Glenn G. Amatucci’s research while affiliated with The State Of New Jersey and other places
What is this page?
This page lists works of an author who doesn't have a ResearchGate profile or hasn't added the works to their profile yet. It is automatically generated from public (personal) data to further our legitimate goal of comprehensive and accurate scientific recordkeeping. If you are this author and want this page removed, please let us know.
Li-metal anodes can provide high-energy-density battery configurations, but their practical use is hindered by safety concerns and poor efficiencies due to non-ideal lithium plating. In utilizing ultra-low areal plating capacities (0.08 mAh/cm2, LCP) within Li-metal half-cells, it was found that the initial formation efficiency of the solid electrolyte interphase (SEI) can be amplified and correlated with initial losses and capacity fade over time under higher areal plating capacities (2.5 mAh/cm2, 4.0 mAh/cm2, and 6.5 mAh/cm2) within an in-situ formed anodeless LCO configuration. Herein, these techniques have been utilized to introduce and optimize novel fluoroganosiyl (FOS) based dual salt electrolytes for use in in-situ formed Li-metal batteries, achieving initial cycling loss of <3% (at 4.0 mAh/cm2). Further characterization of the functional benefit of this electrolyte was elucidated using X-ray photoelectron spectroscopy surface analysis, revealing unique Li-C-N, Li3N, Si, and B-N chemistries that likely contribute to the formation of a robust SEI.
Most research on the electrochemical dynamics in materials for high‐energy Li‐ion batteries has focused on the global behavior of the electrode. This approach is susceptible to misleading analyses resulting from idiosyncratic kinetic conditions, such as surface impurities inducing an apparent two‐phase transformation within LiNi0.8Co0.15Al0.05O2. Here, nano‐focused X‐ray probes are used to measure delithiation operando at the scale of secondary particle agglomerates in layered cathode materials during charge. After an initial latent phase, individual secondary particles undergo rapid, stochastic, and largely uniform delithiation, which is in contrast with the gradual increase in cell potential. This behavior reproduces across several layered oxides. Operando X‐ray microdiffraction (μ\umu‐XRD) leverages the relationship between Li content and lattice parameter to further reveal that rate acceleration occurs between Li‐site fraction (xLi) ≈0.9 and ≈0.5 for LiNi0.8Co0.15Al0.05O2. Physics‐based modeling shows that, to reproduce the experimental results, the exchange current density (i0) must depend on xLi, and that i0 should increase rapidly over three orders of magnitude at the transition point. The specifics and implications of this jump in i0 are crucial to understanding the charge‐storage reaction of Li‐ion battery cathodes.
Most research on the electrochemical dynamics in materials for high-energy Li-ion batteries has focused on the global behavior of the electrode. This approach is susceptible to misleading analyses resulting from idiosyncratic kinetic conditions, such as surface impurities inducing an apparent two-phase transformation within LiNi 0.8Co0.15Al0.05O2 . Here, we use nano-focused X-ray probes to measure delithiation operando at the scale of secondary particle agglomerates in layered cathode materials during charge. After an initial latent phase, individual secondary particles undergo rapid, stochastic, and largely uniform delithiation, which is in contrast with the gradual increase in cell potential. This behavior reproduces across several layered oxides. Operando X-ray microdiffraction (µ-XRD) leverages the relationship between Li content and lattice parameter to further reveal that rate acceleration occurs between Li-site fraction (xLi) ~0.9 and ~0.4 for LiNi0.8Co0.15Al0.05O2 . Physics-based modeling shows that, to reproduce the experimental results, the exchange current density (i0) must depend on xLi , and that i0 should increase rapidly over three orders of magnitude at the transition point. The specifics and implications of this jump in i0 are crucial to understanding the charge-storage reaction of Li-ion battery cathodes.
