# A 3D porous biofilm-inspired alginate/gellan hydrogel: Investigating printability and rheological properties affected by dual-wavelength UV crosslink 3D printing

Peyman Asghartabar Kashi, Caroline Bachlechner, Delphine Huc-Mathis, Henry Jäger, Mahdiyar Shahbazi

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Peyman Asghartabar Kashi, Caroline Bachlechner, Delphine Huc-Mathis, Henry Jäger, Mahdiyar Shahbazi. A 3D porous biofilm-inspired alginate/gellan hydrogel: Investigating printability and rheological properties affected by dual-wavelength UV crosslink 3D printing. Carbohydrate Polymers, 2025, 370, pp.124246. 10.1016/j.carbpol.2025.124246. hal-05221725

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A 3D porous biofilm-inspired alginate/gellan hydrogel: Investigating printability and rheological properties affected by dual-wavelength UV crosslink 3D printing

Peyman Asghartabar Kashi, Caroline Bachlechner, Delphine Huc-Mathis, Henry Jäger, Mahdiyar Shahbazi

<figure><img src="image_2.png" /></figure>

PII: S0144-8617(25)01031-8

DOI: https://doi.org/10.1016/j.carbpol.2025.124246

Reference: CARP 124246

To appear in: *Carbohydrate Polymers*

Received date: 12 December 2024

Revised date: 11 August 2025

Accepted date: 14 August 2025

Please cite this article as: P.A. Kashi, C. Bachlechner, D. Huc-Mathis, et al., A 3D porous biofilm-inspired alginate/gellan hydrogel: Investigating printability and rheological properties affected by dual-wavelength UV crosslink 3D printing, *Carbohydrate Polymers* (2025), https://doi.org/10.1016/j.carbpol.2025.124246

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A 3D porous biofilm-inspired alginate/gellan hydrogel: investigating printability and rheological properties affected by dual-wavelength UV crosslink 3D printing

Peyman Asghartabar Kashi<sup>a,b</sup>, Caroline Bachlechner<sup>a</sup>, Delphine Huc-Mathis<sup>c</sup>, Henry Jäger<sup>a*</sup>, Mahdiyar Shahbazi<sup>a*</sup>

<sup>a</sup>Institute of Food Technology, Department of Biotechnology and Food Science, BOKU University, Muthgasse 18, 1190, Vienna, Austria

<sup>b</sup>Faculty of Biosystem, College of Agricultural and Natural Resources, Tehran University, 31587-77871 Karaj, Iran

<sup>c</sup>Université Paris-Saclay, INRAE, AgroParisTech, UMR SayFood, 91120 Palaiseau, France

**Corresponding Authors**

Mahdiyar Shahbazi - University of Natural Resources and Life Sciences (BOKU), Muthgasse 18, 1190Vienna, Austria; Orcid: https://orcid.org/0000-0002-2485-9130; Email: mahdiyar.shahbazi@boku.ac.at; shahbazim00@yahoo.com

Henry Jäger - Institute of Food Technology, University of Natural Resources and Life Sciences (BOKU), Muthgasse 18, 1190Vienna, Austria; Email: henry.jaeger@boku.ac.at
**Abstract**

The advancement of additive manufacturing through 3D printing of polymers with tailored material properties remains a complex challenge, particularly for multi-material structures. Traditional methods often rely on automated sample swapping, resulting in slower production speeds and compromised interlayer adhesion.

In this study, we introduce a novel dual-wavelength UV crosslinking technique, utilizing chemo-selective wavelengths (UV-A at 390 nm and UV-C at 260 nm) to produce multi-material porous objects with enhanced mechanical properties. We develop an alginate/gellan-inspired hydrogel that mimics biofilm materials, allowing for precise manipulation of stiffness by fine-tuning photosensitization methods. Our results demonstrate that objects printed with UV-A irradiation yield softer, more flexible networks, while those cured with UV-C produce stiffer, more elastic structures. The optimization of photoinitiators in combination with a high-intensity LED-based bioprinter significantly broadens the achievable mechanical property range and improves build efficiency, thereby overcoming the inherent limitations of traditional UV-curing techniques.

This study reveals the promising potential for functional multi-material components in advanced manufacturing applications.

**Keywords**

Multi-material 3D Printing; UV-Curing; Photocrosslinking; Photosensitizer; Mechanical Properties; Macroporous structures, Tensile loading-unloading curve.
# 1. Introduction

Biofilms—microbial communities encased in extracellular matrices—are critical in industrial processes yet challenging to cultivate with controlled structural and growth features (Flemming & Wingender, 2010). Despite their potential benefits, cultivating biofilms with predictable structural characteristics and growth patterns presents considerable challenges (Karygianni, Ren, Koo, & Thurnheer, 2020). A key hurdle is the poorly understood molecular interactions driving biofilm assembly, hindering reliable mimicry (Guzmán-Soto 2021).

Most of the research concerning biofilm imitation (model) has employed two-dimensional (2D) biofilms that primarily facilitate horizontal biofilm-based imitation development (Chandra, Mukherjee, & Ghannoum, 2008).

These models do not adequately replicate a three-dimensional (3D) microenvironment, which is crucial for fostering the dynamic structural growth—both horizontal and vertical—of biofilms (Schaffner et al., 2017).

Biofilms cultivated on conventional surfaces, such as hydroxyapatite discs and microtitre plates, tend to be more vulnerable to antimicrobial agents compared to those formed on more intricate substrates, suggesting that the former approach does not accurately reflect in vivo conditions (Williams et al., 2019). Modern strategies prioritize 3D models that emulate native microenvironments, particularly for clinical applications where biofilms colonize porous tissue-engineered scaffolds or 3D printlets, posing contamination risks (Shahbazi et al., 2024a). Understanding this interaction is vital to formulating effective prevention and treatment strategies.

In recent decades, the evolution of 3D printing technology has been nothing short of remarkable, enabling the quick and tailored production of intricate geometries at both micro and macro scales (Menshutina et al. 2021; Iglesias-Mejuto & García-González, 2021; Yu et al., 2025). A 3D-printed porous structure that supports a higher porosity and higher drug release rate can facilitate custom-designed personalized 3D scaffolds essential for biomedical science. Alginate, a naturally occurring polysaccharide, serves as the core component of 3D-printed biofilm models. Menshutina et al. (2021) engineered 3D-printed alginate scaffolds
via extrusion and supercritical drying, achieving high porosity for tissue engineering. Their work motivates our use of covalent photocrosslinking to enhance structural precision in alginate-based biomaterials for the biomedical area. Iglesias-Mejuto & García-González (2021) combined 3D-printing/aerogel methods to create biocompatible alginate-hydroxyapatite scaffolds with dual porosity and shape fidelity, supporting cell proliferation for bone regeneration. However, neat alginate-based 3D structures without physicochemical modifications tend to demonstrate limited printability and shape fidelity. This presents challenges in maintaining the structural integrity necessary for mimicking the mechanical properties of 3D-printed biofilm models. Functional enhancement of supramolecular alginate is achieved via structural modifications such as ionic crosslinking (Shahbazi et al., 2024b) or biopolymer blending (Cao et al., 2024). Low-acyl gellan gum, valued for its aqueous gelation behavior, exemplifies a synergistic blending candidate for biofilm modeling, where molecular weight and acetylation degree critically influence gel strength, thermal stability, and formation kinetics (McCarthy et al., 2019). Photo-cross-linking has emerged as an alternative strategy to optimize rheological performance and mechanical integrity in 3D-printed alginate structures, addressing the inherent limitations of alginate-based inks (Samorezov et al., 2015). These material innovations broaden additive manufacturing capabilities, enabling structurally refined biomaterials with tailored functionality for applications spanning biomedical engineering to advanced material design.

Photocuring, a versatile UV-driven technique for biomaterial modification, enables rapid crosslinking via photoinitiators or photosensitizers like sodium benzoate, offering simplicity and scalability (Deng et al., 2009; Nemani et al., 2018). Benzoic acid and related compounds are regarded as generally recognized as safe (GRAS) by the FAO/WHO and FDA/USDA for use in food, cosmetics, pharmaceuticals, and hygiene products. Despite being considered non-harmful, some adverse effects have been reported. Benzoic acid is slightly irritating to the skin and eyes, while sodium benzoate is only irritating to the eyes (Del Olmo, Calzada, & Nuñez, 2017). While light-based 3D printing holds transformative potential in biofabrication, its success
hinges on the development of biocompatible, photoactive inks. A key requirement is the integration of non-toxic photoreactive components, such as radical-generating photoinitiators. In Norrish Type II reactions, UV-excited carbonyl compounds undergo γ-hydrogen abstraction to form 1,4-diradical intermediates, leading to α,β-bond cleavage and subsequent enol-keto tautomerization (Norrish & Bamford, 1936; Chatani et al., 2014). However, sodium benzoate, lacking γ-hydrogens, operates via a Norrish Type I mechanism: UV irradiation induces photolytic cleavage, generating benzoyloxy radicals (Ar–COO•) that decarboxylate into phenyl radicals (Ar•) and CO2. These radicals initiate intermolecular hydrogen abstraction from biopolymer hydroxyl groups, driving crosslinking via recombination (Zhao et al., 2013; Delville et al., 2002; Wells, 1972). This process aligns with green chemistry principles, yielding benign byproducts (CO2, benzoic acid derivatives) while eliminating structural constraints of γ-hydrogen dependence. Photopolymerization offers distinct advantages over conventional methods, including rapid curing under low-intensity light, minimal thermal output, and compatibility with aqueous, physiologically relevant conditions—features critical for cell-laden bioprinting (Lim et al., 2020). Radiation-induced crosslinking, applied during or post-printing, enhances biomaterial functionality by reinforcing mechanical properties (Choi, 2019).

Despite progress in hydrogel fabrication, unmet challenges persist in achieving rapid-curing, self-supporting viscoelastic materials for high-fidelity 3D printing. To address this, we developed a dual-wavelength UV crosslinking platform (390 nm and 260 nm, 5 mW/cm² intensity) for direct printing of rigid, UV-curable alginate-gellan blends. Building on prior multi-material printing strategies (Asghartabar Kashi et al., 2024), our system integrates a screw-driven syringe extruder for conductive paste deposition, a vacuum-assisted nozzle for precision component placement, and servo-controlled micro-z-stages for layer-by-layer assembly.
By leveraging wavelength-dependent curing kinetics, we demonstrate accelerated print speeds and structural precision, advancing the development of artificial biofilms that mimic the hierarchical complexity and functional resilience of natural systems.
Accordingly, we created a photo-curable porous hydrogel specifically designed for 3D printing through the UV-irradiation crosslinking of a biofilm mimic. We utilized a Norrish type I photoinitiator during the UV exposure to develop a crosslinked hydrogel ink that served as the basis for our 3D printing efforts. By incorporating sodium benzoate into a mixture of alginate and gellan, we produced a photo-curable ink capable of generating a functionalized scaffold. To evaluate the ink's effectiveness for 3D printed scaffolds, we thoroughly examined different preparation conditions, emphasizing various rheological characteristics as well as mechanical and thermal properties, supported by structural characterization methods.

## 2. Methods and Materials

### 2.1. Materials

Sodium alginate (CAS: 9005-38-3, MW = ~ 80,000 g mol⁻¹, M/G ratio = ~ 1:2) was supplied by Sigma-Aldrich (Steinheim, Germany). Low-acyl gellan gum (CAS: 71010-52-1, purity > 99%) has an acyl content of 3.5% and a carboxylate content of 25.4%, sourced from Gelrite (Carl Roth GmbH, Germany). The average molecular weight of low acyl gellan gum is approximately 1 × 10⁶ Da. Sodium benzoate (CAS: 532-32-1, purity > 99.7%) was also provided by Sigma-Aldrich (Steinheim, Germany) (Section S1, Supplementary Materials). All chemicals were used as received without further purification.

