Datasets:
pmid stringlengths 8 8 | pmcid stringlengths 8 11 ⌀ | source stringclasses 2
values | rank int64 1 9.78k | sections unknown | tokens int64 3 46.7k |
|---|---|---|---|---|---|
38667221 | PMC11048303 | pmc | 1 | {
"abstract": "Friction, wear, and the consequent energy dissipation pose significant challenges in systems with moving components, spanning various domains, including nanoelectromechanical systems (NEMS/MEMS) and bio-MEMS (microrobots), hip prostheses (biomaterials), offshore wind and hydro turbines, space vehicles, solar mirrors for photovoltaics, triboelectric generators, etc. Nature-inspired bionic surfaces offer valuable examples of effective texturing strategies, encompassing various geometric and topological approaches tailored to mitigate frictional effects and related functionalities in various scenarios. By employing biomimetic surface modifications, for example, roughness tailoring, multifunctionality of the system can be generated to efficiently reduce friction and wear, enhance load-bearing capacity, improve self-adaptiveness in different environments, improve chemical interactions, facilitate biological interactions, etc. However, the full potential of bioinspired texturing remains untapped due to the limited mechanistic understanding of functional aspects in tribological/biotribological settings. The current review extends to surface engineering and provides a comprehensive and critical assessment of bioinspired texturing that exhibits sustainable synergy between tribology and biology. The successful evolving examples from nature for surface/tribological solutions that can efficiently solve complex tribological problems in both dry and lubricated contact situations are comprehensively discussed. The review encompasses four major wear conditions: sliding, solid-particle erosion, machining or cutting, and impact (energy absorbing). Furthermore, it explores how topographies and their design parameters can provide tailored responses (multifunctionality) under specified tribological conditions. Additionally, an interdisciplinary perspective on the future potential of bioinspired materials and structures with enhanced wear resistance is presented.",
"introduction": "1. Introduction Sustainable development is closely linked to improvement in tribology (low friction, low wear, and enhanced lubrication), which has a significant and beneficial impact on both the environment and society [ 1 ]. Global energy consumption is increasing year on year, putting pressure on both resources and the environment. Notably, the repercussions of friction and wear extend beyond substantial energy and economic loss (approximately 23% of global energy loss and 1–2% of a nation’s GDP) and serve as a major source of CO 2 emissions [ 2 , 3 ]. Crucially, tribological losses (material degradation), comprising high friction and wear inefficiencies, expenses related to part replacement and remanufacturing, and maintenance costs, affect virtually all moving elements, ranging from industrial machinery to the intricate mechanics of natural bone joints and artificial implants integrated into the human body [ 4 ]. Any effort aimed at curbing these losses yields a direct and positive impact, not only conserving energy but also fostering economic stability, and enhancing the well-being of both individuals and the environment. Consequently, it is entirely valid to assert that tribology, with its inherent capability to diminish friction and wear, characteristically contributes to the noble pursuit of a sustainable and human-centric world, and directly or indirectly contributes to Sustainable Development Goals 3, 6, 7, 8, 9, 11, 13, 14, and 15 [ 5 ]. Figure 1 shows the triangle of sustainability, revealing nature, bioinspired materials, and bioinspired tribology as its three corners. In this context, various methods can mitigate friction and its negative effects, such as wear. Those can be classified into (i) the mechanical approach (e.g., well-lubricated systems, novel composites, microstructural enhancements, and heat treatments, etc.) and (ii) the surface topography engineering approach (e.g., adjusting surface roughness or texture to reduce contact area). According to a model proposed by McFarlane and Tabor [ 6 ], friction force is the sum of two components: a mechanical contribution resulting from surface deformation and an adhesive contribution mainly influenced by surface energy. This model suggests that a minimum coefficient of friction (CoF) can be achieved at a critical value of real contact area [ 7 ]. Consequently, controlling surface topography can lead to CoF control. However, traditionally, the focus has been on mechanical approaches, overlooking the potential of engineered interfaces in friction control due to incomplete understanding of surface texturing (optimal choice and design of textures). Additionally, an innovative approach utilizes a combination of both approaches to enhance tribological properties and reduce friction and wear. The synergy may enable simultaneous tribological benefits, such as favorable effects of wear debris, continuous lubricant replenishment, formation of protective tribolayers or tribofilms, etc. [ 8 , 9 ]. Multifunctionality within tribological materials/surfaces represents a paradigm shift in engineering [ 10 ]. These materials not only mitigate wear and friction but also offer measurable benefits for efficient