We show that Cs intercalated bilayer graphene acts as a substrate for the growth of a strained Cs film hosting quantum well states with high electronic quality. The Cs film grows in an fcc phase with a substantially reduced lattice constant of 4.9 Å corresponding to a compressive strain of 11% compared to bulk Cs. We investigate its electronic structure using angle-resolved photoemission spectroscopy and show the coexistence of massless Dirac and massive Schrödinger charge carriers in two dimensions. Analysis of the electronic self-energy of the massive charge carriers reveals the crystallographic direction in which a two-dimensional Fermi gas is realized. Our work introduces the growth of strained metal quantum wells on intercalated Dirac matter. Cesium atoms that are grown on intercalated bilayer graphene can create an ordered epitaxial film. Here, the authors report that such a strained film can host quantum well states with high electronic quality as characterized through angle-resolved photoemission spectroscopy. Graphene has turned out to be an excellent substrate for van-der-Waals heteroepitaxy and the growth of transition metal dichalcogenides such as MoS2, MoSe2, or TaS2 has been achieved on epitaxial monolayer graphene6. Epitaxial growth of thin films of conventional materials such as simple metals on van-der-Waals substrates is complicated by the notoriously low wetability of van-der-Waals materials9. Improving the wetability of graphene for adsorbed water has been achieved by doping graphene10 and changing the layer under graphene7. Yet, the growth techniques and the characterization of hybrid structures consisting of van-der-Waals materials and metals are unexplored. Its development can have a large impact on electronic structure engineering of 2D matter and extend growth techniques that use van-der-Waals materials as substrates. Here we introduce an epitaxial growth method for the synthesis of crystalline and strained alkali-metal films on top of bilayer graphene. Electronic structure characterization of the bilayer graphene/alkali-metal film heterostructure highlights the joint occurrence of graphene-derived Dirac Fermions and alkali metal-derived Schrödinger Fermions. The strain in the alkali-metal layer is a consequence of the lattice mismatch to graphene and the relatively high ductility of alkali metals. Graphene hosts ordered layers of alkali metals either adsorbed onto11 or intercalated in between individual graphene sheets12 or between the substrate and graphene13. The mechanism which dictates the alkali-metal order is the Coulomb repulsion between ionized alkali atoms. For cesium (Cs) adsorbed on or intercalated underneath monolayer graphene, a perfect order in a (2 × 2) or a . In particular, ℑΣ, the imaginary part of Σ, is proportional to the photohole scattering rate. Thus, ARPES can be used to identify a realization of a Fermi gas in a 2D alkali-metal film.Van-der-Waals heteroepitaxy15. In these conditions, Cs intercalates under and in between the graphene bilayer and even grows on top of bilayer as a crystalline, thin Cs film. Intercalation between the graphene layers has been reported in the literature21. The Cs film that grows on top adopts a highly compressed (~11%) fcc structure which is different from the Cs bcc bulk structure. This observation is reminiscent of the phase transition of many alkali metals to an fcc structure under pressure26 which has been observed for lithium27, potassium28, and rubidium29. Bulk Cs has been reported to maintain the bcc structure (this phase is termed Cs-I) down to 4 K31, and a bcc to fcc (Cs-II) structural transition has been observed at a pressure of 23.7 kbar25. The additional Cs phases termed Cs-III, Cs-IV, and Cs-V appear upon further compression34. Theoretically, ab initio methods predict that the Cs-I and Cs-II structures are degenerate to within several meV of energy and there is no consensus which is the structure of the ground state37. It has been pointed out that electronic and dynamical effects are also important for the bcc to fcc structural transition38. Moreover, the ground state structure of a thin Cs film can also be different from the bulk Cs structure. Electronic and dimensional effects on the structure highlight the usefulness of applying ARPES and ab initio calculations in tandem since the structure of a material is also encoded in its band structure. The ARPES investigation shows the coexistence of a series of parabolic Cs 6s derived quantum well states and a linear graphene-derived dispersion relation. The quantum well states have a very narrow energy broadening and analysis of the self-energy indicates that there are certain crystallographic directions in which the observed broadening is dominated by instrument resolution rather than quasiparticle lifetime. The vibrational structure of pristine and Cs-intercalated bilayer graphene/Ir(111) is characterized. We show how Raman spectroscopy can be used to identify epitaxial pristine and doped bilayer graphene from the position and Fano asymmetry of the Raman G mode of bilayer graphene.In the present work we study the interaction of epitaxial bilayer graphene with high Cs coverages by performing longer Cs evaporation onto the graphene bilayer than the previous reports39. In the upper panel of Fig. z) variation of a scan across the domain boundary. The scan of the z profile in the bilayer domain reveals that the corrugation in the bilayer domain is at least a factor 3 larger compared with the monolayer domain which is the reason for the weak moiré pattern observed in LEED. We note that for ARPES and Raman, the synthesized sample had a complete bilayer coverage while for the STM measurements shown in Fig. π bands (labels π1 and π2) instead of one π band which is observed for monolayer graphene/Ir(111). The two observed parabolic π valence bands are consistent with Bernal (or AB) stacking40 and in agreement with earlier reports on bilayer graphene on Ir(111)14. An equi-energy cut around the K point at a binding energy EB = 0.7 eV is shown in Fig. K point corresponding to the two π bands of bilayer graphene. In Fig. K point of bilayer graphene for photon energies hν = 30 eV, hν = 40 eV, and hν = 50 eV using circular polarized light. Both π bands are observed for hν = 40 eV whereas for hν = 30 eV and hν = 50 eV only one π band can be observed clearly. We attribute this variation in ARPES intensity to a matrix element effect41. In Fig. π bands from the ARPES spectrum of bilayer graphene measured at hν = 40 eV. In the upper panel of Fig. π bands. In the lower panel an energy distribution curve close to the K point is shown. Two peaks that are split by ΔE = 0.32 eV can be observed. The splitting is in reasonable agreement to tight-binding calculations with parameters fit to graphite42 and to ARPES experiments with bilayer graphene on SiC43 which yield a splitting of  ~0.4 eV. In Fig. π band structure using hν = 31 eV and linearly p polarized light. In these conditions, both π bands are visible. The visibility of the second band in these conditions can be explained by the different photon energy and polarization used, as well as by the better resolution set for these measurements44.The low energy electron diffraction (LEED) pattern of epitaxial bilayer graphene on Ir111) is shown in Fig. 11 is sho45. The observed Raman spectra of bilayer graphene in the region of the G and 2D bands show strong differences from the monolayer case performed in identical conditions15. These differences can be used for easy identification of epitaxial bilayer graphene similar to what is routinely performed on exfoliated graphene samples. For bilayer graphene, the overall Raman intensity is increased dramatically compared with monolayer and, at particular laser energies, a strong 2D mode appears only for bilayer graphene/Ir(111)45. We attribute the presence of a Raman signal for bilayer graphene to the larger corrugation of bilayer as discussed above. The corrugation increases the distance to the metal which implies a weaker screening of the electric field of the incident light. The Raman G band spectrum of epitaxial bilayer graphene is peaked at 1588 cm−1 (for T = 5 K). Comparing the bilayer G band peak maximum to the one of monolayer graphene/Ir(111) which is located at 1604 cm−1 , we see that the G band of epitaxially grown bilayer graphene is downshifted by 16 cm−1. The downshift of the C–C vibrational frequency in bilayer graphene points to an increase in the C–C bond length and is also consistent with the larger corrugation as observed by STM. Finally we note also that there is a small shoulder around 1630 cm−1 which we observe only after C evaporation onto the sample and thus we relate it to the presence of bilayer graphene.The Raman spectra of the bilayer graphene sample are displayed in Fig. −1 in Fig. 46 and on Ir15 