The article begins with a general introduction to various energy storage systems and the need for dielectric capacitors as energy storage devices. This is followed by a brief discussion on the mechanism of energy storage in capacitors, ferroelectrics, anti-ferroelectrics, and relaxor ferroelectrics as potential candidates for energy storage. The remainder of this article is devoted to reviewing the energy storage performance of bulk ceramics, multilayers, and thin films of BiFeO3-based relaxor ferroelectrics, along with a discussion of strategies to address some of the issues associated with their application as energy storage systems.Dielectric capacitors have been widely studied because their electrostatic storage capacity is enormous, and they can deliver the stored energy in a very short time. Relaxor ferroelectrics-based dielectric capacitors have gained tremendous importance for the efficient storage of electrical energy. Relaxor ferroelectrics possess low dielectric loss, low remanent polarization, high saturation polarization, and high breakdown strength, which are the main parameters for energy storage. This article focuses on a timely review of the energy storage performance of BiFeO In addition, the depletion of fossil fuel resources, the dominant candidate in the energy market, poses the risk of an energy crisis in a world where the number of consumers is increasing day by day. Considering the serious threats to the lives on Earth and the risk of an energy crisis posed by the use of fossil fuel resources, the transition to clean energy is a serious consideration. However, renewable energy sources such as solar, wind, tides, and geothermal energy are intermittent by nature. Therefore, harnessing and storing renewable energy for future access is a challenging task.Global warming poses potential threats to the planet Earth’s future. The continuous burning of fossil fuels has increased the concentration of COElectrical energy harvested from renewable sources offers enormous opportunities for meeting future energy demands and the feasibility of the transition to clean energy. However, the usefulness of the electrical energy generated depends on its efficient storage, which is necessary for around-the-clock use. An efficient electrical energy storage (EES) system is the heart of the commercial and residential grid-based utilization of electrical energy. Therefore, the development of advanced EES systems is critically important for meeting the growing energy demands and effectively leveling the cyclic nature of such energy sources . For oveA perfect energy storage device is characterized by high energy and power densities. A comparison of the storage efficiency of the technologically relevant candidates for EES systems can be realized from the Ragone plot shown in Dielectric capacitors are one of the key components in modern communication technology, with applications in electronic circuits, warfare, distributed power systems, hybrid electric vehicles, clean energy storage, high-power applications, etc. ,9,10,11.3. The two important figures of a capacitor that determine its energy storage performance are the recoverable energy density (Urec) and energy efficiency (η), which depend on the saturation polarization (Pmax), remnant polarization (Pr), and breakdown strength (BDS) of the materials. Linear dielectric (LD), ferroelectric (FE), and anti-ferroelectric (AFE) materials are widely used for the fabrication of ceramic capacitors. Although the LDs possess excellent values of BDS, their Urec values are quite low due to weak polarization. FE materials, on the other hand, possess quite a large Pmax and Pr. As can be seen in Pmax and Pr result in a low Urec value. Moreover, the ferroelectric materials suffer large hysteresis loss, which has a significant detrimental effect on the energy efficiency of the capacitors. AFE materials behave like LDs in the low field regime, undergoing a field-induced FE state, yielding a high Pmax at a high electric field with large hysteresis loss. Therefore, the issues associated with the AFE for their applications in the high-energy storage application are essentially similar to that of LDs at low electric fields and to that of FEs at high electric fields. However, the problem of hysteresis loss inherent to FE and AFE materials is found to be minimized in RFEs while maintaining a significantly large value of Pmax. This motivates the scientific community to turn their heads towards the RFEs in search of high-energy storage capacitors.The current review discusses the recent progress on the development of high-energy storage dielectric capacitors based on the relaxor ferroelectric (RFE) of BiFeO3 have recently gained interest in the field of energy storage applications because of the high polarization (~90 μC cm−2) predicted in BiFeO3, along with its high ferroelectric Curie temperature (TC) (~830 °C)  direction . Various experiments have confirmed the ferroelectricity in BiFeO3 below TC = 1143 K [3 are displaced to a large extent in comparison with other FE compounds with non-lone-pair-active cations. Therefore, a large value of spontaneous polarization, on the order of 90 μC cm−2, has been predicted in BiFeO3 from ab initio calculations [3, as predicted by theory [−2) in a high-quality single-crystal BiFeO3 at room temperature, as shown in The = 1143 K ,59. The = 1143 K . The lonulations ,60. Howeulations , after ay theory ,65,66,67y theory  measured3 possesses the highest values of spontaneous polarization and TC among Pb-free FEs. As discussed in Pmax and large Pmax –Pr are among the most important factors for obtaining a high storage capacity. Therefore, BiFeO3 with a large Pr in its naturally occurring FE phase is not suitable for energy storage applications, which is a drawback of all FE materials in general. One way to obtain a small Pr is to break the long-range FE order such that it becomes an RFE. Many researchers have reported this method to create PNRs embedded in a non-FE matrix and obtain a significant reduction in the Pr value. In addition, there are always issues pertaining to the leakage current in pure BiFeO3, which eventually limits its high breakdown strength. Various successful methods to enhance the resistivity in phase-pure BiFeO3, such as doping at the A-site and the addition of Mn, have been discussed in the literature. In the following sections, considering these issues, the energy storage performance of BiFeO3-based materials, with special emphasis on the RFEs, are reviewed.BiFeO3−BaTiO3 solid solutions have been widely investigated as promising candidates in the field of ceramic dielectrics-based energy storage materials. A remarkable feature of BiFeO3-based solid solutions is the morphotropic phase boundary (MPB), where the solid solution displays a composition-driven structural transition in its phase diagram. The crystal structure changes abruptly across the MPB, and various physical properties, such as piezoelectric coefficients and polarization, are maximal at the MPB. The MPB of the BiFeO3−BaTiO3 system is shown in the phase diagram in 3 content produces excellent ferroelectric and piezoelectric properties in MPB compositions.BiFeO3-BaTiO3 also possesses a high Pr [2O3, MnO2, and Nb2O5 in BiFeO3 can significantly enhance the electrical resistivity and energy loss density [2O5-modified BiFeO3-BaTiO3, i.e., (1-x)(0.65BiFeO3-0.35BaTiO3)–xNb2O5  by several orders (~1010–1014 Ω cm) compared with the undoped BiFeO3-BaTiO3 (~108 Ω cm). The compounds with x = 0.01 and 0.03 exhibited slimmer P−E loops, similar to RFEs, with x = 0.03 exhibiting the highest Pmax (25.21 μC cm−2) and lowest Pr (5.53 μC cm−2). They obtained a maximum Urec of 0.71 J cm−3 at E = 90 kV cm-1. Zhu et al. [3-0.48BaTiO3 ceramic up to 130–140 kV cm−1 by adding La2O3 and MnO2. Under this condition, they obtained Urec = 1.22 J cm−3 with η = 58% for 0.52Bi0.98La0.02FeO3-0.48BaTiO3 + 0.3 wt.% MnO2 compound, whereas undoped 0.52BiFeO3-0.48BaTiO3 exhibited Urec = 1.08 J cm−3 with η = 49%. They found that the addition of La2O3 and MnO2 increased the amount of the FE phase, reduced the grain size, and facilitated densification, which helped to induce large ∆P (Pmax − Pr) as well as BDS compared with undoped compounds. This phenomenon was found to be more pronounced when Nd was substituted for Bi sites in MnO2-added BiFeO3-BaTiO3 solid solution, as shown in ρr = 95% to 97.6%) 0.75BiFeO3-0.25BaTiO3 ceramics by Nd substitution with 0.1 wt.% MnO2 addition that could endure a high electric field up to 180 kV cm−1, as shown in Ust = 4.1 J/cm3 and Urec = 1.82 J cm−3; however, it had a low value of η = 41.3%. Recently, Chen et al. [3-BaTiO3 binary ceramics that can endure a very high electric field of up to 200 kV cm−1. Sm substitution significantly reduced the grain size and enhanced the density, which is believed to be the reason for the high BDS. The binary solid solutions exhibit excellent Ust and Urec; however, their low efficiency limits practical device applications.However, BiFeO high Pr  and high density ,73, whic density  reportedu et al.  found a n et al.  successf3-MTiO3 (M = Ba and/or Sr) as the end member of the ternary system has been found to be very promising for inducing the relaxor phase and enhancing the energy storage performance [3-BaTiO3 as a result of the substitution of BaMg1/3Nb2/3O3. The relaxor 0.61BiFeO3-0.33BaTiO3-0.06BaMg1/3Nb2/3O3 exhibited a Urec of 1.56 J cm−3 at E = 125 kV cm−1, with η ~ 75%. This compound also exhibited good temperature stability for energy storage and efficiency in the temperature range of 25 °C to 190 °C. In a similar effort, Zheng et al. [Urec of 1.66 J cm−3 at 130 kV cm−1 and η ~ 82% in highly dense (ρr > 97%) 0.61BiFeO3-0.33BaTiO3-0.06LaMg1/2Ti1/2O3 ceramics. Meanwhile, Liu et al. [3-0.34BaTiO3-xBaZn1/3Ta2/3O3 for x > 0. They reported slim P–E loops for x > 0 with the highest BDS of E = 160 kV cm−1 and a high Urec of 2.56 J cm−3 for the x = 0.06 composition (−1 and Urec = 1.62 J cm−3 for 0.85(0.65BiFeO3-0.35BaTiO3)-0.15Ba(Zn1/3Nb2/3)O3 bulk ceramics. Sun et al. [Urec = 2.11 J cm−3 at E = 195 kV cm−1 with η = 84% in a highly dense 0.56BiFeO3-0.30BaTiO3-0.14AgNbO3+5 mol% CuO system prepared by a modified thermal quenching technique.The addition of a third perovskite oxide to binary BiFeOformance ,82,83,84formance  reportedg et al.  reportedu et al.  reportedposition . Tang et3-BaTiO3-Bi(Mg2/3Nb1/3)O3 system. They prepared coarse-grained (grain size ~2 to 4 μm) and fine-grained (~0.55 to 0.9 μm) microstructures using planetary ball milling and high-energy ball milling processes, respectively. BiFeO3-BaTiO3-Bi(Mg2/3Nb1/3)O3 solid solutions with a fine-grained microstructure exhibited higher ∆P (~ 30 μC cm−2) and BDS (~ 110 kV cm−1) than the coarse-grained samples . Under such drastic microstructural evolution, they reported a higher Urec of 1.26 J cm−3 in fine-grained samples, compared with Urec = 0.16 J cm−3 for coarse-grained samples. Yang et al. [2O5) (BCB) as the low melting point additive in 0.1 wt% MnO2-added (0.67−x)BiFeO3–0.33(Ba0.8Sr0.2)TiO3-xLa(Mg2/3Nb1/3)O3 solid solution, they achieved a BDS of 230 kV cm−1 for x = 0.06. With the use of such a high field, the compound exhibited Urec = 3.38 J cm−3 with η = 59%. BCB formed large amounts of liquid phase at the grain boundaries during sintering and significantly reduced the average grain size down to the submicron range, as shown in 2 helped to decrease Fe3+ ↔ Fe2+ valence fluctuations by minimizing Bi2O3 loss during synthesis, which was also critical for enhancing the BDS. This compound exhibited good temperature stability, with Urec = 1.15–1.27 J cm−3 in the temperature range of 30 °C to 170 °C. Using the PVA-assisted viscous polymer process (VPP) route, Liu et al. [ρr ~ 99% 15 mol% Sr0.7Bi0.3FeO3–modified 0.85(0.65BiFeO3-0.35BaTiO3) system with a fine grain microstructure. The ultra-high ρr and uniform submicron grains significantly enhanced the BDS, with a value of 330 kV cm−1, compared with that of the sample prepared without VPP (180 kV cm−1). Under this condition, they obtain an ultra-high Urec of 4.95 J cm−3 with η ~ 73%. The calculated Urec from the