Traditional printed electronics processes have recently been utilized within 3D-printed structures where components and interconnects are introduced during manufacturing disruptions. The dielectric performance of 3D-printed materials has a low-resolution problem, and many technologies have been proposed for direct printing on a 3D curved surface or structure. This paper reports a humidity sensor fabricated with a 3D-printed electrode and cellulose nanofibers on a curved surface. The electrode part of an interdigital electrode (IDE) sensor is printed on a flat glass substrate and a 3D-curved glass substrate using a double blanket reverse offset. Subsequently, a cellulose nanofiber emulsion is coated onto the IDE pattern as a sensing layer with a dispenser. The electrical impedance of the sensor is measured with the relative humidity (RH) changes between 10% and 90% RH. The sensor demonstrates a high repeatability and sensitivity, even on a 3D curved substrate. This technology provides a promising method to integrate humidity sensors and 3D deformable surfaces. Humidity sensors are used in various fields, including environmental monitoring, automobiles, and medical care. Different types of humidity sensors that operate on different transduction principles exist, such as resistive, capacitive, optical, and surface elastic wave types ,2,3,4. SIn general, rigid substrates, such as ceramics, glass, and silicone, are used to fabricate sensors. Flexible and deformable printed sensors have become popular since they offer a cost-effective and versatile alternative to traditional rigid sensors. The ability to print sensor devices onto various shapes and substrates, including plastics, textiles, and paper, has opened new avenues for innovation in humidity sensors. Recently, research has been conducted on insertable devices that can be mounted on light, flexible, and thin substrates using printing . ChemoseFor resistance-type sensors, extensive research is being conducted on the printing process, in which humidity sensors are manufactured using various printing techniques on diverse substrates, including rigid substrates. Most electronic devices and sensors are made on 2D surfaces, robust and brittle substrates suitable for various applications in diverse fields. However, they are incompatible with complex curved surfaces, such as soft 3D living organisms in the health care/monitoring areas or irregular curved wings required for equilateral integration . Thus, wThe measurement accuracy and sensitivity of existing sensors are limited at high temperatures or humidity. As a result, it is very important to implement a humidity sensor with high sensitivity, flexibility, and a wide response range for integrating it with existing sensors. In particular, a specific range can be set, and a humidity sensor can be manufactured with patterned electrodes on a curved surface using various printing methods. In printing humidity sensors, fine-line printing capabilities are required.In this study, the proposed double blanket reverse offset (DB-RO) printing can print fine patterns with a line width of several micrometers on a 3D substrate. DB-RO printing can realize fine line widths without causing damage to the substrate using a soft blanket. It has the advantages of a high resolution and the precision of traditional lithography, which can be a key in fabricating precise and deformable curved sensors. Therefore, the humidity sensor is manufactured in this study using DB-RO on a curved surface, demonstrating a soft blanket hardness, enabling 3D printing.Printing several materials with a single process remains limited and is a major issue associated with producing a curved humidity sensor using an inkjet and dispenser, which are conventional noncontact printing methods. It is not easy to realize a complete sensor device using these methods. When a sensor device is manufactured directly on a 3D curved surface, the structure control range of the device is widened, which enables the device fabrication with high functionality. Curved lithography can be applied to all surfaces with a high pattern resolution, but the challenges are deposition, etching, or material growth/doping in a wide area of the curved surfaces ,15. ThusIdeal humidity-sensor materials require high sensitivities, wide humidity-sensing ranges, fast response times, short recovery times, and inexpensive manufacturing methods . Herein,This study aims to manufacture humidity sensors on a 3D curved substrate through a printing process equipped with a double blanket. An interdigital electrode (IDE) pattern with a line width of 40 µm and a width of 40 µm is fabricated using DB-RO equipment to obtain signals from the humidity sensors. A CNF sensing layer is deposited on the IDE pattern by dispensing CNF reactant. The humidity-sensing behavior of the sensor fabricated using the simple and direct printing process on the 3D curved substrate is investigated. Silver conductive ink  was purchased from ANP in Chungcheongbuk-do, Republic of Korea. A PDMS (polydimethylsiloxane)-based blanket (#CF 0.70) was purchased from Fujikura Composites, Tokyo, Japan. A glass substrate was used as the substrate, and a flat glass with a radius of curvature of 10° was used. The CNF suspension was isolated from a hardwood pulp  by using a 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO) oxidation and aqueous counter collision (ACC) method . All cheThe humidity sensor comprises an IDE pattern fabricated on a 3D curved substrate and coating a CNF sensing layer. The sensing mechanism of the humidity sensor is based on resistance conversion, and the electrical conductivity of the sensing layer increases due to the absorption/desorption of water molecules by the sensitive CNF layer. CNFs have many hydroxyl groups on their surface, and two water molecules exist: adsorbed water and free water molecules. Adsorbed water molecules bind CNFs tightly by making hydrogen bonds, forming a solid CNF layer. On the other hand, free water molecules can be easily absorbed/desorbed to CNFs depending on the environment humidity. Depending on the amount of adsorbed free water molecules, the electrical resistance of the CNF layer can be changed, giving it its humidity-sensing behavior.f double blanket pressure  at an accelerating voltage of 15 kV after coating 3 nm of platinum. Ag IDEs were characterized via the SEM and a 3D profiler system . The manufactured humidity sensor was obtained via real-time monitoring of the temperature and relative humidity (RH) at room temperature inside an environment chamber . The humidity-sensing response was recorded as a change in the resistance with the RH. An impedance analyzer  powered by a custom-designed Labview-based interface monitored the resistance across the IDE in real time . The inp2 dimension. Ag IDEs can be formed rapidly and easily using a printing system. The total sensing area is 15 mm × 10 mm. The large-scale fabrication of the humidity-sensor array was successfully achieved using the proposed process. The CNF-layer thickness can be controlled by the dispenser’s air pressure. The humidity sensing of the fabricated sensor was tested at different RH values ranging between 10 and 90% RH. The sensitivity, repeatability, hysteresis, and response time were evaluated.The humidity sensitivity of the sensor fabricated on the curved substrate was compared with the flat substrate sensor, and The patterned sensor can easily be transferred to arbitrarily shaped substrates without causing any damage. The CNF sensing layer on the IDE, after being transferred to a curved surface, is shown in Stability is an important factor for real applications of humidity sensors. The RH was maintained, and the sensor response was monitored to check the stability. This study examined the challenges of fabricating curved humidity sensors with a 3D-printed electrode and cellulose nanofibers. The interdigital electrode (IDE) part of the 3D humidity sensor was manufactured using the DB-RO printing process with silver conductive ink. The printing process realized a fine line width with a high resolution. A CNF sensing layer was coated on the IDE to form the 3D humidity sensor. The same humidity sensor was fabricated on a flat substrate and compared with the curved humidity sensor. The CNF-based curved humidity sensor detected between 10% and 90% RH levels with low hysteresis.Additionally, the curved sensor exhibited a very fast response and recovery times of 61.5 s and 31.5 s, respectively. The curved humidity sensor showed lower hysteresis and a shorter response time than the flat one. The sensor demonstrates a high repeatability and sensitivity, even on a 3D curved substrate. The proposed method provides a promising strategy for fabricating electronic and sensor devices on various substrates, including curved ones.