Benefited from their tunable porosity, bondability, high voltage stability, fast transport properties, and good electrolyte retention, plastic separators have been an attractive technology option for separators in Li ion batteries. Despite their invention in 1994 at Bellcore, there has been little optimization and further understanding of this plastic separator comprised of PVDF-HFP copolymer resin, a plasticizer (dibutyl phthalate), and nano filler (SiO2). This paper explores the impact of formulation of the Bellcore plastic separator on transport and pore structure to identify the optimal pore architecture for fast ion transport. Each formulation was investigated for pore size distribution by gas adsorption analysis (BET) and He ion microscopy, ionic conductivity and tortuosity by electrochemical impedance spectroscopy (EIS), and mechanical properties by tensile testing to correlate pore size distribution to transport, mechanical properties, and formulation. Bondable, high voltage stable separator formulations were created with ~ 280% greater conductivity than the baseline formulation and an industry standard polyolefin separator.
Despite the fact that the first commercial rechargeable Li battery was based on a Li-Al alloy in the 1970s, extensive investigation of Li alloys for next generation high performance Li-ion batteries has taken place only since Fuji’s initiation of the renaissance in the 1990s [1-2]. Although the majority of this research has been directed towards Si, Sn, Ge and Al [3-4], Bi possesses interesting properties that make it unique amongst the lithium alloys. The heavy pnictogen is of relatively low toxicity with a broad range of applications from Pb replacement in solders to formulations in antacids. In spite of its high molecular weight Li x Bi exhibits very high volumetric capacity and exceptional Li ⁺ diffusion in its second lithiated phase, Li 3 Bi [5]. The 0.8 V alloying potential of Bi vs Li/Li ⁺ allows for a more advantageous selection of electrolyte salts and solvents than the aforementioned Si, Ge, and Al. Coupled with its superionic Li ⁺ diffusivity, the higher alloying potential of Bi also provides a larger safety window against Li plating from overpotentials and enables fast charge. While the considerable volume expansion of Bi impedes sustained capacity retention—as with all other Li alloy metals—the volume change from complete lithiation is significantly smaller than that of the popular Si: 108% for Li 3 Bi compared to 208% for Li 13 Si 4 at a comparable 3.25 Li ⁺ insertion. Combined, these positive attributes present the opportunity to meet the high-rate and volumetric energy capability needs in the next generation of Li-ion batteries.
In this work, we explore the effectiveness of BiF 3 as an electrochemical precursor to generating a Bi nanocomposite with excellent cycling efficiency extending beyond 250 cycles. An understanding of the nanocomposite’s effectiveness is derived from observations of Bi crystallite sizes produced from the conversion of BiF 3 , Bi 2 O 3 , and Bi 2 S 3 . Through electrochemical and physical characterization of these conversion materials, the size of the post-conversion Bi product was found to be not a result of processing conditions but rather the choice of conversion material. The resulting Bi crystallite size post-conversion shows a correlation between the in-situ derived crystallite size and the ionic conductivity of the resulting Li salt matrix. The resulting Bi crystallite size also provides evidence for the inverse relationship between the in-situ derived crystallite size and cyclability that is consistent with theory. For further development, alternative matrices to C were explored in order to preserve the high volumetric capacity of Li x Bi. Metal sulfides proved a more effective and volumetrically efficient C substitute, with the former exhibiting gravimetric capacities comparable to the latter. Electrolyte formulations derived from use in Bi/C nanocomposites were also tested and proved effective in stabilization of the Li–Bi alloying mechanism in our BiF 3 compositions for further cycling improvements. Finally, we demonstrate the ability of BiF 3 -derived Li x Bi alloy thin films to delithiate with an 80% utilization at >100C rate despite the presence of a LiF nanomatrix.
Through the development of BiF 3 nanocomposites using carbon or sulfide matrices in conjunction with an alloy-specific optimized electrolyte, we are able to demonstrate the exceptional cycling of the Li x Bi alloy (Fig. 1 a,b black) relative to using pure Bi metal in a composite (Fig. 1a,b blue). This presentation will outline the optimization of this alloy family from a holistic perspective and discuss future pathways based on the foundations established here, inclusive of the minimization of the LiF matrix, and in context with the other relatively few publications on this interesting alloy [6-9].
