Staging Phenomena in Lithium Intercalated Boron-Carbon. - ACS

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Staging Phenomena in Lithium Intercalated Boron-Carbon. Bora Kalkan, and Engin Ozdas ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b19142 • Publication Date (Web): 10 Jan 2019 Downloaded from http://pubs.acs.org on January 15, 2019

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ACS Applied Materials & Interfaces

Staging Phenomena in Lithium Intercalated Boron-Carbon. Bora Kalkan†#* and Engin Ozdas†‡ †Advanced

Material Research Laboratory, Department of Physics Engineering, Hacettepe University, Beytepe, Ankara 06800, TURKEY. #Advanced

Light Source, Lawrence Berkeley National Laboratory, Berkeley, 94720 CA, USA.

KEYWORDS: Staging, diffraction, domain, refinement, high pressure. ABSTRACT: Hole-doped LiBC has become of great interest as a candidate for the high temperature superconductivity with its structural similarity to MgB2, and as possible cathode for rechargeable lithium batteries with the large graphene-like BC hexagons. The limitation on synthesis studies has detracted from the structural, electronic properties, and application studies of LixBC. Here, we present the successful synthesis of hole-doped Li0.5BC, and its structural and electronic characterization using detailed structural modelling based on a unique stage-2 Daumas-Hérold type domain structure consistent with x-ray diffraction data. The strong intra-layer motion of the guest Li island induce extraordinary structural properties, such as staging and staging disorder and are confirmed by the compressibility results obtained from Li0.85BC (anisotropic cell compressibility with B0 ~ 134 GPa) and Li0.48BC (isotropic cell compressibility with B0 ~ 190 GPa). Lideficiency increases the conductivity; however, the temperature dependent conductivity is dominated by the thermal excitation of carriers in a strongly-disordered regime with well-defined localized states.

Introduction Graphite like Boron-substituted Carbon materials are of wide-ranging importance in material science in the areas of oxidation resistant1, super-hard materials2 and hydrogen storage3. Layered structure of these materials not only provides tuning of doping capability, electronic and mechanical properties, but also of high energy storage capacity for Lithium (Li) ion batteries4,5. Investigation of advanced anode and cathode materials in Li ion batteries for heavy applications, such as electric vehicles, is based on the selection of relevant systems which have weaker and stronger binding strengths with Li, respectively5. Recent experiments as well as ab initio calculations indicate that Li can effectively intercalate the graphite-like BC3 structure and form Li2BC3 stoichiometry, and a large electron transfer from Li to B atoms result in a potential barrier decrease for Li movement4-8, suggesting a promising anode material. On the other hand, the intercalation potential is found to be significantly higher (~2.3 V) in Li0.5BC than that in graphite with the computational simulations based on density functional theory (DFT)9, which predicts it as a potential cathode material. Several attempts have focused on the reduction of the Li content in LiBC by different techniques, such as vacuum annealing de-intercalation reactions10, 11, or oxidation of LiBC in organic solvents12, however, have been resulted in formation of poorly crystalline multiphase LixBC13 and decomposition of LiBC10. Most of the above studies focus on predicted superconductivity of LixBC14. No evidence of superconductivity and Li ordering has been observed in these samples (unlike Li intercalated graphite compounds) 15-17. Experimental results offer extensive information that is yet to be explored, but they do not probe directly the intercalation processes and diffusion paths. The hexagonal structure of LiBC remains stable up to 60 GPa with ~25% unit cell volume change18. However, phase stability, bulk modulus of ~Li0.5BC, and storage of Li in open host structure still remain lacking to date. Here, we present: (i) the successful moderate-temperature synthesis of LixBC phase (x~0.5), (ii) novel structural properties of ordered LixBC phases (x~0.5): formation of stage-2 layer stacking15, puckering in BC layers and staging disorder in a Daumas-Hérold (DH) type domain microstructure16,17, unique to this system11,12,19,20, (iii) higher phase stability and lower compressibility of LixBC phase (x~0.5) compared to LiBC, and relevance to its stage-2 layered structure.

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Experimental Section Synthesis. The starting materials, Li metal flakes (rod, 99.95% Alfa Aesar), amorphous Boron (powder 325 mesh, 99.95% Alfa Aesar) and graphite (rod, 99.999% National Carbon), including 20% excess Li were loaded in a custom made Mo cell placed in a Ta tube under inert Argon atmosphere (Figure 1, see Section 1, Supporting Information for the details of reaction cell design). The Ta tube was sealed by the TIG welding and encased in a sealed quartz tube, then heated at 200 ˚C (slightly higher than the melting point of Li) for 1 hour. LixBC phases were synthesized at various temperatures in the range of 850-1150˚C, lower than previously reported values to establish a better control of Li vapor pressure 11, 12, 19, 20. For the enrichment of the LixBC (x~0.5) phase, the temperatures and the periods in the heat treatments described above were changed systematically in many prepared samples. An intermediate grinding was applied at each heat treatment and the sample was pressed into a pellet form wrapped with a Mo foil. X-ray powder diffraction (XRPD) measurements at room temperature and ambient pressure. The ambient powder diffraction patterns were carried out on a 12 kW rotating anode equipped with a Huber G670 Guiner Imaging Plate Camera. The samples were mounted in 0.5 mm capillaries and the data was collected with 2θ step of 0.005˚ from 5˚ to 100˚ 2θ range. The radiation was CuKα1 and the wavelength, 1.5406(1) Å, was determined using a Si standard. Rietveld refinements of XRPD data collected at room temperature and ambient pressure. Rietveld refinements of each phase were performed with the software package DBWS 21. The background for each pattern was modelled automatically by Sonneveld-Visser method22. The background curve was subtracted from the pattern, and the residual background of each pattern was fitted by a polynomial function of order 6 in the DBWS routine. The peak profile was fitted using the pseudo-Voigt function with two parameters mixing the Lorentzian and the Gaussian terms as a function of 2θ. All refinements were started with scale factor and zero point, followed by background parameters, then lattice parameters of the main phase in the sample. The structural parameters, such as atomic positions or occupancies, were added next, while peak profile parameters and temperature factors were refined towards the end in a refinement cycle using the correlation matrix. The temperature factors and z-coordinates of Li, B and C in stage-2 LixBC were refined collectively in all samples, and their occupancies in both sub-lattices were coupled. The atomic coordinates and occupancies of the secondary phases were adapted from the references 23, 24 and 25 for Li2C2, Li2B6 and LixC6, LiC18, graphite, respectively. The peak shape parameters, zero error, scale factors, the atomic occupancies in stage-2 LixBC phase, and the lattice parameters of stage-1 and stage-2 LixBC phases were refined as sample dependent parameters. The refined scale factors were converted to the relative phase compositions by weight % using the Equation 1; w j %  S j Z j M j Vj

