Idiopathic pulmonary fibrosis (IPF) is a lethal fibrotic lung disease characterized by aberrant remodeling of the lung parenchyma with extensive changes to the phenotypes of all lung resident cells. The introduction of transcriptomics, genome scale profiling of thousands of RNA transcripts, caused a significant inversion in IPF research. Instead of generating hypotheses based on animal models of disease, or biological plausibility, with limited validation in humans, investigators were able to generate hypotheses based on unbiased molecular analysis of human samples and then use animal models of disease to test their hypotheses. In this review, we describe the insights made from transcriptomic analysis of human IPF samples. We describe how transcriptomic studies led to identification of novel genes and pathways involved in the human IPF lung such as: matrix metalloproteinases, WNT pathway, epithelial genes, role of microRNAs among others, as well as conceptual insights such as the involvement of developmental pathways and deep shifts in epithelial and fibroblast phenotypes. The impact of lung and transcriptomic studies on disease classification, endotype discovery, and reproducible biomarkers is also described in detail. Despite these impressive achievements, the impact of transcriptomic studies has been limited because they analyzed bulk tissue and did not address the cellular and spatial heterogeneity of the IPF lung. We discuss new emerging technologies and applications, such as single-cell RNAseq and microenvironment analysis that may address cellular and spatial heterogeneity. We end by making the point that most current tissue collections and resources are not amenable to analysis using the novel technologies. To take advantage of the new opportunities, we need new efforts of sample collections, this time focused on access to all the microenvironments and cells in the IPF lung. Our understanding of idiopathic pulmonary fibrosis (IPF), a chronically progressive scarring lung disease, with a significant genetic component, has dramatically changed in the last two decades. This has happened because after years of formulating hypotheses based on animal models, or analogies from other diseases, pulmonary researchers shifted their focus to analyzing the human lung. The increased availability of well-characterized human tissues and the emergence of high throughput transcriptomic profiling technologies facilitated a new era in IPF research, one in which novel hypotheses are based on observations from human lungs. The sheer size of the data, and its unbiased nature, reintroduced serendipity in pulmonary fibrosis research, and thus led to numerous, previously unexpected observations, novel hypotheses and paradigm shifts. In this perspective, we provide an overview of the impact of transcriptomics on our understanding of IPF. We highlight the timeline of major discoveries Figure  with a fp-value , and reduced activity of matrix metalloproteases (MMP) . The firWhile their exact roles have not been fully elucidated, the initial unexpected observation that MMPs are increased in the IPF lung, has been validated numerous times. It is now well accepted that MMPs affect numerous signaling pathways that together contribute to the profibrotic environment in the IPF lung and may also serve as effective biomarkers (see below).Transcriptomic analysis of bulk tissue depends on follow-up analyses to decipher the cellular origin of differentially expressed genes. One of the most surprising findings in IPF transcriptomics was that cellular origin of large number of genes that distinguish the IPF lung from controls ended up being the alveolar epithelium , 60 Fig; Table 1Other genes increased in IPF and unexpectedly localized to the alveolar epithelium adjacent to fibrotic regions include N-cadherin , HIF-1-aOf particular interest in this context, is a study that demonstrated that IPF patients with increased expression of cilia genes exhibited also increased MMP7 and MUC5B, as well as microscopic honeycombing but not myofibroblast foci on histological examination, suggesting that they represented a distinct IPF endophenotype  , gene expression patterns that correlate with disease clinical attributes. The studies are summarized in Table An early suggestion that lung gene expression can be used to classify disease emerged from comparison of lungs of patients with IPF from those with fibrotic hypersensitivity pneumonitis (HP) using transcriptome analysis . The enrAnalysis of lung samples obtained from patients with sporadic IPF, familial pulmonary fibrosis with a usual interstitial pneumonia (UIP) pattern, and non-specific interstitial pneumonitis (NSIP) revealed similarities on gene expression patterns and pathways and a minimal difference between IPF and NSIP Table  5, 8). , 8. 5, 8A recent study used supervised machine learning algorithms to distinguish lung biopsy samples with UIP from non-UIP  identified a 22 gene signature . This approach was solely based on transcriptional data concordant with UIP pathological findings without integration of clinical information, or comparison to patient-level diagnoses by multidisciplinary teams, the current diagnostic gold standard . The samIdiopathic pulmonary fibrosis has different patterns of progression, from stable disease lasting for long periods of time to rapid progression, and acute exacerbations that are highly lethal. Despite a very small number of samples, differentially expressed genes were found in end-stage lungs obtained from patients with rapid and slow progression defined by length of symptoms Figures ; Table 3The study of acute exacerbations of IPF has been limited, because of lack of tissue availability. Using a unique resource of rapid lung autopsies  investigA more sophisticated effort to identify disease endotypes based on tissue gene expression, incorporated clinical and histological information in the analysis . This deThe availability of large datasets such as the LGRC, allowed also analysis of multiple chronic lung disease in parallel. Recently, applying a novel computational approach named integrative phenotyping framework, investigators discovered novel endotypes of chronic obstructive pulmonary disease (COPD) and IPF . They inGenome scale transcriptome studies facilitated the development of protein-based biomarkers for IPF diagnosis Figures ; Table 4Indeed, MMP7, which emerged out of the first microarray analysis of human IPF lungs, was replicated as predictive of increased mortality in multiple cohorts of IPF patients –92, 95. The transcriptome of the peripheral blood is highly appealing because of information about disease presence and outcome. It represents a safe and accessible alternative to availability of samples from the lung. Microarray gene expression profiles of whole blood RNA  distinguA subsequent study aimed to identify PBMC gene expression profiles predictive of increased mortality in patients with IPF . The autUntil recently considered the dark matter of the genome, the significant role of non-coding RNAs in human health and disease is increasingly appreciated . We willMicroRNAs are small non-coding RNAs that regulate gene expression by either initiating RNA degradation or inhibiting translation through binding to the 3′ UTR of their target gene. Acting as rheostats, many microRNAs regulate the general responsiveness of a cell to a certain stimulus by affecting numerous genes and frequently serving as gate keepers of feed forward loops. The expression of approximately 10% of the microRNAs is different in IPF compared with control lungs , 99. Thein vivo and in vitro , as well as a patent on novel biomarkers in IPF (not licensed). NK consulted Biogen Idec, Boehringer Ingelheim, Numedii, MMI, Pliant, Third Rock, and Samumed. NK has an ongoing collaboration with MiRagen but no fund exchange. The other author declares that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.


Planet Earth is by far a microbial world, and microbial contribution is evidenced by the significant biomass and activities displayed by microbial biofilm communities in all habitats.npj Biofilms and Microbiomes will feature new and exciting findings on biofilms and microbial communities: how member organisms interact at various levels; how they influence their respective environments; and the effects of external perturbations on structure function relationships.Modern biology developed through studying relatively simple model systems, and major findings on key traits such as those of environmental stress responses were accomplished using mono species cultures of common laboratory bacteria. The full extent to which unicellular organisms exhibit multicellular behaviour, initially for populations and more recently for communities, was not realised until several decades later. For example, we now understand that microorganisms use various extracellular signalling mechanisms to establish communities where their collective activities are substantially different from the summation of the activities of individual members. Moreover, it was recently demonstrated that communication mechanisms using electron shuttling are employed by several bacterial species. Individual members of a population, however, can evolve new variants as a response to stress, which may impart benefits for the entire population. The emerging concept of bacteria as social organisms is a disciplinary milestone and reflects the convergence of microbiology and related life sciences, as well as ecology and microbial ecology. In addition, understanding the mechanisms of microbial population and community behaviour, using quantitative systems analysis and real-time imaging technologies, sheds light on complex social–microbial interactions and contributes to both ecological theory and experimental ecology.The largely hidden powers of microbial biofilms hold much promise for applications across a broad scope of systems. These include the control of disease-causing microorganisms and the detection and management of pathogens in the environment; engineering microbial biofilms for effective bioremediation of pollutants; and the discovery of biofilm-generated novel bioactives. While the realisation of several biofilm applications is imminent, others may be reality only after robust tools for reproducible mixed-species microbial biofilm community analyses have been developed. Notably, to realise these opportunities, cross-disciplinary collaboration and systems need to be embraced, ranging from hydrodynamics at microscales to computational systems biology at the ecosystem level.Advancing the fields of microbial community and biofilm biology calls for a change in how microbiology is perceived, practiced and studied. Biofilms are not just accidental mixtures of different types of microorganisms, rather their highly coordinated three dimensional architecture is built to accommodate unique and specific communal interactions. These interactions are often diffusion-limited and dictated by the relative spatial distribution of member organisms. Moreover, the polymeric matrix in which biofilm cells are embedded is a remarkably complex and active component of all biofilms, but remains relatively poorly understood. This matrix holds together microbial populations and communities in polymeric structural scaffolds that contribute to unique biofilm properties, such as stress resilience and long-distance communication mediated by directed electron shuttling. An array of such specific traits and properties is predicted to emerge from the increased efforts of contemporary biofilm research into this integral biofilm component.9 years. They have since evolved along with, and closely associated to, the more recent evolution of higher organisms, establishing a strong interdependence on a single holobiont. We can now explore the biology of host multicellular organisms, as well as the ecology of various habitats, by addressing their interactions within their respective microbiomes.The microbiomes of biofilm communities are now rapidly becoming focal points in studies of all ecosystems, from environmental and industrial to medical, driven by the notion that the well-being of these systems depends on the unique contribution of their microbiomes. In fact, microbial biofilms are the oldest known biological communities, dating back more than 3.6×10npj Biofilms and Microbiomes will welcome manuscripts dealing with large-scale analyses of microbial