References
[1] Y. Idota, M. Nishima, Y. Miyaki, T. Kubota, and T. Miyasaki, Eur. Patent, 0651450A1 (1995).
[2] P. Pereira, G. G. Amatucci, M. S. Whittingham, R. Hamlen, J. Power Sources 280 (2015) 18–22.
[3] U. Kasavajjula, C. Wang, A. J. Appleby, J. Power Sources 163 (2007) 1003–1039.
[4] M. N. Obrovac and V. L. Chevrier, Chem. Rev. 114 (2014) 11444–11502.
[5] W. Weppner and R. A. Huggins, J. Solid State Chem. 22 (1977) 297–308.
[6] J. -S. Bridel, S. Grugeon, S. Laruelle, J. Hassoun, P. Reale, B. Scrosati, J. -M. Tarascon, J. Power Sources 195 (2010) 2036–2043.
[7] W. Xianming, T. Nishina, and I. Uchida, J. Power Sources 104 (2002) 90–96.
[8] A. Finke, P. Poizot, C. Guéry, L. Dupont, P.-L. Taberna, P. Simon, and J.-M. Tarascon, Electrochem. Solid State Let. 11 (2008) E5–E9.
[9] C.-M. Park, S. Yoon, S.-I. Lee, H.-J. Sohn, J. Power Sources 186 (2009) 206–210.
Figure 1
The ubiquitous use of Li-ion batteries is hindered in part by limitations to achievable specific and volumetric energy. In addition to the physical constraints these properties address, they lead to higher cost per Wh. Replacing the intercalation or alloying based negative electrode leads to significant increases in both of the aforementioned energy related properties. In-situ formed or “anodeless” Li metal batteries, where Li is plated directly on the current collector, enable significant cost savings and improvement of energy of Li metal batteries relative to traditional Li metal which contain Li metal upon initial fabrication [1].
Using anodeless configurations, we imparted on an effort to investigate the deleterious capacity consuming phenomena of the plating process by electrochemically isolating the key degradation and performance limiting phenomena. Using ultra low-capacity Li metal plating (0.08mAh/cm ² ) in Li-metal anodeless half cells, the formation and quality of the solid electrolyte interphase (SEI) can be investigated with amplification. The solid electrolyte interphase (SEI) is a protective and passivating layer formed during the initial reduction of electrolyte. A robust SEI protects against excess electrolyte consumption and allows for subsequent stable, high efficiency cycling while being electronically resistive and ionically conductive [2]. The composition and mechanical stability of this dynamic layer influences efficiency and cycle lifetime [2]. At higher lithium plating capacities however, the influence of the SEI is not easily observed. Instead, capacity fade attributed to dendrite formation and mechanical damage to the SEI can be more easily investigated. Using higher capacity lithium plating (2.5mAh/cm ² ) in anodeless cell configurations, dendrite formation and capacity fade over cycle lifetime can be observed. Dendrite formation, as the result of irregular lithium deposition, can hinder the amount of active lithium available as well as lead to battery failure through short circuiting and thermal runaway [3].
In this work we have isolated the study SEI and dendrite formation electrochemically, using Li-metal and LiCoO 2 counter electrode anodeless cell configurations, respectively. As a result, we developed novel non-aqueous, non-ionic liquid electrolytes that have achieved efficiency of 85% at 0.08mAh/cm ² where most of the capacity is associated with the formation of the SEI and 96% at 2.5mAh/cm ² at the more challenging initial cycles. In contrast, a standard electrolyte composition such as 1M LiPF6 EC/DMC shows poor initial efficiency of 52% at 0.08mAh/cm ² and 66% at 2.5mAh/cm ² . High efficiencies were also achieved with optimized ionic liquid electrolytes. This work may provide insight into the initial stages of SEI formation and provide a systematic methodology for electrolyte optimization through SEI capacity loss and dendritic capacity fade observation.
