High quality epitaxial thin films of actinide oxides, carbides, and nitrides: Advancing...

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Coordination Chemistry Reviews 266–267 (2014) 137–154

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Coordination Chemistry Reviews

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igh quality epitaxial thin films of actinide oxides, carbides, anditrides: Advancing understanding of electronic structure of

-element materials

rian L. Scotta,∗, John J. Joycea, Tomasz D. Durakiewicza, Richard L. Martinc,. Mark McCleskeyb, Eve Bauera, Hongmei Luod, Quanxi Jiaa

Materials Physics and Applications Division, Los Alamos National Laboratory, Los Alamos, NM, USAChemistry Division, Los Alamos National Laboratory, Los Alamos, NM, USATheoretical Division, Los Alamos National Laboratory, Los Alamos, NM, USAChemical Engineering Department, New Mexico State University, NM, USA

ontents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1382. Theoretical background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1383. Actinide containing epitaxial thin films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 140

3.1. PAD technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1403.2. Uranium oxide films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1403.3. Transuranic oxide films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1423.4. Uranium carbide films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1443.5. Uranium nitride films. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 144

4. Photoelectron spectroscopy. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1454.1. UO2 and U3O8 ARPES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1464.2. PuO2 ARPES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1484.3. Actinide oxide comparisons . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1494.4. UC2 and UN2 ARPES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1494.5. Theory–experiment comparison . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 150

5. Optical band gaps . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1515.1. UO2 and U3O8 band gap measurements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1515.2. PuO2 and NpO2 band gap measurements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1515.3. Electrical conductivity in UC2 and UN2 films. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153

6. Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153

r t i c l e i n f o

rticle history:eceived 31 May 2013eceived in revised form4 September 2013ccepted 25 September 2013vailable online 4 October 2013

a b s t r a c t

Over the past five years we have developed a solution based technique for synthesizing high qualityepitaxial thin films of actinide materials. These films include oxides (UO2, U3O8, PuO2, and NpO2),UC2, UN2, UC2−xOx, and UN2−xOx. These nearly single crystal quality films have allowed unprecedentedexperimental measurements (ARPES, optical band gap, electrical conductivity, and XRD). These data havebench marked and advanced DFT theoretical predictions and understanding of bonding and electronicstructure in f-orbital actinide materials. The synthesis and characterization of the films, measurement

eywords:ctinide oxidesctinide nitridesctinide carbidesybrid DFTlectronic structure-Elements

of electronic structure, and the resulting validation of theoretical models are reviewed. The potentialimpacts of this work in areas ranging from nuclear energy to environmental fate of actinides in theenvironment are discussed.

∗ Corresponding author at: Los Alamos National Laboratory, MS J514, Los Alamos, NM 8E-mail address: [email protected] (B.L. Scott).

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38 B.L. Scott et al. / Coordination Chem

. Introduction

The thin film synthetic technique known as polymer assistedeposition (PAD) was first reported in 2004 [1]. Since then PADas been used to make a wide variety of high quality, epitaxialhin films with properties ranging from metallic, semiconduct-ng, ferroelectric, ferromagnetic, and/or superconducting [2]. These

aterials include binary metal oxides (TiO2, Eu2O3) [1,3], ternaryetal oxides (SrTiO3, CoFe2O4) [4], and complex metal oxides

YBa2Cu3O7−x) [5]. The growth of nitrides (NbN, MoN, Mo2N, andrTiN2) [6–8], carbides (TiC, SiC, NbC) [2,9,10] and single elementsGe) [11] has also been realized. This technique involves complex-ng metal ions to polymers in solution. The polymer precursors arehen spin coated onto a single crystal substrate that has appro-riate lattice match with that of the desired film. Finally, the film

s heated in the appropriate atmosphere to drive off the polymernd form the epitaxial film. In general, carbides are formed underthylene, oxides under air, and nitrides under ammonia. This rela-ively simple solution based technique offers many advantages overraditional thin film synthetic techniques including physical vaporeposition, chemical vapor deposition, and sol–gel techniques. Twof the most notable advantages offered by PAD are that it requiresmall amounts of material and gas phase techniques are not needed,hich are both important factors when working with hazardousaterials such as the actinides and in particular the transuranic

lements.In the past five years we have built on these early PAD successes

o synthesize several actinide containing thin films. The motivationor this work has been the synthesis of well characterized singlerystal quality materials to enable unprecedented experimentaleasurements, especially in the area of electronic structure. Before

his work single crystal quality materials of purity did not existor many of the basic oxides, nitrides, and carbides of transuraniclements such as U, Pu, Np and beyond. In particular, for PuO2nd NpO2 only a few reports of single crystal growth exist, andhe characterization data indicate impure materials [12–14]. Theadioactivity and toxicity of the actinide elements make chemicalnd physical vapor deposition techniques problematic. The syn-hesis of these materials is also complicated since they have veryigh melting points (2400 ◦C for PuO2, 2547 ◦C for NpO2, 2865 ◦C

or UO2, 2350 ◦C for UC2), precluding growth of crystals from theelt. In this vein, we employed PAD to overcome these synthetic

nd safety hazards to synthesize epitaxial thin films of the followingctinide materials:

1) PuO2, NpO2, UO2, U3O8, ThO22) UC2, UC2−xOx

3) UN2, UN2−xOx

These thin films have provided systems on which to makeround breaking measurements. For example, until our recent workn PuO2 and NpO2 the optical band gap of these materials hadever been measured. For both materials decade’s old values ofand gaps based on indirect measurements of the activation barriero electron conduction were cited. Also, the single crystal quality ofhe epitaxial thin films has allowed angle resolved photoelectronpectroscopy (ARPES) measurements to map out band dispersionnd the Fermi surface. We have used these experimental results todvance understanding and validate new theoretical models of f-rbital interactions, especially with regard to their strong electronorrelations.

Beyond providing experimental bench marking of f-orbital the-
ry, these materials also are relevant to research efforts in nuclearnergy and environmental fate of the actinide elements. For exam-le, Generation IV nuclear reactors fuels that are being considered
nclude uranium carbides, nitrides, and mixed oxide-nitride and

Reviews 266–267 (2014) 137–154

oxide-carbide fuels [15]. These potential new fuel systems mayprovide higher thermal conductivity, providing increased safetyand fuel life. Currently, fuel rods are taken out of service at 5% burnup or less, and understanding electronic structure and f-electronbehavior in these systems is crucial to predicting thermal con-ductivity and potentially materials aging. The aging phenomenoncarries over into how nuclear fuels evolve after taken out of the fuelcycle, and how actinides may enter and move in the environment.There are currently 500 t of separated plutonium and 60,000 t ofuranium in the world-wide stockpile [16,17]. Moreover, there arenumerous radioactive contamination sites spread across the globethat need to be stabilized and/or cleaned up. The fundamental prop-erties of actinide materials lays the foundation for studies involvingtheir interactions with air, water, minerals, natural organic mat-ter, and microbial organisms, and the ability to make high-qualityfilms having a particular crystalline face is crucial to enable studiesaimed at understanding materials aging and environmental fate.Improved predictive models could aid decision making and spawnnew technologies to minimize risks across the nuclear fuel cycleranging from mine waste to accidental releases from reactors tonuclear waste.

In this review article we will describe our synthetic techniquesfor several actinide thin films of the oxides, nitrides, and car-bides. Subsequently, the experimental characterization of thesefilms including X-ray diffraction, optical band gap, electrical con-ductivity, and ARPES will be presented. The validation of theoreticalmodels resulting from these new data will be described, and theimprovement over previous theoretical attempts in these materialswill be highlighted.

2. Theoretical background

The f-element oxides are of interest in a number of applicationcontexts ranging from automobile exhaust catalysts (cerium oxide,or Ce2O3) to nuclear fuels (uranium dioxide, or UO2). In addition,they are also of fundamental interest as members of the set of Mottinsulators, a class of materials long problematic for theory becausewhile they should conduct electricity, they instead act as insulators,especially at low temperatures [18,19].

What came to be known as the Mott insulator problem wasfirst defined at a conference in 1937 in Bristol, England. The Dutchphysicists Jan Henrik De Boer and Evert Johannes Willem Verweypresented electrical conductivity data for a number of first-rowtransition metal oxides, including scandium, ScO (d1); titanium, TiO(d2); chromium, CrO (d4); manganese, MnO (d5); nickel, NiO (d8);and copper CuO (d9). The transition metal sites in these materialshave a formal valence of +2, leading to the anticipated d-electroncount given parenthetically above.