biofilms and communities in various natural or engineered settings, but will prioritise those that provide mechanistic insights, and/or means to make use of or perturb the communities in ways that can be beneficial.Technical advancements such as metagenomic, metatranscriptomic, metaproteomic and metabolomic approaches have provided tools to dissect virtually any natural polymicrobial community and assess responses to a multitude of external and internal perturbations. Many of these communities are highly complex and require sophisticated experimental design, data collection and analysis to attain meaningful results. The analysis of complex interactions in multispecies microbial communities will require increased efforts in computational biology, bioinformatics and overall systems biology approaches. Clostridium difficile infections are treated with particular complex microbial communities from healthy individuals, with promising outcomes. We can expect to see how such interventions will improve with further understanding of the various roles microorganisms have in healthy gut physiology. We are witnessing other exciting developments such as possible links between gut microbial activities and cognitive functions, as well as studies linking gut-derived microbial products and the shaping of the immune system.The human microbiome is a system that has recently attracted significant attention. The resulting wealth of data has established important correlations between the host’s various organs and the associated microbiomes. For example, the gut microbiome composition has been linked to various host conditions such as diabetes, atherosclerosis and autoimmune disease. Indeed, studies are now emerging on the effect of microbial fermentation products from dietary sources on intestinal gluconeogenesis, which subsequently affects the host’s glucose and energy homeostasis, and negative correlations between gut microbe gene richness and the likelihood of acquiring an autoimmune disease. Similarly, differential effects of vaginal versus caesarian birth delivery on shaping the infant’s initial gut microbes have been observed, a finding that is likely to result in directed manipulation of intestinal microbiome. Already today, patients suffering from Beyond the various human microbiomes, we know very little about the myriads of interactions between environmental microbial communities and a spectrum of hosts and surfaces in all ecosystems.npj Biofilms and Microbiomes very much wishes to become the journal for such contributions.Future challenges will also involve a clearer explanation of the emerging host–microbiome interactions in mechanistic terms. npj Biofilms and Microbiomes, I wish this journal to become a platform for cross-disciplinary discussions and scientific exchange that allow for the understanding of the biology and ecology of biofilms, microbial populations and communities, as well as applications so derived across medical, environmental and engineering domains.As the first Editor-in-Chief for


Combining atomically-thin van der Waals materials into heterostructures provides a powerful path towards the creation of designer electronic devices. The interaction strength between neighbouring layers, most easily controlled through their interlayer separation, can have significant influence on the electronic properties of these composite materials. Here, we demonstrate unprecedented control over interlayer interactions by locally modifying the interlayer separation between graphene and boron nitride, which we achieve by applying pressure with a scanning tunnelling microscopy tip. For the special case of aligned or nearly-aligned graphene on boron nitride, the graphene lattice can stretch and compress locally to compensate for the slight lattice mismatch between the two materials. We find that modifying the interlayer separation directly tunes the lattice strain and induces commensurate stacking underneath the tip. Our results motivate future studies tailoring the electronic properties of van der Waals heterostructures by controlling the interlayer separation of the entire device using hydrostatic pressure. Van der Waals heterostructures enable fabrication of materials with engineered functionalities. Here, the authors demonstrate precise control over the interaction between layers by application of pressure with a scanning tunnelling microscopy tip. Previous work has focused on controlling the properties of these systems through the choice and ordering of the materials in the heterostructure, as well as the rotational alignment between layersδ∼1.8%)  is an excellent testbed for this effect, as a long-wavelength periodic interaction emerges when the two crystals are in near-rotational alignment due to their small lattice mismatch 8%) refs , 7, 8. T% refs , refs 1112. Small eHere we demonstrate a path towards achieving control over this degree of freedom by demonstrating that pressure exerted by a scanning tunnelling microscope (STM) tip16171819I as a function of the relative tip-sample separation Δz. The tunnelling current is expected to scale exponentially with Δz asWe first present evidence of the out-of-plane movement of the graphene lattice produced by the tip, depicted schematically in m is the electron mass and ϕ is the tunnel barrier height. This exponential approximation holds well for graphene on SiO2, but fails for graphene on hBN (I(Δz) becomes strongly dependent on the specific tunnelling parameters, with the tunnel current decay growing slower as the tip distance is brought closer to the surface. Furthermore, the decay is initially quadratic rather than exponential. I(Δz) hysteresis between tip approach and tip retraction (2 substrate and is relatively immobile).where e on hBN , indepene on hBN . In the traction . This alz in z−zg(z), where zg(z) represents the movement of the graphene relative to the hBN substrate as a function of the tip position z. We plot the relative movement of the graphene in ϕ=4 eV, as extracted from measurements acquired at large tip-sample separations. The tip initially lifts the graphene away from the hBN as it retracts. After around 2 Å of retraction, the tip is no longer able to continue pulling the graphene, which then begins to slowly relax back towards the hBN substrate, as it is still under the influence of a vdW force from the tipzg>0. The blue and black curves in z, as in To account for the additional out-of-plane movement of the graphene sheet, we substitute Δ2) areas, perform a Fourier transform  as a function of position  and AA (hexagons atop one another) stacking configurations depend on the interlayer separation between the two materials . To underansform , and extposition . Finallyposition .Interestingly, we observe qualitatively similar behaviour in slightly misaligned samples as well. Specifically, we observe the three different strain profiles as a function of tip-sample separation in all moiré areas studied with periods varying from 14 nm  down to about 6 nm . As an example, z from attractive, to repulsive, and to strongly repulsive. In the attractive regime, the graphene under the tip is lifted off the hBN surface, lowering the adhesion potential. The stacking boundaries are then attracted to the scanning tip, and as a result the graphene lattice appears to be expanded along the stacking boundaries  at high pressure, and the stacking boundaries are pushed away  and over tens of pulse cycles per tip. We note that qualitatively similar moiré scale lattice deformations have been observed in graphene on Ir(111) with AFM using a tip intentionally terminated with a carbon monoxide moleculeTungsten tips were prepared by electrochemical etching, and further shaped R around its apex, hovering at height h0 relative to a relaxed reference plane . We assume that the vertical graphene displacement conforms to the tip profile as long as it does not exceed a certain height, hmax, see u(r) of the sample, relative to the relaxed moiré pattern, which we want to determine. We construct a smooth interpolation of the ab-initio adhesion potentials VS(z) between different graphene/hBN stackings, where z is the separation between the two crystals. Using the interpolated potential, we evaluate the total adhesion energy per unit area for a given field u(r). At each r, the value of z is constrained by the tip profile, as described above. To this adhesion energy, we add the corresponding elastic energy associated to u(r). We discretize r, and express the total energy as a function of the finite set of u on the discrete mesh. We minimize the total energy, using conjugate gradient methods, and find the deformation u(r) at equilibrium. We then obtain the dynamical strain as measured by the tip by performing this sample relaxation as the tip moves across the sample at a constant height h0. The model has no unconstrained free parameters, as all can be roughly estimated experimentally.An overview of our theoretical model is as follows .Supplementary Figures 1-13, Supplementary Table 1, Supplementary Notes 1-6 and Supplementary ReferencesMovement of the scanning probe tip across the moiré at 0.8 Angstrom above the equilibrium graphene position.Movement of the scanning probe tip across the moiré at 0.2 Angstrom above the equilibrium graphene position.Movement of the scanning probe tip across the moiré at the equilibrium graphene position.Movement of the scanning probe tip across the moiré at -0.2 Angstrom above the equilibrium graphene position.


Our differential characterization is able to distinguish the thermal conductance of β-W film and β-W/G interface. The cross-plane thermal conductivity (k) of β-W films is determined at 1.69~2.41 Wm−1K−1 which is much smaller than that of α-phase tungsten (174 Wm−1K−1). This small value is consistent with the large electrical resistivity reported for β-W in literatures and in this work. The β-W/β-W and β-W/G interface thermal conductance (GW/W and GW/G) are characterized and compared using multilayered β-W films with and without sandwiched graphene layers. The average GW/W is found to be at 280 MW m−2K−1. GW/G features strong variation from sample to sample, and has a lower-limit of 84 MW m−2K−1, taking into consideration of the uncertainties. This is attributed to possible graphene structure damage and variation during graphene transfer and W sputtering. The difference between GW/G2 and GW/W uncovers the finite thermal resistance induced by the graphene layer. Compared with up-to-date reported graphene interface thermal conductance, the β-W/G interface is at the high end in terms of local energy coupling.We report the thermal conductance induced by few-layered graphene (G) sandwiched between  Although there still remains challenges in wafer-scale deposition and controlling the electronic bandgap, graphene is widely seen as a strong candidate for future microelectronics7. In the applications of graphene, interface thermal resistance (R) or conductance (G) induced by graphene is the most common quantity used to characterize heat dissipation from graphene to its substrate. Early work by Freitage et al.8 characterized the heat dissipation from graphene to substrate for the first time, but didn’t explore graphene-substrate interface thermal resistance. Following work by Chen et al.9 employed the differential 3ω method on the graphene/SiO2 interface and reported a R range from 5.6 × 10−9 to 1.2 × 10−8 Km2W−1 at room temperature. Mak et al.10 employed the ultrafast pump-probe method and obtained a G of the single-layered and multilayered graphene/SiO2 interface varying from 2000 to 11000 Wcm−2K−1. Koh et al.11 performed the time-domain thermoreflectance (TDTR) measurement on the Au/Ti/graphene/SiO2 and Au/Ti/SiO2 sandwiches (graphene layers 1 ≤ n ≤ 10) and reported a G of ~25 MWm−2K−1 at room temperature for the Au/Ti/graphene/SiO2 interface. Similarly, Guzman et al.12 performed the TDTR measurements on the variable metals/single-layered graphene/SiO2 structures and gave a G range of 15-60 MWm−2K−1. Hopkins et al.13 determined the thermal boundary conductance across the Al/single-layered graphene/Si interface from the TDTR measurement and found a two-fold increase with the oxygen functionalization of the graphene.Monolayered or few-layered graphene has attracted remarkable attention over the past several years due to its extremely high electron mobility and thermal conductivity2 interfaces, the metal/graphene/metal interfaces with metal substrates