Growth plate injuries affecting the pediatric population may cause unwanted bony repair tissue that leads to abnormal bone elongation. Clinical treatment involves bony bar resection and implantation of an interpositional material, but success is limited and the bony bar often reforms. No treatment attempts to regenerate the growth plate cartilage. Herein we develop a 3D printed growth plate mimetic composite as a potential regenerative medicine approach with the goal of preventing limb length discrepancies and inducing cartilage regeneration. A poly(ethylene glycol)-based resin was used with digital light processing to 3D print a mechanical support structure infilled with a soft cartilage-mimetic hydrogel containing chondrogenic cues. Our biomimetic composite has similar mechanical properties to native rabbit growth plate and induced chondrogenic differentiation of rabbit mesenchymal stromal cells in vitro. We evaluated its efficacy as a regenerative interpositional material applied after bony bar resection in a rabbit model of growth plate injury. Radiographic imaging was used to monitor limb length and tibial plateau angle, microcomputed tomography assessed bone morphology, and histology characterized the repair tissue that formed. Our 3D printed growth plate mimetic composite resulted in improved tibial lengthening compared to an untreated control, cartilage-mimetic hydrogel only condition, and a fat graft. However, in vivo the 3D printed growth plate mimetic composite did not show cartilage regeneration within the construct histologically. Nevertheless, this study demonstrates the feasibility of a 3D printed biomimetic composite to improve limb lengthening, a key functional outcome, supporting its further investigation as a treatment for growth plate injuries. Normal bone elongation is uniform, with the growth plate becoming progressively thinner until complete closure occurs at skeletal maturation2. However, the growth plate also represents a weak point of the developing pediatric skeleton due to the juxtaposition of relatively soft cartilage against hard bone. Injuries involving the growth plate make up ~15–30% of all childhood skeletal injuries4 and of these up to 30% can lead to lifelong orthopedic problems5. The main concern with growth plate injuries is the formation of unwanted bony repair tissue known as a “bony bar” that leads to bony tethering of the epiphyseal and metaphyseal bones  is a cartilaginous structure that lies between the epiphysis and metaphysis of children’s long bones. It is a dynamic tissue that drives bone elongation through the process of endochondral ossificationnes Fig. . This bo6. On the other hand, when the bony bar is larger, more invasive procedures are warranted 7. To date the most common interpositional material is autologous fat12, but studies have shown poor to fair outcomes in 16–63% of patients treated with this approach13. Current interpositional materials are designed to act simply as a spacer, and thus are limited in that they do not function to replace the native growth plate tissue or in themselves induce growth through the normal biological process of endochondral ossification. If the desired outcome is a fully restored growth plate and new growth at the former site of bony tethering, a tissue engineering and regenerative approach is needed.Current clinical treatment options include resection of the bony bar when it occupies <33% of the growth plate and implantation of an interpositional material to avoid reformation of bony tethering and to prevent subsequent growth deformities15. Studies have also shown that overloading can suppress longitudinal bone deposition18. Thus, a successful tissue-engineered construct would require a structure that is stiff enough to support loading and prevent collapse but also that fosters an environment for stem cell proliferation and differentiation. Previously explored scaffold chemistries for growth plate tissue engineering include chitin20, agarose22, atelocollagen23, poly(lactic-co-glycolic acid) (PLGA)24, and demineralized bone matrix25. The primary purpose of these scaffolds was delivery of biologically active components  at the injury site, rather than provision of robust structural support. For example, PLGA was used to create a porous, degradable scaffold that locally delivered insulin-like growth factor 1, in a rabbit growth plate injury model24. Although these structures resulted in promising indications for growth factor stimulation of chondrogenesis, the scaffold itself was found to fragment during the final stages of degradation in vitro. The development of a treatment approach capable of providing both relevant biologically active cues to support regeneration and robust structural support, would be an important next step toward developing a clinically relevant treatment approach.The growth plate is subject to a complex local mechanical and molecular environment. Dynamic compressive forces have been shown to stimulate epiphyseal cartilage growth, while shear and hydrostatic pressures more likely drive cartilage ossification26 capable of supporting the regeneration of growth plate cartilage. Limitations of many printing methods include resolution and the need to print extra material to support overhangs. Stereolithographic digital light processing (DLP) is one promising printing method that addresses these shortcomings. DLP printing involves rapid liquid-solid transition under light irradiation and thus resolution is a function of the projector, making resolutions <50 µm possible27. Furthermore, DLP printing does not require a support material for overhanging features which allows for more complex designs compared to other 3D printing techniques.Three-dimensional (3D) printing provides a feasible method to fabricate structures that meet the mechanical needs of the biological system, combined with a structural design (low volume fraction) that allows for the inclusion of a bioactive niche31. We investigated the ability of our mimetic composite (gel + 3D print) to induce chondrogenic differentiation of rabbit mesenchymal stromal cells (rMSCs) in vitro. We then evaluated its efficacy as a regenerative interpositional material in a rabbit model of growth plate injury and compared it to an untreated control, cartilage-mimetic hydrogel only, and a fat graft as the clinical gold standard. Limb growth, tibial angle, morphological bone changes, and histological analysis of the repair tissue between the different treatment groups were performed. A second in vivo study compared different iterations of our growth plate mimetic composite with altered structural properties. This study combined 3D printing and a bioactive hydrogel to develop a mimetic composite as a potential treatment of growth plate injuries that addresses both the need for mechanical support and cartilage regeneration following a bony bar removal.The objective for this study was to develop a 3D printed growth plate mimetic composite for use following a bony bar resection to prevent bony bar reformation, support cartilage regeneration, and restore limb growth. To do so, we first characterized the mechanical properties of rabbit growth plate to guide the creation of the 3D printed growth plate mimetic composite. A poly(ethylene glycol) (PEG)-based resin was used in DLP 3D printing to create a stiff 3D printed structure made of vertical pillars Fig. . This siTo guide the creation of a 3D printed structure for the growth plate mimetic composite (gel + 3D print) Fig. , native 32. The resin material used to 3D print the structure had an elastic compressive modulus of 32 ± 2 MPa. When unfilled, the 3D print using Formulation A, B, and C had effective structural compressive modulus of 1.03 ± 0.21 MPa, 0.17 ± 0.03 MPa, and 0.27 ± 0.03 MPa, respectively  with modifications to the volume fraction of printed pillars, application of a thermal post-cure, and addition of a solvent  in the resin formulation Fig. . Represeely Fig. . Once inely Fig. . Howeverely Fig. , suggestTo assess whether the presence of a stiff 3D printed structure affects rMSC viability and morphology, cells were encapsulated in the cartilage-mimetic hydrogel, and were either directly polymerized (gel only) or used to infill a 3D printed structure and then polymerized (gel + 3D print). Samples from both conditions were stained with calcein AM (live cells) and ethidium homodimer (dead cells) to assess cell viability after nine weeks of culture Fig. . MinimalTo assess the chondrogenic differentiation of the rMSCs in the gel only or gel + 3D print conditions, Safranin-O/Fast Green staining and immunofluorescence for aggrecan and collagen types II and X were performed after one day, three weeks, and nine weeks of culture Fig. . Safrani33. During the first surgical procedure, when the rabbit was 6-weeks old, a 5 mm × 5 mm × 1 mm area of the proximal anterior-medial tibia growth plate was removed from the right limb , with no differences observed between control limbs for the different groups    or treatment limbs   Fig. . Limb leoc) Fig. . The tiboc) Fig. . At 8 weInjured limbs were assessed by microCT at 8-weeks post treatment, and all showed a disrupted growth plate at the injury site Fig. . No signTo assess the repair tissue that formed at the site of injury, Alcian Blue-Hematoxylin, Eosin Y, and Orange G (ABH) staining was performed. The uninjured limbs had a continuous cartilaginous (blue) growth plate region Fig. , while tp = 0.047; two-way ANOVA, Tukey’s HSD post hoc)   of time with respect to limb length growth. The average tibial length difference between the control and treatment limb increased from 4.5 ± 3.4 mm to 8.0 ± 4.4 mm and 10.0 ± 4.3 mm from 5-weeks to 8-weeks and finally 12-weeks post treatment. The mean tibial angle of the treatment limbs was 75 ± 8°, 71 ± 9°, and 71 ± 9° for 5-, 8-, and 12-weeks post treatment, respectively. No significant differences were observed between any bone metrics across the three different formulations  played a role in our treatment, we again used our rabbit growth plate injury model Fig.  and evaloc) Fig. . Tibial oc) Fig. . There wons Fig. .Fig. 10VInjuries to the growth plate that result in bony bar formation can lead to significant limb length discrepancies and angular deformities, which can subsequently cause gait disturbances, low back pain, cosmetic deformity, and early-onset osteoarthritis. Current treatments are directed at removing the bone tethering and preventing its recurrence but do not attempt to regenerate the damaged growth plate. Herein, we utilized 3D printing technology to create a growth plate mimetic composite that meets the mechanical needs of the biological system and contains a bioactive niche capable of supporting chondrogenesis. We characterized the mechanical properties of the rabbit growth plate to instruct the mechanical properties of the stiff 3D printed structure that was then infilled with the cartilage-mimetic hydrogel. In vitro we demonstrated that our growth plate mimetic composite induced chondrogenic differentiation of rMSCs. We observed that degradation of the PEG-based hydrogel provided space for local deposition of aggrecan, collagen II, and collagen X indicating new ECM was being produced within our cartilage-mimetic hydrogel both in gel only and gel + 3D print conditions. In vivo, the 3D printed growth plate mimetic composite allowed for significantly more limb growth compared to other treatment options, but limited cartilage-like tissue was observed, and it was not located within the construct. We also demonstrated that the material and structural properties of our composite material are tunable.37. These modulus values, 0.2–1.14 MPa, were used to guide the creation of three tunable composites with compressive moduli, 0.4–1.12 MPa, encompassing that of native tissue. While only one animal was evaluated in this study, subsequent work on a larger