n

S i 1

i

(1)

Z i M i Vi

where S, Z, M and V signify the phase scale factor, number of formula units per unit cell, formula weight and unit cell volume, respectively21. Geometry optimization using Density Functional Theory. The calculations for a fully relaxed stage-2 Li0.5BC structure are performed with the Cambridge Serial Total Energy Package (CASTEP) 26 which is a density functional theory (DFT) quantum mechanical code using first principle pseudopotential planewave. The exchange-correlation energy is calculated within the generalized gradient approximation (GGA) using the scheme of Perdew-Burke-Ernzerhof (PBE)27. Ultra-soft pseudo-potentials are employed to describe the electron-ion interactions. An energy cut-off, 500 eV is used for the expansion of wave functions with 14x14x10 Monkhorst-Pack k-point mesh in Brillouin zone where as for a 2x2x1 primitive super cell containing 20 atoms of Li0.5BC. Further increase of the number of the k-points does not affect the results significantly. Optimization is considered to be converged when the total energy variation ACS Paragon Plus Environment

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between the successive iterations is smaller than 5.0x10-6 eV/atom, the residual forces are less than 0.01 eV/Å, the displacement of atoms during the optimization steps is less than 5.0x10-4 Å. High pressure X-ray powder diffraction (HP-XRPD) measurements. High pressure x-ray diffraction experiments were performed at the high pressure beamline (BL 12.2.2) of the Advanced Light Source at Lawrence Berkeley National Laboratory28. Angle-dispersive diffraction data were collected at energy of 25 keV with a well-focused (10×10 μm spot size) x-ray beam. The sample to detector distance can be as small as 200 mm to allow the detector to match the opening angle of the diamond anvil cell. A MAR345 image plate detector was used to collect diffraction images. The sample-detector distance and the detector tilt angles were measured using powder diffraction from a LaB6 standard. The x-ray beam was 99% horizontally polarized and all geometric and polarization corrections were made during the angular integration using the program FIT2D29. The Celref program30 was used to refine unit cell parameters. High-pressure was generated on samples using a standard symmetric diamond anvil cell (DAC) with 300 μm culet diamonds. Rhenium gaskets were indented to a thickness of about 50 μm. A 100-μm hole was drilled in the center of the indentation and was loaded with sample, 4:1 methanol:ethanol mixture (MEM) as the pressure transmitting fluid and a number of small spheres of ruby as pressure marker. Regarding the pressure calibration in these measurements, the ruby fluorescence spectra showed a sharp doublet throughout our measurements, with no measurable broadening of the peaks. Pressure was measured from each of ruby spheres, and differences in the measured pressure between these spheres were found to be within the experimental accuracy of the pressure scale used (~ ± 0.05 GPa). These results lead us to conclude that sample is under hydrostatic pressure in our measurements. Measurements in the scanning electron microscope (SEM). The morphologies of samples are characterized by scanning electron microscope using TESCAN GAIA3 FIB/SEM equipped with a STEM detector. All images were acquired at an accelerating voltage of 5kV with a spot size of 1.7 nm. X-ray photoelectron spectroscopy (XPS). XPS measurements and qualitative analysis were carried out on samples at room temperature to confirm the presence and the chemical state of Li, B and C. Samples were prepared on Cu-foil and transferred to the XPS system. XPS data were collected using a K-Alpha, Thermo Scientific model electron spectrometer with an Al Kα x-ray radiation. X-ray spot size on the sample is adjusted to 300 µm in diameter. All the binding energies were corrected with C 1s reference line at 284.6 eV. Electrical resistance measurements. The temperature dependence of the resistivity for the bar shaped samples was measured in the 3-300K temperature range using a closed cycle He refrigerator and a Stanford Research SIM900 Resistance Bridge. The electrical contacts have been made using silver paste on the silver pads deposited by thermal evaporation at the sample surface. Results and Discussion The layered hexagonal structure of stoichiometric LiBC with 𝐷46ℎ (P63/mmc, 194) symmetry was first studied by Worle et. al.19 The B and C atoms alternately occupy sites within the graphite-like layers, and the unit cell along the c-axis is doubled with the stacking of two BC layers as shown in the inset of Figure 2a. The adjacent BC layers have an asymmetric type AA' stacking sequence where the boron atoms locate directly above (and below) the carbon atoms. This stacking creates two chemically equivalent interstitial sites at the center of the hexagons of adjacent BC layers. Complete occupation of these sites gives a saturated LiBC phase with two Li layers. In the high temperature disordered phases, the lithium ions randomly occupy these two sites with equal probability20. The presence of an averaged structural symmetry between these two metal layers in both saturated and disordered non-stoichiometric phases produces an l=even extinction rule for (00l) reflections in the XRPD pattern due to 63 screw axis. Figure 2a presents the XRPD pattern of a Li0.83BC sample (17% Li vacancies) that was synthesized at 1050˚C for 12 h with three intermediate grinding. The pattern shows no (00l) (l=odd) reflections, indicating P63/mmc as space group. Figure 2b shows the XRPD pattern of a LixBC phase synthesized at moderate ACS Paragon Plus Environment