De Boer and Verwey pointed out that because each of thesematerials had a partially filled d-band, according to the thenrecently proposed band theory of solids they should all be metallic.In stark contrast, they were all insulators [20]. A spirited discussionfollowed, but the problem waited to be addressed until after WorldWar II when it was revisited by British physicists Mott [21,22] andHubbard [23–26], who established that large electron-correlationeffects not adequately treated in the band theory were responsiblefor the insulating behavior.

Perhaps the easiest way to understand the Mott problem qual-itatively is through the potential energy curve for the diatomichydrogen molecule, H2 (Fig. 1).

At equilibrium, we know that a molecular orbital description is
a reasonably good approximation for the wave function. From twoH 1s atomic orbitals, one on each hydrogen atom, we form bonding(in phase) and antibonding (out of phase) linear combinations. Thebonding molecular orbital lies lower in energy and is occupied with

B.L. Scott et al. / Coordination Chemistry

Fig. 1. Potential energy curve for the diatomic hydrogen molecule, H2. Near theequilibrium geometry a molecular orbital description provides a qualitatively cor-ria

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ect description of the bonding. At long distances, the valence bond approximation,n which one electron is localized on the left hydrogen, and one on the right, is moreppropriate.

wo electrons; the corresponding antibonding molecular orbital ismpty. It is easy to show that this wave function contains equalomponents of ionic (H+H−, H−H+) and Heitler–London (covalent,:H) character. At long distances, however, the physical situation isest described by a wave function in which one electron is localizedn the left atom and the other on the right.

The Hubbard model captures the essential physics. The modelonsists of a hopping term (t), which transfers an electron from oneenter to another and determines the electronic kinetic energy, and

repulsive term (U), which reflects the energy required to put twolectrons in the same atomic orbital. Near equilibrium, the kineticnergy gained by delocalizing (hopping) the electrons easily over-omes the repulsion in the ionic component of the wave functionssociated with the double occupancy, and the molecular orbitalpproximation works well.

On the other hand, at long distances the orbital overlap is weak,nd so the hopping term and kinetic energy are small, and theenalty associated with the double occupancy of the molecularrbital form dominates. The energy is minimized by eliminatinghe ionic piece, leaving only the Heitler–London component. Herehe electronic correlation is strong: when one electron is on the left,he other is on the right; the electrons localize, one to each nucleus.

ost importantly, in this localized wave function the ionic termsre now absent; the covalent piece now dominates.

Consider now a long chain of equally spaced hydrogen atomsFig. 2). If the distances are typical of an H2 bond length, the molec-
lar orbitals will form a continuous band of states, called the crystalrbitals, or Bloch states.
Each hydrogen atom contributes one electron, so the band wille partially filled, and the ground state a metal. If the distance

ig. 2. The nature of the electronic bands in a chain of equally spaced hydrogen atoms chmaller. In the case where the overlap is weak, the electrons will tend to localize. This intrerm U.

Reviews 266–267 (2014) 137–154 139

between hydrogen centers is made longer, or if the size of theatomic orbitals were made smaller, thereby reducing the orbitaloverlap and hence the kinetic energy, at some point the repulsionwill dominate and the system will localize electrons, one to a center.Higher-order interactions generally lead to an antiferromagneticordering of the spins. Note also that this localization will lead toan energy gap, and insulating behavior, Because the lowest excitedstate corresponds to the energy it takes to move an electron fromone center to another. This energy, the Mott gap, is given by theHubbard repulsion term U.

While the basic physics of this situation has been understoodfor some time, the point at which the transition from delocalized tolocalized behavior occurs has been difficult for first-principles the-ories to predict. For example, the workhorse of electronic structuretheory, density functional theory (DFT), generally overestimatesthe electronic delocalization, predicting metallic behavior for theseinsulators [27]. For example, the first-generation functional, thelocal density approximation (LDA), predicts NiO to be metallic; itis in fact an antiferromagnet with a gap of some 4 eV [28]. Second-and third-generation approximations—generalized gradient (GGA)and meta-GGA functionals—do no better. A fourth generation, thehybrid functionals, departs from the conventional local and semi-local approximations by incorporating a component of the fullnonlocal Hartree–Fock exchange term:

Exc = aEHFx + (1 − a)EGGA

x + EGGAc (1)

The constant a is in principle adjustable, but a value in the vicinityof 0.25 has proven to be quite successful. It is generally not variedfrom one material to another, and it is in that sense universal. Thisfunctional, with a = 0.25, when combined with generalized gradi-ent exchange and correlation functionals, is known as PBE0 [29].This refinement largely remedies the problem. A number of studieshave now shown that hybrid density functional theory with a = 0.25yields the proper lattice constants, insulating behavior, band gaps,and magnetic properties of many Mott insulators, when known[30,31].

The success of the hybrid functionals in describing the Mottinsulators is related to improvements in what is known as the self-interaction error of DFT. There are other many-body approacheswhich address this problem either directly (the self-interaction cor-rection, SIC), or indirectly (the DFT + U and DMFT approaches). Ourintent here is not to review the theoretical methods. Each of theseapproximations has strengths and weaknesses; their performancein actinide systems has recently been reviewed [31]. Our princi-pal focus has been in developing and applying hybrid DFT, whosemajor strength is that there are no material-dependent parametersneeded for its implementation.

In what follows we will review the results of calculations for UO2and U3O8 using PBE0 or the closely related screened hybrid func-tional HSE [32], compared with experimental results that probe theoccupied density of states (photoemission). Excellent agreement

anges if the distance between H centers becomes longer or the size of the orbitalsoduces an insulating gap, whose energy is given roughly by the Hubbard repulsion

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ith experiment is observed, and the calculations suggest that UO2s indeed a Mott insulator [33–36], the 5f levels quite localized withery little dispersion. We then turn our attention to plutoniumxide (PuO2), where hybrid density functional theory predicts atrong, counterintuitive, orbital mixing between the Pu 5f and O 2prbitals. This manifests as a large dispersion in the band structure,

prediction most easily tested with angle-resolved photoemissionpectroscopy. Finally, we will compare band gaps calculated usinghese hybrid techniques with the experimentally measured valuesaken from the epitaxial thin films of UO2, U3O8, PuO2, and NpO2.

. Actinide containing epitaxial thin films

The synthesis and subsequent experimental measurements ofctinide containing thin films provide many advantages over otheraterials including polycrystalline powders and films and even sin-

le crystals. First, the single crystal like quality of the films allowor more direct comparison with DFT theory, which assumes andeal crystalline lattice free of defects and surface effects present inolycrystalline materials. Second, the well aligned epitaxial thinlm (typical rocking curves ≤0.08◦) provides a more favorableatio of experimental signal (XRD, optical, PES) to material masshen compared to powders and polycrystalline films, and makesolarized experiments possible. Third, the smaller amount of mate-ial required for measurements reduces the likelihood of healthazards and instrument contamination when toxic and radioactivectinide elements are being studied. Fourth, the 25–100 nm thick-ess of the films allow the transmission of light leading to highuality optical measurements when compared to single crystals.inally, we will show that through choice of substrate the crys-alline face of the film can be controlled. This will allow a singleacet of a material to be measured with respect to reactivity.

.1. PAD technique

PAD is both solution-based and polymer-assisted: It is solution-ased in that, when used for deposition, it employs directpplication of a water-based chemical solution to the surface toe coated (“substrate”). The solution contains water, water-solublealt, and water-soluble polymers such as polyethylenimine (PEI)
nd/or ethylenediaminetetraacetic acid (EDTA). The specific saltelected, i.e. what metallic element it contains—determines whichetal element will be present in the substance synthesized. Once
he substrate is coated, it is heated in a controlled environment

Fig. 3. The step-by-step process of growing epitaxial thin films

Reviews 266–267 (2014) 137–154

to convert the coated precursor film into desired compounds. It ispolymer-assisted in that the solution also contains organic poly-mers such as PEI that bond to metal ions, a key that gives PADits advantages over other chemical solution deposition processesand the capability to synthesize materials with desired structuraland physical properties. The polymers used in the solution (i) pre-vent the metal ions from engaging in unwanted chemical reactions,(ii) maintain homogenous distribution of the metal ions in solu-tion, thus yielding an even, uniform coating, (iii) help maintain thedesired viscosity of the solution, enabling good control of thickness,and (iv) keep the solution very stable. Fig. 3 shows the schematicdrawing of typical steps to grow films by PAD. It should be notedthat the formation of polymer/metal ion complex in the solutionmakes it possible for filtration, a critical step used by PAD.

