Article pubs.acs.org/JPCC
Conversion Reaction of FeF2 Thin Films upon Exposure to Atomic Lithium Sylvie Rangan,*,† Ryan Thorpe,† Robert Allen Bartynski,† Mahsa Sina,‡ Frederic Cosandey,‡ Ozgur Celik,§ and Daniel D. T. Mastrogiovanni§ †
Department of Physics and Astronomy and Laboratory for Surface Modification, Rutgers University, 136 Frelinghuysen Road, Piscataway, New Jersey 08854, United States ‡ Department of Materials Science and Engineering, Rutgers University, 607 Taylor Road, Piscataway, New Jersey 08854, United States § Department of Chemistry and Chemical Biology, Rutgers University, 610 Taylor Road, Piscataway, New Jersey 08854, United States ABSTRACT: The chemical and structural phase evolution of ultrathin (∼5 nm) FeF2 films upon deposition of atomic lithium in an ultrahigh vacuum environment has been studied using X-ray and UV photoemission spectroscopies, inverse photoemission spectroscopy, and transmission electron microscopy in an effort to explore the fundamental properties of the conversion reaction of this promising Li battery cathode material. Spectroscopic measurements show reduction of FeF2 into a metallic Fe0 phase and a LiF phase upon Li deposition. No other phases are detected. Transmission electron microscopy reveals extensive changes in the film’s morphology and material reorganization upon full lithiation. The initial FeF2 film, with grains on the order of 10 nm in diameter, phase separates into smaller (∼3 nm) interconnected Fe0 regions surrounded by LiF. This structural modification is attributed to the large Li+ ionic mobility with respect to Fe2+. The intrinsic nanoscale texture of the final phases is believed to aid in accommodating the extensive structural transformations that occur in this conversion reaction material during an electrochemical cycle in battery applications.
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INTRODUCTION With the rise in popularity of battery-operated portable electronic devices, hybrid vehicles, and plug-in electric vehicles during the past few decades, the need for high capacity secondary (rechargeable) batteries has become apparent. In contrast to microelectronics, batteries have failed to keep pace with Moore’s law and are the limiting factors in many modern devices, motivating the search for new types of battery materials.1 Li-intercalation batteries are the most common energy storage devices in modern portable electronics but are limited by the number of host sites for topotactic Li insertion into the cathode material. Conversion reaction materials offer an alternate approach to intercalation, utilizing the full charge state Mx+ of a cathode metal ion by reducing it to the M0 charge state. This implies not only a larger amount of inserted Li+ ions but also a profound alteration of the initial cathode material, leading to phase separation between a metallic M0 phase and an oxidized Li phase.2,3 Iron fluoride (FeF2 and FeF3) nanoparticles have gained interest as choice materials for conversion reaction-based batteries. Owing to their large band gaps and their ability to store up to three electrons per formula unit (eq 1), batteries using these materials operate at high voltages and high energy densities.4−7 +
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FeFx + x Li + x e → Fe + x LiF © 2012 American Chemical Society