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temperatures (900˚C) for 8 hours with five intermediate grindings. A new reflection at low angle (~12, 2θ), not consistent with possible impurity phases or P63/mmc space group, is observed, and does not correspond to any previously reported structure of LixBC. It appears at a d-value of ~7.1 Å, exactly equals the c-axis of a LixBC unit cell, belonging to the (001) reflection of a trigonal lattice of P3m1 (No:164) space group consistent with symmetry breaking between the BC layers due to ordered Li stacking. In this study, the synthesis temperatures were chosen in the range of 850 ˚C to 1150 ˚C, lower than previously reported values to establish a better control of Li vapor pressure 11,12,19,20. The existence of the (00l) reflection in the XRPD pattern clearly demonstrates that the reaction conditions allow the control of the reaction rate and consequently, Li diffusion into the galleries of the host material. This shows that it is possible to synthesize ordered dilute LixBC phases with different Li stoichiometries in the intercalant layers at temperatures around 1000 C in a well-sealed Ta tube. The appearance of (001) reflection in the XRPD pattern clearly reveals the existence of stage-2 ordering in LixBC as observed in the intercalation of guest species into a layered host material such as graphite17 and transition metal dichalcogenide31. Ordered sequences of host and guest layers lead to “staging” phenomena, and are common features of layered intercalation materials15. For a stage-n compound, each intercalated layer is separated by n layers of the host material. The formula unit of a stage-2 compound in this system is therefore Li0.5BC which has alternating full/empty interlayer spaces as illustrated in the inset of Figure 2b and Figure 3a, and a lower symmetry, 𝐷33𝑑 and the same AA' stacking arrangement of the BC layers as LiBC. Refinement attempts with a pure stage-2 model, however, did not improve the agreement between observed and calculated peak intensities (Supporting Information Figure S1a). The attempts to change the local symmetry by refining the structure in sub- and super-space groups of P3m1 were unsuccessful. The advantage of using of P3m1 space group consists in the possibility to express approximate local Li symmetries and the refinements tested the occupation of all Li sites for a possible mixed-stage phase formation. Although no higher staged (n>2) intermediates could be detected by testing all primitive reflections, the co-existence of stage-1 LiBC (Figure 3a) with stage-2 Li0.5BC (Figure 3b) is evident in the selective line broadening of (0hl) and (00l) reflections. In stage-1 LiBC, every gallery in an ideal stage-1 compound contains a layer of guest species. The Columbic interaction between the alternate layers in only attractive with the ideal stacking, Li/BC/Li/BC/Li of stage-1 LiBC structure. In well-ordered stage-2 structure, every second layer is occupied by Li. The absence of one Li layer changes the sequential stacking to Li/BC//BC/Li ( implies vacant layer) and causes a strong short-range repulsive interaction between the adjacent negatively charged BC layers around the vacant layer. The attractive interaction between Li and BC layers generates an additional interaction in the same direction. These two potentials must be only compensated by an increase in the space between the layers around vacant galleries in order to minimize the lattice energy. However, a mixed-stage model including stage-1 and stage-2 phases generates more intense (001) and (012) reflections (Supporting Information, Figure S1b). The intensities of these reflections can be tuned by the occupation of empty BC galleries belonging to empty Li site, (0,0,1/2) in the unit cell of Li0.5BC. The partial occupation of both galleries in a stage-2 compound includes intercalants between all BC layers. Most likely, the distribution of Li atoms in both galleries is due to the repulsion among intercalated islands in different layers26 to form a minimum energy structure in which the lithium ions move as far apart as possible. An iterative series of fits to the x-ray diffraction data leads to a quantitative structure with an overall composition of Li0.56BC (Figure 2b). This phase, Li0.56BC, belongs to a mixed-stage phase of 24% (w%) stage-1 LiBC and 76% stage-2 Li0.43BC indicating 14% vacancies in Li0.5BC. The refined parameters of both phases are summarized in Table 1. The crystal structure of the stage-2 LixBC phase is determined as trigonal ( P3m1 , 164) with refined lattice parameters a=b=2.7500(1) Å and c=7.1448(2) Å. The structure consists of two host sub-lattices which are occupied independently, and the normalized occupancies of Li ACS Paragon Plus Environment

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atoms in two chemically equivalent sub-lattices, optimized to 0.67(2) and 0.19(2) Li occupation. also allows to refine independently the position of B and C atoms along c-axis.

