Li-insertion in hard carbon anode materials for Li-ion batteries

Electrochimica Acta 45 (1999) 121±130

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Li-insertion in hard carbon anode materials for Li-ion batteries
Edward Buiel, J.R. Dahn*
Department of Physics, Dalhousie University, Halifax, Nova Scotia, Canada, B3H 3J5 Received 7 February 1999; received in revised form 8 April 1999

Abstract `Non-graphitizable' or `hard' carbon anode materials for Li-ion batteries have many advantages and disadvantages when compared to graphitic materials such as mesocarbon micro-beads (MCMB). The advantages include higher capacity (per unit mass) [Yamada et al., United States Patent No. 5,834,138 (1998); Buiel et al., J. Electrochem. Soc. 147(2) (1988) 2252±2257; Buiel and Dahn, J. Electrochem. Soc. 145(6) (1998) 1977±1981], higher cycle life [Omaru et al., United States Patent No. 5,451,477 (1995)], good rate capabilities [Rakotondrainibe et al., In: Proceedings of the 194th Meeting of the Electrochemical Society. Abstract No. 83. 1±6 November 1998] and lower cost of production. The disadvantages that must be resolved before a successful material can be commercialized are the low density, incompatibility with current coating technologies, larger irreversible capacity and hysteresis in the voltage pro?le [Buiel and Dahn, J. Electrochem. Soc. 145(6) (1998) 1977±1981]. Some reports have suggested that the problem of low density can be solved using composite anode materials consisting of a mixture of hard carbon and MCMB. These composites also boast higher rate capabilities and longer cycle life when compared to pure MCMB. In this paper, reducing the hysteresis in the voltage pro?le and reducing the irreversible capacity of hard carbons is the primary focus. In order to achieve this goal, a study of the electrochemistry and structure of promising hard carbon materials is presented and correlated to various parameters that can be adjusted during synthesis. # 1999 Elsevier Science Ltd. All rights reserved.
Keywords: Lithium-ion battery; Mesocarbon micro-beads; MCMB; Lithium insertion

1. Introduction Non-graphitizable petroleum coke based carbons were ?rst to replace Li-metal anodes in Li-ion batteries. Most of the safety problems associated with using Li-metal could be traced to the high reactivity, high surface area, dendritic lithium deposits formed during cycling [3]. This form of dendritic lithium was

* Corresponding author. Tel.: +1-902-494-2312; fax: +1902-494-5191. E-mail address: je?.dahn@dal.ca (J.R. Dahn)

found to be increasingly reactive with electrolyte solutions at elevated temperatures causing signi?cant capacity loss per cycle [4]. Good cycling e?ciencies (98 to 99.5%) could be achieved, in many di?erent electrolyte solutions, provided 5 to 6 times the required amount of Li-metal was incorporated into the cell [5,6]. This leads to an increasingly dangerous situation as more and more lithium is transformed into the highly reactive dendritic state as the cell is cycled. Attempts to prevent this situation led researchers to replace the Li-metal electrode with an intercalation compound such as carbon. These new Li-ion rechargeable batteries are said to be based on the `rocking

0013-4686/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 3 - 4 6 8 6 ( 9 9 ) 0 0 1 9 8 - X

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Fig. 1. Plot of reversible capacity for lithium vs heat treatment temperature for a variety of carbon samples (open symbols, hard carbons; solid symbols, soft carbons). These data are for the second charge±discharge cycle of lithium±carbon test cells. The three regions of commercial relevance are shown. This graph has been taken from the work of Dahn et al. [7].