3.2. Uranium oxide films

The first epitaxial thin films of actinide materials, UO2 and U3O8,were reported by members of this team in 2007 [37]. Until thistime only polycrystalline films of UO2 had been prepared by reac-tive sputtering, laser ablation, and sol–gel approaches. The growthof pure single crystal uranium oxides is greatly complicated by arich phase diagram including multiple phases (UO2, U3O7, U3O8,and UO3) that can easily convert to one another [38,39]. More-over, many of these phases have polymorphs, i.e. U3O8 has twoknown polymorphs (space groups C2mm and P62m). The conver-sion between phases and polymorphs is also dependent on the formof the materials ranging from powders, to polycrystalline pellets, tocrystals. The PAD technique provides a way to stabilize a particu-lar phase, valence state, and/or polymorph through choice of singlecrystal substrate. The reaction chemistry remains the same for allphases, and only the substrate is changed to template the growth ofthe desired phase. The precursor for uranium oxide films are simpleaqueous solutions of UO2(NO3)2 added to polyethylenimine (PEI)polymer to form an aqueous solution with adjusted pH of 8.0. Thissolution is then spin coated onto the desired substrate, which areannealed at 1000 ◦C in air or O2 to form the uranium oxide epitaxialfilms.

The formation of uranium oxide films on the substrate lan-thanum aluminum oxide (LaAlO3; LAO) results in UO2 films. The
LAO lattice is pseudocubic, with a unit cell of dimension 3.789 A, andprovides excellent lattice match for cubic UO2 (a = 5.466 A, Fm3m).The face centered UO2 lattice, when rotated 45◦, matches the LAOlattice spacing to within 2% (Fig. 4).
using the polymer assisted deposition (PAD) technique.

B.L. Scott et al. / Coordination Chemistry Reviews 266–267 (2014) 137–154 141

Fs

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ig. 4. Picture showing lattice match between the epitaxial UO2 thin film and theubstrate, LaAlO3.

The XRD pattern of the film is shown in Fig. 5a, where onlyhe reflections (2 0 0) and (4 0 0) of the UO2 are observed, demon-trating alignment of the film along the a-axis. In addition, XRD-scans (Fig. 5b) show the epitaxy of the UO2 film with the LAO

ubstrate and the 45◦ rotation between the two lattices. Finally,igh-resolution electron microscopy shows excellent alignment
f film and substrate, no detectable second phases or voids, and
high degree of epitaxy as evidenced by electron diffractionFig. 6).

ig. 5. (a) X-ray diffraction of the UO2 film on LAO showing alignment of the filmlong the a-axis, and (b) �-scans showing epitaxy with the LAO substrate.

Fig. 6. TEM images showing the UO2 epitaxial film deposited on LAO substrate. (Top)A high resolution TEM image of a cross-section showing the clean interface between
film and substrate; (bottom) selected area electron diffraction pattern [37].
Angle resolved photoelectron spectroscopy of the UO2 filmswere identical to those of single crystal UO2 samples (vide infra).The films also appear to be immune to conversion to other phasesunder ambient conditions as we have observed stable XRD pat-terns for more than 400 days. This is consistent with the fact thatthe films were formed and annealed at 1000 ◦C; the LAO substrateseems to stabilize UO2 to additional oxidation. Moreover, H2 dos-ing experiments with ARPES on both our epitaxial thin films andfreshly cleaved single crystals of UO2 did not show any appreciabledifferences in the valence band further indicating the resistance ofthe UO2 films to oxidation. We believe the stability of the UO2 filmsto oxidation is the result of the film structure being pinned to theunderlying substrate structure.

Through the proper choice of substrate, we were able to syn-thesize both polymorphs of U3O8 as epitaxial thin films for the firsttime. The hexagonal polymorph of the U3O8 oxide is stabilized onc-cut sapphire (0 0 1), while the orthorhombic phase forms on ther-plane of sapphire (0 1 2) under the same reaction conditions. Thehexagonal polymorph is normally stable above 210 ◦C, but remainsstable when the film is cooled to room temperature. The sapphire(Al2O3) crystal structure is hexagonal (R3c), with lattice parametersof a = 4.759 A, and c = 12.992 A. In contrast, the hexagonal (P62m)crystal structure of U3O8 has lattice constants of a = 6.210 A andc = 4.517 A. While the lattice parameters of U3O8 do not closelymatch those of sapphire, the oxygen–oxygen distances in the 0 0 1
planes are on the order of 2.5 A for both. The TEM and electrondiffraction measurements for the hexagonal phase are presented

142 B.L. Scott et al. / Coordination Chemistry Reviews 266–267 (2014) 137–154

F hire,

t

it

oTomattbs

ac[spspfirbl(sArT

ig. 7. (a) high resolution TEM showing cross-section of U3O8 growth on c-cut sapphe film and substrate interface [37].

n Fig. 7, and again show the high quality of the films with respecto purity, defects, and epitaxy.

The orthorhombic structure of U3O8 has a rectangular lattice ofxygen atoms lying in the (0 0 1) plane of 4.1 A and 4.3 A on a side.he corresponding (0 1 2) plane of sapphire has a similar networkf oxygen atoms that is 4.7 A and 5.1 A on a side. The approxi-ate oxygen–oxygen distance match between sapphire and U3O8

re close enough to provide growth of an epitaxial thin film ofhe orthorhombic phase of U3O8 (not shown). This system showshe importance of not only lattice match, but also that distancesetween atoms should also be taken into account when matchingubstrate to film.

Following our work on the uranium oxides published in 2007 group at the University of Illinois published the synthesis andharacterization of epitaxial thin films of uranium oxides in 201240]. These films were produced via reactive gas magnetronputtering under varying temperatures and partial pressures ofure oxygen. This work used the single crystal substrates r-planeapphire (Al2O3) and yttrium stabilized zirconia (YSZ). Higherartial pressures of oxygen (pO2 = 0.03–0.07 Pa) yielded U3O8lms, while lower pressures (pO2 = 1.2 × 10−5 to 9.3 × 10−6 Pa)esulted in UO2 films. They found that UO2 epitaxial films coulde produced on both the r-plane sapphire and YSZ, but that the

atter gave tighter rocking curves (0.06–0.60◦) vs. r-plane sapphire1.1–2.7◦). The orthorhombic phase of U3O8 formed on r-plane
apphire, but yielded broad rocking curves of approximately 2◦.ll films showed good epitaxy with their substrates, and broadocking curves were attributed to significant dislocation density.his well done and thorough study also looked at film architecture
(b) selected area electron diffraction, and (c) fast Fourier transform image between

using Rutherford backscattering (RBS) and X-ray photoelectronspectroscopy (XPS). Both of these techniques revealed aluminumimpurities when sapphire substrates were used. The correspond-ing YSZ materials did not show any diffusion of Y or Zr into theuranium oxide film layers. Their XPS sputtering experiments alsoshowed a slight oxidation of the UO2 films down to about 10 Aconsistent with formation of U3O7 or U3O8.

3.3. Transuranic oxide films

After developing techniques for the synthesis of uranium oxideepitaxial thin films we moved our focus to plutonium and nep-tunium oxides [41]. The precursor solutions for these preparationsconsisted of EDTA complexes of Np(V) and Pu(IV) bound to PEI poly-mers as homogeneous solutions near pH of 5. Solutions (30 �l) werespin coated onto 1 cm2 substrates and annealed at 1000 ◦C to formthe films. The cubic NpO2 (a = 5.433 A; 3.842 A diagonal) is isostruc-tural to UO2 (a = 5.466 A; 3.865 A diagonal), and the first attempt togrow an epitaxial thin film was on the LAO substrate (a = 3.789 A)that gave quality UO2 films (vide supra). The result was a polycrys-talline NpO2 film (Fig. 8), which shows the importance of substrateto the formation of epitaxial films.

