as anion or cation, respectively.The outer most electrons determine much of the interaction of an atom withsurrounding atoms and are thus most important for the chemical dynamics ofordinary matter.
The spatial distribution of electrons around the core is structured in a veryspecific, highly non-trivial way and it took a large amount of outside-the-boxthinking by a series of geniuses to refine the solar-system-like Rutherford-Bohr model to the modern orbital shell model. NIELS BOHR’s famous postu-lates from 1913 about the behavior of bound electrons, which he assumed tohave discrete orbits associated with quantized energy levels and orbital angu-lar momenta [2], meant a major step forward in the understanding of atomsand one of its successes was the accurate theoretical prediction of the Rydbergconstant for hydrogen, which had been empirically know from spectroscopicexperiments. While Bohr’s postulates turned out to be spot-on and remainedin some form an integral part of more advanced models, this early model hadserious shortcomings. Among them is the fact that it only gives accurate re-sults for hydrogen-like atoms with not more than one electron and more im-portantly that the discrete electron orbits predict unstable atoms as the radiallyaccelerated motion would lead to a continuous dissipation of energy and thusa collapse of the orbiting electron into the core.
It was the incorporation of quantum mechanical principles in the 1920’sinto Bohr’s model that removed many of the remaining contradictions andeventually enabled the much improved theoretical prediction of spectral linesalso for many-electron atoms.
2.1.2 Quantum mechanical principlesOne key aspect of quantum mechanics is the realization that when approach-ing atomic length scales (100 nm and below) entities start showing behaviorknown from waves when probed for certain properties while still behavinglike particles when probed for others. This wave-particle duality was firstshown for light quanta, so called photons, beginning with MAX PLANCK’sexplanation for black-body radiation in 1900 [3] and again with ALBERT EIN-STEIN’s successful explanation of the photoelectric effect in 1905 [4]. It fi-nally gained wide-spread acceptance after ARTHUR COMPTON’s experimentsin 1923. Planck first postulated that the energy E of any light quantum is amultiple of an elementary unit:
E = hν , (2.1)
with the Planck constant h= 6.626×10−34 Js of dimension action and the fre-quency ν of the light quantum. Even though Planck adopted the idea of quan-tized energy packages only reluctantly, it is today considered the birth point ofquantum mechanics. Shortly after the discovery of the Compton effect, which
14
confirmed the particle nature of photons [5], LOUIS DE BROGLIE proposed inhis PhD thesis the extension of the wave-particle duality to electrons, relatingtheir momentum p to a wavelength λ according to the equation
λ =hp. (2.2)
Because of the very small action constant in relation to impulses of macro-scopic objects, the likelihood of observing this property in macroscopic matteris very close to zero, and thus naturally appears alien to everyone only familiarwith the objects of common perception. Nonetheless, scattering experimentsof electron beams that produced interference patterns similar to those knownfrom X-rays confirmed this bizarre property of microscopic matter [6]–[8].This encouraged the development of ERWIN SCHRÖDINGER’s wave formu-lation of quantum mechanics in 1926. The Schrödinger equation [9] in itstime-independent form is stated as
H |Ψ〉= E |Ψ〉, (2.3)
where H is the Hamiltonian operator accounting for the kinetic and potentialenergies of the particles in the system. The system’s state vector |Ψ〉 is aneigenfunction of H and the solution of this equation is the set of eigenvalues Erepresenting the possible, discrete energy states the system can take up. WhileSchrödinger’s formulation marked a revolutionary step in the description ofquantum mechanical systems, equation (2.3) can only be solved exactly forhydrogen-like atoms, for many-electron systems certain approximations mustbe included in the Hamiltonian.
So what is the wave function |Ψ〉? Later the same year in which Schrödingerdeveloped his formulation, MAX BORN presented his statistical interpretationof |Ψ〉 [10]. According to this interpretation, the absolute square of the wavefunction, |Ψ|2, represents the probability to observe an electron in a certainregion in space and time. Repeating this hypothetical experiment many timeswith the same conditions would yield the probability density described by|Ψ|2, meaning that this is a representation of the inner, electronic structure ofatoms.
This probabilistic nature of physics at the atomic level was rooted even fur-ther in reality through the introduction of the uncertainty principle by WER-NER HEISENBERG [11], which came as a result of the matrix mechanics for-mulation of quantum mechanics that he developed in cooperation with Bornand PASCUAL JORDAN in 1925 [12]. In this alternative, but ultimately equiv-alent formulation to Schrödinger’s wave formulation physical quantities orrather observables are mathematically described by matrix operators and ifthese operators do not commute, a simultaneous measurement becomes trou-blesome. This is e. g. true for the two observables position and momentum, aswell as for energy and time. In the latter case the uncertainty principle can be
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stated through the inequality
ΔEΔt ≥ h, (2.4)
where h= h/2π is the reduced Planck constant and ΔE and Δt are uncertaintiesin the determination of the observables energy and time, respectively. Theline width of spectroscopic emission lines is a direct result of this principle,called lifetime broadening, as short-lived excited states (small Δt) decay bythe emission of photons with less defined energy (big ΔE) compared to stateswith longer lifetimes.
2.1.3 The electron cloud atom modelTurning back to equation (2.3) and assuming an electron in a bound statewith a nucleus, which is mathematically reflected by choosing an appropri-ate Hamiltonian, H, the resulting set of possible solutions are described by acombination of only four, so called quantum numbers. Each of these combi-nations describe a possible state a bound electron can take, which is referredto as electron or atomic orbital (AO) and its associated, respective probabilitydensity |Ψ|2 is the quantum mechanical progression of a postulated discreteorbit in the Bohr model. AOs can be categorized into main shells and theirsubshells. The former are assigned a positive integer known as the principalquantum number n= 1,2,3, . . . With increasing n the center of the probabilitydensity shifts progressively further away from the nucleus.
The subshells represent an additional degree of freedom within each mainshell and are assigned the second, so called azimuthal quantum number l =0,1, . . . , n−1. This integer value is related to an observable L =
√l(l+1)h,
which is generally referred to as the electron’s angular momentum stemmingfrom the fact that its mathematical description is largely analogous to its name-sake in classical mechanics. The azimuthal quantum number determines theorbital shape. Subshells with varying l have different binding energies, whichexplains the different spectral series that were categorized into so called princi-pal, sharp, diffuse, and later also fundamental groups. This historical spectro-scopic nomenclature carried over into chemistry through the work of FRIED-RICH HUND [13], who correctly assigned the s, p, d, and f series to l = 0,1,2,and 3, respectively1.
The third, magnetic quantum number ml is connected to the projection ofL onto an arbitrary direction in space via Lz =mlh and determines the numberof orbitals in a given subshell by taking up the integer values from −l to +l.And finally, as WOLFGANG PAULI showed in his profound 1925 paper [14],a fourth, so called spin quantum number, ms, completes the description of anypossible bound state of an electron in an atom. This number is the projection
1Similarly, in X-ray spectroscopy notations are commonly used that assign the letters K, L, andM to the n= 1,2, and 3 main shells with subscripts further differentiating the specific orbitals.
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of an intrinsic degree of freedom of subatomic particles, called spin s. Forelectrons the spin always has the value s = 1/2 and its projection onto thesame arbitrary axis as above can take the two values ms = ±1/2, commonlyreferred to as up- and down-spin, respectively. This allows two electrons,one with up- and one with down-spin, to reside in the same AO, which is theupper limit of occupancy according to the exclusion principle that states thatno two electrons in a given system can have the same set of the four introducedquantum numbers.
-2-1
01
2ml
n = 1
n = 2
n = 3
E
sp
dl
Figure 2.1. Energy level diagram with splitting of main shells and subshells. The pand d subshells have 3 and 5 degenerate energy levels. The schematically depictedorbitals are calculated for the one-electron atom hydrogen with the visualization soft-ware Orbital Viewer.
With this description it is possible to explain the electronic structure of anyatom of the periodic table. The shapes or more precisely equiprobabilisticsurfaces of AOs of a hydrogen atom are depicted schematically in Figure 2.1along with their respective energy levels. The different colors, red and yel-low, represent the positive and negative phases of the wave function, which
17
becomes important when looking at the combination of any two AOs. Theshapes depicted are the surfaces in which an electron occupying the respectiveorbital would be observed 90% of the time. In the remaining 10% of obser-vations the electron would have a position outside of these surfaces, whichleads to a more blurry picture and to the reason this model is often referredto as electron cloud model. Energy level diagrams typically omit the depic-tion of orbital shapes and only show the black lines as representations of therespective AOs in two dimensions, which are also called energy levels.
In many cases the three p orbitals, as well as the five d orbitals have thesame respective energy. These orbitals are referred to as degenerate energylevels or energy levels of the degeneracy 3 and 5.
2.1.4 From atoms to molecules to solidsIn molecules, which are bound states of at least two atoms, the higher re-spective AOs are generally strongly modified by the interaction of the con-stituents’ valence electrons. Molecular orbitals (MOs) were first introduced byHund and ROBERT MULLIKEN in 1927 and 1928. The linear combination ofatomic orbitals (LCAO) approximation for molecular orbitals was introducedin 1929 by JOHN LENNARD-JONES. His ground-breaking paper showed howto derive the electronic structure of the fluorine and oxygen molecules fromquantum principles. LCAO can be used to estimate the molecular orbitals thatare formed upon bonding between the molecule’s constituent atoms. Similarto an atomic orbital, a Schrödinger equation can be constructed for an MO aswell.
(b) Lin n = 1 n = 2 n = 3 n = 4 n = 1023
2p
2s
(a)
2s2s
1s 1s
2py2px 2pz 2py 2px2pz
O AOs AOs O AOs AOs O2 MOs MOs 1s
2s-band
2p-band
1s-band
E
N(E)
EFValence
band
Bandgap
Conduc-tion band
Figure 2.2. (a) Energy level diagram of two O atoms combining their AOs to MOs ina O2 molecule. (b) Incremental densification of MOs to energy bands in a solid.
18
When AOs combine, the resulting MOs can be of three types: bonding,antibonding, or nonbonding. Figure 2.2 (a) schematically shows the resultingenergy diagram of two O atoms in a bound state of an O2 molecule. Bondinginteractions between AOs are constructive (in-phase) interactions and bondingMOs are lower in energy than the AOs that combine to produce them. Anti-bonding interactions between AOs are destructive (out-of-phase) interactions,with a nodal plane where the wave function of the antibonding orbital is zerobetween the two interacting atoms. Antibonding MOs are higher in energythan the AOs that combine to produce them. Nonbonding MOs are the re-sult of no interaction between AOs because of lack of compatible symmetries.Nonbonding MOs will have the same energy as the AOs of one of the atomsin the molecule.
Solids consist of a large number N of bound atoms. As schematically de-picted in Figure 2.2 (b), the addition of an increasing number of atoms to asystem incrementally adds occupied and unoccupied MOs whose energy lev-els become increasingly dense so that in the limit of N → 1023 the originallydiscrete energy levels densify to continuous energy bands.
2.2 Light matter interactionThe microscopic nature of the objects of interest in this thesis makes it im-possible for their inner structure to be observed directly, thus one needs torely on indirect means of investigation instead: spectroscopy. This is doneby disturbing the system with light of known initial energy and intensity, andsubsequently observing the reaction of the system. Not only does the correctinterpretation of the system reaction require a framework to describe matteron its subatomic length scale, which was established in section 2.1, it alsorequires a good understanding of the underlying dynamic processes. A shortintroduction of such light matter interaction follows below.
A first phenomenological approach to this that utilizes the quantum me-chanical concepts introduced in section 2.1.2 is to think of light as a numberof photons given by the intensity. Each photon can be assigned an energy givenby equation (2.1). If certain conditions are fulfilled an absorption process cancommence in which a photon may annihilate and deposit all its energy into asystem. If the deposited energy is high enough it can then excite an electronfrom a core level into a level of higher energy and thus creating a core hole.Four principal processes involving the creation of a core hole are schematicallydepicted in Figure 2.3. In the first two cases (a) and (b), referred to as X-rayphotoelectron spectroscopy (XPS) and normal X-ray emission spectroscopy(NXES), the incident photon energy is high enough to excite the electron intoa continuum state and completely remove it from and thus ionizing the sys-tem. In the other two cases (c) and (d), called X-ray absorption spectroscopy(XAS) and resonant X-ray emission spectroscopy (RXES), the incident pho-
19
Continu-um state
Conduc-tion band
Valence band
Corelevel
E
XPS
E
NXES
E
XAS
E
RXES
(a) (b) (c) (d)
Figure 2.3. Schematic representation of four different core hole spectroscopic pro-cesses: XPS, NXES, XAS, and RXES.
ton energy is smaller so that the core electron is excited into an empty energylevel. In these cases the excited electron is still bound to the system and thusthe processes are electrically neutral. XPS and XAS are first-order optical pro-cesses that involve only one (the incoming) photon, while both XES processesare second-order involving the absorption of one photon and the emission of asecond photon. In the examples of Figure 2.3 (b) and (d) the emission of thesecondary photon is a result from the re-filling of the core hole by an electronfrom the valence band. These four exemplary processes are the basis of thevarious spectroscopy techniques with the same name, which require differentinstrumentation and provide different information about the material.
In each of the four examples the system is in an excited final state andrelaxation back into its ground state will then proceed in one out of many dif-ferent ways, so called deexcitation channels, e. g. by emitting another photonor Auger electron. As a consequence of quantum mechanical principles it isimpossible to predetermine which deexcitation channel the system will takeafter any one specific absorption. Certain channels are, however, more likelythan others and therefore spectroscopy relies on statistics meaning that thesame experiment, i. e. absorption of a photon with certain energy and record-ing of the system reaction, needs to be repeated many times for a statisticallymeaningful observation.