like Cu, Pd and W are lack of attention until recent years16. Huang et al.16 performed the TDTR measurements on the Pd/transferred graphene (trG)/Pd interface and reported a G of 300 MWm−2K−1 at room temperature for the sample with radiofrequency (rf) magnetron sputtering top Pd. This value is seven times larger than that with thermal evaporation top Pd (42 MWm−2K−1). The large enhancement of thermal conductance is attributed to the electronic heat conduction via atomic scale pinholes formed in the graphene during the sputtering process. W, as an excellent radiation tolerance material19, is of great importance for the safety of nuclear reactors. Graphene, because of its impermeability to all standard gases (including He)20, has also been regarded as an effective material in retarding radiation damages during nuclear reaction21. These remarkable properties of W and graphene motivate us to fabricate the W/graphene multilayered system and study the ability of its radiation tolerance and the resulting thermal transport capacity. As a special phase of W, β-W has been extensively investigated due to its high superconducting transition temperature25 since first prepared by Hartmann26 and Fink27. In recent years, investigations of β-W have mainly focused on the nanostructure, phase transformation conditions, residual stress, electrical resistivity and spin Hall effect33. No direct characterization of the thermal transport properties of this unique metastable structure has been reported.Despite the extensive studies of metal/graphene/SiOβ-W films sandwiched with graphene layers by employing the photothermal (PT) technique36. In the PT technique, the sample is heated by an intensity-periodic laser and the thermal radiation signals from the sample surface are detected by an infrared detector. By fitting the phase shift between the radiation signals and the incident laser beam, the cross-plane thermal conductivity (k) of β-W, β-W/β-W interface thermal resistance (RW/W) and β-W/G interface thermal resistance (RW/G) are determined. Furthermore four-probe electrical resistivity study is carried out on the measured samples and compared with the PT measurement results to interpret the experimental observations.In the present work, we systemically study the structure and thermal transport properties of the multilayered 2/Si substrate [see Fig. β-W with the (200) plane parallel to the surface [see Fig. The single-layered W films on 100-nm-SiOsee Fig.  are prepsee Fig. , the W fsee Fig. . Detailsβ-W films as well as the thermal resistance induced by the SiO2 layer and the β-W/SiO2 interface. Given these physical properties, the theoretical phase shift can be calculated. The physical model of the PT technique and the experimental setup are detailed in the Supplementary materials (see S2). In the experimental measurement, we are interested in the phase shift between the thermal radiation and the modulated laser beam. However, the measurement will inevitably include time delay induced by the whole system. This time delay can be eliminated by measuring the phase shift between the reflected laser beam and the irradiating laser beam . The experimental setup for the calibration process has no other difference except for detecting the reflected laser signals instead of the thermal radiation by removing the Ge window. Figure −6 s. The real experimental phase shift (noted as ϕnor) between the thermal radiation and the laser beam can be calculated as ϕnor = ϕraw − ϕcal, with ϕraw being the raw phase shift detected in this experiment.In this work, the PT technique is used to characterize the thermal transport properties of the films. It irradiates the sample surface with an amplitude modulated laser beam, and measures the surface thermal radiation. The laser beam modulation frequency is scanned within a wide range, and the time delay (phase shift) between the thermal radiation and laser beam is measured and used for data fitting. The phase shift is related to the physical properties of the sample such as thermal conductivity, heat capacity and density of the 2 layer and the β-W/SiO2 interface  of the sample, which includes both the thermal resistance induced by the β-W film and β-W film. By studying how Rtotal varies with the W film thickness (L), k of β-W film and The fitting process is operated by using a well-developed program by our lab. By using different trail values of unknown parameters, the theoretical phase shifts are calculated over the specified modulation frequency range. The value that gives the least square deviation between the theoretical phase shifts and the experimental ones is taken as the real property of materials. Here, the thermal resistance induced by the SiOβ-W film of 110 nm thickness. Excellent agreement is observed between the fitting (green line) and experimental data (black circles) at all the modulation frequencies with a fitting residue of 0.30 degree. Rtotal is determined to be 2.02 × 10−7 Km2W−1 for this sample. When Rtotal is taken as 1.93 × 10−7 Km2W−1 and 2.10 × 10−7 Km2W−1, the fitting residue is 0.53 and 0.51, respectively, which are much larger than the experimental uncertainty of the phase shift as shown in Fig. −7 Km2W−1.Figure Rtotal of the samples can be expressed as the following equation:β-W films are prepared using the same substrate and under the same conditions, k of β-W films and Rtotal and L. Figure Rtotal versus L for the single-layered β-W films studied in this work. An obvious linear relation is observed. Based on linear fitting, k of β-W films and −1K−1 and (1.43 ± 0.02) × 10−7 Km2W−1, respectively. In the range of the laser modulation frequency (600 Hz to 20 kHz), the thermal diffusion depth in one period . As there is no record for k of β-W, we will try to explain this low thermal conductivity of the metastable β-W from the Wiedemann-Franz (WF) law assuming that electrons still dominate in thermal conductance of this material.Since all the σ) by the WF law at relatively not-very-low temperatures by the following equation:LLorenz is the Lorenz number and for tungsten LLorenz is 3.04 × 10−8 W·ΩK−2 at 300 K37. σ is calculated at 2.17 × 105 Ω−1m−1 based on our measured k, corresponding to an electrical resistivity of 4.61 × 10−6 Ω·m (or 461 µΩ·cm). This value is comparable with the reported electrical resistivity for this phase of W39. Early work by Petroff et al.38 reported an electrical resistivity of β-W ranges from 150–350 µΩ·cm. Following work by O’Keefe et al.30 reported an even higher electrical resistivity of as-deposited β-W (155–870 µΩ·cm) and found that even after rapid thermal annealing, it can still be as high as 478 µΩ·cm. Recent work by Hao et al.32 measured the electrical resistivity of β-W with different thicknesses and extracted a value of about 195 µΩ·cm. The relatively high electrical resistivity (compared with 5.33 µΩ·cm for the bulk pure α-W32) will be discussed in detail below.Thermal conductivity is related to the electrical conductivity . A current I is fed through the two outer probes and the voltage V is measured between the two inner ones. The sheet resistivity of samples can be expressed as: Rs = F1F2F3V/I, where F1 is the finite shape correction factor, F2 the probe spacing correction factor and F3 the thickness correction factor. F1 can be derived from the existing standard table42. F2 can be calculated by the following equation42:S2 is the spacing between the two inner probes and S the average probes spacing. In our experiment, with equal spacing between adjacent probes, F2 is taken as 1. F3 is also taken as 1 for all the samples, as in our case the thickness (t) of all these films is very small which means t/s << 0.4.As the film is very thin, the cross-plane σ of β-W is calculated as electrical resistivity σ−1 varies from 258 to 316 µΩ·cm which is in the well accepted range (150–350 µΩ·cm) of β-W38. Typically, the relatively higher resistivity in thin metallic films is attributed to the carrier momentum loss along the current flow direction due to surface and grain boundary scattering which is commonly explained by using the Fuchs-Sondhermer (FS) surface scattering model44 and the Mayadas-Shatzkes (MS) grain boundary model46, respectively. However, the finite-size effect and the grain boundary scattering are not sufficient to account for the large resistivity of β-W. According to Hao’s recent work, the electrical resistivity of bulk β-W is calculated at 195 ± 3 µΩ·cm with an effective mean free path (leff) of about only 0.45 nm at 300 K32. This means leff of β-W films may be ten to hundred times smaller than the averaged grain size and the film thickness [see Fig. leff, surface scattering or grain boundary scattering has negligible influence on the resistivity of thin metallic films44. Other scattering mechanism must exist and dominate the electron transport property. One possible explanation is the charge carrier concentration. Recent work by Lee et al.39 obtained the charge carrier concentration of different phase of W by using Hall measurement and found a much low carrier concentration of the β-phase one. Another possible reason is the impurities, where a small amount of oxygen is believed to exist and induce the β-W formation without forming a WxO compound47. Such a dilute bulk of impurities may induce significant electron scattering. This also helps explain the large range of reported resistivity of β-W which may be due to the different concentration of oxygen induced impurities. The local structure may be another reason as β-W is believed to be a mixed phase consisting of ordered and stacking faulted W3W structures38. Considering the relatively larger grain size, local disorders or dislocations of atoms may also play an important role. So far, as there is no clear mechanism for the large resistivity, further work, particularly theoretical study, is needed.In-plane entclass1pt{minimaβ-W films. The β-W films are produced in the vacuum magnetron sputtering system, where the deposited metal films typically show kind of columnar structure. For the tungsten films prepared in this work, the columnar structure can be clearly seen from the SEM images in Fig. σ and k, the validation of WF law is also questionable. It has been widely studied that, for nanocrystalline films, the Lorenz numbers are very different from their corresponding bulk values. For example, Yoneoka et al.48 reported an average Lorenz number of 3.82 × 10−8, 2.79 × 10−8, and 2.99 × 10−8 WΩK−2 for 7.3-, 9.8-, and 12.1-nm Pt films, respectively. Experimental results of Zhang and co-workers50 showed that the Lorenz number of 21–37 nm and 53 nm thick polycrystalline Au films are around 7.0 × 10−8 and 5.0 × 10−8 WΩK−2, respectively. Calculations by Ou et al.51 revealed that the Lorenz number of a 180-nm nickel nanowire is a little higher than the bulk value. Our previous experimental work52 on ultrathin iridium films (0.6–7 nm) shows that the Lorenz number ranges from 5.83 × 10−8 to 7.8 × 10−8 WΩK−2. As there is no literature data for the Lorenz number of β-W nanofilms and how it will deviate from the bulk value is not clear to our knowledge, the use of WF law based on the bulk’s Lorenz number is not suitable in this work. Therefore, considering the unclear anisotropy level of the β-W nanofilms and the deviation of the Lorenz number, we intend to extract cross-plane direction k directly by performing the PT measurements on this group of samples in this section.The discrepancy in the thermal conductivity determined by the PT technique and the four-probe method is mainly attributed to the anisotropy of the β-W films from other resistances. This also holds true for the multilayered β-W samples. Figure Rtotal of each sample in this group is calculated. The inset in Fig. Rtotal versus L of both the single-layered β-W films (black rectangles) and the multilayered β-W films (red circles). We can see that all these points show a linear relationship and that the differences of Rtotal between the single-layered β-W films and the multilayered ones are not significant. This points out that RW/W will be small. It is determined as below.Not like the samples in the first group, the samples in this group were grown for several times, that is, these samples consist of different number of sublayers see Fig. . The SEMRtotal of the multilayered β-W films can be expressed as the following equation:So, we haveR/n versus L/n, these points also show a linear relationship as shown in Fig. k and the intercept