### 2.2. Preparation of alginate/gellan biofilm-inspired printing ink

Alginate (2 wt%) and gellan (1 wt%) stock solutions were prepared separately in deionized water under controlled shear of 56 s⁻¹ (G-force of 367) using a high-speed homogenizer (IKA T18 Ultra-Turrax, Staufen, Germany) at 60°C for 120 min. The binary blend (alginate: gellan = 2:1 w/w) was homogenized, pH-adjusted to 5.2 (to optimize carboxylate reactivity), and stored at 4°C. Sodium benzoate was incorporated into the alginate/gellan blend as a photoinitiator, with stirring at 60°C for 80 min to ensure molecular dispersion. Different concentrations of sodium benzoate (5–30 wt%) were systematically tested to identify optimal
crosslinking efficiency under UV irradiation (Section S2, Supplementary Materials). Optimization revealed that 20 wt% sodium benzoate produced the highest crosslinking density (Fig. S1, Supplementary Materials) and was selected for all experiments.

## 2.3. Custom-built 3D printing system

In this section, we discuss the components of the custom-built UV crosslink 3D printing machine that is designed to mimic the functional properties of biofilms in spatial patterns. Due to the limitation of suitable commercially available printers for multimaterial printing, along with accomplishing faster building speeds and improved durability in the printed structures, we chose to build our own 3D printer. The design of our printer is based on the open-source D-Bot printer with slight modifications (Flowers et al., 2017), and has a total parts cost of under $1000 (Fig. S.2, Supplementary Materials).

The original design, which is described by a detailed build guide and bill of materials, was converted from a Bowden extruder to a syringe-like extruder. This custom-built 3D printer consists of a 3-axis gantry system, where the x and y axes are mounted at a fixed height, and the print-bed serves as the z-axis. The carriage is equipped with a standard syringe-like extruder for paste-link inks, a screw-driven syringe extruder for conductive paste, a pivot-mounted vacuum nozzle for pick and place handling of surface-mount device (SMD) components, and a specific UV chamber. The extruder, coupled with the UV chamber, is mounted on individual, servo-controlled, micro-z-stages and can be positioned below or above the extruder to avoid collisions. We use consumer-grade polylactic acid and Acrylonitrile butadiene styrene materials for the plastic part and a silver-filled polymer ink (#6130F, Methode Electronics Inc.) to print the wires.

Specifically, we replaced a Bowden extruder system with a syringe-like extruder. This is because the Bowden extruder pushes filament through a long tube before it reaches the nozzle, which frequently results in softer filaments bending and getting jammed inside the tube due to the back pressure from the nozzle. A syringe-
like extruder system, in which the extrusion motor sits directly above the nozzle, pushes on a piece of paste-like samples inside the syringe so that bending and jamming do not occur (Fig. S.2, Supplementary Materials). A syringe-like extruder system has enough weight to extrude viscoelastic materials typically used in low-cost 3D printers from a nozzle tip of up to ~800 µm, but the aluminum extrusions used in the D-Bot printer can easily accommodate the increased mass. The print projection system used in this work was a digital projector projected via UV lamps through a steel hexagonal chamber. The projected patterns are constructed such that a UV dose distribution corresponding to the desired print emerges in the inks after an integer number of rotations. The printer operated with two UV irradiation sources, which contained a 390 nm and 260 nm LED light source, respectively. Their intensity amounted to 8 mW/cm². The distance of the light source from the solution was 1 cm with an angle of 90°. The building area of the printer was 80 × 50 × 150 mm³ with a maximum filling of 700 mL. The layer thickness could be adjusted between 10 and 250 µm. The vat could be heated up to 60 °C, but in this study, the printing was carried out at room temperature.

To control the 3D printer hardware, we utilized OctoPrint (Häußge, 2020) as a print server. We developed an OctoPrint plugin called OctoPNP, which integrates vision-based pick-and-place operations into the standard printing workflow. Each component's shape and target position are incorporated into the G-code format through a unique extension. Upon loading a G-code file, OctoPNP retrieves this additional information and assigns a specific component tray box for every component. For the design of the 3D-printed objects, we employed AutoCAD software, exporting the final design as an STL file (AutoCAD; Autodesk, Cupertino, CA, USA). The XYZ pathway instructions required by the printer were generated from the STL file using the open-source CAM software Cura™ (Version 2.3.1; Ultimaker, Geldermalsen, Netherlands). Table S3 (Section S5 in Supplementary Materials) presents the settings employed during the tests for printability and the printing of hydrogel-forming inks. Each printed object consisted of 30 layers, with the height of the nozzle raised by
11 mm after each layer. The air pressure for ink extrusion was determined experimentally, typically ranging from 52 × 10⁵ Pa (22 psi) to 1.72 × 10⁵ Pa (25 psi). Surface photo-crosslinking was performed using various UV exposure times, ranging from 0 to 8 min.

## 2.4. Fourier-Transform Infrared Spectroscopy (FT-IR)

The transmission infrared spectra of 3D printed samples were analyzed using an FT-IR spectrometer (Jasco FT/IR6200, Tokyo, Japan) to verify the crosslinking process. The solid samples required for the FT-IR analysis were prepared in pellet form by mixing approximately 10 mg of each sample with 100 mg of dry potassium bromide (KBr). These mixtures were then pressed into pellets before scanning across a spectral range of 800 to 3800 cm⁻¹, with a total of 50 scans recorded at a resolution of 1 cm⁻¹.

## 2.5. Capillary Viscometer

The coil-overlap concentration (c*) for alginate, gellan, and 3D printed samples based on the alginate/gellan blend was measured using a Cannon-Ubbelohde capillary viscometer (model 9721-K56, size 75, Cannon Instruments Co., Germany), which has a constant k of 0.007690 mm² s⁻². This measurement was conducted with the aid of a thermostatic water bath maintained at a controlled temperature of 25 ± 1 °C. Various concentrations of each sample (ranging from 0.1-5 g/dL) were prepared by dissolving them in DMSO, ensuring coverage from the dilute to the semi-dilute regions. The solvent was created by mixing 9 ml of DMSO with 1 ml of deionized water (90% v/v), followed by stirring at 60 °C for 30 minutes to achieve a uniform solution. For testing, precisely 2 mL of sample solution was diluted, and at least three efflux time readings were recorded for each concentration. The relative viscosity (ηrel), reduced viscosity (ηred), and inherent viscosity (ηinh) were subsequently calculated using equations (1)-(3), respectively:

$$ \eta_{rel} = \frac{\eta}{\eta_0} = \frac{t}{t_0} \cdot \frac{\rho}{\rho_0} \qquad (1) $$
$$ \eta_{\text{red}} = \frac{(\eta_{\text{rel}} - 1)}{c} \qquad (2) $$

$$ \eta_{\text{inh}} = \frac{\ln(\eta_{\text{rel}})}{c} \qquad (3) $$

In these equations, $t/t_o$ and $\rho/\rho_o$ denote the ratios of the efflux time of the solution to that of the solvent and the densities of the solution and solvent, respectively. Additionally, the $(\eta_{rel}-1)$ corresponds to the specific viscosity $(\eta_{sp})$, and c represents the concentration of the biopolymers.

## 2.6. Macromolecular Characteristics via SEC-MALLS

The macromolecular properties of each biosurfactant, including molecular weight, polydispersity index, and radius of gyration, were assessed using size exclusion chromatography (SEC) paired with a multiangle laser light scattering (MALLS) photometer (DAWN EOS, Wyatt Technology Co., Santa Barbara, CA). A refractive index of 1.33 was applied during the analysis. The SEC setup utilized an Ultrahydrogel™ column (7.8 mm × 300 mm, Waters Co., Milford, MA), and an Optilab refractometer (Dawn, Wyatt Technology Co., Santa Barbara, CA) was connected for simultaneous data acquisition. Each experiment involved an injection volume of 75 mL with a flow rate set at 1.5 mL/min. The refractive index increment ($dn/dc$) for the samples was determined using the Optilab refractometer ($\lambda$ = 690 nm) at 25 °C, resulting in a value of 0.140 mL/g.

## 2.7. Rheological measurements

The rheological properties of the 3D printed samples were evaluated using an AR-G2 rheometer (TA Instruments, West Sussex, UK) equipped with a heating circulator (Julabo, F-12, Seelbach, Germany). Each sample was positioned at the center of the Peltier plate and allowed to equilibrate for 10 min to achieve a stable temperature. All assessments were performed at ambient temperature using parallel plate geometry with a 40 mm diameter. A small-amplitude oscillatory shear (SAOS) experiment was conducted to identify the limit of the linear viscoelastic region (LVR). Strain sweep tests were performed at a frequency of 1 Hz to
assess the oscillatory shear linear viscoelastic range specific to each sample. Moreover, frequency sweep tests were conducted on the inks over a range of 0.1 to 100 Hz while maintaining a constant shear strain of 1% (γ = 1.0%) within the linear regime. The rheological parameters, including the storage modulus (G') and the loss modulus (G''), were acquired through the manufacturer's software (TRIOS, TA Instruments, West Sussex, UK) (Chakraborty et al., 2022; Segovia-Gutiérrez et al., 2024).

## 2.8. Large amplitude oscillatory shear (LAOS) test

The large amplitude oscillatory shear (LAOS) test was performed to examine the viscoelastic characteristics of samples within the non-linear viscoelastic (NLVE) regime. During the testing, the printing inks were subjected to a logarithmic ramp strain amplitude sweep, varying from 1% to 1000%, while maintaining a constant frequency of 1 Hz and a temperature of 20 °C. Viscoelastic behaviors in the NLVE were recorded as a function of intercycle strain (γ₀), with shear stress (σ) assessed concerning intracycle strain (γ) and strain rate. To analyze the non-linear rheological response of the inks, Lissajous plots, commonly referred to as Lissajous-Bowditch plots, were constructed. The methodology for this analysis was thoroughly explored by Ewoldt et al. (2010). Variations in the Lissajous plots provide direct insights into the distinct NLVE properties of the inks. The behaviors reflecting intracycle strain stiffening and intracycle shear thickening can be quantified using an S-factor and a T-factor, respectively, as defined in Eqs. (4) and (5).

$$S = (G'_{L} - G'_{M}) / G'_{L} \qquad (4)$$

$$T = \eta'_{L} - \eta'_{M} / \eta'_{L} \qquad (5)$$

In these equations, parameters such as the large-strain elastic modulus (secant modulus), the shear elastic modulus at minimum strain (tangent modulus at zero strain), the viscosity at the maximum shear rate, and the viscosity at the minimum shear rate are represented. Specifically, an S value of 0 denotes linear elastic
behavior, while $S > 0$ indicates intracycle strain stiffening, and $S < 0$ signifies intracycle strain softening. Similarly, a $T$ value of 0 corresponds to linear viscous behavior, with $T > 0$ indicating intracycle shear thickening and $T < 0$ denoting intracycle shear thinning.

## 2.9. Mechanical properties of 3D structures

The mechanical characteristics were evaluated under a pure shear profile using a Zwick Roell Testing Machine (Z005, Zwick/Roell, Ulm, Germany) equipped with a 100 N load cell. Rectangular specimens measuring 60 mm in length, 10 mm in width, and 4 mm in height were secured between two clamps, and uniaxial deformation was conducted at a crosshead speed ranging from 1.44 to 1140 mm/min, which corresponded to a strain rate ($\dot{\epsilon}$) of 0.0012 to 0.92 s$^{-1}$. The nominal stress ($\sigma$) was calculated based on the force applied over the cross-sectional area of the specimen before deformation, while strain ($\epsilon$) was determined by the change in length compared to the original length. The Young's modulus ($E$) was derived from the slope of the initial linear portion (0–50% strain) of the resulting stress–strain curves.

For the cyclic testing, each specimen was elongated to a predetermined strain and then returned to its original length at the same velocity. Hysteresis was assessed by measuring the area enclosed by the loading and unloading curves. Each sample underwent four measurements, and the reported results reflect the average of these trials.

In the tearing test for evaluating the tearing energy of gels, a standardized trousers test was employed. A pre-existing crack, approximately 5 mm long with a sharp tip, was created using a fresh razor blade. The specimens, with two arms, were securely clamped and aligned perpendicularly to the clamps, facilitating crack propagation along the specimen's centerline. With the lower clamp fixed, the upper clamp was pulled at a constant rate until the crack extended fully through the specimen, allowing for the recording of the stretched displacement ($L$) and tearing force ($F$). The tearing energy ($\Gamma$) was calculated using the formula $\Gamma$
= 2F/w, where F denotes the tearing force during crack propagation and w represents the specimen's thickness (Yang, Ma, Zhang, & Xu, 2016).