charge-discharge cycling test was 2.36 J cm−3 at 300 kV cm−1. Likewise, Sm doping has been found to be very effective in reducing the grain size and increasing the density of sintered ceramics. Chen et al. [1-xSmxFeO3-0.33BaTiO3 system with an increase in x, resulting in a BDS of 200 kV cm−1 at x = 0.1. This resulted in an Urec = 2.8 J cm−3 for x = 0.1; however, it had low efficiency, η = 55.8%. In another report, Li et al. [Urec = 3.2 J cm−3 at E = 206 kV cm−1 with high efficiency, η = 92%, in a highly dense (ρr ~ 98%) solid solution of (1−x)Bi0.83Sm0.17Fe0.95Sc0.05O3-x(0.85BaTiO3-0.15Bi(Mg0.5Zr0.5)O3) with x = 0.75. A similar Urec ≈ 3.06 J cm−3 at E = 167 kV cm−1 with η ≈ 92% associated with the improvement of BDS has been reported in (1−x)BiFeO3-x(0.85BaTiO3-0.15Bi (Sn0.5Zn0.5)O3) with x = 0.65 [Yu et al.  studied g et al.  showed tu et al.  obtainedn et al.  reportedi et al.  reportedx = 0.65 .3-rich shells and BiFeO3-rich cores in (0.7-x)BiFeO3-0.30BaTiO3-xBi(Zn2/3Nb1/3)O3+0.1wt%Mn2O3 ceramics that could withstand an electric field as high as 190 kV cm−1. They found that the shells and cores in the solid solutions were different in structure and electrical characteristics. The shell exhibited a pseudo-cubic structure, which is paraelectric in nature, whereas the core parts had a ferroelectric R3c structure. The core-shell structure and the cationic charge disorder at the B-sites helped to establish the relaxor phase in 0.70BiFeO3-0.30BaTiO3 (2/3Nb1/3)O3 further exacerbated the long-range order, thereby inducing a highly disordered RFE phase while maintaining the polarizability as high as Pmax = 36.7 μC cm−2 and ∆P = 32.8 μC cm−2. They obtained Ust = 3.7 J cm−3, Urec = 2.06 J cm−3 at 180 kV cm−1 and Ust = 2.9 J cm−3, Urec = 1.98 J cm−3 at 190 kV cm−1 for 5 mol% and 8 mol% Bi(Zn2/3Nb1/3)O3-doped compounds, respectively. The solid solutions could also successfully deliver discharge energy within 0.5 μs. Wang et al. [3-0.3BaTiO3-xNd(Zr0.5Zn0.5)O3 for x = 0.05, 0.08, and 0.10, with a core-shell structure, and studied the energy storage performance at room temperature both in ceramics and multilayers prepared by a solid-state reaction route. The substitution of Nd(Zr0.5Zn0.5)O3 induced chemical inhomogeneity, revealed as a -rich core for x = 0.05 and -rich cores for x = 0.1, and complex multiphase microstructures with both -rich and -rich cores for x = 0.08. A high degree of chemical inhomogeneity and better electrical homogeneity of grains existed at x = 0.08, which led to a high ∆P and provided a difficult current path for electrical breakdown. As a result, x = 0.08 exhibited Urec ~2.45 J cm−3 at E = 240 kV cm−1 with η = 72%. Lu et al. [3, SrTiO3, Nb2O5, and BiMg2/3Nb1/3O3 exhibiting a core-shell structure. Their study successfully demonstrated that Nb2O5 doping into BiFeO3-SrTiO3 and employing the third perovskite end member BiMg2/3Nb1/3O3 (BMN) could induce an insulating relaxor phase at room temperature without reducing the average ionic polarizability of the solid solution. It was found that the substitution of 1–3% Nb5+ for Ti4+ (B sites) in 0.6BiFeO3-0.4SrTiO3 suppressed the formation of oxygen vacancies and significantly reduced p-type conductivity compared with that of the undoped compound. A similar reduction in the p-type conductivity was observed for the 0.56BiFeO3-0.4SrTiO3-0.04BiMg2/3Nb1/3O3-xNb2O5 (x = 0–0.05) solid solution, and an enhanced BDS of 360 kV cm-1 was obtained at an optimized Nb2O5 content (x = 0.03). Such a high BDS was attributed to two factors: (i) the improved insulating character caused by Nb doping at B-sites related to the suppression of the formation of oxygen vacancies, and (ii) a core-shell microstructure with electrical homogeneity throughout the grains. The x = 0.03 composition exhibited a high performance of Urec = 6 J cm−3 with η = 74.6%. Then, a much-improved Urec value with similar η could be achieved by optimizing the BiMg2/3Nb1/3O3 content in (0.6-y)BiFeO3-0.4SrTiO3-0.03Nb2O5-yBiMg2/3Nb1/3O3 (y = 0.02–0.12) solid solutions. For y = 0.1, the BDS was enhanced further up to 460 kV cm−1, producing Urec = 8.2 J cm−3 and η = 74.6%. A core-shell microstructure design provides a large BDS and large Urec, but at the same time, it has low efficiency because of the not-very-fast response of the PNRs to the electric field.Designing a specific microstructure has been found to be highly effective for enhancing energy storage performance. Microstructure strongly influences the BDS of dielectric materials and their relaxor characteristics. Wang et al.  designed30BaTiO3 a. Substig et al.  reportedu et al.  demonstrUrec and η. Using it, they obtained large Pmax, large ∆P, and superior energy storage performance in 0.57BiFeO3-0.33BaTiO3-0.1NaNbO3 with the addition of 0.1 wt% MnO2 and 2 wt% BaCu(B2O5). Nanodomain engineering produced stripe-like PNRs embedded in featureless nanodomains. The stripe-like PNRs were rich in BiFeO3, whereas the matrix domains were rich in BaTiO3 and NaNbO3. An HR-TEM image of the solid solution is shown in 3 were dispersed in BaTiO3- and NaNbO3-rich featureless nanodomains. This structural heterogeneity at the domain level led to a rapid polarization response to the external E-field, unlike in the core-shell microstructure, and produced a hysteresis-free P−E loop, as well as a large Pmax. Moreover, NaNbO3 substitution increased the bandgap and helped to obtain a uniform and fine-grained microstructure, which was beneficial to enhance BDS up to 360 kV cm−1, producing Urec ≈ 8.12 J cm−3 with η ≈ 90%. This solid solution also exhibited excellent thermal stability, with Urec = 8.12 J cm−3 ± 10% in the temperature range −25 °C to 250 °C, and an ultrafast discharge rate (< 100 ns).Qi et al.  employedIn the past few decades, thin films have gained enormous scientific importance in a plethora of applications. The obvious advantages thin films possess over their bulk counterparts are that they save material and have reduced weight. In addition, it is well known that material properties drastically change when they are deposited in thin film forms, which paves the way for industrial applications. Dielectric thin films with thicknesses on the nano- or submicron scale have shown promising potential in the field of energy storage for low-power small devices. This is because of their extraordinarily high BDS (1 MV/cm) and energy density compared with bulk dielectrics.3-0.6SrTiO3 onto a SrRuO3-electroded (100)-SrTiO3 substrate using the pulsed laser deposition (PLD) method, which can endure an electric field as large as 0.972 MV cm−1. Under such a high electric field, the film exhibited Urec = 18.6 J cm−3 with η > 85%. It also featured a small temperature coefficient of capacitance (TCC < 11%) over a wide range of temperatures up to 200 °C. Pan et al. [−1, when doped with Mn and deposited onto a Nb-doped SrTiO3 (001) substrate by the PLD technique. The high BDS of the Mn-doped 0.4BiFeO3-0.6SrTiO3 thin film was attributed to (i) a dense, uniform, and crack-free microstructure, (ii) high-quality epitaxial growth, and (iii) low leakage current by reducing the Fe3+/Fe2+ valence fluctuation with Mn substitution. They obtained a colossal Urec = 51 J cm−3 and η = 64%, which are comparable to those of lead-based thin films (Urec = 61 J cm−3 and η = 33%) [7 cycles and good thermal stability over a wide range of temperatures, from −40 °C to 140 °C. In another report, Pan et al. [3-xSrTiO3 thin films (x = 0.3–0.75) deposited by the PLD technique and found that the BDS of the films could be enhanced further, up to 4.46 MV cm−1, by increasing the SrTiO3 content. Atomic-scale microstructure analysis based on the STEM of BiFeO3-SrTiO3 films revealed that SrTiO3 disrupted the long-range FE order, and this disruption cascaded with an increase in SrTiO3 content. The incorporation of SrTiO3 could transform the micrometer-scale FE domains into nanoscale PNRs. Paraelectric SrTiO3 acts as a matrix for embedded PNRs and separates the PNRs such that the inter-PNR interaction almost vanishes. Because there are no inter-PNR interactions, the PNRs are very dynamic under an external electric field and produce slim P–E loops with very small Pr values. The size of the PNRs continued to decrease with increasing SrTiO3 content, forming an almost domain-less feature for x = 0.75. This domain evolution reduced the domain switching energy, producing a slim P–E loop, while maintaining a large Pmax value and a colossal BDS. In addition to the domain perspective, SrTiO3 substitution enhanced the insulating character of the films by stabilizing the Fe3+/Fe2+ valence fluctuation and reducing the formation of oxygen vacancies, leading to further enhancement of BDS. For example, the BDS value of (1-x)BiFeO3-xSrTiO3 films increased from 2.77 MV cm−1 to 4.46 MV cm−1 as the SrTiO3 content increased from for x = 0.3 to x = 0.75. At such a high electric field, the films with x = 0.6 and 0.75 exhibited a giant Urec ~ 70 J cm−3.Correia et al.  depositen et al.  demonstrη = 33%) . The thin et al.  utilizedR) and tetragonal (T) domains in a cubic paraelectric matrix. If these polymorphs have competitive free energy, Landau phenomenological theory predicts the weakening of polarization anisotropy and lowering of the energy barrier between the R and T polarization states [3-xBaTiO3-0.45SrTiO3 solid solution (x = 0 to 0.4). Here, BiFeO3 and BaTiO3 were the hosts for the R and T domains, whereas SrTiO3 provided a cubic paraelectric matrix. The relaxor nature and energy storage performance of the (0.55−x)BiFeO3-xBaTiO3-0.45SrTiO3 solid solutions are shown in 3 gradually enhanced the relaxor nature, as can be seen from the wider peaks in the ε–T plots  and Mn addition. The LLD produced highly stoichiometric and perfectly crystalline films compared with standard deposited films, while the addition of Mn reduced the leakage current by creating vacancy trap centers. Using this method, they obtained an ultra-high value of Urec = 80 J cm−3 at a BDS of 3.1 MV cm−1 with η = 78% in 2.5 mol.% Mn-doped 0.25BiFeO3-0.75BaTiO3 thin films.Domain engineering techniques have been found to be more fruitful in thin films that exhibit polymorphic nanodomains, e.g., rhombohedral  and large volume; therefore, it is very promising for practical high-power energy storage applications.Although the energy density of thin films is superior to that of bulk ceramics, the usability of thin films is limited for low-power applications because of their small volume. In this context, multilayered structures have received great attention because the technology behind them is well known and inexpensive. A multilayered ceramic consists of a number of thin ceramic layers with thicknesses on the micrometer scale and internal electrode layers stacked in parallel and connected through terminal electrodes. It features both high BDS (on the MV cm0.85Nd0.15FeO3)-0.25BaTiO3 + 0.1wt.