P3m1

Optimization of the z-coordinates of B and C atoms improves the goodness-of-fit from 1.16 to 1.01, a significant improvement in the refinement at the final level. Figure 4a shows the systematic variation of the goodness-of-fit with the z-coordinates of B and C atoms. Rietveld analysis of x-ray data provides an evidence of shifting and puckering in the BC layers of Li0.5BC, similar to the buckling in MgB2C227. The global minima for C and B atoms occur near zC=0.2350(4) and zB=0.2411(5), respectively that shows well agreement with the results of DFT32 geometry optimization of a stage-2 Li0.5BC structure. Figure 4b demonstrates the minimization of total energy as a function of the fractional coordinates of B and C atoms. The zB is determined as the minimums of the zB-dependent total energy and goodness-of-fit curves calculated for the optimized C positions in the total energy calculations and the refinements, respectively. The total energy curves calculated using GGA, predict the global minimums at essentially close the fractional coordinates of B and C atoms from the x-ray refinements. When taken together, the present results are in good agreement with the experimental values, with maximum deviations of only 0.3% in the fractional coordinates compared to those optimized with the GGA exchange and correlation potential33. The anisotropic electrostatic interaction results in a deformation of the local structure by buckling the BC layers20 due to the inhomogeneous charge distribution in the hexagons since B and C have different charges34 around vacancies. The relatively ‘floppy’ structure of the layers allows the transverse distortions in which three C atoms of a BC hexagon are offset more than ~ 0.05 Å from the median plane defined by the three B atoms in the same ring (Figure 4c). The isotropic interaction stabilizes flatness on BC layers in stage-1 LiBC. However, vacant layers in stage-2 Li0.5BC disrupt the equilibrium on BC layers, and the anisotropic Columbic interaction results a puckering on BC layers due to local charge difference in B and C atoms. Raman measurements show a strong evidence for the existence of BC-puckering even in saturated LiBC compounds as a result of the Li vacancy and/or BC stacking faults35. With diffusion of Li atoms into the layered BC network, the charge transfer from the intercalated ions reduces the space between BC layers around the intercalate layers and creates an average 0.09 Å local distortion at the boundaries of the intercalated BC galleries. If an intercalated sub-lattice is followed by another intercalated one, rather than a vacant one, the possible distortion doubles to 0.18 Å at the boundaries. This considerable distortion renders the structure unstable and favors a domain model over a disorder model36. Domain-A and domain-B, representing two well-ordered sequences in a possible domain model (shown in Figure 4d) simulate the galleries formed by sequential unit cells which have the same sub-lattice or occupied Li site, 1a or 1b. Each gallery occupied by Li leaves the (next) deficient and forms a well-ordered stage-2 structure in each sequence. The 0.18 Å distortion at the domain boundaries assumed to form the vacancies and the empty spheres at the domain boundaries simulate 14% vacancy in the structure (see Table 1 for details). The 6 balls in the each row of domain-A and 2 balls in that of domain-B simulate 67% and 19% Li occupations of two Wyckoff positions, 1a and 1b of Li0.43BC, respectively. The existence of the stage-2 structure in LixBC itself indicates that the combination of elastic and electrostatic forces are sufficient to form well-ordered galleries, but the partial occupation of both Li sites in the unit cell points to a domain model for LixBC as observed in the alkali metal intercalated graphite compounds16,17. The “domain” refers to 3D stacking of the intercalated layers which show a periodic structure perpendicular to the basal plane. Experimental studies in graphite intercalation compounds (GICs) provided strong evidence for the effective screening of the intercalate layers by adjacent graphite layers, and the staging phenomenon is related to a long-range lattice strain interaction rather than the electrostatic effect17. The relation between the staging and strain in graphite intercalation compounds has been studied by Safran, and Hamann37, and extended by Kirczenow38. These authors calculated the longrange interaction energies of two-dimensional islands of intercalant according to the DH domain model, and showed that the compound lowers its strain and also electrostatic energy by forming mixed-stage or randomly staged domain structures. The domain model was first suggested for Li intercalated graphite ACS Paragon Plus Environment

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compounds by Daumas & Hérold 16 and provided a good description for the arrangement of the guest in the intercalation host compounds. In this model, the staged structure consists of microscopic domains in which the intercalate layers with the host layers form a well ordered sequence within any domain, however the intercalate layers occupy the empty galleries between different pairs of host layers in adjacent domains as illustrated for a stage-2 compound in Figure 4d. A considerable separation in c-parameter is expected for the unit cell of Li0.5BC due to the anisotropic strong interaction in the vacant layers. However, the domain structure likely compensates the layer expansion in the opposite direction and results in a unit cell with a c-parameter of 7.1448(2) Å which is smaller than those optimized as 7.312(5) Å by DFT calculations in this study (Supporting Information, Section 2 and Table S1). We have investigated the phase formation of LixBC and depicted the XRPD patterns of the samples prepared at 900 C for 4 (Supporting Information, Section 3 and Figure S3a) and 6 (Supporting Information, Figure S3b) hours and the change in the weight fractions of all observed phases as a function of reaction time (Supporting Information, Figure S3c). The phase content of the sample heated for 4 hours is composed of the well-ordered stage-2 LixBC phase, stage-1 LiBC, unreacted graphite and four intermediate phases: Li2B6, Li2C2, LixC6 and LiC18. With the additional heating, the amount of all intermediate phases decreases and the fraction of the staged LixBC (x ≤ 1) phases rises (Supporting Information, Figure S3c, S3b and S2). The refined structural parameters and relative standard deviations in the phase compositions of all phases are listed in detail in Supporting Information, Section 4, Table S2, S3, S4 and S6. No dramatic shift in the position of the main reflections and no new reflections are observed in the XRPD patterns of the samples with increasing the reaction time. The detailed phase analysis has also clarified that there is no evidence for the formation of possible stable carbon rich phases such as BC3 and LixBC3 in this system which have a smaller cell (c ~ 6.8 Å39) and larger cell (c ~ 7.5 Å39, 7.65 Å20) than LixBC, respectively. The changes in the B/C ratio in BC layers gives an extra chemical degree of freedom in the crystal structure, and changes the lattice parameters dramatically due to charge rearrangement which is also demonstrated by DFT calculations (Supporting Information, Table S1). These results indicate that BC network forms only upon the LixBC phase formation. The Li-diffusion steps in this system proceeds as follows: at the early stages of the reaction, possibly, the Li reacts with amorphous boron and graphite to form Li2B624 and Li intercalated graphite compounds40, which transform irreversibly to lithium carbide, Li2C223 above 750 K40. The stage-2 phase of LixBC forms preferentially, thus the repulsion between intercalated islands in different layers makes the lithium ions move as far apart as possible in the stage-2 stacking41. As Li2C2 decomposes, excess Li becomes available in the material and starts to occupy empty galleries of the well-ordered stage-2 LixBC. This induces the domain structure, followed by the formation of stage-1 LiBC. The structure consists of two sub-lattices occupied independently by Li ions. The more energetically favored sites, sub-lattice-A, are occupied first, followed by the less favored sites on sub-lattice-B. Adding the guest Li ions would expand the domain area, and in consequence, increase the fraction of the saturated LiBC phase in the material (Supporting Information, Figure S3d) and change the stage number from stage-2 to stage-1. Our X-ray diffraction studies indicate that the formation of the saturated LiBC phase is considered to be the process whereby the stage-1 is formed from the stage-2. In Figures 5a-5d and 5e-5h the SEM images of samples, Li0.83BC and Li0.56BC are shown, respectively. Figures 5a-5d show that Li0.83BC crystallizes with plate-like morphology, which is in accordance with previous reports13, 42.The best Li0.83BC sample shows significant homogenous size distribution, and contains platelets of ~ 0.2 µm thickness and ~ 1 µm in diameter, and clean surfaces (Figure 5d). The smooth edges of crystals indicates high crystallinity of Li0.83BC (Figures 5c and 5d). The examination of Li0.56BC reveals that the plate-like morphology is still preserved to some extend (Figure 5e). However, the Figures 5f and 5g indicate that Li0.56BC has inhomogeneous crystallite size distribution, and contains smaller and bigger platelets compared to Li0.83BC. Detailed examination of the particles with high magnification images (Figure 5h) reveals a grainy surface structure in Li0.83BC, which can be attributed to a partial surface oxidation 42. Although SEM samples were handled in inert Ar atmosphere, during ACS Paragon Plus Environment

