Since the late 1990s, the preparation of porous carbon materials with designed micro- and meso-pores has been widely investigated. These porous carbon materials have been applied as absorbents, catalyst supports, and electrode materials because they show superior mechanical and chemical stability, and electrical conductivity1,2,3,4,5,6,7,8,9,10,11,12,13. One promising approach to produce such porous carbon materials with controlled pore sizes is the template method6,14. First, organic carbon sources are loaded into inorganic templates to form nano-composites. Second, calcination of the nanocomposites in an inert atmosphere results in conversion of the organic carbon source to carbon. Finally, the inorganic templates are removed from the composites by chemical etching. After the etching, carbon materials with controlled pore sizes at the nanometer level are obtained as negative replicas of the inorganic templates. This method is useful to obtain carbon materials with controlled pores. However, pore-size control at the Ångstrom level has been difficult owing to the unavailability of an inorganic template that can homogeneously accommodate a sufficient amount of the carbon source in their pores. There are several other classical methods that have been used to make porous carbons, such as chemical vapor deposition on activated carbons15. However, in these methods, the maximum carbon yields after the carbonization are approximately 30%, which indicate that the original carbon precursor structures are completely destroyed during the carbonization process. The extraction of metals from carbides is a well-known technique to make porous carbons with a control of the pore size16. However, the process is costly and generates gases that are toxic and corrosive. Furthermore, these methods depend on a craftsmanship process, which is based on a number of empirical parameters. This craftsmanship process has been the greatest hinderance to obtain porous carbons with controlled pores at the Ångstrom level using the procedures reported by different research groups. Therefore, to produce porous carbons with controlled pores at Ångstrom level, methods that are simple, versatile, and reproducible are highly desired. The carbonization of molecular-based porous frameworks, such as porous coordination polymers (PCPs) or metal–organic frameworks (MOFs), and covalent organic frameworks (COFs) has been attempted17,18,19,20,21,22,23,24,25,26,27. However, in most cases, carbonization induces significant damage to the frameworks, which results in a structural change of the original pores. Previously, we successfully synthesized carbon materials with controlled pore sizes at the Ångstrom level by carbonization of assemblies of pillar[6]arenes28. Carbonization of a two-dimensional sheet assembly constructed from pillar[6]arenes, which are hexagonal prism-shaped macrocyclic compounds, yielded carbons having a pore size identical to the original pore size of the pillar[6]arene molecule. Subsequently, we developed advanced techniques for pore-size control in carbon materials by the molecular design of a precursor. Herein, we report the synthesis of carbon materials with pore-size control at the Ångstrom level by the molecular design of carbon sources (Fig. 1).

Fig. 1: Molecular design of carbon sources.
figure 1

Carbon sources (1 and 2) with three-dimensional (3D) aromatic molecules and thermally polymerizable ethynyl groups. Carbonization of these carbon sources afford carbon materials with controlled pore sizes at the Ångstrom level (C1 and C2).

As observed from the carbonization of pillar[6]arene, a thermally stable three-dimensional (3D) structure is effective to retain the molecular-level structures. Additionally, it has been shown that the introduction of ethynyl29,30 or diacethylene31 into the precursor molecule induces 3D cross-linking simply by heating, and this allows for a high yield of carbonization of the precursor molecule.

Considering the above background, we selected the thermally stable 3D aromatics, tetraphenylmethane (1, Fig. 1) and tetrabiphenylmethane (2), which are functionalized with four thermally polymerizable ethynyl groups. The four ethynyl groups induce isotropic and rigid 3D cross-linking, thus the pore size can be controlled by the length of the phenyl or biphenyl moieties. The carbon material C1 prepared by carbonization of 1 had ultramicropores (4.05 Å). The carbon material C2 from 2 had micropores (4.40 Å), which were larger than the pores of C1. Thus, by molecular design of the carbon sources, we were able to demonstrate the chemically sophisticated pore-size control of porous carbons at the Ångstrom level. The entire process of the present method involves the direct carbonization of the precursor molecules without the use of a template that would need to be removed by using toxic chemicals such as HF and NaOH. The carbonization takes place in the solid-state without significant emission of toxic gases. Furthermore, neither liquid tar nor any waste liquid caused by template removal are involved. Therefore, the proposed method is beneficial from the viewpoints of low environmental impact and sustainability. Interestingly, the porous carbon C1 derived from 1 shows high conductivity and works as an anode electrode material for not only lithium- but also sodium-ion batteries.