is related to RW/W. k determined here (2.05 ± 0.36 Wm−1K−1) is a little higher than (but very close to) that of the single-layered β-W films in the first group (1.98 ± 0.06 Wm−1K−1). Considering the uncertainties of the experiment and data processing, we conclude that the samples from the two groups have negligible structure difference during sample preparation. From this linear fitting, RW/W is determined at (3.57 ± 2.67) × 10−9 Km2W−1. This RW/W investigation serves as a comparison base for the RW/G study that will be described in the next section. At this point, we can conclude that the β-W/β-W interface thermal conductance (GW/W) has an average of about 280 MW m−2K−1.When plotting Δβ-W films has also been measured by using the four-probe method. Although there are some variations in several samples, the electrical conductivity and resistivity of the samples in this group have no significant difference with those of the first group. Considering the large thickness difference between the single-layered films and the multi-layered ones, the previous assumption that the size-effect or surface-scattering play a tiny role in the very high resistivity has also been proved. The electron thermal conductivity in the in-plane direction is also calculated using the WF law, which is larger but still comparable to that in the cross-plane direction. The calculation results are detailed in the Supplementary materials  can be estimated as RG/G = l/kG with l the distance between graphene layers (0.335 nm) and kG the cross-plane thermal conductivity of graphene (5.7 Wm−1K−1)54. Accordingly, RG/G is about 5.9 × 10−11 Km2W−1 which is much smaller than Rtotal. Thus, the intrinsic thermal resistance of the graphene layers can be neglected in this experiment. With k of β-W films taken as 1.69~2.41 Wm−1K−1 and −7 Km2W−1, RW/G of all the samples in this group can be calculated by using the following equation:The samples in this group have the same thicknesses as those in the second group, respectively. Compared with samples in the second group, the difference is that we have graphene layers sandwiched between tungsten sublayers see Fig. . Figure RW/G varies from sample to sample and has a maximum value of 9.67×10−9 Km2W−1. The lower and upper limit uncertainties are also given in Table RW/G of each sample. We can see the uncertainties also vary from sample to sample and has a maximum value of 2.43 × 10−9 Km2W−1. These uncertainties are mainly attributed to the uncertainty of k of β-W films, which varies from 1.69 to 2.41 Wm−1K−1.The calculation results are shown in Table β-W layers). In Table R2W/G represents thermal resistance of two W/G interfaces induced by each graphene layer. Also shown in Table RTW/G) induced by all β-W/G interfaces in one sample. Most of the RTW/G values are much larger than the experimental uncertainty (8~9 × 10−9 Km2W−1). This means the thermal resistance induced by the graphene layers is detectable in our experiment. It is clear that most R2W/G is larger than the derived RW/W in the last section, indicating the introduction of graphene layers indeed gives rise to a finite interface thermal resistance. Taking into consideration of the uncertainties, RW/G is no larger than 11.9×10−9 Km2W−1. This value is comparable to the graphene/SiO2 interface resistance (5.6~12 × 10−9 Km2W−1) reported by Chen9 while is much smaller than that of the epitaxial graphene and SiC interface (5.3 × 10−5 Km2W−1) reported by Yue55. The corresponding thermal conductance (GW/G) is 84 MWm−2K−1 which is also at the high end of the graphene/SiO2 interface (20~110 MWm−2K−1) reported by Mak and Liu10 and that of the Au/Ti/graphene/SiO2 interface (about 25 MWm−2K−1) reported by Koh and Bae11. This GW/G value is much larger than 4~5 MWm−2K−1 reported by Jagannadham on characterizing thermal transport properties of W/graphene/Cu structure14. However, in his measurement, the structure was annealed at high temperature where reaction took place between W and graphene which reduces the interface thermal conductance significantly. Recent work by Huang et al.15 reported measurements on thermal conductance of Al/transferred graphene (trG)/Cu and Al/grown graphene (grG)/Cu interfaces, which is a good representative of the metal/graphene/metal interfaces. The reported G of Al/trG/Cu interfaces is about 20 MWm−2K−1, 35% lower than that of the Al/grG/Cu interfaces (about 31 MWm−2K−1). This lower G of the Al/trG/Cu interfaces is attributed to the lower conformity of trG to Cu substrate, which is further confirmed by the increase of G after the annealing treatment. The different degree of conformity could also be a reason for the variations of our measurement results. Despite the variations, our calculated G2W/G is no smaller than 42 MWm−2K−1, still larger than 31 MWm−2K−1 of the intrinsic value of their Al/G/Cu interface. The relatively large interface conductance of our β-W/G/β-W interfaces is due to the unique structure of the samples. One possible reason is that the graphene transfer and processing will inevitably introduce functional groups or defects to the graphene surface. These functional groups or defects may enhance the energy coupling between β-W films and graphene. Another reason may be the damage of graphene during the β-W sputtering progress, where additional channels of direct heat transport between the β-W films form and significantly enhance the thermal conductance of the interfaces. This damage enhanced phenomenon has been studied in the most recent work by Huang et al.16. In this work, they reported measurements of thermal conductance of Pd/trG/Pd interface with the top Pd prepared by either thermal evaporation or rf magnetron sputtering. The results shown that, G of the sample with the rf magnetron sputtering Pd is 300 MWm−2K−1 at room temperature, seven times larger than that with the thermal evaporation Pd (42 MWm−2K−1). This enhancement is attributed to the electronic heat transport via atomic scale pinholes formed in the graphene during sputtering process, which has also been proved by the AFM characterization. This can also help explain the variations and even negative value of RW/G in Table Note that, for each graphene layer, it has two W/G interfaces  peak of β−W. The pole-figure XRD shows that the β(200) plane is along the in-plane direction of the films. Based on the G and 2D peaks from the Raman spectra, the graphene samples were determined to be 1~2 layers. The thermal conductivity of β-W films is very low (1.69~2.41 Wm−1K−1) compared with that of the bulk α-phase tungsten (174 Wm−1K−1). This low k is mainly due to strong electron scattering of defects. The β-W/β-W interface thermal resistance was determined at (3.57 ± 2.67) × 10−9 Km2W−1, indicating excellent bonding and energy coupling at the W/W interface. The β-W/G interface thermal resistance was also determined. For each β-W/G interface, RW/G did vary from sample to sample, largely due to the inconsistency in sample preparation and unknown graphene structure damage and alternation. Taking into consideration of the uncertainties, the largest RW/G is 11.9 × 10−9 Km2W−1, corresponding to a lower bound of thermal conductance of 84 MWm−2K−1. Compared with the up-to-date reported graphene interface thermal conductance, the β-W/G interface thermal conductance is at the high end. The cross-plane thermal conductivity of β-W is smaller but still comparable to that of in the in-plane direction derived from the four-probe measurements and Wiedemann-Franz law.In this work, we have conducted systematic studies of the thermal resistance and conductance at the β-W/G multilayered film is synthesized by depositing tungsten film and transferring monolayered graphene to the tungsten film. Monolayered graphene films are grown on 25-μm thick Cu foils  by chemical vapor deposition (CVD). The transferring process of graphene is as follows. First, a PMMA (polymethyl methacrylate) layer is spin-coated on one side of the Cu foils and the graphene on the other side of Cu is cleaned by oxygen plasma. Then, the Cu foil is etched by FeCl3 solution. The remained graphene/PMMA is floated on the surface of FeCl3 solution, and is then cleaned by DI-water. The cleaned graphene/PMMA is transferred onto the surface of the tungsten film which is produced in an ultra-high vacuum magnetron sputtering system  at room temperature. Finally, the PMMA layer is dissolved by acetone to form the contact between tungsten and graphene. Through repeating the above process, the tungsten-graphene multilayered film can be obtained. It is worth noting that in one tungsten-graphene cyclic multilayered film, the thickness of each layer of tungsten film is kept the same. For convenience, in the following discussion, at the present of graphene, we name the thickness of each layer of tungsten film  as ‘A’, ‘B’, ‘D’, respectively [see Fig. et al.56 had ever placed a 1-mm-thick tungsten film in an oxygenated environment and investigated the oxidation behavior of tungsten under different temperature. It is found that the tungsten oxide was less than 1 nm even though the sample was exposed to air and heated to 100 °C for hours. While during our sample preparation process, the samples are immediately sent back to the vacuum chamber for sputtering the next tungsten layer after the graphene transfer process. The XRD patterns [see Fig. The see Fig. . For comsee Fig. . In addisee Fig. . In thissee Fig.  also shoα-W. However, a metastable form of β-W can also appear in thin films58. The black line in Fig. β(200) diffraction peaks of tungsten, which means the sample is mostly β-W with no other crystalline or amorphous and the preferred crystalline orientation is the β(200) direction59. The columnar structure shown in Fig. 60. The β(210), β(312) and β(400) peaks can also be found in the black line but are much weaker than β(200). The blue line in Fig. 61. We can see that the (110) diffraction peaks of α-W is the strongest, followed by the α(211) peaks and then the α(200) peaks. The large difference of these two XRD patterns reveals different phase of W. The crystallites or grain size of the tungsten films is also estimated using the Scherrer equation62 from the β(200) peak in the diffractogram, which is determined at 11 nm. Figure α angle  and the β angle . Figure α and β angle. Figure α = 0°.The structure of the tungsten films is explored through XRD to investigate phase, crystallite orientation and grain size. In bulk tungsten, stable structure of tungsten is mainly body-centered-cubic (bcc) −1. A 20× objective lens is used, with an integration time of 10 s and a laser spot size of 2.01 μm. The laser energy reaching the surface of the samples is 44.7 mW. Figure 53, the ratio of the integrated intensities of the G and 2D peaks (named IG/2D) can be used as a parameter to determine the number of layers of the graphene flake. From Fig. IG/2D) of the two samples are calculated at 0.38 and 0.42, respectively which means the graphene used in our experiment is 1~2 layered graphene53.AFM characterization is performed on the transferred graphene and the image shows kind of wrinkles in the graphene sheet while PMMA residues hardly appear. The AFM image is detailed in the Supplementary materials see Fig.  in S5. TSupplementary information