with no or very weak 2D signal, the present epitaxial bilayer graphene sample shows a strong 2D signal when measured with blue or green light excitation (442 and 532 nm). The intensity of the 2D mode for ultraviolet (UV) excitation (325 nm) is much weaker in agreement with previous works47.The second-order Raman mode of epitaxial bilayer graphene appears in the region of 2500–2800 cmG band and 2D band intensities and lineshapes are a highly nonlinear function of the bilayer coverage: already at low bilayer coverages we start to observe the G and 2D Raman modes. Thus, UHV Raman spectroscopy is useful to characterize the growth of epitaxial bilayer graphene on Ir(111) in situ as a quick alternative to more demanding methods such as ARPES or STM.Figure 14. We believe that the flat substrate obtained by Cs intercalation is also important for the subsequent growth of Cs films. Figure ΓKM direction of the original Brillouin zone (BZ). The interesting features that can be observed from the ARPES scan are (1) a large downshift of the Dirac points corresponding to an electron doping, (2) a single π band at the K point which is broadened in the KM) direction, (3) a 2 × 2 zone folding of the band structure yielding a new Dirac cone at the EF). Let us now discuss observations (1)–(4) in more detail. First, the Cs doping causes a partial charge transfer of the 6s electron to graphene and thus a downshift in energy of the full band structure. Second, an important difference to the ARPES spectrum before Cs doping is that in the present case we have a single π band while before doping we have two π bands 16 and what has been reported for oxygen intercalated graphene48. Zone folding causes the appearance of a new BZ with new high-symmetry points that are labeled by 16 for a sketch of the zone-folded BZ). The broadening of the π conduction band of graphene along the KM) direction is a result of covalent bonding of C to Cs and also present in the calculations that we will show later. Fourth, a series of four quantum well states with parabolic dispersion can be observed around the Γ point. Considering the linearity of graphene related bands in Fig. π branch is nonlinear in the regions of k = 1.7 − 2.0 Å−1 and k = 0.5 − 0.86 Å−1. However, a linear dispersion is observed in the other branch, that is for the regions of k = 1.5 − 1.7 Å−1 and k = 0.86 − 1.0 Å−1. Irrespective of the existence of one nonlinear branch, Dirac Fermion behavior is retained for important physical properties such as the optical properties. These are determined by the electronic states far away from the K point, i.e., by states from the nonlinear bands. In the visible spectral range, graphene absorbs about 2.3% of light, independent of the photon energy up to  ~3 eV49. From Fig. Figure see Fig. . The dism1 = 1.0m0, m2 = 1.0m0, m3 = 0.8m0, and m4 = 0.6m0. Here m0 is the free electron mass. Since the bottom of the parabolic band of the quantum well state 1 has a very low ARPES intensity, a parabolic fit of this band is too unreliable to draw conclusions regarding Fermi gas behavior from it. Instead, we will evaluate the self-energy of this band later in the paper to confirm a Fermi gas behavior. The origin of the deviations from the free electron mass of quantum well states 3 and 4 could be related to zone folding of the bulk Cs band structure or the fact the Cs orbitals that make up quantum well states 3 and 4 are adjacent to graphene. Quantum well state 3 comes from Cs orbitals that belong to the first layer of Cs atoms grown on top of the bilayer and quantum well state 4 is from Cs orbitals of intercalated Cs atoms. The proximity of C and Cs orbitals can cause hybridization between them16 which is clearly visible from Fig. π bands and the Cs quantum wells is clearly visible along the π band. Importantly, we have not been able to synthesize structures with more Cs bands than observed in Fig. E = EF. In the ARPES map we observe three circular equi-energy contours around the Γ point that correspond to quantum well states 1, 2, and 4. The quantum well state 3 is not observable at EF due to its weak intensity relative to the other features. The isotropic shape of the quantum well states in wavevector space highlights that these are free electron like states. For comparison, we show in Fig. G band of Cs-doped bilayer graphene has a large downshift and a Fano asymmetric shape which arises from electron doping and is in agreement to previous works on monolayer graphene16. Performing a fit with one Fano line and one line for the feature at 1632 cm−1 that we identified already in the pristine bilayer samples . Second, its band structure must be in agreement with the experimental ARPES spectra. We apply constraints regarding the possible structures that we derive from the Cs amount deposited and the Cs structure observed. We have deposited about five atomic layers of Cs onto bilayer graphene but will consider also structures with four and six Cs layers to take into account Cs loss during the intercalation and inaccuracies of the quartz microbalance. We observed a 2 × 2 Cs order with respect to graphene which enforces an in-plane Cs–Cs distance of 4.94 Å. Since we observed four Cs-derived bands, it is clear that we have at least four Cs layers. More layers are possible if their Cs atoms are fully ionized and hence their bands are above EF. Indeed, a Cs layer intercalated in between graphene and the Ir(111) substrate is fully ionized16 and the electronic band formed by these Cs orbitals is located above EF and hence cannot be measured using ARPES. However, intercalation of a second Cs layer under graphene has not been reported and for Cs-intercalated monolayer graphene, excess Cs forms a layer on top of graphene rather than a second intercalant layer16. For the discussion of possible structures, it is convenient to adopt the following notation for the structure that consists of bilayer graphene and Cs layers which can be intercalated or adsorbed. In our notation, an ℓmn system represents ℓ Cs layers intercalated in between Ir(111) and the bottom graphene layer, m Cs layers intercalated in between the two graphene layers, and n Cs layers grown on top of bilayer graphene. All Cs layers have a 2 × 2 Cs order with respect to graphene. We have analyzed, using DFT, all possible ℓmn structures that fulfill the above mentioned constraints. In particular, we cast the above mentioned constraints as ℓ + m + n = 4,…,6, ℓ = 0, 1, m = 0,…,3, and n = 0,…,6. Applying these constraints results in the following 24 possible ℓmn structures: ℓmn = 004, 005, 006, 013, 014, 015, 022, 023, 024, 031, 032, 033, 103, 104, 105, 112, 113, 114, 121, 122, 123, 130, 131, and 132. We have systematically investigated the stability, the total energy and the electronic energy band structure of these ℓmn structures and find that the 113 phase has lowest total energy. We find that most of these structures can be excluded from being candidates that can describe the experimental results on the basis of having a too high total energy or a misfit of the electronic structure with the experiment.We employ DFT in order to identify the structure of the Cs/bilayer graphene system. The correct structure should fulfill two criteria. First, its energy must be within kℓmn phase and the reference 113 configuration is defined as μCs is the chemical potential of the Cs atom. The number of Cs atoms in the X-phase is given by EX for the six structures having lowest energy, assuming Cs-rich experimental conditions (μCs = ECs-Bulk), where ECs-Bulk is the energy per atom in the stable Cs bulk phase. Varying μCs between ECs-Bulk and ECs-atom (energy of an isolated Cs atom) in order to mimic Cs-rich and Cs-poor experimental conditions, we find that the 114 phase is favored for μCs − ECs-Bulk > 6 meV. Figure ℓmn structures.The formation energy difference between any given entclass1pt{minimaEF and that two Cs band merge with each other below EF. Since the photoemission intensity of the lowest energy Cs state is practically equal to zero around the Γ point, we cannot exclude the presence of a fifth Cs band in its proximity. Given the better agreement of the 114 band structure with the ARPES and the unknown experimental μCs value, we conclude that the experimental structure is likely the 114 phase.The improved agreement to the experiment is particularly true for the Cs-derived quantum well states. Since the 114 phase has five incompletely ionized Cs layers, we obtain five calculated Cs bands. This seems to be inconsistent to the experiment at first sight since the ARPES band structure shows only four Cs quantum well states. Nevertheless, a close look to the comparison of ARPES with the calculated band structures in Fig. 11 and the long-standing literature of intercalated few-layer graphene20 and graphite intercalation compounds19 have made a case that the