In the concrete case of battery materials, one is generally interested in thereactions during charge and discharge and thus one wants to know, e. g. whichelements are the reaction centers and what is their inner structure close to thechemically relevant valence shell. For this purpose core hole spectroscopy isideal, as the energy levels of core electrons from different elements are gener-ally well separated, well known, and sufficiently well defined as the chemicalbonding in the specific material only slightly affects these levels. Binding en-
20
ergies of core electrons depend mostly on the number of protons given by theatomic number of the specific element. For example the O 1s level (K-edge)or the transition metal 2p1/2 (L2-edge) and 2p3/2 levels (L3-edge) are locatedin the energy range of 102 to 103 eV. If incoming photons are in this energyrange they are referred to as soft X-rays. Core electrons have localized wavefunctions so that the transitions upon X-ray absorption promote these electronsto unoccupied orbitals of the same atom. Additionally, the transition proba-bility is heavily effected by selection rules taking the symmetries of the twoinvolved orbitals as well as conservation laws into account. Thus core levelspectroscopies are various techniques to obtain element-specific informationof the electronic structure around an absorption site and are thereby suitabletools to study the chemical state, local geometric structure, nature of chemi-cal bonding, and dynamics in electron transfer processes centered around oneatomic site.
In the case of first-order optical processes, e. g. the XPS and XAS processesdepicted in Figure 2.3 (a) and (c), respectively, the transition probability Γ forthe quantum mechanical system to go from an initial ground state |i〉 of energyEi over into a final state | f 〉 of energy Ef upon the absorption (or emission) ofa photon with energy hν is described by Fermi’s Golden Rule:
Γi, f =2πh|〈 f |H|i〉|2 δ (Ef −Ei±hν). (2.5)
The Dirac delta function ensures energy conservation, while −hν and +hνrepresent the case of absorption and emission, respectively. The perturbationoperator H is a Hamiltonian describing the interactions of the electron withthe photon.
As mentioned earlier electronic transitions are governed by certain, so calleddipole selection rules. In the case of atoms the azimuthal quantum numbermust change by Δl = ±1, the magnetic quantum number must change byΔml = ±1 or 0 and the spin must be conserved Δs = 0. In the molecularcase the so called Laporte rule states that transitions in centrosymmetric sys-tems must involve a change of parity. So as a consequence the probability of atransition between subshells of the same symmetry, e. g. 1s to 2s or 3d to 3d, isseverely reduced, albeit not zero because of a possible coupling to vibrations.This means that e. g. the absorption at the O K-edge (excitation of a 1s coreelectron) is much more sensitive to unoccupied states of 2p character then, say,states of 3s character.
Turning to the second-order optical processes, that involve both the absorp-tion of one photon and the subsequent emission of another photon like e. g.the NXES and RXES processes depicted in Figure 2.3 (b) and (d), respec-tively, the description requires an intermediate state. HENDRIK KRAMERSand Heisenberg developed a description for this case [15], which was broughtto a quantummechanical description shortly after by Dirac [16]. The Kramers-Heisenberg equation contains the following proportionality for the differential
21
cross section of a second-order optical process:
d2σdΩd(hν2)
∝ν2
ν1∑| f 〉
∣∣∣∣∣∑|m〉〈 f |T †|m〉〈m|T |i〉
Ei−Em+hν1+ iΓm2
∣∣∣∣∣2
δ (Ei−Ef +hν1−hν2) (2.6)
Here |i〉, |m〉,and| f 〉 are the initial, intermediate, and final states of the systemwith energy Ei, Em, and Ef , respectively. The incident and emitted photonshave energy hν1 and hν2, respectively, and Dirac delta function again ensuresenergy conservation. The solid angle dΩ is centred in the direction of theemitted photon and T is the transition operator containing the interactions ofelectrons and photons. Finally, Γm is the intrinsic linewidth of the intermediatestate.
2.3 X-ray spectroscopy techniquesThe following section briefly discusses the experimental methods and the cor-responding instrumentation utilized in the research of the different materialsin this thesis. It is an introduction on how the theoretical framework estab-lished in section 2.2 can be utilized in an experimental context of the specifictechniques that were used in the course of this work. Several textbooks wereused as the basis for writing this chapter [17]–[19].
2.3.1 Synchrotron radiationPerforming X-ray spectroscopy with appropriate statistics and resolution re-quires X-rays beams of high intensity and a small bandwidth, i. e. a largeamount photons ideally of the same energy. Additionally in order to sweepover a desired absorption edge in XAS or to target a specific resonance inRIXS it is essential that the X-ray source is able to deliver X-rays that arecontinuously tuneable over a certain energy range. The combination of theserequirements are met by facilities of high technical effort and complexity, socalled synchrotrons, where the primary research covered in this thesis wasconducted. The general synchrotron components necessary for XAS, RIXS,and many other core-hole spectroscopy techniques are the injector, storagering, insertion device, beamline, and endstation.
The injector creates tightly compressed bunches of electrons that are broughtto relativistic speeds in a linear accelerator (linac). In some facilities thelinac is connected to a so called booster ring in which the relativistic elec-tron bunches are further accelerated before injection into the storage ring. Thestorage ring provides a well defined path for the relativistic electron bunches,which is torus shaped in first approximation, but features both curved andstraight magnetic sections. The kinetic energy lost in form of emitted radi-
22
ation due to the radial acceleration in the magnetic sections is constantly re-stored into the bunches by klystrons precisely tuned to the right frequency toensure a stable trajectory without diminishing radius.
The mentioned magnetic sections are part of the so called insertion devices.These magnetic structures can be categorized in three general types: bendingmagnets, undulators, and wigglers. In bending magnets the electrons are ra-dially accelerated which causes them to emit radiation sweeping tangentiallyfrom their curved trajectory. Due to the relativistic speed, the emitted radiationis Lorentz-compressed to a narrow cone as seen from the laboratory frame ofreference. This radiation cone features an intensity which is distributed overa relatively broad solid angle as well as energy spectrum and thus is not suit-able for RIXS. Much higher brightness can be achieved by using undulators,which are placed in the linear sections of the storage ring and consist of anarray of magnets with alternating polarity and relatively low field strength.The electron bunches transversally oscillate with a small amplitude and theresulting radiation cone features a much smaller solid angle as well as spectralbandwidth as in the case of bending magnets or wigglers. In order to optimizephoton flux for certain energy regions the gap between the upper and lowermagnet arrays can be adjusted. All experiments in this work were performedat beamlines that utilize undulator radiation.
Beamlines are a series of ultra-high vacuum tubing that connect precisionoptical components for guiding, monochromatizing, and focusing the X-raysfrom the insertion device to the endstation. Soft X-ray beamlines typicallyuse spherical grating monochromators to further define the incoming photonenergy. A gold mesh is placed in the refocusing mirror directs the X-raysonto the sample with an optimal beam spot size and in order to measure theincoming intensity a is Finally the sample can be mounted in the endstation,which in the case of soft X-rays also needs to be a high vacuum chamber. Thischamber houses a manipulator onto which the sample(s) can be mounted andwhich is used to adjust the measuring geometry like e. g. the angle of inci-dence. The endstation is also connected to the various spectroscopic analyzerelements, which for the cases of XAS and RIXS will be briefly discussed inthe following sections.
2.3.2 X-ray absorption spectroscopyX-ray absorption spectroscopy (XAS) is a nondestructive method to probethe electronic, chemical, and crystal structure of solids and also fluids. It re-lies on the access to intense energy-tunable monochromatic X-rays, which areprovided by synchrotron facilities. A schematic illustration of the underlyingabsorption processes utilized in this technique can be seen in Figure 2.4 (a)and (b) and are described in more detail in section 2.2. This section focusesmore on the technical aspects of XAS.
23
main-edge
pre-edge
EXAFS
EEE
I(E)
(a) (b) (c)
NEXAFS
post-edge
Figure 2.4. Photon absorption and core electron excitation into an unoccupied energylevel of a partially filled valence band (a) and into a continuum state leaving behind acore-ionized intermediate state (b). A typical O K-edge XAS spectrum of a transitionmetal oxide (c). The situations in (a) and (b) give rise to intensity in the pre andpost-edge energy region, respectively.
There are several different detection methods, referred to as modes, that canbe utilized in XAS. One method to directly detect the first-order optical pro-cess that is XAS is transmission mode, where X-rays are transmitted throughand thereby attenuated by a sample. The ratio between outgoing and incomingintensity is proportional to the reciprocal exponential of the absorption coef-ficient times the sample thickness. This absorption coefficient is proportionalto the cross-section described by Fermis golden rule, (2.5), and increases dras-tically once the incident photon energy reaches the binding energy of a corelevel. This sudden increase is called absorption edge and the specific core leveldetermines the name of the edge, e. g. th K-edge stems from the excitation ofa 1s core electron and the L1-, L2-, and L3-edges stem from the 2s, 2p1/2, and2p3/2 levels respectively. As consequence of the aforementioned dipole selec-tion rules (Δl =±1, etc.) K- and L-edges are heavily dominated by transitionsinto orbitals of p and d character, respectively. As a result, the area under anabsorption spectrum maps out the partial density of unoccupied states abovethe Fermi level of a material. Because of the relatively short attenuation lengthof soft X-rays, the direct observation of XAS in transition mode is restricted tovery thin samples and is thus experimentally unpractical for common batterysamples. The more suitable indirect detection modes that rely on (secondary)deexcitation processes that happen after absorption. These processes can becategorized into radiative an non-radiative deexcitation channels. In the lattercategory the core-hole excited system relaxes back into ground state (GS) via
24
reoccupation of the core level with an electron from a higher energy level. Thisis accompanied by the emission of another, so called Auger electron remov-ing all excess energy from the system with it. The high energy Auger electronin turn scatters inelastically and initiates a cascade of lower energy secondaryelectrons that can leave the sample altogether. The number of created coreholes is directly proportional to a drain current that replaces the lost photo-electrons, which can be detected by a nanoampere-precision ammeter. Thisdetection mode is called total electron yield (TEY) and is mostly sensitive tothe sample volume about 2 to 10 nm [17], [20], [21] below the surface becauseof the short attenuation length of scattered electrons.
Another indirect detection mode is called total fluorescence yield (TFY).The underlying mechanism is the second-order optical XES process, see sec-tion 2.2, where the core-excited system relaxes back into GS via the transitionof a valence level electron into the empty core level orbital. This is accompa-nied by the spontaneous emission of a photon carrying the energy difference ofthe two energy levels. The emitted photons can be detected via a photodiodeor a Channeltron. The attenuation length of photons is longer than that of elec-trons so the resulting probing depth of up to 100nm gives information aboutthe bulk material of a sample and is thus complementary to the more surfacesensitive TEY mode. The bulk sensitivity can be varied by choosing the angleof incidence: X-rays coming in at normal incidence can probe deeper volumesof the sample than those with more grazing incidence.
By using an energy resolved detector, it is possible to restrict the sensitivityto a certain range of emission energies. In this case the detection mode iscalled partial fluorescence yield (PFY). It has been shown recently that it iseven possible to use the non resonant emission (NXES) region of a lower lyingabsorption edge from the same element or different element in the compound[22]. Because the absorption cross section of a lower lying edge decreasesin a proportional fashion With increasing absorption at the resonant edge onecan simply integrate over the lower NXES region for each incident photonenergy and then plot the inverse of that against the incident energy to gainan absorption spectrum that is free of the artifacts that might otherwise beintroduced due the so called self-absorption effect [23]. This mode is termedinverse partial fluorescence yield (IPFY).
A typical absorption spectrum measured at the O K-edge of a transitionmetal oxide can be seen in Figure 2.4 (c). XAS spectra can be divided intothree general energy regions: The pre-edge resulting from core-electron exci-tation into unoccupied states close to the Fermi level, the main-edge probinghigher lying quasi-bound states, and the post-edge region resulting from exci-tations above the ionization threshold, i. e. a photoionization or XPS process.In addition to the fundamental light matter interaction process, XAS can beseen as an umbrella term for more specific techniques like X-ray near-edgestructure (XANES) or the essentially equivalent near-edge X-ray absorption
25
fine structure (NEXAFS) 2. Both of these focus on the analysis of the pre- andmain-edge regions. Finally, extended X-ray absorption fine structure (EXAFS)is dedicated to the analysis of the post-edge region, which in some cases canextend up to 1000 ev above the absorption edge and thus can only be used ifhigher absorption edges of other elements are much higher in energy, whichoften is not the case in battery materials. The oscillations here depend uponthe type, position, and number of neighboring atoms and thus can deliver in-formation on the bond lengths and symmetries (crystal structure) of a materialif its elemental composition allows for it.
2.3.3 Resonant inelastic X-ray scatteringResonant inelastic X-ray scattering (RIXS) utilizes the RXES process3 in-troduced in section 2.2. The detection of a RIXS spectrum requires a spec-trometer with the following main components: emission source (slit), grazingincidence diffraction grating for wave length dispersion, and a position sensi-tive detector. In a way the recording of a RIXS spectrum is similar to that ofTFY-mode XAS with two important distinctions: (1) The emitted photons arenot summed up to one intensity data point, but instead are separated by theirdifferent emission energies and projected on a detector with separate channelsfor the different energies. (2) The incident photon energy must be held for alonger time to ensure proper statistics. However, sweeping through incidentphoton energies is possible too; the resulting series of RIXS spectra is referredto as RIXS map. Integration over each RIXS spectrum of the map or certainemission energy regions would yield a PFY-mode XAS spectrum.