However, as a result of the large alteration of the cathode material while cycling, multiple aspects of the local chemistry remain unclear. In particular, the phase transitions with increasing Li concentration and the effect of ionic mobility on the reorganization of the cathode material are still poorly characterized, as are more subtle effects related to real batteries conditions such as nanoparticle surface defects or the effect of the surroundings (carbon matrix, electrolyte...). In this work, we have studied a model conversion reaction by evaporating atomic Li onto polycrystalline FeF2 thin films as a means to explore the fundamental aspects of the conversion reaction mechanisms. In an electrochemical cell, the initial reactants are electrically isolated by an ion-conducting electrolyte. As such, the chemical reaction (eq 1) is expected to be limited by the mobility and conductivity of both Li+ ions and electrons. In particular, although partially compensated by the nanoscale particle size and the conductive carbon matrix used in battery applications, the insulating nature of FeF2 can hinder the electronic pathway and impose a strong kinetic limit to the conversion process. In the work we present here, deposition of atomic Li on an ultrathin FeF2 film should significantly reduce any kinetic limitations and the thermodynamic energetics of the reaction should dominate. Received: January 19, 2012 Revised: April 11, 2012 Published: April 20, 2012
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obtained using a JEOL 2010F operated at 197 kV and equipped with a Gatan GIF 200 spectrometer. The EELS spectra were obtained with a collection half angle of 27 mrad and convergence angle of 10 mrad and with an energy resolution of 1.1 eV. EELS spectrum images (32 × 32 pixels) obtained with a 0.3 nm probe size and a 1.25 nm spatial resolution were collected to obtain information on Fe and Li elemental distribution. Sample Preparation. The FeF2 samples were prepared by exposing an iron substrate cleaned in UHV by Ar+ sputtering to XeF2 gas at a pressure of 10−4 Torr for 20 min. The growth of FeF2 is believed to follow a Mott−Cabrera mechanism, where F− ion diffusion is facilitated by their image charge creation in the metal substrate.9 This is a self-limiting process leading upon growth of an insulating layer, to a thickness previously estimated to be 5 nm of FeF2 on Fe.10,11 Two sets of samples were prepared for this study. For spectroscopic measurement, FeF2 films were grown on a polycrystalline Fe foil. For TEM measurements, a thin layer of Fe was evaporated on a well-degassed Si−O−N TEM membrane (20 nm thickness, SiMPore Inc.) and then exposed to XeF2. The metallic layer was sufficiently thin so that all Fe was reacted with XeF2. Furthermore, no evidence of membrane etching was observed during exposure to XeF2. After fluorination, the films were briefly exposed to atmosphere before measurements. Only a small amount of iron oxidation was induced by this step, and most adsorbed species were removed by annealing the samples for a few minutes in UHV at 295 °C, below the reduction temperature for FeF2, as verified by XPS. Lithiation was performed in UHV by heating a thoroughly degassed Li getter source (SAES getters) while maintaining the pressure in the low 10−10 Torr. The absolute Li and F coverages were estimated using nuclear reaction analysis (NRA).
Using X-ray and UV photoemission as well as inverse photoemission spectroscopies, the occupied and unoccupied electronic states of the resulting samples have been systematically probed as a function of increasing lithium concentration. Transmission electron microscopy (TEM) has been used in parallel to investigate both the film’s morphology and material reorganization upon lithiation. Despite the very different electronic and ionic conduction channels in an electrochemical cell, as compared to those in this work, the final products obtained in this surface science approach are similar to what has been observed in conversion batteries after discharge.4−7 In particular, FeF2 is reduced to metallic Fe0 without intermediate oxidation state, with the concomitant appearance of an LiF phase. The morphological behavior upon lithiation also indicates particle size reduction, as observed in electrochemical cells.