intercalation of Cs in between the bilayer is not surprising. What is rather surprising in the present work is the fact that high-quality Cs quantum wells can be grown on top of such intercalated samples.Thus, let us summarize our findings on the layer-by-layer intercalation of bilayer graphene with Cs. With increasing Cs amount, the Cs will first intercalate between the bilayer graphene and the metal followed by intercalation in between the bilayer. Finally, adsorption on top of bilayer graphene will occur. From an energetic perspective, such a hierarchy of intercalation events is reasonable because intercalated Cs has a higher binding energy than adsorbed Cs. Energetically preferred intercalation of Cs was reported also for graphene monolayer/IrEa/Ec)51. This ratio is generally very low for nonalkali metals (the only notable exceptions are Nd and Sm). The low value of Ea/Ec is ultimately the reason why most of the transition metals, noble metals, and rare earth metals favor a 3D growth mode when deposited on graphene as highlighted in refs. 51. For Cs however, this ratio is relatively higher than for other metals and also higher than for other alkali metals. In particular, the Ea of Cs on graphene is comparably high with respect to other alkali metals52 and its surface energy is among the lowest in the periodic table53 . In fact, the bcc (110) Cs surface energy has been experimentally measured54 to be  ≃0.095 Jm−2. Following the technique illustrated in ref. 55, we calculated the surface energy of the (111) surface of fcc Cs −2. This is the same value reported for the hexagonal (0001) surface in ref. 55, which is not surprising given the similarity between the two stackings.Let us now discuss the reason why Cs metal grows in 2D films on bilayer graphene and does not form clusters like most other metals do. Because of their low cohesive energy, alkali metals generally possess a relatively high ratio between adhesive energy on graphene and cohesive energy . Comparing this angular FWHM to the instrument resolution of 0.1° we see that the measured FWHM is strongly affected by the broadening induced by the instrument resolution. We now proceed to the analysis of the energy dependence of the imaginary part of the self-energy. This quantity is related to the FWHM Δk and the Fermi velocity vF via ℑΣ = vFΔk. By investigation of the energy dependence ℑΣ(E) the various origins  of the lifetime broadening can be extracted. For example, impurity scattering gives a constant contribution to ℑΣ(E), electron–phonon scattering a step-like contribution56 and electron–electron scattering in a 2D Fermi gas has a quadratic energy dependence 58. Matthiessen’s rule states that these contributions to ℑΣ(E) are additive, and because of their distinct functional form, individual contributions can generally be extracted. For example, a detailed analysis of the strength of each of these contributions was performed, using ARPES, for the Sr2RuO4 compound58. The ℑΣ(E) for the Cs-derived quantum well state 1  is constant up to an energy of 150 meV where it suddenly increases. The constant value of ℑΣ(E) for E between the Fermi level and binding energy of 150 meV can be taken as an indication for a very weak or absent electron–electron scattering. The large energy range where ℑΣ(E) is equal to a constant value does not exclude the inclusion of a component due to electron–electron scattering. Nevertheless, it puts an upper limit to the contribution of electron–electron scattering which we extract from the fit shown in Fig. E) of conventional metals and their surface states  in a similar energy range as plotted in Fig. E and a “kink” due to electron–phonon coupling. Qualitative comparison of the contribution of the electron–electron scattering term for the Sr2RuO4 compound58, the surface states of conventional metals59 and the present Cs metal quantum well state 1 suggests that the present system has the lowest contribution of electron–electron scattering to ℑΣ(E). The step-like behavior of ℑΣ around 150 meV of the Cs quantum well states can in principle have two possible origins. One is that this feature looks similar to a “kink” feature that is induced by electron–phonon coupling and the other possibility is the opening of a scattering channel for photoholes to the graphene band. The graphene π band and the Cs-derived quantum well state band number 1  the quantum well states and the Dirac cones do not cross at EF. Therefore, the carriers along the ΓK direction have a weak scattering rate and this direction corresponds to the realization of a Fermi gas. For the other cuts β-δ, the crossing point between the quantum well states and the Dirac cone moves up towards the Fermi energy and scattering of quantum well states at EF increases which can be seen from the width of the quantum well states in the ARPES experiment shown in Fig. E) of quantum well state 2 is significantly higher at all investigated energies E. In particular, for binding energies E < 100 meV, the ℑΣ(E) of quantum well state 2 is equal to 150 meV. This value has to be compared with an ℑΣ(E) ~30 meV for quantum well state 1 in the same energy range. The larger ℑΣ(E) of quantum well state 2 can be attributed to the larger scattering rate in this band due to hybridization with C states . This allows Cs to match the graphene lattice parameter with an 11% in-plane compressive strain. While we can understand from these arguments why alkali metals can wet the substrate, one would expect that a Cs layer forms equally well on monolayer graphene. Nevertheless, we have not been able to grow an ordered Cs thin film on monolayer graphene under the same conditions. We thus speculate this could be related to the adhesion energy which might be lower for the monolayer. This could be due to the fact that Cs intercalates in a In conclusion, we have prepared Cs-intercalated epitaxial bilayer graphene and used it as a substrate for the growth of epitaxially strained Cs quantum wells. This material system is a realization of a vertical heterostructure between a layered material and a metal quantum well. The well-ordered growth of Cs is understood from the improved flatness of bilayer graphene after Cs intercalation and the energy balance of the Cs surface free energy and the adhesion energy to intercalated bilayer graphene. A Cs layer forms on the intercalated bilayer graphene if the surface free energy of Cs metal is smaller than the adhesion energy to the intercalated bilayer substrate. In this case, the Cs metal can wet the intercalated bilayer graphene and an ordered thin film can grow. We have also computed the surface energy of the fcc Cs(111) surface and found that it has a very low value compared with other metals. Furthermore, alkali metals are very soft materials, and in particular Cs has a very low bulk moduluss electrons as well as a Dirac like dispersion of the π electrons of graphene. The electronic structure of this heterostructure thus hosts both, Schrödinger and Dirac like charge carriers. By analyzing the ARPES data we have found that the linewidth of the Cs bands is extremely narrow and ℑΣ(E) is constant as a function of binding energy E until the value of E which corresponds to the crossing point of the quantum well state and the Dirac cone. At the crossing point, the lifetime of the carriers decreases  because of hybridization between the Cs and C bands. The sharpness of the quantum well state bands is a result of the low electron density and the constant value of ℑΣ(E) points towards a small contribution of the electron–electron scattering suggesting that this system is a realization of a 2D Fermi gas. Regarding further photoemission experiments, we expect that this system is suitable for laser ARPES measurements employing the commonly used hν = 6.05 eV. The Cs quantum well state 4  surface state that was probed using this laser source. Thus, part of the observed ARPES spectra could also be measured by laser ARPES. Such a laser ARPES experiment would benefit from better spatial, energy and momentum resolution as compared with the synchrotron light source used here. This experiment could therefore be useful for the investigation of fine details in the quantum well states such as electron–phonon coupling and spin–orbit interactions.Performing ARPES, we investigated the electronic structure and found quantum well states with a parabolic electron energy dispersion coming from the Cs 6see Fig.  has a Fe61 and an interesting future experiment could therefore be to grow Au on top of the Cs/bilayer graphene heterostructure. Another interesting future experiment is the study of the plasmon dispersion relation of the thin Cs film/bilayer graphene heterostructure. The plasmon dispersion can be studied using high-resolution electron energy loss spectroscopy and could reveal the presence of two distinct plasmon