With the booming demands for electric vehicles and electronic devices, high energy density lithium-ion batteries with long cycle life are highly desired. Despite the recent progress in Si1 and Li metal2 as future anode materials, graphite still remains the active material of choice for the negative electrode.3,4 Lithium ions can be intercalated into graphite sheets at various stages like LixC12 and LixC6, providing a high specific capacity of 372 mAh/g (&#;2.5 times higher than LiCoO2) and high volumetric capacity (similar to LiCoO2) corresponding to LiC6.5,6 In addition to its low cost and non-toxicity, graphite has a lowest average voltage (150 mV vs. Li/Li+) and the flat voltage profile, rendering a high overall cell voltage.7 Graphite also displays a very low voltage hysteresis, which in turn results in high energy efficiency.8 For applications in lithium-ion batteries, many properties of the graphite powders must be optimized including crystallinity, particle size, morphology, and surface chemistry. However, most research effort has been placed on the development and characterization of high performance cathodes, while the impact of the graphite anode on full cell performance has been largely unexplored.
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Compared to widely used battery cathodes such as LiCoO2 (140 mAh/g), LiFePO4 (160 mAh/g), LiNi1/3Mn1/3Co1/3O2 (160 mAh/g), and LiNi0.5Mn0.3Co0.2O2 (175 mAh/g),9&#;11 nickel-rich LiNi0.8Mn0.1Co0.1O2 (NMC811) is delivers a higher specific capacity (180&#;220 mAh/g), which increases the battery life on a single charge.12,13 The high capacity arises because Ni is the main redox-active species in the NMC host structure.14 While Ni-rich cathodes could reach an energy density that exceeds 700 Wh/kg, undesirable side-reactions like electrolyte oxidation15 and cathode surface reconstruction16,17 have been identified as major causes of capacity loss, due to the high reactivity of Ni4+.18&#;20
In this study, we show that the choice of graphite also significantly impacts the long-term cycling stability of cells with Ni-rich NMC cathodes. Six commercially-available natural and synthetic graphite samples were tested in full cells with NMC811 cathodes. The variations in cell performance were correlated with the chemical and structural properties of the different graphite anodes.
Cells were discharged to 3.0 V at C/10 before disassembly. Cycled electrodes were harvested in an argon-filled glove box for post-mortem analysis. X-ray powder diffraction (XRD) was collected on a PANalytical X'Pert system with a Cu source (λ = 1.54 Å) operated at 45 kV and 40 mA, and automatic divergence and anti-scatter slits. XRD patterns were collected from pristine powders and cycled electrodes. N 2 adsorption of each graphite powder sample was measured on a Quantachrome gas sorption analyzer, and the BET surface area was calculated using a multiple-point method. X-ray photoelectron spectroscopy (XPS, Thermo Scientific K-Alpha) was used to analyze the surface chemistry of powder samples and electrodes. The X-ray source was monochromated Al Kα with .6 eV photon energy and a spot size of 400 μm. The system used low energy Ar-ions and a low energy electron flood gun for charge compensation. The harvested electrodes were lightly rinsed with DMC solvent. After rinsing, the electrodes were dried in the glove box for an hour to evaporate most of the DMC solvent. Then the electrodes were further dried under vacuum overnight in the glove box antechamber. After drying overnight, the electrodes were taken into the glove box again and loaded into the vacuum transfer container to avoid air and moisture exposure. The transfer container was directly inserted into the XPS chamber with a base pressure of 10 &#;9 Torr. Scanning electron microscopy (SEM, Merlin VP, Zeiss) was used for microscopy analysis. Cross-sectional SEM images of anodes were obtained by polishing the anodes buried in resin substrates. Raman spectra were acquired with an Alpha 300 confocal Raman microscope (WITec, GmbH) using a solid-state 532 nm excitation laser, a 20 × objective, and a 600 grooves/mm grating. Acquisition times for each spectrum were typically 15 s. FTIR spectra were collected using a Bruker Alpha spectrometer housed in an argon-filled glove box. ATR-FTIR spectra were collected using a germanium crystal at 4 cm &#;1 resolution.