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EXPERIMENTAL METHODS Spectroscopic Methods. The spectroscopic measurements presented here were obtained using a single ultrahigh vacuum experimental chamber that housed instrumentation for X-ray and ultraviolet photoemission spectroscopies (XPS and UPS) as well as inverse photoemission spectroscopy (IPS) described in detail elsewhere.8 The base pressure of the chamber was less than 5 × 10−10 Torr. Valence band photoelectrons were excited using a Leybold−Heraeus helium discharge photon source (HeII: 40.8 eV), and core levels were probed using the nonmonochromatized Al Kα line of a SPECS XR50 dual anode source. The energy analysis of the emitted electrons was performed in an angle-integrated mode using a double pass Phi 15-255G cylindrical mirror analyzer (CMA). The axes of the photon sources and the CMA formed a 90° angle, and the sample normal was oriented midway between the two. Inverse photoemission spectra were obtained using a grating spectrometer, which was mounted on the same experimental chamber. Briefly, a well-collimated, monoenergetic electron beam (primary energy ranging from Ep = 20.3 eV to Ep = 48.3 eV in this study) was directed toward the sample along the surface normal. The electrons couple to high-lying unoccupied states and a subset relax via a direct optical transition to low lying unoccupied states in the conduction band, emitting a photon in the process. The photons were dispersed by a concave spherical diffraction grating and detected by a microchannel plate with a position sensitive resistive anode encoder. With this approach, the intensity of photons as a function of photon energy reflects the density of unoccupied states in the conduction band. In our photoemission and inverse photoemission spectra, the valence band maximum (VBM) and the conduction band minimum (CBM) are both measured with respect to the Fermi level of a gold sample in contact with the iron foil. The overall energy resolutions for the UPS and IPS spectra are estimated to be 0.3 and 0.6 eV, respectively. The IPS measurements presented in this work have been taken while keeping a small sample current (0.5−1 μA) for a beam size of 1 mm2 and by sampling several spots on a large 1 cm2 sample with short beam exposure (3−10 min). For such electron doses, no beam damage was observed during IPS nor in subsequent UPS and XPS spectra. TEM. Bright-field TEM images, selected area electron diffraction pattern (SAED), annual dark field STEM images (ADF-STEM), and electron energy loss spectra (EELS) were
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RESULTS AND DISCUSSION FeF2 Thin Films Characterization. Although FeF2 is a large band gap Mott−Hubbard insulator, no sign of sample charging was observed in the present measurements, indicating that the FeF2 film was thin enough for the excess charge to escape via the metal substrate. The valence and conduction band spectra measured, respectively, using UPS and IPS are presented Figure 1. The zero of energy is chosen as the Fermi level of the system, so that occupied states have a negative energy and the unoccupied states a positive energy. In this figure, the conduction band is a composite spectrum obtained for different primary energy of
Figure 1. Valence band spectrum measured in UPS (energy < 0) and conduction band spectrum measured in IPS (energy > 0) of a 5 nm FeF2 film on Fe foil. A linear extrapolation of the band edges leads to a measure of the transport gap of FeF2 of 1.9 eV. The energy scale is referenced to the Fermi level of the system. 10499
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electrons (Ep = 20.3 eV to image the region from 0 to 6 eV and up to Ep = 48.3 eV for higher lying unoccupied states).12 The valence band was measured using a photon energy of 40.8 eV. The occupied electronic structure of FeF2 has been explored earlier using XPS,13−18 UPS,19 and X-ray emission.20 However, owing to the existence of conflicting accounts in the literature, the interpretation of our experimental data will be discussed in detail below. The unoccupied states for this material have been probed using X-ray absorption20−22 and more recently EELS.23 In both of these techniques, the spectra are dependent on the core hole created and are governed by dipole selection rules, resulting in a partial description of the unoccupied density of states. Thus, XAS or EELS spectra of the F K-edge (transition of an F 1s electron to the unoccupied states) describe a conduction band weighted by the extent of hybridization between F 2p and Fe levels.20,21 In contrast, in Fe K-edge spectra, the lowest states of Fe 3d character in the conduction band can only be probed with quadrupole transitions.22 At the Fe L2 and L3 edges (electrons excited from the Fe 2p1/2 and 2p3/2 levels), dipole-allowed transitions imply that the Fe 3d states will be prominent and only fluorine states that are hybridized with these Fe states can be accessed.21,23 In IPS, the unoccupied states of all orbital character are probed, only weighted by the cross section of the electronic states involved in the inverse photoemission process. The band edges of FeF2 can be understood in a simple approach by considering a strong octahedral field (or a ligand field if one wants to include fluorine states). Neglecting any exchange splitting, the 3dn high spin