A novel metamaterial structure consisting of multiple graphene/dielectric layers and metallic substrate is proposed to achieve the broadband absorption response at terahertz (THz) frequencies. Utilizing the phase modulation effect generated by graphene ribbons, the bright-dark field is formed to suppress the reflection based on interference theory in a wide period. By irregularly stacking four graphene ribbons of varying widths on four dielectric layers with unequal thickness in a period, we merge successive absorption peaks into a broadband absorption spectrum successfully. The absorption decreases with fluctuations as the incident angle increases. The position of the absorption spectrum can be dynamically tuned by a small change in the Fermi level of graphene instead of re-optimizing and re-fabricating the device. In addition, the bandwidth of the absorber can be further improved by means of increasing the graphene/dielectric layers. The structure proposed in this paper has potential applications in tunable terahertz photonic devices such as dynamic broadband filters, modulators and sensors. Over the past decades, the rapid developments of MMs have attracted extensive attention from material scientists, physicists and engineers, because various types of fascinating metamaterial devices had been theoretically and experimentally investigated7. The appearance of split ring resonators8, cut wire pairs9 and metallic ribbons10 have greatly enriched the MMs theory12. Various types of exotic features that are unavailable in nature have been theoretically and experimentally investigated. MMs perfect absorber13, which is an important branch of MMs, has attracted a lot of attention currently, because the inevitable losses in metallic plasmonic nanostructures can be used for this research area. The first experimental demonstration of MMs absorber was given by Landy et al. in 200814, in which the electric and magnetic fields of the incident are absorbed by two MMs resonators integrated on the top and the bottom sides of a substrate. Several studies presented designs for broadband absorbers in the terahertz and microwave regions by merging successive absorption peaks with different geometrical parameters resonators in a single unit cell17. So far, most of the MMs absorbers are realized at a fixed operating frequency. Once the devices are designed and fabricated, tuning the absorption window is very difficult to achieve18.Metamaterials (MMs) are a kind of artificial structural materials with exotic properties not easily obtained or completely unavailable in nature19, acting as a prospective, alternative candidate for plasmonics at mid-infrared and THz frequencies22. Because of the fantastic electromagnetic properties, provided by the unique electronic band structure that the energy-momentum relation for electrons is liner over a wide range of energies24, graphene has attracted extensive attention in optics, electricity and magnetism. More importantly, one of the remarkable properties of graphene is that its Fermi level can be dynamically changed by means of external gate voltages25, which makes graphene an ideal material for realizing optical tunable devices. Currently, a variety of graphene-based MMs have been devised for enhancing optical absorption, such as, graphene ribbons26, discs27, even double-layer graphene structures28. However, the narrow absorption spectrum of the majority of graphene-based absorbers has greatly restricted its development.Graphene, a monolayer of carbon atoms arranged in a honeycomb lattice, is the first truly 2-D  material to be observed in nature34. Furthermore, the bandwidth of the absorber can be further improved by increasing the graphene/dielectric layers, which is incredible for previous researches36.In this paper, we propose an absorber, based on multilayer graphene ribbons, to suppress the reflectivity of incidence terahertz light in a broadband spectral range. A series of successive absorption peaks are merged into a broadband absorption spectrum, by irregularly stacking four graphene/dielectric layers on an optically thick golden substrate. This absorber is simulated and optimized using Finite Element Method. Simulation results exhibits that the broadband absorption with general efficiency exceeding 95% (from 9.9 to 10.9 THz) and the full bandwidth at half-maximum (FWHM) of 1.6 THz has been achieved, on the condition of vertical incidence. Meanwhile, this absorber exhibits good absorption stability over a wide angle range of incidence from −30° to +30° at least. The absorption windows of such absorbers can be dynamically tuned by changing the Fermi level of graphene. The absorbers are composed of simple ribbons instead of complex structure as usualR1~Rn), which avoid the electrical isolation caused by insulator. The incident light is transverse magnetic (TM) polarization light.The proposed absorber composed of multiple graphene/dielectric layers stacking on an optically thick metal film is depicted schematically in Fig. Λ is fixed at 17 μm, the thickness of the dielectric H is 7 μm, the width of the graphene ribbon L is 1.7 μm and the Fermi level Ef is 0.64 eV. There is a bridge connect all the ribbons together, which makes the tuning of the Fermi level easier. In Fig. a) appears distinctly with the resonant frequency of fa = 9.84 THz and the absorptivity is nearly 100%. The FWHM of the absorption peak is 0.6 THz (from 9.5 to 10.1 THz). The extra peak (peak b) in Fig. 38. At the position of absorption peak, the Ex field component distribution of the reflected light is shown in Fig. normE) distribution is shown in Fig. π) by the region contain graphene ribbons39, generating the phase matching between the bright and the dark field area. Therefore, the reflected light is suppressed due to the destructive interference of the bright field and dark field.In order to express the principle of reflection suppression more clarity, we take a simple single graphene/dielectric layer structure as example Fig. . The perH). If we change the dielectric thickness, the absorption shows periodic oscillations, and the oscillation period is about 11.5μm. The absorptivity of the system nearly reaches 100% within the scope of ΔH in each oscillation cycle. We define the area of ΔH as Efficient Absorption Range, and the value of ΔH is 5.5 μm . In the structure we proposed, H is fixed in the Efficient Absorption Range to achieve excellent absorption. Meanwhile, it is shown in Fig. H.To comprehend the characteristics of the system, we research the influence of absorption effects caused by the parameters variation based on the structure in Fig. Λ). When Λ < 15 μm, the absorptivity increases with the increase of Λ. When Λ >21 μm, the SPs disturb the absorption spectra forming multiple Rabi splitting analogues. When Λ varies between 15 and 21 μm, the absorptivity keeps high because the equal intensity of bright field and dark field. In this work, the period of each system is fixed at 17 μm.From Fig. Hi) is 5.8 μm in Fig. L and Li) increase, the resonance frequency is pushed toward the direction of the low frequency (red shift). As is well known, the frequency of the absorption peak is associated with the permittivity of the media around the graphene film. Therefore, when L is equal to Li, the corresponding frequency of the absorption peak in Fig. The response frequency of absorption peak is related to the width of the ribbons. In Figs. It should be pointed out that incident light often irradiate with an oblique incidence angle in actual application. In order to investigate the absorption sensitivity to the oblique incident light, we vary the incident angle α from 0° to 60°. Figure Ef, and the absorption effect is pretty good. The absorption peaks tend to exhibit blue shift with the increase of Fermi level, and there is slight effect on the SPs from Fermi level.For the sake of investigating the absorption response of the system on the Fermi level of the graphene, we present the absorption spectra of the system as a function of Fermi level in Fig. Λ, the thickness of the dielectric is H, the width of the graphene ribbons is L1 and L2, the distance between the two graphene ribbons is d and h in x direction and y direction respectively. Two graphene layers are tuned by different gate voltage independently, the Fermi level of the top layer graphene is Ef1 and the other one is Ef2.Through the research of the single layer absorber, we have comprehended the basic principle of reflection suppression. The absorber has excellent absorption effect and tenability. However, the narrow absorption bandwidth limits the application of the absorber seriously. Therefore, expanding the bandwidth has been our primary goal. A feasible method to expand the bandwidth of the absorption spectra is that increasing the number of the graphene ribbons. Then, we put two graphene ribbons in each period instead of one. The schematic of the double layers absorber is shown in Fig. Λ is 17 μm, H is 7 μm, L1 is 1.7 μm, L2 is 1.23 μm, d is 5 μm, h is 0.8 μm, Ef1 and Ef2 are both 0.64 eV. The absorption spectrum is shown in Fig. a, fa = 9.90 THz; b, fb = 10.12 THz) on the curve expanding the absorption spectrum. Figure x-y section at peak a and peak b. I and III are the field intensity and Ex field component distributions of the reflected light at peak a; II and IV are the field intensity and Ex field component distributions of the reflected light at peak b. We can obtain from Fig. The parameters of the double layers absorber are set as follows: x direction and y direction, the width of the graphene ribbons, the Fermi level and the incidence angle. Firstly, the distance between the two ribbons is taken into consideration. The simulated absorption spectra with varying distance between the two ribbons in x direction (d) and y direction (h) are plotted in Fig. h varying from 0.8 to 1.8 μm. When h = 0.3 μm, the absorption curve has changed a lot, for the reason that the graphene below is close to the interface of the dielectric and air, varying the medium condition around the graphene ribbon and inducing mismatching of the phase matching condition. It can be obtained from Fig. d is 5 μm, the absorption effect is the best. The absorptivity decline as d increasing or decreasing, and the intensity of the two peaks become unbalanced. However, when d reduce to 0 μm, great changes have taken place in the curve of the absorption spectrum and new peaks have replaced the original peaks.The absorption effect of the broadband absorber depends on various factors, including the distance between the two graphene ribbons in d is 0 μm, the vertical distance h is 0.8 μm, the thickness of the dielectric is 7 μm, the width of the ribbons are 1.7 μm (L1) and 1.23 μm (L2) respectively. New absorption peaks are produced, because the ribbons are so close that form a new consortium. It can be confirmed in Fig. c (11.18 THz) is formed and the two ribbons are as a consortium in the field distribution at peak c. The other small peak in the curve is the responding of high order mode. But, when the vertical distance h increases, the responding is changing. As shown in Fig. H is 17.5 μm, h is 10.5 μm, L1 is 1.7 μm and L2 is 1.2 μm, two ribbons are in different Efficient Absorption Range. Here, L2 is slightly tuned for better broaden effect. The absorption spectra is shown in Fig. d (fd = 9.88 THz) and e (fe = 10.16 THz). The illustrations in Fig. Ex field component distributions of the two peaks respectively. The two peaks correspond to different ribbon, and there is no interaction between two ribbons. We can obtain that the bandwidth of the absorption spectrum can be further broaden by increasing graphene layers in different Efficient Absorption Range.As shown in Fig. L2) are plotted in Fig. L1) is fixed and the other parameters are the same with that in Fig. x and y direction properly confirmed.The simulated absorption spectra of the double layers absorber with varying width of the lower ribbon  in Fig. 1, Ef1 = Ef2 = 0.64 eV; R2, Ef1 = 0.80 eV, Ef2 = 0.78 eV; R3, Ef1 = 0.50 eV, Ef2 = 0.51 eV. Both the absorptivity and the bandwidth of the absorption spectrum should be taken into account, when we adjust the Fermi level. As we can see, the absorption window shift toward higher frequency (blue-shift) with the increase of Fermi level, and the bandwidth of the absorption spectrum narrow down. The absorption window shift reversely (red-shift) with the decrease of the Fermi level, and the bandwidth of the spectrum become broader. However, in the considered range of the Fermi level, the absorption is always lager than 90% which is the powerful proof of the efficient tenability of the double layer absorber we proposed.The absorption spectra as a function of the Fermi level of the graphene ribbons are shown in Fig. Based on the foundational theory above, the broadband absorber proposed broaden the absorption spectrum by stacking multiple graphene/dielectric layers and merging successive narrow absorption spectrum into a broad one. We can learn from Fig. x-o-y coordinates is set up, offering great help for accurate expression. The period of the ribbons is Λ. The total thickness of all the dielectric layers is fixed as H. There are four ribbons, named as rn  respectively, distributing in each cycle. The width of ribbon rn is assumed as wn. We assume that the coordinates of ribbon’s center is . The Fermi level of each ribbon is Efn.The schematic of the multiple layers absorber is shown in Fig. Λ is 17 μm, the thickness of the dielectric H is 17 μm. The width and center coordinates of each ribbon are: ribbon r1, w1 = 1.680 μm, ; ribbon r2, w2 = 1.235 μm, ; ribbon r3, w3 = 1.130 