When the NpO2 film was grown on a buffer layer of UO2 onLAO, allowing for closer lattice match, an epitaxial film of NpO2was formed (Fig. 9). The rocking curve of this bilayer film shows
both NpO2 and UO2 films (Fig. 9, inset).
While these materials are not suitable for further characteri-zation to measure optical band gaps and electronic structure inNpO2 the result does show that PAD can be used to make multilayer


F[

fifilaaatoasdtPsmor

t

Fig. 10. Rocking curve of PuO2 epitaxial film on YSZ [40].
ig. 8. X-ray diffraction pattern showing polycrystalline thin film of NpO2 on LAO40].
lms of different actinide containing materials. In the end this effortound that cubic YSZ substrates (a = 5.139 A) result in quality epitax-al films of both NpO2 and PuO2. The XRD rocking curve is an excel-ent indication of film quality, and is obtained by fixing the detectort the expected Bragg angle and measuring the diffracted intensitys the sample is independently rotated or “rocked”. The full widtht half maximum intensity of the rocking curve is an indication ofhe dislocation density in the film. The rocking curve (0.07◦) of PuO2n YSZ is shown in Fig. 10, and shows a very crystalline thin film. Inddition, the �-scan of the PuO2 film on YSZ is shown in Fig. 11 andhows a high degree of epitaxy. The �-scan is obtained by fixing theetector at a Bragg angle corresponding to an out-of-plane reflec-ion, and rotating the film through 360◦. For example, in the case ofuO2 on YSZ the [2 2 0] family of reflections of both the film and sub-trate are measured as the film is rotated one revolution. The align-ent of the PuO2 and YSZ [2 2 0] reflections is a very good indication

f how well the film is aligned to the substrate. Taken together the
ocking curve and �-scan show these films to be highly epitaxial.
The YSZ substrate provides advantages over LAO, which canerminate as either Al or La oxide layers and undergoes a phase

Fig. 9. X-ray diffraction pattern of epitaxial UO2 and NpO2 layered on YSZ

Fig. 11. �-Scans of the 2 2 0 reflections of PuO2 (red) and YSZ (black) showingalignment of the film and substrate [40].

transition during the annealing process leading to film strain. Even
though LAO has a more favorable lattice match than YSZ for NpO2,the latter results in better quality NpO2 films suggesting that latticematch is only one consideration when choosing a substrate.
. The inset shows the rocking curve of the bilayer epitaxial film [40].


F�

3

fisaapgsth(r

o(tostoaYT3aofl1ppfpcsat

ig. 12. X-ray diffraction pattern of UC2 film on YSZ substrate. The inset pictures-scans showing film epitaxy.

.4. Uranium carbide films

The synthesis and characterization of uranium carbide, UC2,lms demonstrates the importance of the underlying substratetructure in stabilizing the epitaxial film [42]. At room temper-ture UC2 exists in a tetragonal phase, I4/mmm, with a = 3.633 And c = 6.036 A. A phase transition at 1765 ◦C takes UC2 to a cubichase, Pa3, with a = 5.472 A (3.868 A diagonal). First attempts torow uranium carbide on LAO (a = 3.789 A, pseudocubic) and c-cutapphire (a = 4.758 A, hexagonal) gave only black, highly conduc-ive amorphous materials. Finally, highly reflective films of theigh temperature cubic phase of UC2 were grown on cubic YSZa = 5.139 A; 3.634 A diagonal, Fm3m), and the X-ray diffractionesults are shown in Fig. 12.

As can be seen from the diffraction pattern, the appearancef (0 0 2) and (0 0 4) diffraction peaks of UC2 together with the2 0 0) and (4 0 0) YSZ peaks indicates that the UC2 film is preferen-ially oriented with respect to the substrate normal. The in-planerientation of the UC2 film with respect to the major axes of the YSZubstrate is measured by XRD �-scans on both the UC2 (1 0 1) andhe YSZ (1 0 1) reflections. The inset in Fig. 12 illustrates the �-scanf the (1 0 1) UC2 film. Four peaks of the (1 0 1) UC2 reflections at 90◦

part are reasonable considering the crystal structures of the cubicSZ (a = 0.514 nm) and tetragonal UC2 (a = 0.352 nm, c = 0.598 nm).he lattice mismatch between the YSZ and the UC2 is only about.1% (after a 45◦ rotation of the basal plane). The films were grownfter spin coating a solution of PEI polymer and EDTA complexf uranium adjusted to pH 7. The films were then heated under aow H2/ethylene at 350 ◦C to drive off the organics, and then to000 ◦C for annealing. It is quite remarkable that the high tem-erature phase of UC2 formed at only 1000 ◦C, and that the phaseersists once cooled to room temperature. The best explanationor this is that the YSZ lattice is stabilizing the high temperaturehase in UC2. These films appear to be stable under ambient
onditions for months at a time, keeping their black reflectiveurface. Only when heated to 300 ◦C does the film begin to oxidizes evidenced by in situ X-ray diffraction, and additional heatingo 900 ◦C fully oxidizes and anneals the film to U3O8. This phase
Fig. 13. (Left image) Picture of UC2 film with camera lens reflection; (middle) p

Fig. 14. (a) X-ray diffraction pattern of UN2 film on LAO; (b) �-scan showing epitaxyof the UN2 film.

transition is accompanied by a loss of the black sheen of the UC2film, and is replaced by a light brown color of the very thin U3O8film (Fig. 13).

3.5. Uranium nitride films

Films of UN2 were recently synthesized in our laboratories, andthe characterization data is presented here for the first time. Thecrystal structure of UN2 (a = 5.301 A, Fm3m) is isostructural to UO2,and also grows high quality films on the LAO substrate. However,the nitride films grow on the (0 0 1) facet of hexagonal LAO, and the(1 1 1) face of UN2 is parallel to the (0 0 1) face of LAO (Fig. 14a). The�-scans show the films are epitaxial (Fig. 14b). We have also triedto grow UN2 films on c-cut sapphire, which also has a hexagonal
crystal structure and (0 0 1) face. In this case the UN2 film was poly-crystalline. The oxidation of these UN2 films occur at 200–300 ◦Cin a similar fashion to the UC2 films, forming UN2−xOx films asobserved with in situ X-ray diffraction.
artially oxidized UC2 film, UC2−xOx; (right) fully oxidized U3O8 film [41].

istry

4

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B.L. Scott et al. / Coordination Chem

. Photoelectron spectroscopy

Photoelectron spectroscopy (PES) provides a direct measuref the electronic structure of materials. Within an independentarticle approximation one may associate the density of electronictates with the energy distribution curves that are the measuredutput from angle-integrated photoemission. Angle integratedhotoemission measures the intensity as a function of kineticnergy of the photoemitted electrons. Conversion of the mea-ured electron kinetic energy to binding energy is made throughssessment of the spectrometer work function. The photon energyependence of the photoionization cross section provides the dis-ribution of specific orbital character (s, p, d, and f) in the spectrumFig. 15), and the mean inelastic scattering length informationrovides ways to distinguish bulk from surface features.

Variants on the PES concept provide additional informationncluding resonance photoemission (RESPES) which can be used tonhance specific orbital character via transition through an appro-riate (dipole selection rule) core level region. RESPES is a mostelpful tool in studying f-electron systems. Specifically, the reso-ances most utilized here are 4d → 4f in e.g. Cerium or 5d → 5f inranium systems. Here, the excitation by photon of specific energyan promote an electron directly from the 5d level to the 5f level.his excited state decays with emission of an electron. The finaltate of this process is identical to a direct emission from 5f level,hich allows the two final states to resonate. As a result, for the

pecific resonance photon energy the intensity of the 5f emissionan be substantially increased. In anti-resonance, at a different pho-on energy, the same f-weight is suppressed. The technique is mostseful in identifying the location, in terms of binding energy, of the-orbital character in the valence band.

Angle-resolved photoemission (ARPES) determines both energys well as crystal momentum for a material providing the full
mage of the band structure. The ARPES method allows directomparison of measured and calculated electronic band struc-ure in terms of band dispersion, binding energies, effective
ass, momentum-dependent gapping, etc. From the standpoint of

ig. 15. Photon-energy dependent photoemission of native Pu oxide. The ratios of the pmaller than they are at 115 eV (right panel). Also, the choice of 115 eV is dictated by theesulting photoemission spectra will indicate the location of the 5f spectral weight as a strohoton energy.

Reviews 266–267 (2014) 137–154 145

instrument capability, f-electron photoemission with the highestresolution in the VUV range is best performed at public synchrotronfacilities, with table-top laser sources offering reduced capability inthe resolution-photon flux product at a given photon energy. Veryhigh resolution direct laser excitation systems are not suitable forf-electron work due to the negligible f-electron cross section in the5–7 eV photon energy range. Most of the uranium and thorium PESwork included here was conducted at a public synchrotron. Due toauthorization basis and safety restrictions, transuranic materialsare not allowed at public facilities without multiple encapsula-tion schemes which are incompatible with PES measurements.Because of the encapsulation restrictions on transuranics, PES,ARPES and RESPES measurements on transuranics, are conductedusing the laser plasma light source at Los Alamos National Labora-tory [43–45].