ground state t2g(5)eg(2) gives rise to 3dn−1 and 3dn+1 multiplet final states in UPS and IPS, respectively. In UPS, the lowest energy state accessible has 6 A1 symmetry, followed by 4T1, 4T2, and 4E(G) states. The statistical weight of these states with respect to the 6A1 state is expected to be 5:1.24 Thus, in the valence band spectrum, the peak found around −3 eV is attributed to the 6A1 configuration, whereas the feature found around −6 eV is attributed to the 4 T1, 4T2, and 4E(G) states. The last feature with a peak near −9 eV is attributed mostly to F 2p nonbonding states. This assignment is based on the results of X-ray emission for several metal fluorides20 and consideration of a simple ligand field approach. In the IPS final state, three Fe 3d multiplets are expected: the state the closest to the Fermi level having a 4T1 symmetry and two 4T2 and 4T1 states at higher energy and with half of the weight of the first multiplet.22 Unfortunately, these states are too close in energy to be separated given the resolution of our inverse photoemission apparatus and give rise to the threshold line shape between 0 and 4 eV of the conduction band spectrum.24 Above these Fe 3d multiplet states, Fe sp states hybridized with F 2p levels are expected according to a simple ligand field model. Experimentally, this is justified by the remarkable similarities between absorption spectra at the F K-edge for several metal fluorides, thus probing hybridized F 2p states.20,21 As a last comment, it should be noted that because the spectra obtained from UPS and IPS are both referenced to the Fermi level of the system, the distance separating the band edges is directly related to the electronic transport gap. Using a linear extrapolation of the band edges to the spectra background, a transport gap of 1.9 eV is measured, consistent with a ∼1.5 eV edge found in optical absorption measurements.25 The Fe 2p and F 1s core levels measured on the same FeF2 thin film are presented in Figure 2. The Fe 2p line shape is
Figure 2. X-ray Al Kα photoemission spectra of the Fe 2p and F 1s core levels measured on a 5 nm FeF2 film on Fe.
characteristic of FeF2 and is composed of a sum of high spin Fe2+ multiplets and satellites as well as F 1s plasmon loss peaks.16−18 The centroid of the main Fe 2p3/2 peak is found at 711.1 eV, that is 4.5 eV below the Fe 2p3/2 peak position for metallic Fe. Note that no metallic iron from the substrate is visible, consistent with a homogeneous thin film of at least 5 nm thickness. A single feature from the F 1s core level is found at −685.1 eV. TEM measurements (shown in Figure 3) have been performed in order to characterize both the film morphology
Figure 3. (a) ADF-STEM image of the initial FeF2 thin film on a TEM grid. The dark area is the Si−O−N membrane substrate. (b) SAED pattern indicating a rutile structure (P42/mnm). (c) STEM image and the corresponding elemental maps from the FeF2 thin film for (d) Fe and (e) Si relative intensities.
and the crystalline structure obtained by exposing metallic Fe to XeF2. The FeF2 film of Figure 3a, imaged in ADF-STEM mode, appears as interconnected particles with an average size of 9.1 nm. In this mode, it is most likely that high-Z elements such as Fe will appear as bright features. This interpretation is supported by smaller scale STEM images (Figure 3c) and their corresponding elemental maps for Fe and Si (shown, respectively, in Figure 3d,e). These maps were obtained from the Fe M edge and Si L edge EELS spectra intensity, acquired at each point of the image. In these images, Fe containing particles appear on a Si rich background, attributed to the TEM thin membrane material. The grainy nature of the film might be due to a small initial Fe thickness on the TEM grid before exposure to XeF2 or to poor wetting properties of the evaporated iron on the silicon nitride grids. A reorganization 10500
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The UPS and IPS spectra measured on the same surfaces are shown Figure 5a,b, respectively, and can help clarify the
of the 5 nm initial Fe layer upon XeF2 exposure cannot be excluded. Small area XPS measurements performed on the same FeF2 film also indicate porosity. However, Fe L2,3 electron energy-loss near-edge structures obtained on these films indicate an FeF2 environment.26 The SAED pattern of Figure 3b is also characteristic of a rutile P42/mnm phase and confirms the total transformation of metallic Fe into FeF2. Although the FeF2 films produced on TEM grids appear discontinuous as opposed to the FeF2 films grown on Fe foil, it is clear that both films possess the same stoichiometry and thus can be used as model conversion reaction material. Lithiation of FeF2 Thin Films. In the following section, we discuss the evolution of the FeF2 films upon exposure to atomic lithium: the spectroscopic measurements (XPS, UPS, and IPS) were performed in the same UHV chamber during sequential lithiation, whereas the lithiated FeF2 films for TEM were briefly exposed to atmosphere during transfer to the microscope. Figure 4a,b shows, respectively, the Fe 2p and F 1s core levels, measured on a FeF2 thin film on Fe, as well as their
Figure 5. Evolution of the (a) valence and (b) conduction band spectra of an FeF2 thin film sequentially exposed to atomic lithium. The vertical arrows indicate increasing lithium coverage θLi.