dispersion relations that correspond to 2D massive and massless particles. A knowledge of the plamon dispersion relations would also allow to assess the usefulness of this material for nanoplasmonics. Finally, it would be interesting to vary the alkali-metal type and hence the electron density in the hybrid structure.We believe that the Cs/bilayer graphene heterostructure which we have synthesized can serve as a substrate to grow quantum wells made of conventional metals and semiconductors on top. The role of Cs would be crucial in such an experiment since it could provide the large adhesion energy needed so that the conventional metal or semiconductor wets the surface. The large adhesion energy could come from deposited metals that form alloys with Cs. For example, Cs forms a well-studied alloy with Au39 followed by C deposition via electron beam evaporation onto the monolayer graphene with the sample at a temperature of 1000 °C62. The amount of carbon deposited for the growth of the second layer, i.e., the bilayer coverage, was controlled by the evaporation time. We have investigated samples between monolayer and full bilayer coverage. The optimum growth time was determined by the appearance of two sharp π valence bands in the band structure that we observed using ARPES. The quality of bilayer graphene was investigated using LEED, STM, ARPES, and UHV Raman spectroscopy. The combined application of these methods corroborates the formation of AB  stacked bilayer graphene which is discussed in the following. Additional C 1s core level measurements by X-ray photoelectron spectroscopy are shown in the Supplementary Fig. Bilayer graphene was synthesized on Ir(111) by a two-step procedure which involves growth of monolayer graphene by chemical vapor deposition63 of the SOLEIL synchrotron in St. Aubin (France) and at the BaDElPh beamline64 of the Elettra synchrotron in Trieste . The pristine bilayers  integrated in a homebuilt optical chamber45. The exciting and Raman scattered light were coupled into the vacuum using a ×50 long-working distance microscope objective with an NA of  ~0.4 and a focal distance of 20.5 mm for lasers with wavelength 442 and 532 nm. For the UV laser, UV compatible optical elements have been used. The ×20 UV objective has a focal distance equal to 13 mm and an NA = 0.32. The position of the laser on the sample was checked by a camera in the laser path. All spectra have been calibrated in position to the O2 vibration at 1555 cm−1 (see ref. 65). The O2 vibration is visible in the spectra due to the laser path outside the UHV. In Fig. 2 Raman intensity while in Fig. 2 Raman mode intensity in Fig. UHV Raman measurements were performed with the sample at −10 mbar. Image processing was conducted with the WSxM software66.For STM, a sample with an only partial bilayer coverage was grown, enabling imaging of bilayer and monolayer graphene areas in a single topograph. Sample preparation and imaging was conducted in the UHV system ATHENE with a base pressure below 1 × 1067, adopting PBE (GGA) exchange and correlation functional68. An energy cutoff of 400 eV for the plane wave basis and 14 × 14 × 1 Monkhorst-Pack grid69 for BZ sampling were used in order to ensure a total energy convergence of 1 meV. In our calculations of the Cs layer, we obtain a bulk bcc equilibrium lattice constant of 6.16 Å, ~2% larger than the experimental one30. Our results are in line with those in ref. 37, predicting that Cs-II (fcc) is the ground state (only 1 meV difference with respect to Cs-I). As pointed out in ref. 38, also the hexagonal close packed (hcp) structure, which is very similar to the fcc structure, should emerge as a competing phase at low temperature. Indeed, we find that hcp and fcc phases are energetically degenerate up to less than 1 meV. We found that the lowest energy phase contains only one Cs atom intercalated inside the graphene bilayer, while the other three stay on top of the bilayer. Other configurations, with two, three, or four Cs in between the two carbon layers are higher in energy by 0.054 eV, 0.135 eV, and 0.316 eV per unit cell, respectively. The intercalation process of one layer of Cs atoms inside the graphene bilayer lowers the energy of the system by 0.615 eV per unit cell. In our simulations, we have varied the stacking of the three Cs layers atop of bilayer graphene to find the minimum energy configuration. The three Cs layers on top can be arranged either according to the fcc(111) stacking (ABC) or alternatively according to the hcp stacking (ABA). These two configurations are similar and the energy difference between them is below 1 meV, so it is difficult to establish which one is observed experimentally. Nevertheless, the corresponding band structures are almost identical. Importantly, the in-plane lattice parameter of the Cs fcc trilayer is compressed by 11% with respect to the fcc (hcp) ground state; this, in turn, corresponds to an effective in-plane stress of 11 kbar, while the Cs–Cs out-of-plane distance is compatible with the equilibrium bulk fcc phase in the out-of-plane direction. Regarding the structure of bilayer graphene, we observe, that after the intercalation, graphene changes from AB  to AA stacking for all investigated structures. This is consistent with previous findings on Li-intercalated graphene70.We perform DFT calculations with the projector augmented-wave pseudopotential methodE1 = ECs-adsorbed − ECs-bulk. Here, ECs-adsorbed is the total energy of the Cs atom on the graphene and ECs-bulk is the total energy per atom of Cs in its equilibrium bulk form. The result ΔE1 = −0.092 eV means that the adsorption of Cs onto the bilayer is energetically favored. The same quantity can also be calculated increasing the number of adsorbed Cs layers. In particular, for the trilayer, the formation energy is equal to ΔE3 = (Etrilayer − 3ECs-bulk)∕3 = −0.04 eV. This demonstrates that the trilayer is still favored with respect to the formation of the unstrained bulk Cs. The same calculation for the case of four Cs layers (following the same fcc stacking) on top of the intercalated graphene bilayer results in a small positive value of ΔE4 = 0.005 eV. This would suggest that n = 3 is the limit for the Cs film thickness. Nevertheless, the analysis from the DFT section suggests that the 114 phase with n = 4 is a possibility. The value of ΔE4 is small compared with the energy corresponding to room temperature and the experimental μCs value may be higher than ECs-bulk. Moreover, the calculation does not include external effects such as the film morphology which might alter the value of the computed energies. It is therefore likely that systems up to n = 4 Cs layers can be stabilized.The thickness-dependent stability of the Cs film grown on intercalated bilayer graphene is investigated using DFT. There are two energies that contribute to the total energy of the Cs film. One is the adsorption energy which is negative and the other one is the compression energy which is positive. We show that the adsorption of strained Cs layers on the intercalated graphene bilayer is energetically favored. Let us first consider the binding energy for one Cs atom on the Cs-intercalated graphene in the 2 × 2 unit cell, defined as ΔSupplementary Information


Reaction of either Na4(H-L1)4 (2) with YbCl3 or reaction of H2-L1 with Yb[N(SiMe3)2]3 afforded yellow crystals of the highly distorted octahedral complex Yb(H-L1)3 (3).The reaction of 2-aniline (H 2-L1) with one equivalent of Na[N(SiMe3)2] in toluene afforded pale-yellow crystals of tetrameric poly[bis[μ3-2-anilinido][μ2-2-aniline]tetrasodium(I)], [Na4(C11H13N2O)4]n or [Na4(H-L1)4]n (2), in excellent yield. Subsequent reaction of [Na4(H-L1)4]n (2) with 1.33 equivalents of anhydrous YbCl3 in a 50:50 mixture of toluene–THF afforded yellow crystals of tris[2-anilinido]ytterbium(III), [Yb(C11H13N2O)3] or Yb(H-L1)3 (3) in moderate yield. Direct reaction of three equivalents of 2-aniline (H2-L1) with Yb[N(SiMe3)2]3 in toluene resulted in elimination of hexamethyldisilazane, HN(SiMe3)2, and produced Yb(H-L1)3 (3) in excellent yield. The structure of 2 consists of tetrameric Na4(H-L1)4 subunits in which each Na+ cation is bound to two H-L1 bridging bidentate ligands and these subunits are connected into a polymeric chain by two of the four oxazoline O atoms bridging to Na+ cations in the adjacent tetramer. This results in two 4-coordinate and two 5-coordinate Na+ cations within each tetrameric unit. The structure of 3 consists of a distorted octahedron where the bite angle of ligand L1 ranges between 74.72 (11) and 77.79 (11) degrees. The oxazoline (and anilide) N atoms occupy meridional sites such that for one ligand an anilide nitrogen is trans to an oxazoline nitrogen while for the other two oxazoline N atoms are trans to each other. This results in a significantly longer