Synthesis of porous carbons

Tetraphenylmethane with four ethynyl groups, 1, was synthesized according to the previous work by W. Lu et al.32. Details of the synthetic procedure of the carbon source, tetrabiphenylmethane with four ethynyl groups, 2, are provided in the Supporting Information. The thermal behavior of 1 and 2 was investigated by thermo-gravimetric analysis (TGA, Fig. 2a) and differential thermal analysis (DTA) in N2.

Fig. 2: Carbonization and determination of pore sizes by gas sorption measurements.
figure 2

a Thermo-gravimetric analysis (TGA, 1: blue solid line, 2: red solid line) and differential thermal analysis (DTA, 1: blue dash line, 2: red dash line) under a nitrogen atmosphere (heating rate: 1 °C / min). b CO2 (25 °C) and c N2 (−196 °C) sorption isotherms of the powders of porous carbons C1 (blue circles) and C2 (red triangles). Solid symbols = adsorption; open symbols = desorption. d Micropore volume (MV) of C1 (blue circles) and C2 (red triangles) estimated by the DA method, plotted against kinetic diameter of molecular probes.

The DTA curve of 1 exhibited an intense exothermic peak at 230 oC (blue dash line). An exothermic peak was also observed at 242 oC in the DTA curve of 2 (red dash line). As a reference, the thermal behavior of tetrakis(4-bromophenyl)methane (without ethynyl groups), 3, was also investigated (Supplementary Fig. 6). The exothermic peak was not observed in the same temperature range for the DTA curve of 3. Therefore, the exothermic peak corresponded to the thermal polymerization of ethynyl groups. Similar exothermic peaks from the thermal polymerization of organic compounds with ethynyl groups in the same temperature range have been reported29,30. After the polymerization, both polymerized samples adsorbed CO2 and N2 (Supplementary Figs. 7 and 8); thus the porous structures were developed by thermal polymerization. From TGA, 1 and 2, after calcination at 900 oC, afforded carbon materials (C1 and C2) in high yields of 89% and 86%, respectively (blue and red solid lines). In contrast, 3 showed a large weight loss at around 300 oC, and its carbonization yield after heating at 900 oC was only 2% (Supplementary Fig. 6). From these results, thermal polymerization of the ethynyl groups resulted in the formation of cross-linked polymers, which contributed to the high yield of production of carbon materials C1 and C2 by carbonization. The carbonization yields were quite high compared with previous examples of molecular precursors. For example, carbonization of ionic liquids afforded carbon materials with a high carbonization yield of approximately 22%33. Carbonization yield of pillar[6]arene assemblies was approximately 55%28. The present high carbonization yields of 1 and 2 are comparable to those of porphyrin derivatives with ethynyl groups29,30.

Characterization of porous carbons

We investigated the carbon materials C1 and C2 obtained by calcination. The CHO ratios were precisely analyzed by high-sensitivity temperature-programmed desorption (TPD) up to 1800 oC, which enabled the detection of ppm-level H and O contents. The carbon contents in carbon materials C1 and C2 were approximately 98% (details in Supplementary Table 1), which indicated the formation of carbon-rich materials. The oxygen contents in carbon materials C1 and C2 were less than 2%. The oxygen contents of typical activated carbons are 3–8%, thus the oxygen contents of these carbon materials were relatively small compared with typical activated carbons. The origin of elemental oxygen resulted from oxidation of the dangling bonds after the carbonization. Generally, carbons after carbonization contain dangling bonds (radicals), which are easily oxidized upon exposure to air31. The Raman spectra of C1 and C2 showed carbon D- and G-bands at 1355 and 1590 cm-1, respectively (Supplementary Fig. 9), which indicated the formation of a defective carbon structure. To determine the crystallinity of C1 and C2, the powder X-ray diffraction patterns were measured (Supplementary Figs. 10 and 11). The powders of 1 and 2 showed clear diffraction patterns, which indicated the crystallinity of 1 and 2. After calcination, the sharp peaks disappeared, which indicated that the crystal structures of 1 and 2 became amorphous. At the same time, broad peaks were observed at around 22o and 45o, which corresponded to the carbon (002) and (10) planes, respectively, and are typical in nanoporous carbon materials.

Pore size of porous carbons

The structures of carbon materials C1 and C2 were observed by transmission electron microscopy (TEM), as shown in Fig. 3.

Fig. 3: TEM images.
figure 3

TEM images of porous carbons of a C1 and b C2.