chemical modifications at the surface upon lithiation. Starting with a well-defined energy gap of 1.9 eV, upon lithiation features appear near the Fermi level in both the occupied and the unoccupied states, attributed to metallic Fe. Simultaneously, the valence band develops new electronic states from −12 to −8 eV, that are characteristic of the F 2p levels of LiF.19 In the unoccupied states, while the density of states increases at the Fermi level, the electronic states near 5 eV attributed mostly to hybridized F 2p levels become broader but are still visible. It appears that, upon lithium deposition, Fe2+ is directly reduced to Fe0 with concomitant formation of LiF. We thus obtain a final product that is similar to what has been observed previously during electrochemical discharge of an FeF2-based conversion battery.4−7 In particular, the fact that we cannot see any Fe+ intermediate state confirms that lithium must be highly mobile at the surface of the sample, with respect to iron, and can readily reduce Fe2+ to Fe0.27 The high mobility of lithium with respect to iron should also have an important impact on the morphology of the lithiated phases. TEM measurements (shown in Figure 6) have been performed on the FeF2 films exposed to a large dose of atomic lithium in order to fully reduce the initial material. After lithiation, the film appears in ADF-STEM as smaller interconnected particles whose average size has been reduced from 9.1 nm to less than 3.3 nm as shown in Figure 6a. As mentioned above, most of the signal measured in ADF-STEM can be attributed to Fe in this image. Smaller scale elemental maps in STEM mode indicate that these particles (as seen Figure 6c) are Fe and Li rich. However, there are inhomogeneities in the Li/Fe ratio (Figure 6d). In Figure 6e, an EELS spectrum, averaged over a Li-rich phase (indicated by the arrow) can be decomposed into two contributions: an Fe M-line characteristic of metallic and oxidized Fe and a Li K-line similar to what is observed from a reference LiF sample.28 This phase attribution is also confirmed by the SAED pattern of Figure 6b, which can be indexed as the superposition of bcc Fe (space group, Im3m ̅ , a = 0.2859 nm) with rock salt LiF (space group, Fm3̅m, a = 0.402 nm). There is no sign of the presence of Li2O, LiSix, or LiNx phases as would be expected from reduction of the support membrane. As a small amount of O is
Figure 4. Evolution of the Fe 2p and F 1s core level spectra of an FeF2 thin film sequentially exposed to atomic lithium. The vertical arrows indicate increasing lithium coverage θLi.