Yb—N(oxazoline) distance [2.468 (3) Å] for the bond trans to the anilide compared to those for the oxazoline N atoms trans to one another .Reaction of 2-aniline (H However, the only example of a transition-metal structure containing the deprotonated and unsubstituted anilido parent ligand (H-L1−) is an Ru carbonyl hydride dimer 4]n (2), and its 6-coordinate, monomeric ytterbium complex, Yb(H-L1)3 (3) are reported in this communication. The ytterbium complex 3 can be prepared by either the salt metathesis reaction between 2 and YbCl3 or by the acid–base (protonolysis) reaction of Yb[N(SiMe3)2]3 with three equivalents of H2-L1. The yields and purity of 3 are better for the protonolysis reaction aniline 4 subunits in which each Na+ cation is bound to two H-L1 bridging bidentate ligands . The tetrameric subunits are connected into polymeric chains by two of the four oxazoline oxygens (O9 and O37) bridging to Na+ cations  in the adjacent tetramer –2.3611 (17) Å, respectively]. In sharp contrast, the distance between the oxazoline N and the Na ions is much shorter for the 4-coordinate Na centres  than for the 5-coordinate Na . It is clear that the oxazoline nitrogen is much more weakly coordinating to the 5-coordinate Na cation. In fact, for the O-bridging oxazoline, this distance is so long that it is debatable whether there is a significant bonding interaction. However, this result is consistent with localization of electron density on the bridging oxygen at the expense of the nitrogen atom in the same oxazoline ring.The structure of ds Fig. 3a and 3ber Fig. 4. This re al. 2009 and Mukh al. 2010; Na: Zou al. 2013; Nd: Kan al. 2018]. SignifL1)3 (3) is a distorted octahedron where all three H-L1 ligands are distinct –77.79 (11)°, which sits between that of the diphenylamido-oxazoline  for the bond trans to the anilide compared to those for the oxazolines trans to one another . The Yb—N(anilide) distances [2.234 (3)–2.260 (3) Å] show less variation although the Yb—N(anilide) distance trans to the oxazoline N atom is slightly shorter than those trans to each other. Overall, this is consistent with a stronger trans influence for the anionic anilide nitrogens as might be expected. The torsion angles representing twisting from coplanarity of the oxazoline and benzene units all fall between −9.3 (6) and +8.4 (6)° so the distortions from planarity of the H-L1 ligands in 3 are relatively small.The structure of Yb4 subunits connected through bridging oxazoline oxygen atoms 3 units in the solid state  coordinated to a transition or main group metal. [COZFIH  coordinated to a transition metal aniline was prepared according to Gossage (20093)2]3 was prepared by analogy to the procedure of Bradley et al. (19733)2 and YbCl3 and was recrystallized from a hot mixture of hexane and toluene. NMR spectra were recorded on a Bruker AV III 300 MHz Spectrometer in sealable Teflon-valved tubes and were referenced to residual solvent resonances. The line widths at half maximum (ν1/2 in Hz) were measured for all paramagnetic resonances in 3 and are reported below. Elemental analyses were performed by Canadian Microanalytical Ltd.sage 2009 and puri al. 1973 using NaSynthesis of [Na4(L1)4]n. One equivalent of Na[N(SiMe3)2]  was dissolved in toluene (10 mL) and to this was added 1 equivalent of 2-aniline  in 20 mL toluene under vigorous stirring. The colourless reaction mixture was stirred overnight, filtered through Celite on a sintered glass frit and the solvent removed under reduced pressure to leave a tacky white solid. Recrystallization of the product from a hot mixture of toluene and hexane afforded clear pale-yellow crystals of [Na4(L1)4]n (2). Yield: 0.178 g (84%). 1H NMR : δ 7.568 , 7.067 , 6.635 , 6.62 , 6.459 , 3.954 , 1.315 2); 13C{1H} : δ 163.12 (C=N), 150.79 (arylCNH), 132.31 (5-arylCH), 130.04 (3-arylCH), 115.90 (6-arylCH), 115.15 (4-arylCH), 109.27 (2-arylC—C=N), 77.82 (OCH2), 68.58 [NC(CH3)2], 28.99 [NC(CH3)2].Synthesis of Yb(H-L1)3 (3) Method A: A solution of [Na4(L1)4]n (2)  in THF (10 mL) was added to a suspension of YbCl3  in THF (5 mL) under vigorous stirring. The suspension was stirred overnight at room temperature, filtered through Celite on a sintered glass frit and the filtrate was evaporated to dryness under reduced pressure. The yellow solid was recrystallized from a mixture of toluene and hexane at 243 K overnight. Yield: 0.102 g (63%). Method B: A solution of 2-aniline  in 25 mL toluene was prepared in the glovebox and added by Pasteur pipette to a vigorously stirred solution of Yb[N(SiMe)3)2]3  in 15 mL of toluene. The pale-yellow solution darkened to golden yellow on stirring overnight. The solution was filtered through Celite on a sintered glass frit and the filtrate was evaporated to dryness under reduced pressure. The orange–yellow solid was recrystallized from a mixture of toluene and hexane at 243 K yielding yellow crystals. Yield: 0.301 g (93%). 1H NMR : δ 88.4 , 49.4 , 47.9 , 12.86 , 11.70 , 10.93 , 10.00 , 9.30 , 1.26 , 0.96 , −2.77 , −3.89 , −5.38 , −11.2 , −11.4 , −16.0 , −24.4 , −77.2 . Analysis calculated for C33H39N6O3Yb (%): C, 53.49; H, 5.31; N, 11.35. Found: C, 53.39; H, 5.22; N, 11.11.L1)4 (2), the H atoms on N1, N15, N29 and N43 were located in a difference map and refined with distance restraints of 0.88 (1) Å. Uiso(H) were freely refined. In Yb(H-L1)3 (3), the H atoms on N1, N15 were added geometrically and refined with distance restraints of 0.88 (1) Å, with Uiso(H) = 1.2Ueq(N). H29 was located in the difference map for geometrical considerations and refined with coordinates riding on N29 with Uiso(H) = 1.2Ueq(N). All the H atoms bonded to carbon were refined in geometrically calculated positions, with C—H= 0.95 (methine), 0.99 (methylene), and 0.98 Å (methyl), and with Uiso(H) = 1.2Ueq(C) (methine and methylene) or 1.5Ueq(C) (methyl).Crystal data, data collection and structure refinement details are summarized in Table 410.1107/S2056989020005034/zl2777sup1.cifCrystal structure: contains datablock(s) Na4H-L142, YbH-L133, global. DOI: 10.1107/S2056989020005034/zl2777Na4H-L142sup2.hklStructure factors: contains datablock(s) Na4H-L142. DOI: Click here for additional data file.10.1107/S2056989020005034/zl2777Na4H-L142sup5.molSupporting information file. DOI: 10.1107/S2056989020005034/zl2777YbH-L133sup3.hklStructure factors: contains datablock(s) YbH-L133. DOI: Click here for additional data file.10.1107/S2056989020005034/zl2777YbH-L133sup6.molSupporting information file. DOI: 1996037, 1996036CCDC references:  crystallographic information; 3D view; checkCIF reportAdditional supporting information:


In this study, three-dimensional porous graphene composites loaded with MoS2 were prepared via a facile two-step method. The MoS2 nanoflower particles were uniformly dispersed within the layered graphene matrix, and a three-dimensional porous graphene supported MoS2 nanoflower battery (MoS2/3DG) was demonstrated to have superior performance to that of the pristine pure MoS2 nanoflower battery. At a current density of 100 mA/g, the MoS2/3DG delivered a reversible capacity of 420 mAh/g. What is more, it yielded a reversible specific capacity of 310 mAh/g at 2 A/g, showing an excellent rate of 73.8%. The excellent performance of the novel MoS2/3DG composite are attributed to the promoted infiltration of electrolytes and the hindered volume expansion for the porous structure, good conductivity, and robust mechanical properties of graphene.Sodium-ion batteries (SIBs), as a supplement of lithium-ion batteries (LIBs), are attracting intensive research interest due to their low cost and abundance. Molybdenum disulfide (MoS Moreover, the conductivity of MoS2 material is poor, which severely limits the application at a high current density [2 are usually compounded with different dimensions of carbon materials in order to solve these problems. For example, Choi [2/graphene nanocomposites delivered a discharge capacity and charge specific capacity of 797 and 573 mAh/g, respectively. Most of the preparation operations of the MoS2/carbon composites are complex, such as the use of special high-temperature pyrolysis devices and the removal of templates [As rising stars proceding lithium-ion batteries (LIBs), sodium-ion batteries (SIBs) have the advantages of being low in cost, rich in resources, and environmentally-friendly, which play an increasingly important role in the development of power supply, portable electronic equipment, and national defense technology ,2,3,4. I72 mAh/g . However72 mAh/g ,14. Therfriendly ,16,17. Tfriendly ,19. Althch ~300% . At the  density ,22,23,24le, Choi  preparedemplates ,27.2 nanoflower composites by a mild solvothermal method 6Mo7O24.4H2O  and 0.93 g thiourea CS(NH2)2  were dissolved in 70 mL deionized water and stirred for 1 h, then the mixed solution was transferred to 100 mL hydrothermal reactor and