μm, ; ribbon r4, w4 = 1.065 μm, . The absorption simulation result is shown in Fig. a (fa = 10.00 THz), b (fb = 10.30 THz), c (fc = 10.59 THz), d (fd = 10.79 THz) respectively. It is significant in Fig. a, b, c, d correspond to r1, r2, r3, r4, respectively.The parameters of the absorber in our simulations are as follows. The period of the absorber In Fig. 1, S2 and S3) in Fig. 1, Ef1 = Ef2 = Ef3 = Ef4 = 0.64 eV; S2, Ef1 = 0.75 eV, Ef2 = 0.73 eV, Ef3 = 0.71 eV, Ef4 = 0.70 eV; S3, Ef1 = 0.53 eV, Ef2 = 0.54 eV, Ef3 = 0.54 eV, Ef4 = 0.54 eV. The bandwidth of S2 is 0.8 THz (from 10.7 to 11.5 THz) with the absorptivity exceeds 80%, and the bandwidth of S3 is 1.2 THz (from 8.8 to 10.0 THz) with the absorptivity exceeds 90%. As shown, the absorption window shift toward higher frequency (blue-shift) with the increase of Fermi level, and the absorption window shift reversely (red-shift) with the decrease of the Fermi level. Through the research above, the dynamically tenability of the multiple layers broadband absorber is obtained.To confirm the tenability, we calculate the absorption spectra of normal incidence at different Fermi level, as shown in Fig. In conclusion, tunable broadband absorbers based on multi-layer graphene ribbons has been proposed and investigated. This structure is composed of several graphene/dielectric layers stacking on an optically thick golden substrate, in which the graphene ribbons behave as phase modulators. Because of the phase matching between the regions with and without graphene ribbons, the reflected light is suppressed in a broadband spectrum range. It is significant that the bandwidth of the absorber can be broaden unlimitedly as long as the added graphene layers are properly located in Efficient Absorption Range. Furthermore, the advantage of the graphene-based broadband absorber is that the Fermi level of graphene can be dynamically tuned by applying external gate voltages, without re-fabricating a new structure. By adjusting the Fermi level of the graphene ribbons, the absorption window can be adjusted dynamically. Meanwhile, the presented device exhibits good absorption stability over a wide range of incidence around ±30° at least. Thus, such a broadband absorber provides a novel way for the fabrication of nanophotonic devices for optical absorption in the terahertz region.40,e is the electron charge, ℏ is the reduced Planck’s constant, ω is the angular frequency, τ is the electron relaxation time, and Ef is the Fermi level. The electron relaxation time τ is calculated from τ = μEf/evf2. We take the carrier mobility μ = 104 cm2/Vs, the Fermi level Ef = 0.64 eV and the Fermi velocity vf = 106 m/s. The Fermi level of the graphene ribbons is associated with gate voltage. Because of the dependence of conductivity on the Fermi level, the resonance frequency (fres) of the absorber can be adjusted by changing the Fermi level, and has a relation 41. In the simulation, we apply transition boundary condition that allocates the conductivity to a single interface with effective thickness tg = 0.5 nm instead of a film with finite thickness. Such replacement can greatly decrease meshing load and simulation time. The permittivity of graphene can be calculated from εg = 1 + iσg/(ωε0tg), where ε0 is the vacuum permittivity. The permittivity of the dielectric εd is 1.90. Furthermore, the Drude model is applied to fit the experimental data of permittivity of gold for the fully reflecting mirror.The optical conductivity of graphene in terahertz and infrared region is dominated by the intraband transition, which can be described by a simplified semi-classical Drude modelThe datasets generated and analysed during the current study are available from the corresponding author on reasonable request.


Within the hybrid film, Co3O4 nanoparticles with size of 40~60 nm uniformly in-situ grew on the surface of graphene, forming a novel porous and interleaved structure with strong interactions between Co3O4 nanoparticles and graphene. Such fascinating microstructures can greatly facilitate interfacial electron transportation and accommodate the volume changes upon Li ions insertion and extraction. Consequently, the binder-less hybrid film demonstrated extremely high reversible capacity (1287.7 mAh g−1 at 0.2 A g−1), excellent cycling stability and rate capability .A self-supporting Co Among these materials, Co3O4 is especially noticeable due to its ease of synthesis and relative large theoretical specific capacity (~890 mAh g−1)6. Nevertheless, the practical application of this material is limited by its extremely low electrical conductivity as well as the rapid capacity attenuation resulted from the volume expansion/contraction as well as materials pulverization during cycling. One effective strategy to solve these problems is to design hybrid anode materials, which can not only provide active sites but also accommodate the volume change during charging/discharging process, thus attributes to higher capacity and cycling stability.The increasing demand for high performance energy storage applications makes it urgent to develop new lithium ion batteries (LIBs) which possess enhanced capacity, cyclic stability and rate capability. Transitional metal oxides (TMOs), have been considered as prominent anode materials, thanks to their high theoretical specific capacity. In this context, various anode materials with different novel structures, such as Co2 g−1), is envisioned as excellent candidates for substrate materials. The unique structure as well as distinctive electrical and mechanical performance is favorable to endure the strain caused by volume change and provide a continuous electron conducting networks, thus could greatly improve the electrochemical performance of active materials9. Recently, much effort has been conducted on integration of Co3O4 with graphene as anode materials for LIBs11. For example, Co3O4/graphene nanoflowers has been synthesized through a two-step process12, involving a hydrothermal reaction of Co(NO3)2·6H2O and CTAB with graphene oxide (GO) followed by a calcination. The reversible capacity of the as-synthesized composites is as high as 1120.8 mAh g−1 at 1 A g−1. Wu et al. employed a surfactant-assisted hydrothermal route to synthesize Co3O4/nitrogen-doped graphene, within which Co3O4 nanoparticles with size of ~15 nm homogeneous distributed in the macropore-walls formed by graphene. The obtained composites showed excellent electrochemical performance13. Yang et al. reported the preparation of porous Co3O4 nanofibers coated with graphene layer, which delivers a capacity of 900 mAh g−1 at 1A g−114. These results demonstrated the integration of graphene with Co3O4 can remarkably improve the electrical conductivity and alleviate the volume changes of Co3O4 during charge/discharge reactions, thus significantly enhanced the electrochemical performance. Even though, most Co3O4/graphene composites are powder materials, which need a complex slurry coating process in the conventional electrode preparation method. Moreover, the electrochemical-inactive polymer binder and carbon black are required to prepare electrodes, which would increase the extra weight and decrease the energy density of the actual battery16.Graphene sheets (GS), with high conductivity, excellent mechanical flexibility and ultra-high specific surface area  in-situ grew on the surface of GS, forming a novel interleaved structure with strong interfacial interactions. The hybrid electrode showed extremely high reversible capacity (1287.7 mAh g−1 at 0.2 A g−1), good cycling stability (capacity retention of 85.5% after 100 cycles) and excellent rate capability, owing to the fascinating microstructure and binder-free electrode nature, demonstrating great potential to be used in energy storage filed.Herein, we report the synthesis of a self-supporting Co3O4/GS hybrid films. The surface of GO, functionalized with oxygen-containing functional groups17, was measured to be a negatively charged surface  in this study. Co(OH)2 colloid, which was tested to be with a zeta potential of +22 mV, was a positively charged dispersion with high stability. Learning from colloid science, there are strong electrostatic attractive interactions between two colloid with opposite charges. As a result, a flocculent solution was produced immediately when mixing GO colloid with Co(OH)2 colloid, indicating an excellent self-assembly of the two dispersions driven by the strong electrostatic attractive interaction18. Afterwards, the flocculent solution can be easily vacuum-filtrated to construct a self-supporting Co(OH)2/GO hybrid film. Further two-step heat treatments will lead to the construction of a porous and free-standing Co3O4/GS hybrid film which can be used as anode electrode directly.Figure 3O4 during preparation procedures. As shown in XRD pattern of Co(OH)2/GO hybrid films . No obvious peaks of GO were detected, because graphene oxide sheets were highly separated by Co(OH)2 nanosheets. After annealed in Ar atmosphere at 600 °C for 1 h, Co(OH)2/GO was first converted to Co/GS as shown in Figure 3O4 2/GO 2/GO also exhibited stretching vibrations bands of Co-O bond centered at 489.8 cm−1 and O-OH bond in Co(OH)2 (~3631 cm−1)22. After thermal treatment, the intensities of peaks belonging to oxygen-containing functional groups significantly decreased for Co3O4/GS, suggesting the effective reduction of GO due to the thermal decomposition of these oxygen-containing functional group24. Meanwhile, the characteristic peak of υ(O-OH) in Co(OH)2 disappeared, instead with two distinctive peaks appeared at 563.1 and 657.6 cm−1, respectively. The first band υ1 at 563.1 cm−1corresponds to the BOB3 vibrations in the spinel lattice, in which B is associated with the Co cations in an octahedral position, i.e. Co3+ ions25. The second bands υ2 at 657.6 cm−1 is resulted from the ABO3 vibrations . These two stretching band can be well assigned to the characteristic peaks of Co3O4, further confirming the successful formation of Co3O4.The chemical changes from precursors to CoO Figure , the vis3O4 and Co3O4/GS composites. The three specific peaks at 458, 503 and 640 cm−1 can be assigned to Eg, F2g1 and Ag1 vibration modes of pristine Co3O4, respectively27. While the spectrum of Co3O4/GS displayed two distinctive peaks at 1347 and 1592 cm−1 apart from those for Co3O4, corresponding to the D and G band of graphene29, respectively. This further proves the successful preparation of Co3O4/GS composites, in accordance with the XRD and FTIR results. As shown in Figure 3O4 adsorption bands in the hybrid film showed significant blue shifts compared with those of pure Co3O4. This suggests that there are significant interfacial interactions between Co3O4 and graphene within Co3O4/GS hybrid films, which would greatly influence the interfacial lithium ion and electron transport.Figures 2 shows a sheet-like morphology and will be well dispersed on the matrix of GO via electrostatic attractive force e Figure 22. Afters Figure . More ime Figure . The diae Figure , with a e Figure . The nov3O4/GS hybrid electrode. As shown in the CV curves, a strong cathodic peak at about 0.6 V is observed in the first discharge process, corresponding to the multi-step electrochemical reduction between Li ions and Co3O4. Two anodic peaks appeared at ~1.5 and 2.3 V, owing to the oxidation of Co atoms and in consistent with the reported literatures31. The reactions can be described as:The hybrid film was directly used as electrodes, in which no polymer binder or carbon additives was added. Figure In the subsequent cycles, the cathodic peaks shifted to ~0.75 V, showing a tendency toward stabilization. The anodic peaks were almost the same and overlapped together, indicating the good electrochemical reversibility of the hybrid electrode.3O4/GS film was characterized by galvanostatic charge/discharge at 0.2 A g−1  of 74.9%, which was comparable to that of Co3O4/graphene hybrid electrodes reported previously32. The relative low initial CE value may be ascribed to the inevitable formation of a solid electrolyte interface (SEI) film over the electrode during the charge/discharge process, which led to the insufficient release of capacity. A well-established conductive network may help to further increase the CE32, such as surface modification13, pre-doping lithium metal33 and encapsulation34. Compared with the theoretical capacity of graphene (372 mAhg−1) and Co3O4 (890 mAh g−1), the extra capacity