Fig. 2. The voltage pro?les of a typical hard carbon anode material and MCMB, a graphitic material. An electrochemical cell with an Li-metal negative and a carbon positive electrode was used to collect this data. The voltage plotted is the voltage of the carbon electrode measured with respect to the Li electrode [1].

chair' concept because both electrodes are lithium intercalation compounds. Many di?erent carbon materials with good capacity and low binding energy with respect to Li-metal have been identi?ed as suitable replacements for Li-metal. The advantages of carbon based anodes over Li-metal anodes are signi?cant: 1. The surface area of the electrode, and hence safety characteristics, remains approximately constant during cycling. 2. The reaction rate between the anode and the electrolyte is limited by the di?usion of lithium ions in the anode material. Therefore reaction rates at high temperature are much lower than those of Li-metal. 3. Carbon based anode materials do not require excess anode capacity and thus the potential for violent behaviour during worst-case scenario thermal runaway incidents is decreased. 4. The anode does not melt under normal conditions and does not produce the extremely reactive form of molten-Li produced at 1808C with Li-metal anodes. Carbon materials can be classi?ed in many di?erent ways, disordered versus graphitic, hard versus soft, hydrogen versus non-hydrogen containing and there are no doubt many other classi?cations as well. Dahn et al. [7] showed that three classi?cations were signi?cant for the use of carbon in Li-ion batteries as shown in Fig. 1. These regions include: 1. Graphitic materials. 2. Hydrogen-containing materials. 3. Single-layer hard carbons. Graphitic carbons show good capacity at low voltage,

very little hysteresis and good cycling ability. For these reasons graphitic carbons such as MCMB represent the predominant choice for battery manufacturers with few exceptions. The major drawbacks are the high cost of arti?cial graphite and various safety concerns with natural graphite. Hydrogen containing materials show good capacity but su?er from large hysteresis between change and discharge. Lithium insertion in these materials was correlated to the hydrogen content by Zheng et al. [8,9] and modelled analytically by Fischer and his colleagues [10,11]. The large hysteresis is a signi?cant detriment because it requires a higher voltage for the removal of lithium from the carbon. For a real Li-ion battery, this correlates to reduced cell potential (e.g. the lithium battery will charge at say 4 V but will discharge at only 3 V). As a result, these carbons are not suitable for practical applications. Single-layer hard carbon is the third category of carbon materials that show promise for Li-ion batteries. These materials have a disordered structure very di?erent from the neatly aligned graphene layers in graphite (note: a graphene layer is a single layer of carbon atoms arranged in an aromatic or honey-comb-like fashion). The disordered structure naturally leads to an increase in capacity because, in a nutshell, there is more space to insert lithium ions when compared to graphite. In Fig. 2, a typical hard carbon voltage pro?le is compared to MCMB which represents a standard anode material found in Li-ion batteries today. Hard carbons show larger reversible capacity but su?er from three major de?ciencies: low density, large irreversible capacity and hysteresis between charge and discharge. The larger irreversible capacity must be

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Fig. 3. Voltage-capacity pro?les for the speci?ed samples. The data has been shifted sequentially by 3 V for clarity. The irreversible capacity in hard carbon is thought to contain two components; A: related to the formation of the SEI layer and B: the component related to the reaction of Li atoms with species adsorbed from air in the nanopores of the sample.

compensated by adding excess cathode material which is both expensive and leads to cells with lower capacity. Reducing the irreversible capacity is thus one of the most important problems with hard carbon anode materials. The small hysteresis between charge and discharge must also be reduced because it leads to lower cell potentials and a more sloping voltage pro?le. In some situations these properties may be advantageous but in general it should be limited as much as possible. As our results will show, it is di?cult to reduce the irreversible capacity without similarly a?ecting the reversible capacity. In this paper, an explanation for the origins of the large irreversible capacity will be given. Ways to maximize the reversible capacity while minimizing the irreversible capacity will be explained. 2. Irreversible capacity The irreversible capacity is one of the most important factors preventing the widespread commercialization of hard carbon for Li-ion batteries. Weibing Xing et al. [12] showed that there are two mechanisms that contribute to the irreversible capacity in hard carbon prepared from sucrose: (1) reaction of lithium with electrolyte to form a passivating layer and (2) the reaction of lithium with surface functional groups or absorbed molecules that result from exposing the sample to air after pyrolysis. The electrolyte/lithium reaction will always occur to some extent unless elec-