## 2.10. Microstructure evaluations of 3D structures

The 3D printed objects were analyzed using a field-emission scanning electron microscope (FE-SEM, S-4700, Hitachi, Japan) to obtain high-resolution images with a greater depth of field. Firstly, each 3D construct was trimmed down to a precise size of (15 × 15 × 15) mm³. The 3D printed samples were sectioned and then placed on a stage cooled by Peltier technology, which maintained a temperature of – 10 °C to prevent thermal damage. The image gas nitrous oxide was used at a pressure of 50.7 Pa for capturing the microstructures of each 3D construct. An accelerating voltage of 20 kV was applied to obtain the microstructure images through a solid-state backscatter detector.

## 2.11. Statistical analysis

In the wake of the above-mentioned procedure, measurement analysis of a completely randomized design was carried out with the analysis of variance (ANOVA) followed by Duncan's test procedure in SPSS software (Version 19, SPSS Inc., Chicago, IL). A value of *p* ≤ 0.05 was considered statistically significant. The measurements of all the experiments were analyzed in triplicate, and the mean and standard deviations of the data were reported.

# 3. Results and discussion

## 3.1. Network design and photocuring kinetics: FT-IR and ¹H NMR study

Sodium benzoate, a Norrish Type I photoinitiator, facilitates hydrogel crosslinking via UV-induced decarboxylation. This generates phenyl radicals (Ar•), which abstract hydrogen atoms from hydroxyl (-OH) and carboxyl (-COOH) groups on alginate and gellan gum chains. The resulting polysaccharide macro-radicals recombine to form covalent crosslinks, creating a stable 3D network (Zhou et al., 2009; Deng et al.,
2009; Nemani et al., 2018). Upon exposure to UV light, these photoinitiators undergo decomposition, generating reactive radicals that initiate a photocrosslinking mechanism through a radical pathway, primarily via hydrogen abstraction from the substrate. This process significantly enhances the efficiency of UV-induced crosslinking in the modification of polymeric materials. In our study, we specifically investigated how the implementation of two distinct photoreactions activated by different wavelengths of light could lead to localized control over mechanical properties during dual-wavelength 3D printing. To that end, we utilized a dual-wavelength printing prototype developed by our team, which is equipped with two high-intensity LED light sources operating at 390 nm and 260 nm, each providing an intensity of 8 mW/cm². This configuration facilitates layer-by-layer printing using these two distinct wavelengths, as illustrated in Fig. 1a,c.

In this investigation, we employed a photocurable blend of alginate and gellan that undergoes free radical crosslinking of a photosensitizer via photolysis when exposed to UV light. According to the results of rheological measurements and printability results (data not shown), the alginate/gellan with a ratio of 0.5/0.5 wt.% offered a superior blend system with the highest mechanical properties and the greatest tensile loading-unloading curve. It is important to note that the two reactions are not entirely orthogonal since the utilized radical photoinitiator (sodium benzoate) absorbs light in both the UV-A and UV-C regions (the absorption spectra for the photoinitiator can be found in Fig. S3 (section S4 of the Supplementary Materials). Consequently, dual-wavelength UV irradiation (UV-A at 390 nm and UV-C at 260 nm) enables spatially controlled crosslinking of the alginate/gellan blend. Under UV light, sodium benzoate undergoes homolytic cleavage of the C=O bond, generating radicals. Sodium benzoate's carbonyl group absorbs strongly at 260 or 390 nm, making this pathway dominant over Type II (which requires γ-H abstraction, absent here). Moreover, oxygen-free conditions prevent radical quenching, ensuring efficient initiation. Sodium benzoate is not incorporated into the network—it regenerates post-reaction (Fig. 1b).
<figure><img src="image_3.png" /><figcaption>**Fig. 1.** (a) Schematic diagram and (c) image of the dual-wavelength UV crosslinker prototype functioning at 390 and 260 nm. (b) Schematic and curing mechanism illustrations of the wavelength-selective formation of soft and rigid domains. Soft materials are produced through radical-induced crosslinking of photosensitizer via photolysis when exposed to UV irradiation. Exposure to 260 nm results in rigid structures due to the development of a highly crosslinked interpenetrating photopolymer network, involving simultaneous radical graft crosslinking of alginate/gellan. (d) Printing performance of different 3D structures after 6 weeks at 37°C/80% humidity.</figcaption></figure>

Among photoinitiators, sodium benzoate has gained prominence for hydrogel crosslinking due to its biocompatibility and efficient radical generation under UV irradiation. Sodium benzoate operates via a Norrish Type I mechanism, where UV irradiation (λ = 254–390 nm) is absorbed by the carbonyl (C=O) chromophore, inducing homolytic cleavage of the C=O bond (Zhou et al., 2009; Deng et al., 2009; Nemani et al., 2018). This generates a benzoyloxy radical (Ar–COO•), which rapidly decarboxylates to yield a phenyl radical (Ar•) and CO₂. The phenyl radical abstracts hydrogen atoms intermolecularly from hydroxyl (-OH) and carboxyl (-
COOH) groups on polysaccharide chains (alginate/gellan), initiating crosslinking via radical recombination (Delville et al., 2002; Wells, 1972; Kashi et al., 2024; Ihmels & Scheffer, 1999).

This distinction is critical: Norrish Type II reactions are restricted to carbonyl compounds with a γ-hydrogen in a four-carbon chain (e.g., ketones like benzophenone), enabling intramolecular 1,4-diradical formation. Sodium benzoate, with its aromatic carboxylate structure, cannot undergo this pathway. Its efficacy as a photoinitiator arises from the Type I cleavage and intermolecular hydrogen abstraction, a mechanism well-documented in biopolymer-based systems (Ghosh & Gangopadhyay, 2000; Delville et al., 2002). Despite this reaction's efficacy, the curing process often faces diffusion limitations, leading to slow gelation times that can result in irregular network structures and significant shrinkage stress in photopolymers. This can lead to the development of brittle networks characterized by low fracture toughness, especially when high crosslink densities are utilized. To address this issue in the design of printable photocurable biopolymers, our group (Asghartabar Kashi 2024) found that reducing crosslink density could mitigate brittleness. They specifically developed a UV-irradiated photocurable hydrogel by incorporating sodium benzoate into a photo-curable PVA hydrogel with a bifunctional acrylate crosslinker, resulting in soft polymer networks with favorable swelling properties and enhanced tensile strength. While commercial gellan and alginate may contain trace salts (e.g., Na⁺, Ca²⁺) from production, their low concentration (<0.1 wt% in final hydrogels) ensures no interference with covalent photocrosslinking. Control experiments confirmed identical crosslinking efficiency in dialyzed vs. undialyzed materials (Data not shown). Moreover, the minimal concentration of proteins (less than 0.9%) in gellan gum allows us to regard their effects as negligible in the context of our experiments (Table S1, Supplementary Materials). Importantly, since the gellan content was kept constant throughout all treatments, the primary variable affecting the crosslinking reaction was the duration of UV irradiation. Moreover, the presence of sodium benzoate serves a crucial role in initiating the photocrosslinking mechanism through the generation of free radicals without directly participating in the crosslinking reaction
itself. This dual consideration of impurities and the integral role of sodium benzoate enhances the rigor of our hypotheses and interpretations, supported by structural and rheological characterizations, thereby providing a more comprehensive understanding of the photocrosslinking process.

Hybrid systems require optimized polymers/photosensitizers miscibility and photo-crosslinking kinetics to control phase separation and network morphology, which critically determine the mechanical properties and printability of fabricated objects (Fig. 1d). We first evaluated the orthogonality of radical photoreactions under different irradiation wavelengths and durations using FT-IR (Fig. 2, Top), ¹³C NMR (Fig. 2, Bottom), and ¹H NMR (Fig. 3). The FT-IR spectra of precursor materials (Fig. S4, Supplementary Material) established critical baselines. Neat alginate showed characteristic bands at 3,463 cm⁻¹ (O–H stretch), 2,987 cm⁻¹ (C–H stretch), 1,594 cm⁻¹ (asymmetric COO⁻ stretch), 1,405 cm⁻¹ (symmetric COO⁻ stretch), and 1,016 cm⁻¹ (C–O stretch) (Cao et al., 2024). Gellan gum exhibited analogous vibrations at 3,351 cm⁻¹ (O–H bend), 2,894 cm⁻¹ (C–H stretch), 1,593 cm⁻¹ (carbonyl/acylated groups), 1,410 cm⁻¹ (COO⁻ stretch), and 1,032 cm⁻¹ (C–O stretch) (Cao et al., 2024). Sodium benzoate displayed diagnostic peaks at 3,466 cm⁻¹ (O–H), 3,051 cm⁻¹ (aromatic C–H), 1,649 cm⁻¹ (asymmetric COO⁻), 1,458 cm⁻¹ (symmetric COO⁻ + C=C), and 778 cm⁻¹ (monosubstituted C–H bend), confirming its aromatic carboxylate structure (Kumar et al., 2017).

Alginate/gellan and alginate/gellan/benzoate blends (UA0/UC0) retained these signatures but revealed consequential shifts (Fig. S4 in Supplementary Material), which contain an alginate COO⁻ peak shift from 1,454 cm⁻¹ to 1,445 cm⁻¹ and O–H stretch decrease from 3,463 cm⁻¹ to 3,395 cm⁻¹, indicating altered hydrogen-bonding environments (Naji-Tabasi et al., 2023; Thinkohkaew et al., 2024). Photo-crosslinked samples exhibited three definitive spectral modifications in Fig. 2 (Top). This includes the appearance of a new peak at 1,437 cm⁻¹ corresponding to sodium benzoate's aromatic C=C stretch (Fig. S4, in Supplementary Material). This confirms UV-induced radical generation from the photoinitiator (Kumar, 2014).
Moreover, a COO⁻ stretch shifted from 1,614 cm⁻¹ to 1,600 cm⁻¹, consistent with covalent bond formation altering carboxylate electronic environments during crosslinking (Zain et al., 2018). Finally, a significant reduction of the characteristic band at 3,395 cm⁻¹ indicates the hydroxyl consumption, collectively demonstrating successful intermolecular crosslinking (Zhang et al., 2017; Zain et al., 2018).

<figure><img src="image_4.png" /><figcaption>**Fig. 2.** (Top): FT-IR spectra of 3D printed alginate/gellan samples treated with UV-A (a) and UV-C (b). (Down): Solid-state ¹³C NMR spectra of unmodified sodium alginate, gellan gum, unirradiated alginate/gellan/benzoate blend (UA0/UC0), and UV-irradiated crosslinked samples (UA8/UC8).</figcaption></figure>

Solid-state ¹³C NMR (see Section S7 for experimental details) provided complementary molecular evidence, with unmodified alginate showing resonances at δ = 177 ppm (C6 carboxylate), δ = 101–104 ppm (C1 anomeric), and δ = 70–84 ppm (C2–C5 ring carbons) (Fig. 2, Bottom). The absence of signals below δ = 65
ppm is consistent with alginate's exclusive composition of uronic acids, which lack aliphatic $-CH_2OH$ groups. Gellan gum featured additional peaks at $\delta = 64$ ppm (glucosyl C6) and $\delta = 18$ ppm ($rhamnosyl CH_3$), while unirradiated blends (UA0/UC0) displayed superimposed benzoate signals at $\delta = 128-136$ ppm (aromatic carbons). UV-irradiated samples exhibited transformative changes, including complete disappearance of benzoate aromatic resonances, emergence of a new alkyl carbon peak at $\delta = 25$ ppm assigned to covalently incorporated benzoate fragments, reduced intensity at $\delta = 70-84$ ppm indicating restricted chain mobility, and retention of the $rhamnosyl CH_3$ signal at $\delta = 18$ ppm confirming gellan's structural integrity. These observations collectively demonstrate radical-mediated ether bond formation at C-OH sites (notably glucosyl C6) and covalent grafting of benzoate fragments into the polymer network.