RIXS is a complementary technique to XAS, where the density of occu-pied states rather than the unoccupied portion of the valence band is probed.RIXS is a photon-in/photon-out technique where photons inelastically scat-ter off matter and trigger certain intrinsic excitations by depositing parts oftheir energy and momentum. By measuring the change in energy and possiblymomentum and polarization of the scattered photons information about theseexcitations can be gathered. As is the case with XAS, elemental and orbitalspecificity can be achieved by tuning the energy of the incident photons to aspecific absorption edge. By further fine-tuning to the exact energy of a cer-tain absorption feature a resonant condition is created which enhances specifictransitions. This in turn enables to precisely probe e. g. various lattice sites de-fined by different chemical bonds with varying neighbors or oxidation statesof the selected atomic species [24].
2The difference between XANES and NEXAFS is more of historical than of conceptual char-acter. NEXAFS may be used relating to the analysis of light elements, e. g. C, N, or O, whileXANES may be more commonly used in relation to heavier elements. Because both light andheavier elements, i. e. O and TMs respectively, are analyzed in this thesis the technique is re-ferred to by the more general term XAS.3The terms RIXS and RXES are often used interchangeably.
26
Due to the long attenuation length of photons, RIXS is a bulk sensitivemethod that does not rely on super clean sample surfaces. The typical probingdepth for soft X-rays is in the order of 10 to 100 nm depending on the scatter-ing geometry. Due to the high brilliance of modern X-ray sources and becauseof the relatively high photon-matter interaction only small sample volumes arerequired. This enables the investigation of e. g. thin films if a grazing scatter-ing angle is chosen.
Eloss
I(Eloss)
(b)
Ground state2p63dn
L-edge XAS
(a)
dd-excitation
2p63dn*
"Elastic"scattering
2p63dn
Ligand CT
2p63dn+1LCore-hole
excited state
2p53dn+1
E
tInitial Intermediate Final
Figure 2.5. (a) Schematic view of the steps in a RIXS process shown in the total en-ergy representation with three principal de-excitation channels. The transition fromthe initial to the intermediate state is based in the XAS process. The various possibletransitions to the final state are of radiative nature. By using an energy resolved ana-lyzer the emitted photons give rise to a RIXS spectrum. A schematic illustration of atypical energy loss spectrum from a transition metal oxide is shown in (b). The colorcoded energy regions are associated with the different de-excitation channels from (a).
As depicted in Figure 2.5 (a) and described in section 2.2, equation (2.6)RIXS can be understood as a two-step process with a resonant absorption of aphoton of known incident energy and then emission of another photon of vary-ing energy. The energy difference between incoming and outgoing photons isdeposited into the system and responsible for a variety of possible excitations.The RIXS process leaves the system at an excited final state, but is chargeneutral. Transitions to unbound continuum states involving photo-ionizationsare by definition not considered as true RIXS processes, but are referred toas NXES instead. The latter do, however, show up in the spectra as so calledfluorescence contributions that are dispersive in energy loss scale with varyingincident photon energies.
Figure 2.5 (b) shows a simplified picture of a RIXS spectrum with the threeprincipal regions typically observed in a TM oxide. The color coded regionstem from three different principal de-excitation processes. The ground state
27
configuration is excited via an L-XAS process to a core-hole excited interme-diate state and subsequent de-excitation can commence via different channels:3dn → 2p53dn+1 → {
3dn;3dn∗;3dn+1L}, with n as the ground state occupa-
tion of the 3d band.The first channel with identical initial and final states 3dn is associated with
the so called elastic peak (red region in Figure 2.5 (b)) with zero or next to zeroenergy loss. Small energy loss features close to the elastic peak, which requirevery high resolution to be detected, are connected to excitations like excitons,polarons, phonons, and other depositions of momentum into the material.
In the second channel another 3d electron from the valence band annihilateswith the core hole, leaving the initial photoelectron reside in a reorganized, ex-cited 3dn∗ configuration. This process is called dd-excitation, which is oftenreferred to as a Raman-like feature, because of its non-dispersive behaviorin energy loss scale, which was previously known from Raman-spectroscopy.The first dd-excitation in n = 6 and higher TMs with octahedral local sym-metry is usually a reordering of the 3dn
[t62ge
n−6g
]to the 3dn∗
[t52ge
n−5g
]final
state. Thus, its detection enables an experimental way to quantify the energylevel separation between the t2g and eg sub-bands due to crystal field splitting.Subsequent dd-excitations are often associated with a spin-flip [25].
And finally the core hole can also be annihilated by a valence electron orig-inating from the ligand, creating a hole state L in the ligand band with a finalstate configuration of 3dn+1L. Due to the broad character of the O 2p in TMoxides this CT excitation can show a fluorescence-like behavior, i. e. a featurethat disperses with varying incident photon energy [26].
28
3. Batteries
In the course of this work, X-ray spectroscopy has been utilized to investigatethe electronic structure of energy related materials. The range of materialscovered in this thesis consists of promising, emerging or entirely novel cath-odes for lithium- or sodium-ion batteries. The following chapter presents themost important concepts related to the function of batteries. It also introducesthe novel and emerging materials investigated in this work along with theirrespective challenges and various research approaches to improve their func-tional aspects.
3.1 Working principleA battery is a device made up of one or more galvanic (or voltaic) electro-chemical cell(s), which stores potential energy in the form of chemical bondsand spontaneously releases the usable portion of it in form of electrical en-ergy once the necessary conditions for it are met. The general makeup ofany battery is based on two half-cells each including an electrode in contactwith an ion conducting electrolyte, which may or may not be the same in bothhalf-cells. One example for an electrochemical cell is schematically depictedin Figure 3.1. In order to prevent short circuits, the electrodes need to beelectrically separated often by means of an ion permeable separator materialbetween their front sides. During use, the electrical current is led over an ex-ternal circuit with an electrical load, which often connects to the backsides ofthe electrodes. The general working principle during discharge is based onspontaneous electrochemical reactions that are enabled by both the release ofelectrons (into the external circuit) and cations (into the electrolyte) from thenegative electrode, i. e. anode, and the simultaneous uptake of electrons (fromthe external circuit) and cations (from the electrolyte) by the positive electrode,i. e. cathode1. This process is driven by the difference in Gibbs free energy ofreactants and products in both half-cells, which include the sum of differentlattice cohesive energies in the electrodes as well as ionization and solvationenergies in the electrolytic environments [27]. In rechargeable, so called sec-ondary batteries the reverse reactions can be driven by a power source in the
1Technically the positive and negative electrodes should only be called cathode and anode dur-ing discharging and switch names during charging. However, for the sake of simplicity through-out this thesis the terms cathode and anode are used synonymously for the positive and negativeelectrode, respectively, irrespective of the state of charge.
29
external circuit in order to recreate conditions for the spontaneous reactionsto take place anew. Ideally this charging process restores both electrodes to astate without irreversible side-reactions degrading any of the battery parts.
+ -
Discharge
e-
Charge
e-
Charge
Discharge
ElectrolytePositive electrode Negative electrodeFigure 3.1. Illustration of the working principle of a rechargeable lithium- or sodium-ion battery.
The range of battery types is extensive, including a plethora of materials, allwith slightly to vastly different operating mechanisms and specifications. Allmaterials studied in this work are suitable for use as cathodes in lithium-ion orsodium-ion batteries (LIBs or SIBs). The most commonly used electrolytes inthese types of batteries are non-aqueous and can be either liquid (e. g. LiPF6or NaPF6 dissolved in organic carbonate-based solvents [28]) or solid (e. g.amorphous Li3PO4 [28] or Na3PS4 [29]). An air-tight casing is needed toenclose the cells due to the high reactivity of battery compounds with e. g.H2O. In case of a liquid electrolyte, an ion-permeable but electrically insulat-ing membrane is used as electrode separator [30]. The electrodes are typicallycomposites of an active material and functional additives. The former is ca-pable of accepting and releasing Li+ or Na+ cations and is often ground intomicron-sized particles in order to create a large reaction surface. Betweenthese particles small amounts of additives are mixed like carbon black for in-creasing the electric conductivity and a typically polymeric binder materialfor attaching the active particles together and connecting to them to a currentcollector. The current collectors (e. g. Al, Cu foil) are used for connection toan external circuit. A commonly used anode material with market maturity isgraphite, C6 [31].
The working principle of LIBs and SIBs is comparatively simple and basedon the transfer of Li+ or Na+ cations between the two electrodes during charg-
30
ing and discharging as depicted in Figure 3.1. There are many criteria govern-ing the performance of a battery with a hierarchy of importance determined bythe specific field of application. Among the most important specifications isthe energy that can be stored per unit weight and/or volume, i. e. gravimetricand volumetric energy density, sometimes referred to as specific energy andenergy density, respectively,
Eg,v =
T∫
0
EcellI
m,Vdt (3.1)
= EcellQg,v, (3.2)
where Ecell = Ecathode−Eanode is the cell voltage given by the difference be-tween electrode potentials of the cathode and anode, I is the current, and mor V are electrode mass or volume, respectively. The integral goes over thetime, t = T , it takes for the cell to reach the desired cut-off potential. Since thecell voltage often changes with the state of charge, one can approximate theenergy density by using the average operating voltage or discharge mid-pointvoltage Ecell. Qg,v is the charge stored per unit mass or volume of the activeelectrode materials, referred to as specific capacity or capacity density, whichis inversely proportional to the specific rate at which the battery is dischargedrelative to its maximum capacity, i. e. the C-rate.
The theoretical maximum limit for storable charge in any alkali metal (AM)based battery is the number of Li+ or Na+ cations either electrode can ac-commodate. Compared to anode materials, cathode materials can generallystore much lower charge per unit mass or volume and are therefore the bot-tleneck for increasing the energy density of batteries. All materials studiedin this work have high or positive electrode potentials vs. standard hydrogenelectrode (SHE) and are therefore suitable for use as the positive electrodesor cathodes. The chemical processes that the electric storage mechanism isbased on are called oxidation and reduction, referred to as the portmanteauredox reactions. These are explained in the following section.
3.1.1 Cationic versus anionic redoxTypically, the pristine or uncycled state of AM-ion batteries is the unchargedstate. The charge process encompasses the extraction of alkali cations fromthe cathode and incorporation into the anode. The underlying reactions on ei-ther electrode can be of different type, e. g. conversion, alloying, or insertion.All materials studied in this thesis are insertion or intercalation type cath-odes, meaning that the active material exhibits a crystal lattice with vacanciessuch as interlayers into which AM-ions can diffuse into and out of withoutmajor disruption to the host structure [32]. This reaction shall be exploredmore deeply with the help of a prominent example. One extensively studied
31
and largely commercially successful example of insertion cathodes is LiCoO2(LCO) [33], which is a member of the layered oxides (see section 3.2.1) featur-ing alternating layers of CoO2 and Li. This layered structure provides lithiumwith a 2D-pathway for (de)intercalation. In commercial batteries LCO is typi-cally combined with a graphite anode. The chemical reactions taking place inthe respective electrode, as well as the complete battery cell can be summedup as follows:
cathode : LiCoO2 −−→←−− Li1−xCoO2+ xLi++ xe− (3.3)
anode : xLi++ xe−+C6 −−→←−− LixC6 (3.4)
full cell : LiCoO2+C6 −−→←−− Li1−xCoO2+LixC6 (3.5)
Here the left-to-right reactions (→) are associated with charging and reactionsin the opposite direction (←) are associated with discharging of the battery.Removal and insertion of the positively charged Li+ ions from or into eitherelectrode must always be charge balanced by the simultaneous removal or in-sertion of an equal number of negative charges. These charges are ideallyelectrons, e– , that can be used in the electrical circuit. However, generally un-favorable side reactions, involving e. g. the irreversible removal of highly reac-tive O– from the cathode lattice, which often leads the production of gaseousspecies, can occur as well.
Focusing on the cathode material during charging, the delithiation of thematerial is accompanied by the oxidation of Co cations, where an electronclose to the Fermi level is removed from their 3d-bands with each extractedLi+, successively oxidizing each cobalt cation from its initial Co3+- to its Co4+-state. Reinsertion of Li+ during discharging then reduces the Co4+-speciesback to the initial Co3+-state and the overall-process is called cationic redox.
With increasing removal of lithium from Li1– xCoO2 during charging, alsoreferred to as delithiation, the active cathode material undergoes a series ofphase transitions that are reversible only for x ≤ 0.5. Further delithiationor overcharging introduces irreversible structural changes that severely lowerspecific capacity, effectively limiting the operation of LCO cathodes to ap-proximately half of its theoretical specific capacity and a cut-off potential of4.2V vs. Li/Li+ [33]. Research into the origins of these irreversible changesis extensive, results include the observation of Co dissolution into the elec-trolyte as well as lattice O loss. One controversially discussed proposition toexplain further delithiation despite the loss of Co redox centers due to disso-lution was the onset of a second redox process centered at the oxygen anions[34], [35]. This explanation is based on the theoretical framework introducedearlier by JEAN ROUXEL for highly covalent chalcogenides [36] that predictscharge transfer from anions to metal sites once a crystal split metal 3d-subbandenergetically shifts down into the anion sp-band, see Figure 3.2.
32
E
N(E)
E
N(E)
EF
EF
e-
e-e-
O 2p
M 3d
M--O
(a) (b)
Figure 3.2. Schematic band diagram illustrating the anionic redox process.
This is called anionic redox and is often associated with a degradation ofcathode materials due to dimerization of anions, forming e. g. O2
2– , that maylead to irreversible capacity fade. However, if irreversible phase transitionscan be avoided in the presence of anionic redox, it can lead to an access ofotherwise untapped capacity as well as higher operation potentials. This ideabecomes especially relevant in the context of lithium-rich layered oxides sincea large portion of their capacity cannot be explained by charge compensationvia cationic redox alone. This is a central theme throughout the work of thisthesis.