[1] S. H. Park, D. Jun, G. H. Lee, S. G. Lee, and Y. J. Lee, J. Materials Chemistry A 9 (2021) 14656-14681.
[2] S. J. Park, J. Li, C. Daniel, D. Mohanty, S. Nagpure, and D. L. Wood, Carbon 105 (2016) 52-76.
[3] J. B. Goodenough, J. Solid State Electrochem. 16 (2012) 2019-2029.
Figure 1
Li x Bi alloys are a unique path to enable extraordinarily high rate, safer, and high volumetric energy density batteries. The effectiveness of BiF 3 as an electrochemical precursor to a Li x Bi nanocomposite with excellent cycling efficiency extending beyond 250 cycles is explored to identify key factors critical to future successful development. The relationships between the converted crystallite size and subsequent electrochemical properties were reported with a specific focus on cycling efficiency. Through electrochemical and physical characterization of post-converted BiF 3 , Bi 2 O 3 , and Bi 2 S 3 , the size of the post-conversion Bi product was directly correlated with the ionic conductivity of the in-situ formed Li salt matrix and subsequent cycling stability. Further key areas for development were introduced, including volumetrically dense conductive alternatives to C and electrolyte formulations, which demonstrate significant improvements in cycling stability. In addition, we demonstrate the ability of BiF 3 -derived Li x Bi alloy thin films to delithiate with an 80% utilization at >100C rate despite the presence of a LiF nanomatrix.
Metal fluorides, promising lithium-ion battery cathode materials, have been classified as conversion materials due to the reconstructive phase transitions widely presumed to occur upon lithiation. We challenge this view by studying FeF3 using X-ray total scattering and electron diffraction techniques that measure structure over multiple length scales coupled with density functional theory calculations, and by revisiting prior experimental studies of FeF2 and CuF2. Metal fluoride lithiation is instead dominated by diffusion-controlled displacement mechanisms, and a clear topological relationship between the metal fluoride F⁻ sublattices and that of LiF is established. Initial lithiation of FeF3 forms FeF2 on the particle’s surface, along with a cation-ordered and stacking-disordered phase, A-LixFeyF3, which is structurally related to α-/β-LiMn²⁺Fe³⁺F6 and which topotactically transforms to B- and then C-LixFeyF3, before forming LiF and Fe. Lithiation of FeF2 and CuF2 results in a buffer phase between FeF2/CuF2 and LiF. The resulting principles will aid future developments of a wider range of isomorphic metal fluorides.
... In addition to binder gradient [45] conductive additive content gradient [69] and particle size gradient [70,71] following the principles discussed above which all show great improvement when implemented, porosity gradient is beneficial too. [72,73] Porosity is larger near separator and smaller near current collector, helping ion transform more homogeneous and reducing overpotential to a large content, as shown in Figure 12. While single gradient brings merits to SOC homogeneity without doubt. ...
... Electrochemical-based and data-driven methods are two common approaches for estimating battery SOH [20][21][22][23]. Electrochemical-based methods estimate the SOH directly through capacity testing, impedance measurement, and other techniques. ...
... 2. These results indicate that thinner films result in higher measured capacities. A solid electrolyte interface (SEI) formation during first discharge cycle in FeF 2 films has been reported [20] , and a capacity in excess of the theoretical value of FeF 2 at C/10 was reported and attributed to the SEI layer formation [5]. The large capacity in the first discharge cycle observed in this work compared to subsequent cycles irrespective of the FeF 2 thickness is also likely due to the formation of an SEI layer. ...
... However, the electrochemical performance of Bi-based electrodes is hindered by its high volumetric expansion during alloying reactions and the unstable SEI films induced by side reactions at the interface of the electrolyte and electrode. To address these challenges, various modification methods have been introduced into Bi anodes, including carbon matrix modification, nanosizing and electrolyte engineering [66][67][68]. ...