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transfer into the microscope chamber a short contact with air cannot be avoided with the instrumentation available. Nevertheless some trends could be observed. Surface composition was characterized by XPS. Figure 6a-6e compare 1s signals from Li, B, C and O of Li0.83BC and Li0.56BC. The survey spectra of the samples presented in Figure 6a confirm the presence of Li, B, and C with three characteristic peaks at ∼55 eV, ∼192 eV, and ∼284 eV corresponding to Li 1s, B 1s, and C 1s, respectively. The O 1s peak located at ∼532 eV (Figure 6e) indicates the surface oxidation of the samples, showing good agreement with SEM results. The Li 1s spectra are compared in Figure 6b where the peaks are observed at ∼55 eV and∼57.8 eV for Li0.83BC and Li0.56BC, respectively. The change in Li concentration induces dramatic change in both shape, and position of the Li 1s peak. The peak shift implies significant drop in the charge transfer rate from lithium atoms in Li0.83BC. The resulting peak is broader in Li0.56BC and is probably a result of superposition of at least two peaks, indicating contribution of different Li species such as Li 1s of LiC6 (57.6 eV 43) and some lithium-boride phases 44. The B 1s peaks at ∼191.3 eV and ∼194.1 eV for Li0.83BC and Li0.56BC (Figure 6c), respectively are evidence for the existence of B species. With a change-over from Li0.83BC to Li0.56BC the bond energy for 1s electrons of boron increases from ∼191.3 eV to ∼194.1 eV, suggesting dominant covalence nature of B-bonds 44. The C 1s XPS spectra of Li0.83BC and Li0.56BC are displayed in Figure 6d. The Li0.83BC presents an intensive peak at 284.5 eV, corresponding to the graphite-like C-atoms. The low intensity peak at ∼289 eV indicates a small degree of oxidation of the samples since C-O and C=O groups are usually located at ∼286 eV to ∼290 eV 45. The Li0.56BC peaks show substantial shift and are characterized by the binding energies of ∼286 eV and ∼292 eV. Meanwhile, the peak corresponding to the graphite-like C-atoms shows a significant decrease whereas the intensity of the lower energy one increases. These changes in the C 1s and O 1s spectra create a clear general picture of the surface oxidation, especially dominant in Li0.56BC, probably occurred during XPS measurements. Possible changes in the stacking of (002) and (001) layers is expected to be favorable for the existence of different compressibilities in both phases, especially along the c-axis. This is tested with high pressure synchrotron based XRPD measurements performed on samples which are stage-1 LiBC and mixed-stage LixBC (x < 1) for increasing pressures up to 21.3 GPa and 24 GPa, respectively. Rietveld refinements on XRPD data of the samples at 0.3 and 0.4 GPa were performed with variable stoichiometry allowed for the Li site (Supporting Information, Section 5, Figure S4). The Li site occupancy and lattice parameters were found to be 0.850(8) and a=2.7421(8), c=7.0900(9) Å, respectively for LiBC (Supporting Information, Figure S4a). Similar attempt on mixed-stage LixBC (x < 1) (Supporting Information, Figure S4c) resulted in two phases, 27.5% Li0.85BC (P63/mmc) and 72.5% Li0.48BC (P-3M1) coexist together, respectively. The refined lattice parameters were obtained as a=2.7450(9), c=7.0858(8) Å and a=2.7454(7), c=7.1600(7) Å for Li0.85BC (at 0.3 GPa) and Li0.48BC (at 0.4 GPa), respectively. Overall, we can argue that lab-scale and synchrotron based results are in good agreement within the experimental errors. The compression cycle results in a simple shift of diffraction Bragg peaks to higher angles, indicating a decrease in the unit cell dimensions for Li0.85BC (Supporting Information, Figure S4b). The results clearly reveal that Li0.85BC remains stable up to 21.3 GPa. The profile of (002) peak in the XRPD pattern of mixed-stage LixBC (x < 1) becomes highly asymmetric (Figure 7a) and it splits into two peaks (Figure 7c and Supporting Information, Figure S4d) upon compression, suggesting nonequivalent compressibility of both phases along c-axis. The lower intense right peak shifted to higher angles more, compared to higher intense left peak indicating more and less compressible c-axis for Li0.85BC and Li0.48BC, respectively. The pressure dependence of (001) peak provides us direct reliable information about the c-axis compressibility for Li0.48BC (Figure 7b). Taken together, mixed stage nature of the LixBC (x < 1) and good experimental resolution on data motivates us to identify the change of lattice parameters of both phases under compression. Ambient pressure lattice parameters were not measured in compression cycle since the DAC was loaded at ~ 0.3-0.4 GPa, and so, these values were calculated from the Birch-Murnaghan-equation of state (BM-EOS) fits on a/a0 and c/c0 for both samples (Figure 7d) and listed in (Table 2), also compared ACS Paragon Plus Environment