All cells were tested on a Maccor Battery Cycler in an Espec environmental chamber at 30°C. The cells were cycled between 3.0 and 4.2 V using the standard cycling protocol adopted in Reference 22 . The protocol includes a cycling test interrupted periodically to obtain reference performance data, which includes a DC resistance test (hybrid pulse power characterization (HPPC)) after every 50 cycles. Cycled electrodes were extracted from cells in an argon-filled glove box without wash and re-assembled into half cells with fresh lithium foil and electrolyte for further electrochemical analysis. The discharge rate performance was evaluated in half cells by intercalating lithium at a constant rate of C/5 and de-intercalating lithium at variable C-rates for the graphite anode, and by de-intercalating lithium at a constant rate of C/5 and intercalating lithium at variable C-rates for the cathode. Half cells with fresh graphite anodes were also built to evaluate the reversible capacity of the different graphites. Bio-Logic potentiostats/galvanostats (VSP) and EC-Lab software version 11 were used to obtain and analyze electrochemical impedance spectra (EIS).
Electrode fabrication and cell build were completed at the U.S. Department of Energy (DOE) Battery Manufacturing R&D facility at Oak Ridge National Laboratory. 21 Electrodes were prepared by slot-die coating slurries onto foil current collectors (aluminum foil for the cathode and copper foil for the anode). The cathode slurry contained 90 wt% LiNi 0.8 Mn 0.1 Co 0.1 O 2 powder (Targray, Lot #), 5 wt% acetylene carbon black (Denka Black), and 5 wt% polyvinylidene difluoride (Solvay ) in N-methyl-2-pyrrolidone (NMP). Six graphite powder samples (ConcoPhillips A12, GrafTech APS19, Showa Denko SCMG-BH, Hitachi MAGE, Hitachi MAGE3 and Superior SLC T, summarized in Table S1) were provided by suppliers and used as received. The anode slurry contained 92 wt% of each graphite powder, 2 wt% carbon black (Timcal Super C65), and 6 wt% polyvinylidene difluoride (Kureha ) in NMP. Graphite samples were coated onto Cu foil to obtain the same areal capacity of 2.6 mAh/cm 2 . The NMC811 cathode was coated with an areal capacity of 2.2 mAh/cm 2 , yielding a negative to positive capacity ratio (N/P ratio) of 1.2 in the full cells. All electrodes were calendared to 35% porosity after primary drying. The electrodes underwent secondary drying under vacuum at 120°C prior to cell assembly. 1.2 M LiPF 6 dissolved in ethylene carbonate: ethylmethyl carbonate (EC:EMC = 3:7 by weight, SoulBrain) was used as the electrolyte. -type (Hohsen) coin cells were assembled in an argon-atmosphere glove box, and single layer pouch cells (&#;100mAh) were assembled in the dry room (with a dew point of less than &#;50°C and relative humidity (RH) of 0.1%) and cycled with pre-loaded pressure of 5 psi.
Half cells with lithium metal counter electrodes were first used to evaluate the specific capacity of each graphite. Most graphite samples showed a high reversible capacity close to the theoretical value of 372 mAh/g at C/3 intercalation rate, while SCMG-BH graphite exhibited a relatively lower capacity of 322 mAh/g (Figure S1). Figure 1 shows capacity-voltage profiles from the full cells with each graphite anode and NMC811 cathode during the first four cycles at C/10 rate (also known as formation cycles).23 During the first lithiation of the graphite anode, electrolyte reduction occurs on the graphite surface forming a solid-electrolyte-interphase (SEI), which is ionically conductive but electronically insulating. The SEI acts as a protective layer to impede continuous electrolyte decomposition and solvent co-intercalation into graphitic layers during subsequent cycles.24,25 The large charge capacity and low coulombic efficiency in the first formation cycle seen here is directly associated mainly with anode SEI formation as well as irreversible capacity loss on the cathode (Figure S2). For example, the NMC811/SLC T cell delivered a charge and discharge capacity of 224 mAh/gNMC and 192 mAh/gNMC, respectively, yielding a coulombic efficiency (CE) of 86.1% for the first formation cycle. During the second formation cycle, a charge capacity of 194.4 mAh/gNMC was observed. The discharge capacity increased to 194.0 mAh/gNMC and coulombic efficiency increased to 99.8%, indicating that a stable SEI layer was formed in the first cycle. The first formation cycle charge/discharge capacity for full cells with different type of graphite anodes are summarized in Table S2, and the corresponding coulombic efficiencies are displayed in Figure S3a. The cells with the MAGE graphite exhibited the highest first cycle coulombic efficiency of 86.1%, followed by SLC T. On the other hand, MAGE 3 had the lowest first cycle CE of 82.2%, indicating more active lithium loss or parasitic reactions for this graphite.26 To further evaluate the active lithium loss during formation and the effectiveness of the passivation layer, the compounded CE for all four formation cycles was calculated and displayed in Figure S3b. MAGE 3 exhibited the lowest cumulative coulombic efficiency, followed by SCMG-BH. This indicates that more active lithium was lost due to SEI formation for MAGE3 and SCMG-BH graphites. MAGE and SLC T graphites, by contrast, demonstrated the highest cumulative coulombic efficiency, consistent with the formation of an effective passivation layer.