trolytes can be found that do not react with lithium. This reaction takes place on the surface of the electrode to form a passivating layer called the solid electrolyte interphase (SEI). The surface area is more important than surface chemistry and the formation of the SEI usually contributes about 50 mAh/g to the irreversible capacity. The second component is speci?c to hard carbon and contributes 150+ mAh/g; the research presented here focuses primarily on reducing this latter component of the irreversible capacity. Two processes have been found to lead to lower irreversible capacity in hard carbons prepared from sucrose. The ?rst is the chemical vapour deposition (CVD) of a carbonaceous substance from ethylene gas at temperatures >7008C and the second is lowering the dewatering temperature used to reduce the water content of the sucrose before pyrolysis. The former technique modi?es the surface of the hard carbon before it is exposed to air and this produces a large reduction in the irreversible capacity from more than 150 mAh/g to less than 70 mAh/g. Electrochemical tests shown in Fig. 3 indicate that the irreversible capacity for the cell made from the sample pyrolysed in ethylene is substantially lower than that of the vacuum pyrolysed sample (52 versus 220 mAh/g) while the reversible capacity remains essentially the same (516 versus 511 mAh/g). The second curve (no ethylene treatment) displays the `bench-mark' charge/discharge characteristics for untreated hard-carbon materials that is consistent with numerous authors [8,12±14], i.e. showing large irreversible capacity. The irreversible capacity in hard carbon contains two contributions labelled `A' and `B' on Fig. 3. `A' corresponds to the formation of the SEI layer [15] and is present for all anode materials leading to between 30 and 50 mAh/g of irreversible capacity [12]. Component `B' is particular to hard carbon materials and can lead to anywhere from 0 to 150 mAh/g of additional irreversible capacity. Again, this component has been correlated to some interaction of the sample with air. The ethylene treatment has the e?ect of modifying the surface of the hard carbon to eliminate this component of the irreversible capacity. The carbonaceous material deposited from ethylene is a soft carbon material with high hydrogen content. The voltage pro?le of this material is shown as the ?rst curve in Fig. 3. It was deposited directly onto a nickel current collector and subsequently made directly into a cell. Such a material has low commercial value however the low irreversible capacity is very signi?cant. In many materials the irreversible capacity is strongly correlated to surface properties, e.g. the adsorption/formation of surface functional groups. Therefore coating materials with this low irreversible capacity carbonaceous material serves to reduce the irreversible capacity. This, we believe, is the mechanism

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Table 1 Measured parameters of sucrose samples dewatered and pyrolysed to the speci?ed temperature De±H2O temp. Yield after De±H2O Yield after pyrol. BET SA R 20.05 Rg IRREV. capacity REV. capacity ? ? (8C) 2108 (%) 21% (%) 21% (m2/g) 25% (A) 25A (mAh/g) 25 mAh/g (mAh/g) 210 mAh/g 140 155 170 93 87 66 23 24 24 40 9.0 16 1.88 1.94 1.88 6.61 6.98 6.66 67 93 102 555 520 560

for the reduction of irreversible capacity in hard carbon. The second method that has been found to lower the irreversible capacity in hard carbon involves lowering the dewatering temperature. The dewatering process consists of heating raw sugar in an oven at temperatures between 150 to 1808C. During this heating process the sugar will ?rst melt, decompose and solidify in a way very much like the process used to form caramel from sugar. This pre-conditioning step drives o? most of the water contained in the sugar and allows for higher yields and easier sample handling through the high-temperature pyrolysis phase of the process. Since most of the properties of the hard carbon are linked to the high-temperature pyrolysis phase of the process, it is surprising that lowering the dewatering temperature has a signi?cant impact on reducing the irreversible capacity. Samples were prepared by dewatering sucrose at temperatures between 135 and 1908C and subsequently pyrolysing to 10508C in vacuum. The results of these experiments are tabulated in Table 1. The empirical parameter R represents the 002 Bragg peak to background ratio as measured by wide angle X-ray scattering (WAXS). Essentially this parameter measures the degree of disorder between graphene sheets where higher values of R indicate a more graphitic nature. Rg