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with the lab-based x-ray results. The Li0.85BC is highly anisotropic, with compression along c-axis several times greater than compression along a-axis due to the strong in-plane covalent bonding in BC layers and weak inter-plane bonding, such as van der Waals interactions. Interestingly, the pressure dependence of the normalized cell parameters of Li0.48BC, a/a0 and c/c0, shows that the linear compressibility along two directions are very similar. Although c length is about 2.6 times longer than a, the difference in the compressibility is very small, suggesting a stiffening of the c-axis in Li0.48BC in comparison with Li0.85BC and nearly isotropic cell compression in the pressure range of ambient to 24 GPa. The evolution of the c/a lattice parameter ratios for both phases is shown in Figure 7e, which indicates that Li0.48BC is almost invariable with respect to pressure compared to Li0.85BC. The observed changes may correspond to rearrangement of Li subnetwork in the structure of Li0.48BC, while the in-plane covalent bonding in BC layers does not show significant difference. The measured pressure-volume (P-V) curves for Li0.85BC and Li0.48BC are shown in Figure 7f. It should be noted that the hydrostatic conditions were estimated at this relatively low pressure range in a MEM based on a previous report and our pressure measurements. Therefore, we can conclude that there is no overestimated trend on the value of B0. To better understand the compressibility of LixBC (x < 1) the experimental volume compression as a function of pressure is plotted with the results for MgB246, graphite-like BC47, diamond48, and LiBC18 determined via experimental data and first-principles calculations (Figure 7f). Our results for Li0.85BC phase show a remarkably good agreement with the calculations of Lazicki et al18. The MgB2 is less compressible, since larger Mg ions contribute to larger c-axis stiffening compared to LiBC47. The higher compressibility of graphite-like BC48 indicates the strong influence of the in-plane boron on the inter-layer interaction compared to graphite and LiBC. In the case of Li0.48BC, Li-deficiency causes to increase of interlayer interactions, probably repulsive forces in ionic character, along the BC layers and makes Li0.48BC phase harder than Li0.85BC and MgB2 (Figure 7f). The Bulk modulus, B0, was calculated in two different ways. First, it was calculated using the 3rd-order BM-EOS fits. Then, a linearized version of the equation of state, f-F plots were analyzed to compute the bulk moduli of Li0.48BC and Li0.85BC phases (Figure 7g). All of these methods for calculating the Bulk modulus, B0 produced the same general trend: B0 was enhanced in Li0.48BC phase (190-202 GPa) compared to Li0.85BC phase (126-134 GPa). The 3rd-order BM-EOS fits were obtained with the parameters B0=134(5) GPa (single phase) and B0=126(2) GPa (mixed phase) for Li0.85BC, and B0=190(7) GPa for Li0.48BC, with the pressure derivative of the bulk moduli of B0'~3.7. In f-F plots, the value of FE at fE=0 is equal to the ambient pressure bulk moduli B0, which are equal to B0=132  3.2 GPa for Li0.48BC and to B0=202  4.6 GPa for Li0.85BC, with the pressure derivative of the bulk moduli of B0'~3.5 and ~4.0, respectively. The bulk modulus values of the LixBC (x < 1) phases obtained with f-F analyses compare well with the previous reports and our 3rd-order BM-EOS fits, and confirm the validity of this method. In the end, we listed all the B0 values in Table 2. These results give us confidence that the trend we report is reasonable. Li deficiency improves the electrical conductivity; however, LixBC samples do not show superconductivity above 3 K. The general trend for both phases, Li0.56BC and Li0.83BC, is semiconducting behavior (Figure 8a). The overall temperature dependence for Li0.83BC is more dramatic with a residual resistivity ratio of 9, the resistivity shows only an overall change of about a factor of 2 for Li0.56BC. The room temperature resistivity from 2.81 .cm (the reciprocal of the electrical conductivity, 0.36 -1.cm-1) to 0.045 .cm (22.22 -1.cm-1 as the electrical conductivity) shows drastic changes with increasing of the hole doping on the BC network. The electrical conductivity was analyzed in terms of band and hopping mechanisms46 since it seems different above and below ~50 K, allowing us to develop a qualitative model assuming three distinct mechanisms acting in parallel; =1/ρ=LH