sample size  has since been published and values are in agreement with those presented here38. The open 3D printed structure allows for ease of infilling with a softer and cell instructive hydrogel. Our biomimetic composite is an improvement to other approaches that have been explored in growth plate applications that utilize either a soft hydrogel only , which have compressive modulus in the 0.001–0.150 MPa range39 far less than the mean native tissue modulus value of 0.48 MPa found in this study, or that have more physiologically relevant mechanical properties  but lack bioactivity.Lacking literature-based values at the time of this work, it was necessary to characterize the native tissue to aid in the design of our composite. Our testing methods determined a Hertz contact modulus which represents an aggregate tissue response and does not account for potential anisotropy. We observed differences in Hertz contact modulus with depth which is in agreement with growth plate studies in other speciesIn vitro, rMSCs encapsulated in both the gel only and gel + 3D print conditions were well dispersed. There was positive staining for newly synthesized aggrecan, collagen II, and collagen X confirming both chondrogenesis and new cartilage-like ECM production independent of the presence of the 3D structure. The encapsulated rMSCs secreted MMPs, but at levels that did not lead to rapid degradation of the cartilage-mimetic hydrogel, allowing for local ECM deposition and growth. While additional quantitative measures for assessing ECM production including gene and biochemical levels should be considered in future studies, the current data suggest that our composite design allows for independent control over the properties with the incorporation of a tunable stiff 3D printed structure and provides an inductive environment for chondrogenesis afforded by the soft cellular niche.Impaired bone elongation is a major clinical concern with growth plate injury and bony bar formation. Using a rabbit model of growth plate injury that develops a bony bar and limb length discrepancy, we were able to demonstrate that resection of the bony bar and treatment with a 3D printed growth plate mimetic composite (gel + 3D print) allowed for significantly more limb growth within 8-weeks post treatment compared to all other treatments. This increased bone elongation may be a result of the 3D print alone acting as an interpositional material limiting bony tethering and providing improved mechanical support compared to the other treatments. Thus, a major limitation of the current study is the absence of an empty scaffold condition. Since limb lengthening was improved with the gel + 3D print compared to the gel only condition it is clear the 3D printed scaffold is beneficial, but we are unable to discern the exact mechanism of in vivo efficacy. Additional work is needed to determine the role of the bioactive hydrogel in promoting cartilage regeneration in vivo both independently and as an infill in the 3D printed scaffold.While the untreated group had bone across most of the defect region, and the fat graft condition displayed bony tissue along with pockets of fibrous tissue particularly along the periphery of the defect, the gel only and gel + 3D print treatments had bony infill in the center of the defect but also had cartilage-like tissue with presence of sGAGs at the periphery. In the gel only condition, cartilage-like tissue was limited to the medial and peripheral most aspect of the treatment site and contained few chondrocyte-like cells. The 3D printed growth plate mimetic composite led to more intense staining for sGAGs with cells showing distinct chondrocyte morphology and some cells staining positive for collagen II. Given that the cartilage-rich regions were not within the 3D print but rather near its periphery and at the proximal edge of the injury area, it is possible that the gel only and gel + 3D printed growth plate mimetic may preserve surrounding growth plate cartilage. However, due to the extensive injury that occurred during the second surgical procedure that extended proximal and distal from the growth plate, as well as the organization of the tissue, it is also possible that these cartilage-like areas, particularly those seen in the gel + 3D print, are neo-tissue formation. Since histology was performed at the end of the 8-week study, it is unknown whether a more robust cartilage tissue formation occurred early on and underwent endochondral bone formation to support bone lengthening. Future studies with earlier time points and an unfilled 3D print could help elucidate the origin of the cartilaginous tissue and the mechanism that led to greater bone elongation with the 3D printed growth plate mimetic composite.25, we observed some degree of bony infill in all treatments. Since microCT was performed at euthanasia and partial growth plate closure had occurred, it is difficult to ascertain if the mineralization in the defect region was normal bone closure, or bony bar reformation. Previous studies have pointed to most interpositional materials being only acutely effective in preventing bony bar reformation. A study by Jin et al., found that their autologous tissue-engineered composite only prevented bony reformation in the first two weeks25. Based on in vitro analysis we did not anticipate the extent of mineralization observed in our gel only or gel + 3D prints in the in vivo studies. Future studies might consider evaluating earlier time points to determine the rate of mineralization within the scaffold. Additionally, design modifications to the scaffold including the lattice structure  might also prevent ingrowth of calcified tissues. We also noted minimal trabecula occupying the space distal to the growth plate. Given that we did not perform microCT on the contralateral control limbs in this study, we are unable to attribute that bone loss to normal age-related processes or a result of our injury model. However, we did not observe differences in the degree of metaphyseal bone reduction across groups, so we do not believe this is related to any treatment modality.A goal of interpositional materials is to prevent bony bar reformation, however, as reported in other tissue-engineered composite and gel studies11, the threshold for corrective osteotomy usually varies from 10 to 35° and is partially dependent on the location and extent of growth plate damage6. Moreover, it is not yet clear if the rabbit growth plate injury model is capable of a natural correction. Since we did not perform an osteotomy, the angular deformity present at the time of treatment remained. Experimental growth plate injury and treatment models have varied in the time treatment is delivered. Studies that treated immediately following injury have reported a higher instance of significant reductions in angular deformities and leg length discrepancies42 compared to those that treated 3 weeks post injury43. In our current model, we attempted to simulate a more clinically relevant condition where the injury is treated after presentation rather than at time of injury, but earlier intervention might result in more significant reduction in deformities long term due to the rapid rate of growth in rabbits at this age. Despite this challenge in the rabbit model, our gel + 3D print composite displayed positive effects in limb lengthening and warrants further investigation, even if a follow-up osteotomy may be necessary.Angular deformity is also a concern with growth plate injuries as it can lead to growth disturbances. In this study, tibial angle of the injured limbs continued to decline despite any of the treatments evaluated. This is likely due, in part, to the timing of the bony resection and subsequent treatments. On average, in the three weeks between the original injury and bony bar formation, when resection and treatment occurred, the tibial angle of the injured limbs changed by 9 ± 5°. While natural correction of angular deformities of up to 35° have been reported in humansMinimal differences in measured outcomes were observed between the three formulations for the 3D printed structures in the second in vivo study. When comparing the 8-week post treatment time point in the initial study to that of the second, all three formulations had increased tibial growth and reduced tibial angle changes compared to the untreated condition. This finding suggests that a 3D printed structure with a range of pillar diameters and composite moduli from 0.4 to 1 MPa does not affect the positive outcome. Thus, the benefits observed from our 3D printed growth plate mimetic composite are seen across a range of structural properties and appear to result in improved outcomes compared to softer treatments such as fat grafts or soft hydrogels.In conclusion, we have demonstrated the ability to 3D print a growth plate mimetic composite that led to increased tibial lengthening in a rabbit model of growth plate injury. The composite ultimately did not result in a robust layer of regenerated cartilage-like tissue 8 weeks post treatment. It also remains unclear if the improved tibial lengthening is mainly a result of the additional mechanical support provided by the scaffold or if the scaffold and cartilage mimetic hydrogel infill had a synergistic effect. In addition to improved limb lengthening, the mimetic composite has the advantage of being an off-the-shelf product, while the current clinical gold standard fat graft treatment requires a second surgical incision thus increasing patient risk for infection and recovery time. Further design modifications are necessary to improve in vivo chondrogenesis and reduce mineralized infill, but the results from this study are promising for moving towards using a regenerative interpositional material to treat growth plate injuries.44 to account for surface roughness. Hertz contact modulus, E*, was calculated at each point by fitting Hertz’ equation for spherical indentation of a half-space to the loading curve. We made no assumptions about the material’s Poisson’s ratio and therefore report Hertz contact modulus instead of Young’s Modulus. Indents exceeding 10 MPa were considered outside the soft tissue region of the growth plate and were discarded. Results were also confirmed using a nonlinear poroelasitc model, Hertz Biphasic Theory46.A single tibia sourced from a 9-week-old New Zealand white rabbit  was used to assess the mechanical properties of native growth plate cartilage. Sections were taken from the sagittal plane of the proximal tibial growth plate and polished to 2500 grit . Arrays (17×1) of micro-indents spanning the epiphyseal bone, growth plate cartilage, and metaphyseal bone were performed  at eight evenly spaced locations spanning the sagittal slice. For each array, indents were made with a diamond R = 50 μm spherical probe and spaced 50 μm apart. The probe was first lifted off the sample surface, then indented to a depth of 15 μm. A virtual contact point was determined47. Briefly, the PEG amine, dissolved in dimethylformamide (DMF) and dichloromethane (DCM), was reacted overnight, under argon, with eight times molar excess of 5-norbornene-2-carboxylic acid in the presence of three molar excess 1-[Bis(dimethylamino)methylene]-1H-1,2,3-triazolopyridinium 3-oxide hexafluorophosphate (HATU) and four molar excess N,N-diisopropylethylamine (DIPEA). The product was recovered and purified. Comparison of the area under the peak for the allylic hydrogen closest to the norbornene hydrocarbon (δ = 3.1–3.2 ppm) and the peak for the PEG backbone methyl group (δ = 3.4–3.85 ppm) in an 1H-NMR spectrum were used to determine a ~100% conjugation of norbornene to the 8-arm PEG. Thiolated chondroitin sulfate (ChS-SH) was then synthesized3. Briefly, Dimethyl 3,3′-dithioldipropinoate (DTP) was reacted with eight molar excess hydrazine hydrate. The product (DTP dihydrazide) was reacted with