is a measure of the radius of gyration for electron density ?uctuations, measured by small-angle X-ray scattering, and correlates to a measure of the size of nanopores in the sample. The BET surface area (BET SA) is a measured by nitrogen adsorption at 77 K. The voltage pro?les of these materials are shown in Fig. 4. The results of WAXS, SAXS and CO2 gas adsorption experiments shows no signi?cant changes in the bulk or surface properties of the samples and therefore cannot explain the trend in irreversible capacity. These results show that only the ?nal HTT, and not the dewatering process, has an impact on structure as measured by the analytical tools available to us. Other gas adsorption studies showed no sign of oxidation or burn-o? that may have been present if prolonged dewatering at higher temperatures lead to the oxidation and subsequent burn-o? of the sample surface. It is possible that some type of burn-o? is still present but that it does not a?ect the surface structure enough to be detected by the CO2 gas adsorption and other studies presented here. This would be consistent with the results of oxidation shown by Xue et al. [14] where

Fig. 4. Voltage pro?les of sucrose dewatered to the speci?ed temperature followed by vacuum pyrolysis to 10508C.

Fig. 5. Thermal gravimetric analysis (TGA) of sugar carbon tablets under argon gas. The tablets had been pyrolysed at 10508C and exposed to air for di?erent periods of time as indicated before the TGA measurements [12].

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Fig. 6. Thermal gravimetric analysis (TGA) of various samples with di?erent irreversible capacity prepared from pyrolysed sucrose.

Fig. 7. Voltage pro?les of hard carbon prepared by pyrolysis of sucrose in argon gas. Heat treatment temperatures are indicated [1].

3. Reversible capacity The large reversible capacity of hard carbon makes it a very attractive material to study with good potential for use in commercial cells. The small hysteresis in the voltage pro?le degrades real cell performance and has been correlated to the residual hydrogen content [16] after pyrolysis (<0.5% by mass). The hydrogen content, and hence hysteresis, can be reduced by heating the materials to higher temperatures during pyrolysis. However, there is a critical temperature above which the hard carbon begins to show a loss in reversible capacity. Hard carbons exhibit two distinct components that contribute roughly equally to the reversible capacity: the sloping higher voltage region and a low voltage plateau near the potential of lithium metal. The former is typical of intercalation between turbostratically disordered graphene sheets found in materials such as petroleum coke derived battery carbons. The latter makes hard carbon very interesting, however, the origins of this low voltage plateau is not well understood. We believe that the low voltage plateau (capacity below 40 mV) results from the formation of small lithium metal clusters inside the nanopores found in hard carbon. This would suggest that the lithium insertion mechanism has more of an adsorption character very di?erent from the intercalation process that occurs earlier at higher voltage (1000 to 40 mV). The observed reversible capacity loss occurs primarily in the low voltage plateau component of reversible capacity. Our experimental results have shown that the loss in reversible capacity is associated with the closure of nanopores [2]. A series of samples was prepared from dewatered sucrose pyrolysed at temperatures between 900 and 14008C in argon gas. The results of electrochemical measurements on these samples is