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where  L and H are the residual, low temperature and high temperature conductivities, respectively. Figure 8b shows four models including thermally activated hopping (TAH), σ=σ0exp-(Ec/kBT), and/or Mott-type variable range hopping (VRH) conduction, σ=σ0exp-(T0/T)1/4 49. Ec is the critical energy separating localized from non-localized states, and T0 is material dependent characteristic temperature. The high temperature conductivity terms of both phases show VRH character similar to c-conductivity of exfoliated highly oriented pyrolytic graphite (HOPG) above 50 K17,50 frequently explained with hopping conduction induced by phonon- and/or impurity-assisted hopping between localized states formed by stacking faults, in-plane defects and/or impurity sites50-52. The conductivity fits to both phases indicate that the hopping mechanism is also characteristic of the transport properties at lower temperatures. Moreover, a TAH process with a Ec of 2.2 meV gives a marginally better fit than a VRH process below 50 K in Li0.56BC. The conductivity data of both phases fit better to a unique model with a temperature-independent residual conductivity term (σ0>0) (Figure 8b). The residual resistivity is mainly a measure of the imperfections attributed to structural defects and impurities or disorder relevant for superconductivity53. DH domain-wall scattering mechanism may provide a considerable contribution to the residual resistivity54 of Li0.56BC. Two orders of magnitude improvement in the conductivity of Li0.56BC indicates that the Li deficiency enhances the conductivity. An exhaustive study is necessary to investigate the relation between the domain size, the electronic structure of the domain boundaries and the transport properties of LixBC system. The observation of superconductivity on CaC6 and YbC655 reveals the importance of the perfection of host structure on transport properties of layered materials. This demonstrates the effect of crystal imperfections and disorder in the material and provides a natural explanation for the weakening of the superconductivity with increasing disorder in graphite compounds. In the case of 2D systems, the transport properties are extremely sensitive to the imperfections56,57 and all carriers are believed to be localized even for the smallest possible level of disorder58. The irregular elastic strain associated by the puckering in BC layers can promote this sensitivity and generate inhomogenously broken crystal symmetry. This would be sufficient to suppress bulk superconductivity. It is hoped that the data presented here will stimulate further theoretical and experimental interest to resolve the non-metallic, and non-superconducting properties of the hole doped forms of BC compounds; LixBC, Mg1-xLixBC59 and most lately, Ca1-x(BC)6 and Ca160 x(BC)8 . Conclusion We have synthesized hole doped stage-2 LixBC (x~0.5) at lower temperatures as the next stable compound of Li-BC family. It is possible to synthesize ordered mixed-stage LixBC phases in a solid solution form with different Li concentrations in the host layers as a result of staging and staging disorder. The structural stability of Li0.85BC and Li0.48BC phases have been studied with single and biphasic samples as a function of pressure up to 24 GPa using synchrotron x-ray diffraction. We have explored that hexagonal form of these phases remain stable in the studied pressure range. Stage-2 structure of Li0.5BC phase is accompanied by the change in the stacking of (00l) layers which has also resulted in different elastic and electrical properties. The bulk modulus for Li0.85BC determined from EOS agrees quite well with the recent spectroscopic results and theoretical predictions. Li0.85BC indicates a large anisotropic cell compression, whereas the Li0.48BC is almost isotropic and c/a ratio is nearly insensitive to pressure. The linear compressibility along the lattice axis suggest significant stiffening of the c-axis in Li0.48BC. The bulk modulus of Li0.48BC was determined to be B0=190(7) GPa, which is, to date the only and most accurate experimental data available. These results also present experimental evidence shows that Li0.48BC is more durable against phase transition under pressure at least. A two orders of magnitude improvement in the electrical conductivity of Li0.56BC indicates that the Li deficiency enhances the conductivity as expected. However, it is not ACS Paragon Plus Environment

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sufficient for metallic behavior and superconductivity due to the microstructure of Li0.56BC which has substantial parallel high-conductivity BC layers which are frequently separated by local barriers as a natural result of the domain structure and stacking faults which all contribute to the transport properties. We hope that the knowledge gained through this work will provide a new basis for the significantly broader applications of hole doped LixBC which may be used for hydrogen storage or as a potential cathode material in high capacity LixBC-Li ion batteries. ASSOCIATED CONTENT Supporting Information. Design and operation of inert reaction cell (Section 1); the steps of the Rietveld refinement of dilute LixBC sample(Figure S1-S2); geometric optimization (Section 2, Table S1); LixBC phase formation (Section 3, Figure S3); structural parameters of the secondary phases and LixBC phases from the multiphase refinements (Table S2-S6); high pressure x-ray diffraction patterns (Figure S4). This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * [email protected]

Present Addresses ‡IMTEK

Cryogenics, Kahramankazan, Ankara, TURKEY.

Author Contributions The synthesis routes were developed by E.O. and carried out by B.K. The ambient and high pressure x-ray data were analyzed by E.O. and B.K., respectively. The electrical conductance data was refined by B.K. and E.O. The DFT calculations were done by E.O. The manuscript was written by E.O. and B.K., and both authors commented on the paper. E.O. directed and coordinated the project. All authors have given approval to the final version of the manuscript.

Notes

The authors declare that they have no competing financial interests.

ACKNOWLEDGMENT The authors would like to thank T. Siegrist for critical readings and useful discussions of the manuscript. The electron microscopy investigations were performed at the electron microscopy center of HUNITEK, Hacettepe University. BK acknowledges Evren Cubukcu for experimental support during SEM measurements. The authors are also acknowledged Bilkent University, National Nanotechnology Research Centre for XPS measurements. Support for BK was provided by The Scientific and Technological Research Council of Turkey fellowship under Contract No. 114C120. This work was supported by the State Planning Organization under contract DPT-03K120570-05-6; by Hacettepe University Research Council, 01G011, 04D11602001 and 0801602008; by Turkish Scientific and Research Council, TBAG-AY355. The Advanced Light Source is supported by the Director, Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. This research was partially supported by COMPRES, the Consortium for Materials Properties Research in Earth Sciences under NSF Cooperative Agreement EAR 11-57758.

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Table 1. Structural Parameters of stage-1 LiBC and stage-2 Li0.43BC from Rietveld Refinements of Li0.56BC Atom Li C B

Site 2a 2c 2d

LiBC† x y 0 0 2/3 1/3 1/3 2/3 Li0.43BC†† x y

Atom Site Sublattice-A Li1 1a C1 2d B1 2d Sublattice-B Li2 1b C2 2d B2 2d †Space group, P63 mmc

0 2/3 1/3

0 1/3 2/3

z 0 1/4 1/4

N 1 1 1

z

N

0 0.2350(4) 0.2411(5)

0.67(2)* 0.77(1) 0.77

0 0 0.5 0.19(2) 2/3 1/3 0.2650 0.23 1/3 2/3 0.2589 0.23 (No. 194); a = b = 2.7488(1) Å,

c = 7.0930(2) Å; Z=2 formula units /unit cell; calculated density, 2.130 g/cm3. Refined Displacement parameters; Uiso (Li) = 0.0141(3) Å, U11, 22 (B,C) = 0.0149(1) Å, U33(B,C) = 0.0380(5) Å. ††Space

group, P 3m1 (No. 164); a = b = 2.7500(1) Å, c = 7.1448(2) Å; Z=2 formula units per unit cell; calculated density, 1.832 g/cm3; Uiso (Li) = 0.0158(4) Å, U11,22 (B,C) = 0.0386(1) Å, U33(B,C) = 0.0461(8) Å. Converged figure of merit parameters, Rp (%)=4.59, Rwp(%)= 6.31 and S (or 2)=1.01 *Estimated standard deviations of the last digits in the parentheses are from the refinement.