chondroitin sulfate A (Sigma Aldrich) in a two-molar excess at a pH of 4.75, followed by a reaction with 1-ethyl-3-(3-dimethylaminopropyl)carbodiimide. The reaction was stopped by adjusting the pH to 7.0. Dithiothreitol was added in a 6.5 molar excess and the pH was adjusted to 8.5. The product was recovered and purified. Comparison of the area under the peak for the methylene groups of DTP (δ = 2.5–2.6 and 2.6–2.8 ppm) and the area under the peak for the methyl group of the acetyl amine side chain of the chondroitin sulfate backbone (δ = 1.8–2.0 ppm) in an 1H-NMR spectrum were used to determine a ~18% conjugation of thiol groups per ChS molecule. For the cell-laden in vitro cell studies and the in vivo studies, TGFβ3 (Peprotech) was functionalized by addition of a thiol group by reaction with four molar excess Traut’s reagent in buffer, for one hour. Once functionalized, TGFβ3-SH was tethered to the PEGNB macromer via thiol:norbornene photoclick reaction.A 10 kDa, 8-arm poly(ethylene glycol) (PEG) amine (JenKem) was modified with a norbornene group at the end of each armPoly(ethylene glycol) diacrylate  and pentaerythritol tetrakis(3-mercaptopropionate)  were mixed at 99:1 wt% with 0.85 wt% diphenyl-phosphine oxide , as the photoinitiator, 0.05 wt% of 2,2′-Azobis(2-methylpropionitrile)  as the thermal initiator and 0.8 wt% 2-(2-hydroxyphenyl)-benzotriazole derivative  as the photoabsorber. The photo-resin used for Formulation C differed in that 50 wt% water was added and the Tinuvin Caroboprotect photoabsorber was replaced for 0.08 wt% Quinoline Yellow (Aldrich) due to solubility. AIBN was added to the resin prior to performing the experiment to avoid degradation.2 and since the pillar array was less dense in Formulation B and C the open regions on the lattice in these designs were 256 × 167 μm2. Formulation A design was sliced into sixty 25 μm layers, Formulation B and C designs were sliced into thirty 50 μm layers, to form a structure with a height of 1.5 mm before swelling. Each layer was exposed for 6 s at an irradiation intensity of 14.4 mWcm2 as measured by a power meter at λ = 405 nm where the power was 0.9 +/− 0.1 mW for a 0.0025 cm2 square image . After printing, Formulation A structures were subsequently post-cured in an oven at a temperature of 105 +/− 5 °C under vacuum for 1 h. All were soaked in 100% ethanol for a minimum of 72 h.Three different formulations were used in this study to assess the impact of structural properties. All were created using CAD designs consisting of vertical 200 μm diameter pillars connected by a lattice on each side. Formulation A had 98 pillars, which yields a theoretical pillar volume ratio of 12.3% while Formulation B and C had 61 pillars corresponding to a theoretical 7.6% volume fraction of solid pillar structure in the CAD design. To connect these pillars Formulation A had a lattice with open regions of 256 μmThe biomimetic soft hydrogel was prepared by photopolymerization of 9 wt% poly(ethylene glycol) norbornene (PEGNB), 50 nM TGFβ3-SH, 1 wt% ChS-SH, 0.1 mM GCRGDS (GenScript), MMP-sensitive peptide , and 0.05 wt% photoinitiator Irgacure 2959 (I2959) in phosphate-buffered saline (PBS) yielding a final thiol:ene ratio of 1:1. The hydrogel precursor solution was either polymerized directly or infilled into the 3D printed structure and then polymerized under 352 nm light for 8 min.n = 8) and the 3D print structure with the pillar array for each formulation both unfilled and filled (n = 5/group/condition) were assessed in unconfined compression . The 3D printed structures were infilled with the cartilage-mimetic hydrogel precursor solution and then polymerized to create the composite (gel + 3D print) or were tested without the cartilage mimetic hydrogel infill (unfilled 3D print). All samples were swollen to equilibrium in PBS prior to testing. Following a 3 mN pre-load, a constant strain rate of 0.03 mm/s was applied to each sample. Single pillars were loaded until they failed while the 3D printed structures were strained to 12.5%. To account for the change in cross-sectional area with compression the modulus was derived from a fit of the linear region of the true stress-true strain curve (assuming a Poisson’s ratio of 0.5). For analysis of the composites and unfilled scaffolds, it is assumed that the contact force was evenly distributed throughout the contact surface. Cross-sectional area of the unfilled 3D prints was based on the theoretical pillar volume fraction. Atomic force microscopy  was used in fast force mapping (FFM) mode to measure the indentation modulus along the surface of a sectioned 3D printed structure and evaluate the efficacy of thermal post-curing to reduce material property fluctuations as a function of layer depth. A cryo-ultramicrotome  was used to generate a flat cross-section of the structure for accurate FFM. A silicone doped probe with a tip radius of curvature <10 nm  and a spring constant of 3.44 nN/nm was used as it is specifically tailored for FFM. To account for adhesion, AFM indentation modulus measurements were analyzed using a Johnson–Kendall–Roberts contact mechanics model48. Pillar diameter measurements were taken from SEM images. Images were scaled, and multiple measurements along the pillar height were taken and averaged for a single value of pillar diameter. A total of n = 13, 12, and 21 pillars were measured for Formulations A, B, and C, respectively.A 3D printed structure of a single pillar representing the polymerized resin material  and the stiff 3D printed structure infilled with the cell-laden cartilage-mimetic hydrogel . Rabbit MSCs were encapsulated at 50 million cells mL−1 of filter-sterilized (0.22 µm) hydrogel precursor solution. For the gel only condition, the precursor solution was transferred to a 5 mm × 5 mm × 1 mm square Teflon mold and photopolymerized. For the gel + 3D print condition, the cell-laden precursor solution was pipetted into sterile 5 mm × 5 mm × 1 mm 3D printed structures and allowed to completely fill the interior. Both conditions were polymerized under 352 nm light at 5 mW cm−2 for 8 min. Cell-laden constructs were cultured for up to 9 weeks in chondrogenic differentiation medium  under standard cell conditions of 37 °C with 5% CO2. Medium was exchanged every other day.Femoral bone marrow from New Zealand white rabbits was harvested and nucleated rMSCs were isolated using density gradient centrifugationn = 2 per time point) were embedded in glycol methacrylate (GMA) using the Technovit® 7100 kit (EMS) per manufacturer instructions. Samples were sectioned at 5 µm using a tungsten blade on an automated microtome. GMA-embedded sections were stained with hematoxylin and eosin (H&E) to visualize cell morphology and non-specific tissue deposition using light microscopy . The remaining half (n = 2 per time point) were dehydrated and embedded in paraffin following standard protocols. Paraffin-embedded samples were sectioned to 5 µm. Sections were stained with Safranin-O/Fast Green to visualize sulfated glycosaminoglycans (sGAGs) using light microscopy . Immunohistochemistry was completed as follows. Deparaffinized and rehydrated samples were pretreated with specified enzymes  for aggrecan, pepsin (280 kU) for collagen X) for one hour at 37 °C and antigen retrieval (collagen X and aggrecan). Sections were blocked with 10% Normal Goat Serum (NGS) and 2% Bovine Serum Albumin (BSA) in Tris Buffered Saline (TBS)  or 1% BSA in PBS (aggrecan) and permeabilized with 1% BSA and 0.25% Triton-X 100 in TBS (collagen II and collagen X) or 1% BSA and 0.25% Triton-X 100 in PBS (aggrecan). Sections were subsequently treated with primary antibodies against aggrecan , collagen type II , collagen X , and PEG (1:50 Academia Sinica Anti-PEG 6.3) followed by secondary antibodies with conjugated AlexaFluor 488 or 546 probes and counterstained with DAPI for nucleus detection. Negative controls were prepared by omitting the primary antibody. A laser scanning confocal microscope  was used to acquire images at ×40 magnification (Plan Fluor 40X oil objective). Representative images were selected from those acquired.Samples were taken from both conditions at nine weeks to visualize the encapsulated cells for assessment of viability and distribution within the hydrogel. Structures were incubated at 37 °C in a solution of 4 nM calcein AM (live cells) and 2 nM ethidium homodimer (dead cells). Stained samples were then immediately imaged on a confocal microscope . At day 1 and at 3- and 9-weeks, samples were fixed in 4% paraformaldehyde. Half of the samples (n = 9). Briefly, 8 mM APMA enzyme activating solution was added to each of the media samples to activate MMPs and incubated for one hour at 37 °C. MMP substrate solution was then added to activated samples and incubated for one hour and fluorescence (490/520 nm) was measured. Total MMP in solution was calculated using collagenase type-II as a standard. Measurements for each week, within a given sample, were combined to yield the total MMPs produced during the duration of the study.Total MMP in media was measured with the Anaspec Generic MMP Assay kit (AS-71158) following manufacturer instructions  were assigned to one of four treatment groups: (1) untreated (n = 10), (2) autologous fat graft (n = 9), (3) cartilage-mimetic hydrogel (gel only) (n = 10), and (4) 3D printed growth plate mimetic composite (gel + 3D print) . To evaluate the effect of scaffold design, a second in vivo study was performed using twenty-four 6-week-old New Zealand white rabbits  equally assigned to one of three gel + 3D print formulation groups: Formulation A (n = 8), Formulation B (n = 8), and Formulation C (n = 8). Right limbs were injured, and the left limbs served as a non-operative control.All procedures complied with the Guide for the Care and Use of Laboratory Animals and were approved by the Institutional Animal Care and Use Committee at the University of Colorado Denver. The first in vivo study was used to evaluate the impact of treatment modality. Thirty-nine 6-week-old New Zealand white rabbits , gel only, or a gel + 3D print condition. For these studies, a lithium acylphosphonate (LAP) photoinitiator was used in the hydrogel precursor solution and a handheld 405 nm light was applied for 1 min. In the gel + 3D print condition, the 3D print was infilled with the hydrogel precursor solution and placed in the defect. An additional amount of precursor solution was added to ensure complete infilling, followed by light polymerization. All animals were allowed normal cage activity, exercise, and food and water at-lib. Animals were euthanized via overdose of isoflurane and intracardiac injection of pentobarbital (150 mg/kg) at 8 weeks post treatment for the first in vivo study evaluating different treatments and at 12 weeks post treatment for the second in vivo study evaluating scaffold formulation.An initial 5 × 5 × 1 mm proximal tibial physeal injury was created from an anterior-medial approach using a 1 mm drill bit and intraoperative radiographic images for defect positioningAnteroposterior radiographic images using 50 kV and 0.4 mA  were taken of the lower limbs before growth plate injury, during treatment, and at 4- and 8-weeks post treatment for the first study, and at 5-, 8-, and 12-weeks post treatment for the second study. These images were used to measure tibial length and tibial angle of both the injured and control limbs with ImageJ (v 1.53c with Java 1.8.0_172). Tibial length was measured as the distance between the tibial distal growth plate and the proximal tibial plateau at 50% of the full width. Tibial angle was assessed as the angle between the leg length measure and the average angle across the entire plateau. Tibial length change was assessed as change from time of treatment.50 was implemented to separate ranges of the histogram and binarize the data set. Segmentation of cortical and trabecular bone was performed