3.2 Emerging cathode materialsThe range of different battery materials and research strategies for their im-provement is truly extensive and can appear confusing. The following sectionpresents short introductions to the specific materials studied in this thesis andbriefly reviews the advantages as well as scientific challenges of these novelbattery cathodes.
3.2.1 From layered to Li-rich layered transition metal oxidecathodes
The majority of cathode materials investigated in this thesis are part of the lay-ered oxide class. The general formula is AxMO2, with A representing an alkalimetal, x≤ 1, and M representing one or more (mainly transition) metal(s) pos-sibly in different oxidation states. The various crystal structures all have layers
33
of edge-sharing MO6 octahedra in common, in between which an alkali cationinterlayer is placed with different possible oxygen environments. The clas-sification of possible polymorphs can be indicated by a special notation [37]including a prefix made up of a letter representing the specific oxygen environ-ment of the alkali cation (O for octahedral, T for tetrahedral, or P for prismatic)followed by the number of metal layers necessary to achieve lattice periodicity.The aforementioned LCO is one notable cathode material in this class and canbe written as O3-LiCoO2 according to the notation introduced above, becauseits Li-ions are surrounded by oxygen octahedra and three layers of CoO2 areneeded to reproduce a periodic unit cell. This material was introduced as thefirst intercalation LIB cathode in 1980 [38], became widely used in portableelectronics after its commercialization by Sony in 1991, and remains a bench-mark material especially in portable electronics applications due to its highscores in the performance matrix containing gravimetric and volumetric ca-pacity density, operating potential, rate capability, as well as high-temperatureand cycle performance [33].
Li[Co]O2LCO: 150 mAhg-1
Co Li[NiMnCo]O2NMC: 180 mAhg-1
CoMnNi Li[LiNiMnCo]O2LrNMC: 270 mAhg-1
CoMnNiLi
Partially replaceCo by Ni and MnCo MnNi
Excess Li onTM sites
Li
Figure 3.3. Evolution of layered oxide cathode crystal structures.
In the pursuit of achieving higher energy as well as power densities andalso to replace rare elements with more abundant and cheaper ones, a widerange of spin-off materials from the established CoO2 cathode were investi-gated by means of completely or partially substituting Co with other metalsstarting with Mn, Ni (binary) or a combination of both (ternary) arriving at theLi[Ni1– y– zMnyCoz]O2 (NMC) compounds with varying reversible capacitiesof up to 180mAh g−1. The research groups of Thackery and Dahn succeededfirst in 2006 in the synthesis of NMC cathodes that incorporate Li not only inbetween the TM layers but also replacing TMs inside their layer, see Figure3.3, thus far exceeding the capacities for conventional NMC cathodes. Theselithium-rich NMC (LrNMC) materials can reach capacities of 270mAh g−1,which is a huge step forward from the benchmark 150mAh g−1 reached bythe established LCO batteries. This intriguing extra capacity, which cannotbe explained by cationic redox on the TM ions alone, comes with the disad-vantage of voltage decay upon repeated cycling. Naturally this poses a majorobstacle to commercialization and thus both effects have been at the center
34
of a lively scientific debate about their origin. Among proposed explanationsare migration of O2– and/or TM cations to the surface or into the Li layerand even over-oxidation of TMs. Research by Bruce’s group [39] resultedin strong evidence suggesting that a major portion of the extra capacity inLrNMC derives from reversible anionic redox of oxygen. The work of thisthesis directly connects to and solidifies these findings with a series of furtherexamples of anionic redox active cathode materials starting with the Co-freeLi[Li0.2Ni0.2Mn0.6]O2.
Layered Li2MnO3Another example of the LrNMC layered cathode materials investigated in thisthesis is the Ni- and Co-free end member Li2MnO3 (LrMO). It too exhibitsa rocksalt structure consisting of alternating layers of Li ions and Li and Mnions in a 1:2 ratio, so in the layered compounds notation it is stated as O3-Li[Li0.33Mn0.67]O2. LrMO has been considered electrochemically inactive,because further oxidation of its Mn4+ cations is energetically strongly sup-pressed due to the octahedral ligand site symmetry. However, it was discov-ered that after an initial delithiation to a high potential of 4.5V vs. Li/Li+ thematerial becomes electrochemically active [40]. Thus, the initial cycle hasbeen dubbed activation phase. The remarkable theoretical charging capac-ity of 460mAh g−1 and experimentally confirmed initial capacities of over300mAh g−1 [41] combined with rapid voltage and capacity fading have con-tributed to the strong scientific interest in LrMO as has the elusive reactionmechanism responsible for the electrochemical activity of LrMO.
The theoretically questionable oxidation to Mn5+ was experimentally ruledout early by XPS [42] and later again by XAS [43]–[45]. Alternative explana-tions encompass O loss from the lattice, possibly in combination with Li+-H+
exchange with decomposed electrolyte [41], structural conversion to a materialwith e. g. MnO2 phases [46] or spinel phases [47], and O hole state formation[48].
In this thesis XAS and RIXS has been used on LrMO thin film cathodesto further the knowledge of the underlying mechanisms that are at play in theactivation cycle.
3.2.2 Spinel LiNi0.5Mn1.5O4LiNi0.5Mn1.5O4 (LNMO) is an insertion material with spinel crystal structuremade up of metal oxide octahedra, MO6, and lithium cations in interconnectedtetrahedral oxygen environments, LiO4. The associated three-dimensional Li+
diffusion pathways enable high insertion- and extraction-rates (ionic conduc-tivity), which in combination with a sufficiently high electronic conductivitymakes this material suitable for high-power applications [49], [50]. The ma-terial is based on the spinel LiMn2O4 (LMO), which already reached mar-ket maturity e. g. in the electric vehicle sector [51]. Nickel substitution in
35
LiNiyMn2– yO4 is a strategy of addressing some of the problems associatedwith the end-member material LMO, which suffers from capacity fading dueto manganese dissolution attributed to under-coordinated surface Mn3+ cationsand comparatively low energy density [52]. As nickel substitution approachesy = 0.5, the Mn3+-to-Mn4+ ratio decreases and the operating voltage shiftsfrom 4.0V vs. Li/Li+ associated with a charge compensation via the Mn3+/4+
redox couple to a higher voltage plateau of 4.7V [53], [54]. This higher volt-age plateau increases the energy density from 440 to 686Wh kg−1 [55] and ishitherto associated with the Ni2+/3+ and Ni3+/4+redox couples [56].
The in itself favorable shift of electrochemical potential to 4.7V placesLNMO into the class of high-voltage battery materials at the drawback ofthe potential leaving the stability range of conventional electrolytes. Whilehalf-cell setups with Li metal anodes show satisfactory long-term stability,full cells with graphite anodes suffer from severe capacity fading especiallyat high temperatures, which is associated with Li ions being trapped or pas-sivated in the solid electrolyte interface (SEI) forming on the graphite anode[57]. The SEI formation process involves the dissolution of Mn2+ and Ni2+
ions from the cathode [58], similar to what is observed in LMO, and researchefforts have been directed at cathode surface modification by thin film coatingas well as cation doping [50]. Significant attention is also directed towardsoptimizing the crystal structure as LNMO can crystallize in varying degreesof disorder of its TM octahedra: from predominantly ordered with Ni2+ sur-rounded by six Mn4+ nearest-neighbors (space group P4332) to predominantlydisordered with a random distribution of both TM cations over their octahe-dral sites (space group Fd3m). This degree of disorder has an effect on theelectrochemical performance as well as cycle stability, which is explored aspart of this thesis.
3.2.3 Sodium-ion battery cathodesThe research into sodium-ion batteries (SIBs) started shortly after the discov-ery of LIB intercalation compounds in the 1980’s and initially proceeded inparallel with the LIB research. With the commercialization of LCO in the1990’s research activity of SIBs shifted almost completely away towards LIBsmainly for two reasons: SIBs exhibit lower energy densities for the same hostmaterial than their LIB counter part, e. g. NaCoO2 vs. LiCoO2, due the lowerelectrochemical potential and higher mass of Na compared to Li, and the in-sufficient intercalation of Na+ cations into conventional graphite on the anodeside [59].
The steadily rising demand for energy storage solutions in electric vehiclesand future projection scenarios by e. g. institutions like the International En-ergy Agency that place a large emphasis on electric mobility fueled specula-tion about possible future Li supply shortages and thus price hikes [60]. Con-
36
cerns like this make the much more abundant and thus currently and prospec-tively cheaper Na based SIBs an attractive alternative to LIBs. More impor-tantly, SIB research gained new momentum with e. g. the discovery of hard-carbons as suitable anode materials for Na intercalation [61] and recently Fe-based cathode materials have been found that exhibit energy densities compa-rable to those found in LIB cathodes [62]. Another advantage of Na next to itsabundance is that fact that, unlike Li, it does not form alloys with Al, whichenables the replacement of the heavier Cu as the charge collector material andthus reducing the overall weight of battery cells. While these developmentsincrease suitability of SIBs for portable electronics and electric vehicles, pos-sible future applications for SIBs include areas like mid-to-large-format en-ergy storage for load leveling intermittent energy supply infrastructures, wherevolumetric and gravimetric energy density are secondary concerns behind pro-duction costs.
Key SIB research activity is dedicated to the discovery of new electrodematerials exhibiting high energy densities and especially cycle life. Recently,sodiated layered transition metal oxides, phosphates and organic compoundshave been introduced as cathode materials for SIBs [63].
In the course of this thesis, a portion of the investigated cathode materi-als are SIB cathodes that also crystallize in the layered oxide framework andexhibit capacity predominantly linked to anionic redox of oxygen.
3.2.4 Single crystal thin film cathodesA wide variety of oxides, including layered and spinel structures can be syn-thesized with high precision of stoichiometry and crystallinity by means ofe. g. pulsed laser deposition. This line of material synthesis is explored es-pecially for all-solid-state batteries, but it can also serve as a source of high-quality, single-crystal samples to investigate the active materials in batteryelectrodes without the influence of any additives like binder and conductivematerials. However, because film-thicknesses are typically in the range oftens of nanometers, nanosize effects might play a role in the cycling behav-ior of these films that are otherwise a non-factor in compound materials withmicron-sized active particles [64], [65].
37
4. Summary and discussion of results
This section summarizes the most important results of all papers and manu-scripts that emerged from this thesis. While the focus lies mostly on my owninput encompassing the results from from XAS and RIXS, limiting to theseonly would take away important aspects of the main message of some papers.So where it is appropriate for the main message of the respective paper, Iwill also present key results that originated from the work of the respectivecollaborators to whom I am indebted to.
A large portion of the work has been done prior to the spectroscopy mea-surements at the various synchrotrons. My collaborators from the differentchemistry groups were responsible for synthesis of the various cathode ma-terials, incorporating these into battery cells for electrochemical cycling, andperforming a large number of additional experimental characterization tech-niques including e. g. X-ray diffraction (XRD), transmission electron micros-copy (TEM), inductively coupled plasma optical emission spectroscopy (ICP-OES), operando electrochemical mass spectrometry (OEMS) and many more.Battery cell assembly as well as disassembly was performed in an inert-gasenvironment (glovebox). The cathode materials were cycled different states ofcharge (SoCs), meaning that a series of electrodes with different levels lithi-ation or sodiation were prepared for direct comparison. After electrochemi-cal cycling the cathode materials were rinsed with dimethyl carbonate severaltimes in order to remove any salt residues. After drying, the cathodes werecut into suitable sizes and attached to strips of adhesive Cu tape. In order toavoid any air exposure prior to the spectroscopic measurements the sampleswere then transported in air-tight bags to the synchrotron facilities, where theywere typically unsealed in a glovebox and transported to the endstation in anair-tight suitcase that could be mounted directly onto the transfer chamber.
4.1 Lithium-ion battery investigations: spinel materialsDepending on the specific synthesis conditions, spinel LNMO crystallizeswith varying degrees of disorder regarding the lattice positions of its Mn andNi cations. The conditions for the formation of the highly disordered phasetypically also result in the incorporation of oxygen deficiencies. Both, thedegree of disorder [66] and oxygen deficiencies [67], affect long-term electro-chemical performance. A novel synthesis method enabled the production ofdisordered LNMO while avoiding the typical formation of oxygen vacancies
38
and the resulting material shows enhanced long-term cycling stability com-pared to the ordered material [68]. The question of the electronic origin ofthese differences mark the starting point of the research that culminated in Pa-pers I and II. In order to answer this, a wide variety of experimental methodswas used for electrochemical and structural characterization and, more specif-ically as part of this thesis, XAS and RIXS were employed to investigate theelectronic structure of the different materials. Material variations include twosets of oxygen non-deficient samples with low and high TM cation disorder,for the sake of simplicity denoted as TM ordered and TM disordered respec-tively, that were cycled to different states of lithiation at the first and tenthcycle.
Both XAS and RIXS measurements reveal overall very similar behaviorsof the TM ordered and disordered materials during the first lithiation cycle.Very little variation upon delithiation is seen in the Mn L-XAS pointing to adominant Mn4+ oxidation state throughout the whole cycle. Minor changescan be observed between the pristine and BoP spectra that are reversed uponrelithiation to EoD. This points to a redox of the Mn3+/4+ couple on only asmall number of Mn sites associated with the short voltage plateau at 4.2Vvs. Li/Li+ of the charge-discharge curve. The majority of redox activity iscentered around Ni with some contribution from O, evidenced by strong vari-ations in the in the Ni L-XAS and the pre-edge region of O K-XAS especiallybetween samples BoP to EoC that are associated with the high voltage plateauat 4.7V. So far, these observations line up with what has been reported before[56], [69]–[71]. However, O K- and Ni L3-RIXS show evidence for an activerole of oxygen in the redox process not previously discussed on this level ofdetail. The surprising nature of this discovery prompted us to focus Paper Ion the evolution of electronic structure of the ordered material during first cy-cle. A comparison of both materials after repeated cycling, including a widerarray of characterization methods is presented in Paper II. The subtle differ-ences between the TM ordered and disordered materials will be discussed aftera detailed summary of the findings of Paper I.