operation in specific applications. For instance, the wear rate of modern hip implants has been reduced by several orders of magnitude [ 11 ]. However, to enhance their longevity and minimize the necessity for revision surgeries, it is equally crucial for biomaterials to positively carry dynamic fluctuating high loads (during human movement), facilitate cell transfer and growth (biological interaction), and ensure seamless tribological integration within the human body (chemical interaction), even leveraging the positive effects of wear debris (e.g., hydroxyapatite-based biomaterials can generate wear debris that stimulates bone growth and integration around the implant, while optimized surface roughness/texture can lead to improved osteoconductivity or lubrication) [ 11 , 12 , 13 , 14 ]. Similarly, in the domain of photovoltaic (PV) solar cells, materials/surfaces are engineered to reduce wear (erosion) while simultaneously improving light absorption and self-cleaning properties [ 15 ]. This innovation can result in remarkable, up to 15%, enhancement in the overall efficiency and lifespan of solar panels [ 16 ]. In hydro-bearings and hydrofoils, meticulously designed surface textures that reduce drag lead to elevated energy conversion rates [ 17 , 18 ]. In certain instances, these materials have demonstrated substantial efficiency gains exceeding 20%, thereby translating into increased renewable energy production. In other cases, such as surgical instruments or soft robotics, it is advisable that the surface demonstrate higher friction or higher adhesion in order to improve the grip strength while accommodating high load transfer (pick up/drop) [ 19 , 20 ]. In space shuttles and satellites, the surfaces are designed to resist impact (wear) from space debris (meteors) while being lightweight, anti-weathering, etc. [ 21 , 22 , 23 ]. Space rover wheels are designed to navigate easily through harsh terrain, requiring enough flexibility while delivering tribological robustness [ 24 ]. Highly durable surfaces that demonstrate water-repellent properties, efficient water harvesting and spreading, or exceptional resistance to oils play a pivotal role in various applications [ 25 , 26 , 27 ]. These surfaces are essential for ensuring rapid lubricant dispersion, enabling rainwater harvesting systems, and powering nanogenerators, among other uses [ 28 ]. For instance, in the context of lubrication, surfaces with superoleophobic characteristics facilitate swift and uniform distribution of lubricants in machinery, reducing friction and wear. Moreover, surfaces engineered for efficient water spreading enable collection and storage of rainwater for various purposes while being resistant to enduring impact, erosion, etc. In summary, the creation of wear-resistant surfaces with specific functionalities, such as water harvesting and superoleophobic properties, has far-reaching implications across a range of applications, from industrial machinery to sustainable water management and energy harvesting technologies [ 29 ]. Frictional anisotropy is a critical tribological feature that spans from the molecular scale to the macroscale, characterized by directional asymmetry in friction response during sliding [ 30 ]. This feature is particularly valuable for microbots designed for friction-based locomotion, offering high maneuverability and precise targeting in confined three-dimensional (3D) spaces. Additionally, it ensures stress compatibility with soft living matter by limiting interface stresses [ 31 ]. As the demand grows for microbots in various applications like robotic assembly, drug delivery, lab-on-chip technology, sensing, microsurgery, and cancer treatment, there is a practical need to construct artificial prototypes [ 32 ]. These prototypes, exploiting geometric features to induce controllable anisotropic motion responses under external oscillating loads, serve as models for understanding and predicting the dynamics of future industrial microactuators [ 33 ]. For instance, in medical microbots navigating complex biological environments, anisotropic friction-based locomotion coupled with physical propulsion holds promise as an enabling technology for their development [ 34 ]. Table 1 shows the essential functional features required in key tribologically challenged fields. Recently, significant attention has been paid to the exploration of materials, surfaces, and architectures inspired by nature [ 25 , 26 , 62 ]. The fact that natural objects have survived the harshest conditions through multifaceted evolution inspires engineers to create or imitate the construction of similar structures. Nature-inspired materials are man-made materials that mimic the structure, properties, or functions of natural materials or living organisms and offer the potential for sustainable synergy between tribology and multifunctionality [ 63 ]. The properties of these materials and surfaces result from complex interplay between the surface structure and the morphology and physical and chemical properties. Moreover, many materials, surfaces, and devices with such designs provide multifunctionality. Various terms such as bioinspiration, biomimicry, biomimetics, nature inspiration, and nature mimicry are commonly used interchangeably by researchers [ 63 ]. Biomimetic materials imitate the evolutionarily developed structural features, leading to adapted architecture, especially