$^1$H NMR spectroscopy (Fig. 3, Left) revealed neat alginate peaks at $\delta$ 4.53 ppm (H1 of $\beta$-D-mannuronic acid) and $\delta$ 5.12 ppm (H1 of $\alpha$-L-guluronic acid). Additionally, Peaks at 3.89 ppm and 4.09 ppm correspond to H4 protons of mannuronic (M) and guluronic (G) acid units, respectively (Makhado et al., 2017). Gellan gum showed characteristic signals at 5.05 ppm ($rhamnose H1$, R1), 4.28 ppm (glucuronic acid H1, GA1), 4.40 ppm (glucose H2, G21), and 1.14 ppm ($rhamnose -CH_3$, R6). Upon blending alginate with gellan (without sodium benzoate or irradiation process), the spectrum remains largely unchanged, though minor peaks near $\delta$ 1.1 and $\delta$ 2.0 ppm emerge, likely reflecting $-CH_3$ protons from gellan's acetylation motifs (Kamer et al., 2021). Sodium benzoate introduction (UC0/UA0) produced a distinct aromatic signal at $\delta$ 7.25-7.88 ppm (Fig. S5), with quantitative HPLC analysis (Section S12) establishing pseudo-first-order depletion kinetics ($k = 0.42 \text{ min}^{-1}$, $R^2 > 0.99$).

In UV-treated samples (UA2/UC2 and UA4/UC4), the aromatic peak persists at $\delta$ 7.25-7.88 ppm with reduced intensity, corresponding to 58% and 22% residual sodium benzoate respectively per HPLC quantification (Fig. 5b). By contrast, UA6/UC6 and UA8/UC8 exhibit no detectable aromatic protons,
indicating complete sodium benzoate consumption (<5% residual) at higher irradiation times, as conclusively confirmed by HPLC (LOD = 0.2 mg/L). Critically, the alkyl peak increase is diagnostic of benzoate fragment incorporation, while the static δ 1.14 ppm signal (gellan rhamnose –CH₃) confirms backbone integrity during curing. Electron spin resonance (ESR) spectroscopy (Section S12) also detected no persistent radicals post-curing; the inverse correlation between aromatic depletion and alkyl growth directly evidences hydrogen abstraction; and the absence of diagnostic peroxy ester IR bands (1710-1770 cm⁻¹) rules out alternative pathways. These convergent findings support a radical-mediated crosslinking mechanism (Fig. 3, Right) where UV dissociation generates transient radicals that enable intermolecular bonding through recombination and termination (Zain et al., 2018; Cao et al., 2024; Kumar, 2014).
<figure><img src="image_5.png" /><figcaption>**Fig. 3.** (Left): ¹H NMR spectra of pure Na-alginate and neat gellan, as well as treated photo-crosslinked samples. (Right): Formation of reactive radicals from sodium benzoate and potential mechanism in photo-cross-linking.</figcaption></figure>

## 3.2. Coil-overlap and SEC-MALLS results

A comprehensive analysis of the viscosity for various samples across the full concentration spectrum—including dilute and semi-dilute ranges—was conducted by plotting the specific viscosity ($\eta_{sp}$) against the concentration of samples on a log-log scale. The curve distinctly reveals two prominent segments, each characterized by different slopes (Figs. 4a,b). In the dilute region (when $c < c^*$), the initial slope exhibited a linear trend, with a slope between approximately 0.6 and 1.8, indicating that the individual polymeric chains exist as separate coils. As the concentration of samples increased, there was a significant rise in specific
viscosity, corresponding to the entanglement of the separate chains. This transition in the conformation of
biopolymers at the coil-overlap concentration was observed, for example, the c* values for UVA2 and UVA8,
as well as UVC2, UVC8, being 1.22, 0.63, 1.14, and 0.41 g L⁻¹, respectively (Table S4 in section S8
Supplementary Materials). At concentrations below c*, the individual coils remain largely separate, resulting
in minimal mutual interference and demonstrating Newtonian flow characteristics (Morris et al., 1981). In
contrast, as the concentration exceeds c* (c > c*), the entanglement of biopolymeric chains becomes more
pronounced, shifting the dynamics of chain relaxation from Rouse/Zimm behavior to a regime described by
the reptation model (Doi & Edwards, 1988). In the semi-dilute range, the entanglement of the surface-active
biopolymer is primarily driven by chain-chain interactions. Consequently, the entangled polymer chains,
influenced by reptation motions, lead to conformational changes due to enhanced interchain interactions (Doi
& Edwards, 1988). This phenomenon can significantly impact certain functional properties of system-
containing samples, such as mechanical, rheological, thermos-mechanical, and printing performance,
indicating that systems with lower c* values could experience more pronounced changes in rheological and
mechanical characteristics.
<figure><img src="image_6.png" /><figcaption>Fig. 4. Specific viscosity plotted against the concentration on a log-log scale (a,b). SEC-MALLS chromatograms (c,d) of different 3D printed objects.</figcaption></figure>

Further understanding of the macromolecular characteristics was achieved by assessing the molecular weights of different sample variants utilizing SEC-MALLS (as illustrated in Figs. 4c,d). Size exclusion chromatography effectively separates biopolymers based on their hydrodynamic volume, which relates to the length, swelling, and hydration state of the biopolymeric chains. Figs. 4c,d presents the elution profiles corresponding to the molecular weight distributions of the various samples as determined via SEC-MALLS. The pertinent molecular parameters for all samples are compiled in Table S4 (section S8 Supplementary Materials). Notably, an increase in retention time was observed for samples with a lower irradiation time, indicating that these samples are retained longer within the column due to a decrease in their hydrodynamic volume (Yang & Zhang, 2009). As a result, UA2/UC2 exhibited lower average molecular weights (*Mw*) in contrast to UA8/UC8 (Figs. 4c,d). The data derived from SEC-MALLS corroborated findings from previous coil-overlap experiments. The polydispersity index (*DP*), defined as *Mw/Mn*, is a critical parameter for
evaluating the distribution of molecular mass within a biopolymer. A higher *DP* signifies a broader molecular weight distribution, which could lead to a less effective gel network and diminished structural integrity. According to Table S4 (section S8 Supplementary Materials), the *DP* ranged from 1.93 to 3.15, with UA0 showing the highest *DP* and UC8 the lowest. An elevated *DP* combined with a reduced Mw can contribute to decreased elastic modulus with lower tensile strength, which may correspond with lower printing performance of samples. Additionally, Table S4 (section S8 Supplementary Materials) summarizes data on the radius of gyration (*R*<sub>g</sub>) for the sample variants. Generally, *R*<sub>g</sub> reflects the mass-weighted (root mean square) distance of each point in a molecule from its center of gravity, related to the spatial occupation of the chain (Yang & Zhang, 2009). UC8 displayed the highest *R*<sub>g</sub> at 65 nm, followed by UA8 at 55 nm, while the *R*<sub>g</sub> values for UA2 and UC2 were recorded at 31 and 32 nm. This indicates that UC8 and UA8 may contribute to more structured 3D-printed samples.

### 3.3. Quantitative crosslinking assessment in printed irradiated hydrogels

Fig. 5a illustrates variations in equilibrium swelling ratio (ESR) and gel fraction (GF) for hydrogels irradiated at 260 nm (UV-C) and 390 nm (UV-A) (see Supplementary Material, Section S2, for experimental details). Irradiated samples exhibited significantly reduced water uptake, with ESR values decreasing by up to 97% (UC8: 0.09 ± 0.01) compared to non-irradiated controls (3.8 ± 0.2). This confirms enhanced photo-crosslinking and reduced hydrophilicity, attributed to covalent intermolecular network formation between alginate and gellan chains, consistent with NMR-detected alkyl fragment grafting (δ 1.1–2.0 ppm). Critically, the radical-mediated mechanism was validated by quenching experiments with 2.0 wt% TEMPO (a characteristic radical scavenger). The TEMPO-containing formulations failed to develop a crosslinked network under both UV wavelengths, exhibiting ESR >0.048 and GF < 0.02%–values indistinguishable from non-irradiated control (Fig. 5a).
The results indicated that increasing irradiation time (in samples with sodium benzoate) led to a dose-
dependent effect on crosslinking, as evidenced by a reduction in ESR and an increase in GF, with the UV-C
treatment showing more than 40% greater GF in UC8 compared to the control. Specifically, the UV-A samples
demonstrated a gradual decline in ESR until 6 min (UA6: 0.35 ± 0.03), which was followed by a sharp
decrease to 0.23 ± 0.02 at 8 min (UA8), a threshold that corresponds to complete depletion of sodium
benzoate (HPLC >95%) and indicates radical saturation for crosslinking (Fig. 5b). Furthermore, the UV-C
irradiation produced a significantly higher crosslinking density, as reflected by the ESR values (0.09 ± 0.01
for UC8 vs. 0.23 ± 0.02 for UA8), resulting in about 60% higher GF in UC8. The UC8 sample's 76% ESR
reduction and >60% GF increase quantitatively validate covalent network formation. This contrasts sharply
with TEMPO-quenched systems, which showed no significant ESR/GF changes (p > 0.05 vs. control) (Fig.
5a). These crosslinked networks restrict polymer chain mobility, diminishing water absorption affinity (ESR <
0.35) while enhancing mechanical stiffness, critical for maintaining structural integrity during layer-by-layer
printing deposition. According to Fig. 5c, TEMPO-containing inks (Fig. 5c, iii and iv) exhibited immediate layer
spreading (>300% area expansion) and <30% shape retention after deposition, while the sodium benzoate-
based system (Fig. 5c, ii) achieved dimensional stability with ≤100 µm resolution and >90% geometric
integrity.
<figure><img src="image_7.png" /><figcaption>**Fig. 5.** (a): Crosslinking density quantification: Equilibrium swelling ratio (ESR, dimensionless, solid symbols) and gel fraction (GF, %, hollow symbols) for UV-A and UV-C irradiated hydrogels. Error bars = ±SD (n=3). The green triangle shows the samples that contained TEMPO with no sodium benzoate. (b): Analytical characterization of residual sodium benzoate in hydrogels by HPLC. Chromatographic quantification methodology: Comparative HPLC chromatograms of UC2 (i) and UA2 (ii) samples. UC8 (iii) and (iv) UA8 hydrogels post-irradiation. UV-Vis spectra of analytes (λmax = 260 nm). (c): Structural fidelity comparison of 3D-printed hydrogel lattices using different radical sources. (i) Target CAD model of cubic lattice (800 µm features). Printed structures: (ii) Sodium benzoate-based UC8 without TEMPO, (iii) UA8 with TEMPO (260 nm irradiation), (iv) UC8 with TEMPO (390 nm irradiation). Scale bars = 5 cm. (d): Cytotoxicity assessment protocol for hydrogel extracts. Fibroblast suspensions (2 × 10⁴ cells/mL) exposed to extracts for 3 h, 18 h, and 36 h.</figcaption></figure>

Radical quenching with 2.0 wt% TEMPO abolished network formation under both UV wavelengths, as quantitatively evidenced by unchanged ESR (0.48 ± 0.2) or GF (0.02 ± 0.001%) values during 0-8 min irradiation times (p > 0.05). Previously, the printing performance results revealed that the inks containing TEMPO exhibited immediate layer spreading upon deposition, which indicated insufficient structural cohesion for self-support, especially under 260 nm irradiation (Fig. 5c, *iv*). This instability in the network led to poor shape fidelity, with spatial resolution falling below detectable thresholds. Similarly, at a wavelength of 390
nm (Fig. 5c, iii), the TEMPO formulations failed to create cohesive geometries, resulting in unstructured constructs with poor printing quality. In stark contrast, hydrogels based on sodium benzoate, especially UC8, produced dimensionally stable, self-supporting structures with precise feature resolution, maintaining a deviation of no more than 100 µm (Fig. 5c, ii). These radical-preserved hydrogels exhibited excellent geometric integrity without cracking or distortion during printing and hydration cycles, confirming that the radicals derived from sodium benzoate are the exclusive drivers of covalent crosslinking in alginate and gellan. Consequently, this data provides direct experimental evidence supporting the radical-mediated mechanism, corroborated by spectroscopic analysis using FTIR and NMR.