Figure 1. Voltage profiles of NMC811 cells with different graphite anodes during formation cycles a) A12 graphite anode. b) APS 19 graphite anode. c) SCMG-BH graphite anode. d) MAGE graphite anode. e) MAGE 3 graphite anode. f) SLC T graphite anode.
Figure 2a shows the capacity retention of each full cell with different types of graphite anodes during the aging cycles at C/3 rate. Here we used an aggressive cycling protocol that includes a three-hour long voltage hold at the top of each charge to accelerate the effects of high voltage on cell degradation.22 As expected, a rapid decrease of capacity was observed in all cells from this accelerated aging test. At the end of long-term cycling (300 aging cycles), only the cells made with MAGE and SLC T graphite retained more than 80% of initial capacity. Moreover, different fading trends were observed for cells made with different graphite anodes. While the capacity fading rate didn't change for cells with MAGE and SLC T, those with A12 and APS 19 showed much greater capacity decay after 200 cycles. On the other hand, cells with MAGE3 and SCMG-BH graphite anodes showed higher capacity loss initially (Figure S4). In addition, cells with MAGE3 and SCMG-BH anodes also exhibited a much lower initial discharge capacity compared to cells made with other graphites. This is a direct consequence of their relatively lower coulombic efficiency during formation cycles compared to other graphite samples, indicating more active lithium loss from parasitic reaction. As cycling progressed, the rate of capacity fade became less steep, in accordance with the CE increase and gradual stabilization of the SEI. In the end, the cells with MAGE 3 and SCMG-BH anodes still exhibited the lowest discharge capacity and the worst capacity retention after 300 cycles, mostly due to the severe loss of active lithium during formation cycling. Initial and final discharge capacities of cells with different graphite anodes and their corresponding coulombic efficiencies during aging cycling are summarized in Table S3.
Figure 2. a) Cycling performance of NMC811 full cells with different graphite anodes during C/3-C/3 aging cycles from 3.0 to 4.2 V at 30°C. b) Differential capacity vs. cell voltage of NMC811- SLC T graphite cell. c) Rate capability of NMC811 half cells with a fresh cathode and cathodes extracted from full cells with different graphite anodes after aging cycles. d) Rate capability of graphite half cells with anodes extracted from full cells after aging cycles. The capacity is normalized to the capacity of the corresponding fresh anode at C/10 discharge rate.
To further investigate the origin of capacity loss, differential capacity vs. voltage profiles (dQ/dV) were analyzed. A representative dQ/dV plot for the NMC811/SLC T cell is displayed in Figure 2b, and the corresponding capacity-voltage profile is depicted in Figure S5. During the first formation cycle, there is a small anodic bump from 2.97 to 3.05 V, which arises from EC reduction on the graphite anode to form the SEI layer.27 During the second formation cycle, the first prominent peak at 3.44 V is attributed to lithium intercalation into the graphite anode, and the other three peaks stem from phase transitions in NMC811.28 The first phase transition from a hexagonal to a monoclinic (H1 &#; M) lattice was observed at 3.63 V, followed by a small anodic feature around 3.91 V, which stems from the M &#;H2 phase transition.29 The last anodic peak observed at 4.13 V belongs to the H2-H3 phase transition, which leads to a significant unit cell volume reduction and contributes to structural instability.30 Recent studies also suggest that oxygen could be redox-active during this phase transition.17 As the cycling proceeded, all redox peaks from NMC811 shifted further apart, which is consistent with increasing impedance and polarization. Note that the peaks corresponding to H1 &#; M and M &#;H2 transitions also became smaller as the cycle number increased, consistent with the overall capacity fade during cycling.