K-edge confirm the presence of a magnetic signal in the graphene layer,and the sum rules give a magnetic moment of up to ∼0.47 μB/C atom induced in the graphene layer. For a more preciseestimation, we conducted PNR measurements. The PNR results indicatean induced magnetic moment of ∼0.41 μB/C atomat 10 K for epitaxial and rotated-domain graphene. Additional PNRmeasurements on graphene grown on a nonmagnetic Ni9Mo1 substrate, where no magnetic moment in graphene is measured,suggest that the origin of the induced magnetic moment is due to theopening of the graphene’s Dirac cone as a result of the strongC pz-Ni 3d hybridization.We report the magnitude of the induced magnetic momentin CVD-grownepitaxial and rotated-domain graphene in proximity with a ferromagneticNi film, using polarized neutron reflectivity (PNR) and X-ray magneticcircular dichroism (XMCD). The XMCD spectra at the C  This is a result of the overlap of the graphene’svalence band with that of the metal substrate due to the breakingof degeneracy around ED in a partiallyfilled d-metal, as discussed in the universal model proposed by Voloshinaand Dedkov.21 Their model was supportedby density functional theory calculations and proven experimentallyusing angle-resolved photoemission spectroscopy.26 Furthermore, a small magnetic signal was detected in the X-ray magneticcircular dichroism (XMCD) spectra of the graphene layer in proximitywith a FM transition metal (TM) film, suggesting that a magnetic momentis induced in the graphene.28 However, no direct quantitativeanalysis of the total induced magnetic moment has been reported.Grapheneis a zero-gap semiconductor because the π and π*bands meet at the Fermi energy  substrate was used asthe FM because it has a small lattice mismatch of −1.2%, abond length of 2.03 Å, and their d orbitals are positioned ∼1.1eV below EF .29 Epitaxial and rotated-domaingraphene structures were investigated because rotated graphene isexpected to interact more weakly with the TM film underneath. Thisis a result of the loss of epitaxial relationship and a lower chargetransfer from the TM due to the missing direct Nitop–Cinteraction and the smaller region covered by an extended graphenelayer as a result of the 3° rotation between the graphene andNi.32 Therefore, a smallermagnetic moment is expected to be induced in rotated-domain graphene.It is widely accepted that graphene’s C atoms are assembledon top of close-packed (111) surfaces in what is known as the top-fccconfiguration, where the C atoms are placed on top of the atoms ofthe first and third layers of the TM substrate.We have studied the structural, magnetic, and electronic propertiesof epitaxial- and rotated-domain graphene grown on Ni films, confirmedthe presence of a magnetic moment in graphene by element-specificXMCD, and measured the induced magnetic moment at 10 and 300 K forrotated and epitaxial graphene using polarized neutron reflectivity(PNR).K-edge signal due to thecontribution of the carbon contamination in the beamline optics. Weattribute the presence of an induced magnetic moment in graphene tothe hybridization of the C pz orbitalwith the 3d bands of the TM, which is supported by additional PNRmeasurements on graphene grown on a nonmagnetic Ni9Mo1(111) substrate, where no magnetic moment is detected in thegraphene layer.To our knowledge, our attempt is the first reported approachinusing PNR and XMCD sum rules combined to estimate the total inducedmagnetic moment in graphene. This is due to the thinness of graphene,which is close to the resolution of the PNR technique, and the difficultyin processing the XMCD C 34 For thesemeasurements, the graphene layer was first transferred by a chemicaletching process from the metallic film onto a Si substrate with athermally oxidized SiO2 layer similar to that reportedin ref . The variationin the spectra of each sample, the presence of second-order and defect-inducedpeaks, the large FWHM of the 2D peak, and the high I2D/IG ratiocould be a result of the chemical etching and transfer process   as a result ofthe occasional overlap of the graphene domains.38 The Raman spectra for the graphene/Ni9Mo1 sample, as well as the full list of the peak positions andthe 2D average FWHM of all of the measured samples, are provided inthe Supporting Information. seethe  or the c see the .The broXAS) and the XMCD response (μXMCD) at the Ni L2,3-edge in therotated graphene/Ni sample are shown in 40 The region between the L3 and L2 edges with a constant negative intensity isknown as the diffuse magnetism region, μdiff, andit has been observed and reported for the Co, Ni, and Fe XMCD spectra.42 μdiff is expected to arise as a result of the oppositespin directions for the 4s and 3d electrons, interstitial and sp-projectedmagnetic moments, and the fact that it couples antiferromagneticallyto the sample’s total magnetic moment in 3d elements.41 Although μdiff has been reportedto contribute to about −7% to the total magnetic moment inNi,43 because the sum rule does not account forμdiff, the integration range over the L3 was stopped just before μdiff for thecalculation of the orbital magnetic moment, mo, and the spin magnetic moment, ms (858.7 eV). However, the main L3 peakand the shoulder, μshoulder, are due to multipleinitial-state configurations, 3d8 and 3d9, respectively,and therefore they were accounted for in the sum rule calculations.41The X-rayabsorption spectra . Thistailhas been excluded from the sum rules as it is considered part of thenonresonant contribution. The calculated mo and ms values are 0.084 and 0.625 μB/Ni atom, respectively, and thus mtotal is 0.709 μB/Ni atom . Considering the20% accuracy of the XMCD technique in estimating the magnetic momentsof materials, the results obtained for Ni are consistent with thevalues reported in the literature.45Although the nonresonant contribution was subtractedfrom the μK-edge spectra for the rotated graphene/Nisample are shown in K-edge, higher sourcesof errors are expected in the XMCD estimation attributed to the difficultyof applying the sum rules to the C K-edge spectrain comparison with that for the Ni L2,3-edge. For instance, various studies have been reported for Ni,48 which can be used asreferences for our measurements, but the application of the sum ruleshas not been reported for graphene before. Also, the number of holes, nh, has not been measured for C previously. Moreover,the gyromagnetic factor (g) of the graphene was foundto be different depending on the underlying substrate,51 and it has not been reported for graphene on Ni. It is also noteworthyto mention the difficulty associated with measuring the C K-edge due to the C contamination of the optical elements,which appear as a significant reduction in the incoming intensityat this particular energy.TheC XMCD|, which is shown in z–Ni 3d hybridization,14 a smallmagnetic signal is observed at the 1s → σ* transitionpeak as well; a similar behavior was reported for graphene/2 ML Co/Ir(111).25 The calculated upper bound mo for graphene is 0.062 μB/C atom using nh = 4, which corresponds to an ms of 0.412 μB/C atom, using g = 2.3, which is the value reported for graphene grown on SiC.49 Therefore, the upper limit of mtotal of the rotated-domain graphene grown on Ni is ∼0.474μB/C atom .Nonetheless, we can obtain an upperlimit to the orbital momentof the graphene layer by integrating the modulus of the dichroic signal,|μDespite the large