reacted at 200 °C for 24 h. Then, the product was filtered in vacuum and washed with deionized water and absolute ethanol for several times, and then placed in a vacuum oven and dry for 12 h to obtain the MoS2 nanoflower particles.1.07 g ammonium molybdate  was added for another 10 min ultrasonic treatment, then the sealed reaction flask was place in the mild solvothermal of 80 °C for 8 h to obtain the monolithic hydrogel. The hydrogel was dialyzed with deionized water for 2 days, then freeze-dried to obtain the MoS2/3DG composite, placed in a tube furnace under the Ar atmosphere, and thermal treated at 400 °C for 2 h.The graphite oxide powder  was dispersed in deionized water (concentration of 4 mg/mL), and the graphene oxide suspension was obtained by ultrasonic peeling for 1 h. Then 20 mL of the graphene oxide suspension was transferred into the reactor flask with 160 mg of pre-prepared MoS2/3DG was analyzed using the Raman tester . The X-ray photoelectron spectroscopy (XPS) was tested on Nexsa of UK. A thermogravimetric analyzer  was used to character the composition ratio of MoS2 and graphene in the MoS2/3DG composite. The micro-morphology of samples was characterized by the Quanta 200F scanning electron microscope (SEM) and Tecnai G2 transmission electron microscope (TEM) of Philips-FEI company . The MoS2/3DG composite powder, carbon black (conductive agent) and PVDF (binder) were mixed with a mass ratio of 8:1:1, then painted on a pre-treated Cu foil current collector and dried in a vacuum oven at 100 °C for 12 h. The electrode sheets were assembled into the CR2032 button cell in the glove box. The sodium sheet  was used as the counter electrode, the glass fiber  porous membrane was used as the battery separator, and 1 M sodium perchlorate (NaClO4) dissolved the mixture of ethylene carbonate (EC) and dimethyl monocarbonate (DMC) with a volume ratio of 1:1 was used as the electrolyte. Electrochemical impedance spectroscopy (EIS) analysis within a frequency range of 100 kHz to 10 mHz with a voltage amplitude of 5 mV and cyclic voltammetry (CV) at different scan rates were performed through VMP3 , and the cycle stability and rate performance were characterized through the Wuhan Land battery test system.The XRD test was carried out on a diffractometer manufactured of Hitachi Company  with Cu-K α as the radiation source, the scanning speed was 10°/min with a step of 0.01°, the scan range is 5° to 80°. The composition of MoS2 nanoflower particle and MoS2/3DG composite were studied by XRD. As shown in 2 almost focus on 2θ = 14.2° (002), 33.6° (100), 39.8° (103), and 59.3° (110), corresponding to the hexagonal crystal structure (JCPDS 37-1492) [2 particles has been obtained by hydrothermal reaction in the first step. Meanwhile, the diffraction peak positions of the MoS2/3DG composite corresponds well with pure MoS2 particles, indicating that the crystalline structure of MoS2/3DG was not changed by three-dimensional graphene. It is worth noting that a less obvious peak at 25° can be observed, which is attributed to the reduced graphene oxide, for the initial graphene oxide was finally transferred to reduced graphene oxide during the solvothermal and thermal reduction treatment of the second step. The Raman spectra of pure MoS2 nanoflower particle, and MoS2/3DG composites are shown in 2 can be observed in the low wavenumber band (300–500 cm−1), which are attributed to the vibration of the atoms in the hexagonal MoS2 [−1) of MoS2/3DG composites, representing the D-band and G-band of graphene, respectively [2/3DG; meanwhile, the inset high-resolution XPS spectrum reveals two strong peaks at around 232.8 and 229.6 eV, which can be attributed to the Mo3d3/2 and Mo 3d5/2 orbitals of Mo4+ of the MoS2, respectively, while the peak appearing at 226.9 eV is attributed to the S 2s orbital [2 and MoS2/3DG composites in air, the mass loss of pure MoS2 is due to the evaporation of water molecules during the temperature rise from 25 to 300 °C, the mass loss at 300 to 450 °C is due to the conversion of MoS2 to MoO3 [2/3DG composites from room temperature to 300 °C is also due to the evaporation of water molecules and the decomposition of residual oxygen-containing functional groups in graphene. The mass loss of MoS2/3DG composites at 300 to 600 °C is due to the decomposition of graphene and the conversion of MoS2 to MoO3, the residual mass fraction was 68 wt % at 650 °C. Hence, the mass of graphene and MoS2 in MoS2/3DG composites are 22 and 78 wt %, respectively.The crystalline properties of pure MoS37-1492) . There anal MoS2 . There aectively . The XPS orbital . Figure  to MoO3 , and the2 nanoflower particles with the pure MoS2 nanoflower particles composed of stacked nanosheets and closely clustered together. A columnar aerogel of the MoS2/3DG composite is shown in the inset of 2 nanoflower particles are dispersed in the graphene sheet as shown in 2 nanoflower particles and reduce the agglomeration. 2/3DG composite, graphene sheets are tightly wrapped on the surface of MoS2 nanoflower particles, which is benefit to improve the conductivity of the MoS2-based composite and alleviate the volume expansion effect of MoS2 material during the charge and discharge process. At the same time, the three-dimensional porous structure of the composite will be conducive to the penetration of electrolyte, providing the path for sodium ions and electrons transmission. The further analysis of EDS dot-mapping test reveals the well distribution of Mo, S, and C elements in the MoS2/3DG composite as shown in +. As shown in + into MoS2 (MoS2 + xNa+ + xe− = NaxMoS2 (x < 2)), while the reduction peak at 0.65 V is attributed to the formation of Mo and Na2S (NaxMoS2 + (4 − x)Na+ + (4 − x)e− = 2Na2S + Mo) [2/3DG composite, the peak positions of the CV curves are similar to that of pure MoS2, and the coincidence of CV curves at different scan rates are better than that of pure MoS2, indicating the good cyclic stability of MoS2/3DG composite. 2 and MoS2/3DG composites are about 0.8 V and below 0.5 V and the charging platforms are about 1.7 V, which consistent well with CV results. The MoS2/3DG delivers a discharge capacity of 961 mAh/g and a charge capacity of 538 mAh/g for the first cycle, while the discharge and charge capacity of pure MoS2 are 541 mAh/g and 428 mAh/g, respectively. After 50 cycles, the specific discharge and charge capacity of MoS2/3DG composite are 455 and 444 mAh/g, while pure MoS2 deliver the specific discharge and charge capacity of 262 mAh/g and 258 mAh/g, respectively. The excellent electrochemical performance of MoS2/3DG composite can be attributed to its unique microstructure, as the graphene sheets lap together to construct a three-dimensional porous conductive network structure, and MoS2 nanoflower particles are well-dispersed in graphene sheet layers. What is more, the excellent mechanical properties of graphene can effectively alleviate the volume expansion effect of MoS2 in the process of charge and discharge, hence showing much better cycling performance.The cyclic voltammetry (CV) curves and galvanostatic charge and discharge curves were further tested at the voltage range from 3.0 to 0.01 V vs. Na/Na2S + Mo) . Figure 2/3DG and pure MoS2 electrodes with a current density of 100 mA/g was evident in 2 electrode decreases continuously after 20 cycles, with the discharge specific capacity decreases from 541 to 133 mAh/g within 100 cycles. Meanwhile, the specific discharge capacity of MoS2/3DG keeps at 420 and 363 mAh/g after 10 cycles and 100 cycles. 