Unlike conventional activated carbons, graphene stacking structures were not found. Instead, numerous micropores with a surprisingly uniform size of less than 1 nm (white dots) were observed in both samples. Thus, the origin of the pores in carbon materials C1 and C2 was not the inter-sheet spaces called slit pores which are typical in general porous carbon materials, but spherical spaces derived from the developed 3D cross-linking of 1 and 2. However, we were unable to observe a clear difference between the pore sizes of carbon materials C1 (Fig. 3a) and C2 (Fig. 3b) owing to the limitation of the TEM resolution. To investigate the porosity of the obtained carbon materials C1 and C2 in detail, we performed gas sorption measurements. Figure 2b shows the CO2 (diameter: 3.3 Å) gas sorption isotherms of the powders of C1 and C2. Both samples showed CO2 uptake, which indicated that C1 and C2 had pores larger than CO2 ( > 3.3 Å). C2 (red triangles) adsorbed a much larger amount of CO2 gas than C1 (blue circles), owing to the larger pore size and volume that were developed by the biphenyl moieties. When nitrogen gas was used (Fig. 2c), the powder of C2 adsorbed N2 (diameter: 3.7 Å, red triangles), but that of C1 did not (blue circles). C2 adsorbed ethane (diameter: 4.0 Å) and n-butane (diameter: 4.3 Å), but C1 hardly adsorbed ethane and was unable to uptake n-butane gas (Supplementary Fig. 12). The pore sizes of C1 and C2, which were determined by the molecular probe technique, were 4.05 and 4.40 Å, respectively (Fig. 2d). Therefore, we were able to successfully control the porosity of porous carbons by the molecular design of carbon sources. We also investigated the surface area by using N2 gas and the Brunauer−Emmett−Teller (BET) method (Supplementary Fig. 13). The surface area of C1 was much lower than that of C2 (8.21 vs 437 m2 g−1). C1 would not have continuous pores, thus N2 molecules were unable to access the inner pore spaces of C1 at −196 °C. In contrast, the size of the C2 pore was larger than that of N2 molecules, thus N2 molecules were able to access the continuous inner pore spaces of C2. Therefore, the surface area of C1 calculated by the N2 molecule probes was too small compared with that of C2.

We prepared porous carbons from another aromatic precursor spirobifluorene with four acetylene groups (4). In this case, the obtained carbon adsorbed CO2 and ethane (Supplementary Fig. 14), and had pores of 4.07 Å (Supplementary Fig. 15), which were a similar size to those of C1 because the acetylene-modified spirobifluorene 4 has a 3D structure, that is similar to 1. From the result, it was revealed that carbonization of 3D structures with acetylene groups is a useful way to prepare porous carbons with controlled pores at the Ångstrom level.

The thermal stability of these porous carbons is a very important point for practical applications. When heated again at 900 oC under nitrogen atmosphere (Supplementary Fig. 16a), the porous carbon C2 which was synthesized by a simple template-free pyrolysis, did not show any loss of porosity. Such a stable behavior is comparable to that of conventional activated carbons such as MSC-30 (Supplementary Fig. 16b), and is superior compared with the high-performance materials synthesized by template techniques. For example, zeolite-templated carbon (ZTC), an ordered microporous carbon known as a high-performance material, was not stable at 900 oC. The surface area of ZTC was decreased by ca. 40% after heating (Supplementary Fig. 16c).

Application of porous carbon materials for lithium- and sodium-ion batteries

Porous carbons that can adsorb CO2 gas and cannot adsorb N2 gas are good candidates for anode active materials; thus, we used C1 as an anode active material for a battery. When C1 was used in a lithium-ion battery (LIB; Fig. 4a), C1 showed a capacity comparable to that of graphite. The ionic volume of a lithium ion (1.84 Å3) is small, and therefore graphite can intercalate and de-intercalate lithium ions, which is the reason why graphite works as an anode active material. C1 also has a porous structure, thus C1 can absorb and desorb lithium ions (Supplementary Fig. 17), and works as an anode active material. Interestingly, when C1 was used in a sodium-ion battery (SIB; Fig. 4b), C1 exhibited a capacity approximately 1.7 times higher than that of graphite. The ionic volume of a sodium ion is 4.44 Å3, which is larger than that of a lithium ion (1.84 Å3). Owing to the large size of a sodium ion, graphite can hardly intercalate sodium ions, which results in a low capacity. However, C1 has pores, thus larger sodium ions can move in the pores of C1, which results in the high capacity. We also used C2 in lithium- and sodium-ion batteries (Supplementary Fig. 18). The capacity of C2 was approximately 0.8 times smaller than that of C1. From the gas sorption experiments, the cavity size of C2 was larger than that of C1. The cavity of C2 would therefore be too large for the lithium and sodium cations. Therefore, the pore size control by molecular design of the carbon sources in this study is very important. For application of the porous carbon as an electrode material, the conductivity is an essential property. Therefore, we measured the conductivity of C1 (Fig. 4c). C1 showed good conductivity compared with Gramax (synthetic graphite material used for lithium-ion batteries) and YP50F (activated carbon used for electric double-layer capacitors).