% MnO2 and Pt internal electrodes, compared with its bulk ceramic counterpart, as shown in −1 with Ust ~ 8.75 J cm−3, Urec~ 6.74 J cm−3, and η~77%. Yan et al. [3-0.33BaTiO3)-xNa0.73Bi0.09NbO3 multilayered ceramics with A-site cation vacancies. For an optimized content of Na0.73Bi0.09NbO3, i.e., x = 0.12, the multilayer exhibited Urec = 5.57 J cm−3 at E = 410 kV cm−1 with η = 83.8%. A multilayer ceramic with 7-μm-thick 0.61BiFeO3-0.33Ba0.8Sr0.02TiO3-0.06La(Mg2/3Nb1/3)O3 dielectric layers was reported to exhibit a much higher BDS, > 740 kV cm−1 [Urec of 10 J cm−3 with η ~ 72% at E = 730 kV cm−1. Wang et al. reported a large ∆P of ~ 34 μC cm−2 and a high BDS of 700 kV cm−1 in a multilayered ceramic composed of 16-μm-thick electrically homogeneous 0.62BiFeO3-0.3BaTiO3-0.08NdZr1/2Zn1/2O3 ceramic layers with a core-shell microstructure [2, Urec as high as 10.5 J cm−3 with η = 87% was obtained. Lu et al. [−1 in a multilayered ceramic with 10 mol% Bi(Mg2/3Nb1/3)O3 and 3 mol% Nb2O5-doped 0.6BiFeO3-0.4SrTiO3 with a core-shell microstructure. The device, with 8-μm-thick dielectric layers, exhibited Urec = 15.8 J cm−3 under E > 1 MV cm−1 with η = 75.2%. Wang et al. [−1 in a multilayered ceramic with 8-μm-thick 0.57BiFeO3-0.3BaTiO3-0.13Bi(Li0.5Nb0.5)O3 layers . While the bulk ceramic featured a Urec ~ 3.64 J cm−3 at E ~ 260 kV cm−1 with η ~ 75%, the multilayered ceramic exhibited a Urec = 13.8 J cm−3 with η = 81% because of the large enhancement of the BDS. The multilayered ceramic showed good temperature stability (< 10%) and frequency independence (< 5%) of Urec from 0.01 to 1 Hz, as well as fatigue resistance (< 5%) during 104 cycles of unipolar P−E loop tests in the temperature range from RT to 100 °C at E = 400 kV-cm−1.Wang et al.  reportedn et al.  reported kV cm−1 . This detructure . With seu et al.  obtainedg et al.  reported3 layers  having a3-based RFEs suffer from a low efficiency (<85%) despite having ultra-high discharge energy densities. Leakage-related high conductivity in BiFeO3-based RFEs produces apparent large values of Pr and wide hysteresis loops, which lead to a higher loss density. The conductivity usually increases at high temperatures, limiting their potential for high-temperature applications.High discharge energy density and energy efficiency are the primary requirements of an EES system. However, most BiFeO3-based compounds is associated with the loss of Bi2O3 during its synthesis at elevated temperatures (950–1050 °C), which creates ionic vacancies and Fe3+/Fe2+ valence fluctuation [3-based compounds has been effective for reducing the leakage current density by several orders of magnitude [2O3 loss. Adopting a low-temperature sintering method is expected to reduce the Bi2O3 loss and the associated leakage current densities by a few more orders. The addition of appropriate glass additives and/or low melting point compounds is expected to significantly reduce the sintering temperature [2O3 volatilization. Low-temperature sintering is also beneficial for the enhancement of BDS, as discussed in The high leakage current in BiFeOctuation ,104,105.perature , therebyUrec as a function of BDS for BiFeO3-based binary and ternary solid solutions fabricated in bulk ceramics, ceramic multilayers, and thin films. A high BDS usually provides a higher recoverable energy density. Therefore, enhancing the BDS is one of the most effective ways to increase the energy storage performance of RFEs. Microstructure engineering, such as the design of core-shell structures and domain engineering, has been found to be very effective for enhancing the BDS and realizing a dynamic relaxor ferroelectric phase. The BDS depends on the grain size (BDS ∝ 1/√G (G = average grain size)) and size distribution. Therefore, it is imperative to design a fine-grained microstructure with a uniform size distribution to obtain a high BDS. Another way to improve the BDS is to prepare a highly dense microstructure. Pores or cavities present in microstructures are usually filled with gaseous or liquid phases with lower permittivity than those of solid dielectrics. As a result, the voltage across the cavities or pores (Vc) is enhanced as per the following relation [Vapp is the applied voltage, εc and εr are the permittivity of the cavity and dielectric, respectively, and t and d are the size of the cavity and the thickness of the dielectric, respectively. A small pore can create a large electric field across it and cause local breakdown and internal discharge even at a low external voltage. Therefore, a highly dense microstructure is essential. Because BiFeO3-based materials are prone to the loss of Bi and Fe valence fluctuations during sintering at high temperatures, resulting in high electrical conductivity, sintering at low temperatures could be very beneficial. However, a detailed investigation of the microstructure control for the purpose of enhancing the energy storage performance employing the low-temperature sintering technique in BiFeO3-based dielectrics is yet to be conducted. The two-step sintering (TSS) method modified by Chen and Wan [In relation :(10)Vc=V and Wan  is one oEg is the energy bandgap, and ωmax is the maximum phonon frequency. Pure BiFeO3 has a direct bandgap of 3 eV [3-based RFEs via modulation of the energy bandgap. Qi et al. [3 (Eg ~ 3.28 eV) in 0.67BiFeO3-0.33BaTiO3 solid solutions. They demonstrated that the bandgap of the solid solutions (0.67-x)BiFeO3-0.33BaTiO3-xNaNbO3 (0 ≤ x ≤ 0.15) increases monotonically with increases in x from ~ 2.6 eV for x = 0 to ~ 2.95 eV for x = 0.15. The enhancement of BDS followed a similar trend with an increase in NaNbO3 content, where the average BDS increased from 230 kV cm−1 for x = 0 to 420 kV cm−1 for x = 0.15. Further studies on bandgap engineering can be explored to enhance the BDS and energy storage performance of BiFeO3-based relaxor ferroelectrics.Electronic breakdown is a crucial intrinsic mechanism for the breakdown of a dielectric in a large electric-field regime. Above a certain electric field, electrons in the valence band gain sufficient energy to jump to the conduction band. This results in an increase in the electron density in the conduction band and leads to a large current discharge. Therefore, an insulator exhibiting a wide forbidden energy gap can withstand a high electric field. The empirical relation between the BDS and energy band gap is given by the following relation :(11)EB=2 of 3 eV . Only a i et al.  reported3-based relaxor ferroelectrics is the electromechanical breakdown. Because a high electric field is required to obtain a large recoverable energy density, the dielectrics are under extreme electrostrictive strain, increasing the mechanical breakdown. This issue should be taken very seriously from the device viewpoint, as the devices undergo a large number of charge-discharge cycles. Although high recoverable energy density with high efficiency is obtained at high electric fields under laboratory conditions, the same samples might not be suitable for application purposes where they totally collapse because of mechanical failure under a large number of charge-discharge cycles. However, systematic studies on the issues of electromechanical breakdown in BiFeO3-based dielectrics are lacking. Rare-earth-ion doped BiFeO3 compounds with their composition lying across the morphotropic phase boundary could be interesting in this regard as these compounds feature high electromechanical strain. For example, Walker et al. [3 polycrystalline samples, associated with the electric field-induced phase transition and ferroelectric/ferroelastic domain switching.Another issue that has not gained much attention but is very critical in BiFeOr et al.  reported3-based relaxor ferroelectrics are projected to be potential Pb-free candidates for application in the field of high-energy-density storage and high-power-delivery systems. They are mostly fabricated in the form of ceramics, multilayers, and thin films from binary or ternary solid solutions with other perovskite oxides. BiFeO3-BaTiO3 and BiFeO3-SrTiO3, with their compositional ratios lying in the morphotropic phase boundary, are the two most widely studied BiFeO3-based binary solid solutions for high-energy-density storage. However, the typical value of recovering energy density of bulk ceramics of BiFeO3-based binary solid solutions is well below 3 J cm−3, with low efficiency . Ternary solid solutions exhibiting specific microstructures, such as core-shell structures and stripe-like nanodomains, have shown significant enhancement of the energy density properties. For example, 0.5BiFeO3-0.4SrTiO3-0.03Nb2O5-0.1BiMg2/3Nb1/3O3 bulk ceramic with a core-shell microstructure exhibited excellent energy storage properties, with Urec ~ 8.2 J cm−3 and η ~ 74%. The nanodomain-engineered bulk ceramic 0.57BiFeO3-0.33BaTiO3-0.1NaNbO3 with the addition of 0.1 wt% MnO2 and 2 wt% BaCu(B2O5) also featured a similar value of Urec (~ 8.12 J cm−3); however, it had a higher value of η ~ 90%. The microstructure design approach has been found to be attractive for enhancing the energy storage properties of thin films and multilayers. The BDS values of thin films and multilayers can be as high as MV cm-1, which is very helpful for further increasing the recoverable energy density. Thin films of a (0.55−x)BiFeO3-xBaTiO3-0.45SrTiO3 solid solution (x = 0–0.4) containing polymorphic nanodomains showed a recoverable energy density of ~ 110 J cm−3 at a BDS ~ 5 MV/cm with an efficiency of ~ 80%, which is the highest among BiFeO3-based systems. The achievement of such excellent energy storage properties is very encouraging for applications in low-power, small-size devices. However, for large-scale applications, it is necessary to focus on the fabrication of multilayers, as they deliver high power and are easy to fabricate. Although relatively few studies have been conducted on BiFeO3-based multilayers, their energy storage performance is very encouraging (Urec ~ 15 J cm−3), with a high BDS of ~ 1 MV/cm. More studies should be conducted in the field of BiFeO3-based multilayers to optimize structural parameters such as the thickness of dielectric layers and number of layers so that a better energy storage device that can deliver high output power can be fabricated at a low cost.BiFeO


Escherichia coli (STEC) O26:H11 is an emerging foodborne pathogen of growing concern. Since current strategies to control microbial contamination in foodstuffs do not guarantee the elimination of O26:H11, novel approaches are needed. Bacteriophages present an alternative to traditional biocontrol methods used in the food industry. Here, a previously isolated bacteriophage vB_EcoM_VR26 (VR26), adapted to grow at common refrigeration temperatures (4 and 8 °C), has been evaluated for its potential as a biocontrol agent against O26:H11. After 2 h of treatment in broth, VR26 reduced O26:H11 numbers (p < 0.01) by > 2 log10 at 22 °C, and ~3 log10 at 4 °C. No bacterial regrowth was observed after 24 h of treatment at both temperatures. When VR26 was introduced to O26:H11-inoculated lettuce, ~2.0 log10 CFU/piece reduction was observed at 4, 8, and 22 °C. No survivors were detected after 4 and 6 h at 8 and 4 °C, respectively. Although at 22 °C, bacterial regrowth was observed after 6 h of treatment, O26:H11 counts on non-treated samples were >2 log10 CFU/piece higher than on phage-treated ones (p < 0.02). This, and the ability of VR26 to survive over a pH range of 3–11, indicates that VR26 could be used to control STEC O26:H11 in the food industry.Shiga toxin-producing  Escherichia coli (E. coli) is a gram-negative, facultatively anaerobic, rod-shaped bacterium of the genus Escherichia, family Enterobacteriaceae, that is commonly found in the gut of humans and warm-blooded animals [E. coli are non-pathogenic, however, a number of strains, such as Shiga toxin-producing E. coli (STEC), can cause severe foodborne disease. The Food and Agriculture Organization (FAO) together with the World Health Organization (WHO) estimated that more than 1.2 million new cases resulting in 128 deaths and nearly 13,000 Disability Adjusted Life Years (DALYs) occur each year because of contamination by the foodborne STEC [ animals . Since t animals ,3. Most rne STEC . Based orne STEC . The assrne STEC ,6,7,8. Arne STEC ,10,11,12Listeria monocytogenes on foods [A promising alternative in the food industry for the elimination of STEC from fresh produce is bacteriophages (phages): viruses that infect and kill bacteria . There aon foods . Since ton foods ,15. Moreon foods ,15. Withon foods ,15,16,17The most common STEC serotype associated with human disease is, indeed, O157:H7 ,4,18. HoE. coli phage vB_EcoM_VR26 (VR26) and showed that none of the predicted VR26 proteins exhibited sequence homology with known integration-related proteins, antibiotic resistance determinants, or virulence factors. [E. coli phages, all belonging to the family Myoviridae, have been shown to be adapted to lyse their hosts at low temperatures (4–7 °C) [Previously, we presented the complete genome sequence of a T4-related factors. . VR26 isfactors. ,30. To d(4–7 °C) ,30. Due E. coli MH1 was a gift from Prof. Kenneth N. Kreuzer [Bacterial strains used in this study are listed in  Kreuzer . For inoFirst, the ability of phage VR26 to infect different serotypes of STEC at 22 and 8 °C was determined using a quantitative spot dilution test as described in . The hos10 