even the most limited burn-o? can cause a dramatic increase in irreversible capacity, a decrease in reversible capacity and generally has a very negative e?ect of the shape of the voltage pro?le. In any event, the samples seem to be identical and yet the irreversible capacities are di?erent. This is a surprising result considering the ?rst component of irreversible capacity is strongly dependent on many of these parameters. In fact, for graphite the irreversible capacity is generally proportional to the BET surface area. However in this case, there seems to be undetectable changes occurring in surface morphology and chemistry that are a?ecting the interaction of the surface with air. TGA analysis has provided some evidence to link this phenomenon with the adsorption of water. Hard carbon samples exposed to air for di?erent periods of time were analysed in argon gas to temperatures of 9008C and show a substantial decrease in mass beginning at temperatures of about 1008C (see Fig. 5) [12]. This decrease in mass is more pronounced with longer exposure times and occurs at temperatures much lower than the ignition point for hard carbon (I4008C). Further TGA studies found that the irreversible capacity in these samples was proportional to mass loss in the range from 100 to 5008C as shown in Fig. 6. Judging by the temperature and the dependence on exposure time, the observed mass loss is very likely the desorption of water from the sample. This strongly suggests that the second contribution to the irreversible capacity (observed in hard carbon) is the reaction of lithium with water adsorbed on the surface of the sample to form species such as lithium hydroxide; the elimination of which remains paramount in reducing the irreversible capacity of hard carbon and producing useful materials for Li-ion batteries.

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Fig. 9. Voltage of the peak in di?erential capacity vs HTT for both charge and discharge [1].

Fig. 8. The di?erential capacity vs voltage for samples pyrolysed to di?erent temperatures between 900 and 14008C. The solid lines are guides to the eye [1].

shown in Fig. 7. By observation, the hysteresis in the voltage pro?le and irreversible capacity decreases as the HTT increases. The reversible capacity remains essentially constant up to about 11008C and then begins to decrease predominantly in the low-voltage region of the voltage pro?le. In Fig. 8, the di?erential capacity is plotted versus voltage to elucidate the capacity at low voltage. The vertical lines at about ?0.02 and 0.02 V represent the plating and stripping of lithium metal on the surface of the carbon anode. For either charge or discharge, this corresponds to the lithium metal chemical potential measured under constant current conditions. The observed shift is due to the non-zero internal impedance of the cell. The peak in both curves corresponds to the low voltage plateau or the insertion or removal of lithium at a potential near that of lithium metal. The area of this peak is a measure of capacity. The potential of the peak is a measure of the binding energy for the adsorption of lithium metal in the nanopores of the hard carbon. It is very close to the chemical potential of lithium metal; evidence for the formation of small lithium metal clusters. The observed shift in peak position with HTT is dif-

?cult to explain. It is not a result of changing cell impedance because this would shift the lithium insertion peaks in opposite directions between charge and discharge when the current is reversed. In Fig. 8, parallel lines are drawn through the discharge and charge peaks showing that the shift is in the same direction for both. These peak positions have also been plotted versus HTT in Fig. 9 and, by observation, are nearly linear and parallel. This shows that the shift in binding energy for lithium insertion is inherent to changes in the binding energy of lithium in the sample and not a result of changing di?usion kinetics of lithium into the hard carbon structure. X-ray di?raction techniques have been used to show that these materials show only a very limited change in structure when heated between 900 and 14008C. The WAXS measurements show that arrangement of graphene sheets within the sample is essentially independent of HTT. For brevity, Fig. 10 shows only the two extreme cases for samples pyrolysed at 900 and

Fig. 10. WAXS scattering pro?les for samples pyrolysed to 900 and 14008C [1].

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Fig. 13. The falling card model for hard carbon; solid lines represent graphene sheets [1].

Fig. 11. SAXS curves for Fluka graphite and samples pyrolysed between 900 and 14008C. Smooth solid lines represent the best ?t calculated intensities. The ?t for graphite contains no nanopore contribution and the observed peak at q = 0.4 represent an artifact of the measurement.

14008C. The sharp peaks are due to corundum and are artifacts of the alumina boat used to contain the sample. The work by Dahn and co-workers [17] showed that these materials are comprised of predominantly 50% single, 30% double and 20% triple layers at 9008C. The SAXS curves shown in Fig. 11 show a continuous trend to larger nanopores as the HTT is increased. Fig. 12 shows that the empirical parameter R (002 peak height to background ratio) increases from about 1.6 to about 2.1 between 900 and 14008C. This can be explained by the conversion of about a ?fth of the single layers to double and triple layers in a very mild graphitization process. Both the limited graphitization and increase in nanopore size can be explained with the simplistic `falling cards' model [18] depicted in Fig. 13. Two nanopores separated by a single graphene sheet coalesce when a