To evaluate the degree of cathode degradation after long-term cycling in the full cell, half cells were built with the cycled cathodes and fresh lithium chips as counter electrodes. All cycled cathodes exhibited an inferior rate performance to the pristine cathodes as shown in Figure 2c, likely due to the continuous impedance growth during long-term cycling (Figure S6a). Electrochemical impedance spectra (EIS) were collected on each half cell to evaluate the cathode impedance growth. Cathodes just after formation cycles were also harvested and assembled into half cells to measure their impedance for comparison (Figure 3). In addition to the high frequency semicircle observed after formation cycles, a second semicircle appeared at lower frequency for all cathodes after aging cycles, which could be attributed to the development of charge-transfer impedance of slower process.31,32 As summarized in Figure S6b, all cycled cathodes had significantly higher impedance after aging cycles. The strong cathode impedance growth may come from side reactions that occur at the cathode surface during cycling33,34 and the formation of surface reconstruction layers.16 Significant efforts have been devoted to improve the electrochemical performance of nickel-rich NMC cathodes through various approaches including doping,35 surface modification36,37 and careful structural design such as core-shell geometries38 and concentration gradients of the transition metals.39,40 The capacity retention of each cycled cathode in the half cell varied from 81.0% to 86.2% compared to the fresh cathode measured at C/3 rate (Figure S7a). The capacity retention in the half cells was higher than the full cells (Figure S7b), consistent with loss of active lithium in the full cells from SEI formation and growth at the graphite anode.
Figure 3. Impedance spectra of half cells built with cycled NMC811 cathodes from cells with a) A12 graphite anode. b) APS19 graphite anode. c) SCMG-BH graphite anode. d) MAGE graphite anode. e) MAGE3 graphite anode. f)SLC T graph anode.
Cells with cycled anodes were also assembled into half cells to evaluate the degradation from the anode side. The discharge capacity of each cycled anode at C/5 was compared with fresh anodes (Figure S8). Only a small capacity decrease (within 3%) was observed in each cycled graphite anode at C/5. However, the capacity fade became obvious as the C-rate increased, and the difference between each cycled graphite anode became more significant as well (Figure 2d). Cells made with cycled SCMG-BH, A12, and MAGE3, which had relatively lower capacity retention in the full cells, showed the most dramatic capacity loss at high C-rate. Studies have shown that Li+ desolvation and its subsequent slow solid-state diffusion through the SEI layer could be the rate limiting step at higher discharge rates. Thus, more significant capacity loss in some graphite samples at high C-rate could indicate greater polarization due to a thicker and more resistive SEI layer.41 EIS measurements were also conducted on anodes both after formation and aging cycles to evaluate the impedance growth (Figure 4). Note that the impedance obtained from each graphite after aging cycles was much smaller than that of the corresponding cathode (Figure S6b), further proving that the cathode is the main contributor to the full cell impedance after the aging cycles. Though the contribution from the anode side is smaller, the impedance growth on different graphite samples varied dramatically. For instance, the impedance of SCMG-BH, A12, and MAGE3 anode more than doubled, consistent with their inferior capacity retention. On the other hand, MAGE and SLC T, which demonstrated the best capacity retention after the aging cycles, exhibited a much smaller impedance rise, consistent with less SEI growth and active lithium loss.
Figure 4. Impedance spectra of half cells built with a) Cycled A12 graphite anode. b) Cycled APS19 graphite anode. c) Cycled SCMG-BH graphite anode. d) Cycled MAGE graphite anode. e) Cycled MAGE3 graphite anode. f) Cycled SLC T graph anode.