### 3.4. HPLC method development and optimization:

Cinnamic acid (CA) was chosen as an internal standard due to its structural similarity to sodium benzoate and suitable UV absorption. Common internal standards (e.g., phthalate, hendecanoic acid) were unsuitable due to hydrolysis issues or poor UV response. The detection wavelength was optimized to maximize sensitivity for sodium benzoate and CA. The mobile phase composition (methanol: ammonium acetate buffer, 30:70 v/v) achieved baseline separation of sodium benzoate and CA within 8.1 min, with retention times of 3.65 min and 8.01 min, respectively. The calibration curve for sodium benzoate exhibited excellent linearity ($R^2$ = 0.9996) over 20–100 mg/L (data not shown). LODs were 0.2 mg/L for sodium benzoate and 0.5 mg/L for CA. Precision studies showed relative standard deviation (RSD) < 3% for both intra-day and inter-day analyses, indicating high repeatability. Recovery rates in spiked hydrogel samples ranged from 96.1% to 104.3%, confirming the method's accuracy and minimal matrix interference. UV spectra indicated maximum absorbance at 260 nm for both analytes (Fig. S3 in Supplementary Materials).

The developed method successfully quantified residual sodium benzoate in printed UC2/UA2 hydrogels (Fig. 5b). Compared to UA2, peak intensities of UC2 were notably lower (Fig. 5b). It is noteworthy to note that the
chromatograms for UA4 or UC4 were not provided; however, their results were similar to those of UA2/UC2.

Furthermore, the results confirmed an efficient removal of sodium benzoate during irradiation with the higher irradiation times, i.e., regarding UA6/UC6 (data not shown) and UA8/UC8 (Fig. 5b), with residual levels below 0.5 mg/g in all samples. Regarding UA6 and UC6, their chromatograms were also not provided, and their results were similar to those of UA8 and UC8. This validated HPLC protocol provides a rapid, sensitive, and environmentally friendly approach for monitoring sodium benzoate in biomedical hydrogel products, addressing safety concerns for medical or food applications.

**3.5. Cytocompatibility Profile**

All hydrogels demonstrated excellent biocompatibility, with cytotoxicity levels significantly below ISO 10993-5's 30% threshold ($p < 0.05$) (see Supplementary Material, Section S14, for experimental details). Maximum cytotoxicity was observed at 3 h exposure (13.7 ± 0.5% for UA2/UC2 formulations), decreasing to 1.9 ± 0.4% after 36 h for UA8/UC8 hydrogels (Fig. 5d). This time-dependent reduction correlated strongly with irradiation duration ($R^2 = 0.94$) (data not shown); mirroring sodium benzoate decomposition kinetics quantified by HPLC (Section 3.4). Critically, cytotoxicity values for highly irradiated samples (UA6/UC6 and UA8/UC8) were statistically indistinguishable from negative controls at all timepoints ($p > 0.05$), confirming complete elimination of cytotoxic residues. The inverse irradiation time-cytotoxicity relationship (UA2/UC2: 13.7 ± 0.5% vs. UA8/UC8: 1.9 ± 0.4% at 36 h; $p < 0.05$) validates the purification protocol's efficacy.

These direct cytotoxicity assessments address a critical gap in biocompatibility validation for biomedical and food-contact applications. The consistently low LDH release (<14% across all formulations and exposure durations) demonstrates exceptional cytocompatibility, satisfying ISO 10993-5 and EU 10/2011 safety thresholds (Standard, 2009). The observed time-dependent cytotoxicity reduction provides compelling evidence that extended UV irradiation effectively eliminates residual sodium benzoate, as corroborated by our HPLC quantification (Fig. 5b). The absence of cytotoxicity escalation beyond 18 h exposure suggests
confirmed by NMR-detected covalent crosslinks (40% increase in alkyl fragment grafting). Beyond y > 3%,
the sharp G' decline reflects disruption of labile H-bonds in UV-A networks versus covalent crosslink fracture
in UV-C systems. As outlined in the FT-IR Section, the hydroxyl groups present in alginate and gellan, along
with the symmetric stretching vibrations of the COO⁻ group, played a significant role in the reaction induced
by UV irradiation, leading to the formation of crosslinked structures. This mechanism is substantiated by the
absorption bands observed at 3395 cm⁻¹, 1600 cm⁻¹, and around 1439 cm⁻¹ in the FT-IR spectrum (Fig. 2,
Top). Therefore, this improved rheological property is mainly owing to crosslinking via hydrogen bonds, which
agrees well with FT-IR spectroscopy to offer a good indication of photo-crosslinking. It is anticipated that
these free radicals in neighboring chains are sufficiently close to facilitate crosslinking. This reinforces the
significance of forming a strong polymeric gel-like network, which likely arises from enhanced intra- and
interchain interactions within the copolymers. Such structure-rheology relationships enable rational design of
carbohydrate-based materials for applications requiring tunable viscoelasticity (e.g., injectable therapeutics,
edible scaffolds).
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** (Top) Oscillatory strain sweep plots and (Bottom) Oscillatory frequency sweep. The results are shown for UV-A (a,c) and UV-C-treated (b,d) UV-C treated 3D printed structures. Here, *G'* is solid symbols and *G''* is open symbols. The data presented in our figures are derived from the average of three or more independent measurements, with standard deviation bars indicating the variability within these measurements.</figcaption></figure>

Additionally, the mechanical properties of the 3D objects were analyzed through an oscillation frequency sweep test (Figs. 6c,d). For all samples assessed, *G' (ω)* surpassed *G'' (ω)*, confirming the elastic or gel-like characteristics of the systems. The frequency sweep results revealed that the mechanical properties of the 3D-printed biofilms were largely invariant concerning changes in oscillation frequency (Lim et al., 2020; Chakraborty et al., 2022). In this case, *G' (ω)* values consistently remained greater than *G'' (ω)* throughout the entire frequency range, indicating that the hydrogels exhibit quasi-solid behavior and possess a well-organized, interconnected network. Moreover, an increase in the irradiation duration time corresponded with
enhanced *G'* (ω) values, which underscores the development of a robust polymeric gel-like network, likely attributable to increased intra- and interchain interactions within the polymers in accordance with FT-IR results. Additionally, a strong gel-like viscoelastic structure achieved through UV-C irradiation presents clear advantages for 3D printing applications. This enhanced structure allows the printed objects to retain their shape when subjected to shear forces during the 3D printing process, thereby improving both the printability and accuracy of the final 3D-printed products.

### 3.7. LAOS evaluation

Complex polymeric systems with similar LVR can exhibit distinct nonlinear behaviors. Traditional viscoelastic moduli obtained from the first harmonic stress responses may not accurately reflect the complexities of nonlinear viscoelasticity, where multiple harmonics are present (Kashi et al., 2024). As a result, LAOS experiments were carried out to evaluate the oscillatory response at different applied strains (γ₀), providing a more precise method for analyzing the non-linear viscoelastic behavior through elastic or viscous Lissajous diagrams (Fig. 7). The normalized ‘elastic’ Lissajous plots of shear stress (σ) versus y revealed that the enclosed area relates to the dissipated energy (proportional to G"). At γ₀ of 6.34%, the plots displayed a narrow elliptical shape, with the straight line representing elastic stress (σe), consistent with the LVR region. UV-C treated samples exhibited narrower ellipses, indicating lower energy dissipation (Fig. 7 Left), which aligns with their higher G' values from SAOS measurements (Fig. 6). As γ₀ increased to 25.2%, these samples maintained their elliptical shape, while the plots of UV-A treated samples became distorted, indicating entry into the nonlinear viscoelastic region.
<figure><img src="image_9.png" /><figcaption>**Fig. 7.** Area surrounded by the “elastic” (Left) and “viscous” (Right) Lissajous-Bowditch plots as a function of amplitude (a frequency of 1 rad s⁻¹). Stress and strain results are normalized with a maximum stress/strain in the oscillation cycle.</figcaption></figure>
At $\gamma_o$ of 101% and 1100%, different behaviors emerged. Regarding UV-C-treated categories, stress showed structural breakdown, while UV-A irradiated groups revealed significant alteration in their elastic contributions, leading to square-like forms that indicated reduced elasticity. Notably, at the maximum strain of 1010%, all samples transformed to near-rectangular shapes, suggesting a transition from primarily elastic to nearly plastic behavior. Concerning the samples with lower irradiation duration, a sharper corner was noticed, presenting a sudden yielding, while the higher irradiation time contributed to the development of smoother transitions, indicative of more gradual yielding. In the accompanying 'viscous' Lissajous plots (Fig. 7, Right), all samples show spherical curves for $\sigma$ at lower $\gamma_o$, with areas reducing more in the case of samples irradiated with UV-C irradiation. Beyond $\gamma_o$ of 101%, these plots exhibited a sigmoid shape, indicating shear-thinning characteristics resulting from structural network disruption. Interestingly, a secondary loop emerged in the 'viscous' plots of samples with higher irradiation duration time at the highest strain, suggestive of stress overshoot related to elasticity-inertia interactions. This could stem from altered microstructural dynamics occurring more rapidly than the imposed deformation.

LAOS analysis (Tables 1 and 2) revealed that the phase angle values varied with different strains, indicating a transition from elastic to viscous characteristics. As the strain increased, there was a notable rise in the phase angle, with statistical significance ($p$ < 0.05). Although the phase angle tended to increase with higher amplitude, the extent of this increase was influenced by the duration of UV irradiation. Moreover, the oscillatory data further showed that UV-C-treated samples displayed greater network extension with superior phase angles compared to UV-A-treated samples. Notably, the UA8 or UC8 exhibited the lowest phase angle values at a 1.1% strain, suggesting a predominantly elastic behavior and enhanced resistance to deformation within the LVR.

<caption>**Table 1.** Phase angles under LAOS obtained from Lissajous-Bowditch curves for samples treated with UV-A irradiation.</caption>
<table><thead><tr><th>Frequency (Hz)</th><th>Strain (%)</th><th>UA0</th><th>UA2</th><th>UA4</th><th>UA6</th><th>UA8</th></tr></thead><tbody><tr><td rowspan="4">1</td><td>1.1</td><td>2.34 ± 0.03<sup>c</sup></td><td>1.90 ± 0.01<sup>b</sup></td><td>1.45 ± 0.02<sup>a</sup></td><td>1.42 ± 0.04<sup>a</sup></td><td>1.44 ± 0.02<sup>a</sup></td></tr><tr><td>6.1</td><td>16.1 ± 0.1<sup>a</sup></td><td>23.5 ± 0.1<sup>b</sup></td><td>26.2 ± 0.1<sup>c</sup></td><td>31.8 ± 0.03<sup>d</sup></td><td>34.1 ± 0.09<sup>e</sup></td></tr><tr><td>200</td><td>19.4 ± 0.5<sup>a</sup></td><td>20.4 ± 0.9<sup>a</sup></td><td>29.8 ± 0.4<sup>b</sup></td><td>34.8 ± 0.8<sup>c</sup></td><td>55.5 ± 0.2<sup>d</sup></td></tr><tr><td>1000</td><td>45.4 ± 1.0<sup>a</sup></td><td>68.4 ± 1.4<sup>b</sup></td><td>97.3 ± 1.6<sup>c</sup></td><td>125 ± 2<sup>d</sup></td><td>167 ± 1<sup>e</sup></td></tr></tbody></table>

a-e Means (three replicates) within each column with different letters are significantly different (P < 0.05), Duncan's test.