uncertaintiesexpected for the estimated graphenemoments, the XMCD results demonstrate the presence of magnetic polarizationin graphene.Next, we turn to PNR to combine all of the informationfrom complementarytechniques, XMCD, Raman, and SEM, to obtain a more precise value forthe magnetic moment in the graphene.PNR experiments werecarried out to measure the magnetic properties of each layer of thesamples individually and to determine the value of the induced magneticmoment in graphene quantitatively.SA = [R+ – R–]/[R+ + R–], where R+ and R– are the spin-up and spin-down neutron specular reflectivities,respectively). SA scales with the magnetic signal.A flat SA line at zero, shown as a blue dashed line,represents no net magnetic induction present in the system. Panel(d) displays the nuclear scattering length density (nSLD) for thesample structure. This structure is shared at both temperatures inthe co-refinement, and panel (e) shows the magnetic scattering lengthdensity (mSLD) for each temperature.The PNR results for the rotated-domaingraphene/Ni(111) and epitaxial graphene/Ni(111) samples, measuredat 10 and 300 K, are displayed in Supporting Information, which contains the variousmodelsused to fit the rotated graphene/Ni sample. The 10 and 300 K datashown in Supporting Information shows the importance ofprior knowledge of the sample properties to obtain the best PNR fitto estimate the induced magnetic moment in graphene. We used the informationobtained from the structural characterizations (SEM and Raman spectroscopy)to obtain a lower bound on the graphene layer thickness. The fit tendsto shrink the graphene thickness to less than one monolayer, if unrestrained,which does not agree with the SEM and Raman scans. This could be dueto the limited Q (wavevector transfer) range measuredin the time available for the experiment, as the fit is found to relystrongly on the high Q statistics. It should be notedthat in model 16, shown in Supporting Information. It is noteworthy that the improvement in the figure of merit (χ2) for the unconstrained to the constrained graphene thicknessis within the error bar of each other. However, the results show thatthe graphene thickness has only a subtle influence, if any, on theamount of the induced magnetization as can be deduced from the valuesof the measured magnetic moments. To get greater certainty on themodel selection given the data, we used a nested sampler53 to calculate the Bayesian evidence taking into account the structuralinformation, providing a high degree of confidence in the final fitmodel as shown in the Supporting Information.The fitting procedure is fully described in the 2C layer, respectively, at theinterface between the Ni and graphene, but this led to poorer fits,and the best results, shown in Supporting Information).Additional scenarios were also tested. For example, oxidationofthe Ni layer and the formation of Ni-carbide were examined by embeddingan intermediate NiO and Ni–6 Å–2)54 and reduces in size toward the surface, which is requiredto fit the lower Q features and allow the higher Q features to converge, paramount to getting certainty onthe thin graphene layers. Consequently, the mSLD also has a gradientthat oppositely mirrors the nSLD at 300 K and reaches the full momentfor Ni  near thesurface and being slightly reduced near the substrate. At 10 K themagnetic gradient vastly reduces, with the Ni layer becoming almostuniformly magnetic; we speculate that this may be a result of thelower temperature overcoming any change in the Curie temperature dueto strain effects across the film. Similar effects have been observedin other systems due to doping and strain both versus temperature.57 The Ni(111) used for the epitaxial graphene displays a much weakerstructural gradient, being almost uniform across the Ni thickness,but has the same general trends. We attribute the origin of the differencein the nSLD profiles to the fact that the samples were deposited atdifferent times and have different strains, following the growth recipedescribed in the The results of the fits to the PNR dataare summarized in 9Mo1 sample was used toclarify whetherthe induced magnetic moments in graphene are due to the C pz-Ni 3d hybridization, which opens the Dirac coneas a result of the degeneracy breaking, as postulated in refs , , a, a9Mo1 s in refs , or beca in refs −60. For d in ref . Because9Mo1 sample are shown in 9Mo1 film. At 10 K, a small butdetectable spin splitting and a minute variation in the SA are observed. Surprisingly, a higher magnetization is detected ingraphene  μB/C atom) than in theNi9Mo1 film  μB/C atom), but the large values of the 95% confidence, reflectingthe limited Q range and low counting statistics inthe data, suggest that this effect may not be real. This is visiblein the mSLD curve in 9Mo1 and a nonzero moment in the graphene.The origin of the small residual moment at 10 K could arise from clustersof unalloyed Ni throughout the layer that become ferromagnetic atlow temperature, which then polarize the graphene as per the Ni(111)samples. At 300 K, both the alloy and the graphene have effectivelyzero moment within the 95% confidence intervals. Therefore, the resultssupport the hypothesis of the universal model, whereby the measuredinduced magnetic moment in graphene is due to the opening of the ED rather than the distortion of the d band.This is because no magnetic moment is detected in the Ni9Mo1 nor in the graphene layer. If the induced magneticmoment in graphene was due to the electron transfer and surface reconstruction,we would have detected magnetism in graphene grown on Ni9Mo1.The results of the PNR measurements of the graphene/NiB/C atom was induced in both the rotated-domainand the epitaxial graphene grown on Ni films. These results indicatelarger moments than previously reported.27 In ref (B/C atom when grown on a hcp Co(0001) film.They obtained this value by comparing the XMCD signal of grapheneto that of the underlayer Co film.27 Weseret al. suggested a similar value (0.05–0.1 μB/C atom) for the magnetic moment induced in a monolayer of graphenegrown on a Ni(111) film.14 However, theirassumption was based on comparing the graphene/Ni system with otherC/3d TM structures, such as a C/Fe multilayer with 0.55 nm of C,63 and carbon nanotubes (CNTs) on a Co film,64 where magnetic moments of 0.05 and 0.1 μB/C atom were estimated, respectively. However, it is difficultto compare graphene-based heterostructures with other C allotropes/TMsystems. This is because the Dirac cone is a characteristic featureof graphene and CNTs of the carbon allotropes. Therefore, one cannotexclude that a different mechanism other than the break of degeneracyaround ED may be responsible for the magneticmoment detected in the C layer of a C/Fe multilayer system. However,although Dirac cones exist in CNTs, for the CNTs/Co heterostructurea direct quantitative analysis of the induced magnetic moment wasnot possible from the MFM images reported in ref  on different TM substrates. Induced magnetic momentin rotated-domain graphene as a result of the proximity effect inthe vicinity of a FM substrate was detected by element-specific XMCDmeasurements at the C The sample preparation procedureinvolved two stages: the growth of the TM films using magnetron sputteringand the growth of graphene by CVD.2O3(0001) substrates usinga CEVP magnetron sputtering chamber with a base pressure of (1.2–2)× 10–8 mTorr. The thick substrates were usedto reduce the possibility of sample deformation, which could affectthe reflectivity measurements. The deposition of the TM films wasperformed using 99.9% pure Ni and Ni9Mo1 targets.A DC current of 0.1 A and a constant flow of pure argon of 14 sccmwere used to grow 80 nm of highly textured Ni(111) and Ni9Mo1(111) films at a rate of 0.02 nm s–1 in a plasma pressure of 2 mTorr (3 mTorr for Ni9Mo1). 