evolution during sequential metallization with Li. Immediately after Li deposition, a new feature appears 4.5 eV above the main Fe2+ peak centroid, attributed to metallic Fe0. This feature increases with increasing Li deposition. During lithiation, no signal coming from other Fe oxidation state (such as Fe+) is apparent: each spectrum can be produced by a linear combination of the spectra of FeF2 and metallic Fe0. The lithiation was stopped intentionally when the Fe appeared fully reduced. In contrast to the significant changes to the Fe 2p spectrum, the F 1s spectrum changes little upon lithiation, exhibiting a slight shift to higher binding energies. Furthermore, the relative amount of F with respect to Fe is not altered during lithiation. Due to the low XPS cross section of the Li core level, NRA was used to determine an absolute coverage of fluorine and lithium after full reduction of the FeF2 film. The fluorine concentration was found to be 20 × 1015 atoms cm−2 (with an estimated 26% relative error), compatible with a 5 nm starting thickness for the FeF2 film. Furthermore, a lithium concentration of 23 × 1015 atoms cm−2 (with an estimated 12% relative error) indicates that full reduction of Fe2+ into Fe0 is obtained when the amount of deposited lithium is comparable to the initial amount of fluorine. 10501
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morphology disruption: smaller interconnected Fe particles are the result of Fe precipitation with fast LiF formation. Despite the differences between an UHV surface science experiment and an electrochemical cell, the results presented here are similar to studies of conversion reaction in FeF2-based batteries.4−7 This approach opens the possibility of probing fundamental properties of conversion materials. In particular, it should be possible to directly examine the more complex reaction path for Li reactivity with FeF3, in order to verify if an Fe2+ intermediate state appears for low lithium coverage, followed at higher lithium concentration by a fully reduced Fe0.27 As the reactivity of crystalline nanoparticles might depend on crystallographic orientations of the nanocrystals, another important aspect that could be treated using our surface science approach is the Li reactivity of different surface orientation for the same material.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected].
Figure 6. (a) ADF-STEM image of a lithiated FeF2 thin film on a TEM grid. The dark area is attributed to the Si−O−N TEM membrane. (b) SAED pattern indicating complex structure due mainly to the overlap of a cubic (Im3m) Fe0 and a cubic (Fm3m ̅ ) LiF pattern. (c) STEM image and the corresponding elemental maps (d) of the Li−K over Fe-M intensity ratio from the lithiated FeF2 thin film 32 × 32. The EELS spectrum extracted from a Li-rich region can be decomposed into a Fe-M edge and a Li-K edge.
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was supported by the Northeastern Center for Chemical Energy Storage, an Energy Frontier Research Center funded by the U.S. DOE, BES under Award DE-SC0001294. R.T. was supported by National Science Foundation Grant 0903661: Nanotechnology for Clean Energy IGERT. The authors would also like to thank Prof. Eric Garfunkel for access to the XeF2 exposure chamber, Dr. Leszek Wielunski for NRA measurements, and Prof. Glenn G. Amatucci for fruitful discussions.
measured by EELS, the presence of the cubic rock salt FeO (space group Fm3̅m, a = 0.469 nm) cannot be ruled out due to its overlapping ring pattern with LiF. This could be due to the brief exposure to air during sample transfer to the microscope. A weak reflection observed in the SAED diffraction pattern with d = 0.401 nm is at present unidentified. These results confirm the similarity between the reaction of atomic Li in UHV on a clean FeF2 film, and its electrochemical counterpart where Li+ ions are inserted into FeF2 in a conversion battery. The reduced size of the particles is attributed to the difference in mobility between Li and Fe: with Li being more mobile, Fe2+ is immediately reduced to Fe0 which precipitates locally without diffusing. The result is an interconnected array of Fe particles mixed with an LiF phase. This is an important aspect of the FeF2-based conversion batteries as this interconnected array is though to facilitate ionic and electronic transport during the reconversion process, where Li+ ions are extracted from the LiF + Fe0 phase.
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REFERENCES
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CONCLUSION FeF2 films were prepared by exposing either a polycrystalline Fe foil or an Fe film deposited on a Si−O−N TEM membrane to XeF2 in a UHV environment. The films prepared on an Fe foil were flat with a thickness estimated to 5 nm, and their electronic structure was probed using XPS, UPS, and IPS. The FeF2 films prepared on TEM grids were composed of an interconnected array of particles with an average lateral size of 10−15 nm and height on the order of 5 nm. Upon deposition of atomic lithium on these FeF2 films, a spontaneous conversion reaction was observed: FeF2 was transformed into a mixed phase composed of metallic Fe0 and LiF. No intermediate Fe+ oxidation state was observed during the conversion reaction, indicating a high mobility for Li at the surface, allowing access to at least two Li for one Fe2+ site. Another consequence of the high mobility of lithium is the 10502
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