2/3DG and pure MoS2 electrodes at different current densities. For the MoS2/3DG composite electrode, the revisable specific capacity gently decreases from about 420 to 380 mAh/g, 360, 340, and 310 mAh/g with the current density gradually changing from 100 to 200 mA/g, 500 mA/g, 1A/g, 2A/g, and the reversible specific capacity of the composite rapidly restoring to 420 mAh/g when the current density is restored to 100 mA/g, showing extremely excellent rate of 73.8%. The reversible specific capacity of pure MoS2 is 370, 350, 300, 250, and 200 mAh/g with the current density of 100, 200, 500 mA/g, 1 A/g, and 2 A/g, respectively. The reversible specific capacity of pure MoS2 decreases much obviously with the current density returns to 100 mA/g, for the structure of pure MoS2 is destroyed and powdered due to the huge volume expansion effect during the charging and discharging process, while the graphene in MoS2/3DG composite can effectively relieve the volume tension of MoS2 and improve the structure stability.The long-term cycling performance of MoS2 and MoS2/3DG composites after 20 cycles of charge and discharge were tested to further reveal their outstanding rate performance. The analysis of impedance spectra is based on the equivalent circuits shown in 2 is 120.5 Ω, while MoS2/3DG composite only delivers a total impedance of 74.5 Ω, which is much smaller than that of pure MoS2. This indicates that the insertion and desorption of sodium ions in the MoS2/3DG composites are much easier, and the redox reactions of charge and discharge can be carried out faster, which can be attributed to the fact that graphene endows the composite with higher conductivity, hence the MoS2/3DG composite presents much better electrochemical performance.The impedance spectra a of pure2 composite MoS2/3DG was prepared by mild solvothermal reaction. At the current density of 100 mA/g, the reversible discharge specific capacity of MoS2/3DG composite was 420 mAh/g. The electrochemical performance of MoS2/3DG composite was greatly improved compared with pure MoS2 for its unique three-dimensional porous structure which helped it to permeate the electrolyte, at the same time, the abundant pore structure also helped to shorten the transmission path of sodium ions in the material and improve the transmission rate. In addition, the graphene sheet wrapped on the pure MoS2 nanoflower also improved the conductivity. Hence, the capacity of MoS2/3DG shows a good rate of 73.8% with the current density increased from 100 mA/g to 2.0 A/g, and the capacity keeps well with original capacity when the current density returned to 100 mA/g. This work provides a new idea for the preparation of high-performance graphene-based electrode materials for sodium ion batteries.In this work, three-dimensional porous graphene supported MoS


Long interspersed element 1  is a retrotransposon that constitutes ~ 17% of the human genome. Although ~ 6000 full-length L1s spread throughout the human genome, their biological significance remains undetermined. The L1 5′ untranslated region has bidirectional promoter activity with a sense promoter driving L1 mRNA production and an antisense promoter (ASP) driving the production of L1-gene chimeric RNAs. Here, we stimulated L1 ASP activity using CRISPR-Cas9 technology to evaluate its biological impacts. Activation of the L1 ASP upregulated the expression of L1 ASP-driven ORF0 and enhanced cell growth. Furthermore, the exogenous expression of ORF0 also enhanced cell growth. These results indicate that activation of L1 ASP activity fuels cell growth at least through ORF0 expression. To our knowledge, this is the first report demonstrating the role of the L1 ASP in a biological context. Considering that L1 sequences are desilenced in various tumor cells, our results indicate that activation of the L1 ASP may be a cause of tumor growth; therefore, interfering with L1 ASP activity may be a potential strategy to suppress the growth. Full-length L1 consists of a 5′ untranslated region (UTR), two open reading frames (ORFs) that encode the proteins ORF1p and ORF2p, and a 3′ UTR with a polyadenylation signal. Although most L1s in the genome are truncated and are therefore have no retrotransposition activity3, some intact L1s are still capable of retrotransposing in the genome. During retrotransposition, L1 RNA is transcribed from the L1 5′ UTR promoter, which is followed by reverse transcription of the L1 mRNA and insertion of the L1 cDNA sequences into the genome5. L1 retrotransposition requires both ORF1p, an RNA-binding protein with nucleic acid chaperone activity6, and ORF2p, a protein with endonuclease and reverse transcriptase (RT) activity3. Approximately 6000 full-length L1s spread throughout the human genome7; however, the biological significance of these widespread L1 sequences remains undemonstrated.The human genome contains many transposable element-derived sequences, such as endogenous retroviruses and long interspersed element 1 . L1 is one of the major classes of retrotransposons, and it constitutes ~ 17% of the human genome10. Some of these L1-gene chimeric transcripts have been specifically detected in breast and cancer specimens11, suggesting a possible role of the L1 ASP in cancer development. Recent advances in transcriptional profiling revealed genome-wide characterization of these L1 ASP-driven L1-gene chimeric transcripts10. By gene ontology (GO) analysis, Criscione and colleagues showed that these transcripts are involved in diverse cellular processes, including vesicle-mediated transport, intracellular protein transport, mitosis, morphogenesis, and protein modifications10. The human-specific L1HS subfamily and the primate-specific L1PA2-8 subfamily contain an ASP-driven ORF named ORF0, which does not share any extensive homology with known genes12. ORF0-encoded protein (ORF0p) localizes in close proximity to promyelocytic leukemia protein nuclear bodies (PML-NBs) in the nucleus and stimulates L1 retrotransposition12. Since ORF0-proximal exon fusion products are detected in the cells12, ORF0-fusion proteins may change their original protein properties and impact some biological processes. Although L1 ASP-driven transcripts, including ORF0 transcripts, are believed to play a role in gene regulation and/or cell signaling, most studies of the L1 ASP transcripts have been gene expression profiling, so this hypothesis currently remains unevaluated by wet experiments.The L1 5′ UTR contains both sense promoter activity, which drives L1 mRNA transcription, and antisense promoter (ASP) activity, which generates L1-gene chimeric transcripts that include neighboring exon sequencesHere, we developed a novel method to stimulate L1 ASP activity using CRISPR-Cas9 technology and demonstrated that L1 ASP activation enhanced cell growth. We further found that the overexpression of L1 ASP-driven ORF0 also enhanced cell growth. Taken together, L1 ASP stimulation enhances cell growth at least through ORF0 expression. This is the first report demonstrating the role of L1 ASP activity in a biological context.RP , which was then transiently expressed in 293T cells together with sgRNAs targeting the ASP of L1lls Fig. 14, furth10. GO analysis of these genes  were cultured in Dulbecco’s modified Eagle’s medium  supplemented with 5%, 10%, and 5% fetal bovine serum (FBS), respectively.293T cells (a human embryonic kidney cell line from ATCC), HeLa cells , and OL cells 23. The reporter plasmid for L1 ASP activity, pGLuc-L1-ASP, was generated by subcloning the ASP of L1, which belongs to the L1HS subfamily, from the pYX014 plasmid into a pGLuc-Basic plasmid . A parental pGLuc-Basic plasmid was used as a mock reporter. For plasmids expressing sgRNAs against L1 ASP (sgL1ASP #1–#3), a pair of oligos  were annealed and inserted into the AflII sites of a gRNA_Cloning Vector (Addgene #41824). These oligos were designed using the CRISPRdirect server (http://crispr.dbcls.jp/) and the L1 ASP sequence as a query. The L1 ASP sequence was divided into three parts  and a designed guide RNA was chosen from each part. A parental gRNA_Cloning Vector was used as a plasmid expressing mock sgRNA. The expression plasmid of dCas9-VP64, pcDNA-Cas9m4-VP64, was provided through Addgene (#47319). For the ORF0 expression, One-Strep-FLAG-tagged ORF0 was cloned into the pCAG plasmid similarly to the previous study13. The pCAG plasmid expressing only the One-Strep-FLAG tag was used as a mock vector.The luciferase-based L1 retrotransposition reporter plasmid, pYX014, was kindly provided by Dr. Wenfeng An 23 with some modifications. Briefly, 293T cells were cotransfected with pGLuc-5′-UTR, pCMV-CLuc (New England BioLabs), and pcDNA-Cas9m4-VP64, together with a plasmid expressing mock sgRNA or sgL1ASP using Lipofectamine 2000 reagent . At 2 days after transfection, the Gaussia and Cypridina luciferase activities were measured using Gaussia or Cypridina Luciferase Assay Kits (New England BioLabs) according to the manufacturer’s instructions. The Gaussia luciferase (GLuc) activity was normalized to the corresponding Cypridina luciferase (CLuc) activity.L1 promoter and ASP assays were conducted as described23. Briefly, 293T cells were transfected with pYX014 and pcDNA-Cas9m4-VP64, together with a plasmid expressing mock sgRNA or sgL1ASP using Lipofectamine 2000 reagent. At 4 days after transfection, Firefly and Renilla luciferase activities were measured using a Dual-Luciferase Reporter Assay System  according to the manufacturer’s instructions; readings were collected with a single-well luminometer . Renilla luciferase was constitutively expressed from the reporter construct and used for normalization of transfection efficiency.The L1 retrotransposition assay was conducted as previously described31. Total RNA was extracted from 293T cells expressing dCas9-VP64 together with mock, sgL1ASP #1, or sgL1ASP #3 using an miRNeasy kit . Library preparation was performed using a TruSeq stranded mRNA sample prep kit  according to the manufacturer’s instructions. Sequencing was performed on an Illumina HiSeq 2500 platform in 75-base single-end mode. Sequenced reads were mapped to the human reference genome