9Mo1 films.The TM films were depositedat RT on 1 mm thick Al9Mo1(111)films on ∼2 cm × 2 cm substrates. Growth recipes similarto those reported by Patera et al.30 wereadapted to obtain epitaxial and rotated-domain graphene directly onthe Ni film. For the Ni9Mo1 film, rotated-domaingraphene was grown by first introducing pure H2 gas ata rate of 200 sccm to the CVD chamber with a base pressure of 2.7× 10–6 mbar. The CVD growth chamber was heatedto 650 °C for 12 min, and the sample was then exposed to C2H4 with a flow rate of 0.24 sccm for 40 min beforeit was cooled to RT in a vacuum. This approach reduces any oxidizedTM back to a clean metallic surface before the growth of graphene.The samples were then transferred into a CVD systemfor the growthof graphene directly on Ni(111) and NiThe SEM images shown in 2 layer. This approachovercomes the fact that closely lattice-matched films lead to a lossin the resonance conditions for observing Raman spectra as a resultof the strong chemical interaction between the graphene π orbitaland the d-states of Ni and Ni9Mo1, which alsoalters the graphene’s pz orbitals.Furthermore, the increase in the C–C bond length to match thelattice of the FM leads to significant changes in the graphene’sphonon spectrum.65 For the transfer process, thesamples were cleaved into ∼5 mm × 5 mm squares, and HNO3, diluted to 5% for Ni9Mo1 and 10% forNi, was used to etch the metallic films slowly while preserving thegraphene layer.Room-temperature Raman scans weretaken at three different regions of each sample using a 532 nm excitationlaser wavelength (50× objective lens and ∼1 μm laserspot size). Before the measurements, the graphene was first transferredby a chemical etching process from the metallic films onto a Si substratewith a 300 nm thermally oxidized SiOK-edge due to the presence of carbon on the surface of theX-ray optical components.We carriedout element selective XMCD measurements to detect and distinguishthe magnetization in the graphene from that of the FM layer. The XMCDexperiments were performed at 300 K at the SIM end station of theSwiss Light Source (SLS) at the Paul Scherrer Institut (PSI), Switzerland,using the total electron yield (TEY) detection mode with 100% circularlypolarized light. The rotated-domain graphene/Ni sample was set toan incident angle of 30° from the incoming X-ray beam. An electromagnetwas fixed at 40° to the incoming X-ray beam, and a magnetic fieldof 0.11 T was applied for 30 s in-plane to the surface of the samplesto align the film magnetization along the beam direction. It was thenreduced to 0.085 T during the X-ray absorption spectroscopy measurements.The intensity of the incident X-ray beam was measured with a clean,carbon free, gold mesh placed just before the sample position. Thisis particularly important for normalizing the signal at the C Refl1D.The PNR measurementswere conducted at 10 and 300 K, under an in-plane magnetic field of0.5 T, using the Polref instrument at ISIS spallation neutron source(UK). The fitting of the data was done using the Refl1D software


Patients with colorectal cancer (CRC) and diffuse peritoneal metastasis (PM) are not eligible for surgical intervention. Thus, palliative treatment remains the standard of care in clinical practice. Systemic chemotherapy fails to cause drug accumulation at the lesion sites, while intraperitoneal chemotherapy (IPC) is limited by high clearance rates and associated complications. Given the poor prognosis, a customized OxP/R848@PLEL hydrogel delivery system has been devised to improve the clinical benefit of advanced CRC with diffuse PM. This system is distinguished by its simplicity, security, and efficiency. Specifically, the PLEL hydrogel exhibits excellent injectability and thermosensitivity, enabling the formation of drug depots within the abdominal cavity, rendering it an optimal carrier for IPC. Oxaliplatin (OxP), a first‐line drug for advanced CRC, is cytotoxic and enhances the immunogenicity of tumors by inducing immunogenic cell death. Furthermore, OxP and resiquimod (R848) synergistically enhance the maturation of dendritic cells, promote the expansion of cytotoxic T lymphocytes, and induce the formation of central memory T cells. Moreover, R848 domesticates macrophages to an anti‐tumor phenotype. OxP/R848@PLEL effectively eradicates peritoneal metastases, completely inhibits ascites production, and significantly prolongs mice lifespan. As such, it provides a promising approach to managing diffuse PM in patients with CRC without surgical indications. Patients with colorectal cancer (CRC) and diffuse peritoneal metastasis (PM) are not eligible for surgery. A customized OxP/R848@PLEL hydrogel delivery system has been developed. The PLEL hydrogel forms a drug depot within the abdominal cavity. OxP can enhance the immunogenicity of tumors. OxP/R848 exhibits a “tumor vaccine‐like” effect that elicits specific immunity. OxP/R848@PLEL effectively eradicates PM and significantly prolongs mice lifespan. At the initial diagnosis, distant metastasis is present in 44% of patients. Peritoneal metastases (PM) accounts for 25% of all cases. PM is the most intractable form of metastasis, resulting in carcinogenic consumption, intestinal obstruction, and cancerous ascites. Cancerous ascites develop rapidly, leading to abdominal distension, pain, poor appetite, and other symptoms that substantially impair patients' quality of life. Alarmingly, the median overall survival (mOS) of patients with CRC and PM is 16.3 months, whereas the mOS of patients with malignant ascites does not exceed 4 months.Colorectal cancer (CRC) is the most prevalent malignant tumor of the digestive system and is characterized by high morbidity and mortality rates.1 The  In this scenario, palliative systemic chemotherapy is adopted to maximize survival. However, the dosage of systemic chemotherapy is restricted by toxicity and the patient's physical condition, resulting in suboptimal drug concentrations at tumor sites and mediocre therapeutic outcomes. Thus, intraperitoneal chemotherapy (IPC) has attracted attention from oncology researchers due to its ability to increase drug concentration in abdominal lesions and reduce the incidence of systemic adverse reactions. Unfortunately, IPC has certain limitations, including limited drug retention time, uneven distribution of drugs, and potential complications such as bleeding and intestinal perforation. Therefore, developing novel drug formulations that can effectively eradicate PM and suppress ascitic fluid production is imperative.For patients diffuse PM, surgery fails to prolong patient survival and poses a risk of complications, such as enterocutaneous fistula, wound infection, and wound dehiscence.10 In  Compared to nanocarriers, hydrogels made from suitable materials are biodegradable and less toxic, making them crucial in drug delivery, 3D cell culture, tissue engineering, and other medical fields. Hydrogels can be customized to facilitate subcutaneous, rectal, and