PFU/mL) was added to the tubes containing sterile PBS (90 μL) with pH values ranging from 2 to 13, adjusted with NaOH or HCl. The tubes were allowed to incubate at 22 °C for 60 min, then a double agar overlay plaque assay was performed [10 PFU/mL).To determine the effect of pH on the stability of phage VR26, 10 μL of phage suspension  was purchased at a grocery store IKI in Vilnius (Lithuania). Outer leaves of the lettuce were removed whereas the inner leaves were cut into pieces of ~1 by ~1 cm with a sterile, stainless steel scalpel. The samples were then transferred to the sterile Petri dishes and spot inoculated with 10 μL of STEC O26:H11 (4 log10 CFU/cm2). Inoculated lettuce was then treated with the phage VR26 suspension (8 log10 PFU/cm2) or PBS (control) and the samples were stored at 4, 8 and 22 °C. After 2, 4, 6, and 24 h of incubation, one sample per treatment was withdrawn, transferred to 2 ml Eppendorf tube with 1 mL PBS, and then homogenized with sterile bars and vortexed. Homogenates or serial dilutions of homogenates were plated on LB agar, and bacterial viable counts were determined after overnight incubation at 37 °C.Iceberg lettuce (t-test (two-tailed) was employed to calculate the statistical significance (p < 0.05) between treatments and control.All experiments were run in triplicate and bacteria/phage counts were determined by duplicate plating. Colony counts are presented as mean values with standard deviation. Unpaired E. coli, representing different STEC serotypes: O157:H7, O103:H2, O111:H8, O145:H128, and O26:H11. A quantitative spot dilution test  in the viable titer was observed only after the exposure to extremely acidic (pH ≤ 2) or basic (pH ≥ 12) conditions.Since STEC strains are known to tolerate a somewhat broad range of pH (4.0–9.0) , bacteriTaken together, the results of the pH-inactivation assay suggested that VR26 sustained its lytic property through a broad range of pHs, indicating greater stability of this phage compared with other O26:H11-infecting  or even The ability of phages to survive acidic (pH ≤ 5) or alkaline (pH ≥ 7) conditions is a desirable trait in phage-based applications. It has been suggested previously that acid-tolerant phages could be utilized to control pathogen colonization in the gastrointestinal tract (pH 1–5) of food animals if used as animal feed and/or water supplements ,38,39,40p < 0.01) of O26:H11 in LB broth was observed after the addition of VR26, at both 22 °C and 4 °C  at 22 °C, and by approximately 3 log (p < 0.01) at 4 °C. VR26 was able to reduce (p < 0.01) numbers of O26:H11 to <10 CFU/mL after 4 h of incubation at 4 °C, whereas, after 24 h, the viable cells were not detected (After 2 h of incubation VR26 reduced O26:H11 counts by more than 2 log (detected a. At 22 detected b.10 CFU/mL (p < 0.001) at 37 °C but the efficacy of the phage (p < 0.001) was reduced after 7 h of incubation and regrowth of bacteria was seen thereafter. Mangieri et al. [As discussed above, two reports describing phages capable of reducing O26:H11 levels in a broth culture system under standard laboratory conditions may be found in literature ,28. In aThe results presented here are hardly comparable with those obtained during the aforementioned studies since bacterial challenge with phage VR26 assay was performed at 4 and 22 °C, not at 37 °C. However, in the food industry, foods are hardly ever exposed to 37 °C. Hence, the ability of VR26 to effectively reduce O26:H11 at both refrigeration and room temperature conditions indicates that this bacterial virus shows potential in the decontamination of foods, especially at post-harvest.p < 0.02) the viable counts of O26:H11 by ~2 log10 CFU/piece after 2 h of treatment at all temperatures. No survivors were detected after 6 and 4 h of treatment at 4 and 8 °C, respectively.The ability of the phage VR26 to reduce O26:H11 on lettuce was examined at 4, 8, and 22 °C. As seen in 10 CFU within 6 h and, although bacterial regrowth was observed afterward, O26:H11 counts on non-treated lettuce pieces were >2 log10 CFU/piece higher than on phage-treated samples (2.3 ± 0.06 log10 CFU/piece) (p < 0.02).At 22 °C, the pathogen was reduced by more than 2 logBacterial regrowth observed during phage treatment at room temperature was also reported by a number of previous studies ,42,43,44E. coli virus VR26 in the biocontrol of foodborne pathogens was evaluated. Our results clearly demonstrate that bacteriophage VR26 retained its lytic activity after being exposed to a range of pH from 3 to 11 and efficiently reduced O26:H11 counts, both in broth culture and on fresh lettuce, at all temperatures tested. Moreover, the greatest efficacy of VR26 on food was observed at common refrigeration temperatures (4 and 8 °C), suggesting that this phage could potentially be used as a biocontrol agent against STEC O26:H11 in the food industry.Here, for the first time, the efficacy of low-temperature


Thermoanaerobacterium saccharolyticum (TsaBgl). TsaBgl exhibited its maximum hydrolase activity on p-nitrophenyl-β-d-glucopyranoside at pH 6.0 and 55 °C. The crystal structure of TsaBgl showed a single (β/α)8 TIM-barrel fold, and a β8-α14 loop, which is located around the substrate-binding pocket entrance, showing a unique conformation compared with other structurally known Bgls. A Tris molecule inhibited enzyme activity and was bound to the active site of TsaBgl coordinated by the catalytic residues Glu163 (proton donor) and Glu351 (nucleophile). Titration experiments showed that TsaBgl belongs to the glucose-tolerant Bgl family. The gatekeeper site of TsaBgl is similar to those of other glucose-tolerant Bgls, whereas Trp323 and Leu170, which are involved in glucose tolerance, show a unique configuration. Our results therefore improve our knowledge about the Tris-mediated inhibition and glucose tolerance of Bgl family members, which is essential for their industrial application.β-Glucosidases (Bgls) convert cellobiose and other soluble cello-oligomers into glucose and play important roles in fundamental biological processes, providing energy sources in living organisms. Bgls are essential terminal enzymes of cellulose degradation systems and attractive targets for lignocellulose-based biotechnological applications. Characterization of novel Bgls is important for broadening our knowledge of this enzyme class and can provide insights into its further applications. In this study, we report the biochemical and structural analysis of a Bgl from the hemicellulose-degrading thermophilic anaerobe  Biomass-derived fuels are believed to play major roles in the global energy future, and technological development driving down costs are important for competing with petroleum and other alternative fuels . LignoceRetaining Bgls have been found in the glycoside hydrolase (GH) families GH1, GH3, GH5, GH9, GH30, and GH116 of the CAZy database . The GH1Thermoanaerobacterium saccharolyticum is a hemicellulose-degrading thermophilic anaerobe and a biological catalyst for the conversion of cellulosic biomass to ethanol [T. saccharolyticum, including endoxylanase [Thermophilic anaerobic bacteria include species with the natural ability to digest and ferment polysaccharides that constitute lignocellulosic biomass ,20. Ther ethanol . This or ethanol . This sp ethanol . Variousxylanase , β-xylosxylanase , glucuroxylanase , and amyxylanase ; howeverT. saccharolyticum (TsaBgl). We investigated the substrate specificity, optimum temperature and pH, and thermal stability of TsaBGl and the effects of metals on this enzyme. A high-resolution crystal structure of TsaBGl was determined at 1.7 Å, showing the binding of the inhibitor Tris to the active site. Moreover, a titration experiment of TsaBgl using Tris and glucose molecules was performed. Our results showed that the favorable properties of TsaBgl could be used in the field of biotechnology and provide useful information to broaden our structural knowledge of the Bgl family.In this study, we report the biochemical and structural analysis of Bgl from E. coli host expression, was overexpressed in E. coli BL21 (DE3). The protein was eluted at a molecular mass position of approximately 50 kDa during size exclusion chromatography, which was almost identical to the theoretical molecular weight of TsaBgl (51.728 kDa), indicating that TsaBgl was a monomer in solution  as substrates. pNPG and pNPC were selected based on the fact that Bgls often possess hydrolyzing activity towards soluble cellodextrins, including cellobiose and cellotriose. In addition, α-bonded- pNPαG was chosen to examine the preferred linkage type in TsaBgl. Our results show that TsaBgl exhibited the highest hydrolase activity on the artificial substrate pNPG with a specific activity of 23.6 U/mg. Moreover, the relative hydrolase activity of TsaBgl toward pNPC was 19.71% of that toward pNPG; however, no enzyme activity was observed on pNPαG  and 55 °C was defined as 100%, where the specific activity was 23.6 U/mg. The hydrolase activity of TsaBgl was investigated in the pH range of 4.0–10.0 using a four buffer system (p < 0.05)  .The enzyme activity of TsaBgl was then determined over a temperature range of 30–70 °C d. Maximu2+, Zn2+, Cu2+, and Fe2+ compared with that in the absence of any metal. The inhibition trend was as follows: Cu2+ > Zn2+ > Cd2+ and Fe2+, with relative activities of 8.7%, 25.2%, 68.2%, and 71.2%, respectively  exhibited a considerably higher B-factor than the whole protein (17.37 Å2).TsaBgl contained 16 α-helices and 13 β-strands, forming a single (β/α)5 × 50 Å a. During5 × 50 Å , which w5 × 50 Å c. Intera5 × 50 Å , and the5 × 50 Å a. Temper5 × 50 Å b. In parHalothermothrix orenii , Exiguobacterium antarcticum , Clostridium cellulovorans , and Thermotoga maritima . Superimposition of the TsaBgl structure with those of the homologous Bgls showed high similarity, with r.m.s. deviations of 0.49–0.61 Å  [Oryza sativa  [Micrococcus antarcticus  [Humicola insolens  [Oryza sativa subsp. indica (Rice)  [Spodoptera frugiperda  [In the Protein Data Bank, we found eight crystal structures of Tris-bound Bgls: Bgl from e: 3AHY) , Oryza sl, 3PTK) , Microcol, 3W53) , Humicoll, 4MDO) , Oryza snd 4RE3) , Spodoptl, 5CG0) , and frol, 5CG0) ,35. The On the other hand, in the structures OsaBgl2, the Tris molecule was located in the binding channel or at the surface of the TIM-barrel.Glucose produced by the Bgl-catalyzed decomposition of cellobiose is known to act as a product inhibitor . This isT. saccharolyticum. The growth pH and temperature of the T. saccharolyticum strain were 3.85–6.35 and 30–66 °C, respectively. Purified recombinant TsaBgl showed high activity at pH 6.0 and 50–55 °C, but the activity decreased to less than 60% at other pH and temperature values. In addition, TsaBgl showed the highest enzyme activity at 55 °C, although thermostability experiments revealed that the residual activity was reduced by 37% after 10 min. As a result, the in vitro enzymatic activity of purified TsaBgl was not as consistent over a broad pH and temperature range as the growth of the strain.Bgls are essential terminal enzymes of the cellulase system and attractive targets for lignocellulose-based biotechnological applications. To discover a novel Bgl and broaden the knowledge of this enzyme class at the molecular level, we performed biochemical and structural analyses of Bgl from the hemicellulose-degrading thermophilic anaerobe Thermoanaerobacteriumthermosaccharolyticum (Thermoanaerobacteriumthermosaccharolyticum (named TthBgl). Despite the close relationship between TsaBgl and TthBgl, there are several notables differences in their enzymatic activity. Specifically, the optimum temperature of TsaBgl and TthBgl for enzyme activity are 55 and 70 °C, respectively. Moreover, the activity of TsaBgl was not affected by Mn2+ and inhibited by Fe2+ while the activity of TthBgl was activated by both Mn2+ and Fe2+. Such a discrepancy may arise from a difference in the reaction buffer between the two studies (sodium phosphate vs. imidole-potassium). Since phosphate often forms an insoluble complex with metal ions, in our study, effects of Fe2+ and Mn2+ might