single layer becomes mobile at high temperature and aligns with neighbouring layers. This process is similar to graphitization but on a very small scale. The importance of these results is that the graphene sheets in hard carbon are, to some degree, mobile at temperatures above about 10008C. Another important result from SAXS is that the nanopore structure is always present in the sample. There is no evidence for the loss of this structure; only an increase in the average nanopore size. Therefore the capacity loss cannot be attributed to a loss of the nanopore structure. CO2 gas adsorption provides three distinct sources of evidence for nanopore closure: 1. The dramatic decrease in adsorption with increasing HTT (Fig. 14). 2. The hysteresis that appears and then disappears with increased HTT is evidence for presence of partially closed nanopores. 3. Kinetics of adsorption data (Fig. 15) shows that the rate of adsorption decreases dramatically for the range of HTTs where partially closed nanopores are present in the samples. The surprisingly large adsorption of 60 cc per gram of hard carbon corresponds to roughly one CO2 molecule for every ?ve aromatic rings (C6). This is a considerable amount of adsorption and is a good indication of

Fig. 12. The empirical parameter R and the e?ective pore size Rg determined from WAXS and SAXS data respectively, vs HTT [1].

Fig. 14. CO2 gas adsorption isotherms for samples pyrolysed between 900 and 14008C. HTTs are indicated [1].

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Fig. 15. Kinetics of adsorption data for samples pyrolysed between 900 and 14008C. HTTs are indicated [1].

signi?cant penetration of CO2 inside the hard carbon sample. This suggests that a well developed nanopore network is present within the sample that creates signi?cant accessibility throughout the material for gas molecules and perhaps also Li-ions. A simple model explaining the process of nanopores closure is shown in Fig. 16. At 9008C we believe that the nanopores in the sample are predominantly of Type 1. These nanopores are easily accessible by CO2 and as a result there is no hysteresis between adsorption and desorption and the kinetics of adsorption are relatively fast. At temperatures higher than 9008C, we believe that a combination of all three nanopores exist in hard carbon. Type 2 nanopores have an `ink bottle' shape which is believed to result in the observed hysteresis between adsorption and desorption. Many di?erent researchers have proposed models for this behaviour, however, a consensus has yet to be reached. Some of the ?rst

Fig. 16. Model of nanopore closure depicting open nanopore (Type 1), partially closed nanopores (Type 2) and fully closed nanopores (Type 3) [1].

models [19±22,28] were based on applying the capillary condensation theory proposed by Kelvin [23] to `ink bottle' shaped pores. Many researchers believe that such a model cannot be applied to nanopores because they are essentially too small to support the surface tension arguments proposed by Kelvin. The validity of these arguments is discussed by Brown et al. [24] and molecular dynamic simulations in nanopores ranging from 2 to 4 nm in size have supported the presence of capillary condensation in nanopores. Other researchers have proposed new models such as porous networks containing restrictions [25] and more advanced models of single nanopores with `ink bottles' shapes. With all the variations in the models and theories the one unifying characteristic is that the hysteresis phenomena requires the presence of some kind of restriction, albeit in the openings of single nanopores or in connections between nanopore networks. The data shown in this paper suggest that restrictions can be generated by heat treating hard carbons to temperatures of around 11008C. These restrictions lead to the presence of hysteresis between adsorption and desorption and have a further impact on adsorption kinetics. The formation of Type 2 nanopores forces CO2 molecules to di?use through narrowing constrictions which has the e?ect of reducing the rate of adsorption and desorption by the hard carbon material. Examining the kinetics of adsorption data (Fig. 15), we notice that the sample pyrolysed to 12008C exhibits the slowest rate of adsorption. In Fig. 16, we also notice that this sample shows the largest amount of hysteresis between adsorption and desorption. Since both of these phenomena are attributed to Type 2 nanopores, it is likely that this sample contains the largest number of these nanopores. The third type of nanopore is a closed nanopore not capable of adsorbing CO2 in its interior. The presence of this type of nanopore acts only to reduce the total amount of adsorption. We believe that samples pyrolysed to temperatures of 14008C and above show primarily this type of nanopore. In Fig. 15 the kinetics of adsorption between this sample and the sample heat treated to 9008C are very similar. From the model we would not expect to see a signi?cant di?erence in the kinetics of adsorption for samples comprised of primarily Type 1 or 3 nanopores because both have relatively open surfaces for adsorption when compared to a Type 2 nanopore. The total amount of adsorption is, of course, dramatically di?erent between these two samples because in one case the nanopores are open and in the other the nanopores are closed. This simple model of nanopore closure explains the gas adsorption results quite well on a qualitative basis. As the samples are heated to higher HTTs, the nanopores close and form what we call `embedded fullerenes'. Based on results of lithium intercalation in C60