intraperitoneal delivery according to pathological characteristics, indicating their potential use in localized drug delivery for a variety of solid tumors. Furthermore, intelligent hydrogels can detect and react to environmental stimuli for precise drug delivery. Due to its excellent biocompatibility, facile modifiability, and high drug encapsulation efficiency, extensive research and development have been conducted on PEG‐based hydrogels. In our previous study, we successfully synthesized a biodegradable PDLLA‐PEG‐PDLLA (PLEL) hydrogel with an appropriate phase‐transition temperature. The PLEL hydrogel exhibits a flowable sol state at room temperature (RT), facilitating drug loading and injection. After in vivo administration, it undergoes a phase transition to form a drug‐loaded gel depot for sustained drug release. Moreover, PLEL hydrogels can ensure the uniform distribution of drugs in the abdominal cavity and effectively prolong drug retention time. Their injectability and biodegradability considerably reduce complications, thereby improving patient compliance.Hydrogels with biocompatibility and tunable mechanical properties may present an alternative solution for overcoming the limitations of IPC.20 Com Oxaliplatin (OxP) is the preferred first‐line agent for treating advanced metastatic CRC. Apart from to cytotoxicity, OxP can induce immunogenic cell death (ICD), which promotes the release of damage‐associated molecular patterns, boosts exposure to tumor‐associated antigens (TAAs), and thus improves tumor immunogenicity. However, the immunity triggered by ICD effects is insufficiently robust in most cases. With advancements in vaccines, immune adjuvants have become essential components of cancer vaccines. As the most capable antigen‐presenting cells (APCs), the maturation of dendritic cells (DCs) is crucial for initiating anti‐tumor immunity. The immune adjuvant resiquimod (R848) can stimulate toll‐like receptor 7/8 (TLR7/8) and promote DCs maturation, thus inducing cytotoxic T lymphocytes (CTLs) response. Furthermore, R848 polarizes tumor‐associated macrophages (TAMs) into pro‐inflammatory phenotypes (M1), alleviating the immunosuppressed tumor microenvironment (TME). Therefore, OxP combined with R848 exhibits a “tumor vaccine‐like” effect that elicits specific immunity and produces an efficient and long‐lasting anti‐tumor response.In addition to chemotherapy, immunotherapy has brought possibilities for treating refractory cancers.29 OxaFigure Herein, we further explored the potential benefits of a chemo‐immunotherapy strategy based on thermosensitive hydrogels. We fabricated an OxP/R848@PLEL drug‐loading system via simple physical mixing techniques (22.12). The structure and molecular weight were characterized by 1H NMR spectroscopy .We synthesized PLEL block copolymers via the ring‐opening polymerization of polyethylene glycol (PEG) and D, L‐lactide  catalyzed by stannous octoate (Sn(Oct)Figure  The in vitro gelation of OxP/R848@PLEL was assessed using the tube inversion method  to the membrane and facilitating antigen uptake by DCs. Using immunofluorescence, we detected CRT expression in CT26‐luc cells treated with low, medium, and high concentrations of OxP. As depicted in Figure  Additionally, ATP can recruit immature DCs and macrophages to ICD sites. Subsequently, the concentrations of both HMGB1 and ATP in the supernatant were measured using ELISA  pathway. This study yielded the anticipated flow cytometry results verified that R848 alone or combined with OxP could polarize Raw264.7 cells to the M1 phenotype . Combining the two drugs further increased the proportion of CD8+ T cells in the spleen  were procured from Sigma‐Aldrich . D, L‐LA was obtained from Daigang Chemicals . OxP and R848 were supplied by Macklin Biochemical Co., Ltd. . RMPI 1640 medium, DMEM, fetal bovine serum (FBS), and penicillin‐streptomycin were purchased from Gibco, Inc. (USA). The CCK‐8 was purchased from MedChemExpress . Sulfo‐Cy5.5 was provided by Beijing Fluorescence Biotechnology Co., Ltd . Recombinant mouse GM‐CSF, IL‐4, and IL‐10 were purchased from Peprotech. Mouse IFN‐γ ELISA kit, Mouse IL‐2 ELISA Kit and antibodies for flow cytometry were obtained from BioLegend (USA). d‐Luciferin potassium salt was acquired from the Shanghai Yeasen Biotechnology Co., Ltd. . Anti‐mouse CD8α (clone 2.43) was purchased from Selleck (USA). All other chemicals used in this study were of analytical grade and used without further purification.PEG (Mn = 1500) and Sn(Oct)CT26‐luc and Raw264.7 were cultured in RPMI 1640 medium and DMEM, respectively, supplemented with 10% FBS, 1% penicillin‐streptomycin. The cultures were maintained at 37 °C under a relative humidity of 95%.BALB/c mice  aged 6—8 weeks were purchased from Gempharmatech Co., Ltd . The mice were housed under specific pathogen‐free conditions and provided ad libitum access to food and water. All animal experiments were approved by the Experimental Animal Ethics Committee of the State Key Laboratory of Biotherapy at Sichuan University (No: 20 220 531 049).2. A dry three‐neck flask was filled with PEG, D, L‐LA, and Sn(Oct)2 , followed by vacuum pumping to remove water. Subsequently, the reaction mixture was heated to 140 °C in an oil bath under a nitrogen atmosphere for 10 h. The crude product obtained was dissolved in ultra‐pure water at RT, followed by heating to 80 °C for precipitation of the target product and removal of unreacted monomers and low molecular weight products. The precipitate was freeze‐dried until a constant weight was obtained. The structure and molecular weight of PLEL copolymers were characterized by 1H‐NMR  and gel permeation chromatography .Triblock copolymer PLEL was synthesized via ring‐opening polymerization of PEG and D,L‐LA catalyzed by Sn(Oct)2.27 A d−1, R848: 0.1 mg mL−1) with a PLEL micellar solution at RT. The drug‐loaded hydrogel was sterilized using a 0.22 µm filter membrane and stored at 4 °C for later use.First, the PLEL block copolymer was dissolved in PBS (pH 8.0) at RT to prepare blank PLEL micelles at 20 wt%. Second, OxP/R848@PLEL was obtained by mixing a certain amount of OxP and R848 . The sample precooled at 4 °C was placed on a circular flat fixture with a diameter of 20 mm. Low‐viscosity silicone oil was applied to the edge of the sample to prevent moisture volatilization owing to the rising temperature. The sample's temperature was increased from 10 to 60 °C at the rate of 1 °C min One milliliter of drug‐containing PBS solution or drug‐loaded PLEL hydrogel  was placed into a dialysis bag and incubated at 37 °C for 30 min. Then the dialysis bags were transferred to centrifuge tubes containing 10 mL of pH 7.4 PBS and incubated on a shaker at 100 rpm and 37 °C. The release medium was extracted at predetermined time points, followed by adding an equal volume of preheated PBS. The collected release medium was stored at −20 °C and subsequently analyzed for drug release using high‐performance liquid chromatography.The release behavior of OxP and R848 from the PLEL hydrogel was investigated using an improved dialysis method.59 OneTo evaluate the sustained‐release performance of the PLEL hydrogel as a drug delivery carrier, the fluorescent dye Sulfo‐Cy5.5 as a model drug to investigate drug retention in the abdominal cavity was utilized . BALB/c mice were intraperitoneally injected with 200 µL PBS solution containing Sulfo‐Cy5.5 or PLEL