... Since the structural degradation of cathode materials is directly related to the transition metal migration, 26 the overincrease in interlayer spacing would decrease the migration barriers of transition metal, thus accelerating the structural deterioration from layered to spinel and eventually rock-salt structures. 27,28 To correlate the interlayer spacing and the electrochemical properties, the initial charging/discharging measurements were performed at 0.1 C rate between 2.8 and 4.3 V. As shown in Figure 3a, the initial discharge capacities are near values of 183.9 mA h g −1 for NCM811, 184.9 mA h g −1 for Mg-doped NCM811, 184.7 mA h g −1 for Ca-doped NCM811, and 188.1 mA h g −1 for Mg/Ca-doped NCM811. ...
... It is also seemingly higher than the intercalation plateau of FeF 3 (3.0-3.3 V) 19 . However, this may be due to the pronounced overpotential in the reactions of FeF 3 . ...
... Research in the LiMn 2 O 4 and LiFePO 4 space has showed significant improvement in electrochemical performance for these less costly positive electrode materials. [1][2][3][4][5] Additionally, Na-ion cells have gained much popularity in the last years as researchers turn to more abundant, inexpensive elements in anticipation for the continuously growing market demand for batteries. [6][7][8] These alternative positive electrode active materials-albeit interesting for applications where gravimetric or volumetric energy are less of a concerncan only provide a fraction of the energy density provided by Ni-rich layered oxides. ...
... At such SOCs and high voltages, Ni-rich cathodes lose oxygen from the surface [32][33][34]. Consequently, the cathode surface "reconstructs" to compensate for the O deficiency by transforming from the layered to cubic (rock-salt and/or spinel) phases, creating a surface layer with reduced transition-metal (TM) species [20,25,30,35,36]. These species can also lead to TM dissolution from the cathode to the anode, which can adversely affect the SEI layer and contribute to capacity fade [26,27,[36][37][38]. ...
... Therefore, using RIXS to ascertain the role of O-O dimerization in the bulk charge compensation mechanism and its ability to provide ''excess'' capacity necessitates additional scrutiny, especially when considering the difficulty in separating it from possible surface contributions. 11,56 Long-term high-voltage cycling of NMC811-graphite pouch cells Typically, high-voltage O-redox activity is thought to be intimately linked with bulk structural transitions. In Li-rich NMC cathodes, these changes are irreversible, leading to poor longterm electrochemical performance. ...
... Unsurprisingly, the19 F NMR spectra (Figure S15) of the pristine ALD NMC811 and EC/EMC-soaked material show no resonances; i.e. no fluorine-containing surface species are observed. In the electrolyte soaked ALD NMC811 (point I in Figure 3 (a)), the resonances are assigned to residual LiPF6 salt (-76 --80 ppm), 8,10,81 difluorophosphate species from hydrolysis of the salt (-88 ppm)10 , an aluminium oxyfluoride species (-184 ppm),83 and LiF (-206 ppm).83 Fluorination of the Al2O3 coating could explain the decrease in27 Al chemical shifts of the electrolyte-soaked ALD NMC811(Figure 2 ...