using the Buie Method51, excluding the fibula, and evaluating equal volumes above and below the growth plate to accommodate any variability in scaffold position. The bone analysis plugin for Dragonfly  was used to evaluate bone volume fraction (BV/TV) as well as trabecular thickness (Tb.Th.) and trabecular spacing (Tb.Sp.) using a sphere fitting algorithm52.Following euthanasia, proximal tibias from treatment limbs were harvested, wrapped in parafilm to prevent dehydration, and imaged. Animals in the first study were imaged with an Xradia 520 Versa  at a voxel size of 26 μm . Automated centering and beam hardening corrections were applied using the Scout-and-Scan Control System Reconstructor software (version 14.0.14829.38124). Animals from the second study were imaged with SCANCO μCT50  at a voxel size of 17.2 μm . Reconstructions were imported into Dragonfly  and an Otsu’s algorithm−1) for one hour at 37 °C. Sections were blocked with 10% Normal Goat Serum in Tris Buffered Saline (TBS) and subsequently treated with a primary antibody against collagen type II  overnight at 4 °C. The secondary antibody, peroxidase-conjugated goat anti-mouse IgG (Jackson ImmunoResearch), was diluted 1:500 and incubated at room temperature for 1 h. Slides were developed using the DAB substrate kit (Vector Laboratories). Negative controls were prepared by omitting the primary antibody. Slides were digitized with the Axio Scan.Z1 (Zeiss).Tissues were fixed in 10% neutral-buffered formalin for 7 days and then decalcified in 14% ethylenediaminetetraacetic acid for up to 4 months. After decalcification, samples were embedded in paraffin, coronally sectioned (5 μm), standard Alcian Blue Hematoxylin was performed to assess repair tissue. For immunohistochemistry, samples were pretreated . A two-sample t-test was used to assess differences between the gel only and gel + 3D print for MMPs (two-sided). Comparisons between groups in all in vivo analyses were evaluated using a two-way ANOVA with Tukey’s HSD post hoc test for multiple comparisons (two-sided). All statistical tests were evaluated as two-sided, and significance was set at p < 0.05. All statistical analysis was performed using GraphPad Prism .All data reported are mean ± standard deviation. Pillar diameter measures were not homogeneous or normally distributed so a Kruskal–Wallis test (one-sided) was used to evaluate differences across formulations. Compressive modulus across the three formulations was evaluated using a one-way analysis of variance (ANOVA) with Tukey’s HSD post hoc test to correct for multiple comparisons (two-sided). Differences between empty and infilled scaffolds were assessed with a Further information on research design is available in the REPORTING SUMMARY


Nature 10.1038/nature09687 Published online 26 January 2011Correction to: 1 found that eight of the ten samples were mislabelled, and we wish to correct the original publication according to their findings. In addition, Banes et al.1 find that one of the orang-utans that we had labelled as a male Sumatran is in fact a female Tapanuli orang-utan ). While we here correct the original tables to display the correct names and IDs of the individuals studied, we do not correct the analyses in the original paper since the Tapanuli species was discovered only after we had concluded our study. We refer the reader to Banes et al.1 for further information.In this Article, we reported genome sequencing data from five Sumatran and five Bornean orang-utans. In re-analysing these data, Banes et al.four Sumatran, five Bornean and one Tapanuli, using a short read sequencing platform ”.In the main text, the sentence “We also sequenced the genomes of 10 additional unrelated wild-caught orang-utans, five Sumatran and five Bornean, using a short read sequencing platform ” should have read “We also sequenced the genomes of 10 additional unrelated wild-caught orang-utans, four Sumatran and one Tapanuli) (Supplementary Information section 4).In addition, the sentence “We then called single nucleotide polymorphisms (SNPs) from the alignment of all short read data from 10 individuals  (Supplementary Information section 4)” should have read “We then called single nucleotide polymorphisms (SNPs) from the alignment of all short read data from 10 individuals  should be recalculated for the Sumatran and Tapanuli separately in the future. Our Bornean estimates in Fig. 5 are unaffected by this sample mislabelling, as noted in Banes et al.1. The original Article has not been corrected online.None of these corrections affect the results and conclusions of the original Article with the exception of Fig. 5 in the main text. The estimates of SNPs in panel Supplementary Table 1A corrected version of the original Supplementary Table S4-1 (Next generation sequence data summary).Supplementary DataA corrected version of the original Supplementary Data (Accession table for Bornean and Sumatran orang-utan short read data).


Here, we calibrate a mechanistic model for gas exchange in modern lycophytes and constrain atmospheric CO2 levels 410–380 Ma from related fossilized plants with bound uncertainties of approximately ±100 ppm (1 sd). We find that the atmosphere contained ~525–715 ppm CO2 before continents were afforested, and that Earth was partially glaciated according to a palaeoclimate model. A process-driven biogeochemical model (COPSE) shows the appearance of trees with deep roots did not dramatically enhance atmospheric CO2 removal. Rather, shallow-rooted vascular ecosystems could have simultaneously caused abrupt atmospheric oxygenation and climatic cooling long before the rise of forests, although earlier CO2 levels are still unknown.The emergence of forests on Earth  Dahl et al. present new evidence based on leaf gas-exchange in primitive vascular plants and fossil remains of some of their earliest ancestors. This alters our thinking on how plants impacted the Earth System and climate. Its variation in the past informs us about the natural long-term sources and sinks. In the absence of anthropogenic fossil fuel combustion, the dominant atmospheric CO2 source is volcanic outgassing, and this source is balanced mainly by the removal that occurs when CO2-bearing fluids chemically react and weather silicate rocks followed by deposition of carbonate in the oceans4. The dissolution of silicate minerals in the weathering zone occurs via interactions between the terrestrial ecosystem and geological processes that make fresh rock available at the surface for reaction. Yet, the role of biology and the CO2-sensitivity of the feedbacks governing global CO2 removal is debated9. Enhanced continental weathering is suggested to have caused a decline in atmospheric CO2 pressure (pCO2) from a level ~10 times higher than today’s concentration11 to near modern levels linked to the Devonian-Carboniferous transition from greenhouse to icehouse conditions in response to the afforestation of the continents. This process is accompanied by burial and preservation of organic matter that also influence atmospheric CO2 levels and, in turn, acts as the main long-term source of atmospheric O2. However, recent geochemical evidence and Earth system models12 suggests atmospheric oxygenation occurred well before trees evolved on the continents ~393–383 Ma3. Further, the temporal correlation between plant colonization and the Permo-Carboniferous glaciation has been disputed8. There is compelling evidence that Earth also transitioned into a glaciated state in the Ordovician-Silurian13. Yet, the link between glaciation and atmospheric CO2 is complicated, and palaeoclimate models shows that glaciations could persist even at 12–14 times pre-industrial atmospheric levels 14. Therefore, a precise reconstruction of atmospheric pCO2 in relation to plant evolution is key to assess the impact of the terrestrial biota on Earth’s climate. Here, we show that atmospheric pCO2 was markedly lower than previously thought when trees and forests appeared on our planet.Atmospheric CO2 concentrations were one order of magnitude above pre-industrial levels in the early Palaeozoic17, although more recent studies suggest levels much closer to today18. Traditionally, palaeo-CO2 estimates from proxy data come with large and sometimes unbounded uncertainty20. For example, in the Late Ordovician (~445 Ma), there is evidence for high pCO2 levels of 17 ± 4 PIAL  from CO2 hosted in pedogenic goethite from the Neda Formation in Wisconsin, USA19. This paleosol shows a coupling between CO2 content in goethite and its isotope composition interpreted to reflect variable mixing of atmospheric CO2 and soil respired carbon substituted into the goethite mineral lattice. The range above reflects all analytical errors propagated through the calculation , and we note that the error could still be larger if the CO2 surface adsorption properties21 on natural goethite deviate from that of phosphated goethite grown in the laboratory22. This has never been verified in modern soils, and the proxy has also never been applied at any other time in Earth history.In the canonical view, atmospheric CO2 levels through the Devonian from 5 to 0.7 PIAL has been inferred from the carbon isotope compositions of pedogenic carbonate17. Although the declining trend may be real, the absolute atmospheric CO2 level reported from pedogenic carbonates from this time interval have been adjusted further down23 and are likely systematically overestimated because of lower productivity in early Palaeozoic soils relative to modern soils17. Propagating the uncertainty associated with assumed model parameters that cannot be independently constrained from the rock record, such as the proportion of soil-respired CO2 in the soil and its isotope signature, shows that the absolute palaeo-pCO2 estimate obtained this way carries one order of magnitude uncertainty 20 comes from low stomatal density (mm−2) in some fossil plants . These anatomical features were interpreted as evidence that these plants had adapted to minimize water loss in a high CO2 atmosphere, but it is essential to compare to plants with similar gas exchange anatomy and behavioural control over water loss rate and CO2 uptake. This becomes very problematic with this group of enigmatic early vascular plants that lack living descendants. In comparison fossils of lycophytes that co-occur with these extinct plant groups that do have modern relatives with similar physiology display similar stomatal density as their modern descendants  and Mid-Devonian (~700–1400 ppm) based on carbon isotope data of marine phytoplankton recorded via a diagenetic product of chlorophyll, phytane18. No phytane-CO2 data is reported between ~432 Ma and 390 Ma . When applied to modern phytoplankton, the phytane proxy predicts a wide range of atmospheric CO2 levels even today (~300–1200 ppm)26. The accuracy of this proxy depends on variables that are not easily detectable from the geological record, including the ratio of atmospheric CO2 to dissolved CO2 of the seawater in which the phytoplankton grew, the nature of the phytoplankton species, other sources of phytane, the growth rate of phytoplankton and, thus, the isotopic response to ambient dissolved CO2 concentration of the phytoplankton from which phytane was derived26. Given the many uncertainties, the errors associated with the phytane CO2-proxy are likely large.Recently, atmospheric COleaf) is sensitive to the internal CO2 concentration in the substomatal cavities that, in turn, depends on water availability and ambient pCO2 in the environment27. Recently, the combination of isotope data and stomatal parameters has led to the development of a mechanistic proxy for prediction of palaeo-pCO2 based around leaf-gas exchange27. Applying this approach to the fossil record shows that post-Devonian atmospheric CO2 levels were <1000 ppm most of the time, but CO2 estimates from the Lower and Middle