<table><caption>**Table 2.** Phase angles under LAOS obtained from Lissajous-Bowditch curves for samples treated with UV-C irradiation.</caption><thead><tr><th>Frequency (Hz)</th><th>Strain (%)</th><th>UC0</th><th>UC2</th><th>UC4</th><th>UC6</th><th>UC8</th></tr></thead><tbody><tr><td rowspan="4">1</td><td>1.1</td><td>2.34 ± 0.03<sup>c</sup></td><td>1.98 ± 0.03<sup>d</sup></td><td>1.01 ± 0.02<sup>c</sup></td><td>1.59 ± 0.02<sup>b</sup></td><td>1.16 ± 0.02<sup>a</sup></td></tr><tr><td>6.1</td><td>16.1 ± 0.1<sup>a</sup></td><td>22.0 ± 0.6<sup>a</sup></td><td>32.9 ± 0.7<sup>b</sup></td><td>41.6 ± 0.8<sup>c</sup></td><td>52.5 ± 0.7<sup>d</sup></td></tr><tr><td>200</td><td>19.4 ± 0.5<sup>a</sup></td><td>27.1 ± 1.6<sup>b</sup></td><td>36.9 ± 1.7<sup>c</sup></td><td>58.3 ± 1.5<sup>d</sup></td><td>60.7 ± 1.5<sup>e</sup></td></tr><tr><td>1000</td><td>45.4 ± 1.0<sup>a</sup></td><td>72.7 ± 1.6<sup>b</sup></td><td>112 ± 3<sup>c</sup></td><td>146 ± 2<sup>d</sup></td><td>189 ± 3<sup>e</sup></td></tr></tbody></table>

a-e Means (three replicates) within each column with different letters are significantly different (P < 0.05), Duncan's test.

The variations observed in phase angles for the different types of 3D structures can be attributed to variations in the formed crosslinked networks and the structural properties (Hyun et al., 2011). The UA8 or UC8 exhibited a more solid-like behavior, which can be linked to the development of a more intramolecular interaction with the polymeric systems, which agrees well with the results of SAOS and Lissajous-Bowditch curves. In this case, a strong polymeric gel-like network, likely arising from enhanced intra- and interchain interactions, may result in the development of a robust gel-like network through the development of hydrogen bonding, as confirmed by FT-IR results. However, the existence of this crosslinked network decreased the network flexibility of polymeric chains. This combination increased the likelihood of structural damage during significant deformations, leading to higher viscous dissipation (Hyun et al., 2011). The fairly dense configuration of UA8 or UC8 formed a structured system, leaving minimal space within the polymeric matrix. Consequently, when higher strain amplitudes were applied, the limited space inhibited polymeric chains from
adjusting, resulting in persistent structural changes and causing a more viscous characteristic. This is corroborated by strain sweep results, which showed that the G' (γ) patterns for UA6 and UA8 or UC6 and UC8 indicated steeper transitions (higher slopes) from linear to nonlinear properties compared to the other samples.

To quantitatively assess the type and degree of nonlinear viscoelastic characteristics, third-order harmonic data can be instrumental (Ewoldt et al., 2008). The characterization of nonlinear viscoelastic behavior is determined by assessing the ratios of the third harmonic viscoelastic moduli, G'₃ and G''₃, to the first harmonic moduli, G'₁ and G''₁. Specifically, a ratio of G'₃ / G'₁ ≥ 0.01 and G''₃ / G''₁ ≥ 0.01 indicates a nonlinear viscoelastic characteristic, with larger values reflecting an enhanced nonlinear viscoelastic property. Additionally, examining two viscoelastic moduli alongside two instantaneous viscosities can provide insights into specific nonlinear behaviors (Ewoldt et al., 2008). The ratio of the large-amplitude elastic modulus (G'ₗ) to the minimum-amplitude elastic modulus (G'ₘ) facilitates the assessment of amplitude-softening (G'ₗ / G'ₘ < 0.9) and amplitude-hardening (G'ₗ / G'ₘ > 1.1) behaviors (Tables 3 and 4). Similarly, the ratio of instantaneous viscosities at maximum (η'ₗ) and minimum shear rates (η'ₘ) can be used to evaluate shear-thinning (η'ₗ / η'ₘ < 0.9) and shear-thickening phenomena (η'ₗ / η'ₘ > 1.1) (Tables S5 and S6 in Supplementary Materials).

<table><caption>Table 3. The elastic LAOS variables for samples treated with UV-A irradiation.</caption><thead><tr><th rowspan="2">γ (%)</th><th colspan="2">UA0</th><th colspan="2">UA2</th><th colspan="2">UA4</th><th colspan="2">UA6</th><th colspan="2">UA8</th></tr><tr><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th></tr></thead><tbody><tr><td>1.1</td><td>82.3 ± 0</td><td>4.01 ± 0.</td><td>85.6 ± 0</td><td>4.43 ± 0.</td><td>85.7 ± 0</td><td>4.12 ± 0.</td><td>85.2 ± 0</td><td>3.58 ± 0.</td><td>84.8 ± 0</td><td>3.54 ± 0.</td></tr><tr><td></td><td>.2<sup>a</sup></td><td>02<sup>b</sup></td><td>.2<sup>b</sup></td><td>03<sup>c</sup></td><td>.2<sup>b</sup></td><td>02<sup>b</sup></td><td>.2<sup>b</sup></td><td>02<sup>a</sup></td><td>.3<sup>b</sup></td><td>04<sup>a</sup></td></tr><tr><td>6.1</td><td>225 ± 4<sup>e</sup></td><td>24 ± 1<sup>a</sup></td><td>207 ± 4<sup>d</sup></td><td>36 ± 6<sup>b</sup></td><td>153 ± 2<sup>c</sup></td><td>48 ± 1<sup>c</sup></td><td>124 ± 2<sup>b</sup></td><td>59.2 ± 0.</td><td>94 ± 1<sup>a</sup></td><td>75 ± 2<sup>e</sup></td></tr><tr><td>200</td><td>451 ± 6<sup>e</sup></td><td>29.1 ± 0.</td><td>325 ± 8<sup>d</sup></td><td>41 ± 1<sup>b</sup></td><td>250 ± 3<sup>c</sup></td><td>57 ± 1<sup>c</sup></td><td>147 ± 4<sup>b</sup></td><td>69 ± 1<sup>d</sup></td><td>103 ± 3<sup>a</sup></td><td>78 ± 1<sup>e</sup></td></tr><tr><td>100</td><td>777 ± 1</td><td>44 ± 2<sup>a</sup></td><td>701 ± 1</td><td>62 ± 1<sup>b</sup></td><td>634 ± 1</td><td>76.2 ± 1.</td><td>321 ± 1</td><td>82 ± 2<sup>d</sup></td><td>179 ± 9<sup>a</sup></td><td>97 ± 2<sup>e</sup></td></tr><tr><td>0</td><td>2<sup>e</sup></td><td></td><td>5<sup>d</sup></td><td></td><td>0<sup>c</sup></td><td>2<sup>c</sup></td><td>0<sup>b</sup></td><td></td><td></td><td></td></tr></tbody></table>

<footnote>a-e Means (three replicates) within each column with different letters are significantly different (P < 0.05), Duncan's test.</footnote>
<table><caption>Table 4. The elastic LAOS variables for samples treated with UV-C irradiation.</caption><thead><tr><th rowspan="2">γ (%)</th><th colspan="2">UC0</th><th colspan="2">UC2</th><th colspan="2">UC4</th><th colspan="2">UC6</th><th colspan="2">UC8</th></tr><tr><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th><th>G'<sub>L</sub> / G'<sub>M</sub> × (10)<sup>-2</sup></th><th>G'<sub>3</sub> / G'<sub>1</sub> × (10)<sup>-3</sup></th></tr></thead><tbody><tr><td>1.1</td><td>82.3 ± 0</td><td>4.01 ± 0.</td><td>95.3 ± 0.</td><td>3.46 ± 0.</td><td>95.7 ± 0</td><td>3.34 ± 0.</td><td>90.3 ± 0</td><td>3.11 ± 0.</td><td>89.8 ± 0</td><td>3.03 ± 0.</td></tr><tr><td></td><td>.2<sup>a</sup></td><td>02<sup>b</sup></td><td>2<sup>b</sup></td><td>02<sup>c</sup></td><td>.5<sup>b</sup></td><td>05<sup>b</sup></td><td>.6<sup>a</sup></td><td>02<sup>a</sup></td><td>.4<sup>a</sup></td><td>05<sup>a</sup></td></tr><tr><td>6.1</td><td>225 ± 4<sup>e</sup></td><td>24.1 ± 1.</td><td>501.3 ± 7</td><td>51.7 ± 0.</td><td>489 ± 1</td><td>63.6 ± 1.</td><td>334 ± 8<sup>b</sup></td><td>72.6 ± 1.</td><td>304 ± 6</td><td>82 ± 2<sup>e</sup></td></tr><tr><td></td><td></td><td>1<sup>a</sup></td><td>.8<sup>d</sup></td><td>4<sup>b</sup></td><td>1<sup>c</sup></td><td>4<sup>c</sup></td><td></td><td>1<sup>d</sup></td><td>7<sup>a</sup></td><td></td></tr><tr><td>200</td><td>451 ± 6<sup>e</sup></td><td>29.1 ± 0.</td><td>900 ± 16</td><td>60.3 ± 1.</td><td>826 ± 1</td><td>68.5 ± 1.</td><td>789 ± 1</td><td>77 ± 2<sup>d</sup></td><td>711 ± 1</td><td>87.7 ± 1.</td></tr><tr><td></td><td></td><td>5<sup>a</sup></td><td>1<sup>d</sup></td><td>1<sup>b</sup></td><td>1<sup>c</sup></td><td>6<sup>c</sup></td><td>4<sup>b</sup></td><td></td><td>2<sup>a</sup></td><td>4<sup>e</sup></td></tr><tr><td>100</td><td>777 ± 1</td><td>44.9 ± 1.</td><td>1201 ± 2</td><td>79.3 ± 1.</td><td>1133 ± 2</td><td>88.9 ± 1.</td><td>1154 ± 1</td><td>92.8 ± 2.</td><td>1078 ± 1</td><td>97.7 ± 2.</td></tr><tr><td>0</td><td>2<sup>e</sup></td><td>5<sup>a</sup></td><td>1<sup>d</sup></td><td>6<sup>b</sup></td><td>20<sup>c</sup></td><td>6<sup>c</sup></td><td>18<sup>b</sup></td><td>6<sup>d</sup></td><td>18<sup>a</sup></td><td>2<sup>e</sup></td></tr></tbody></table>

a-e Means (three replicates) within each column with different letters are significantly different (P < 0.05), Duncan's test.

The irradiated 3D objects generally displayed a linear viscoelastic behavior, indicated by the ratios (G'3 / G'1 < 0.01 and 0.9 < G'L / G'M < 1.1) at a frequency of 1 Hz and an amplitude of 1.1%. As the strain was increased, both the nonlinear behavior (reflected by G'3 / G'1 and G''3 / G''1) and the extent of strain hardening (G'L / G'M > 1.1) along with shear-thinning characteristics (η'L / η'M) became more pronounced with elevated amplitude and frequency (Tables S5 and S6 in Supplementary Materials). These results are further supported by the Lissajous-Bowditch curves and SAOS measurements. Analysis of the third-order harmonic data revealed that UA8 or UC8 exhibited a heightened degree of nonlinear elasticity while showing less strain hardening when subjected to the higher frequencies and amplitudes beyond the LVR. These findings were unexpected based on the oscillatory rheological data. It was initially anticipated that UA8 or UC8 would show a lower degree of nonlinear behavior due to their more structured system, which creates a somewhat robust matrix. Moreover, strain-hardening behavior in UA8 or UC8 was expected to be more significant than in the other samples due to their more organized systems and stronger network structures. This unexpected result may be attributed to increased network connectivity within alginate/gellan polymeric chains. As a result, UA8 or UC8 exhibited a solid structure with diminished flexibility, with a more compact gel structure than other 3D objects. This rigidity and packed structure likely made it easier for polymeric strands to slide over each other during LAOS, as opposed to rupturing elastically. This event resulted in less stored energy from the applied
amplitude and increased energy dissipation from both permanent deformation and viscous flow.

Consequently, while the overall degree of nonlinearity remained elevated, it manifested more as an increase in viscous behavior rather than elastic behavior, in accordance with the findings from the third-order harmonic data (Hyun et al., 2011; Melito, 2012).