First cycle behavior of TM ordered LNMOSo for now we focus on the first cycle behavior of the TM ordered materialto see if we can answer the question regarding the character of the hole statesthat develop around Ni and O upon delithiation. The charge-discharge curveis shown in Figure 4.1 (a) along with the labels for the different SoCs: par-tially delithiated to beginning of plateau (BoP) and end of plateau (EoP), fullydelithiated to end of charge (EoC), and fully relithiated to end of discharge(EoD). The pristine material was not cycled. The unoccupied states with O2p-character of the valence and conduction band can be probed by means of OK-XAS. The O K-XAS spectra, shown in Figure 4.1 (b), exhibit a significant,reversible variation in the low-energy portion of the pre-edge region, which arehighlighted by light green (increase) and pink (decrease) shadings. A closer
39
inspection of the nature of this reversible shoulder is depicted in Figure 4.1 (c)showing the respective RIXS spectra at an incident photon energy of 528.5 eV.Increasing delithiation leads to an accumulation of spectral weight on the highenergy side the spectra highlighted by turquoise shading. Closer inspection ofO K-RIXS maps for varying incident photon energies reveal a Raman-like be-havior of the peaks of the pristine and EoD spectra marked by a vertical blackline (meaning they track the elastic peak with increasing incident energy). Onthe other hand the peak of the delithiated samples in the same region behavesfluorescence-like in the O K-RIXS maps (meaning disperse away from theelastic peak upon increasing incident energy). Additionally there is a strongenhancement upon delithiation in the yellow shaded region close to the elasticpeak, which can be ascribed to O pp-excitations across a small band gap ofsome tenth of an electron volt. This all points to a significant reorganizationof electronic structure upon delithiation.
(a)(b) (d)
(c)
Figure 4.1. Paper I – (a) Charge-discharge curve of TM ordered LNMO. Cathodeswere cycled to the depicted SoCs. (b) O K-XAS pre-edge region. (c) O K-RIXS atincident photon energy hνin = 528.5 eV. (d) Ni L3-XAS. Reference spectra for Ni2+
(NiF2), Ni3+ (K3NiF6), and Ni4+ (KNiIO6), respectively, are as reported by Wang etal. [72] with a respective compression factor of 0.75 and 0.5 for Ni3+ and Ni4+.
We now turn to the results from Ni L3-XAS. The spectra depicted in Figure4.1 (d) exhibit a significant and reversible evolution of the spectral shapes thatare usually attributed to Ni redox reactions of the Ni2+/3+ and Ni3+/4+ couples
40
[56], [70]. This mainly takes place over the charge-discharge curve voltageplateau at 4.7V vs. Li/Li+. However, it must be noted here that there aredistinct differences to the shapes of the Ni3+ and Ni4+ reference spectra offormally highly ionic character compounds K3NiF6 and KNiIO6 [72], respec-tively.
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Figure 4.2. Paper I – (a) Ni L3-RIXS maps of TM ordered LNMO. The incidentphoton energies range from 852.0 eV to 855.0 eV (left graph) and from 855.5 eV to857.0 eV (right graph) with increments of 0.5 eV. The approximate positions of reso-nant excitation for Ni2+, Ni3+, and Ni4+ are indicated on the right axes (bubbles). Allspectra are normalized to their respective maximum. (b) Pristine, EoP, and EoC RIXSspectra at incident photon energy hνin = 855.0 eV in direct comparison with literaturereferences. The brown Ni 3d8 (NiO) and yellow Ni 3d8Ln (NdNiO3) reference spectraare taken from Ghiringhelli et al. [25] and Bisogni et al. [73], respectively, and weremeasured at similar incident energies but higher resolution.
The reason for this apparent discrepancy can be elucidated by closer in-spection via Ni L3-RIXS. The two panels of Figure 4.2 (a) show completeRIXS maps for all SoC with incident photon energies spanning the wholeL3-absorption edge, i. e. the region framed by black dashed vertical lines inFigure 4.1 (d). The bubbles on the right incidence photon energy axis markthe energies of maximum intensity for the pristine, EoP, and EoC XAS spec-tra, respectively, where a resonance with the labeled Ni oxidation states wouldbe expected. The RIXS map of the pristine material exhibits two features at1.4 and 3.1 eV that are constant in energy loss over varying incident energies,
41
i. e. behave Raman-like. These are marked by black dashed vertical lines inFigure 4.2 (a). A significant evolution of the RIXS maps is observed upondelithiation, especially beyond BoP, characterized by the development of ahigh energy loss shoulder, which is already seen at the lowest incident photonenergy 852 eV, and the development of an asymmetric, broad RIXS spectrumwith a single peak that appears above 854.5 eV incident photon energies. Re-lithiation to EoD finally largely restores the pristine state again.
To reiterate the initially asked question: What is the nature of the hole statesforming upon delithiation? The RIXS maps of the lithiated materials (pristineand EoD) exhibit a shape that is overall very similar to results reported for theNi2+ compound NiO [25]. The Raman-features at 1.4 and 3.1 eV do appear inNiO as well, small shifts notwithstanding, and are attributed to dd-excitationsfrom a 3d8 GS configuration, which is to be expected from a Ni2+ oxidationstate compound. These features can still be seen in the maps of the delithiatedsamples at low incident energies, which is remarkable as oxidation to Ni3+ andNi4+ would entail a change of GS configurations that are dominated by 3d7
and 3d6, respectively, if the hole states would be centered in the Ni 3d bandand thus a strong modification of the RIXS spectra would be expected. Un-fortunately, there no Ni L-RIXS studies of highly ionic compounds with Ni3+
and Ni4+ oxidation states. However, L-RIXS studies of compounds with Co3+
[74] and Fe4+ [75] in octahedral site symmetry and thus with predominantly3d7 and 3d6 GS configurations, respectively, are available in the literature. Forboth these configurations a dd-excitation pattern is reported that deviates fromthat of delithiated LNMO, casting further doubt on the proposition of 3d-bandcentered hole states. Instead, we propose that Ni largely retains its 3d8 GS inthe delithiated LNMO and that the broadening of the RIXS spectra at low in-cident photon energies is due to the formation of and CT excitations to ligandhole states L. In other words, Ni 3d8Ln with n = 1 or 2 are the dominant GSconfigurations in delithiated LNMO.
To conclude the analysis of the RIXS maps we now consider the spectrawith incident photon energies above 854.5 eV. This region shows the most se-vere changes in spectral shape upon delithiation. Incidentally CT excitationsbecome strongly enhanced in this region, which can be seen by the more pro-nounced tail above 4 eV, which is seen the lithiated materials as well as in NiO[26]. This tail region is significantly enhanced upon delithiation. Additionallythe spectra of the fully delithiated sample EoC exhibits a Raman-like feature at2.4 eV, which is marked by a red dashed vertical line in Figure 4.2 (a). One in-dication leading to the identification of this feature to a transition from a 3d8Ln
GS comes from a study on the negative charge transfer nickelate NdNiO3 byBisogni et al. [73]. Figure 4.2 (b) shows a direct comparison of Ni L3-RIXSspectra of the three LNMO samples pristine, EoP, and EoCmeasured at 855 eVincident energy with spectra of NiO and NdNiO3. While the pristine LNMOspectrum shows a striking resemblance with the NiO spectrum, the delithiatedSoC sample spectra share more similarities with the NdNiO3 spectrum. For-
42
mally NdNiO3 is a Ni3+ compound and therefore loosely comparable to EoPLNMO. However, Bisogni et al. show evidence for disproportionation of Ni3d8Ln configurations with n = 0 and 2 alternatingly residing on neighboringlattice sites.
The similarities of delithiated LNMO to the negative CT compound NdNiO3are also reflected in the low energy loss features seen in the O K-RIXS, whichare highlighted by the yellow shading in Figure 4.1 (c). These likely originatefrom pp-excitations across a small band gap in the partially filled O 2p band,which is also seen in NdNiO3.
So in conclusion, the experimental evidence in the form of O K- and NiL-XAS in combination with RIXS presented in Paper I is pointing to an O 2pcharacter of the hole states created upon delithiation of LNMO.
A comparative study of TM ordered and disordered LNMOPaper II is a comparative study between the already discussed TM orderedLNMO and its TM disordered counterpart. It utilizes a previously introduced,novel synthesis method that avoids the formation of oxygen vacancies typi-cally associated with disordered LNMO. A large number of characterizationmethods, including Raman spectroscopy, XPS, and highly sensitive electro-chemical testing to study electrochemical performance differences and cath-ode degradation mechanics are employed. It is found that both materials ini-tially perform very similar, however, disordering helps to maintain electro-chemical performance over extended cycling. This is because the orderedmaterial develops a higher impedance after repeated cycling, which can beattributed to a thicker SEI film as evidenced by XPS.
Can we also draw a connection between the more stable long term perfor-mance of the disordered material to the results from the electronic structureinvestigations via XAS and RIXS?
Figure 4.3 (a) shows the evolution of Ni L3-XAS of the TM ordered and dis-ordered (labeled TMord and TMdis, respectively) in direct comparison. Theoverall behavior is very similar. Two striking differences are seen here. Firstly,the comparison between the 1-EoP samples of the first cycle shows that Niions in the disordered material reach their highest oxidation state earlier in thedelithiation cycle compared to the ordered material. Secondly, the compar-ison of fully delithiated samples of the eleventh cycle (11-EoC) reveals thatthe spectrum of the disordered material develops a bigger ratio between thehigh energy feature at 855.5 eV and the first feature at 853.0 eV. The spectralshape of the ordered material remains largely unchanged after repeated cy-cling. Turning to the Ni L3-RIXS, the materials show two different changesupon repeated delithiation that are most pronounced at the incident photonenergy of 854.5 eV depicted in Figure 4.3 (b). The disordered material losesintensity at the encircled region closed to the elastic line. This is not seen inthe disordered material. Instead the spectral weight shifts to the low energytail, which is associated with CT excitations from the surrounding ligand.
43
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Figure 4.3. Paper II – (a) PFY mode Ni L3-XAS of cycled LNMO electrodes atdifferent SoCs (numerals in front of the labels denote the cycle number). (b) Ni L3-RIXS of fully delithiated samples, showing the changes after repeated cycling in bothmaterials. The vertical dashed line marks the incident photon energy; normalized tothe respective maximum. (c) O K-edge RIXS with incident photon energy of 528.5 eV,showing the changes after repeated cycling in both materials (upper spectra). Thelower spectra compare the 1-EoP SoC of the TM ordered and disordered materialsdirectly. O K-RIXS spectra are normalized for matching high energy flanks.
The latter is connected to increased hole states on O [26], [76], so we nowturn our description to the O K-edge. There are no noteworthy differences inthe XAS spectra. Both materials show the same reversible development of alow energy shoulder at the pre-edge upon delithiation. Figure 4.3 (c) showsselected comparisons between O K-RIXS spectra using an incident energy of528.5 eV. This probes excitations into the developing pre-edge shoulder whereone could expect the strongest differences. A subtle difference can be observedin the elastic peak region of the 1-EoP spectra depicted at the bottom, whichare marked by arrows. The asymmetry of the ordered material’s elastic peakindicates more occupied states within about 1 eV of the Fermi level comparedto the disordered material. This coincides with the observation in the Ni XAS,in which we observe a more advanced Ni oxidation in the disordered materialat this particular SOC. As seen above, this effect equalizes when reaching theend of the first charge (1-EoC) yet it has a long term effect on the material (11-EoC). This kind of long term effect is also seen in the upper O K-RIXS spectra
44
comparing 1-EoC and 11-EoC, where the ordered material clearly undergoesa development (see arrows pointing out the difference in the ratio between thetop of the band and the main band) whereas the disordered material shows lessspectral change. Note also that the elastic peak of the ordered material showsa similar asymmetry at both 1-EoC and 11-EoC as at 1-EoP.
To summarize the XAS and RIXS results it is observed that the tendency ofthe disordered material to reach a higher degree of Ni oxidation at an earlierSoC in the first cycle and to have a more stable long term electronic structurein the delithiated state than the ordered material, which acquires holes on theO ions instead of on the Ni ions. This could all affect the electrochemicalperformance directly or through additional side reactions. The latter couldalso result in the thicker film formation observed by XPS.
4.2 Lithium-ion battery investigations: Li-rich materialsThe Li-rich cathode materials investigated in this thesis include thin films ofthe layered LrNMC end-member material Li2MnO3 in two different compo-sitions, composite material of layered, Co-free Li[Li0.2Ni0.2Mn0.6]O2, as wellas a composite of the polyanionic Li1.9Mn0.95O2.05F0.95 with disordered rock-salt structure. The results presented here focus on the development of theelectronic structure during the first lithiation cycle and highlight the roles thatoxygen plays in the redox of these cathodes.
4.2.1 The effect of excess Li on covalency in Li2MnO3 thin filmsIn order to elucidate the reason for the remarkably different long-term cyclingbehaviors observed for layered, Li-rich Li2MnO3 (LrMO) thin film cathodeswith different initial Li/Mn-ratios [77], rLi/Mn, we have studied the evolutionof the electronic structure of films with two slightly different stoichiometriesduring the initial lithiation cycle. For this purpose, LrMO thin film cathodeswith near-perfectly stoichiometric as well as over-stoichiometric Li/Mn-ratios,rLi/Mn = 2.07 and 2.26 respectively, were electrochemically cycled to fullyde- and relithiated SoC in half-cell setups with Li-metal anodes. The thinfilm cathodes can be epitaxially grown with high control of composition, film-thickness, and crystallinity by means of pulsed laser deposition [64] and en-able the study of the active cathode material in absence of any conductor orbinder additives.