Devonian using this model show considerable variation with estimates ranging from 530–2853 ppm with one outlier reaching 7320 ppm27.By the same fundamental principle plus an opportunity to constrain all involved parameters, carbon isotope fractionation in plant tissue from terrestrial plants  obtained from the fossils to predict palaeo-pCO2. Both characteristics are an expression of how the plant, when alive, was adapted to its local environment. Consequently, there is a strong relationship between these different “recorders” of plant climate interactions. For example, the carbon isotope signature of leaf tissue has been demonstrated to be strongly influenced by water availability28 and the utility of mechanistic proxies underpinned solely by isotope data31 to be questioned36. Due to the co-dependency of isotope and stomatal parameters it is recommended that these data are ideally extracted from the same sample whenever possible or from contemporaneous sedimentary deposits. In taking this approach, the accuracy of the prediction should be improved because it simultaneously accounts for variable CO2 assimilation rate and water conductance in response to changing water availability27. Here, we calibrate a leaf gas-exchange model for lycophytes and estimate atmospheric CO2 levels with uncertainties by applying it to some of the oldest representatives found in the fossil record.The mechanistic model by Franks et al.2 records, we calibrated a mechanistic leaf-gas exchange model27 in living representatives of the most ancient vascular plant taxa (Lycopodiaceae) constrained by their stomatal density, stomata pore size, and carbon isotope compositions, and applied it to fossil remains of Lower and Middle Devonian lycophytes with the same stomatal anatomy37. When trying to use the plant fossil record to estimate lower Palaeozoic CO2 levels at the time when vascular plants underwent their initial radiation, it becomes challenging to use paired sets of stomatal and isotope data. This situation arises due to the mode of fossil preservation, the need for destructive analysis for carbon isotopes, and the intrinsic scientific value of curated specimens. Therefore, we have further developed the mechanistic model27 to incorporate a solver routine that allows for the propagation of a wide range for each of the observable parameters required to deliver palaeo-pCO2 estimates. This approach has been developed as it allows for all variation between local environments and/or a temporal mismatch between data sources to be folded into the model predictions. Again, to ground truth that our solver approach delivers reasonable estimates of palaeo-pCO2, we have benchmarked the approach by applying the technique to our extant lycophyte dataset.To anchor the early atmospheric CO27 using two lycophyte species (Huperzia phlegmaria and H. squarrosa) presumed physiologically similar to the Devonian lycophytes and grown at known ambient CO2 levels for ~8 years in natural light within a soil substrate and under optimal Relative Humidity (RH) of ~80% in a greenhouse of the Botanical Garden in Copenhagen. Measurements of the leaf carbon isotope fractionation (Δleaf), stomatal density (SD), and stomatal pore length (p) (Supplementary table 2 assimilation rate at reference CO2 (A0) was indistinguishable from that of other modern lycophytes . Using this average A0 value as representative for lycophytes, and other parameters fixed, such as the operational stomatal conductance efficiency (ζ) with small effect on pCO2 estimates (see supplementary text for details), we were able to reproduce measured CO2 levels of 448 ± 51 ppm  in the greenhouse with a CO2 prediction of Initially, we calibrated the gas exchange model2 by developing a dataset based on non-paired data; i.e., where stomata and isotope data comes from H. phlegmaria and H. squarrosa from different growth conditions to mimic fossil samples that come from geographically/environmentally distinct sedimentary deposits of similar age. Based on cultivated lycophytes grown under dry and humid conditions, we find that unpaired data could carry up to ~100 ppm additional uncertainty, if stomatal and isotopic data were derived from plants that lived under vastly different conditions; i.e., stomatal data from slow-growing or even semi-epiphytic plants under extremely dry conditions with isotope data from plants grown under optimal, humid conditions or vice versa  or from different localities. Consequently, by applying the mechanistic model27 to lycophyte data, we are now able for the first time to resolve atmospheric pCO2 accurately and precisely prior to the emergence of forests.Following this initial screening we further explored the mechanistic model to derive estimates of CO2 levels at 410.8–382.7 Ma (Pragian through Givetian) were reconstructed from 66 fossil lycophytes representing three distinct genera  from 13 geological deposits at nine geographically distinct localities. From this, we derive four pCO2 estimates from paired data and six predictions from unpaired data  and there is a degree of similarity in Δleaf . Using paired data pCO2 predictions are constrained to between pCO2 estimates constrained to between 2 regardless of the method used to compile the raw data supporting the predictions  than the Δleaf (20.1 ± 0.9‰) of plants grown under optimal humidity  and a similar pore length (22.8 ± 1.0 µm vs. 21.6 ± 2.7 µm). Although, we do not know that these lycophytes had fully adapted to the drier greenhouse conditions, the pCO2 estimate derived from such plant material yields an under prediction of ambient glasshouse pCO2 level by -ζ) than plants grown under ideal  conditions .Two potentially confounding factors on the magnitude of leaf carbon isotope fractionation are humidity and atmospheric O2, at the lower end of the range of our new CO2 estimate. The model was run to better constrain conditions where the early lycophytes lived (see supplementary information for details) and shows a temperate planet with mean tropical surface air temperatures of 24.1–24.6 °C. The lycophyte floras in Australia, Germany, and China grew in the monsoonal belt with high mean annual rainfall (MAR) exceeding ~1000 mm/yr. But, Asteroxylon  and Drepanophycus floras from Maine, USA; New Brunswick and Gaspé, Canada are found in the slightly drier subtropical zones. Due to its coarse spatial resolution, our model cannot accurately capture local variations in rainfall and humidity at the fossil sites, but the model outputs do suggest growth in either humid or relatively humid environments. Importantly, we see no coupling between Δleaf or predicted pCO2 from each fossil localities and the modelled MAR, RH%, or palaeolatitude. Also, the Δleaf data is not suggestive of arid growth conditions. Given these data, there is no indication that the Devonian lycophytes lived under semi-arid/ arid conditions or the palaeo-pCO2 proxy reported here should deliver a biased prediction as a function of a down regulation in stomatal conductance due to growth in an arid environment.To assess this humidity effect further, we used an atmosphere-ocean palaeoclimate model of intermediate complexity (CLIMBER-3α) to evaluate the Devonian climate with 500 ppm of atmospheric CO2 can also affect the leaf carbon isotope fractionation because photosynthetic CO2 fixation competes with photosynthetic O2 fixation on the Rubisco enzyme39. This effect is more pronounced in modern plants grown at sub-ambient CO2 and super-ambient O2 levels40. The Devonian atmospheric pO2 levels were likely ~15 to 20 atm%12 and not much lower than today, so any O2 effect on our palaeo-pCO2 estimates will likely be small. Controlled growth experiments with vascular plants, including one species of lycophytes (Selaginella kraussiana), show a small positive shift in Δleaf (0.5 ± 0.4‰) when grown at sub-ambient O2 levels (16 atm%)40. This effect has been ascribed to O2-sensitivity of the photorespiratory compensation point41 and has a negligible effect on past atmospheric CO2 estimates compared to the O2-insensitive parameterisation of the mechanistic proxy  allows us to explore ecophysiological performance of early land plants that have no living relatives. This approach enables us to contextualize the very low stomatal density of these plants. Model results suggest that these plants had substantially lower CO2 assimilation rates than contemporaneous lycophytes (A0 < 3.7 μmol m−2 s−1); Aglaophyton (0.40 ± 0.10 μmol m−2 s−1), Rhynia (0.62 ± 0.13 μmol m−2 s−1), Horneophyton (0.84 ± 0.21 μmol m−2 s−1), Sawdonia (0.89 ± 0.04 μmol m−2 s−1) and Nothia (1.27 ± 0.29 μmol m−2 s−1) . Progymnosperm trees towering up to ~30 m diversified and became widespread from palaeo-equatorial to palaeo-boreal latitudes and from seasonally dry to wet habitats43. In contrast to the earlier lycophyte-dominated flora of typically ~0.1 m in height and with rhizoids penetrating only few cm into the ground, the evolving trees had both deeper roots for anchoring and a highly advanced root system akin to modern seed plants1. Still, atmospheric CO2 levels remained rather low and may have declined17 by, at most, a few hundred ppm during the Devonian emergence of forests.The atmospheric rve Fig.  illuminapCO2 decline is still poorly constrained, and there is a conflict between some low pCO2 estimates obtained from Mid-Ordovician marine phytane records and the high pCO2 estimates from pedogenic goethite18. If we trust that the Late Ordovician atmosphere had a higher pCO2 than today, at least outside glacial maxima  and a dramatic shift in the physical weathering regime driven by the evolving terrestrial ecosystems is recorded by the retention of fine grained sediment in continental deposits47. We propose that the earliest vascular vegetation promoted the exposure of more mineral surface area to weathering fluids and amplified global silicate weathering on the continents (far more than subsequent deep-rooted ecosystems could do) owing to a greater nutrient loss from less developed soils, and therefore, forcing a higher weathering demand.The absolute magnitude of the early Palaeozoic pCO2 and pO2 trajectories (see Methods). Our model predicts a massive atmospheric CO2 decline from ~2500 ppm to ~500 ppm in only ~30 Myr in response to enhanced silicate weathering by early vascular plants  outgassing rates is then ~1.0–1.2 times modern levels (in contrast to ~1.5 in previous models) characteristic of ages when supercontinents assembly.To simulate the effect of the establishment of early vascular plant ecosystems, we used a dynamic model (COPSE Reloaded) for the coupled biogeochemical cycles to predict atmospheric pO2 consistent with multiple lines of geochemical evidence for Earth’s oxygenation12 including wild fire evidence supporting atmospheric O2 levels above 15 atm% since ~420 Ma55 and fossil roots56 62. Polar sea ice and snowfall on Gondwana are predicted during winters , where the continent is predicted permanently ice-free. One exception might be the early Lockhovian Cooksonia-rich flora from the Paraná basin in Brazil positioned at high palaeolatitudes (~70°S)70. However, even at 500 ppm CO2, we find that the snow cover on Gondwana was not always perennial .Direct evidence for glaciation occurs in Late Devonian – Carboniferous deposits2 record is compatible with all palaeoclimatic constraints from the geological record, with global palaeoclimate models, and with global models for the coupled biogeochemical cycles that include a larger weathering demand of the earliest terrestrial flora and where early afforestation played only a minor role on global CO2 sequestration.In summary, the revised atmospheric COBaragwanathia flora in Victoria, Australia. Plant macrofossils from the fossil collection of Museums Victoria were originally collected by Isabel Cookson from Mt Pleasant Rd, Victoria, Australia71. The fossils are preserved as