![Operando μ$\umu$‐XRD of LiNi0.8Co0.15Al0.05O2 secondary particle agglomerates during first charge and discharge. Diffraction peaks corresponding to a) (003) and b) (104) reflections. c) Galvanostatic charge/discharge profile. d) a and c lattice parameters for a single secondary particle, P1, refined by Le Bail method. e) Extent of (de)lithiation corresponding to refined cell parameters[⁸] for two secondary particle agglomerates.](https://www.researchgate.net/publication/373799676/figure/fig1/AS:11431281244394041@1715899926675/Operando-mumu-XRD-of-LiNi08Co015Al005O2-secondary-particle-agglomerates-during_Q320.jpg)
![Figure 1: Operando µ-XRD of LiNi 0.8 Co 0.15 Al 0.05 O 2 secondary particle agglomerates during first charge and discharge. Diffraction peaks corresponding to (a) (003) and (b) (104) reflections. (c) Galvanostatic charge/discharge profile. (d) a and c lattice parameters for single secondary particle, P1, refined by Le Bail method. (e) Extent of (de)lithiation corresponding to refined cell parameters[8] for two secondary particle agglomerates.](https://www.researchgate.net/publication/369080568/figure/fig1/AS:11431281125389622@1678292198410/Operando-XRD-of-LiNi-08-Co-015-Al-005-O-2-secondary-particle-agglomerates-during_Q320.jpg)
![X-ray diffraction and PDF analyses of m-FeF3
a, High-resolution X-ray diffraction data for the cycled m-FeF3 collected at selected states of charge. The (003) and (011¯)\documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$(01\bar 1)$$\end{document} reflections of the A-LixFeyF3 phase are marked by red and blue triangles, respectively. Asterisks mark the hydrated FeF3 impurity introduced during sample preparation. Top inset: the intensity variations of the (012) reflection with state of charge. Right blue inset: the two components in the broad feature at around 26°–28° values of 2θ. dis, discharge; cha, charge. b, Bragg contribution of FeF3 (green), A-LixFeyF3 (red) and B-LixFeyF3 (blue) derived from Rietveld refinement shown in selected 2θ range. Key reflections of each phase are marked by dashed lines with indices labelled. The inset shows the evolution of the phase mole fraction within the first insertion cycle. c, PDF results of m-FeF3 collected within the insertion cycle. Low-r atom pairs in FeF3 are labelled. Top inset: the phase evolution of each component derived from refinement. r, is the real-space distance and G(r), the pair distribution function. d, The breakdown of the PDF contributions obtained from refinement of the PDF results from the red dashed rectangles are shown in d where experimental, calculated and difference PDF patterns are shown in black circles and orange and purple lines, respectively. These results are compared to the calculated PDF patterns from the related structures plotted using the same colour code as that used in b.](https://www.researchgate.net/publication/348664724/figure/fig4/AS:982646607847424@1611292616030/X-ray-diffraction-and-PDF-analyses-of-m-FeF3-a-High-resolution-X-ray-diffraction-data_Q320.jpg)
![ED analyses of the partially discharged m-FeF3 ‘Li 0.25’ sample
a–f, Zone axis ED analysis. The diffraction pattern recorded parallel to the [001] direction of the crystal (a); b and c are simulated ED patterns for the same orientation of FeF3 and X-ray-derived A-LixFeyF3, respectively. Experimental pattern recorded parallel to [100] (d); additional reflections from FeF2 in the electron beam are marked by red circles on one side of the pattern; e and f are simulations, respectively, for FeF3 and A-LixFeyF3. Reciprocal lattice parameters a*, b* and c* are denoted by asterisks. g, A projected potential map solved from the pattern in d using charge flipping. h, Structure model (P312) based on the heavy atom positions in g and derived by DFT. i–k, Scanning electron diffraction analysis (i). The red and green colours correspond to A-LixFeyF3 and FeF3, respectively. The diffraction pattern collected at position ‘j’ is shown in j where the strong (013)/(01¯3)\documentclass[12pt]{minimal} \usepackage{amsmath} \usepackage{wasysym} \usepackage{amsfonts} \usepackage{amssymb} \usepackage{amsbsy} \usepackage{mathrsfs} \usepackage{upgreek} \setlength{\oddsidemargin}{-69pt} \begin{document}$$(013)/(0\bar13)$$\end{document} reflections are assigned to the A-LixFeyF3 structure. The diffraction pattern collected at position ‘k’, in k, shows reflections consistent with the pristine FeF3 structure. l, Direct-beam intensity mapping of the sample shows a substantial amount of the less-crystalline FeF2 phase, whose lack of long-range ordering is reflected by the diffraction pattern in the inset.](https://www.researchgate.net/publication/348664724/figure/fig5/AS:982646612045824@1611292617189/ED-analyses-of-the-partially-discharged-m-FeF3-Li-025-sample-a-f-Zone-axis-ED_Q320.jpg)