hydrogel loaded with Sulfo‐Cy5.5. To visualize the in vivo behavior of the drug, fluorescence images were acquired using an IVIS. The fluorescence intensity data were recorded using the LivingImage software for subsequent image analysis.The cytotoxicity of free OxP and R848 in CT26‐luc cells using a CCK‐8 assay was first evaluated. The cells were inoculated in 96‐well plates with 100 µL complete medium per well and incubated overnight. The original medium was replaced with a fresh complete medium or varying concentration of OxP or R848. After treatment for 24, 48, and 72 h, the cells were subjected to CCK‐8 assay and the absorbance was measured at 450 nm using a Bio‐Rad instrument. The relative cell activity in the experimental group was expressed as a percentage of that in the control group.Next, the cytotoxicity of OxP@PLEL against CT26‐luc cells using a transwell co‐culture system was assessed. CT26‐luc cells were seeded into 24‐well plates and incubated overnight. Subsequently, 100 µL of PBS and free OxP were added directly to the upper chamber of the transwell apparatus. Currently, 100 µL of blank PLEL hydrogel or OxP@PLEL was added to the same upper chamber and left for 30 min to solidify from sol phase into gel phase before insertion into its corresponding position within the co‐culture system. The co‐culture system was incubated and refreshed with 500 µL preheated fresh complete medium every 24 h. After culturing for 24, 48, and 72 h, CCK‐8 treatment was administered to the cells, and the absorbance at 450 nm was measured. The relative cell activity in the experimental group was calculated as a percentage of that in the control group.−1) for an additional 6 h. The culture medium was collected, and the cells were fixed in 4% paraformaldehyde at RT for 15 min. After fixation, the samples were sealed with a closed buffer solution for 60 min. The diluted CRT primary antibody was added and incubated overnight at 4 °C. Subsequently, a fluorescein‐conjugated secondary antibody was added and incubated in the dark at RT for 2 h. Finally, the nuclei were stained with hoechst33342. The slides were sealed with an anti‐fluorescence attenuation tablet and immediately observed under a fluorescence microscope. Three visual fields with a cell count >50 were randomly selected for each slide, and the percentage of CRT+ cells was calculated.CRT expression was detected through immunofluorescence. CT26‐luc cells were inoculated onto polylysine‐coated coverslips and incubated for 12 h before treatment with fresh medium containing varying concentrations of OxP  to induce bone marrow cells to differentiate into macrophages (BMDMs). The BMDMs were treated with PBS, R848, or IL‐4&IL‐10 to induce the polarization of BMDMs into M0, M1, and M2 states, respectively. Subsequently, a transwell co‐culture system . The weight and daily observations of the mice's state  and survival were recorded.An acute toxicity test of OxP@PLEL was conducted on 15 healthy BALB/c mice acclimated for 1 week. The mice were randomly divided into five groups of three animals each. Each group received a single intraperitoneal injection of either 200 µL PBS or OxP@PLEL  per group. The drug was administered at 4‐day intervals for three doses, and the weights of the mice were measured twice daily. Daily records of hair, diet, defecation, and survival rate were kept.Subsequently, OxP@PLEL were repeatedly administered to 25 healthy BALB/c mice. Then subjects were randomly allocated into five groups and administered intraperitoneal injections of 200 µL PBS or OxP@PLEL  intraperitoneally. Four days after inoculation, live imaging was performed on small animals. Mice with low or no fluorescence signals were excluded from the study. The successfully modeled mice were randomly divided into seven groups, each consisting of four mice. Each group was intraperitoneally injected 200 µL of PBS, OxP@PLEL  or OxP/R848@PLEL . The drug was administered at 4‐day intervals for three doses. Mouse weight was measured twice daily, and the status  and survival were recorded daily. On the sixteenth day post‐inoculation, the mice were euthanized under isoflurane anesthesia. Abdominal tumors were excised and weighed, and the peritoneal fluid volume was measured.In the two‐drug combination dose screening, the dose of OxP was determined based on the results of toxicity experiments, whereas the dose of R848 depended on its solubility in PLEL micelles. To establish the PM model of CRC, BALB/c mice were injected with CT26‐luc cells (8 × 10−1). The medication was administered at 4‐day intervals for three doses. The mice were weighed twice daily. On day 16, five mice from each group were euthanized under isoflurane anesthesia for the surgical extraction of abdominal tumors and major organs, and ascites volumes were recorded. The remaining 10 mice from each group were observed for survival.Mice with PM were randomly assigned to each of the eight groups. Each group received an intraperitoneal injection of 200 µL of either PBS, blank PLEL, OxP@PBS, R848@PBS, OxP/R848@PBS, OxP@PLEL, R848@PLEL or OxP/R848 @PLEL (OxP: 6 and R848: 1 mg kg−1) every 4 days for three doses. Mice exhibiting no fluorescence signal on in vivo imaging on days 16 and 44 were considered to have achieved complete tumor regression. On the same day, CT26‐luc cells (8 × 105 cells per mouse) were intraperitoneally injected into both mice with tumor‐regressed and untreated mice to simulate residual tumor cells. On days 48 and 54, tumor growth was observed using an IVIS. On day 54, the mice were euthanized with isoflurane anesthesia and underwent surgical removal of the abdominal tumors and spleens. The central memory T cell content in the spleen was analyzed using flow cytometry.Ten PM mice were intraperitoneally injected with 200 µL of OxP/R848@PLEL (OxP: 6 and R848: 1 mg kg−1). The medication was administered at 4‐day intervals for three doses. On day 16, the mice were euthanized under isoflurane anesthesia, and their abdominal tumors and spleens were surgically excised. Proliferation and apoptosis of tumor tissues were analyzed by immunohistochemical staining, while T cell activation in the spleen and the proportion of M1‐TAMs in the TME were assessed via flow cytometry. In addition, the levels of IFN‐γ and IL‐2 in the serum and TME of mice in each group were quantified by ELISA. In order to prove the role of CD8+ T cells in OxP/R848@PLEL anti‐tumor therapy, OxP/R848@PLEL &Anti‐CD8α group was set up in addition to the model group and OxP/R848@PLEL treatment group. The biological activity of CD8+ T cells were restricted by binding to mouse CD8α .Mice with PM were randomly divided into four groups of six mice per group. Each group was intraperitoneally injected with 200 µL of PBS, OxP@PLEL, R848@PLEL, or OxP/R848@PLEL (OxP: 6 and R848: 1 mg kgP values for survival curves were determined by the Gehan‐Breslow‐Wilcoxon test, while Student's t‐test was employed for comparing between two groups. For statistical analysis among multiple groups, one‐way ANOVA with Tukey's test was utilized. Statistical significance was defined as *p < 0.05, while **p < 0.01, and ***p < 0.001 indicated extreme significance.Data were obtained from three independent experiments and analyzed using GraphPad Prism 8.0. All date were presented as the mean ± SD. The The authors declare no conflict of interest.Supporting InformationClick here for additional data file.