in living species [ 64 , 65 , 66 ]. Often, the architectures exhibit a graded structure spanning multiple scales, including macro-, micro-, and nanoscales. For example, bone’s complex porous structure, characterized by complex ligament shapes and variations in density, allows it to achieve superior mechanical (stiffness, energy absorption, stress distribution, etc.) and biological characteristics (e.g., porosity leads to facilitating nutrient exchange and cell proliferation) compared to most man-made materials (biomaterials) [ 65 ]. Friction and adhesion are common in nature. Notable examples involve leveraging nature-inspired surface textures found in animal scales and skins, such as those of sharks (for drag reduction and hydrophobicity), snakes (for erosion resistance), pangolins and turtles (for flexibility combined with erosion/abrasion resistance), and gecko feet (for improved adhesion) [ 24 , 63 , 67 ]. Some species feature adaptive systems that enable changes in color, pattern, or texture for defense, signaling, temperature regulation, or reproduction (also called the ‘chameleon mimetic system’) [ 68 , 69 ]. Nacre, an organic material found in mollusk shells, exhibits remarkable strength and resilience [ 70 ]. Spider silk, renowned for its exceptional mechanical properties and supercontraction abilities, holds great potential for various structural applications [ 71 ]. Examining beetle adhesion systems at the nanoscale has uncovered a diverse array of intricate multiscale architectures serving crucial functions like wing fixation, crawling, mating, and external protection, mainly utilized regarding locomotion for microrobots (anisotropic friction and earthworm-inspired) [ 30 , 34 , 72 ]. These adhesion systems utilize different mechanisms; some rely on liquid secretion (capillary force and lubrication), while others operate through direct interlocking of high-density microfibers or contact of mushroom-shaped hairy structures (van der Waals force) [ 70 ]. A comprehensive analysis of materials (structures) inspired by nature reveals their unique functionality, encompassing various domains of sustainable science [ 63 ]. Bioinspired structures combined with surface topography (or texturing) of these materials demonstrate promising outcomes in customizing friction, wear, and other captivating properties such as antifouling [ 73 , 74 , 75 ], self-lubrication [ 8 , 61 , 76 ], self-adaptation [ 77 , 78 ], hydrophobicity [ 27 ], self-cleaning [ 16 , 25 , 73 ], cell growth [ 12 , 79 ], drug delivery [ 80 , 81 , 82 , 83 , 84 , 85 , 86 ], antibacterial [ 79 , 87 , 88 ], color manipulation [ 68 ], anti-reflection [ 89 ], anisotropic friction in MEMS/NEMS/microrobots/microactuators [ 30 , 70 , 72 ], and enhanced adhesion [ 24 , 25 , 70 , 90 ]. Figure 2 shows various biological organisms (animals/plants) for inspiration regarding creating an efficient multifunctional tribological material. Achieving such complex nature-inspired geometries in materials or on their surface through conventional materials and technologies often proves challenging and expensive [ 92 , 93 , 94 , 95 ]. Historically, progress regarding materials (tribological) has primarily depended on modifying their chemical composition to alter their tribological and mechanical properties. Although this approach has yielded positive results, the journey from discovering new tribomaterial to its commercial availability has typically been time-consuming [ 92 ]. To this end, additive manufacturing (3D printing or laser surface texturing) offers promising prospects, especially in terms of generating complex bioinspired surfaces/materials and also allowing for faster fabrication and up-scaling production [ 94 , 95 , 96 , 97 , 98 , 99 ]. Any bottom-up approach to developing complex, multifunctional, and multiscale mimetics that can provide multiple robust functionalities requires a multidimensional strategy. The research in this multifaceted tribology domain is still at an early stage. The purpose of this review is to provide comprehensive information on the latest advances and future prospects in the field of “nature inspired” or “biologically inspired” tribological materials, with particular emphasis on their design, manufacturing processes, sources of inspiration, as well as friction and wear, performance, and other additional functions. The key focus is to establish a correlation between various tribological scenarios, bioinspired material characteristics, and their resilience to specific environmental conditions, aiming to formulate guidelines for the creation of bioinspired wear-resistant materials and systems. We have attempted to sum up relevant tribological information to a large extent. Nevertheless, due to some missing parameter specifications from the relevant reference publications that do not clearly and unambiguously characterize the physical conditions or tribological stress or the fact that tribological behavior is a system response and mere representation of some value such as force, rotational frequencies, etc., it is not in itself (alone) suitable for characterizing the stress level of the tribological system in question. Also, mixed or incorrect use of data/units for mass and forces should be avoided or carefully analyzed. For this, we recommend that readers defer to the relevant cited source(s). Figure 3 illustrates the research methodology framework for the current work."