Fig. 4: Electronic property.
figure 4

Application of C1 as an anode active material for a lithium- and b sodium-ion batteries (1C = 372 mAh g−1). As references, capacities of graphite are also shown. c Conductivity of the porous carbon C1, Gramax (commercial anode material for a lithium-ion battery) and YP50F (commercial activated carbon for a supercapacitor electrode).


We were able to prepare carbon materials with pore sizes controlled at the Ångstrom level by carbonization of the designed carbon sources, which were 3D aromatic molecules with ethynyl groups. Introducing ethynyl groups contributed to high carbonization yields because of the formation of thermally stable cross-linked polymers. Furthermore, the carbon materials had pore sizes controlled at the Ångstrom level. The pore size of the carbon materials was able to be increased by using large 3D aromatic molecules. To date, there have been few examples of controlling the pore size of carbon materials by molecular design of the carbon sources. The novelty in this study is the preparation of porous carbons resulting from molecular design of carbon precursors. In this method, we successfully obtained porous carbons with pore-size control at the Ångstrom level by simple carbonization of the carbon precursors. Owing to the high carbonization yields, the porous carbon structures reflect the structure of the original precursor, and Ångstrom level pore-size control based on the molecular structure design is achieved. Therefore, porous carbons controlled at the Ångstrom level are produced by a simple carbonization process of the designed carbon sources. The carbonization yields of these 3D aromatic molecules are quite high; thus these carbon sources can be used as carbon sources instead of phenolic resin and poly(acrylonitrile). Compared with the template carbonization method, this synthetic method does not require inorganic templates. However, by using these 3D aromatic molecules as carbon sources in the template carbonization method, porous carbons with two different pore-size control levels at the Ångstrom and nanometer scales can be produced. We are conducting further investigations to improve on these findings.



All solvents and reagents were used as supplied. Tetrakis(4-ethynylphenyl)methane 1 was synthesized according to the previous paper32. Two reference carbon materials were used for conductivity measurement: YP50F and Gramax. YP50F is a water-vapor-activated carbon made from palm shells and was kindly provided by Kuraray Chemical Co. Ltd. YP50F is used as an electrode material for commercial electrical double-layer capacitors owing to its large surface area (1700 m2 g‒1) and sufficient conductivity. Gramax is used as an anode material for commercial lithium-ion batteries and was kindly provided by Osaka Gas Chemicals Co. Ltd.


Synthetic routes and characterization of the carbon source 2 are shown in Supplementary Figs. 14. The synthesis procedures of 2, MSC-30, and ZTC as well as the LIB and SIB preparation and evaluation, are also shown in Supplementary methods. Carbonization procedure of the carbon sources is shown in Supplementary Fig. 5.


The 1H NMR spectra were recorded at 500 MHz with a JEOL-ECA500 spectrometer. TGA was performed using a TG/DTA6200 analyzer, SEIKO Instruments, Inc. under nitrogen atmosphere. The TPD measurement was performed on a sample (ca. 1 mg) by using a home-made high-vacuum apparatus allowing quantitative detection of gas evolutions such as H2, H2O, CO, and CO2, during a heating of the sample up to 1800 °C with a heating rate of 10 °C min–1 34. Raman spectroscopy was performed on a JASCO NRS-3300FL spectrometer (laser: 532.2 nm). The electric conductivity of the carbon powders was measured by a two-probe method reported elsewhere35. Gas and vapor sorption isotherms were obtained by a BELSORP-max (BEL Japan Inc., Osaka, Japan) sorption analyzer. The surface area was calculated by the BET method. Powder XRD patterns were obtained on Smart Lab (Rigaku) diffractometer with CuKα radiation (tube voltage, 40 kV, tube current, 20 mA). The nanostructures of the samples were directly observed by a transmission electron microscope (JEM-2010, JEOL Ltd.) with an accelerating voltage of 200 kV.

Molecular probe method

Adsorption of the various probes was measured at 298 K. Molecular probes of different kinetic diameters, including CO2 (0.33 nm at 298 K), ethane (0.40 nm at 298 K), and n-butane (0.43 nm at 298 K) were used for adsorption to evaluate the pore size of the adsorbents36,37,38.