may owe to the formation of SEI film or interfacial Li-ion storage36. After fifty charge/discharge cycles, the discharge capacity still remains up to 1184.2 mAh g−1, demonstrating the good structural stability of the Co3O4/graphene anode.The lithium-storage performance of the binder-free Co1 Figure . In the 3O4/GS anode at 0.2 A g−1. The hybrid anode retains a reversible capacity of 1095.1 mAhg−1 after 100 cycles with the capacity retention of 85.5%. The rate capacity of the self-supported Co3O4/GS electrode has also been studied in Figure 3O4/GS heterostructures were about 1480, 1300, 1110, 920, 800, 620, 530 and 410 mAh g−1 at the current densities of 0.1, 0.2, 0.5, 0.8, 1.0, 1.3, 1.5 and 2.0 A g−1, respectively, manifesting an excellent rate capability. Importantly, the capacity re-increased to 1601.2 mAh g−1 when the current density returns back to 0.1 A g−1, further demonstrating the good reversibility. Compared with the capacities and cycling performance of other Co3O4-based anode materials reported previously  uniformly and tightly decorated on both surface of GS. The formed porous and interleaved microstructure exhibits critical characters as desired anode materials for LIBs, such as strong interfacial interactions, short transport length for Li-ion, and sufficient space for stress relaxation. Consequently, the constructed Co3O4/GS hybrid film as binder-free electrode exhibited a high capacity of 1287.7 mAh g−1 at 0.2 A g−1, good cycling stability (capacity retention of 85.5% after 100 cycles at 0.2 A g−1) and superior rate capability, making the hybrid films competitive as LIBs anode materials.In summary, a vacuum filtration procedure combined with two-step heat treatment has been developed to construct self-supporting Co2 was synthesized based on previous report and further diluted to a homogeneous dispersion (0.5 mg mL−1)41. Graphene oxide (GO) was prepared using a modified Hummers method with natural graphite as raw materials42, which was also diluted to 0.5 mg mL−1 for further use.β-Co(OH)3O4/GS film, the β-Co(OH)2 and GO dispersion were mixed together by sonication; and then vacuum filtered to form a self-supporting Co(OH)2/GO hybrid film. Afterwards, two-step heat treatment approach was carried out: First, the obtained Co(OH)2/GO film was heat-treated in Ar at 600 °C for 1 h to produce Co/GS. Second, Co/GS was further oxidized to Co3O4/GS by being calcined in air .For preparation of CoThe morphologies and microstructures of the samples were characterized via field-emission scanning electron microscopy(FE-SEMJSM-6700F). The crystalline phase of the materials was analyzed from X-ray diffraction measurements (Rigaku D/max 2550 V diffractometer). Raman spectroscopy was conducted on DXR Raman Microscope . Fourier transform infrared spectroscopy (FTIR) were examined on Nicolet 7000-C.3O4/GS hybrid films were directly used as an electrode without adding any binder or additive. The cell was assembled in glove box, and Li foil was adopted as the counter electrode. The 1 M LiPF6 in a mixture of ethylene carbonate (EC)/dimethyl carbonate (DMC) was used as electrolyte. The electrochemical properties were measured using a CT2001 battery tester at room temperature. Cyclic voltammetry (CV) was performed via electrochemical workstation (CHI760E) at a scan rate of 0.5 mV s−1 within a voltage range of 0–3.0 V. The mass loading was about 0.8 mg for each Co3O4/GS electrode.In this paper, the as-synthesized self-supporting CoSupplementary Information


Research on attentional bias modification has increased since 2014. A recent meta-analysis demonstrates evidence for bias modification for substance disorders, including alcohol and tobacco use disorders. Several pharmacological trials have shown that pharmacological agents can attenuate and modify such attentional bias. The pharmacological trials that have appeared to date have produced mixed results, which has clinical implications. Developments in Internet and mobile technologies have transformed how attention bias modification is currently being achieved. There remains great potential for further research that examines the efficacy of technology-aided attention bias interventions. Field et al.  and Cox Based on the dual-process model, it has been proposed that the chronic use of substances will increase the automatic processing of substance-related cues . The duaWhile Field et al.  highlighSince 2014, there have been further advances in the field of experimental psychology, with more research conducted on attentional bias modification. Attentional bias in addictive disorders has now been extensively investigated in a variety of addiction disorders ,7. Thus,To date, attentional bias modification interventions have been evaluated for a variety of addiction disorders, ranging from substance addiction to behavioral addiction. A prior meta-analytical study has synthesized the evidence for cognitive bias modification based on randomized trials published until 2015 . The revApart from the existence of further studies evaluating the effectiveness of bias modification, the advances in experimental psychology has also led to the creation of newer methods for bias assessment. Leeman et al.  highlighIn summary, the current evidence highlights that attentional biases are present in substance addictions. These attentional biases are amenable to modification using conventional interventions for attentional bias modification. While attentional biases were reduced following bias modification, the reduction in attention bias was not associated with improvement in other symptoms, such as cravings. The advances in the methods of assessing attentional biases and analyzing the involvement of neural circuitry have several implications for further research. Dopaminergic neurotransmission has also been responsible for unconscious attentional processes. In their systematic review, Luijten et al.  synthesiWhile the dopaminergic neurotransmission potentially could account for these unconscious attentional processes, other neurotransmitter systems have been implicated, such as the noradrenaline system and the glutaminergic neurotransmission system. Atomoxetine , a pharmLiu et al.  evaluateThe findings from these studies demonstrate that pharmacological agents are helpful in the modulation of attentional biases when these medications are given acutely. No studies have determined whether there will be further reduction in attentional biases if dopaminergic or other medications are administered for a longer duration. The evidence arising from Liu et al. demonstrates that the chronic administration of serotonergic medications does not alter attentional biases. There are clinical and research implications arising from these results, which will be discussed subsequently. The findings from the current studies also highlight that researchers seeking to evaluate attentional bias and the effectiveness of bias modifications need to consider the pharmacological agents that the sample individuals are already on, as it might confound the results of the psychological intervention. The introduction of digital technologies into healthcare has revolutionized various healthcare practices. E-health refers to the use of Internet-based technologies, telephone, and text-based messaging to support health interventions . M-healtWeb-based attentional bias modification programs have also been evaluated in adolescents with anxiety and depression, with demonstrable efficacy of the web-based program in reducing both anxiety and depressive symptoms . The stuOther studies have also described how smartphone devices have been utilized in the delivery of attentional bias modification intervention ,27. In tThere are clinical and research implications that arise from recent advances in attentional bias modification. The findings from the meta-analysis demonstrate that attentional biases could be subjected to modification. Targeting attentional biases is pertinent as according to the dual-process theoretical model, the repeated usage of an addictive substance would increase the automatic processing of substance-related cues, including attentional biases and automatic approach tendencies for substance-related cues . FurtherIn our literature review, we have not managed to identify studies involving bias modification in individuals with dual diagnoses. However, in the clinical context, it is common for individuals to present with comorbidities or with dual diagnoses. To date, researchers have largely focused on samples with an isolated diagnosis or disorder. It is important for there to be further research evaluating the efficacy of attentional bias modification for individuals with dual diagnoses and comorbidities. The advances in assessment of attentional biases also imply that researchers are no longer confined to a single measure, which was previously reaction time, in the determination of attentional biases. More recently, electrophysiological measures have been used to demonstrate the presence of attentional biases. Thus, it is pertinent for researchers who are planning further studies in this area to consider both an indirect and direct measure, which can potentially improve the accuracy of detection of attentional biases. Direct measures, such as electrophysiological measures, are ideal for demonstrating the neural circuits that are affected and potentially altered using bias modification. Direct measures could also inform clinicians of real-time changes in neural circuitry or changes in brain activation following each administration of bias modification. It would be crucial for future research to examine how technology-aided delivery of an attention bias modification task could be coupled with another direct objective measure, such as electrophysiological measurements. Furthermore, given how repetitive the task might be, it will be wise to consider pairing attentional bias modification intervention with motivational interventions . In the More research is needed to evaluate web-based interventions to establish their efficacy. There are advantages to web-based interventions, as individuals would not need to be confined to a particular environment to receive the interventions for attentional bias modification. Moreover, the advances in technology might also allow for users to partake in these interventions immediately when they are at risk of relapse.The findings from the pharmacological trials to date show that pharmacological agents are only efficacious in the acute reduction of attentional biases. While a diverse range of pharmacological agents has been evaluated, there remains a need for further studies evaluating the longer-term efficacy of pharmacological agents that modulate dopaminergic and other neurotransmitters. Current findings are important clinically as if these agents are useful in the chronic reduction of attentional biases, they could be considered for usage to prevent lapse and relapse in individuals. However, the usage of these pharmacological agents needs to be carefully considered due to their high propensity for inducing side effects. This includes extra-pyramidal side effects, amotivation, and apathy for medications, such as haloperidol. To minimize these side effects, it would be ideal if there were similarly efficacious psychological interventions for attention bias modification for which participants could be recruited. Researchers do need to take into careful consideration the pharmacological agents that individuals are on when they are evaluating attentional bias interventions, as some of these medications are potential confounders. In conclusion, there remains a need for further research synthesizing the evidence and effectiveness of attention bias modification for other substance disorders. It is also important for future research to evaluate the efficacy of attentional bias modification for individuals with dual diagnoses. Future research should also consider pairing an indirect measure with a direct measure for the assessment of biases. Motivational interventions and gamification strategies need further evaluation to determine their effectiveness in helping to increase intrinsic motivation of participants undertaking bias modification interventions. Studies should also evaluate web-based bias interventions to determine their efficacy. Future pharmacological trials should evaluate whether chronic administration of pharmacological agents is effective in reducing attentional biases.