incrustations in fine-grained sandstone sandwiched in a 130 m thick stratigraphic section of mostly siltstone and shale  confines the flora to the Pragian or earliest Emsian, corresponding to ~409.1 ± 1.5 Ma according to GTS202073.Carbon isotopes, stomata density and stomata pore size from fossil leaves was compiled from the literature (see supplementary information) along with new data collected from the Baragwanathia longifolia and Zosterophyllum australianum, were selected for analyses based on the preservation of black organic matter contained within the brown mineralized fossils  which also contain larger fragments of 74. The Syracuse University nano-EA comprises an Elementar Isotope Cube elemental analyzer coupled to an Isoprime 100 continuous-flow stable isotope mass spectrometer via an Isoprime Trace Gas analyzer. Though the presence of carbonate was not suspected, nor observed via testing of sample powders with 6 N HCl under a binocular scope, we decided to decarbonate the sample materials prior to analysis to ensure that sample materials were carbonate free. Sample powders were fumigated in ashed glass vials in the presence of neat hydrochloric acid within an evacuated glass bell jar for 24 hours and then were dried in an oven at 40 °C. For isotopic analysis, sample materials were transferred to silver or quartz cups  and nested in a small amount of quartz wool to ensure retention of sample materials within the cup. The cups and quartz wool were ashed at 480 °C for 8 hours. Sample materials were loaded into cups within a Class 100 laminar flow isolation cabinet with HEPA filtered air to minimize the potential of particulate contamination.Carbon isotopic analyses for were performed in the Syracuse University GAPP Lab using an automated ‘nano-EA’ system adapted from that described in Polissar et al.2 pulse was set for 45 seconds. Carbon dioxide generated during sample combustion was trapped within the EA in a molecular sieve trap. Following passage of the N2 peak, the primary EA trap was heated and carbon dioxide was released to a secondary, silica gel-filled cryotrap which was immersed in liquid nitrogen. Trapping duration was calibrated using the EA thermal conductivity detector data to ensure complete collection of the CO2 peak. Following collection of CO2, the cryotrap gas flow was switched to a lower-flow He carrier gas (~1 mL/min) via an automated Vici Valco 6-port valve. The trap was warmed, and sample gas was released to the IRMS through an Agilent CarboBond capillary chromatography column (25 m x 0.53 mm x 5 µm). The resulting raw carbon isotope data are blank-corrected using direct blank subtraction and normalized to the VPDB scale using the two-point correction scheme75 with the international reference materials NIST 1547-Peach Leaves (δ13C = 26.0 ± 0.2‰) and IAEA C6-sucrose (δ13C = −10.45 ± 0.03‰) which are run as solids. The reference materials USGS 61- Caffeine (−35.05‰) and USGS 62-Caffeine (−14.79‰) are dissolved in UV-treated MilliQ water, dispensed in known quantities, and have long-term laboratory reproducibility of ±0.47‰ for USGS 62 and ±0.39‰ for USGS 61 over a 20–90 nanomole range. Reproducibility of the δ13C values of reference materials with carbon contents greater than 50 nanomoles is ±0.3‰ (1 sd) and is equivalent to that previously reported74.During isotopic analysis, sample cups were evacuated and purged with helium prior to introduction into the EA. Reaction conditions were as follows: oxidation and reduction reactor temperatures were 1100 °C and 650 °C, respectively; helium carrier gas flow was 158 ml/min and the OThe morphology and chemical composition of the fragments were characterized non-destructively by SEM and TOF-SIMS to confirm the presence of organic tissue well suited for carbon isotope analysis. Time-Of-Flight Secondary Ion Mass Spectrometry (TOF-SIMS) was used to produce semi-quantitative maps of the elemental composition in sample #15153 . This releases secondary ions that are accelerated in an electric field and their time of flight to the detector in vacuum is a function of their mass and sample depth2 level of 500 ppm. CLIMBER-3α encompasses a modified version of the ocean circulation model (MOM380) with a horizontal resolution of 3.75° x 3.75° and 24 vertical levels, a dynamic/thermodynamic sea-ice model81 with the same resolution and a fast atmospheric model82 of 22.5° longitudinal and 7.5° latitudinal resolution. The model does not explicitly model ice sheet growth on the continents, but snow cover on the continents is considered. The model was run for Lower Devonian boundary conditions (415 Ma) in terms of continental configuration, solar luminosity and vegetation cover83. Based on previous results67, three different orbital configurations were explored: the standard configuration  as well as cold  and warm  orbital configurations . This forward modelling approach enables hypothesis testing of mechanistic cause-effect relationships in the Earth system. A set of coupled differential equations describing the dynamic evolution of the C, O, P and S cycles were solved using an inbuilt variable timestep solver for ‘stiff’ Ordinary Differential Equation systems in Matlab®. The most recent version of the model , plant evolution (E), selective P weathering (F), volcanic outgassing (D) in a manner to simultaneously fit the effect of shallow vascular ecosystems on weathering processes and produce outputs consistent with the palaeorecords. Figure pO2 and pCO2 trajectories predicted by the revised version of CR with forcings shown in supplementary Fig. We used the Carbon-Oxygen-Phosphorous-Sulphur Evolution (COPSE) model to predict the histories of atmospheric Supplementary informationPeer Review FileDescription of Additional Supplementary FilesSupplementary Dataset 1Supplementary Dataset 2


This extraordinary enhancement in thermal conductivity is shown to be due to a network of continuous graphene sheets over long−length scales, resulting in low thermal contact resistance at bends/turns due to the graphene sheets being covalently bonded at such junctions. Solvent casting offers the advantage of preserving the porous structure of expanded graphite in the composite, resulting in the above highly thermally conductive interpenetrating network of graphene and polymer. Solvent casting also does not break down the expanded graphite particles due to minimal forces involved, allowing for efficient heat transfer over long−length scales, further enhancing overall composite thermal conductivity. Comparisons with a recently introduced effective medium model show a very high value of predicted particle–particle interfacial conductance, providing evidence for efficient interfacial thermal transport in expanded graphite composites. Field emission environmental scanning electron microscopy (FE−ESEM) is used to provide a detailed understanding of the interpenetrating graphene−polymer structure in the expanded graphite composite. These results open up novel avenues for achieving high thermal conductivity polymer composites.We demonstrate in this work that expanded graphite (EG) can lead to a very large enhancement in thermal conductivity of polyetherimide−graphene and epoxy−graphene nanocomposites prepared via solvent casting technique. A  One of the most promising approaches has been the percolative effect, where graphene–graphene contact is used to bypass the polymer, leading to significantly higher thermal conductivity enhancement [k values [−1−1 at just 10 wt% filler concentration. This value represents the highest k value measured in a polymer composite at this low weight fraction of filler. Presented results provide new avenues for achieving efficient thermal management in a wide range of applications.Increased thermal diss0ipation in modern electronic devices has led to a demand for thermally conductive materials with superior thermal conductivity (k) ,2. The lk) ,4. Howevm−1⋅K−1) ,6 of polm−1⋅K−1) ,11,12,13m−1⋅K−1) ,18,19,20ancement . More rek values ,25,26,27k values ,29,30, w−3) and exhibits high mechanical strength (10 MPa), thermal conductivity (25–470 W−1−1), and electrical conductivity (106–108 S−1) [Intercalation of acid molecules and oxidizing agents into graphite followed by rapid heating (900 °C) causes intercalated agents to gasify, producing enough pressure to push the adjacent graphite layers apart, resulting in expanded graphite a. Figure S⋅cm−1) . As a re S⋅cm−1) , phase−c S⋅cm−1) ,34, elec S⋅cm−1) ,36, elec S⋅cm−1) , fuel ce S⋅cm−1) ,39, batt S⋅cm−1) ,41 and s S⋅cm−1) ,43.−1−1 for EG particles. Tao et al. [−1−1 at 10 wt% EG composition. Zhao et al. [−1−1 in EG/paraffin wax composites at 25 wt% composition. Song et al. [−1−1 at 20 wt% composition in EG/MgCl2·6H2O composite. Wei et al. [−1−1 using ~25 vol% expanded graphite. Che et al. [−1−1 at ~20 wt% composition in high−density polyethylene (HDPE) composites.Significant research has been performed into the use of EG in enhancing the thermal conductivity of polymer graphene nanocomposites. Wu et al.  measuredo et al.  preparedo et al.  measuredg et al.  measuredi et al.  created e et al.  employedIn this work, we show that expanded graphite can lead to a remarkable increase in the thermal conductivity of polyetherimide composites prepared via the approach of solvent casting. Expanded graphite has a highly porous structure with interconnected graphitic walls, as seen in the SEM image of expanded graphite in the air in k~2000 W−1−1) [ink~2000 W−1−1, and outk~10 W−1−1) [The resulting interpenetrating polymer graphene structure from the solvent casting method allows for very efficient heat transfer. This can be seen by observing that graphitic nanosheets in this network are continuous for very long−length scales , as shown in m−1⋅K−1)  of graphm−1⋅K−1) .−8 K/W, which is comparable to graphene−polymer interfacial thermal resistance [k enhancement over percolation effect is that typically very high volume concentrations are required to achieve particle–particle contact for percolation. In the case of solvent−cast expanded graphite/polymer composites, however, the continuous graphitic networks are present at all volume fractions, enabling achieving high k even at low particle concentrations.A unique advantage of this continuous network can be understood by noticing that even at bends/turns, the graphitic sheets are still continuous, allowing for a very high thermal conductance  at such junctions. The above effect of low interfacial resistance at bends/turns in an expanded graphite network offers an advantage over the effect of percolation ,51,52,53The polymers chosen for this work, namely polyetherimide and epoxy, enable the advantage of superior thermal interaction with graphene ,56. BothWe further investigated the effect of sonication parameters on thermal conductivity enhancement and showed that lower sonication time and power result in larger expanded graphite particles in the composite, which allows for heat to be conducted efficiently along longer lengths scales, resulting in higher thermal conductivities. Finally, we also compare the measured thermal conductivity results with a recently introduced effective medium model that takes graphene–graphene contact into account for thermal conductivity prediction.Graphite intercalated compound (GIC) or expandable graphite (EPG), with an average particle size of ≥180 te store . Graphente store . The epote store . Commercte store  and Sigmte store . The orgThe fabrication procedure of the EG/PEID composite is illustrated in To prepare epoxy/expanded graphite composites, resin was added to 90 mL N,N−dimethylformamide (DMF) solution and stirred while heating at 150 °C to obtain a homogeneous mixture. Expanded graphite was added to this solution and stirred for one hour. The solution was then tip sonicated for 40 s followed by stirring at 150 °C until the solvent completely evaporated. After the solvent