The crystallinity of cathodes harvested from cycled cells was analyzed by PXRD. Two representative XRD patterns from cathodes cycled with SLC T and MAGE3 anodes are depicted in Figure S9. Note that these cells exhibited among the best and worst capacity retention, respectively. No new peaks or significant peak shifts were identified in either cathode, indicating that the bulk structure remained unchanged. In addition, the I(003)/I(104) ratio remained higher than 1.40 in both patterns, which indicates that no significant cation mixing occurred between Li+ and Ni2+ in the lithium layer.42 However, even at a low discharge rate of C/10, the capacity from cells built with the cycled cathodes and fresh Li counter electrode was still around 10% lower than the capacity from cells built with fresh cathodes (Figures 2c). Because of the low current density, impedance can be ignored, and the capacity loss reflects the intrinsic property of the electrode. Though no significant bulk structural change was observed from XRD, some structural changes may have occurred on the surface, which contribute to the irreversible capacity loss and impedance growth.16,43 In particular, Ni4+ is known to be highly reactive with the electrolyte, which leads to surface reconstruction and the buildup of surface reaction layers. In addition, formation and growth of cracks in the cathode particles degrade their ionic and electronic connectivity and further lead to a sluggish, incomplete reaction.44
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To shed light on the differences between graphites, the pristine graphite powders were analyzed. Distinctly different particle morphologies could be seen in SEM micrographs (Figure S12). In general, SCMG-BH, MAGE, and MAGE3 graphite particles were potato-shaped, while A12 and APS 19 displayed a flake-like morphology. SLC T exhibited the most spheroidized morphology with relatively smooth, round particles. These images also correlate qualitatively with the BET surface area of these graphite samples, which are summarized in Table S1. Spherical particles are known to minimize the surface area, and the SLC T sample showed the lowest surface area of 2.1 m2/g. In contrast, lower density flake morphologies could lead to a larger surface area, which is in good agreement with the largest surface area of 6.0 m2/g for APS19 among all six samples. The crystallite size of the graphite sheets could also be seen in cross-sectional SEM images (Figure S13), which varied significantly between each graphite sample. Among them, SCMG-BH graphite had the smallest graphite sheets. Since solvent decomposition occurs at a much higher rate on edge plane graphite than on basal plane, a morphology that minimizes the graphite edges (non-basal surface area) would reduce irreversible capacity loss during formation cycles.41 For example, spheroidized SLC T had the highest coulombic efficiency in the first formation cycle. The strong correlation between particle surface area and first cycle irreversible capacity loss due to SEI formation is also well known,45,46 and their impact on the full cell cycling performance was clearly revealed in Figures 5a and 5b. Here, MAGE and SLC T, which have the two lowest surface areas, demonstrated obviously high coulombic efficiencies during the first formation cycle and ended with a better full cell capacity retention. On the other hand, APS19 and MAGE3 have the two highest surface areas and showed obviously lower coulombic efficiencies during the first formation cycle (Figure 1 & Figure S3a).
Figure 5. Correlations between full cell capacity retention and a) Surface areas of graphites. b) First cycle coulombic efficiency. c) Cycled graphite grain size.
The crystallinity of all six fresh graphite samples and cycled anodes were also examined by XRD (Figure S10). All pristine graphite powders exhibited a distinctive diffraction peak around 26.42° assigned to the graphite (002) reflection. This diffraction peak corresponds to a spacing of 0.337 nm between two adjacent carbon layers in an ideal hexagonal graphite structure with an ABAB stacking sequence.47 The coherence length of the graphite crystallites along the c axis LC can be estimated from the broadening of the (002) reflection using the Scherrer equation as summarized in Table S4.7 As seen in Figure 6a, SCMG-BH graphite had the broadest (002) diffraction peak among the pristine samples, corresponding to the smallest crystallite size of 53 nm and consistent with our observation in the SEM images. The grain size of SLC T was approximately three times larger at 162 nm. MAGE and MAGE 3 had the largest crystallite sizes, but LC could not be quantitatively determined since their peak widths were close to the instrumental broadening (FWHM &#; 0.081°). In addition to the prominent (002) diffraction peak, an additional peak at 43.45° was clearly observed in four of the six graphite samples (Figure S11). This peak is assigned to the (101) reflection in the rhombohedral phase of graphite.48 The relative amount of the hexagonal and rhombohedral phases can be estimated by comparing the integrated intensities of (101) and (002) diffraction peaks, using a structure factor of 0.754.7 While the rhombohedral phase does not significantly impact the electrochemical potential for lithium insertion, its presence and grain boundaries with the hexagonal phase may reduce irreversible capacity loss.7 However, there seems no optimal rhombohedral phase ratio for long-term cycling. While A12, MAGE and SLC T have similar rhombohedral phase fractions (Table S5), their capacity retention in long-term cycling varied considerably. Raman spectroscopy is another powerful tool to characterize the structure of graphite and estimate the crystallinity. The Raman spectrum of graphite is dominated by the strong G band around cm&#;1, which is the only first-order Raman band allowed in a perfect single crystal. The D band near cm&#;1 is a disorder-induced band assigned to the breathing mode of aromatic rings.49 The D-band to G-band integrated area ratio (ID/IG) is often used as a parameter for evaluating carbon structural properties.50 The increased D-band relative to the G-band is typically associated with a higher content of disordered carbon or smaller graphite crystallite size. While A12, MAGE, and MAGE3 had similar well-ordered graphitic surfaces, SCMG-BH and SLC T had very disordered surfaces (Figure 6b). Nevertheless, these differences of surface structure observed by Raman spectroscopy were not reflected in long-term cycling stability
Figure 6. a) (002) diffraction peaks of pristine graphite powders (solid lines) and cycled graphite anodes (dashed lines). b) Raman spectra of pristine graphite powders. c) XPS spectra of pristine graphite powders. d) FTIR spectra of cycled graphite anodes.