As mentioned in section 3.2.1, LrMO undergoes a so called activation pro-cess during the initial lithiation cycle, in which the material exhibits irre-versible structural changes that set the basis for successive, more reversiblecycling with considerably high capacities. Decades of extensive research ef-forts produced a wide array of plausible, yet competing explanations of theinitial redox mechanism and our study adds experimental evidence about the
45
electronic structure of the Mn–O bonds as well as the role of excess Li in theTM-layer.
In our study, we employed XAS and RIXS over the O K- as well as the MnL-edge. Figure 4.4 (a) shows Mn L3-edge XAS of films with stoichiometric(rLi/Mn = 2.07, upper spectra) and over-stoichiometric (rLi/Mn = 2.26, middleof plot) compositions at different SoC (pristine; delithiated, EoC; relithiated,EoD). The blue dashed lines are simulated spectra from linear combinationsof the reference spectra from polycrystalline manganese oxides: Mn2+ (MnO),Mn3+ (Mn2O3), and Mn4+ (Li2MnO3). The assumed percentages of referencespectra and therefore Mn oxidation states are given in brackets above eachrespective spectrum. While the pristine materials are dominated by spectralfingerprints of Mn4+ , de- and relithiation adds a significant portion of thelower two oxidation states to both compositions.
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Figure 4.4. Paper III – (a) TEY-signal Mn L3-edge XAS of Li2MnO3 (LrMO) epi-taxial thin films in two compositions and at different SoC (pristine; delithiated, EoC;relithiated, EoD); reference spectra as reported by Qiao et al. [78]. (b) TEY-signal OK-edge XAS. (c) Overview of key-findings from XAS.
Selected observations from the XAS experiments are condensed in Figure4.4 (c). The left axis (black traces) depicts the average intensity over theO K-pre-edge from 527.5 to 535.0 eV and the right axis (blue traces) repre-sents the average manganese oxidation state gathered from the relative mixof Mn4+-, Mn3+-, and Mn2+-reference spectra necessary for the best fit sim-
46
ulations of Mn L-edge spectra. The overall electronic structural trends dis-played by both film compositions, i. e. stoichiometries, can be summarizedas a continuing reduction of the initially dominating Mn4+ cations over thewhole first cycle accompanied by an increase and subsequent decrease of holestates in hybridized O2p–Mn3d orbitals, upon de- and relithiation respectively.This means that initial delithiation is in fact not charge balanced by the usualcationic redox, which was to be expected as oxidation states above Mn4+ areenergetically unfavorable in an octahedral oxygen environment. Instead, anaverage of approximately 0.4 e– per unit cell localize on and partially reducemanganese to Mn3+ and even Mn2+, while the majority remains in the Mn4+
state upon delithiation. The actual distribution of Mn oxidation states based onbest fit simulations yields stronger differences between the two compositions,see brackets in Figure 4.4 (a).
Regardless of the detailed and apparently rather intricate manganese ox-idation state distribution, its approximate average reduction by 0.4 e– upondelithiation requires a different source for electrons to compensate the posi-tive charge carried away by 2Li+ ions per unit cell. The most likely source isthe O 2p-band, which accumulates hole states as seen by the strong increaseof O K-pre-edge intensity.
Figure 4.5 shows O K-RIXS of the over-stoichiometric films at incidentphoton energies close to the Fermi level. The vertical dashed line at 525.7 eVmarks the centroid of the main emission band that is nearly stationary againstvarying incident photon energies, which therefore can be assigned to band-like states of O 2p character. Marked by blue arrows are features that aredispersive with incident photon energy with a loss energy of 2.4 eV versus theelastic peaks marked by black arrows. This Raman-like feature correspondsto an excitonic ligand-to-metal charge transfer transition leaving behind a dd-excited final state. In the pristine film spectra this feature is connected to theelastic peak via a pronounced shoulder. This shoulder intensity is lost upondelithiation and does not reappear upon relithiation implying an irreversibleloss of occupied states close to the Fermi level induced by the lithiation cycle.All observation so far can be made for both film stoichiometries (including thestoichiometric films not shown here).
Differences start with after delithiation. While the EoC spectra of the over-stoichiometric film exhibits a spectral broadening especially on the low emis-sion energy side, no such broadening is observed in the EoC spectra of thestoichiometric film. Another observation unique to the over-stoichiometricfilm is a pronounced narrowing of the emission band upon relithiation to EoD,which leads to aforementioned excitonic peak to stick out more prominentlycompared to the same feature in the EoD spectra of the stoichiometric film.Additionally, a low energy loss feature close appears for incident photon en-ergies close to the Fermi level, which is marked by blue circles in Figure 4.5.
47
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Figure 4.5. Paper III – O K-edge RIXS maps of over-stoichiometric LrMO thin filmswith rLi/Mn = 2.26 comparing the available SoC.
Figure 4.6. Paper III – Selected Mn L3-edge RIXS maps of over-stoichiometric LrMOthin films with rLi/Mn = 2.26 at de- and relithiated SoC, upper and lower graph respec-tively. The pie charts on the left show the Mn oxidation state distributions as extractedfrom the best-fit simulations of the Mn L-XAS TEY spectra.
48
Figure 4.6 shows the Mn L3-edge RIXS maps of the de- and relithiatedSoC, upper and lower panel respectively, of the over-stoichiometric films. Therespective distribution of Mn oxidation states, stemming from the best-fit sim-ulations shown in Figure 4.7, is presented as pie charts on the left side. Theselection of these particular two maps highlights the strong decrease upon re-lithiation of intensity relative to the respective maximum in the areas markedby the red polygon and ellipse. The features above 5 eV energy loss high-lighted here are associated with charge transfer transitions across the Mn–Obonds, which are revealed to be severely weakened over the course of thewhole initial cycle in the over-stoichiometric composition. This cannot be ob-served in the stoichiometric composition, so we conclude that the latter hasmore stable Mn–O bonds compared to the over-stoichiometric films.
In summary, we observe that both manganese and oxygen ions participatein the activation process of LrMO as a cathode material. However, the detailedmechanisms depend strongly on their Li/Mn-ratio. While the O K-XAS showsthat the stoichiometric material undergoes moderate changes in the conduc-tion band, the over-stoichiometric material shows a more pronounced changeof occupancy between de- and relithiated states. On cursory inspection, MnL-XAS seems to show similar behavior for the films but more detailed fit-ting analysis shows that, depending on the Li/Mn-ratio, different paths arefollowed during cycling. Importantly, we find by using Mn L3-RIXS that theover-stoichiometric cathode material is prone to a rapid loss of Mn–O hy-bridization whereas the hybridization is significantly more stable in the stoi-chiometric material. This is concomitant with a restructuring of the O-bondingat the top of the valence band (revealed by O K-RIXS) at the initial delithiationof both materials. While this leaves the stoichiometric material passive uponrelithiation, the over-stoichiometric material shows further restructuring at thetop of the oxygen band. O K-XAS reveals that holes are formed in the con-duction band of both materials. However, in contrast to the localized speciesthat forms in the iso-structural NMC-materials, the created holes in the oxygenband of LrMO, regardless of Li/Mn-ratio, have delocalized character.
4.2.2 The role of oxygen in the lithiation ofLi[Li0.2Ni0.2Mn0.6]O2
Figure 4.7 (a) shows O K-edge XAS of Li[Li0.2Ni0.2Mn0.6]O2 for differentpoints in the charge curve, which are shown in the upper panel of (b). Since inO K-XAS O 1s electrons are excited into empty states above the Fermi level,the spectra effectively represent the density of unoccupied states. The lowerpanel of (b) shows the integrated intensity under the shaded pre-peak regionin (a), which corresponds mainly to hybridized O 2p and Mn 3d states. Uponinitial charging (i to ii) Ni is oxidized from Ni2+ to Ni3+. On the plateau be-tween ii and iii further charging can be attributed to oxidation of O, which can
49
be seen by the increase in relative intensity of the pre-peak region as high-lighted in the lower panel of (b). This process is reversed during discharge (ivto v) as the empty states disappear again.
Figure 4.7 (c) shows O K RIXS excited at the energy correspondent to theenergy of the aforementioned extra hole states at 531.5eV. No elastic peakis seen for the pristine sample. The main peak corresponds to nearly pureO 2p states, with the exception of the high energy shoulder correspondingto hybridized states most likely with Mn 3d. Between iii and iv an elasticpeak and an extra feature in the main band appear. This has a number ofimplications. Firstly, a new O species appears here. Secondly, due to therelative prominence of the elastic peak we can conclude the occurrence of alocalized state. As seen as well in XAS, this state disappears again.
Figure 4.7. Paper IV – (a) TFY mode O K-edge XAS of different states of charge.Upper panel of (b) shows the charge-discharge curve with lower case Roman numeralsmarking the points at which samples were harvested. The lower panel of (b) showsthe evolution of the relative integrated O K-pre-edge intensity (shaded region in (a)).(c) O K-edge RIXS at an incident photon energy of hνin = 531.5 eV.
Combining these results with findings from further techniques the paperconcludes that upon charging across the ii-iii plateau 0.43 electrons per for-mula unit are extracted from the cathode due to a reversible oxidation of O2−.This leaves localized and relatively ionic O–(Mn4+/Li+) bonds and as a resultunstable O ions in the lattice. This irreversible O loss accounts for the extrac-tion of a further 0.06 electrons per formula unit. In terms of increasing thepractical charge capacity of O redox materials like Li[Li0.2Ni0.2Mn0.6]O2 it isimportant to find a way to shift the balance between the reversible O redoxchemistry and the irreversible O loss towards the former.
4.2.3 Localized oxygen hole states in novel, oxygen redoxactive, disordered Li1.9Mn0.95O2.05F0.95
Section 4.1 already showed an example in the form of spinel LNMO of amaterial without a layered structure, that nonetheless shows evidence for an-
50
ionic redox involvement and where TM disorder shows beneficial effects onelectrochemical performance. Paper V showcases a Li-rich material in thenon-layered, disordered rocksalt structure, Li1.9Mn0.95O2.05F0.95 (LMOF). Aspecial synthesis method yields a crystallization in the Fm3m space group,depicted in 4.8 (a). While Li and Mn cations are distributed over the 4a sites,O and F anions occupy the 4b sites of the disordered rocksalt lattice.
(c)
(a)
Li
Mn
O
F
(b)
Volatagevs.Li/Li+[V]
Capacity [mAhg1]
Photon energy [eV]
(e)
Photon energy [eV]
Intensity[arb.units]
Photon energy [eV]
(d)(c)
Figure 4.8. Paper V – (a) Schematic crystal structure of Li1.9Mn0.95O2.05F0.95. (b)Galvanostatic first cycle charge-discharge profile. Color coded circles mark positionswhere samples were harvested. (c) Mn L-edge XAS of all samples, including refer-ence spectra. (d) O K-edge XAS. (d) O K-edge RIXS at an incident photon energy ofhνin = 531.5 eV.
This novel cathode material exhibits a promising electrochemical perfor-mance featuring a C/10 first cycle charge and discharge capacity of 291 and283mAh g−1, respectively, at an average potential of 3.4V vs. Li/Li+, see Fig-ure 4.8 (b). The discharge capacity retention is approximately 60% after 50cycles. The charge process can be divided into two regions characterized by atransition at about 4.2V between different slopes in the curve, i. e. the rate ofcharge passed with change of voltage, dQ/dV . The average oxidation state ofMn in the pristine material was determined by iodometric titration to approxi-mately Mn3.3+, which agrees well with nominal oxidation state considerationsas per stoichiometry assuming Li+, F– , and O2– .
51
This is further corroborated by the Mn L-edge XAS spectrum of the pristinesample, shown in Figure 4.8 (c), which exhibits dominant Mn3+ features. Thespectra shown here are IPFY signals recorded with a photon energy sensitivedetector that enabled a bulk sensitive measurement free of the branching ratioeffect often encountered in TFY measurements at the Mn L3-edge of oxides.The delithiated sample (4.8V, green trace) shows the decrease of the Mn3+
feature and the increase of the typical Mn4+ double feature in the L3-edgeas well as a shift of the L2-edge centroid to a higher energy coinciding withthat of the MnO2 reference sample. The oxidation to Mn4+ accounts for abit less than 0.7 e– per unit cell, however, the amount per unit cell of Li+
extracted at 4.8V is a little under 1.3 according to ICP-OES. The resultingcharge extracted beyond cationic oxidation of 0.6 e– per unit cell cannot beexplained by O loss either as OEMS measurements show only evidence forCO2 evolution starting at potentials above the stability limit of the electrolyte,but virtually no O2 evolution. So there is no evidence for significant O lossfrom the active material.
On the other hand, turning to the O K-edge XAS, shown in Figure 4.8 (d),we observe a clear increase in the pre-edge intensity upon delithiation. More-over, the O K-edge RIXS reveals the exact same behavior already observed forLrNMO in section 4.2.2 (Paper IV), namely the emergence of a new emissionpeak at 525 eV when the incident photon energy is tuned to 531.5 eV. Uponrelithiation the absorption pre-edge intensity decreases again and the 525 eVRIXS emission peak disappears as well. This again shows that O oxidationtakes place in LMOF and is responsible for the reversible capacity beyondwhat can be associated with cationic redox of Mn.
52
4.3 Sodium-ion battery investigations: layered materialsIn the realm of sodium-ion batteries (SIBs), three different cathode materialswere investigated in the course of this thesis, which culminated in the PapersVI, VII, and VIII. Here again the role of oxygen redox and the connection tooxygen loss from the lattice are explored.