In this section we provide results from photoemission measure-ments on thin films of actinide materials. We demonstrate thatthe films obtained by polymer-assisted deposition utilizing latticedpinning are of high quality with single crystal characteristics andthat their properties are consistent with bulk single crystals. Thehigh quality of the sample is essential for ARPES measurementas the technique relies on lattice uniformity of the sample overthe surface area illuminated by the photon beam or the imagedarea of the electron energy analyzer. In the case of our ARPES sys-tem, this area is of order 1 mm. The PAD samples show full latticeregistry with the single crystal substrates over this measurementarea for the ThO2, UO2 and PuO2 samples. For UO2 we show thefirst example of a dispersive 5f band in a Mott insulator, in verygood agreement with hybrid functional calculations. Using the PADPUO2 samples, have demonstrated the first example of band disper-sion measurement in any transuranic material. In UC2 and U3O8we measure the distribution of the 5f spectral weight within thevalence band and we also provide initial results for UN2. Results
for a thin film, mixed stoichiometry, Np oxide as well as a Pu oxideare also presented. The overall agreement between the photoemi-ssion experiments compared to the hybrid functional calculationsis excellent and eliminates the shortcoming of previous actinide
hotoionization cross sections Pu5f/O 2p at 40.8 eV photon energy (left panel) are vicinity or the 5d–5f resonance, which significantly enhances the 5f emission. Theng increase in Pu5f intensity measured at 115 eV photon energy, relative to 40.8 eV

1 istry Reviews 266–267 (2014) 137–154

op

4

iipbiogtdsIPwa

uwgosacergtodUPllmcbtlobe

aepSppadsicnfldsetac

Fig. 16. Angle-integrated spectrum of the UO2 valence band. Most of the U5f orbital


xide calculations showing metallic character when experimentalroperties were observed to be insulating.

.1. UO2 and U3O8 ARPES

While the connected phase diagram for actinide metals clearlyndicates a maximum in complexity for Pu metal [46], the situationn the actinide oxides reaches maximum complexity with the U–Ohase diagram. Most actinide oxides, including UO2 are predictedy density functional theory to be metallic. Experimental evidence

ndicates UO2 is an antiferromagnetic insulator with a band gapf over 2 eV. The f-orbital character of the excitations across theap places UO2 in a Mott insulator category. Our first results forhe electronic structure of an epitaxial film of UO2, provided evi-ence for the dispersive nature of 5f bands. We also observed hightability of the PAD UO2 film over time and ambient environment.n the lattice-pinning scheme which seems appropriate for theseAD samples, the crystalline nature of the film is preserved all theay up to the topmost layers even after prolonged exposure to

tmospheric conditions.Hybridized, dispersive bands are common in itinerant, metallic

ranium compounds [47–49]. One finds hybridization of f-orbitalsith conduction states to be common in f-electron systems at cryo-

enic temperatures. Such 4f or 5f bands may reside in the vicinityf the Fermi level and participate in the construction of the Fermiurface. Additionally, hybridization with conduction bands maylso enable the dispersive character of an f-conduction band. Thease is somewhat different in an insulator like UO2, where onexpects a more localized electronic structure, where f-bands areelatively flat and shifted away from the Fermi level by the bandap energy scale [50,51]. An assessment of the extent of localiza-ion or hybridization in UO2 may be determined by measurementsf the band dispersion in a material. To test the hybridization pre-ictions from hybrid functional theory, the electronic structure ofO2 was investigated in both single crystal samples and epitaxialAD films. One of the advantages of PAD, contrary to other methodsike physical vapor deposition, is the possibility to pin the crystalattice of the epitaxial overlayer by the substrate lattice. The ARPES

easurements will only show dispersive bands for a true singlerystal system, where the momentum information is not obscuredy mixing domains of different orientation. Our comparison of arue bulk single crystal with a thin epitaxial film sample where theattice pinning is realized by a lattice matching within a few percentr less, allows a direct comparison of the electronic structures foroth samples, thus verifying the true single-crystal quality of thepitaxial PAD film.

For the UO2, U3O8, UC2, UN2 and ThO2 ARPES setup we used public synchrotron as the light source for determining thelectronic structure, with angle-resolved photoemission resultsresented here obtained in the Synchrotron Radiation Center intoughton, Wisconsin. Samples were generally kept at room tem-erature to avoid or minimize charging effects. Experiments wereerformed in a vacuum of 3 × 10−11 Torr, and samples were bakedt 80 ◦C for 48 h before measurements. This bake regime wasesigned to remove the residual gases physisorbed on the stainlessteel elements of the vacuum chamber and apparatus, while keep-ng the bake temperature at a level predicted to be below chemicalhanges for the PAD samples. We have also measured samples witho further treatment, and after 40 min of sputtering with a weakux of Ar+ ions. Spectra before and after sputtering are only slightlyifferent in terms of better peak-to-valley ratio for the sputteredample data. As described elsewhere in this work, calculating the
lectronic structure of the insulating f-electron compounds proveso be challenging, with conventional density functional methods,s well as generalized gradient approximation leading to metallicharacter [34,52,53].
character resides in the peak located around 1.5 eV below the Fermi level. Broadstructure of O 2p-related bands is located between 3 eV and 8 eV.

We have directly compared the electronic structure of UO2using hybrid functional calculations with ARPES data [54,55]. Themomentum-integrated spectrum shown in Fig. 16 illustrates theposition of the main 5f peak around a binding energy of 1.5 eV, andthe oxygen O 2p manifold at a higher binding energy between 4and 8 eV.

Two novel and interesting features have been observed in theangle-resolved data for UO2, shown in Fig. 17.

First, the thin epitaxial film of UO2 grown by PAD is extremelyrobust, as demonstrated by electronic structure of the film beingentirely consistent with the one obtained for a bulk single crystalsample cleaved in vacuo. Panels 1 and 2 in Fig. 17 show the com-parison of the electronic structure of UO2 between a single crystal(panel 1) and thin film (panel 2). The two figures look very sim-ilar, with the peak shape, binding energy and dispersion of theorder of 100 meV (marked by S-like arrows). This direct compar-ison between PAD and bulk UO2 electronic structure demonstratesthat the quality of the thin film is extremely high and comparable toa freshly cleaved bulk crystal. The top-most layers of the UO2 PADfilm, in spite of being exposed to air for 6 months, still maintainedboth the crystallinity and stoichiometry of UO2 without significantsurface degradation. This high quality of PAD films and the pos-sibility of growing them on slightly curved substrates, opens newpossibilities for industrial UO2 applications.

Second, dispersion of approximately 80 meV within the 5f dom-inated band was observed. As illustrated in Fig. 17, the dispersion ofboth oxygen and uranium-related states is significant and consis-tent with a single crystal lattice over the entire photon illuminationarea. This finding may not be surprising for the oxygen-dominatedbands, but it is most unusual and interesting for the uranium 5fstates in a Mott insulator, where the 5f states are expected to belocalized and thus decoupled from the crystal lattice potential.The existence of the dispersion in this part of the ARPES spec-trum demonstrates that there is a small fraction of Uranium 5felectrons that are indeed coupled to the periodic potential of thelattice, in spite of the insulating nature of UO2. This is the firstobservation of a dispersive band in an actinide insulator compound,where the lower Hubbard band was expected to be localized and
non-dispersive.
A portion of the calculated electronic structure for UO2 is shownin Fig. 18. The normal direction in the fcc Brillouin zone (BZ) was


Fig. 17. Results of the angle-resolved valence band photoemission study of UO2. First panel illustrates the dispersion of the dominant 5f electron feature in the valence band,measured in the single crystal UO2, cleaved in vacuo. Band dispersion of the order of 80 meV is indicated by arrows. This electron band dispersion is commensurate withthe reciprocal lattice in terms of symmetry, which indicates single crystalline character of the investigated surface and its very good coupling to bulk properties. Panel 2illustrates the same electronic structure feature as in panel 1, but measured in the UO2 thin film epitaxially grown on LaAlO3 substrate. From the similarity of the dispersionsand binding energies of characteristic features between single crystal and thin film, we have high confidence in the excellent single-crystal quality of the PAD film. Thef d locaf in Brilo

scd

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ull valence band of thin film UO2, including dispersive peaks within the O 2p banor the normal emission direction and increased angle corresponds to directions

f X point [37].

elected and as a consequence due to the high symmetry of therystal structure, the measured electronic structure in the normalirection is identical to the in-plane direction.