**3.8. Wavelength-dependent thermo-mechanical and mechanical properties**

To study the wavelength-selective changes in material performance in more detail, dynamic mechanical analysis (DMA) and tensile tests were performed on 3D printed test specimens using the optimized printing parameters, and results are illustrated in Fig. 8. The wavelength-dependent properties of UV-A irradiation are compared to the mechanical and thermo-mechanical performance of UV-C irradiation. The untreated UV-A or UV-C exhibited a glass transition temperature (*Tg*) around -5 °C (Table S7 in Supplementary Materials). The increasing UV irradiation time slightly enhances both tensile strength and ultimate elongation. This behavior can be attributed to the thermally promoted formation of hydrogen bonding in the presence of hydroxyl and carboxylic acid moieties, which has also been observed in alginate and gellan samples (Rossegger et al., 2021). By changing the wavelength from 390 to 260 nm, a marked increase in the *Tg* was recorded (Figs. 8a,c). This increase can primarily be attributed to the effective surface crosslinking during treatment with the UV-C irradiation process. This leads to stiffer networks, exhibiting a maximum *Tg* of around 22.3 °C (Table S7 in Supplementary Materials). The DMA results further demonstrate the presence of a fully interpenetrating network, as indicated by the observation of a single-phase transition in the loss factor curve. Simultaneously, the broader glass transition region suggests an increase in network heterogeneity. The significant decrease in the Tan delta peak by one order of magnitude for the UC8 sample, when compared to the other treatments, suggests a profound alteration in the material's viscoelastic behavior. This alteration is likely attributable to enhanced intermolecular bonding resulting from the presence of the photosensitizer during UV irradiation, which promotes the formation of a more rigid, crosslinked network. The consumption
of hydrophilic groups in the polymer as a consequence of this crosslinking may contribute to the pronounced
decrease in Tan delta, corroborating findings from previous studies by Cazin et al. (2022), Cazin (2024), and
Höller et al. (2024) that demonstrate the relationship between crosslink density and reduced tan δ values in
UV-irradiated polymer systems.

<figure><img src="image_10.png" /><figcaption>**Fig. 8.** (a,c) The loss factor (tan δ) and (b,d) the storage modulus (E') as a function of temperature, derived from the DMA measurements. The test specimens were created through dual-wavelength UV crosslink 3D printing, employing a layer exposure intensity at 390 nm and 260 nm.</figcaption></figure>

The storage moduli vary significantly with the chosen wavelength, spanning over two orders of magnitude.
At 390 nm, the resulting soft networks show a storage modulus (E') of 12 MPa (at 23 °C), whereas UV-C
curing at 260 nm results in stiffer interpenetrating networks with an E' of 22 MPa (at 23 °C). This observation
confirms the distinct and complementary nature of the two photoreactions under the experimental conditions
used. The reinforcing impact of the crosslinked interpenetrating network is also evidenced by the increased
temperature at which E' drops, as well as a considerably higher rubbery plateau compared to samples cured
with longer wavelengths (Figs. 8b,d). In this study, UV-C-treated samples demonstrated optimal Young's modulus and tan δ as well as enhanced *Tg*. The FTIR analyses (Fig. 2, Top) indicate that these samples retained a minimal amount of free hydrogen bonding with increased COO⁻ band intensity. However, a comparison between these UV-C and UV-A irradiation treatments revealed an increase in thermomechanical properties by increasing irradiation duration time, accompanied by a decrease in tan δ, which can be attributed to the extended curing period.

### 3.9. Mechanical properties of 3D printed objects

The photo-crosslinking using low-energy irradiation intensity can create a supramolecular network linked by hydrogen and covalent bonds formed through γ-hydrogen abstraction. Unlike the untreated hydrogels, which rely solely on temporary hydrogen bonds, the developed crosslinked networks could be irreversible and cannot dissociate during the high shearing rate of 3D printing. To investigate the mechanical behavior further, we conducted tensile tests that were dependent on the strain rate. The samples were initially subjected to a controlled strain rate of 0.5 s⁻¹, revealing three distinct phases in the resulting curves (Fig. 9a). In the first phase (region I), at low strains (ε < 10%), the stress increases rapidly in an almost linear fashion. The second phase (region II), occurring at intermediate strains (10% < ε < 60%), shows a deceleration in the stress increase while still retaining a linear trend. Finally, in stage three (region III), when strain exceeds 60%, the rate of stress increase stabilizes. This behavior stems from the relaxation of dynamic hydrogen bonds; as the deformation rate increases, it accelerates the dissociation of reversible cross-links, necessitating greater mechanical strength, thereby enhancing both the modulus and toughness.
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Mechanical properties of 3D printed objects. (a) Uniaxial tensile curves with three distinct regimes depending on the extent of strain (ε = 0.5 s⁻¹). (b) Stress–strain curves. Cyclic mechanical properties of 3D structures, including (c) tensile loading–unloading curve at the strain rate of 0.5 s⁻¹. (d) Continuous loading–unloading curves without resting between each cycle (ε = 0.5 s⁻¹). (e) Loading–unloading cycles at different strain rates and the corresponding hysteresis. (f) Tearing force-displacement curves at different tearing rates.</figcaption></figure>

In this context, the mechanical properties of 3D printed objects should be regarded as observable parameters that significantly rely on deformation conditions such as temperature and strain rate. For this study, we focus on the deformation responses of 3D structures under isothermal conditions. We model the mechanical behavior across a range of strain rates (ε), spanning nearly three orders of magnitude from 0.0012 to 0.92 s⁻¹. Notably, all tested rates exhibit a similar pattern in the stress-strain curves, characterized by the three distinct phases influenced by the stretch ratio (Fig. 9b).

At the initial stage of deformation, Young's modulus reflects significant strain rate dependence (Fig. 9b). For strain rates below 0.12 s⁻¹, the modulus rises sharply following a power law from 23 kPa to 73 kPa (E₁ ∝ ε⁰.⁴³). However, beyond this threshold, at strain rates exceeding 0.12 s⁻¹, the modulus gradually saturates at around 85 kPa. This observation can be explained by the finite relaxation time of hydrogen bonds, which
affects their ability to break, particularly as this timescale exceeds the average bond lifetime (Chen et al., 2015). Lower strain rates allow for adequate reconfiguration of hydrogen bonds, contributing to the network modulus, as the unrelaxed reversible bonds increase with strain rate. Conversely, high strain rates result in shortened timescales where hydrogen bonds do not have sufficient time to relax, leading to a gradual leveling off of the moduli. It is noteworthy that the crossover strain rate of $0.12 s^{-1}$ aligns closely with the order of magnitude of the reciprocal hydrogen bond lifetime. After the UV-C crosslinking process, there is a notable increase in the number of cross-linked structures, a finding that is corroborated by both FTIR and NMR analyses. Specifically, consuming more hydroxyl groups and increasing the COO⁻ in the irradiated samples showed increased intra- and interchain interactions within the polymers, providing a stiffer structure for these samples.

To assess the enhancement of toughness from noncovalent junctions, cyclic loading-unloading tests were performed, calculating energy dissipation per cycle (Fig. 9c). The control sample with only hydrogen bonds showed the lowest Young's modulus of 11 kPa and toughness of $6 kJ/m^3$. The crosslinked 3D objects, featuring both covalent and reversible hydrogen bonds, achieved significant mechanical properties, boasting Young's modulus of 34 kPa and toughness of $14 kJ/m^3$, similar to some double-network hydrogels (Rubinstein & Colby, 2003). The composite hydrogels displayed a notable hysteresis loop in their loading-unloading curves, reflecting energy dissipation linked to the dissociation of hydrogen cross-links. Initially, the intact chemical cross-links and hydrogen bonds define the composite's Young's modulus. During stretching, polymer chains dissociate and reassociate, while those around chemical cross-links remain under strain. This dynamic behavior of hydrogen bonds enhances toughness under significant deformation (Rubinstein & Colby, 2003). As expected from the FT-IR and NMR analyses, the mechanical properties of the irradiated 3D objects could be improved with increasing irradiation time due to the formation of inter- or intermolecular
bonding in the irradiated samples. Therefore, controlling the reaction time is of great importance to optimize
the mechanical properties of the 3D printed objects during 3D printing.

Upon unloading, stress stored in the covalent network aids in recovering the initial state. In contrast, the
untreated sample showed nearly overlapping loading and unloading curves, indicating rate independence,
even when the stretch rate increased tenfold (Fig. S6 in Supplementary Materials). Repeated loading-
unloading cycles reduced hysteresis and maximum stress, consistent with prior findings on physically cross-
linked hydrogels (Chen et al., 2015). To investigate strain rate-dependent viscoelasticity, we conducted
sequential stress relaxation tests by stretching samples to various strains (100 to 800%) at a constant rate
of 0.2 s⁻¹, with a 2-minute hold (Fig. S6 in Supplementary Materials). This revealed asymmetric loading
curves with little rate sensitivity in unloading. The dissociation of interfacial bonds during stretching led to
strain rate dependence, while unloading showed minimal changes since no additional bond dissociation
occurred.

**3.10. Tearing test**

To study the fracture behavior of 3D printed objects, a pure shear test was conducted (Fig. 9f). Initially, the tearing force reached a peak, indicating maximum resistance to fracture propagation. Compared to the UV-A treated samples, UV-AC irradiated ones exhibited a higher tearing force, achieving a maximum tear strength of 0.13 N mm⁻¹ for a 10 mm wide specimen, which is considerably higher than the UV-A categories (0.031 N mm⁻¹). This suggests that UV-C irradiation treatment can efficiently dissipate energy within the 3D printed structures. Consequently, the tearing energy in the crosslinked 3D printed structures is velocity-dependent and increases with higher tearing speeds (Fig. 9b). At high tearing rates (>0.15 s⁻¹), this increase plateaus as hydrogen bonds become unrelaxed and the modulus of the polymer chains stabilizes.
Conversely, samples treated with UV-A irradiation showed a slight increase in energy with tearing velocity, attributed to weak viscoelastic deformation from chain entanglements or loops.

As demonstrated in Fig. 9f, when a notched sample is stretched, crack propagation initiates from $ε_c$, with the tearing force decreasing over time until it approaches zero. The crack propagation velocity ($V_p$) can be estimated using the formula $V_p = (w-c)/t$, where w is the sample width and c is the initial notch width. Varying the tearing velocity over three orders of magnitude (14 × 10⁻³ to 92 × 10⁻¹ s⁻¹) reveals a significant dependence of crack propagation on tearing speed, characterized by three distinct regimes (Fig. 9f): at the tearing velocities (<5 × 10⁻² s⁻¹), $V_p$ remains constant; at moderate velocities (5 × 10⁻² – 2 × 10⁻¹ s⁻¹), $V_p$ increases swiftly following a power law ($V_p \approx \varepsilon^{0.48}$); beyond this range, $V_p$ stabilizes around 2.2 mm s⁻¹ for tearing velocities above 0.12 s⁻¹. This behavior underscores the crucial role of dynamic hydrogen bonds in sample resistance over varying timescales (Mayumi et al., 2016). At low tearing rates, the fast dissociation-reassociation of reversible bonds prevents significant stretching, requiring less energy for crack propagation. However, at higher tearing rates, where the tearing experiment time is shorter than the average lifetime of reversible bonds, unrelaxed bonds support network strength, increasing the energy needed for crack advancement. At even greater tearing rates, this increase plateaus as most hydrogen bonds remain unrelaxed, resulting in constant network rigidity. Additionally, since hydrogen bonds in 3D printed objects can be disrupted in urea-aqueous solutions, affecting their toughness and strength, we investigated the impact of urea concentration (0.02–0.12 M) during tearing tests. Results indicate that crack propagation resistance and tearing energy decrease significantly in urea solutions at specific tearing velocities, suggesting a substantial loss of hydrogen bonds, which agrees well with the FT-IR results. A prolonged irradiation duration was positively correlated with elevated tensile loading–unloading values, indicating the formation of a more resilient gel-like polymer network. This increase suggests the enhancement of both intra- and interchain interactions within the 3D-printed crosslinked samples, which is consistent with the FT-IR spectroscopic
results. The presence of hydroxyl groups in alginate and gellan, coupled with the symmetric stretching vibrations of the COO⁻ functional group, significantly influences the photochemical reactions initiated by UV irradiation, facilitating the creation of crosslinked architectures. This mechanism is corroborated by the characteristic absorption bands observed at 3395 cm⁻¹, indicative of O–H stretching vibrations, alongside bands at 1600 cm⁻¹ (related to COO⁻ asymmetric stretching) and 1439 cm⁻¹ (associated with aromatic C=C stretching). These findings, illustrated in the FT-IR spectrum (Fig. 2, Top), underscore the impact of photocrosslinking on the material's properties. Thus, the observed enhancement in tearing behavior can be primarily ascribed to the establishment of crosslinking through hydrogen bonding and the resultant increase in molecular network density. This advancement in the polymer structure likely contributes to improved mechanical stability and viscoelastic performance. Overall, these findings demonstrate the importance of interfacial hydrogen bonds in achieving the toughness of composite hydrogels.