4.3.1 No oxygen loss in the highly anionic redox activeNa0.67[Mg0.28Mn0.72]O2 cathode
Paper investigates VI Na0.67[Mg0.28Mn0.72]O2 (NMMO), which features theP2-type layered structure with trigonal prismatic site symmetry NaO4 layersand alkali metal (AM) free layers of alternatingMgO6 andMnO6 octahedra. Aschematic illustration of the crystal structure is shown in Figure 4.11 (b). Thismaterial features a first cycle charge-discharge curve, see Figure 4.9 (a), that isvery similar in its shape to the one observed for LrNMO (Figure 4.7 (b)). It canbe divided into three principal regions associated with different reaction mech-anisms: the desodiation of approximately 0.14 Na+ ions per unit cell in regionI (samples (i), pristine to sample (iii), beginning of plateau) is associated withan oxidation of manganese, which changes from an average initial oxidationstate of Mn3.81+ in the pristine state (determined by field cooling magnetizationmeasurements) to Mn4+ at the beginning of the plateau. Region II (samples(iii) to (v), end of charge) exhibits a near constant voltage plateau at 4.2V vs.Na/Na+ with a spike in voltage to 4.5V in the last 15%. While the material isnot completely desodiated – based on passed charge considerations the com-position at (v) is Na0.14[Mg0.28Mn0.72]O2 – approximately 0.38 Na+ ions perunit cell are deintercalated over this region and here no further cationic oxi-dation is observed between samples (iii) and (v), as Mn K-pre-edge centroidsshow only marginal shifts relative to that of the tetravalent reference sample.So charge compensation could be due to oxidation of oxygen or oxygen anionloss from the lattice. The possibility of oxygen loss was carefully ruled outwith a variety of experimental methods including operando thermogravimet-ric analysis mass spectroscopy and nuclear magnetic resonance spectroscopy.Since the results from these methods show no evidence for oxygen loss fromthe lattice, the charge compensation mechanism must be based on electronhole creation on oxygen. Resodiation or discharging is summarized as regionIII (samples (iv) to (vii), end of discharge).
The corresponding OK-edge spectra for all samples are shown in Figure 4.9(b). Besides changes in the main-edge shape above 533 eV there is a strongmodification of the pre-edge in close vicinity above the Fermi level, which isrelated to empty, hybridized O 2p and Mn 3d states. The change in pre-edgeintensities is quantified by integration from 528 to 533 eV shown in Figure 4.9(c), which reveals a clear increase in hole states over region II and a refillingof states upon resodiation in region III. In addition to this oxygen reduction
53
during resodiation, some manganese sites are also reduced between samples(vi) and (vii) from an average of approximately Mn4+ to Mn3.7+. The resodia-tion capacity is estimated slightly higher than that of desodiation, with a totalamount of 0.6 e– charge passed per unit cell, out of which 0.2 and 0.4 e– arecompensated by Mn and O reduction, respectively.
Figure 4.9. Paper VI – (a) Galvanostatic first cycle charge-discharge profile ofNa0.67[Mg0.28Mn0.72]O2; lower case roman numerals label positions where sampleswere harvested. (b) TFY-signal O K-edge XAS of all samples. (c) Relative integratedpre-edge intensities. (d) O K-edge RIXS of all samples at an incident photon energyof hνin = 531.5 eV.
O K-RIXS spectra at a selected incident photon energy of hνin = 531.5 eVare shown in Figure 4.9 (d). Very similar to the results from LrNMO, weobserve a split of spectral weight appearing between the beginning and mid-dle of the plateau, (iii) and (iV), peaking at the end of charge, (v), and thendisappearing again towards middle of discharge, (vi). Here again, this the re-versible appearance of this emission peak at 523 eV is accompanied by thereversible increase of the elastic peak, which is not shown in Figure 4.9 (d). Inan upcoming publication the periodic, low energy loss tail of the elastic peaks,which seems to be connected to the formation of O2 molecules trapped insidethe lattice, is analyzed in detail.
One intriguing result from Paper VI is that in order to access reversible Oredox, no extra AM ions are needed in the TM layer, like is the case for Li-rich
54
materials. Another highly interesting result is the lack of oxygen loss from thelattice during the sodiation cycle, which has been seen in e. g. LrNMO of theprevious paper. This shows that O oxidation cannot be the reason for oxygenloss in anionic redox cathodes, which immediately raises the next question: Isthe suppression of oxygen loss in NMMO due to the lack of AM ions in theTM layer or could the additional bond to Mg2+ ions have a stabilizing effect onlattice oxygen? These questions are investigated in detail and with extensiveseries of experimental techniques in Paper VII.
4.3.2 The reason for oxygen loss in anionic redox active cathodesAs mentioned in the previous section, Paper VII addresses the question of whattriggers oxygen loss in anionic redox active layered cathode materials andwhy it is not seen in NMMO. For this purpose an electrochemically and struc-turally analogous material to NMMO was prepared: Na0.78[Li0.25Mn0.75]O2(NLMO). Figure 4.10 (a) shows a similar charge-discharge curve shape tothat of NMMO with more sloping “plateau” towards the 4.5V cut-off volt-age and slightly higher charge and lower discharge capacity, cf. Figure 4.9.The underlying redox mechanism for NLMO was again determined by MnK-edge XAS, which showed the same behavior with predominant oxidationstate of Mn4+ throughout the cycle and reversible, partial redox. The majorityof charge compensation in NLMO originates again from oxygen redox, whichis evidenced by O K-edge XAS featuring the same evolution of pre-edge in-tensity, both shown in Figure 4.10 (b) and (c). Additional evidence for ananalogous electrochemical behavior during the first sodiation cycle to NMMOis presented through the O K-edge RIXS spectra of Figure 4.10 (d). Again weobserve the same reversible increase of emission intensity upon charging seenin NMMO and LrNMO.
Turning to the crystal structure, a series of experimental methods were usedto arrive at the picture presented in Figure 4.11 including ICP-OES, as wellas powder X-ray and neutron diffraction. Both NLMO and NMMO exhibitthe P2-type layered structure with a difference in the stacking order of the TMlayers: While in NMMO the TM layers are aligned without offset, subsequentTM layers are offset against the c-axis in NLMO. The more important differ-ence for the underlying research question is the presence of an AM ion in theTM layer of NLMO. Here, Mg2+ is replaced by more mobile Li+ ions, whichtested by means of solid-state nuclear magnetic resonance indeed show diffu-sion from their initial TM layer sites into the alkali layer upon desodiation to4.5V.
A final commonality of NLMO and NMMO is the lack of any oxygen lossduring desodiation below a cut-off potential of 4.5V, which was determined byOEMS of 18O-labeled samples. This gives the first, negative result regardingthe origin of oxygen loss: The lack of AM ions or rather the lack of vacancies
55
Figure 4.10. Paper VII – (a) Galvanostatic charge-discharge curve ofNa0.78[Li0.25Mn0.75]O2; color-coded circles label positions where samples were har-vested. (b) TFY-signal O K-edge XAS of samples with different SoC. (c) Relativeintegrated O K-pre-edge intensities. (d) O K-edge RIXS at an incident photon energyof hνin = 531.5 eV.
Figure 4.11. Paper VII – Crystal structures of (a) P2-Na0.78[Li0.25Mn0.75]O2 (NLMO)and (b) P2-Na0.67[Mg0.28Mn0.72]O2 (NMMO) showing the different layers as well asthe honeycomb ordering in the TM layers.
56
that appear upon charging in the TM layer of NMMO is not the reason forits absence of oxygen loss. However, the situation is not directly comparableto that of delithiated Li-rich cathodes that show oxygen loss, because, despitethe migration of Li+ ions into the TM layers upon charging of NLMO, aninvestigation by ICP-OES showed that only Na+ leaves the cathode with noevidence for Li+ deintercalation.
To drive Li+ out of NLMO samples, potentiostatic charging at a step of5V was applied. This was done for NMMO as well and the concomitantOEMS now clearly showed oxygen loss for NLMO but still none for NMMO.Connecting this with the confirmation of delithiation on top of desodiation inNLMO by ICP-OES, the second, positive result regarding the initial questioncan be stated: Oxygen loss is triggered once the average coordination numberof oxygen falls below 3. While O in desodiated and delithiated NLMO iscoordinated by only 2Mn4+ and therefore loses the underbonded lattice O, theO in desodiated NMMO stays coordinated by 2Mn4+ and 1Mg2+, due to themuch lower mobility of the latter compared to Li+.
This ultimately shows that oxygen loss is not an inevitable consequence ofO-oxidation or charging to high potentials and implies the stabilization of lo-cally underbonded oxidized O species should be taken into consideration whenaiming at designing highly reversible anionic redox active cathode materials.
4.3.3 The charge compensation mechanism of theNa0.56Mg0.04Ni0.19Mn0.70O2 cathode
Another Mn-rich layered oxide SIB cathode with additional Ni and less Mgis investigated in Paper VIII. Na0.56Mg0.04Ni0.19Mn0.70O2 (NMNMO) crys-tallizes in the P2-type layered structure with both TMs, Mn and Ni, in anoctahedral local site symmetry surrounded by six oxygen anions and a smallamount of Mg doping in the TM layer.
Figure 4.12 (a) shows the charge-discharge curve of NMNMO and definesthe labels of for all samples at different SoCs. The evolution of the Mn L3-and L2-XAS spectra upon de- and relithiation is depicted in Figure 4.12 (b)and (c), respectively. The surface sensitive TEY spectrum of the pristine sam-ple exhibits strong spectral fingerprints of Mn4+ as well as an admixture con-tribution of Mn2+ with a strong feature at 639.7 eV and a weaker feature at641.1 eV, which means that the surface of the pristine material is dominatedby Mn4+ sites and features a smaller portion of sites with a Mn2+ oxidationstate. Throughout the whole desodiation process to 1-Ch: 4.5V the spectralshapes show minute changes, indicating that the surface Mn is not active incharge compensation during this part of the cycle. Conversely, upon resodi-ation to 1-D: 2.0V the spectrum changes dramatically to only exhibit a Mn2+
signature. Subsequent desodiation in the second cycle to 2-Ch: 4.5 V shows
57
a return of the Mn4+ signature with a slightly higher Mn2+ contribution thanobserved over the first cycle.
The bulk sensitive total fluorescence yield (TFY) signal is not presentedhere because the differential-fluorescence-yield effect at theMn L3-edge strong-ly distorts the spectrum. Instead, we can utilize the TFY at the L2-edge, whichis depicted as dashed traces in Figure 4.12. Though the shorter 2p1/2 corehole lifetime broadens the spectrum, making the determination of the oxida-tion state more difficult, we can draw some conclusions by the shift of thepeak centroid. Overall, the bulk sensitive TFY L2-spectra follows the behav-ior of the TEY L2-spectra. However, one striking difference is seen in thepristine spectrum with a centroid at a lower energy marked by a vertical linelabelled A. This line also coincides with the centroid of the resodiated samplespectrum 1-D: 2.0 V, suggesting that the bulk material starts out with a Mnoxidation state dominated by Mn2+ while its surface is strongly Mn4+. Uponcharging, the bulk rapidly (already at 3.5 V) attains a similar oxidation state asthe surface, i.e. a mixed Mn2+/4+ mixed oxidation state. Already upon desodi-ation to 1-Ch: 3.5 V the peak centroid shifts by 1.1 eV towards higher energies(line B).
Na0.56[Mg0.04Ni0.19Mn0.7]O2
Time [h]
Volta
ge [V
] vs.
Na/
Na+
(c)
Figure 4.12. Paper VIII – (a) Charge-discharge curve of Na0.56Mg0.04Ni0.19Mn0.70O2showing the respective SoCs of all samples. Mn L3- and L2-XAS, (b) and (c) re-spectively. Both edges are shown in TEY mode (solid traces) and the L2-edge isadditionally shown in TFY mode (dashed traces).
Figure 4.13 (a) shows TFY- and PFY-signal XAS of all samples over theO K-edge. In the pre-edge region spanning from 528.5 to 533.0 eV we ob-serve a small variance in the spectral shape connected to a gradual increaseof 2p character hole states during the first desodiation (bottom four spectra).Figure 4.13 (c) shows an analysis of the evolution of O hole state density.We integrated the respective pre-edge intensities of TFY, PFY, and TEY XASmodes at all SoCs and normalize it to that of the pristine material. We find thatboth the surface- and bulk-sensitive signals exhibit a similar trend, namely that
58
there is an increase of holes during desodiation and a decrease of holes duringresodiation.
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Figure 4.13. Paper VIII – (a) O K-XAS pre-edge region in TFY and PFY modeof Na0.56Mg0.04Ni0.19Mn0.70O2. (b) O K-RIXS at incident photon energy hνin =531.0 eV; (c) Relative integrated pre-edge intensities of TFY, PFY, and TEY (notshown) mode XAS.
Figure 4.13 (b) shows O K-RIXS at an incident photon energy of 531.0 eV.The RIXS maps reveal no occurrence of strong localized states, instead weinterpret this as the evolution of the valence band with O 2p character. Thetwo vertical arrows show the development of a low-energy signature at theend of the first desodiation. This peak disappears upon resodiation and re-appears again at second desodiation. This is again similar to the observa-tion in the LrNMO cathode as well as the previously discussed Na-compoundNa0.67[Mg0.28Mn0.72]O2. However, here we do not observe a strong simulta-neous increase of the elastic peak seen in the other materials, whose originis attributed to localized oxygen species that undergo a reversible anionic re-dox process. This may indicate that the emerging feature is related to moredelocalized O 2p states. Comparing the RIXS-spectra of the pristine material
59
(black trace) with the resodiated material (blue trace) reveals a more symmet-ric spectral weight distribution in the resodiated material than in the pristinematerial, which corroborates that initial activation processes also affect thebulk oxygen ions.