A gradual increasein the width of the graphene device resulted in the increase of thefull width at half-maximum (FWHM) of field-effect characteristicsand a decrease in the field-effect mobility. Our work presents a simplifiedfabrication approach to realize high-mobility suspended CVD graphenedevices, beneficial for developing CVD graphene-related applications.High-field-effect mobility and the two-dimensional natureof graphenefilms make it an interesting material for developing sensing applicationswith high sensitivity and low power consumption. The chemical vapordeposition process allows for producing high-quality graphene filmsin a scalable manner. Considering the significant impact of the underlyingsubstrate on the graphene device performance, methods to enhance thefield-effect mobility are highly desired. This work demonstrates asimplified fabrication process to develop suspended, two-terminalchemical vapor deposition (CVD) graphene devices with enhanced field-effectmobility operating at room temperature. Enhanced hole field-effectmobility of up to ∼4.8 × 10 It is unique for its linear energy-momentumrelation with the conduction and valence band intersecting with zeroband gap.2 Ballistic charge carrier transportexceeding several micrometers has been demonstrated in graphene devices.Such unique electronic transport properties are beneficial for exploringnovel applications such as ballistic electronic devices, rectifierswith high voltage responsivity, graphene-based Josephson junctionsin quantum circuits, negative refraction effects, and quantum metrology.9Single-layer graphene films offer a uniquecombination of electronic,mechanical, thermal, and optical properties.13 Chemical vapor deposition (CVD) and epitaxial growth techniquesare the two widely utilized techniques to produce high-quality syntheticgraphene films on an industrial scale.10 Improving the intrinsic quality of synthetic graphene films hasbeen widely studied. “Quality” herein refers to performancemetrics such as field-effect mobility, sheet resistance, crystallinity,layer thickness, and the defect-free nature of graphene films. Manystudies reported the growth methods of single-crystalline CVD graphene(CVDG) films on large-area copper foils.17 This is promising for realizing wafer-scale production of high-performanceCVDG applications. In addition, unique graphene integration methodsto “transfer” the graphene films from their growth substratesto target substrates have been widely explored. Some of the well-knowntransfer methods are polymer-assisted wet etching (WE), electrochemicaldelamination (ED), lamination-assisted (LA), and dry transfer (DT).21Given the unique properties of graphene films, there is animmenseinterest in large-scale production methods and the commercializationof graphene-based applications. Graphene production techniques haveseen significant progress during the last decade, focusing on productionmethods, the quality of synthetic graphene films, and the integrationof graphene films.4 cm2/Vs, and an average hole mobility of >104 cm2/Vs across all devices can be achieved in suspended CVDG devicesoperating at room temperature, realized by standard cleanroom processingmethods.Such transfer methods play a vital role in determining theperformancemetrics of the CVDG devices and their applications. Here, we presentour study on suspended CVD graphene devices with high-field-effectmobility operating in ambient conditions. Our work utilizes standardcleanroom fabrication processes and the CVD method to fabricate thesuspended CVDG devices. We characterize the suspended CVDG devicesin ambient conditions and study the influence of increasing CVDG filmwidth on mobility and impurity scattering. We demonstrate high-field-effectmobility of up to ∼4.8 × 102) is heavilyinfluenced by the charge traps and corrugations in the oxide layer,thereby limiting their field-effect mobility significantly.22 Type 2 and type 3 approaches have explored methodsto isolate the graphene film from the oxide. Type 2 approach is focusedon encapsulating graphene films within multilayer hexagonal boronnitride (hBN) films.24Several graphene device fabricationapproaches have been exploredover the last few years. The exfoliation method was the first approachto demonstrate the successful isolation of graphene and electricaltransport in the material. CVD and epitaxial approaches were lateridentified as viable methods to produce graphene films on a largerscale beyond the limits of exfoliated graphene flake size. In addition,several device architectures have been explored with a focus on boostinggraphene device performance. Three prominent device architecturesare (i) on-substrate (type 1), (ii) encapsulated (type 2), and (iii)suspended (type 3). Charge transport in “type 1” graphenedevices fabricated on silicon dioxide (SiO29 Prominent performance metrics such as field-effect mobility (μ),charge neutrality point (CNP), full width at half-maximum of the field-effectcharacteristics (Δn), residual charge carrierdensity, and minimum conductivity at the CNP provide valuable insightsinto the graphene device performance. Type 2 and 3 graphene devicessignificantly improve such performance metrics compared to type 1devices.29Type 3 approaches have explored “suspending”thegraphene film, wherein the graphene film is supported at the contactregion and freely suspended across the channel region.30 Chen et al. reported onthe influence of depositing potassium atoms on the conductivity ofa graphene device, resulting in decreasing field-effect mobility dueto charged impurity scattering.31 Goslinget al. reported on broadening of the field-effect characteristic dueto increasing impurity density on the graphene layer and thereby decreasingmobility.32 Exploring the influence ofcharge impurity scattering in graphene devices helps validate thequality of graphene devices and their fabrication approaches.Charge impurity scattering, defects, grain boundaries, artifacts, and temperature significantly influence the field-effectmobility (μ) of graphene films. Many studies have shown theinfluence of charge impurity scattering in graphene films. Adam etal. reported a carrier transport theory using self-consistent randomphase approximation (RPA)-Boltzmann transport for impurity scatteringin graphene.Graphene synthesis via chemical vapor deposition is a scalablemethod to produce graphene films on a large scale. The CVD methodallows for synthesizing graphene films on various growth substratesand layer thicknesses. We synthesize graphene on commercially availablecopper foils (thickness ∼25 μm) using CVD. Polymer-assistedwet etching (WE) process was used to transfer CVDG film onto the targetsubstrate. The WE process consists of three steps.First, apolymer film  PMMA) was adheredto CVDG on copper (PMMA/CVDG/Cu) using spin coating. Second, the growthsubstrate wet etching was performed. Third, the PMMA/CVDG layer wastransferred to deionized water (DIW) to reduce contamination frometchant exposure. The free-floating PMMA/CVDG layer on the DIW surfaceis picked up using the preprocessed and structured target substratesee: 1a andle2 interface are attributed to impuritydensity ∼ >1011 cm–2, influencingthe charge transport in graphene films.33 Methods to isolate graphene films is a robust approach to improvedevice performance. Type 2 approach requires high-quality, exfoliatedhBN films for graphene encapsulation. However, the type 3 approachrelies on suspending the graphene films such that the field-effectregion does not interface with the oxide surface. Considering thefeasibility of the CVD process, we present a simplified fabricationapproach to realize type 3 graphene devices using standard microfabricationmethods.Charge traps in the oxide surface and scatteringdue to the surfacephonons at the SiO2 layer. Next, a critical point drying (CPD) process was performedto suspend the CVDG devices.A schematic description of the suspended CVDG devicefabricationis presented in Figure S1.Refer to the Figure S2. This allowed for the preparation ofultra-clean, polymer-free suspend CVDG films for TEM characterization.In 35 The TEM imageof the CVD graphene region wherein the SAED image was collected ispresented in Supporting Information Figure S3. In Supporting Information Figure S3,the elemental analysis of the CVD graphene region was performed usingenergy-dispersive X-ray analysis (EDX). The presence of residual contaminationfrom carbon, oxygen, and sulfur elements is observable in EDX characterization of the suspendedgraphene film region presented in Supporting Information Figure S4.To characterize the crystallinity of the CVDG film, a customizedTEM grid process was followed to suspend CVDG directly on copper foils(growth substrate). Scanning electron microscope (SEM) image of thecustom TEM grid sample is presented in Supporting Information L = 2 μm, W = 6 μm) is presented. In VSD = 5 mV) were performed