have been either insignificant or inhibitory in phosphate buffer. Furthermore, in the current study, TsaBgl was incubated with metal ions in the presence of the substrate while TthBgl was pre-incubated with metal ions prior to substrate addition. Since the latter had a longer time for direct interaction with metal ions that might explain the discrepancy between the two enzymes regarding activities in the presence of Mn2+.A previous study had classified the Bgl family into five clades: Clade I (mainly mesophilic bacteria), Clade II (GH1 Bgl from fungi), Clade III (GH3 Bgl from bacteria), Clade IV (GH3 Bgl from fungi), and Clade V (thermophile or Bacillus GH1 Bgl) . Taxonomolyticum . In addimK value of TsaBgl, which represents binding affinity, was slightly lower or comparable to that of other Bgls, indicating for TsaBgl, in general, to have a higher or similar binding affinity relative to GH1 Bgls. While the catk values, the turnover number of substrate, were similar across TsaBgl, EanBgl, and CceBgl, TsaBgl showed higher values than TrBgl2 and NkoBgl by approximately 3 and 13 fold, respectively, whereas this value was 24-fold lower than that of TmaBglA. The overall catalytic efficiency of TsaBgl for pNPG (cat/Kmk) compared to that of EanBglA was higher than those reported for TreBgl2 and NkoBgl, but 7- and 24-fold lower than those reported for CceBglA and TmaBglA, respectively.Kinetic parameters of TsaBgl were compared to those of other Bgls belonging to GH family 1. The parameters of different Bgls showed distinct values . The Km cat/Kmk) of TsaBgl was similar to that of EanBglA, another glucose-tolerant GH1, but much higher than those of TrsBgl2 and NkoBgl  thoPrevious studies have identified several key amino acids that influence glucose tolerance, most of them at the substrate entrance site ,45,46. ITrichoderma harzianum Bgl (ThaBgl) and the high glucose-tolerant Humicola insolens Bgl (HinBgl). TsaBgl showed amino acid sequence identities of 38.0 and 36.0% with ThaBgl-Mut and HinBgl, respectively. Structural superposition of ThaBgl-Mut and HinBgl to TsaBgl showed r.m.s. deviation of 0.723 and 0.678 Å, respectively. In a previous study of ThaBgl, double amino acid replacements L167W/P172L in wild-type ThaBgl (ThaBgl-WT) dramatically increased the properties of glucose-tolerant ThaBgl (ThaBgl-Mut) [cat/Kmk) of ThaBgl-Mut (7.85 ± 0.22 s−1/0.41 ± 0.30 mM) was approximately 5-fold higher than that of ThaBgl-WT (4.92 ± 0.54 s−1/1.17 ± 0.66 mM). In the crystal structure, ThaBgl-WT showed a wider aglycone channel measuring ~11 Å in width between Trp338 and Pro172  . The catd Pro172 a, wheread Pro172 b. This ad Pro172 c. These d Pro172 d, similad Pro172 e. To betIn conclusion, our biochemical and structural analysis of TsaBgl is expected to improve our understanding of the Bgl family and provides new insights into the utilization of Bgl for industrial applications.Escherichia coli BL21 (DE3). Cells were grown at 37 °C in LB medium with 50 mg/mL ampicillin until reaching an OD600 of 0.4. Protein expression was induced with 0.5 mM isopropyl-d-1-thiogalactopyranoside, and the cells were cultured at 18 °C for 18 h. After cell harvesting by centrifugation, the cells were resuspended in lysis buffer containing 50 mM Tris-HCl (pH 8.0), 200 mM NaCl, and 20 mM imidazole. The cells were lysed by sonication on ice, and cell debris was removed by centrifugation at 18,894× g for 30 min. The supernatant was loaded onto a Ni-NTA column . The column was washed and equilibrated with lysis buffer, and the protein was subsequently eluted with a buffer containing 50 mM Tris-HCl (pH 8.0), 200 mM NaCl, and 300 mM imidazole. To remove the N-terminal hexahistidine-tag, 10 units of thrombin  were added per milligram of TsaBgl fraction and incubated overnight at 20 °C. The protein was concentrated using a Centricon filter , applied onto a Sephacryl-100 column , and eluted with a buffer containing 10 mM Tris-HCl (pH 8.0) and 200 mM NaCl. Protein concentration was determined based on the Bradford assay by measuring the absorbance at 595 nm, and subsequently checking against the calibration curve constructed using bovine serum albumin (BSA) as the standard.A construct containing codon-optimized full-length TsaBgl (accession number: MW655874) with an N-terminally fused hexahistidine-tag was synthesized and cloned into the pBT7 vector . Recombinant DNA was transformed into p-nitrophenol (pNP) from the enzymatic hydrolysis of p-nitrophenyl-β-d-glucopyranoside (pNPG), p-nitrophenyl-β-d-cellobioside (pNPC), and p-nitrophenyl-α-d-glucopyranoside (pNPαG) . In general, enzyme reactions were performed with 7.5 µg/mL TsaBgl in 100 mM sodium phosphate (pH 6.0) containing 2.5 mM pNPG at 55 °C for 5 min in a total volume of 100 µL. Enzyme reactions were terminated by adding 66 µL of 1 M Na2CO3. The amount of pNP produced was measured at 405 nm using a Synergy H1 microplate reader  at 25 °C and calculated from a calibration curve using pNP as the standard. One unit (U) of enzyme activity was defined as the amount of enzyme required to release 1 μmol of pNP per min from 2.5 mM pNPG.Enzyme activity of TsaBgl was investigated by measuring the formation of To examine the effects of pH on the activity of TsaBgl, enzyme activity was measured in the pH range 4.0–10.0 at 55 °C using a four buffer systems: 100 mM sodium citrate (pH 4.0–6.0), 100 mM sodium phosphate (pH 6.0–7.0), 100 mM Tris-HCl (pH 7.0–9.0), and 100 mM glycine-NaOH (pH 9.0–10.0).pNPG in 100 mM sodium phosphate buffer (pH 7.0) for 5 min at different temperatures ranging from 30 to 70 °C. For thermal stability analysis, residual enzyme activity was measured after pre-incubating TsaBgl at specified temperatures in the absence of substrate. In detail, 7.5 µg/mL TsaBgl was pre-incubated at temperatures ranging from 30 to 80 °C for 10 min in 100 mM sodium phosphate (pH 6.0) before running the enzyme reaction for 5 min by adding 2.5 mM pNPG under standard conditions (55 °C and pH 6.0). Relative activity (%) represents the residual enzyme activity relative to that determined without pre-incubation.To investigate the effects of temperature on the activity of TsaBgl, the enzyme reaction was examined using 2.5 mM 4, MgCl2, CaCl2, MnCl2, FeCl2, FeCl3, CoCl2, NiCl2, CuCl2, ZnCl2, CdCl2, and CeCl2.To investigate the effects of metal on the activity of TsaBgl, enzyme activity was measured under optimal conditions (55 °C and pH 6.0 in 100 mM sodium phosphate) in the presence of 1 mM of various metal ions, such as LiSOmK, catk, and catk/mK) were determined by conducting the enzymatic reaction with 7.5 µg/mL TsaBgl in various concentrations (0 to 10 mM) of pNPG at the optimal condition (pH 6.0 and 55 °C). Data fitting was performed using the Michaelis-Menten equation based on SigmaPlot 12.3 software .Kinetic parameters  or glucose (0–830 mM) from 1 M (pH 7.0) or 2 M stock, respectively. In case of Tris inhibition, the enzyme assay was performed at pH 7 (100 mM sodium phosphate) to prevent the potential effect of Tris on pH change in the reaction solution. The relative activity (%) represents the enzyme activity relative to that determined in the absence of Tris or glucose.w/v) PEG 4000, and 0.2 M MgCl2. Crystallization conditions were further optimized by the hanging-drop vapor diffusion method at 22 °C. Rod-shaped TsaBgl crystals were obtained within 1 week by mixing 2 µL protein solution with 2 µL reservoir solution, followed by equilibration against 500 µL of 0.1 M Tris-HCl (pH 7.0), 27% (w/v) PEG 3350, and 0.2 M MgCl2.Purified TsaBgl was concentrated to 30 mg/mL using a Centricon filter . Initial crystallization conditions of TsaBgl were screened using the sitting-drop vapor diffusion method at 22 °C using a Crystal Screen Kit . Microcrystals were obtained in reservoir solution containing 0.1 M Tris-HCl (pH 8.5), 30% (v/v) ethylene glycol and then flash-frozen in a liquid nitrogen stream at 100 K. Diffraction images were indexed, integrated, and scaled using HKL2000 [X-ray diffraction data were collected on a Beamline 11C at the Pohang Light Source II  . The cry HKL2000 .The structure was solved by molecular replacement with MOLPEP , using t


However, there is a need for improved minimally invasive, void-filling materials to provide mechanical support, biocompatibility, and surface erosion characteristics to ensure consistent tissue support during the healing process. Herein, soft, elastomeric aliphatic polycarbonate-based materials were designed to undergo photopolymerization into supportive soft tissue engineering scaffolds. The 4D nature of the printed scaffolds is manifested in their shape memory properties, which allows them to fill model soft tissue voids without deforming the surrounding material. In vivo, adipocyte lobules were found to infiltrate the surface-eroding scaffold within 2 months, and neovascularization was observed over the same time. Notably, reduced collagen capsule thickness indicates that these scaffolds are highly promising for adipose tissue engineering and repair. Shape memory scaffolds are needed for minimally invasive tissue repair and void filling. Here the authors report on the development of 4D printed polycarbonate-based scaffolds with surface degradation properties which fill voids without deforming tissue and allow for tissue ingrowth with reduced immune response. Tissue scaffold design has the potential to revolutionize patient care, with supportive materials being essential to resolve wounds that are otherwise untreatable. Porous materials in particular offer tissues the mechanical support that is necessary for them to rapidly infiltrate a space. This concept has been successfully demonstrated in applications such as siloxane maxillofacial implants, polyurethane cardiovascular occlusive devices, and collagen-derived adipose tissue scaffolds11. However, current porous materials leveraged for these applications are limited by their pore morphological distribution, in addition to other issues such as post-processing requirements prior to use21. 3D printing has emerged as one of the most promising methods to overcome these processing limitations as a consequence of the reproducible, interconnected pore features with micron-scale resolution that can be manufactured in precise designs. However, a significant current limitation in the field is the lack of suitable materials that can be processed efficiently by additive manufacturing techniques, particularly in the layer-integrating photopolymerization methods, coupled with long-term biocompatibility in vivo27. The primary 3D printing material focus has remained on acrylate- and epoxide-containing polymers, which have low toxicity thresholds, while the primary degradable biomaterial focus has been directed towards poly(l-lactic acid) (PLLA), which is limited by its poor processability in photopolymerizations and its acidic degradation products29. An alternative synthetic materials platform that can balance minimally invasive behavior with nutrient/waste diffusion and support tissue regrowth as they degrade in vivo to nontoxic byproducts without succumbing to the limitations of current tissue engineering therapies is needed.Nature has created tissue designs that are ideally suited for their purpose: bone possesses high strength with enough flexibility to not be brittle; arteries are elastomeric without being flimsy; and adipose tissue is soft and yielding while being durable. Despite these evolutionary materials templates, tissue engineering is currently limited by an urgent, unmet need for improved restorative and healing techniques to address tissue defects. An example of this is in soft tissue repair, such as following breast cancer treatment, a mentally and physically debilitating cancer affecting millions of patients globally31. Shape-changing polymers open avenues into printing minimally invasive medical devices and scaffolds as well as biomimetic designs, as smaller footprints result in lower surgical trauma without compromising, and in some cases enhancing, the patient outcomes35. While the current state-of-the-art in minimally invasive biomaterials, such as foams, display the desired reduced