This means that the electronic structure measurement withRPES can be performed along a very well defined surface normalirection by changing photon energy, and calibrated with dataccumulated in plane, by changing the electron energy analyzerollection angle. Similarity between the two measurements is aompelling argument for the high quality crystalline nature of theamples, as well as evidence for the lack of surface-related effects.n this particular direction, called � –X in the fcc BZ, we find veryood agreement between the calculated and measured band struc-ures. Both the crystal momentum dispersion and binding energiesf the bands are in very good agreement, which serves to confirmhe role of hybridization in the formation of the electronic structure
f UO2.
A direct comparison of the electronic structure for UO2 and U3O8s made using photoemission. Fig. 19 illustrates the differences inhe valence band between these two oxides.

ted around a binding energy of −5 eV, is shown in panel 3. In all panels, 0◦ standslouin zone moving away from normal emission toward and through the vicinity

It may be noted that the dominant 5f peak is very similarbetween the two in terms of width and binding energy, with onlya slight shift of the order of 100 meV toward higher binding ener-gies in U3O8 with respect to UO2. The O 2p manifold is significantlystronger in U3O8 compared with UO2 which would be expectedfrom the stoichiometry of the compounds to some extent but themagnitude in the resonance PES data may indicate additional influ-ences on the data. Due to a slight charging effect, the binding energyscale for the U3O8 spectrum needed to be calibrated by extrapo-lating the binding energy of the dominant 5f peak to zero photonflux (zero charging). This calibration was accomplished by col-lecting seven spectra for different photon flux values. The chargingbehavior also indicates that the thin films of U3O8 are slightly lessconducting at room temperature than UO2.

The relative change in f–p hybridization between the two ura-
nium oxides may be extracted from the difference curve shownin Fig. 19. In both resonance and anti-resonance, the p orbital fea-tures are of very similar intensity. However, the 5d–5f resonanceenhances the 5f features at 108 eV compared to the lower photon


Fig. 18. Band structure of UO2. Comparison of the calculated and measured bandsps

eovioO

Fetd

tructure along the surface normal direction in UO2. The location of the measuredhotoemission peaks are shown with red dots, while the calculated spectrum ishown with lines.

nergy. By subtracting the data from the two photon energies webtain a difference curve that greatly enhances the 5f feature in the
alence band. Comparison between UO2 and U3O8 shows the bind-ng energy and width of this feature is very similar between the twoxides. At higher binding energy, there is more 5f mixing into the
2p manifold in U3O8 than is observed in UO2. The position of the

ig. 19. Resonant photoemission results. Left panel shows the comparison of the 5d →nhanced, and similarities between the two oxides are seen, except of differences in

he resonance–antiresonance difference spectra, effectively showing that there is no siistribution.

Fig. 20. Left panel: Angle-integrated PES (blue) and HSE calculations (red) for PuO2.Right panel: orbital character breakout for the HSE calculations on PuO2.

dominant 5f feature in both oxides determines the lower boundaryof the Mott gap – however, it should be noted that the size of theband gap depends on the position of the lowest unoccupied bandwhich is not measured by photoemission and in principle, theredoes not have to be symmetry between the position of the Fermienergy with respect to the highest occupied and lowest unoccupiedbands in a Mott insulator. In practice, the Fermi energy in an insu-lator is often pinned near the mid-gap level by defects at or nearthe surface unless intentionally doped to be near one of the bandedges. Thus, while we cannot conclude from PES that the Mott gapis the same in both oxides, it is likely to be similar.

4.2. PuO2 ARPES

The general trend in the actinide series is a contraction of the5f electron radial wavefunctions with increasing atomic numbersimilar to the Lanthanide contraction in the 4 fs. This is due tothe incomplete shielding of the nuclear charge with larger Z assuccessive electrons are placed in the f-orbitals; i.e. each electronin the f orbital sees a more positive effective charge as Z increase.
This contraction tends to further localize the 5f levels, but alsoleads to an orbital energy stabilization; that is, each additionalelectron becomes more difficult to ionize. With our observation ofsmall hybridization and dispersion in the ARPES data between the
5f resonance spectra collected with 108 eV photons. The 5f spectral weight isspectral weight related to oxygen stoichiometry. The right-side panel illustratesgnificant difference between the two oxides in terms of the U5f spectral weight


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4.4. UC2 and UN2 ARPES

For the thin films of UC2 and UN2 the photoemission mea-surements we conducted at the Synchrotron Radiation Center and

ig. 21. PES vs. HSE for Th through Pu oxide. The overall agreement between PES aoward PuO2 is observed through the sequence.

ranium 5f and O 2p levels of about 100 meV, one might expectess or even no hybridization going toward Np and Pu oxides.his trend is, however, interrupted by a degeneracy in the bindingnergy of the Pu 5f and O 2p levels.

The PES data for PuO2 illustrates the point clearly in Fig. 20,ith the merging of the 5f and 2p energy levels. Now, instead of

wo separated orbital energy intervals we have three peaks in bothhe calculation by the HSE method as well as three peaks in the PESata for PuO2.

To further illustrate the point the right panel of Fig. 20 showshe orbital character of the energy levels and it is observed that the

2p leaves are spread evenly throughout the three peak struc-ure while the Pu 5f character is also dispersed throughout theybridization region with a stronger showing in the first 2 eV ofhe 4 eV hybridized region. The extent of the hybridization is docu-

ented in the ARPES data for PuO2 where there is 3 eV of dispersionn the PuO2 data compared with 100 meV in the UO2 data [56].

.3. Actinide oxide comparisons

The details of the bonding and hybridization in actinide oxidess dependent on more than just the constituent elements in thectinide oxide insulators. In Fig. 21 we show a sequence from Th tou with the dioxides as well as Pu2O3.

The PES data is indicated by blue points and the HSE calculations the red line. These data were collected at a photon energy of about0 eV where the photoionization cross sections for the O 2p (6.8 Marns) and actinide 5f (U = 6.0 M barns) is approximately equal. Alsot this energy, the actinide 6d (Th = 1.5 M barns, U = 0.68 M barns)evel is at least a factor of five below the 5f/2p cross sections andhe actinide 7 s cross section (Th, U = 0.07 M barns) is too low toe observed in this photoemission data [57]. Thus the PES data inig. 21 shows primarily the O 2p level with some of the 6d for thehO2 while the U and Pu oxides show 2p and 5f levels. Across theeries we observe that the O 2p bands are about 4 eV wide whilehe 5f states are between 1.5 and 2 eV wide. It is interesting to notehat the strong hybridization in PuO2 does not yet exist in Pu2O3o that the conditions for strong hybridization are dependent notnly on the constituent elements but also their oxidation state, andrystal structure/lattice spacing.

It is also clear from the data that the electronic structure of PuO2s distinctly different from that of Pu2O3. In particular, the Pu2O3ata still has a significant separation between the Pu 5f and O 2p lev-ls in the electronic structure. We make use of this fact in analysisf other thin film and native oxide layers in the actinide materials.n Fig. 22 we show the known end point materials UO2 and PuO2
long with two native oxide structure for Pu and Np which occurn the actinide metals. From the data in Fig. 21 showing both PuO2nd Pu2O3 one may observe that the thin film data in Fig. 22 for Puxide is a superposition of both PuO2 and Pu2O3.
E is very good and in particular, the reduction in O 2p vs. 5f level gap as one moves

It is known that both Pu and Np form both a sesquioxide anda dioxide on the metal surface. From this thin film data it appearsthat the Np oxide in Fig. 22 is also a superposition of the Np dioxideand sesquioxide. The data in Fig. 22 also then serve to reinforce thenotion that the PAD films are of high quality and single phase withpure AnO2 character as observed in the photoemission.

Fig. 22. Mixed oxides of Np and Pu with the endpoints of UO2 and PuO2 to showthe transition in thin films from sesquioxide to dioxide as well as the hybridizationoverlap in PuO2 compared with UO2. The Pu and Np oxides are not PAD films.