evaporated, a thick mixture of EG/epoxy was obtained. This mixture was then spread over a PTFE sheet and kept in a vacuum oven at 140 °C for 15 h to remove any residual solvent present in the mixture. On cooling the mixture to room temperature, a hardener was added to it and mixed to obtain a homogeneous viscous paste. This paste was then transferred to aluminum molds and cured at 90 °C for 20 h.k): k of EG/polymer composites was measured by the laser flash technique. A Netzsch LFA 467 Hyperflash  laser was used to measure the through−thickness thermal diffusivity of the samples. A total of 8–12 samples of 12.5 mm diameter and 0.3–0.4 mm thickness were used to measure the thermal diffusivity (α) at room temperature (23 °C). The samples were coated with graphite spray before the measurement to efficiently absorb heat from a flash lamp, and an average of 6–8 measurements was reported. This laser flash technique induces heat by a laser pulse on one surface of the sample, and the temperature rise is captured on the other surface of the sample as a function of time. α is determined by LFA using the following equation: α = (0.1388 d2)/t1/2, where α is the thermal diffusivity (mm2/s), t1/2 is the time to obtain half of the maximum temperature on the rear surface, and d denotes the sample thickness (mm). The thermal conductivity was calculated using k = αk, ρ, and Cp represent the thermal conductivity, density, and specific heat constant of the sample, respectively. In this work, the density and specific heat of the composite samples were calculated using a gas pycnometer  and differential scanning calorimetry (DSC) .Thermal Conductivity (Scanning Electron Microscopy (SEM): Morphological characterization of EG filler and EG/polymer composites was carried out by high−resolution field emission environmental scanning microscopy . This SEM was operated in secondary electron (SE) mode at an accelerating voltage of 20 kV. To prepare the samples for SEM imaging, liquid nitrogen was used to crack down the composite to image over the cross−sectional area. Montage large area mapping of EG fillers has been captured using MAPS software of FE−ESEM.Raman Spectroscopy (RS): Raman spectroscopy (RS) was performed using a DXR3 SmartRaman Spectrometer  to collect the data over the range from 3250 to 250 cm−1, laser wavelength λL = 633 nm, spectral resolution = 0.16 cm−1, and imaging resolution = 702 nm for the EG and GIC samples. An Olympus BX 41 microscope with 5× objective, 10 s exposure time for 15 accumulations, and 3 scans per sample were used to collect the spectra.X−ray Diffraction (XRD): A PANalytical Empyrean Diffractometer  produced the information regarding the crystal structure of EG compared to GIC using Bragg−Brentano focusing geometry at room temperature. A total of 3 kW Cu Kα radiation (λ = 1.5406 Å) with a scan range of 2θ = 5 to 80°and step size of 0.013°.X−ray Photoelectron Spectroscopy (XPS): X−ray Photoelectron Spectroscopy (XPS) was performed for GIC and EG samples by Thermo Scientific K−alpha XPS . A Kα gun source was used to excite the sample, and the data were collected for an acquisition time of ~70 s at 400 µm spot size. The passing energy of 200 eV was used to find the carbon (C), oxygen(O), and sulfur (S) peak in this analysis spectrum. The atomic percentage of C, O, and S were investigated using the Avantage software. To determine the atomic percentage, this software was used to perform the curve fitting in accordance with Gaussian and Lorentzian functions.−1−1 at 10 wt% filler composition. This value represents a remarkable 2770% enhancement over the k (0.23 W−1−1) of pristine PEID, providing new avenues for high thermal conductivity composites. Similarly, the measured thermal conductivity of EG/epoxy composites also shows remarkable enhancement at very low loadings of EG. At just 2 wt% EG composition, a thermal conductivity value of 1.74 W−1−1 is achieved for the epoxy composite, representing an enhancement of 1025% over pristine epoxy (0.16 W−1−1). The measured thermal conductivities of EG/PEID and EG/epoxy composites are shown in k value is achieved even at low graphene loading, as opposed to the case of percolation, where typically much higher particle concentrations are required to achieve particle–particle contact. Kargar et al. achieved around 6 W−1−1 in graphene/epoxy composites at a high graphene loading of 35 vol% through the percolation effect [−1−1 in Boron–Nitride/epoxy composites [k of 6.6 W−1−1 is achieved in this work through the use of just 10 wt% EG content. Even in percolative k enhancement, the thermal interfacial resistance at particle–particle contact can be relatively high. Konatham et al. reported a thermal boundary resistance of 5.5 × 10−8 K/W, almost as high as at the graphene−polymer contact [k value of EG/PEID composite has been achieved due to such continuous graphitic network of EG as presented in We first compare the measured values against previous results reported for the percolation effect based on establishing a graphene–graphene network. A key advantage of expanded graphite is that a large enhancement in n effect . Percolamposites  at a vol contact . Expandek achieved in this work is also due to the use of solvent casting, which offers the advantage of preserving the porous structure of expanded graphite within the composite  is higher by 572% relative to GnP/PEID composite (1.0 W−1−1). Similarly, at 2 wt% composition, k of EG/epoxy composite (1.74 W−1−1)) is higher by ~600% relative to GnP/epoxy composite (0.25 W−1−1). At the low graphene content used in this work, GnPs are well separated by polymer; high interfacial thermal resistance between GnPs and polymer then restricts the enhancement achievable through the use of graphene nanoplatelets. These results highlight the large advantage of expanded graphite over graphene nanoplatelets in enhancing composite thermal conductivity.We also compare measured k enhancement. We further investigated the effect of sonication parameters used during composite preparation on While the above images show the expanded graphite embedded in the polymer, we also show expanded graphite before it is embedded into the polymer in k values of EG/PEID composite with theoretical predictions based on a recently introduced effective medium model by Su et al. [In the next section, we compare measured u et al. . CompariThe effective medium model introduced by Su et al.  includes0 and 0k is the thermal conductivity of an interlayer surrounding graphene sheets. The role of this interlayer is to include the effect of interfacial thermal resistance at graphene−polymer and graphene–graphene contacts.In the above equations, The effective thermal conductivities In the above equations, −1−1 from our measurements (in suitable agreement with the literature). The in—plane . This interlayer thermal conductivity of 0.17 W−1−1 at graphene–graphene contact corresponds to a graphene–graphene interfacial thermal resistance of 1.17 × 10−8 m2K/W. This value is significantly smaller than the interfacial resistance predicted for regular graphene–graphene contact [−8 m2K/W), providing evidence that graphene sheets are in superior contact  at bends/turns in an expanded graphite network, compared to the contact between discrete graphene particles. The analysis points to the advantage of the continuous graphitic network achieved through the use of expanded graphite via solvent casting technique in enhancing the thermal conductivity of polymer composites. We further show the effect of lower graphene–graphene contact conductance by decreasing the graphene–graphene contact thermal conductivity in the model. It is seen that as the graphene–graphene contact thermal conductivity is decreased from 0.17 to 0.10, the effective composite thermal conductivity is decreased from 6.6 W−1−1 to 4.4 W−1−1 , indicating the importance of superior graphene–graphene contact conductance for achieving overall high composite k values. We next discuss the characterization of EG filler and EG/PEID composites.A key parameter in the above model is the interfacial resistance at graphene–graphene contact. This value was determined by fitting the measured values to the effective medium model. A suitable agreement between measured thermal conductivities and predicted values was obtained by using a graphene–graphene interlayer thermal conductivity of 0.17 W⋅K−1 see , which i contact   images enable understanding of morphological differences between graphene nanoplatelets (GnPs), Graphene Intercalation compound (GIC), and thermally expanded graphite (EG) a. ThermaX-ray diffraction (XRD) analysis was performed to determine the crystal structure and interlayer spacing of GIC and EG. −1 and ~1350 cm−1 for graphitic material [2 carbon of hexagonal ring due to in−plane tangential stretching of the carbon−carbon bonds [3 hybridized carbon [−1 and represents a second−order two−photon process [D/IG ratio is typically used to characterize the defective state of graphene. The absence of a D band in EG suggests the presence of a highly ordered defect−free graphite structure in EG. This high degree of the ordered structure of graphite, even after thermal expansion, has a strong beneficial impact on the thermal conductivity enhancement of polymer graphene composite as it preserves the intrinsic high thermal conductivity of graphene itself.Nondestructive Raman analysis was also performed to further analyze the crystal structure before and after the thermal expansion. Raman spectroscopy of material . The G bon bonds , and thed carbon . The 2D  process . Raman sXPS analysis was further performed to investigate the concentration of carbon (C), oxygen (O), and sulfur (S) elements before and after the thermal expansion of GIC, as presented in −1−1 is measured, representing an enhancement of 2770% over pristine polyetherimide. This large enhancement in thermal conductivity is found to be due to a network of continuous graphene sheets over long−length scales achieved through the solvent casting technique, which preserves the interconnected porous structure of expanded graphite within the composite. Even at bends/turns in graphene sheets in such a network, the sheets are covalently bonded, which minimizes the interfacial thermal resistance at junctions, enhancing heat transfer. Overall, the resulting structure allows highly efficient heat conduction over long−length scales along the continuous graphitic sheets. This results in the observed high thermal conductivity of the composite through the use of EG. The thermal conductivity of EG/PEID composite is also found to dramatically exceed that of graphene nanoplatelet (GnP)/PEID composites by 572% for 10 wt% filler composition. At low filler loading, GnPs are well separated by polymer and the resulting high graphene−polymer interfacial thermal resistance results in low effective GnP/PEID thermal conductivity. Presented results highlight the advantage of expanded graphite in enhancing the thermal conductivity of polymer composites and can lead to novel avenues for achieving efficient thermal management in a wide array of technologies.In summary, we demonstrate that expanded graphite (EG) can lead to a large enhancement in the thermal conductivity of EG/PEI composites prepared through solvent casting. At 10 wt% EG composition, a high thermal conductivity of 6.6 W