therapeutic footprint, their uneven pore distribution restricts nutrient diffusion. Furthermore, improvements in reproducibility of shape-changing behaviors and increased materials choice to improve biocompatibility and biodegradation would represent significant advances36. Despite the potential advantages of a 3D-printed, non-inflammatory, resorbable, and shape responsive polymer, there are few examples of 4D printing of minimally invasive, clinically relevant material designs37.Recently, the concept of 4D materials has emerged, where a 3D-printed material displays behavior such as shape memory, swelling, or controlled degradation in a 4th dimension, typically timeHerein, we describe an approach to deliver soft tissue engineering constructs by developing 4D printable resin inks that can be photopolymerized into patient-specific, self-fitting scaffolds that can be printed with a wide range of surface morphologies and display tunable shape memory with high strain recoveries and low expansion forces. Furthermore, the materials degrade by a surface-erosion profile to nonacidic products and display excellent cytocompatibility and biocompatibility. By focusing on the design of a material with a unique combination of features, we have been able to achieve a minimally invasive 4D structure that could reduce surgical impact while enhancing rates of healing and patient recovery.Mn) ca. 2 kDa and a dispersity, ƉM, of 1.1  and Fourier-transformed infrared (FT-IR) spectroscopy confirmed the presence of carbonyl, hydroxyl, and alkene groups which, in addition to size exclusion chromatography (SEC) analysis, confirm oligomer synthesis  of aliphatic cyclic carbonates was selected. This process yielded homo- and co-oligocarbonates from allyl- and norbornene-containing monomers  and 2-norbornene-5,5-bis(hydroxymethyl) trimethylene carbonate (NTC), respectively) with a targeted number-average molar mass . In line with rapid solidification of the resin upon exposure to light, photorheological analysis revealed a peak loss factor ratio at 2 s after irradiation and a dramatic increase in both storage modulus and complex viscosity, from 179.6 ± 17.5 Pa to 1.5 ± 0.4 MPa and 3.1 ± 0.1 Pa·s to 23.1 ± 8.3 MPa·s, respectively, followed by a plateau even upon further irradiation s (PCUs) Fig. , or by tion Fig. . 1H NMR ion Fig. . The che27, a stereolithographic-type process, all compositions could be used to produce porous scaffolds with potential for void-filling devices, without the need for additional processing to remove foam cell membranes or additives typically found in porous biomaterials15. To demonstrate the ability to print a range of scaffold geometries, scaffolds were printed with pore sizes ranging from 200 to 1500 µm, and surface areas of between 1 and 3 cm2. Analysis of the resulting scaffolds by µCT revealed that the measured pore size values match the theoretical pore sizes calculated from the theoretical porous structure renderings within 5% error 41. Attempts to manufacture scaffolds with pore sizes below this threshold were limited by irreproducible pore interconnectivity as a consequence of increased viscosity that prevents resin mobility out of the crosslinked pores and resulted in overcuring.Cytocompatibility screening was performed using 2D surfaces, in order to assess compositional factors prior to final scaffold development, and in 3D scaffolds as a more realistic model. No significant differences were found regarding proliferation or morphology when assessed over a 7-day period  in both direct and indirect contact assays based upon ISO 10993 protocols. All cell types, including macrophages, adipocytes and fibroblasts (murine and human), representative of those found in native adipose tissue, displayed good cell spreading and adhesion Fig. –36. No s15. The gas-blown scaffolds displayed a high degree of cytocompatibility over 7 days, without the 3D proliferation that was observed in the 3D-printed structures. In the foams, pre-osteoblasts proliferated along the top outer layers of the foam with some also being found along the bottom surface but critically, in contrast the 3D-printed structures  of cast films examined with dynamic mechanical analysis (DMA). Similarly, the mechanical performance of the materials could also be controlled over a wide range by modulation of the resin composition . The carbonate monomer ratio and the presence of urethane linkages were used to tune the glass transition temperature across a range of more than 100 °C in both dry and solvated conditions Fig. . It was ion Fig. , Table 1MPa Fig.  showing MPa Fig. .Fig. 4Th9. By comparison, the stiffer PNTCTX (compressive elastic moduli of 12.3 MPa) displayed a gradual reduction in recovered strain after each cycle under ambient conditions, decreasing initially by ∼25% before stabilizing by cycle 15 at ~30% of the original strain; increased delays between compression cycles resulted in further recovery of the material when unloaded.Generally, all of the polycarbonate scaffolds undergo compression of up to 85% without catastrophic failure Fig. , and abo43. Cyclic compressive testing of alginate gels that contain 3D-printed scaffolds, similar to the testing of bare scaffolds, was further used to examine the scaffold migration and risk of soft tissue damage that result from the scaffold’s presence42. After a mock surgical opening using an eye-shaped incision, minimal changes in mechanical behavior were recorded for the compression of gels that contained scaffolds lds Fig. . This inTg (both wet and dry) provided direct correlations with strain recovery behavior for compressed polycarbonate scaffolds (Table Tg (tan δ peak) and 20 °C above Tg, respectively. Conversely, increasing NTC content reduced scaffold elasticity and the recovery speed, presumably as a function of scaffold Tg, thereby reducing void filling in irregularly shaped rigid voids  to achieve recovery within 10 mins in the alginate void. PNTCTX scaffolds conform to the soft void with 100% void filling and 90% strain recovery (measured at the center of the scaffold), displaying a low expansion force attributed to minimal polymer chain reorientation as a result of their high Tg. Unlike the PTMPTCX scaffolds, which display only decreasing expansion force with immersion  and an initial relaxation rate of 1.3 mN·s−1 , PNTCTX displays an increasing tan δ and storage moduli at 37 °C in PBS, followed by a gradual decrease corresponding with the material’s creep response that is indicated by a peak expansion force of 0.71 N ± 0.19 N (at ∼3 mins), and an average relaxation rate of 0.3 mN·s−1 . This is ideal behavior because it requires activation of the shape memory response by a surgeon but also enables self-fitting in the soft void. In vivo, this will allow for the void-shape fixing without personalization of the scaffold , and over time the ingrowth of clotting factors and tissue would hold the polymer in the final, void-fitted shape  within 45 s, which distorted the alginate by ∼15% (maximum strain) and decreased void-filling efficiency to ~90% as a function of scaffold shape. By comparison, the PNTCTX scaffold displayed slower shape recovery. Passive shape recovery at 37 °C required ∼50 min for full 100% strain and had to be stimulated using H44, this trend was similar to what is found using static gravimetric analysis with both films and printed, as well as decreased concentrations of NaOH and PBS solution, porous scaffolds, albeit with surface erosion occurring more rapidly as a consequence of the surface deformation caused by mechanical loading as well as the hydrolytic solution to as previously demonstrated in polyesters. All of the materials degraded through a surface-erosion behavior that is demonstrated by the gradual reduction in strut cross-sectional area in printed scaffolds , which by extrapolation indicates that total mass loss would most likely occur for the materials within 20 months. Comparatively, the PLLA control materials did not display significant mass loss, which indicates that minimal chain fragmentation is taking place in this same time period. The degradation displayed by the PTMPTCX and PNTCTX would provide sufficient support for more than a year, a seemingly ideal time frame that allows for mature tissue ingrowth before the mechanical support of the scaffolds is sufficiently reduced via degradation45. Spectroscopic analysis of the implanted samples by FT-IR spectroscopy supports this claim, where minimal shifting of the carbonyl peak indicates less than 30% mass loss has occurred by semi-quantitative analysis of the carbonyl change11. This is further supported by 1H NMR spectroscopic analysis of extracted samples.In vivo analysis was performed over 4 months in murine subcutaneous implant studies comparing PLLA disks, nonporous PTMPTCX disks, and printed, porous disks of PTMPTCX and PNTCTX with 500 µm pores. Material degradation was evaluated using swelling Fig.  and gel 48. For nonporous polycarbonate-derived material disks, lobules were found within 100 µm of the material surface. Adipocyte shape in vivo further reflects the positive response to the surface, as the characteristic round morphology is found within both the lobules as well as individually51. Our results indicate that nearly 40% of the infiltrated tissue is represented by adipocyte lobules, with fibroblasts representing another majority of the tissue; PLLA did not display this type of integration over the same time period. Capsule formation around all examined implants was less than 200 µm thick, well below the 500 µm threshold used for biocompatibility in other studies52. Importantly, the capsule formation was reduced with increased surface area; the porous implants displayed approximately half the capsule thickness (∼50 µm) as solid polycarbonates, which suggests that there may be a benefit to the 3D structure for tissue engineering implants; PLLA displayed ∼120 µm capsules. Macrophage presence  and Trichrome stains Fig. –50 revea55. One of the main failures in contemporary clinical techniques for restoring soft tissue, such as in adipose repair using autologous fat transplantation, is the 40–60% loss of graft volume as a result of poor graft vascularization post-implantation, and aspirated adipocytes are easily damaged by the mechanical force of the procedure, ultimately leading to cyst and localized necrosis that causes immune response and loss of the graft. Neither calcification was found in the implants, nor was necrosis.Vascular bud formation and vascularization occurred by 2 months, with several small, mature vessels found at 4 months in the surrounding tissue but no additional budding. Vascular budding allows for healing to occur, and then ideally will be reduced to match the original tissue as seen here, as adipose tissue is typically not heavily vascularized57. Of particular concern are traumas where tissue is removed, leaving a void that if left unfilled will result in deformity and cosmetic changes. 