sequences (hg19) using TopHat v2.0.13 in combination with Bowtie2 ver. 2.2.3 and SAMtools ver. 0.1.19. The fragments per kilobase of exon per million mapped fragments (FPKMs) were calculated using Cuffdiff version 2.2.1 with a strand specific mode (-library-type fr-firststrand). A total of 230 genes were upregulated by > 1.5-fold (SgL1ASP #1 to #3 and sgL1ASP #3 to mock) (see Table https://david.ncifcrf.gov/) and the significance was evaluated by the Bonferroni adjusted p-value. Access to raw data concerning this study was submitted under Gene Expression Omnibus (GEO) accession number GSE152634.RNA-seq analysis was conducted as previously described23. Total RNA was extracted from the indicated cells using TRI Reagent  and was reverse transcribed using a Verso cDNA Synthesis Kit . Quantitative real-time RT-PCR assays were carried out using Fast SYBR Green Master Mix (Thermo Fisher Scientific) and gene-specific primers with a QuantStudio 6 (Thermo Fisher Scientific). HPRT1 mRNA was quantified and used to standardize the total amount of cDNA. For real-time RT-PCR specific for ASP-driven ORF0 transcripts, cDNA was synthesized by reverse transcription using the ORF0-reverse primer (OU306) and RT-PCR assays were conducted using the ASP-forward (OU89) and ASP-ORF0-reverse primers (OU703). The gene-specific primers used in this study are as follows:GLuc-forward primer, 5′-AGA GAT GGA AGC CAA TGC CC-3′,GLuc-reverse primer, 5′-CAG ATC GAC CTG TGC GAT GA-3′,CLuc-forward primer, 5′-CTG TGA TCT GAC CCC CAA CC-3′,CLuc-reverse primer, 5′-CTG TTG TCC CCT CAG GCA AT-3′,OU306 primer, 5′-CAC GGA TCC TCA AAG AAA GGG GTG ACG GAC G-3′,OU89 primer, 5′-GTG GAA TTC CTG CAG AGG TTA CTG CTG TC-3′,OU703 primer, 5′-GGG GGA GGG GCG CCC GCC AT-3′,CLASP1-forward primer, 5′-AAG CAA TAC GAT TGG CCG GA-3′,CLASP1-reverse primer, 5′-GTG GGG ATG GAG TTA GGC TG-3′,MAPK12-forward primer, 5′-ACC CTG GAT GAC TTC ACG GA-3′,MAPK12-reverse primer, 5′-GGC AGC GTG GAT ATA CCT CAG-3′,CHFR-forward primer, 5′-CTC CTC CGC TCT CGT GTT G-3′,CHFR-reverse primer, 5′-GTT GTG GCT TCC CAG CAT TG-3′,32, 5′-GGA CTA ATT ATG GAC AGG ACT G-3′, and,HPRT1-forward primer (HPSF-F)32, 5′-GCT CTT CAG TCT GAT AAA ATC TAC-3’.HPRT1-reverse primer (HPSF-R)Real-time RT-PCR was performed as previously described4 cells/well) were replated into a 24-well plate. For the ORF0 expression experiment, 293T cells (4 × 104 cells/well) were seeded in a 24-well plate. The cell number was manually counted every day and the cell viability was evaluated by trypan blue staining. The cell proliferation was also evaluated using the CellTiter-Glo Luminescent Cell Viability Assay (Promega) according to the manufacturer’s instructions. The values of the sgL1ASP- or ORF0-expressing wells were normalized to those of the mock-expressing wells.293 T cells were transfected with plasmids expressing sgL1ASP and dCas9-VP64 or that expressing ORF0 using Lipofectamine 2000 reagent. At 24 h after transfection, the cells stimulated by CRISPR-Cas9-based system  according to the manufacturer’s instructions. The stained cells were analyzed with a CytoFLEX  and Kaluza analysis software (Beckman Coulter).33 with some modifications. Briefly, the cell homogenate was subjected to SDS-PAGE and transferred onto polyvinylidene difluoride membrane . The membrane was then blotted with a mouse anti-FLAG (Thermo Fisher Scientific) or a mouse anti-tubulin  antibody. After three washes, a horseradish peroxidase-conjugated secondary antibody was applied. The bound antibodies were detected using a Clarity Western ECL Substrate (Bio-Rad Laboratories).Western blot was performed as previously describedt-test with a threshold of p < 0.05. Other statistical assessment was used, when indicated.Statistical significance was assessed using a two-tailed Student’s Supplementary Figures.Supplementary Table S1.


Hylotelephium spectabile , a widespread ornamental plant was found infested by larvae of Cacoecimorpha pronubana  a highly polyphagous pest of a wide range of crop, as well as ornamental plants. To the best of our knowledge, this finding suggests that H. spectabile should be considered a new host plant for C. pronubana. Moreover, molecular characterization of the Greek pest population place it along with European species clade.Cacoecimorpha pronubana (Hübner)  is a highly polyphagous pest of a wide range of crop and ornamental plants. It is of Mediterranean origin and widespread in European and Mediterranean Plant Protection Organization (EPPO) region. For the first time, infestations of Hylotelephium spectabile (Boreau) Ohba (syn.: Sedum spectabile Boreau)  ornamental plants by C. pronubana larvae, in private gardens in urban area of Drama, Greece, were found. Species identification was conducted based on morphology of female genitalia. In addition, due to reports on occurrence of cryptic C. pronubana species within Europe, DNA barcoding was carried out to determine the molecular status of the pest. This communication reports a new host of C. pronubana and places the Greek pest population along with European species clade. Cacoecimorpha pronubana (Hübner)  is a highly polyphagous and locally important pest of a wide range of crop plants. In addition, C. pronubana infests also ornamental plants  Ohba (syn.: Sedum spectabile Boreau)  is an ornamental plant species that is widely distributed throughout the warm and temperate zones, originating from Korea and Northern China [Hylotelephium spectabile has sparsely toothed leaf margins and flower stamens longer than petals, which differentiates it from other species in the genus [rn China . It is ahe genus .C. pronubana [In this study, we report a carnation tortrix infestation on butterfly stonecrop plants. Based on an extensive literature review, to the best of our knowledge no Crassulaceae member-plants have been previously reported as host of this species. Our results support that the particular plant species can be considered a new host of ronubana ,15.In mid-August 2020, caterpillars of an unknown species have been detected feeding on potted butterfly stonecrop plants located in private gardens in the broader area of Drama . Infestation had caused considerable damage on foliage as well as on blossoms, reducing the ornamental/aesthetical value of the butterfly stonecrop plants. Leaves of the infested plants were chewed and knitted together with silken webbing, in which feeding caterpillars inhabited a. This oC. pronubana [C. pronubana species within Europe [® Genomic DNA kit  and following the manufacturers’ instructions modified during the tissue grinding process [® PCR Purification Kit (Invitrogen) following the manufacturer’s protocol and purified products were sequenced in the automated sequencer ABI3730XL of CeMIA Company , using the same primers as in PCR. Finally, sequences were initially visualized with Chromas Lite® version 2.01  and after removing the ambiguous nucleotides at both ends of the sequences, a final product of 443 bp long was obtained for each individual. These sequences were then blasted in the NCBI GenBank database (https://www.ncbi.nlm.nih.gov/genbank/ (accessed on 13 March 2021)), placing them among the sequences of C. pronubana deposited there. In particular, both these sequences were 100% identical with C. pronubana from France (Accession Number: KX041548) and Italy (Accession Number LC031966), and 99.77% similar with the sequences from the United Kingdom (Accession Number: KX043977), Germany (Accession Number: KX041034) and Italy (Accession Number LC031971). Given the distinct separation that has been previously revealed among the populations of C. pronubana worldwide [C. pronubana sequences, deposited both in NCBI (see Accession Numbers above) and BOLD (BOLD BINs: AAD3477 and AAL5782) into an integrated large dataset and constructed a Neighbor-Joining phylogenetic tree. To do that, we used MEGA6 [C. pronubana sequences, remaining distinctly separated from those of Spain . The external morphological characters of eggs, larvae, pupae and the emerged adults resembled those of ronubana ,11,16,17n Europe , DNA was process . Polymer process . PCR pro process . The purorldwide , we compof Spain . Interes molesta ). Nevert molesta , this fi molesta .Cacoecimorpha pronubana adults are nocturnal and males are much more active than females with a flight that is erratic and agitated [Cacoecimorpha Obraztsov 1954 is a monotypic tortricid genus with adult females having a wingspan of 18–22 mm, with pale orange brown forewings that are reticulated with darker brown coloration, while hindwings are mainly orange. Adult males are generally smaller (wingspan of 12–17 mm) and their forewings are orange brown with reddish-brown and purplish-black markings while their hindwings are bright orange with blackish border  [C. pronubana in Europe.Hereby, a new host plant was detected for . (2020) , placing