E. Buiel, J.R. Dahn / Electrochimica Acta 45 (1999) 121±130 Table 2 Electrochemical and physical properties of the pyrolysed sucrose samples HTT (8C) 2108 900 1000 1100 1200 1300 1400 BET SA (m2/g) 25% 121 35.3 13.1 9.0 9.9 9.0 IRREV. capacity (mAh/g) 210 mAh/g 164 128 105 38 30 26 REV. capacity (mAh/g) 210 mAh/g 527 497 494 462 408 299

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Capacity ratio rev./irrev. 3.2 3.9 4.7 12.2 13.6 11.5

[26,27], this reduces the lithium insertion capacity because Li-ions cannot penetrate into the interior of the fullerene and the surface area for lithium adsorption is decreased. The gas adsorption results show that nanopore closure is occurring at the same HTT where the capacity for lithium insertion is being reduced. Hence we believe that this is at least partially responsible for the observed reduction in capacity. The loss of reversible capacity in the low voltage plateau region is almost certainly tied to the closure of nanopores. The long low voltage nature at potentials very near the chemical potential of lithium metal is clear evidence from the formation of small lithium metal clusters in the nanopores of the hard carbon. A molecular sieve-type interaction explains why the irreversible capacity is a?ected before the reversible capacity is reduced. From the last section, the irreversible capacity was tied to water adsorption on the surface of the carbon and we believe that as the nanopores close, the adsorption of water is reduced before the nanopore openings are too small to a?ect the insertion of Li-ions. Hence this creates an optimum range between about 1100 and 13008C where the highest reversible to irreversible capacity ratios are achieved as shown in Table 2. 4. Conclusion The irreversible capacity in hard carbon has been studied and two methods have been described to reduce the irreversible capacity. This work has led to hard carbons with irreversible capacities that are similar to other anode materials such as MCMB. Both the ethylene surface treatment and dewatering temperature experiments suggest that the irreversible capacity is correlated to the surface of the hard carbon sample and most likely the adsorption of water in the nanopore structure. Reducing the hysteresis in the voltage pro?le by heating the sample to higher HTTs results in the loss of capacity above a critical temperature of about 11008C. X-ray di?raction analysis has shown that a

type of restructuring, similar to graphitization but on a very small scale, occurs in hard carbon at HTTs above about 10008C. Measurements of di?erential capacity versus cell potential show that there is a shift in the lithium insertion potential towards the chemical potential for plating of lithium metal on the surface. At temperatures greater than 12008C it appears that a signi?cant proportion of the lithium insertion binding energies are lower than this threshold. At this same temperature, gas adsorption results show that the nanopores in the sample begin to close and form what we call `embedded fullerenes'. At present, it is unclear how the shift in lithium insertion potential and nanopore closure are related however a signi?cant proportion of the insertion capacity does depend on the accessibility of Li to these micropores.

Acknowledgements This work was supported by NSERC and 3M Canada Co. under the auspices of the NSERC/3M Canada Co. Industrial Research Chair in Materials for Advanced Batteries.

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