3D-printed synthetic scaffolds offer a unique opportunity to resolve tissue voids without the limitations of biological materials. Many synthetic materials suffer from degrading to acidic byproducts, cannot be processed well, and do not possess sufficient elasticity and strain recovery to fit to irregular voids. Despite these limitations PLLA and other polyesters are frequently used in tissue engineering applications3. However, our carbonate-based, shape memory, 4D scaffolds are nonacidic upon chain scission and can be 3D printed into highly porous scaffolds that provide structural support on a mechanically suitable substrate where surface chemistry may be modified. This allows production of self-fitting scaffolds, which will take on soft tissue void geometry in a minimally invasive surgery without deforming or applying pressure to the surrounding tissue. Over implantation time, our data indicates that the scaffold will hydrolyze and erode away, primarily through a surface-erosion mechanism, with minimal swelling, allowing for a slow, continuous tissue infiltration without acidosis and mechanical failure that is characteristic of long-term PLLA degradation58. Unlike previous work which requires aspiration or other biological factors, such as bone morphogenic protein-2, the advantage of our 4D printed scaffolds is the relative simplicity of the system, where the printed material only acts to support the ingrowth of native tissues rather than introducing cellular species, which may increase the host’s immune response48. This type of behavior is ideal, as the material will provide mechanical support in the void space until sufficient time has passed (and tissue remodeling cycles have taken place).Tissue engineering is necessary for many soft tissue deformities, including damage to the head, neck, torso and limbs, and breast cancer treatmentUltimately, the unique combination of printability, shape memory, surface-erosion, and excellent biocompatibility, in concert with the ability to tune the thermomechanical properties of the resultant polycarbonates enables them to demonstrate significant potential for application in tissue engineering for various applications. This rationally designed materials allow for synthetic polymeric tissue scaffolds to be easily synthesized and efficiently processed using photopolymerization additive manufacturing. This opens a wide variety of applications in healthcare, manufacturing, and therapeutics geared towards improving patient quality of life. The versatility of these polycarbonate-based systems could enable their applications to be extended to laparoscopic device design or tuned for use in applications such as catheters as well as for use in other minimally invasive surgical techniques for cardiovascular, nephrological, or hepatological surgeries where injectable or compressible stents and patches are required.1H and 125 MHz for 13C) were recorded on a Bruker 400 spectrometer and processed using MestReNova v9.0.1 . Chemical shifts were referenced to residual solvent peaks at δ = 7.26 ppm (1H) and δ = 77.16 ppm (13C) for CDCl3 and δ = 2.50 for (1H) and δ = 39.52 ppm (13C) for d6-DMSO. Size exclusion chromatography (SEC) was performed using an Agilent 1260 Infinity II Multi-Detector GPC/SEC System equipped with both RI and ultraviolet (UV) detectors (λ = 309 nm); PLGel 3 μm (50 × 7.5 mm) guard column and two PLGel 5 μm (300 × 7.5 mm) mixed-C columns with CHCl3 with 5 mM triethylamine as the eluent  were used for separation. A 12-point calibration was developed using poly(methyl methacrylate) standards  and applied for determination of molecular weights and dispersity (ƉM). An Anton Paar rheometer  fitted with a detachable photoillumination system with two parallel plates  was used for rheology studies using RheoCompass software (v1.20.496). Uniaxial tensile testing was performed using a Testometric MCT-350 equipped with a 100 kgf load cell  and manual tension clamps. Dynamic mechanical analysis was performed using a Mettler-Toledo TT-DMA system  fitted with an immersed static water bath with external recirculating heater system, and samples analyzed using Mettler-Toledo STARe v.10.00 software. 3D printing scaffolds and templates were processed using Solidworks 2019  and printed using a custom digital light processing system that has been previously reported59. Micro-computed tomography analysis was performed using a Skyscan 1172 Micro-CT  using CT-analyser software V1.15.4.0 (CTAn)  at an isotropic pixel size of 7-13 μm, a camera exposure time of 500 ms, a rotation step of 0.4°, frame averaging of 5 and medium filtering with a flat field correction. Image reconstruction was performed using a NRecon 1.6.2 .All starting reagents were commercially available, purchased from Sigma–Aldrich (unless otherwise stated), and used without purification. Solvents were of ACS grade or higher. NMR spectra : δ = 0.94 , 1.55 , 3.47 , 3.88–4.05 , 4.23 , 4.52 , 5.21–5.42 , 5.78–5.90  ppm. 13C NMR : δ = 148.59 (C = O), 134.02 (CH), 117.49 (CH2), 72.81 (CH2), 72.42 (CH2), 68.24 (CH2), 35.45 (C), 23.31 (CH2), 7.37 (CH3).Trimethylolpropane allyl ether  was added to a round bottom flask with 200 mL tetrahydrofuran (THF), and cooled to 0 °C for 1 h. Ethyl chloroformate  was added as a single volume to the solution and allowed to again cool to 0 °C for 15 min. Triethylamine  was added dropwise over the course of 1 h, at which time the solution was allowed to slowly return to ambient temperature over a 12 h period. The precipitate was filtered off and the solute concentrated to a slightly yellow oil and dissolved in ethyl acetate. The organic layer was washed twice with 1 M HCl and once with brine, and concentrated to a colorless, slightly viscous oil. The oil was distilled to achieve TMPAC . Characterization matched previously reported materials62. 1H NMR : δ = 6.17 , 5.93 , 4.51 , 4.06 , 3.89–3.83 , 3.61–3.58 , 2.85 , 2.78 , 2.22 , 1.75 , 1.31–1.17 , 0.74 . 13C NMR : δ = 148.52 (C = O), 138.00 (CH), 132.48 (CH), 107.38 (CH), 71.49 (CH2), 70.37 (CH2), 68.91 (CH2), 49.28 (CH2), 43.57 (CH), 42.20 (CH), 31.53 (CH), 28.34 (CH2).Pentaerythritol  was added to a round bottom flask containing deionized water (500 mL), and was subsequently heated to 80 °C with stirring. Once the solids had dissolved, the colorless mixture was then cooled to 20 °C, at which time concentrated HCl  was added, followed by 5-norbornene-2-carboxaldehyde . The mixture was then stirred for 8 h, and the resulting orange precipitate was isolated using vacuum filtration before being recrystallized from hot toluene/isopropyl alcohol (80/20) as white crystals to yield 2-norbornene-5,5-bis(hydroxymethyl)-1,3-dioxane (NHD). NHD  was dissolved in THF (400 mL) in a round bottom flask and cooled to 0 °C, at which point ethyl chloroformate  was added as a single volume, followed by dropwise addition of triethylamine  1 h while maintaining a 0 °C temperature. Upon completing the addition, the reaction was allowed to come to 20 °C and was allowed to stir for 12 h, after which the precipitate was filtered and concentrated to yield white crystals. The white crystals were recrystallized in hot cyclohexane/THF (90/10)  as the NTC monomer. Characterization matched previously reported materials3 and cyclic monomer(s) were added followed by 1,8-diazabicyclo[5.4.0]undec-7-ene (DBU). For PolyTMPAC, TMPAC  was dissolved in 100 mL CHCl3. DBU  and water  were added as a single unit. The resulting solution was stirred for 24 h at 20 °C, after which the DBU was quenched with the addition of Amberlyst A15 H+ acidic resin, precipitated into ice cold hexanes, and was then filtered through a silica plug in ethyl acetate. The solution was concentrated in vacuo to yield a viscous, colorless liquid .Ring opening polymerization of the cyclic monomers was used to obtain oligomers. To an open round bottom flask, CHCl1H NMR : δ = 0.82 , 1.45 , 3.32 , 3.87 , 4.04–4.21 , 5.11–5.32 , 5.79–5.93 , 6.88 . 13C NMR : δ = 155.15 (C = O), 134.67 (CH), 116.73 (CH2), 72.27 (CH2), 69.46 (CH2), 67.78 (CH2), 41.87 (C), 22.58 (CH2), 7.45 (CH3). SEC  Mn: 2.1 kDa, ĐM = 1.2.1H NMR : δ 6.15 , 5.93 , 5.83 , 5.26 , 4.41 , 4.10 , 3.40–4.00 , 3.32 , 2.93 , 2.30 , 1.81 , 1.49 , 1.44–1.15 , 0.89 . 13C NMR : 155.14 (C = O), 134.68 (CH), 132.92 (CH), 116.73 (CH2), 106.93 (CH), 72.27 (CH2), 69.46 (CH2), 67.78 (CH2), 49.26 (CH2), 43.60 (CH), 41.87 (CH), 37.30 (C), 28.38 (C), 22.59 (CH2), 7.46 (CH3). SEC  Mn = 2.0 kg·mol−1, ĐM = 1.311H NMR : δ 6.14 , 5.92 , 5.83 , 5.25 , 4.41 , 4.10 , 3.40–4.00 , 3.32 , 2.93 , 2.30 , 1.81 , 1.49 , 1.44–1.15 , 0.89 . 13C NMR : 154.87 (C = O), 137.71(CH), 134.68 (CH), 132.66 (CH), 116.73 (CH2), 106.93 (CH), 72.27 (CH2), 69.46 (CH2), 67.78 (CH2), 49.26 (CH2), 43.60 (CH), 41.87 (CH), 37.30 (C), 28.38 (C), 22.59 (CH2), 7.46 (CH3). SEC  Mn = 2.5 kg·mol−1, ĐM = 1.37.1H NMR : δ 6.16 , 5.93 , 5.83 , 5.27 , 4.41 , 4.10 , 3.40–4.00 , 3.32 , 2.93 , 2.30 , 1.81 , 1.49 , 1.44–1.15 , 0.89 . 13C NMR : 154.87 (C = O), 137.71(CH), 134.68 (CH), 132.66 (CH), 116.73 (CH2), 106.93 (CH), 72.27 (CH2), 69.46 (CH2), 67.78 (CH2), 49.26 (CH2), 43.60 (CH), 41.87 (CH), 37.30 (C), 28.38 (C), 22.59 (CH2), 7.46 (CH3). SEC  Mn = 2.5 kg·mol−1, ĐM = 1.34.1H NMR : δ 6.14 , 5.92 , 4.42 , 3.40–4.00 , 2.80–2.93 , 2.30 , 1.81 , 1.44–1.15 , 0.83 . 13C NMR : 154.62 (C = O), 137.08 (CH), 132.68 (CH), 106.84 (CH), 72.04 (CH2), 69.49 (CH2), 67.00 (CH2), 49.25 (CH2), 43.61 (CH), 42.22 (CH), 37.30 (C), 28.38 (CH2). SEC  Mn = 2.6 kg·mol−1, ĐM = 1.37.co-hexamethylene diurethane), PolyTMPAC  was dissolved in a round bottom flask containing dry THF at 60 °C under N2, to which hexamethylene diisocyanate (HDI)  was added. The mixture was allowed to stir for 48 h, during which time the viscosity visually increased dramatically. At 48 h, the temperature was increased to 80 °C and allowed to stir for 12 h, at which time the entire solution was added to 50 mL MeOH. The solution was concentrated, washed with 1 M HCl twice and once with saturated brine solution, and collected as a highly viscous, transparent oil . 1H NMR : δ 5.77 , 5.17 , 4.01 , 3.85 , 3.24 , 3.06 , 1.54, 1.14–1.45 , 0.87 . 13C NMR : δ 155.15 (C = O), 134.67 (CH), 116.73 (CH2), 72.27 (CH2), 69.46 (CH2), 67.78 (CH2), 41.87 (C), 40.82 (CH2), 29.88 (CH2), 26.17(CH2), 22.58 (CH2), 7.45 (CH3). SEC  Mn = 3.5 kg·mol−1, ĐM = 1.81.In a representative synthesis of the poly: δ 5.85–5.94 , 5.18–5.30 , 4.86 , 4.53–4.55 , 3.79–3.82 , 2.91-2.92 , 1.67–1.74 , 1.17–1.21 , 0.83–1.05 , 13C NMR : 156.69 (C = O), 155.51 (C = O), 132.94 (CH), 117.72 (CH2), 65.59 (CH2), 54.89 (CH2), 47.05 (CH), 45.62 (CH2), 44.66 (CH2), 41.85 (CH2), 36.41 (C), 35.04 (CH3), 31.86 (C), 29.70(CH3), 27.63(CH3), 23.24(CH3). Mass spectrometry (ESI); m/z = 338.22 (M+). Elem. anal. Calcd for C18H30N2O4: C, 63.88; H, 8.93; N, 8.28%. Found: C, 63.81; H, 8.77; N, 8.24%.Isophorone diisocyanate  was added by canula transfer to a round bottom flask  followed by dry THF (40 mL). Freshly distilled allyl alcohol , stored over molecular sieves, was added dropwise to the solution while stirring at 300 rpm. Upon complete transfer of the allyl alcohol, the reaction was heated to 50 °C and held isothermally for 24 h, at which point residual diisocyanate was quenched with water (at 50 °C). Crude urethane was obtained after dissolving the reaction mixture in ethyl acetate, washing with 1 M HCl (3 washes) and brine (1 wash) and concentrating the product. A viscous clear oil was collected after column chromatography (25:75 EtOAc:Hexane) and concentrated in vacuo to yield a colorless oil . Oligomer, reactive diluents and other diluents were added to a vial, along with the 4 arm tetrathiol (pentaerythritol tetrakis(3-mercaptopropionate) (PETMP)). After mixing, photointiator and photoinhibitor were added. As an example, the polyTMPAC resin consisted of isophorone di reactive diluent , polyTMPAC , 1,3,5-triallyl-1,3,5-triazine-2,4,6-trione as a second reactive diluent species , PETMP , and propylene carbonate as an unreactive diluent  mixed together for 8 h at ambient temperature. To this was added Irgacure 819 , and paprika extract  in a dark room with little ambient light, followed by 1 h of stirring. After homogenization of the resin, the resin was placed in a brown glass container and stored at room temperature in the dark.3 at ambient conditions, exposed to λ = 340–430 nm light for discrete timepoints prior to storage in brown glass vials.Conversion of alkenes in the oligomeric and monomeric reactive components by reactions with the thiols in PETMP with 1% wt photoinitiator and no inhibitor were performed to study crosslinking kinetics. Experiments were performed in 0.5 mL CDClδ), storage moduli, loss moduli, complex viscosity, and film thickness to determine the gelation time(s)63. Resin samples were subjected to oscillatory shearing between two parallel plates (500 µm gap), one made of glass and transparent, at 1 Hz for 50 sec without irradiation, at which time the resins were irradiated with λ = 430–520 nm light and measurements were taken every 0.2 s over the course of 2 min. The inflection points of the moduli plots (storage and loss), and the peak tan δ values, were used to determine the time to gelation of the resin. Sample shrinkage was determined by measuring the gap between the plates at the same sampling rate as the other metrics.Crosslinking kinetics