The identification of the TAR DNA-binding protein 43 (TDP-43) as the ubiquitinated cytoplasmic inclusions in frontotemporal lobar degeneration (FTLD) and amyotrophic lateral sclerosis (ALS) confirmed that these two diseases share similar mechanisms, likely to be linked to the abnormal hyperphosphorylation, ubiquitination and cleavage of pathological TDP-43. Importantly however, a quantitative analysis of TDP-43 inclusions in predilection cortical regions of FTLD, FTLD-ALS and ALS cases has not been undertaken. The present study set out to assess this and demonstrates that distinct TDP-43 inclusion morphologies exist in the anterior cingulate cortex, but not the motor cortex of FTLD and FTLD-ALS. Specifically, in the anterior cingulate cortex of FTLD cases, significant rounded TDP-43 inclusions and rare circumferential TDP-43 inclusions were identified. In contrast, FTLD-ALS cases revealed significant circumferential TDP-43 inclusions and rare rounded TDP-43 inclusions in the anterior cingulate cortex. Distinct TDP-43 inclusion morphologies in the anterior cingulate cortex of FTLD and FTLD-ALS may be linked to heterogeneity in the ubiquitination of pathological TDP-43 inclusions, with the present study providing evidence to suggest the involvement of distinct pathomechanisms in these two overlapping clinical syndromes. Neuronal cytoplasmic aggregates containing the nuclear TAR DNA-binding protein 43 (TDP-43) are characteristically observed in the cortical neurons of frontotemporal lobar degeneration (FTLD) with TDP-43 pathology (FTLD-TDP), and in the lower motor neurons of cases with amyotrophic lateral sclerosis (ALS). The presence of this shared pathological protein has reinforced the theory that FTLD and ALS represent two ends of a disease continuum, with recent studies attributing the manifestation of this pathological protein in two distinct clinical syndromes to its accumulation in key brain regions , 25, 30.n = 23) also fulfilled pathological criteria for ALS) and 15 ALS cases (Table C9ORF72 and GRN genes using previously published methods [C9ORF72 expansions (n = 20) and 17% GRN mutations (n = 8)). This research project was approved by the Human Research Ethics Committees of the University of New South Wales.All cases with a pathological diagnosis of FTLD-TDP and/or ALS-TDP were selected from a neuropathological series collected by the Sydney Brain Bank through regional brain donor programs in Sydney, Australia. The brain donor programs hold approval from the Human Research Ethics Committees of the University of New South Wales and comply with the statement on human experimentation issued by the National Health and Medical Research Council of Australia. Patients were diagnosed during life by experienced clinicians using standard clinical diagnostic criteria , 26, 34  methods , 28 and The selection of the anterior cingulate cortex and motor cortex as the cortical regions-of-interest in this study was based on these regions being the predilection sites and early cortical regions targeted by TDP-43 aggregates in FTLD-TDP and ALS pathology that also demonstrate differentiation and overlap in TDP-43 between these two phenotypes  and that are rarely/not found to demonstrate TDP-43 in cognitively normal individuals and AD cases , 30, 35.Formalin-fixed, paraffin-embedded tissue blocks for each region-of-interest were sectioned at 10μm and immunostained with the anti-phospho TDP-43 monoclonal antibody . All slides were counterstained with haematoxylin for quantitation of cortical neuronal populations, as described previously . Brieflyp-value <0.05 taken as significant. Specific analyses were conducted to determine:The region with the greatest burden of TDP-43 morphologies using a univariate analysis with inclusion morphology and region as factors covarying for age and disease duration;Any differences between cohorts  in the regional density of different TDP-43 inclusions using ANOVA and posthoc tests for each region.Any relationships between all variables using principal component factor analysis, including regional density of different TDP-43 inclusions (4 variables), neuronal density, cohort group, genetic mutation and demographic variables (age and disease duration).Statistical analyses were performed using SPSS (IBM SPSS statistics version 24), with n = 13) and all other cases, and in the motor cortex of 93% ALS (n = 14), 86% FTLD (n = 24) and 100% FTLD-ALS cases (n = 18). Two prevalent morphologies were identified – rounded TDP-43 neuronal inclusions and circumferential TDP-43 neuronal inclusions , with significantly more TDP-43 pathology identified in the anterior cingulate cortex compared to the motor cortex  = 5.03, p < 0.05) across all cases. The same distribution of inclusions was observed in both regions with no interaction between morphology and regional burden of TDP-43 pathology.TDP-43 pathology was identified in the anterior cingulate cortex of 87% ALS n = 3 and allp < 0.05) with more circumferential TDP-43 inclusions in FTLD-ALS cases compared to ALS and FTLD cases (posthoc p < 0.001) and significantly more rounded TDP-43 inclusions in FTLD cases compared to ALS and FTLD-ALS cases (posthoc p < 0.005) , interaction  > 0.2, p > 0.3)). In contrast, no significant cohort differences were observed in the densities of the different TDP-43 morphologies in the motor cortex  > 0.7, p > 0.2).Previous studies  have sho05) Fig. . The prePrincipal component factor analysis was used to assess all relationships between variables across all cases and loading scores >0.6 was considered significant. Three separate factors accounted for >65% of variance. There were significant relationships between clinicopathological group (0.84 loading), the presence and type of genetic mutation −0.71 loading), disease duration (0.61 loading) and % rounded TDP-43 inclusions (>0.63 loading) . This supports the construct that rounded inclusions differentiated the FTLD group Fig.  which ha loading,The present study identified two prevalent TDP-43 inclusion morphologies in the anterior cingulate cortex and motor cortex of FTLD, FTLD-ALS and ALS cases – circumferential TDP-43 neuronal inclusions and rounded TDP-43 neuronal inclusions Fig. . QuantitThe present quantitative analysis of TDP-43 inclusion morphologies in predilection cortical regions implicated in FTLD, FTLD-ALS and ALS cases demonstrates a significant amount of circumferential TDP-43 inclusions in the anterior cingulate cortex of FTLD-ALS cases only. Importantly, these circumferential TDP-43 inclusions demonstrate a striking resemblance to: 1) the TDP-43 granular neuronal cytoplasmic inclusions recently described in the frontal cortices of FTLD-TDP type B cases, all with co-existing ALS ; 2) theThe present findings of significant amounts of rounded but not circumferential TDP-43 inclusions in the anterior cingulate cortex of FTLD cases, and significant amounts of circumferential but not rounded TDP-43 inclusions in the anterior cingulate cortex of FTLD-ALS cases also suggest that difficulties in reliably distinguishing between FTLD-TDP type A and B, even when attempted by experienced neuropathologists , other gFurther to previous studies demonstrating no significant difference in the presence and severity of TDP-43 pathology in the motor cortices of FTLD, FTLD-ALS and ALS cases , 25, 30,In summary, the present non-biased quantitative analysis in a large series of TDP-43 proteinopathy cases has shown distinct TDP-43 inclusion morphologies in the anterior cingulate cortex of FTLD and FTLD-ALS cases. Circumferential TDP-43 neuronal inclusions were predominantly identified in the anterior cingulate cortex of FTLD-ALS cases, and rounded TDP-43 neuronal inclusions were primarly observed in the anterior cingulate cortex of FTLD cases. The present results converge with recent findings demonstrating heterogeneity in the ubiquitination of pathological TDP-43 protein in FTLD cases associated with ALS , 22 to s