in air, under ambient pressure, and in standardcleanroom conditions . Additional details related tothe measurement acquisition are provided in the In 2/h . The residual charge carrier density was estimated ∼at1.35 × 1010 cm–2 by using a linearfit  conductance (red dotted curve) intersects with the minimumconductance of ∼8.4 eL = 2 μm) and varying widthswere fabricated. The width of the devices was varied from W = 2 μm (S11) to 4 μm (S12), 6 μm (S22),and 8 μm (S21). The field-effect characterization of the differentsample types is presented in To study the influence of the device aspectratio (L/W), CVDG deviceswith fixed lengths (h) and electrons(μel), CNP, and the full width at half-maximum (FWHM)in charge carrier density (Δn) of all devicesis presented. By considering a Gaussian peak fit to the field-effectcharacteristics, the FWHM in gate voltage and, thereby, in chargecarrier density (Δn) is estimated.The CNP of all devices is within the measurementwindow ±5 V gate voltage and > 0 V in all of the devices, indicating thepresenceof p-doping. In h (μel) in sampletype S11 is ∼31000 cm2/Vs (∼14500 cm2/Vs), μh (μel) in S12 is∼16700 cm2/Vs (∼10600 cm2/Vs),μh (μel) in S22 is ∼16400cm2/Vs (∼5100 cm2/Vs), and μh (μel) in S21 is ∼12800 cm2/Vs (∼5300 cm2/Vs). The average μh (μel) across all of the sample types is >19000cm2/Vs (> 8900 cm2/Vs), see: h and μel with the increase in sample widthis observable. The hole mobility μh is higher thanthe electron mobility μel in all of the samples.The μh/μel ratio for each of thesample types is S11 ∼2.1, S12 ∼1.5, S22 ∼3.2,and S21 ∼2.4, with the overall average μh/μel across all sample types being ∼2.3. Several factors,such as electrostatic potential fluctuations at the contact regions,strain, charged impurities, and electron–phonon scattering,can introduce such an asymmetry in field-effect characteristics.37 Studies to selectively suppress hole and electron conductance viamolecular doping (charged impurities) have been reported.37The average μ0 across all samplesis ∼3.2 × 1010 cm–2. P-typedoping of the samples can occur during fabrication and on exposureto photoresist and ambient measurement conditions. Residual p-typedoping correlates well with the observation of a larger hole conductancecompared to the electron conductance in n) parameter from the field-effect characteristics can beused to study the impact on mobility.32A minor shift in the CNP is observablein all of the devices see: 3. In Fign has been reported by considering impurityscattering as a primary factor influencing sample mobility.32 In n is increasing, S11nΔ ∼6 × 1010 cm–2, S12nΔ ∼ 6.6 × 1010 cm–2, S22nΔ ∼ 1 × 1011 cm–2, and S21nΔ ∼ 8.4 × 1010 cm–2. In n increasesby an average ∼1.9 × 1010 cm–2 with the increase in sample width. The increase in Δn and minimum conductance values and the decreasing mobilitycan be collectively attributed to the increase in impurity or defectdensity (nimp) in these samples.32 Considering the exposure to polymeric contamination during the fabricationprocess, current annealing characterization was performed before field-effectcharacterization. Current annealing is a useful technique for removingresidues on suspended graphene samples.39 The influence of current annealing on CVD graphene characteristicsis presented in Supporting Information S6. Other groups have reported current density >108A/cm2 during the current annealingprocess at low temperatures and under high-vacuum conditions.38A model with an inverse power law dependence between mobilityandΔFigure S7,the current annealing characteristics of S11D4 are presented. Notethat in our study, the current annealing was performed in air andat room temperature, which is a typical operating environment forseveral sensing applications. A current density of ∼0.43 ×108 A/cm2 was observable in S11D4at a VSD ∼1V . In Supporting Information Figures S8S–11, the current annealingcharacteristics of additional device types are presented. The currentdensity (at VSD ∼1V) is ∼0.43× 108 A/cm2 in S11D4, ∼0.12× 108 A/cm2 in S12D14, 0.18× 10 8A/cm2 in S22D21, and∼0.11 × 108 A/cm2 inS21D19. Due to the inverse relation between the current density andsample width, a decreasing current density trend with the increasein sample width is observable at a fixed VSD ∼ 1V.In Supporting Information 8 A/cm2, whichis at least an order of magnitude larger than the current densityvalues observed in our devices. Note that the breakdown temperature(TBD) of graphene samples in the air  is significantlylower in comparison to TBD in vacuum .40 In our case, the measurements in higher currentdensity regimes with VSD > 2V werelimitedto minimize the potential risk of device degradation from thermaloxidation in air. The gradual decrease in the field-effect mobilitywith the increasing width can be explained by the decreasing currentdensity trend across the CVDG device types, i.e., higher residualcharge impurities remain on the samples due to the limited currentannealing of larger CVDG devices. EDX characterization of suspendedCVD graphene on custom grids was performed to characterize the elementalresidues on CVD graphene. In Supporting Information Figure S4, the elemental presence of carbon, oxygen, and sulfurwas identified on the suspended CVD graphene films. The presence ofsuch residual elemental composition can be attributed to processing-relatedconditions. Considering the presence of oxygen and sulfur atoms inthe copper etchant (ammonium persulfate), residues from the copperetching process support the presence of residual oxygen and sulfurgroups on the CVD graphene surface.The reported current density inhigh-mobility suspended graphenedevices by other groups is >10In Note that the current annealing conditions vary considerablybetweenS11D4 and the exfoliated graphene-based type 3 devices in 4 cm2/Vs at room temperatureand ambient pressure using standard cleanroom processing methods.A width-dependent study of the suspended CVD graphene devices showedan increase in the FWHM of the field-effect characteristics and adecrease in mobility. This observation is attributed to the increasingcharge impurity scattering in larger suspended CVD graphene devices.Further processing conditions such as thermal annealing, contact passivation,and measurement under inert conditions can be beneficial in furtherimproving the performance metrics of the suspended CVDG devices anduniformly lowering the impurity density across the samples. The demonstrationof high-field-effect mobility in suspended CVD graphene devices operatingin ambient conditions strengthens the prospects of developing scalable,high-performance sensing solutions with synthetic graphene films.In summary, we presented a fabrication approachto realize suspendedCVD graphene devices with high hole field-effect mobility of up to∼4.8 × 10CVD synthesis of graphene was performedusing a low-pressure, hot-wall CVD reactor from Graphene Square, Inc.The copper foils were first thermally annealed at 1030 °C forfour hours. The CVD synthesis of graphene was performed at 1000 °Cusing a methane/hydrogen ratio of ∼ 1:13. After a growth timeof ∼45 min, the chamber was cooled to room temperature in argonand hydrogen gas flow.A PMMA 50K layer was spin-coated on graphene grown on the top side of copper foil.The graphene on the bottom side of the copper foil was removed inoxygen plasma . Ammonium persulfatesolution (50 mM) was used to etch the copper layer. After the etchingprocess, the PMMA/CVDG layer was transferred to the DIW solution.The PMMA/CVDG layer was then picked up using the target substrateand let to dry. The PMMA layer was dissolved in acetone and rinsedin IPA.Two-terminal etch mask patternwas realized using the AZ1512  photolithography process.The exposed graphene layer was removed using an oxygen plasma process. After graphene patterning,the etch mask was removed in acetone and rinsed in IPA.2. After the BOE process, the sample was rinsed in DIW solution (∼15min). The DIW rinse step was repeated 3x and then moved to IPA. Next,the etch mask is removed in acetone (∼15 min) and then movedto IPA. To prevent stiction-related challenges, a Tousimis criticalpoint dryer was used for the suspension process. In total, 17 activedevices were characterized during this study with a fabrication yieldof ∼17%.An etch mask pattern for graphenesuspension was realized using a second photolithography step usingAZ1512  photolithography process. Buffered oxide etching(BOE) step was performed (∼15 min) to remove the SiOThe field-effect characterizationof the devices was performed in air and under ambient cleanroom conditions using a wafer probe station (Cascade Summit 12000B-AP)equipped with an Agilent B1500A semiconductor parameter analyzer.Selected area electron diffraction,scanning transmission electron microscopy (STEM), and energy-dispersiveX-ray analysis (EDX) were carried out on a TFS Talos F200X runningat an 80 kV acceleration voltage.