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5. Conclusion
The electronic structure analysis at the atomic scale is an indispensable aspectof the improvement of functional materials like battery cathode materials. Itallows a shift from a trial-and-error approach, where a variation of virtuallylimitless elemental constituents is systematically tested by electrochemicalmeans for higher performance, to a much more target-oriented approach. Ifthe specific behavior and thus function of the different elemental constituentsof a material are well-understood at the atomic scale, a selective and targetedreplacement of those elements can be done, which speeds up the developmentof new, improved cathode materials.
The work of this thesis shows that the element specific and chemical sitesensitive X-ray core-hole spectroscopy is a powerful experimental tool for thispurpose. XAS and RIXS have been used to prove the important role of anionicredox in various cathode materials and delivered important experimental evi-dence that this can be used to improve battery performance.
Oxygen redox processes have been observed for different materials includ-ing Li-rich layered as well as sodium-ion battery layered cathode materi-als. One common characteristic of the anionic redox is seen in the O K-pre-edge intensity of XAS. Charging leads to a broadening and increase ofintensity in this region, which is reversible upon discharging. The second key-feature for identifying O redox comes in the form of two reversible features(dis)appearing in the O K-edge RIXS upon (dis)charging, namely an elasticpeak and a low energy emission feature. A notable exception in this class ofmaterials is the lithium-rich end-member LrMO where we see no evidence forreversible O in thin film cathodes. On the other hand the initial lithiation cycleis governed by strong changes in the average Mn oxidation state.
New and profound aspects to the underlying redox mechanisms in spinelLNMO have also been discovered in this thesis. While the changes of Ni L-edge XAS observed upon de- and relithiation would suggest that the redoxcenter is localized on the Ni sites of the lattice, our O K- and Ni L-edge RIXSanalysis reveals a more complicated picture. The broadening of O K-RIXSspectra as well as the strong charge transfer related intensities in Ni L-RIXSreveal a strong anionic contribution to the overall redox processes.
The work of this thesis has shown that the combination of X-ray absorptionspectroscopy with resonant inelastic X-ray scattering is a powerful experimen-tal tool for understanding the electronic structure of battery materials.
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6. Populärvetenskaplig sammanfattning
I vårt samhälle, där vi vant oss vid konstant tillgång till energi, är batterieren allt viktigare komponent. Bärbar elektronik och elbilar är exempel på om-råden där batterier idag spelar stor roll. I framtiden kan batterier exempelviskomma att spela en allt större roll även för storskalig energilagring i de el-distributionsnät som baseras på en större andel oregelbundna energikällor såsom vindkraftverk och solceller. Med dessa nya och växande användningsom-råden ställs också högre krav på batteriernas effektivitet och lönsamhet. Atthitta bättre batterier kräver utveckling av nya material och förståelse kring hurmaterialen fungerar.
+-e-
Uppladdning
Li+
+-
e-
Urladdning
Li+
XAS
Anod KatodElek-trolyt
RIXS
Energi Energi
Litiumion batteri(a) (b) (c)
Röntgenspektroskopi
Figure 6.1. (a) Skiss av upp- och urladdningsprocesser i en LIB. Turkosa kulor repre-senterar Li+ katjoner. XAS (b) och RIXS (c) spektra av syre i en urladdad (blå) ochuppladdad (orange) LIB katot.
Batterier består bland annat av elektroder, vilka lagrar den energi som om-vandlas via redoxreaktioner, samt en så kallad elektrolyt, vilken är en vätskasom kan leda joner, se Figure 6.1. Många moderna elektrodmaterial har enkristallstruktur som kan ta upp och avge positiva joner (så kallade katjoner) ien process som kallas interkalering. Vanligast är litiumjoner, vilka gett namnåt de numera vanliga litiumjon batterierna (LIBs), men det finns även materialsom interkalerar natrium (SIBs) och andra katjoner. Under urladdningspro-cessen vandrar katjoner från den negativa elektroden (anod), genom elektroly-ten till den positiva elektroden (katod), och samtidig överförs elektroner frånanod till katod via en extern krets. Elektronerna har högre energi när de befin-ner sig i anoden jämfört med i katoden och denna energi skillnad används för
62
att driva exempelvis en lampa i den externa kretsen. I uppladdningsprocessenkrävs en energikälla för att driva den omvända processen, i vilken elektroner-na och katjonerna transporteras i den motsatta riktningen från katod till anod.Genom sådan uppladdning så lagras kemisk energi i batterierna. Upprepadeladdningscykler kan till viss del bryta ner elektrodmaterialen och försämrarprestandan hos batteriet. Målet för mycket av den batteriforskning som görsidag är att utveckla material som kan lagra mer energi i mindre volym ellermassa och som uppvisar en lång livstid med många upp- och urladdningscyk-ler utan betydlig prestandaförlust. Andra mål kan vara att utveckla batteri-material bestående av mer vanligt förekommande grundämnen (och därmedbilligare) eller miljövänligare komponenter.
Något förenklat så sker de kemiska processerna, dvs redox reaktionerna,under upp- och urladdning, genom att olika atomer i materialen tar upp ochavger elektroner, atomerna reduceras respektive oxideras. En grundläggandeförståelse för batteriers funktion kräver därför också forskning på atomär ni-vå. Sådan förståelse är grundläggande för förbättringen av batterimaterial, inteminst vad gäller förmågan att ladda upp och ladda ur batteriet många gångerpå ett effektivt sätt. Processer som sker på atomär nivå är osynliga för kon-ventionella mätmetoder, och därför används avancerade analysmetoder för attstudera batterimaterialens funktion. Olika röntgenspektroskopimetoder har enunik förmåga att undersöka material och i den här avhandlingen visas att detvå specifika metoderna, röntgenabsorptionsspektroskopi (XAS) och resonantinelastisk röntgenspridning (RIXS) är mycket användbara.
Katodmaterial av LIBs och SIBs består ofta av metaller som reagerat medsyre för att forma metaloxider. Dessa innehåller utöver litium eller natriumolika metalliska grundämnen som exempelvis kobolt, nickel eller mangan. Närpositiva litium- eller natriumkatjoner (Li+ eller Na+) går ur respektive går ini katoden, måste laddningsbalansen bevaras. Det sker genom att en lika stormängd negativa partiklar – i bästa fall elektroner (e−) i form av elektrisk ström– samtidigt går ur respektive går in i materialet. Dessa processer sker genomredoxreaktioner som involverar katodens olika atomer.
En viktig grundläggande fråga är att förstå vilka och på vilket sätt olikaämnen är delaktig i redoxreaktionerna. Detta är två av de återkommande frå-gor som varit centrala i den här avhandling och för att svara på dessa frågori detalj undersöks elektronstrukturen hos olika batterimaterial. I den här av-handlingen är avancerad XAS och RIXS de kraftfulla experimentella spektro-skopiska metoder som främst använts för att karaktärisera materialen. Rönt-genspektroskopi utförs genom att störa systemets elektronstruktur med hjälpav röntgenstrålar; högenergetiskt ljus, med känd initial energi och intensitet,belyser materialen och därefter observera systemets reaktion. Olika grundäm-nens elektronstruktur är unik vad gäller antalet elektroner och elektronernasspecifika bindningsenergier. De hårdast bundna elektronerna har bindningse-nergier som är vitt åtskilda mellan olika grundämnen. Genom att låta ljus medtillräckligt hög och känd energi (röntgenstrålar) interagera med dessa elektro-
63
ner är det möjligt att adressera specifika grundämnens roll i redoxreaktionernasom sker i batterimaterialen vid upp- och urladdning. En detaljerad tolkningav systemets reaktion vid avancerad röntgenspektroskopi kräver ofta både ettkvantteoretiskt ramverk för att till fullo beskriva materialen samt en god för-ståelse för de underliggande dynamiska processerna.
I konventionella batterier är redoxprocesserna i katoderna centrerade till degrundämnen i kristallstrukturen som har positiv jonladdning, så kallade katjo-ner. En ny strategi för att öka kapaciteten hos nya katodmaterial är att utnyttjaredoxreaktioner där negativa joner i katodens kristallstruktur, så kallad anjo-ner, deltar. För oxidmaterialen representeras anjonerna av syreatomerna i ma-terialet. Resultaten i denna avhandling har visat att detta sker i några av deundersökta materialen. Ett vanligt kännetecken för den anjoniska redoxreak-tioner ses i intensiteten hos XAS mätningar som fokuserar på syreatomerna,se Figure 6.1 (b). Uppladdning (de-interkalering av katjoner ur katoden) le-der till en breddning och ökad intensitet av det lägsta energiområdet i syretsXAS spektra, vilket är reversibelt vid urladdning. Ett annat kännetecken föratt identifiera anjoniska redoxreaktioner kommer i form av två reversibla in-tensitetsvariationer som visas i RIXS mätningar på syreatomerna vid upp- ochurladdning av batteriet; en elastisk topp runt 0 eV och en topp runt 8 eV påenergiförlustskalan, se Figure 6.1 (c).
I den här avhandlingen har syrets (anjonens) deltagande i redoxproces-serna observerats för olika katodmaterial inklusive litium- och natriumrika,skiktat kristallstrukturer Li[Li0.2Ni0.2Mn0.6]O2, Na0.67[Mg0.28Mn0.72]O2 ochNa0.78[Li0.25Mn0.75]O2. Ett anmärkningsvärt undantag i denna materialklassär det litiumrika ämnet Li[Li0.33Mn0.67]O2 där reversibelt anjoniska redoxre-aktioner i tunnfilmskatoderna inte kunde påvisas.
Nya och djupa aspekter på de underliggande redoxmekanismerna i spinelLiNi0.44Mn1.56O4 har också upptäckts i denna avhandling. Ändringarna somobserverats i XAS spektra vid upp och urladdning antyder att redoxcentret ärlokaliserat på nickelatomerna, men detaljerade RIXS-analyser avslöjar en merkomplicerad bild med starkt anjoniskt bidrag till redoxprocesserna.
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7. Acknowledgements
In this section I have the chance to acknowledge all the people without whomthis work would not have come to fruition.
I want to start by thanking my three supervisors Dr. Laurent Duda, Prof.Dr. Håkan Rensmo, and Dr. Maria Hahlin. Thank you all for your support,guidance, and your great patience. I am truly indebted to all of you.
Laurent, I will never forget our first of many joint beamtimes at the ALSin Berkeley: A 72 hours monster shift in the noisy environment of the ISAACbeamline with delicate instrumentation that requires a near-constantly highlevel of attention to detail – and all of that was two-manned by a seasonedexperimentalist and a rookie beamline user who had no idea that mere tenhours of sleep stretched out over three hard work days on old office chairs orclearly too small couches would be in the realm of the humanly possible. Andeven though the liquid flow cell did not want to cooperate, we at least got wiserto book rooms at the guest house well in advance for the next couple of visitsto Berkeley. But banter aside. I really want to extend my sincere gratitudetowards you, Laurent. I could not have come this far without your patient andkind guidance on experimentation, theory, and the presentation of scientificresults.
And Håkan, I could always count on your leadership in times were I neededto regain focus and to find my way forward again. So without that kind ofpersistent support this thesis would still not be written and I would be angry atmyself for not trying harder. Instead, I now experienced the incredible feelingof holding my first ever book in my hands, which is a feeling that is hard todescribe and cannot be bought. It was a tough road and you showed me theway when I got lost. Thank you so much!
Last but not least, Maria. I cannot estimate too high the value of yourtruly conscientious feedback on version after version of manuscripts for ourpapers and also this thesis. You always ask the right questions, which not onlyundoubtedly increased the quality and readability of the final products but alsotaught me a valuable lesson in science in general: It is always commendable toask questions; a true scientist values truth over saving face. Thank you, Maria!
I want to once again express my gratitude towards the many contributorsand co-authors that provided all the excellent battery samples and a large rangeof expertise to all the different projects of this thesis. Burak Aktekin and LeAnh Ma from the Department of Chemistry, Uppsala University. Kun Luo,Matthew Roberts, Urmimala Maitra, and Robert House from the Departmentof Materials, University of Oxford. And last but not least Kazuhiro Hikima
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from the Department of Chemical Science and Engineering, Tokyo Instituteof Technology. It was a great pleasure to meet and work together with all ofyou.
And I would also like to extend my appreciation beyond the work envi-ronment. First and foremost I want to thank my whole family, especially myparents. Ohne eure Liebe und Unterstützung wäre ich niemals bis zu diesemPunkt gekommen. Ihr habt mich stets ermuntert meinen eigenen Weg zu gehenund wart für mich da, wann immer ich euch am meisten brauchte. Danke!
I must also thank my wonderful, loving partner Kristin who expertly tookcare of the home front. This was a huge relief for me during the last coupleof months. Your wholesome food kept me healthy and energetic during a timewhen I had very little time for preparing proper food myself. I am sure youcontributed a lot to my health, especially by keeping me from getting chubby,which I have seen happening to many PhD students during the final period.Thank you cutie!
I can think of many more individuals that I could mention here, but it is timeto finish. However, I feel the need to express my gratitude towards one groupof people that made this project possible in the first place: The hard-workingand welcoming people of Sweden. Not only did this remarkable people es-tablish a great infrastructure for high-end science, they also opened their greatnation to this German student of the natural sciences and for that I will be for-ever thankful. Du gamla, du fria, you will always have a special place in myheart.
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Acta Universitatis UpsaliensisDigital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 1835
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A doctoral dissertation from the Faculty of Science andTechnology, Uppsala University, is usually a summary of anumber of papers. A few copies of the complete dissertationare kept at major Swedish research libraries, while thesummary alone is distributed internationally throughthe series Digital Comprehensive Summaries of UppsalaDissertations from the Faculty of Science and Technology.(Prior to January, 2005, the series was published under thetitle “Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology”.)
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