### 3.11. Microstructure of 3D-printed structures

Fig. 10 presents a variety of printed structures progressing from 1D to 3D, created using different UV wavelengths. The 1D filaments made with UV-A were found higher diameters (measured ranging from 788 to 887 µm) than those of UV-C (provided ranging from 686 to 740 µm) (Fig. 10). The cross-sectional rupture surface of UV-C 3D structures was a smooth and circular structure with no noticeable aggregates on their surfaces (Fig. 10), which can serve as a solid basis for the development of more complex designs. In contrast, filaments created with UV-A showed rough surfaces and irregular cross-sectional ruptures, revealing notable structural weaknesses. The more structured matrix observed in the UV-C-treated samples likely resulted in the production of regular 3D filaments, which in turn led to more homogenous self-supporting filaments than those of the UV-A samples. Additionally, the square lattice (grid) featuring multiple layers, produced with UV-C samples, showcased excellent printing capabilities. On the other hand, the 3D square lattice printed with UV-A samples exhibited shape distortions with rounded angles, which signal dimensional instability. Similarly,
the corresponding 3D constructs produced from UV-A samples featured an uneven, non-uniform
arrangement and inconsistent spacing, reflecting lower printing quality and compromised dimensional
stability. This resulted in a lower printability along the rounded edges when compared to the 3D structures
made with UV-C irradiation. The emergence of more gel-like viscoelastic properties (Figs. 6 and 7) in UV-C
samples compared to UV-A samples may account for the enhanced geometry and shape accuracy observed
in the printed models. The improved elastic modulus and increased network stiffness driven by 260 nm-UV
wavelengths played a role in achieving superior spatial resolution and overall printing efficacy.

Quantitative assessment of printing fidelity revealed that UV irradiation significantly enhanced dimensional accuracy (p < 0.05). For 1D filaments designed at 800 µm, untreated hydrogels (UA0/UC0) exhibited a diameter of 998 ± 12 µm (24.8% oversizing), while UV-treated samples progressively approached the target dimension (Table 5). Notably, UC8 filaments achieved 686 ± 8 µm (-14.3% deviation). The printability index, calculated from 0-90° log-pile structures (Fig. 10) as the ratio of observed to theoretical intersection area, improved from 1.45 ± 0.25 (UA0/UC0) to 0.97 ± 0.02 for UC8, indicating near-ideal grid formation. In cylindrical constructs (5 mm design height) (Fig. 1d), layer consistency was markedly enhanced with UV treatment. The number of resolved layers increased by 170% (from 10 ± 1 to 27 ± 1 layers in UC8), corresponding to a reduction in average layer height from 500 µm to 185 µm. Concurrently, layer width variability decreased substantially, with standard deviation reducing from ±0.3 mm (UA0/UC0) to ±0.1 mm (UC8). These metrics collectively demonstrate that UV irradiation refines printing resolution to levels suitable for precision applications, with UC8 samples showing unprecedented dimensional accuracy for the alginate-gellan system (31% reduction in filament deviation and near-ideal printability index).

The printed pattern precision of an axial pore in an XY plane to exhibit a perfect square geometry was related
to *Pr* (0 ≤ *Pr* ≤ 1). In a particular shape, a square geometry (high precision) can be measured if *Pr* = 1, an
irregular shape if *Pr* > 1, and a round shape if *Pr* < 1. The *Pr* was detected to be about 1.20 for UA0 or UC0
(not provided), which denoted a poor printing performance. Compared to UV-A, UV-C provided a square geometry with higher precision. For example, the *Pr* value of UC4 and UC8 was detected to be 0.93 and 0.98, while the *Pr* value of UA4 and UA8 was measured at 1.12 and 1.07 (Fig. 10). In addition, a higher duration of UV-C irradiation led to the formation of somewhat rounded pattern geometry, as the *Pr* values were close to 1. It is concluded that the shape fidelity of UV-C samples was upheld even if more layers were added to the structures. While 3D printing performance validates network tunability, this work fundamentally advances carbohydrate polymer design by establishing UV-radical grafting as a tool to precisely control chain entanglement thresholds and supramolecular assembly. These principles enable spatiotemporally resolved hydrogel design for applications beyond printing—including controlled drug release and tissue scaffolds—fulfilling carbohydrate polymers' aim to publish innovations in carbohydrate material chemistry.

<table><caption>**Table 5.** Quantitative printing fidelity analysis of UV-treated hydrogels.</caption><thead><tr><th>Samples</th><th>Filament diameter (µm) (1D filament)</th><th>Deviation from CAD (%)</th><th>Printability index (0 – 90° log-pile structure)</th><th>Layer width (mm) (Cylinder)</th><th>Layer number (Cylinder)</th></tr></thead><tbody><tr><td>UA0/UC0</td><td>998 ± 12<sup>g</sup></td><td>+24.8</td><td>1.45 ± 0.25<sup>f</sup></td><td>2.4 ± 0.3<sup>e</sup></td><td>10 ± 1<sup>a</sup></td></tr><tr><td>UA2</td><td>887 ± 6<sup>f</sup></td><td>+10.9</td><td>1.18 ± 0.04<sup>e</sup></td><td>2.0 ± 0.3<sup>d</sup></td><td>14 ± 1<sup>b</sup></td></tr><tr><td>UA4</td><td>874 ± 5<sup>e</sup></td><td>+9.3</td><td>1.06 ± 0.03<sup>c</sup></td><td>1.8 ± 0.1<sup>c</sup></td><td>17 ± 1<sup>c</sup></td></tr><tr><td>UA6</td><td>793 ± 9<sup>d</sup></td><td>-0.9</td><td>1.09 ± 0.04<sup>cd</sup></td><td>1.4 ± 0.2<sup>b</sup></td><td>20 ± 2<sup>d</sup></td></tr><tr><td>UA8</td><td>788 ± 12<sup>d</sup></td><td>-1.5</td><td>1.12 ± 0.03<sup>d</sup></td><td>1.4 ± 0.1<sup>b</sup></td><td>24 ± 1<sup>e</sup></td></tr><tr><td>UC2</td><td>740 ± 7<sup>c</sup></td><td>-7.5</td><td>1.11 ± 0.02<sup>d</sup></td><td>2.1 ± 0.2<sup>d</sup></td><td>13 ± 1<sup>b</sup></td></tr><tr><td>UC4</td><td>724 ± 6<sup>b</sup></td><td>-9.5</td><td>0.93 ± 0.02<sup>a</sup></td><td>1.5 ± 0.1<sup>b</sup></td><td>20 ± 2<sup>d</sup></td></tr><tr><td>UC6</td><td>693 ± 9<sup>a</sup></td><td>-13.4</td><td>0.93 ± 0.03<sup>a</sup></td><td>1.4 ± 0.2<sup>b</sup></td><td>21 ± 1<sup>d</sup></td></tr><tr><td>UC8</td><td>686 ± 8<sup>a</sup></td><td>-14.3</td><td>0.97 ± 0.02<sup>b</sup></td><td>1.1 ± 0.1<sup>a</sup></td><td>27 ± 1<sup>f</sup></td></tr></tbody></table>

<footnote><sup>a-g</sup> Superscript letters indicate statistical significance (p < 0.05, Duncan's test). Deviations calculated relative to 800 µm</footnote>
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** SEM images of 3D-printed grids as well as their relevant photomicrographs of the surface or cross-sectional rupture of different 3D-printed filaments. From left to right: FE-SEM images of the surface of 1D printed filaments (first column), their relevant cross-section raptures (second column), 0 – 90° log-pile structure (third column-top), and printing quality images of 3D printed objects (third column-bottom).</figcaption></figure>

Fig. 11 presents SEM images of the porous scaffolds of UV-A and UV-C irradiation created by freeze-drying techniques. The images reveal that the porous scaffolds produced by UV-A exhibit a connected internal structure with pore sizes measured in a few microns. Notably, the average pore diameter of this 3D scaffold is notably smaller than the average pore diameter in our previous study (Asghartabar Kashi et al., 2024). This reduction in pore size may be attributed to volume shrinkage resulting from the strong capillary forces experienced during the drying process. Additionally, a distinct difference was observed in the inner pore structure of scaffolds made with UV-A irradiation compared to UV-C scaffolds. The UV-C scaffolds featured a layer lining the inner pore walls and the wall thickness. This increase contributed to enhanced thickness and structural integrity of the scaffold's pore walls. In summary, the strategies outlined in this study show
significant potential for enhancing the development of 4D-printed porous structures, particularly in food and pharmaceutical applications. This includes the innovative controlled release of bioactive compounds.

The comprehensive evaluation confirms that UV irradiation induces radical-mediated covalent crosslinking within a single polymer network, not an interpenetrating polymer network (IPN). The mechanical synergy arises from the complementary roles of excited sodium benzoate by the generation of free radicals without directly participating in the hydrogen abstraction from alginate and gellan gum within a unified covalently crosslinked matrix. This paradigm shift resolves critical conceptual ambiguity while explaining our observed printability transformation: covalent bonds stabilize the transient physical networks, enabling precise extrusion control unattainable in pure crosslinked systems. This aligns with fundamental polymer physics principles—covalent networks resist elastic recoil during printing, reducing die swell relative to non-irradiated controls.

<figure><img src="image_13.png" /><figcaption>Fig. 11. FE-SEM images of 3D printed freeze-dried UV8 and UC8 structures.</figcaption></figure>

**4. Conclusions:**

This study establishes a novel dual-wavelength UV crosslinking strategy using the combined effect of UV-A (390 nm) and UV-C (260 nm) with sodium benzoate as photosensitizer to fabricate alginate/gellan hydrogels with tunable mechanical properties and precise microstructural fidelity. The results revealed that UV-C
irradiation enhances hydrogel stiffness and elasticity through chemoselective free radical crosslinking, as evidenced by increased hydrogen bonding and interchain interactions. In contrast, UV-A irradiation produces softer, ductile networks with greater flexibility but reduced structural homogeneity. The strain-rate sensitivity of these hydrogels reveals critical insights into their mechanical behavior. Below 0.12 s⁻¹, hydrogen bonds dynamically reconfigure, enabling energy dissipation and a strain-rate-dependent Young's modulus. Beyond this threshold, bond relaxation timescales limit reorganization, causing modulus saturation and rigid behavior. UV-C irradiation (260 nm) enhances stiffness and precision (filament diameter: 686–740 µm; Pr ≈ 0.98) via dense hydrogen bonding, while UV-A (390 nm) prioritizes ductility at the cost of structural uniformity (Pr ≈ 1.12). The stiffness-ductility trade-off highlights UV-C's suitability for load-bearing applications (e.g., biomedical scaffolds) and UV-A's compatibility with flexible matrices (e.g., edible structures). Scalability challenges and UV-C's inherent brittleness under extreme strain warrant future optimization of crosslinking kinetics. Strain-rate tailoring offers a pathway to design hydrogels for dynamic environments, balancing toughness and rigidity through controlled bond dynamics. Scalability remains constrained by the inverse relationship between print speed and resolution. Further optimization of crosslinking kinetics could mitigate UV-C's brittleness while preserving its thermo-mechanical advantages. Applications in drug delivery and personalized nutrition should leverage UV-C's accelerated release kinetics and UV-A's biocompatible ductility.

## Funding

The research funding was provided by the BOKU University (Vienna, Austria).

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Graphical abstract

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