Though no new peaks were identified in any diffraction patterns of the cycled anodes, obvious peak broadening was observed in all samples (Figure 6a). The peak broadening indicates grain size reduction and particle fracturing that occurred during extensive cycling, which is strongly correlated with capacity fade. As demonstrated in Figure 5c, SCMG-BH, A12, and MAGE3, which have relatively lower capacity retention, demonstrated much more significant grain size reduction and had the smallest grain size after cycling (Table S4). On the other hand, MAGE and SLC T retained the largest grain size, in line with their high capacity retention. Reduction of grain size and fracture could be the result of strain and stress during repeated lithiation and delithiation of the graphite. As a result, impedance could rise due to intraparticle resistance and loss of connectivity. Moreover, the mechanical degradation and particle fracture can also expose fresh surfaces to the electrolyte, resulting in more SEI formation and active lithium loss.51 As discussed above, MAGE and SLC T demonstrated the lowest impedance growth (Figure 4 and S6b) during the aging cycles, consistent with the small degree of graphite particle fracture. By contrast, anodes that exhibited higher impedance growth during aging like A12 and MAGE3 showed a much higher degree of grain size reduction.
XPS is able to identify elements on the surface (<5&#;10 nm) and their chemical state with excellent sensitivity, and thus was employed to investigate the surface chemistry of pristine graphite powders and the SEI layers formed after long-term cycling. From the survey scan, only C and O elements were identified in each pristine graphite powder, confirming the purity of all samples in the study. Each graphite exhibited a C 1s peak at 284.6 eV assigned to the C-C bond in graphite (Figure 6C). Moreover, all O 1s spectra showed broad overlapping peaks around 532 eV due to C-O and C=O bonds on the surface of each graphite. However, the oxygen content in each graphite varied significantly, which was reflected in the different intensities of the O 1s peak. Among them, SLC T had the highest surface oxygen content at 4.2%, and a highly disordered surface was observed in the Raman spectrum as well (Figure 6b). On the other hand, A12 had the lowest oxygen content at 2.3%. Electrolyte decomposition and SEI formation depend not only on surface crystallinity and surface area, but also on functional groups at the surface of carbonaceous materials.52 The presence of oxygen groups at the graphite surface is beneficial for the formation of a robust SEI layer, since oxygen-containing groups are electrochemically reducible and serve as nucleation sites for effective SEI formation. The lack of oxygen groups during the first electrochemical cycle hinders the electrolyte reduction process, prevents effective surface passivation, and consequently causes exfoliation of the graphite.52 Fresh anodes were also analyzed for comparison. No Li, P, or transition metals were detected on pristine anodes (Figure S14a). Figure S14b displays the C 1s, O 1s, and F 1s spectra from pristine electrodes. In addition to the prominent peak at 284.6 eV coming from C-C bonds in graphite and carbon black, two more features at &#;286 eV (CH2-CF2) and &#;290.5 eV (CH2-CF2) were due to the PVDF binder. Fluorine from PVDF was also identified in the F 1s spectra at 687.5 eV.53 Broad overlapping oxygen peaks corresponding to C-O and C=O bonds at 532 eV were present in the O 1s spectra, consistent with our observation from pristine graphite powder samples54 The oxygen content in each graphite anode increased significantly after cycling, consistent with the formation of electrolyte decomposition products such as alkyl carbonates and alkoxide species (Figure S15).55 At the same time, the intensity of the C 1s signal from C-C bonds decreased in each cycled anode, accompanied by an increase in features coming from C=O and C-O bonds. Moreover, new peaks at 686.2 eV in the F 1s spectra were also identified in all cycled anodes, which suggests the formation of inorganic SEI components likes LiF and POyFz from LiPF6 salt decomposition.56,57 This was further confirmed from the slightly asymmetric Li 1s peak58 and the presence of P. Of particular interest is that the SEI components were distinctly different in each cycled graphite anode, though no strong correlation was revealed with capacity retention. SLC T, MAGE, and MAGE3 appeared to have more inorganic species like LiF and POyFz than SCMG-BH, APS, or A12. Notably, cycled SLC T, which was among the graphites with the best capacity retention, showed no presence of Mn from transition metal crossover from the cathode (Figure S15). Mn dissolution and crossover has been blamed for impedance rise and capacity fade, possibly by making the SEI more electronically conductive and serving as centers for further electrolyte reduction.59