of resin samples were examined by measuring the dampening or phase ratio . Seven samples were run per composition.Printed modified ASTM Type IV dogbones were examined in uniaxial tensile testing at ambient moisture and temperature. Samples were placed in the tension clamps and allowed to vibrationally equilibrate for 10 min, at which point each sample was extended at 5 mm·min−1. Samples were equilibrated at −30 °C for 5 min and were heated to 200 °C at a rate of 10 °C·min−1. The peak ratio between the loss and storage moduli  was defined as the Tg.Rectangular dynamic mechanical analysis (DMA) samples were prepared via 3D printing sample bars (2.0 × 0.5 × 0.2 cm). Samples were analyzed in tension mode using the standard autotension mode, with a testing frequency of 1 Hz, a preload force of 1 N, and a static force of 0.1 N. Thermal sweeps were conducted at 2 °C·min3) were placed in compression and deformed 10 µm, 1 Hz with a preload of 0.1 N at ambient conditions (not submerged) for ~60 sec. At this time, the scaffold was then immersed in the PBS solution and held isothermally at 37 °C for 60 min. Storage moduli and tan δ values were recorded as a function of time to determine the behavior of the polymer during initial submersion/introduction to biologically mimicking conditions. Expansion forces were measured using the same method in creep mode.Relaxation kinetics studies of the printed scaffolds were conducted using submerged samples at 37 °C in phosphate buffered saline (PBS) solution, in compression mode. Printed porous scaffolds  and cooled to −20 °C. Once the sample was isothermal with the cooled chamber, the load was removed, and the sample expansion was monitored as a function of force and displacement of the compression clamp as the sample was heated to 60 °C at 10 °C·min64. Resins were added in 10 mL quantities to the resin tray, allowing for complete and even coverage of the printer optical window on the surface of the printing plate. Porous scaffolds were printed by exposing resins to λ = 405 nm light using a custom-built digital light processing unit and printing parameters were individually determined for each resin composition through optimization of irradiance, irradiation time, resulting film thickness, and semi-quantified feature resolution , and were further optimized in the printing vat as necessary59. The z-stage step transition was set to 100 µm, and each slice was exposed for 6 s, on average. Print resolution was determined through image analysis (Image J) of the theoretical structure, and pore size analysis using microscopy from the printed structure. The final structures are rinsed with acetone to remove residual resin and photoinhibitor, as denoted by colour removal.Scaffolds based upon previously reported geometries were printed from polycarbonate resins using optimized conditions (dependent upon composition)17. For dynamic degradation studies, films were tested using the DMA in tension mode with the autotensioner, loaded with a 0.1 N preload and 10 Hz oscillation. Films were immersed in 5 M NaOH solution at 37 °C for the duration of testing. Samples were tested until failure, with the phase ratio and the storage moduli recorded over the course of the study.Porous scaffolds and nonporous scaffolds were immersed in degradation solution, following previously established protocols for static degradation analysis3 or DMSO-d6. Scaffold swelling ratio was determined by Eq. mi is the original mass of the scaffold (dry) and mf is the mass of the scaffold after swelling (but blotted dry to remove droplets or excess solvent). The crosslink density, and therefore the remaining mass of the material, was determined by Eq. mf is the final scaffold mass (dry) and mi is the original scaffold mass (dry).For in vivo degradation, samples were removed from subcutaneous tissue and sterilized using EtOH. Tissue was removed and scaffolds were extracted with hexane or methanol over a 48 h period, after which the extracted solutions were concentrated down and dissolved in either CDClentclass1pt{minimaA hexagonal void was produced in Solidworks, and the cross-sectional area was varied to produce irregular voids, one which is sharply irregular and the other possessing rounded edges. The voids were printed and used for studying void-filling behavior, using cross-sectional area of the void and the printed scaffold (cube) to determine void filling as a qualitative function of shape.−1, to which was added 5 mL of calcium chloride dihydrate (0.1 mg·mL−1). The two components were mixed until gelation, and 10 mL of H2O was added as the gels were incubated at 37 °C overnight. Gel mechanical properties were matched adipose and glandular tissue using literature protocol42. Gels were cut with an eye-shaped opening, in the same manner as a lumpectomy surgery. Cubic scaffolds were shape fixed at 60% strain and inserted into the opening, where void filling and gel deformation were examined optically using the same cross-sectional analysis described for the “Printed Void Filling” section. The shape fixation behavior of the scaffold was further examined upon removal of the scaffold from the gel, and the shape recovery efficiency compared with the void-filling behavior, as well as the deformation of the alginate. The thin-walled computational molds previously described were then examined using determined loading forces and compared with the deformation found in alginate gels. An interior force of 1 N was initially applied uniformly to the interior (cut) surface of the gel in the same manner as the scaffold would be in contact and expand. The force was then scaled until deformation matched experimental results.Alginate was dissolved in water at a concentration of 10 mg·mLn = 4) were prepared by spin coating a solution of 0.4 wt% polymer in CHCl3 on a glass coverslip (1 min at 1000 rpm). Spin-coated glass coverslips were then sterilized by immersion in a 70% ethanol solution, fully dried before use, and placed into 12-well plates. NOR-10 (murine fibroblasts), Hs 792 (human fibroblasts), IC21 (murine macrophages), and D16 (murine adipocytes) cell lines were purchased from ATCC UK and cultured in DMEM (NOR-10 and Hs 792), RPMI-1640 (IC21), and DMEM/F12 (D16) media supplemented with 10% FBS (20% for NOR-10) and 1% pen/strep, at 37 °C and 5% CO2. 1% L-Alanyl-L-Glutamine was added in DMEM/F12 medium. MC3T3 (murine pre-osteoblasts) were purchased from Public Health England and cultured in Alpha Minimum Essential Medium with ribonucleosides, deoxyribonucleosides, 2 mM L-glutamine and 1 mM sodium pyruvate, but without ascorbic acid, 10% FBS and 1% pen/strep, at 37 °C and 5% CO2.Samples for cell culture studies  and measuring metabolic activity at selected timepoints . Cell proliferation was evaluated by using a PrestoBlue® metabolic assay following the supplier’s instructions. Briefly, after removing the medium, 1 mL of PrestoBlue® solution (10% in cell culture medium) was added to each well, followed by incubation at 37 °C for 1–4 h. 100 µL of solution was taken from each well and placed in triplicate into a 96-well plate. The fluorescence intensity (FI) was detected in a BMG Labtech Fluostar Omega Microplate Reader at wavelengths of 590 nm for excitation and 610 nm for emission.Cell proliferation assays were performed on spin-coated glass slides by seeding the above cell lines (n = 4) at 4000 cells·cm−2. After 72 h, cells were fixed using a 4% paraformaldehyde solution for 10 min, permeabilized using 0.5% Triton X-100 in cytoskeleton stabilization (CS) buffer  8000 MW polyethylene glycol) at 37 °C for 10 min, rinsed thrice for 5 min each in CS buffer, and incubated in 0.1% sodium borohydride in PBS at ambient temperature for 10 min to quench aldehyde autofluorescence. Samples were then blocked in 5% donkey serum for 20 min at 37 °C and incubated overnight at 4 °C with mouse primary anti-vinculin antibody . Samples were then washed three times with 1% donkey serum for 5 min each, and then incubated with Alexa Fluor 647 Phalloidin for cytoskeleton staining (1:200) for 1 h followed by Alexa Fluor® IgG-594 secondary antibody . DAPI was used to stain the cell nuclei. Cells were imaged with a FV3000 Olympus confocal fluorescence microscope using 350, 594, and 633 nm excitation filters and a ×20 or ×40 objectives65.Cells were seeded on spin-coated coverslips  were seeded on top of the scaffolds (n = 3) and incubated at 37 °C, 5% CO2 for 3 h. After this time, 2 mL of culture medium was added and the cells were incubated again at 37 °C, 5% CO2 for the selected timepoints . A live/dead assay (Invitrogen) was performed at each of the selected timepoints. Briefly, scaffolds were washed with PBS (3 × 2 mL) and incubated with a calcein/ethidium homodimer solution at 25 °C for 20 min, following the supplier’s instructions. Scaffolds were then washed with PBS (3 × 2 mL) and placed on a microscope slide for fluorescent imaging. Cells were imaged with a FV3000 Olympus confocal fluorescence microscope using 488 nm and 594 nm excitation filters and a ×4 air objective65. Image J was used for analysis.3D-printed scaffolds were sterilized by immersion in 70% ethanol, dried, and then placed in 24-well plates, and incubated for 24 h in cell culture medium at 37 °C, 5% COad libitum.Experiments were performed in accordance with the European Commission Directive 2010/63/EU  and the United Kingdom Home Office (Scientific Procedures) Act (1986) with project approval from the institutional animal welfare and ethical review body (AWERB). Anaesthesia was induced in adult male (8 weeks old) Sprague Dawley rats (200–300 g) with isofluorane  in pure oxygen (BOC). Animals were placed prone onto a thermocoupled heating pad , and body temperature was maintained at 36.7 °C. The experimental material and control material (PLLA) were implanted over either the spinotrapezius or lateral aspect of the external obliques. Following an incision of ∼3 cm, the skin was separated from the muscle with large forceps, and any excess fat was removed. The implants were tunnelled under the skin and placed in direct contact with the muscle, at sites distal to the incision. The order of the implants was randomized but constrained so that each implant appeared in each location bilaterally at least once. The wounds were sealed with a subcuticular figure of 8 purse string suture with a set-back buried knot using 3-0 vicryl rapide suture (Ethicon). The surgical procedure was performed under the strictest of aseptic conditions with the aid of a nonsterile assistant. Post-surgical analgesia was administered, and rats were placed into clean cages with food and water At 1-month and 2-month timepoints, samples were excised from the subcutaneous tissue and fixed with 4% paraformaldehyde for 24 h. After fixation, samples were washed with increasing percentages of ethanol (70–100%) for 30 min each, washed thrice with xylene, and embedded in paraffin wax blocks for sectioning. Slices (10–30 µm thick) were cut using a Leica Biosystems microtome for histological analysis before being stained using hematoxylin and eosin stains or Masson’s Trichrome staining using protocols available through Sigma–Aldrich. Analysis was performed using light microscopy  and image stitching was performed in ImageJ . Brightfield images were analyzed and qualitatively assessed for general inflammation compared to PLLA control samples. Samples were also analyzed for a number of inflammatory cells utilizing a modified scoring system designed by the International Organization for Standardization (ISO 10993-6 Annex E). Scoring was based on a scale from 0 to 4 .t-test, with probabilities of 0.05 used to assess the probability of differences between compositional behaviors. Specific reproduction numbers and experimental iterations are included in Figure captions and methodology text.Statistical analysis of thermal, thermomechanical, and degradation results was performed using a standard one-way Student’s Further information on research design is available in the Supplementary InformationDescription of Additional Supplementary FilesSupplementary Video 1Supplementary Video 2Supplementary Video 3Reporting Summary