Here we demonstrate the formation of dipolar exciton polaritons in bilayer MoS2 resulting in unprecedented nonlinear interaction strengths. A ten-fold increase in nonlinearity is observed for the interlayer dipolar excitons compared to the conventional A excitons. These highly nonlinear dipolar polaritons will likely be a frontrunner in the quest for solid-state quantum nonlinear devices.Realizing nonlinear optical response in the low photon density limit in solid-state systems has been a long-standing challenge. Semiconductor microcavities in the strong coupling regime hosting exciton-polaritons have emerged as attractive candidates in this context. However, the weak interaction between these quasiparticles has been a hurdle in this quest. Dipolar excitons provide an attractive strategy to overcome this limitation but are often hindered by their weak oscillator strength. The interlayer dipolar excitons in naturally occurring homobilayer MoS 2 resulting in highly nonlinear dipolar polaritons.Dipolar excitons enable large nonlinear interaction but are usually hampered by their weak oscillator strength. Here, the authors demonstrate the strong light-matter coupling of interlayer dipolar excitons having unusually large oscillator strength in bilayer MoS To further extend the span of their utilities, platforms to achieve and implement strongly interacting photons in solid-state systems are highly desirable3. Conventional materials do not exhibit nonlinear response at power levels associated with single photons. In this context, remarkable advances have been made in cold atom systems to realize interacting photons at the single-particle level4. A strong contender for the generation of strongly interacting photons in the solid-state are exciton-polaritons formed via non-perturbative coupling of cavity photons with excitonic resonances. Although they can be modeled as non-interacting quantum fluids at low densities, beyond a critical density, saturation and short-range exchange interactions become significant and give rise to various phenomena like the appearance of spontaneous coherence, parametric down-conversion, and superfluidity5. Strong spatial confinement along with Coulomb interactions give rise to even stronger correlations between polaritons that can no longer be described using a mean-field theory. Such interactions give rise to non-Poissonian statistics of laser transmission, dubbed as polariton blockade. Preliminary evidence of such non-classical correlation was recently observed in confined polariton systems in GaAs7 but a small ratio of interaction to dissipation rates resulted in only a weak violation of classical correlations.Photons are becoming ubiquitous in emerging quantum technologies like quantum communication and metrology due to the ability to propagate long distances while being robust against decoherence8. Furthermore, their two-dimensional (2D) nature allows for the formation of heterojunctions or homojunctions with arbitrary twist angles resulting in emergent properties12. Their strong binding energy and large oscillator strength render them capable of forming polaritons at room temperature13, which has also been shown to retain the intriguing properties of the 2D excitons16. Polariton interactions in TMDCs under different configurations are being studied extensively. Realization of Fermi polarons17, trion polaritons19, and excited-state Rydberg exciton-polaritons20 have shown great potential in their abilities to harness strong polariton interactions. Moiré exciton-polaritons were recently demonstrated using heterobilayers of WS2/MoSe2 where the electronic confinement potential arising from the twisted heterostructure was shown to enhance the nonlinearity9. However, all the platforms above rely on short-range exchange interactions or phase-space filling that pose a bottleneck in realizing few polariton non-linearity under current experimental capabilities.Transition metal dichalcogenides (TMDCs) have garnered much attention for their exceptional optoelectronic properties and have been used to demonstrate a wide array of fundamental phenomena and technological applications21 that can have a permanent dipole moment and hence support highly interacting dipolar polaritons. Formation of polaritons with excitons that possess a permanent dipole moment has been shown to enhance the polariton interactions both in resonant23 and non-resonant excitation schemes24. However, their inherently low oscillator strength creates an impediment in reaching the strong coupling regime without hybridization with a direct intralayer exciton. Furthermore, the large in-built interfacial electric fields make electrical tuning of the energies of the IE far more difficult. Bilayer MoS2 provides a highly attractive platform to solve these issues28. The IE in naturally stacked 2H bilayer MoS2 remarkably has both an out-of-plane (static) dipole moment and an in-plane (oscillating) dipole moment. Due to the in-plane dipole moment, the IE in MoS2 bilayer has an oscillator strength of approximately 36% of that of the intralayer A exciton along with strong absorption that is visible up to room temperatures. Moreover, a strong response to DC electric fields has been demonstrated in these systems, thus confirming their dipolar nature28. In this work, we achieve strong coupling of microcavity photons with the IEs  in bilayer MoS2. The IE polariton shows 10 fold enhancement of the polariton nonlinearity compared to the intralayer A exciton owing to its dipolar and phase space filling contributions. Such enhanced nonlinear response makes them appealing to realize strongly interacting polaritons in condensed matter systems. In addition, the ease of fabrication and realization of multi-polariton species in this system makes it a practical and fundamentally interesting material for studying polariton physics.In this context, a very attractive possibility is the use of spatially indirect interlayer excitons (IE) in TMDC heterostructures2. IE1 and IE2 excitons in which electrons are in layer 1, and layer 2 respectively, are energetically degenerate at zero external bias while the hole is delocalized among both layers. The black arrows in Fig. 1 and IE2. Owing to the spatial separation of the electron and the hole, it acquires a permanent dipole moment, the nature of which has been a topic of much interest recently28. IEs in MoS2 homobilayer can be thought of as an admixture of B intralayer exciton with an optically dark but electric field tunable charge-transfer exciton, which is typically found in TMDC heterobilayers29. As a result, IEs in MoS2 homobilayer acquire both a strong oscillator strength and electric field tunability29. Figure 2 bilayer flakes after mechanical exfoliation . The relative shift of A exciton and IE so obtained are plotted in Fig. 2 by embedding it in a microcavity. Figure 2 encapsulated by thin (20 nm) hexagonal boron nitride (hBN), which is sandwiched between a bottom dielectric distributed Bragg reflector (DBR) mirror and a top silver mirror. See SI Note 2 flake sits close to an electric field anti-node, allowing us to observe five dispersive modes associated with the different polaritonic states as shown in Fig. EA1s = 1.9323 ± 0.0003 eV, EIE = 2.0014 ± 0.0001 eV, EA2s = 2.078 ± 0.002 eV, EB = 2.111 ± 0.001 eV and the Rabi splittings ΩA1s = 40.4 ± 0.3 meV, ΩIL = 21.4 ± 0.1 meV, ΩA2s = 13 ± 0.5 meV, ΩB = 51 ± 0.4 meV. The exciton energies obtained through the fit agree well with the experimentally determined exciton energies with slight shifts due to the strain and changes in the dielectric environment in the cavity. In our device, the bare cavity mode is positively detuned from the A1s exciton by ∥ where the cavity mode is resonant with the A exciton (yellow), IE (red) and B exciton (blue), respectively. We notice that the Rabi splitting of the IE polariton remains observable even at 77 K . Two Lorentzian fits at each density is used to obtain the energy of the upper and lower polaritons. SI Note ∣∣ along with the reduction of Rabi splitting. This suggests the presence of both exciton energy renormalization and saturation effect. SI Note LP with respect to density, we calculate the strength of the nonlinearity, gLP. Since the rate of blue shift saturates at high density  and phase space filling effect. For intralayer excitons with a small Bohr radius only the exchange term and phase space filling effect is importantexc-exc) versus phase space filling (ΔEsat) we carry out the following analysis. Figure k∥. This is used to obtain the shift in polariton energy, C), see Fig. gLP = 4gSAT∣C∣∣X∣3, where X is the Hopfield coefficient for the exciton. Although this is a multi-exciton system, the contribution of the other excitons close to the zero-detuning k∥ of the IE is small. Thus, we can write 33 which monotonically decreases with increasing ∣C∣  and TiO2 (64.7 nm) layers and were deposited by radio frequency sputtering on an intrinsic Si chip. MoS2 and hBN were exfoliated from bulk crystals (from 2Dsemiconductor Inc.) using blue tape (Nitto) and scotch tape respectively. MoS2 was exfoliated onto PDMS substrate and hBN was exfoliated onto a 300 nm SiO2/Si substrate. Bilayer MoS2 flakes were identified by reflection spectroscopy – the additional dip at ~639 nm in the reflection signal at room temperature is the hallmark of bilayer MoS2. hBN/bilayer MoS2/hBN heterostructure stacking and transfer were done using the well-known polypropylene carbonate transfer technique36. The chosen top and bottom hBN layers were of similar thickness ~20 nm. The final stack was then transferred onto the DBR at a temperature of 150 °C. The chip was kept in chloroform for 12 h to remove the polypropylene carbonate residue. 240 nm polymethylmethacrylate (PMMA) (495 A4 from Michrochem) was spin-coated to form a 3λ/2 cavity. Finally, a silver layer (40 nm) was deposited by electron-beam evaporation for the top mirror of the microcavity. Optical images of the device and the bare cavity response are given in SI Note The DBR used in our experiment is made of 8 pairs of SiO2 on a uniform area of the sample of dimension ~500 μm2.We recorded the angle-resolved reflection spectra using the Fourier space imaging technique. A broadband halogen light source was used for reflection measurements. A supercontinuum pulsed light source  was used to study the power dependence of the bare excitons and polaritons. Appropriate long pass and short pass filters were used in the input to excite a narrow band around the excitons/polaritons under study. For the IE polariton, the excitation bandwidth of the supercontinuum laser was chosen to be ~40 meV around the IE energy in the power-dependent measurements. The set-up was coupled with a Princeton Instruments monochromator with a CCD camera. A 50X, 0.65 numerical aperture objective was used for all the measurements at 7 K in a closed cycle Montana cryostat. The polariton dispersion was revealed by imaging the back aperture of the microscope objective (Fourier plane) onto the camera. The spot size of the laser was 1 μmFor the data presented in Fig. 2We first calculate the absorbed power density by the bilayer MoSPlaser(E) is the intensity profile of the incident supercontinuum laser, L(E − EX) is the absorbance of X exciton around its energy EX, and Abeam is the size of the focused laser beam on the sample. We run the integration within 74 meV (53 meV) band around the A exciton (IE) peak energy, respectively. Beyond these bands, the reflection contrast of A exciton and IE are not detectable. The absorbed photon density from the absorbed power density is calculated taking into account the life time of these excitons (45 fs for A exciton 53 fs for IE). We assume all of the absorbed photons form excitons.Here Transient reflection spectroscopy measurements presented in SI Note 37. However, the relative enhancement of the nonlinearity of IE with respect to A remains unaffected due to the above errors as they exist for both the IE and A exciton polariton.Currently, the density error bars consider the error in measuring power, determining the energy and linewidth of the polariton. However, one primary source of the error is the lifetime of the polaritons. To address this issue we have done pump-probe experiment which validates the use of the linewidth of the polaritons to determine the polariton lifetime. Results of our pump-probe experiment is included in the SI note Supplementary Information