FTIR is well-suited to investigate organic species and was used to further characterize the chemical composition of the SEI layers.60 A representative FTIR spectrum from a pristine SLCT graphite anode is shown in Figure S17. Besides the peaks coming from graphite and PVDF binder, -OH bending and C-O-C stretching modes were also identified, consistent with the high oxygen content on the surface of SLC T detected by XPS. By contrast, no -OH bending and very weak C-OH peaks were seen in the A12 electrodes, which also had the lowest oxygen content on the surface from XPS analysis.61,62 After cycling, both inorganic and organic SEI species were identified in the resulting spectra (Figure 6d). Note that all cycled electrodes analyzed here were recovered in an Ar-filled glove box, and the FTIR spectra were also taken inside the glove box to avoid any air exposure. To facilitate comparison, the peak at cm&#;1 was normalized to unity for each sample. Features from PVDF binder were still seen in the cycled SLC T, but not in the other anodes. This suggests that the SEI layer that formed on SLC T could be thinner, resulting in less impedance growth and active lithium loss. Li2CO3 is a product of EC reduction and has characteristic FTIR bands around 863 cm&#;1 and cm&#;1.63 Consistent with analysis of the C 1s and O 1s XPS spectra, FTIR also showed Li2CO3 as an important component in the SEI. Features around cm&#;1 could be attributed to alkoxides like CH3OLi.64 Vibrations around cm&#;1, cm&#;1 and cm&#;1 reflected the presence of alkyl carbonates like (CH2OCOOLi)2.65 Again, components in the SEI from different cycled anodes varied considerably, consistent with the XPS analysis. For example, the feature belonging to (CH2OCOOLi)2 was stronger in the cycled SLC T graphite anode, while both MAGE and MAGE3 graphite showed much stronger Li2CO3 signals (Figure 6d). In A12, APS, SCMG-BH, and SLC T, CH3OLi showed the strongest signal. Though inorganic species in the SEI layer make it more insulating and make Li diffusion more difficult, the actual thickness of the SEI layers was difficult to estimate. The function of each component in the SEI layer and its impact on long-term cycling performance has not been fully investigated. However, in general no strong correlation was found between the SEI composition and capacity retention during long-term cycling.
While these results from coin cells are informative, the impact of different graphite anodes was also evaluated in form factors that are more relevant to large-scale applications.66 Therefore, single-layer pouch cells were built using two representative graphite anodes and NMC811 cathodes. SLC T and A12 graphite were selected because they demonstrated among the best and worst performance in the coin cell tests, respectively. Pouch cells with A12 graphite anodes exhibited rapid capacity fade compared to the group with SLC T anodes (Figure 7a). After about 180 aging cycles, cells with A12 anodes retained a discharge capacity of 155 mAh/gNMC, while cells with SLC T anodes had a discharge capacity of 181 mAh/gNMC. Greater capacity fade with A12 anodes is also evident in the dQ/dV plots (Figure 7b), consistent with the coin cell study. Overall, the results presented here show that the long-term cycling stability of high-energy cells with NMC811 cathodes can be greatly improved by selection and optimization of the graphite anode.
Figure 7. a) Specific discharge capacity and coulombic efficiency of NMC811-SLC T and NMC811-A12 pouch cells. b) Differential capacity vs. cell voltage of NMC811-SLC T and c) NMC811-A12 pouch cells at different cycle numbers.
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