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Table 2. BM-EOS fitting parameters for both phases, LiBC and Li0.5BC.

Phase

B0 (GPa)

Bc0 (GPa)

Ba0 (GPa)

Li0.85BC

134±5

51±2

318±13

65±7

229±10

126±2

132 (B0 calculated as 3.5) ’

Li0.48BC

190±7

Phase

202 (B0’ calculated as 4.0) V0 (Å3) c0 (Å)

a0 (Å)

Li0.85BC

46.22±0.10

7.093±0.001

2.743±0.001

46.06±0.12

7.059±0.018

2.745±0.001

46.13±0.04

7.103±0.006

2.736±0.001

Li0.48BC

179±11

207±10

Calculated from BM-EOS single phase BM-EOS mixed phase f-F plot BM-EOS mixed phase f-F plot Calculated from BM-EOS single phase BM-EOS mixed phase BM-EOS mixed phase

B0’, Bc0’ and Ba0’ are fixed as 3.7 throughout the BM-EOS fittings.


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Figure 1. Design of the inert reaction cell mounted on a keyless chuck and simulation of thermal gradient during spin welding. The rotational speed of the cell is fixed and ~4 rpm.


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112

022 114 016

112

202,022 114 016,106

110

004

100,010 011,101

102, 012

110

004

012 010 011 002 001

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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002


Figure 2. XRPD patterns of disorder Li0.83BC (a) and Li0.56BC (b) and the corresponding single- and multi-phase Rietveld refinements. The dots and red lines are observed and calculated patterns, respectively, and ticks denote allowed Bragg reflections. The bottom trace (green line) shows difference between the observed and calculated patterns. The inset of (a) shows the unit cell of saturated stage-1 LiBC. Li occupies all possible sites in a hexagonal structure of 𝐷46ℎ symmetry (P63/mmc). The inset of (b) shows the dilute stage-2 Li0.5BC structure; which has an alternation of full and vacant metal layer spaces with a lower symmetry, 𝐷33𝑑 ( P3m1 ).The region (20.0˚- 23.5˚) including the impurities, Li2B6 (solid squares) and LiC6 (solid dots), were excluded from the refinement.


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Figure 3. The Rüdorff model of staging. (a) stage-1 LiBC (b) wellordered stage-2 structure. The lines represent BC layers while the violet balls symbolize Li atoms. The vertical and horizontal (10 Li atoms in one row) arbitrary scales are selected for an easy comparison of the models.


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Figure 4. The fractional coordinates of B and C in LixBC, and Daumas-Hérold (DH) domain model. (a) the goodness-of-fit, S (2) from the refinements (b) the total energy of Li0.5BC as a function of the c-axis fractional coordinate of B atom, zB at the C atom positions from zC=0.22 to zC=0.25 in Li0.43BC. The dots show the data points, and each solid line is a polynomial fitting to the corresponding data set. The minimum energy (marked with the bars) and the goodness-of-fit values from the fits for each zC value are plotted as a function of the zC in the insets. (c) The deformation on the hexagons. (d) Daumas-Hérold (DH) domain model for the refined Li0.56BC structure; with the vacancies at the domain boundaries shows the coexistence of a mixed-stage structure of stage-1 (stg-1) and stage-2 (stg-2) which has two different domain sizes and shows layer steps at the domain boundaries.


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Figure 5. SEM images of Li0.83BC (a-d) and Li0.56BC (e-h) with different magnifications.


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Figure 6. XPS (survey (a), Li 1s (b), B 1s (c), C 1s (d), and O 1s (e)) profiles for Li0.83BC and Li0.56BC.


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Figure 7. (a) (002) reflection in XRPD patterns collected from single stage and mixed-stage LixBC samples, at 0.3 GPa (black) and 0.4 GPa (red), respectively. (b) (001) reflection of Li0.48BC (denoted as ~Li0.5BC) as a function of pressure. (c) Splitting of (002) reflection into two peaks under pressure. (d) Pressure dependence of normalized lattice parameters, a/a0 and c/c0 for Li0.85BC (denoted as LiBC) obtained from single phase (black dots and circles) and two-phase unit cell refinements (gray dots and circles). The red solid and empty dots represent the pressure dependence of a/a0 and c/c0 for Li0.48BC. The solid curves show the 3rd order BM-EOS fits. (e) The pressure dependence of hexagonal lattice ratio, c/a for both phases, Li0.85BC and Li0.48BC. The solid lines are linear fits guide to eye. (f) Volume compression of ACS Paragon Environment these phases as a function of pressure withPlus third order BM-EOS fits and comparison with LiBC18, MgB246, graphite-like BC47 and diamond48. (g) Plot of strain fE against normalized pressure FE, for Li0.85BC and Li0.48BC phases.

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a

b

Figure 8. Temperature dependence of the electrical resistivity of Li0.83BC and Li0.56BC phases. (a) Resistivity data in 3-290 K range (b) The measured (circles) and calculated (lines) temperature dependent conductivities of Li0.56BC, and Li0.83BC. TL , H is material dependent characteristic temperature for low (L), and high (H) temperature (T>50K) regimes, respectively, and given by T L , H =1.5α 3 /k B N(E F ) L , H 49 where α is the decay rate of the localized wave function.


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Table of Contents artwork


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Figure 1. Kalkan et al.


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112

022

114 016

112

202,022

114 016,106

110

004

012 010 011

110

004

100,010 011,101

102, 012

002 001

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


002

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ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47

5 µm

a

5 µm

e

b

2 µm

2 µm

f


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2 µm

1 µm

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g

2 µm

0.2 µm

d

h

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a

b



Staging Phenomena in Lithium Intercalated Boron-Carbon. - ACS

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