Fig. 23. Photoemission spectrum of a PAD film of UC2. Top panel: Two spectra, at92 eV and 98 eV photon energy, correspond to the 5d–5f anti-resonance and reso-nit

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Fig. 24. Resonance photoemission of a PAD film of UN2. Light bake in vacuum

ance conditions, respectively. Enhancement of the 5f emission at resonance allowsdentification of the dominant 5f feature centered around 1 eV binding energy. Bot-om panel: identification of a metallic Fermi level (see text for details).

overed the photon energy range of 108–92 eV, to take advantageor the U 5d–5f resonance/anti-resonance (98 and 92 eV) includinghe second resonance energy for uranium at 108 eV. The samplesere cleaned in situ by thermal desorption or Ar ion sputtering.

he measurement temperatures were room temperature or 8 K. Theata was taken with a variable polarization undulator set to verti-al photon polarization. Data was also collected in the horizontalonfiguration which showed a substantially smaller U 5f comparedo the C 2p emission. The difference in the data collected from hor-zontal and vertical polarizations gives additional solid indicationf the high quality of the PAD samples.

Fig. 23 shows the ultraviolet photoemission (UPS) valence bandpectrum for UC2. The spectrum is dominated by a manifoldikely composed of primarily C 2p spectral weight extending frompproximately 3 to 10 eV below the Fermi energy (EF).

As determined by the spectral weight distribution at resonancetop panel), the U 5f weight is centered at approximately 0.6 eVelow EF but extending up to the Fermi energy. There is also a sig-ificant component of C 2p in the first 2 eV from the Fermi energy,hich indicates hybridization in UC2 while small, is present, simi-

arly to the case of UO2 discussed above. We also note that the U 5f
inewidth is ∼1 eV wide while the C 2p width is at least 4 eV wide,
hich indicates substantially stronger 2p bonding compared to 5fonding in UC2 again consistent with the finding in UO2. A Fermi

evel cut-off is observed in the near – EF region, as shown in the

removes the surface contamination from the thin film. With clean surface, the Fermilevel, as well as the characteristic 5f feature are both recovered from the air-exposedPAD film.

bottom panel of Fig. 23, and it is an indicative of a metallic 5f fea-ture cut by EF. This observation confirms the metallic character ofUC2.

The initial data on UN2 shown in Fig. 24 were provided by pho-toemission measurement at the 5d–5f resonance energy of 108 eV.Here a cleaning progression is observed, with the cleaner samplealready showing a metallic Fermi level (EF) and a dominant uranium5f feature.

This system was found to be the most surface sensitive of all thePAD materials discussed herein, being prone to surface contami-nation including but not limited to oxidation and chemisorptionof vacuum contaminants. A light overnight bake at 100 C in vac-uum seems to be a good starting point to obtaining a metallicsurface. While additional studies of the UN2 system are requiredand underway, it is clear that there are systematics which showcommon actinide-ligand bonding character throughout the UN2,UC2 and UO2 systems with the broader ligand feature centered wellbelow the narrower and less hybridized uranium 5f feature whichforms the highest occupied band state in these materials. Finally,it is noted that both UN2 and UC2 PAD films showed some spec-tral intensity at the Fermi level which demonstrates a PAD metallicactinide capability. It is not clear if this spectral weight at the Fermilevel in UC2 and UN2 contains any significant 5f character or if it isof spd valence band character. Also, these first measurements showa small intensity at the Fermi level which may be an indication ofeither a poor metal or a semiconducting state in these materials.These materials will certainly be studied in greater depth as thereis much interest in the nitrides and carbides for next generationnuclear fuels. By using PAD samples, we have the opportunity toexplore regions of materials phase space that are not accessible bybulk sample growth techniques.

4.5. Theory–experiment comparison

Overall there is very good agreement between experiment andtheory, particularly when the indeterminate nature of the bandgap pinning position in photoemission is considered with respectto the relative energy references of the calculated reference point.
The PAD samples on AnO2 materials enabled a critical test of theapplication of hybrid functional calculations to actinide oxides. Inparticular, while UO2 samples have been available for some time,there have never been PuO2 single crystals of sufficient size to

istry Reviews 266–267 (2014) 137–154 151

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nable ARPES experiments. With high quality PAD PuO2 samplesvailable for the first time, the large difference in hybridization ofhe 2p and 5f levels between UO2 and PuO2 was observed in ARPES

easurements and provided solid evidence that the hybrid func-ional calculations capture the essential properties of actinide oxidelectronic structure.

We find that the polymer-assisted deposition of epitaxial filmsroduces actinide oxide films of very high quality and high sta-ility. Additionally, the entire range of actinide compound films

ncluding the UC2 and UN2 hold great promise. Dispersion of theominant 5f bands are demonstrated for the first time in an actinideott insulator system, pointing toward hybridization as a leading

rocess in establishing the electronic structure of actinide oxides.ybrid density functional theory now allows one to reliably calcu-

ate the electronic structure of actinide oxides in good agreementith experiments.

. Optical band gaps

One of the key fundamental parameters predicted by theory ishe band gap of a material. Experimental verification of the bandap provides an important validation of theory. In some cases, suchs the DFT + U approximation, an experimental band gap can fix the

parameter and then additional experimental parameters such asefractive index, or lattice constants can be used to validate theheory with the same fixed U parameter. The value of U can be crit-cal in predicting properties, for example, as U is varied from 0 to

the behavior of NpO2 goes from metallic to a band gap >3.0 eV.he value for U also affects dynamic mean-field theory (DMFT),hich builds upon a DFT + U base [58,59]. Experimental validation

equires high quality solids such as single crystals or epitaxial films.pitaxial films on transparent substrates offer the advantage thathey can be made thin enough to perform direct optical transmis-ion spectroscopy to determine the band gap.

.1. UO2 and U3O8 band gap measurements

For UO2 there have been direct optical measurements of theand gap based on single crystals. In 1978 Schoenes determinedhe band gap to be 2.1 eV based on polished single crystals [60].or U3O8 the experimental work is more challenging. Single crys-als can be difficult to prepare as U3O8 decomposes into UO2 at200 ◦C prior to melting. The U3O8 system is further complicatedy multiple phases. The room temperature stable orthorhombic

phase undergoes a transition at 210 ◦C to the high temperatureexagonal phase. Epitaxial films of UO2 and U3O8 have been pre-ared from the PAD solution route and used to measure the opticaland gaps. As a control a standard film of Fe2O3 was depositedy the PAD process on c-cut sapphire and the measured band gapetermined to be 2.25 ± 0.05 eV as previously reported in the lit-rature [61]. The optical band gap of UO2 determined from opticalransmission of the thin film was 2.2 ± 0.1 eV, in agreement withchoenes. Epitaxial films of orthorhombic (C2mm) U3O8 have beenbtained on R-plane sapphire and hexagonal (P62m) U3O8 on c-ut sapphire substrates (vide supra). The optical transmission for3O8 on c-plane sapphire has been measured and is best fit to an

ndirect band gap of 0.6 ± 0.1 eV. In the case of an indirect band gapemiconductor

(�ω) ∝ (�ω − Egap)2

�ω

nd a plot of sqrt(ε × hω) vs. hω yields a straight line with an
ntercept of the band gap Eg. The plot for U3O8 is shown in Fig. 25.he value of 0.6 eV is consistent with the black color of bulk U3O8.
screened hybrid and DFT + U study of U3O8 has been carriedut on both the ground state C2mm phase and also the high

Fig. 25. Optical band gap of U3O8.

temperature P62m phase [62]. Both HSE and PBE + U predict thehexagonal phase of U3O8 to be a metal, in contrast to the measuredvalue of 0.6 eV for the band gap. This most likely reflects the addi-tional challenge of mixed oxidation states to the DFT localizationproblem, and additional studies are needed for such systems.

5.2. PuO2 and NpO2 band gap measurements

For NpO2 and PuO2 experimental work is complicated by thehigh melting points of the oxides; 2547 ◦C and 2400 ◦C respectively.There are only two reports of PuO2 single crystals prepared from aflux comprising a mixture of Li2O and MoO3. The color of the crys-tals obtained from this process was black in contrast to the olivegreen color of bulk PuO2 powders [13,14,63]. Polishing of such crys-tals to the thickness needed for direct transmission experiments todetermine the optical band gap has not been done. For neptuniumoxide there are a few reports of very small single crystals border-ing on polycrystalline material [64,65], but only one report of largecrystals that have metal impurities [12]. However, the direct opticalband gaps have not been measured on these crystalline materials.

For both NpO2 and PuO2 most references citing band gaps actu-ally refer to papers in which the experimental measurement wasthe activation barrier to electronic conduction. In these cases, theelectronic conductivity of a pressed pellet is measured as a functionof temperature [66,67]. Activation barriers to electronic conductiondo not always contain data near room temperature, because theconductivity values for insulators such as NpO2 and PuO2 drop offsignificantly as the temperature is lowered. These measurementsare not a direct measurement of the optical band gap and are verysensitive to defects and vacancies. For example, copper selenidethin films with an optical band gap of 2.03 eV have been reportedto have an activation barrier to electronic conduction as low as0.20 eV [68].

The first direct measurements of the optical band gaps for NpO2or PuO2 have been reported recently by us using epitaxial thinfilms made from the PAD technique [41]. Optical band gaps ofboth PuO2 and NpO2 were measured and replicated with differ-ent film thicknesses by coating multiple times. For a direct bandgap semiconductor:

˛(�ω) ∝√�ω − Egap

�ω

Plots of (ε × hω)2 vs. hω, used for direct band gap semiconductors,show linear fits with intercepts of the band gap at 2.80 and2.85 ± 0.1 eV for PuO2 and NpO2 respectively (Fig. 26).

These direct measurements of the optical band gap differ signif-icantly from the literature cited values of 1.8 eV and 0.4 eV for PuO2and NpO2 respectively [66,67]. For PuO2, the only the data greaterthan 200 ◦C was fit to a straight line and the measurements were


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Fig. 27. The temperature dependence of the resistance of a typical epitaxial UC2 filmon the YSZ substrate. To convert the measured resistance into resistivity, one needsto consider the geometric factor. A value of 20 ± 10 × 10−6 cm has been estimatedbased on the sample width, length, and film thickness. This would correspond to a

in PuO2 [71]. The experimental results suggest U should be fixedat 3 eV or less.

ig. 26. Optical band gap of PuO2 [40] and NpO2 based on absorption with 3 (solidine) and 5 (dashed line) coats of PuO2 and NpO2 on YSZ.

one under vacuum. The value of 1.8 eV obtained from this dataork has been referenced as a band gap value and has become a

enchmark for all theoretical work used to model PuO2. The articleotes “The lack of single crystal specimens also places somewhatf a limit on the interpretation of results” [67]. The value of 1.8 eVomes from a fit of data from 200 to 1000 ◦C. At temperatures below00 ◦C the data rapidly falls far below the extrapolated line. The datarom room temperature to 200 ◦C are nearly linear. A fit to this loweremperature data gives an electronic conduction of approximately

eV in much closer agreement to the optical band gap measuredy optical transmission. In the case of NpO2 measurements wereone over a temperature range of 150–300 K, and once again theata do not fit a straight line over the entire energy range. Thealues of 2.80 eV for PuO2 and 2.85 eV for NpO2 based on opticaland gap measurements values are consistent with the olive greenolor of the bulk oxides. Any material with a band gap as low as.8 eV (680 nm) or 0.4 eV (3100 nm) as cited in the past would bexpected to appear black at any reasonable crystallite size. In addi-
ion, we have been able to determine the real and imaginary opticalonstants of the refractive index as a function of wavelength usinglipsometry with values of n = 2.1 and k = 0.012 at 500 nm.
resistivity of 300 �� cm. The high error is associated with estimating the geometricfactor for a thin film [41].

These results of experimental measurements of optical bandgaps on PAD films provide a new benchmark for theoretical calcu-lations. The screened hybrid DFT approximation (HSE) yields gapsof 3.1 eV and 2.7 eV for NpO2 and PuO2 [69], in excellent agree-ment with our measurements of 2.85 and 2.80 eV, respectively.The hybrid DFT calculations were spin polarized (unrestricted) andfound an antiferromagnetic coupling solution as the ground state.Although these early HSE results did not examine the influenceof spin-orbit coupling on the computed band gaps, recent work[55] using an approach identical to that in Peralta et al. [70] showsthe effect is minimal. Inclusion of spin-orbit coupling within HSEyields slightly reduced values of 2.99 and 2.58 eV for NpO2 andPuO2. A GGA + U calculation, specifically PBE + U, U = 4.0 eV, findsvalues of 2.6 and 1.6 eV for the band gaps of NpO2 and PuO2. Thevalue for PuO2 is significantly lower than our experimental result.A material with a band gap this low would be black in the bulk asopposed to the pistachio green color observed for PuO2. The SICapproximation yields 2.3 eV and 1.2 eV, respectively, again lowerthan the experimental value [58]. To our knowledge, the gap forNpO2 from DMFT has not been reported, but for PuO2 the DMFTresult depends upon the value chosen for the parameter U. DMFTyields either 3.5 eV (U = 6 eV) [71] or 2.5 eV (U = 3 eV) for the gap

Fig. 28. Temperature dependence of electrical resistivity of UN2 on c-cut sapphire.

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[37] A.K. Burrell, T.M. McCleskey, P. Shukla, H. Wang, T. Durakiewicz, D.P. Moore,C.G. Olson, J.J. Joyce, Q. Jia, Advanced Materials 19 (2007) 3559.

[38] D.A. Andersson, G. Baldinozzi, L. Desgranges, D.R. Conradson, S.D. Conradson,


.3. Electrical conductivity in UC2 and UN2 films

The conductivity of UC2 and UN2 epitaxial thin films pre-ared using the PAD technique was measured by a standardour-probe technique. For UC2 the resistance of the film increasesith decreasing temperature, indicating a semiconductor-like

haracteristic (Fig. 27). This is consistent with our HSE screenedybrid DFT calculations that show UC2 to be a semiconductor,ith a narrow gap of 0.4 eV [72]. The UN2 results also show a

emiconductor-like characteristic (Fig. 28). However, the resis-ance increases at a much faster rate at temperatures below 40 K,nd further characterization is needed to fully understand thelectrical conduction in UN2. Theoretical calculations employingybrid DFT show a semiconducting ground state for UN2.

. Conclusions

Over the last five years we have used polymer assisted depo-ition to successfully synthesize several epitaxial thin films ofctinide oxides (UO2, U3O8, ThO2, PuO2, and NpO2), and also UC2nd UN2. The quality of these films in terms of purity and singlerystal nature have been demonstrated through measurementsncluding high resolution X-ray diffraction and TEM, and angleesolved photoemission spectroscopy (ARPES). The ARPES mea-urements have yielded unprecedented data and understanding ofhe ligand p and actinide f hybridization in all of these materials.he amount of dispersion in the 5f bands of these materials wasarger than expected. Even though ARPES has been measured onO2 single crystals before, the data presented herein on PAD films

s the first to show f band dispersion in an actinide insulator. ARPESas measured on PuO2 for the first time, and the p–f hybridizationas much larger than in UO2 leading to a larger f-band dispersion

3 eV vs. 100 meV for UO2). In addition, the thin films were usedor electrical conductivity and optical band gap measurementsor some of the materials, with more measurements currentlynderway. These films allowed the optical band gaps of PuO2 andpO2 to be measured for the first time, and showed great improve-ent over previously published values based on indirect methods,

llowing for detailed comparison with theory. For the actinideioxides studied in this work, the agreement between experimentnd hybrid density functional theory is excellent. In UN2 andC2, hybrid DFT finds semiconducting ground states, consistentith the optical spectroscopy and conductivity measurements.

n contrast, the PES measurements show semiconducting to pooretallic behavior. Thus, additional characterization work on theseaterials is warranted. For very strongly correlated metals, such

s UN, agreement with experiment is not so gratifying. Althoughot discussed here, screened hybrid DFT correctly predicts metallicehavior for the mono-nitride UN, in agreement with experi-ent, but has difficulties reproducing details of the experimental

hotoemission spectrum [31].In conclusion, new synthetic and characterization tools have

een applied to the study of the electronic structure of actinideaterials. The results challenge existing theoretical approxima-

ions and allow for a critical assessment [31]. The capability toynthesize actinide materials with selected facets opens the dooro further studies into their reactivity, aging, nuclear fuel behavior,nd environmental fate.

cknowledgements

This work was supported by the LANL Laboratory Directedesearch and Development program, and the Department of Energyasic Energy Sciences program. QXJ acknowledges the support fromhe Center for Integrated Nanotechnologies, a US Department of

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Reviews 266–267 (2014) 137–154 153

Energy, Office of Basic Energy Sciences user facility ay Los AlamosNational Laboratory.

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High quality epitaxial thin films of actinide oxides, carbides, and nitrides: Advancing...

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