} | 4,020 |
36071037 | PMC9452534 | pmc | 2 | {
"abstract": "Light-induced microbial electron transfer has potential for efficient production of value-added chemicals, biofuels and biodegradable materials owing to diversified metabolic pathways. However, most microbes lack photoactive proteins and require synthetic photosensitizers that suffer from photocorrosion, photodegradation, cytotoxicity, and generation of photoexcited radicals that are harmful to cells, thus severely limiting the catalytic performance. Therefore, there is a pressing need for biocompatible photoconductive materials for efficient electronic interface between microbes and electrodes. Here we show that living biofilms of Geobacter sulfurreducens use nanowires of cytochrome OmcS as intrinsic photoconductors. Photoconductive atomic force microscopy shows up to 100-fold increase in photocurrent in purified individual nanowires. Photocurrents respond rapidly (<100 ms) to the excitation and persist reversibly for hours. Femtosecond transient absorption spectroscopy and quantum dynamics simulations reveal ultrafast (~200 fs) electron transfer between nanowire hemes upon photoexcitation, enhancing carrier density and mobility. Our work reveals a new class of natural photoconductors for whole-cell catalysis.",
"introduction": "Introduction Living cells have been incorporated with quantum dots and nanostructures for fluorescent labelling and drug delivery for over two decades 1 . However, light-absorbing nanostructures have not been used to drive catalytic reactions inside of cells due to lack of biocompatibility and high cytotoxicity of foreign materials, such as photosensitizers, inside the cell which often limits the operational efficiency 1 . Furthermore, inherent defects in synthetic photosensitizers cause several problems such as photocorrosion, photodegradation and the generation of photoexcited radicals, which results in low stability, irreproducibility and lack of sustainability of biohybrid materials 2 . Some bacteria produce light-absorbing centers but suffer from low electron transfer efficiency and a lack of durability 1 . Natural electron transfer proteins such as azurins, myoglobin and c-type cytochromes do not show photoconductivity 3 , 4 due to picosecond carrier lifetimes of the heme iron which typically inhibits any charge separation 5 . Covalently linking artificial photosensitizers to these proteins yields low electron transfer rate on the 10 ns timescale or slower, greatly limiting their applications 6 . Furthermore, it is not feasible to use longer excited state lifetimes, such as electron injection from the triplet states due to rapid degradation caused by reactive oxygen species produced in these processes. Therefore, there is an urgent need to develop novel biomaterials capable of ultrafast primary electron transfer to achieve efficient charge separation, followed by sequential secondary electron transfer for long-lived charge separation and charge accumulation 6 . To evaluate the use of engineered living materials as living photoconductors, we chose the electroactive soil organism Geobacter sulfurreducens because it has evolved the ability to export electrons, derived from metabolism, to extracellular acceptors such as metal oxides and electrodes in a process called extracellular electron transfer (EET) 7 , 8 . Bacteria establish direct electrical contact to electron acceptors via micrometer-long, polymerized cytochrome nanowires, called OmcS, which eliminates the need for diffusive redox mediators 7 , 8 (Fig. 1b ). Hemes in the OmcS nanowire form a parallel, slipped-stacked pair with each pair perpendicular (T-stacked) to the next pair, forming a continuous chain over the entire micrometer length of the nanowire 7 (Fig. 1d ). The minimum edge-to-edge distances is 3.4–4.1 Å between the parallel-stacked hemes and 5.4–6.1 Å between the T-stacked pairs. Fig. 1 Living photoconductors. a Measurement schematic. Biofilms are grown on transparent fluorine-doped tin oxide (FTO) electrodes. b Transmission electron microscopy of CL-1 cells producing OmcS nanowires. Scale bar, 200 nm. c AFM height image of a single OmcS nanowire on mica (left) and respective height profile (right) shown where the red line is indicated. Scale bar 50 nm. d Hemes in OmcS stack seamlessly over the entire micrometre-length of nanowires. Edge-to-edge distances are in Å. e UV-Visible spectroscopy of biofilm on FTO electrode with the excitation wavelength of 408 nm marked as a purple triangle. f Current voltage response of biofilm with the laser on and off. Percentage increase in conductance value represents mean ± standard deviation (S.D). of two biological replicates. Source data are provided as a Source Data file. Owing to this evolutionarily optimized OmcS nanowire structure with seamless stacking of hemes, G. sulfurreducens can transfer electrons over distances of one hundred times their size by forming more than 100 µm-thick highly-conductive nanowire networks in biofilms 9 , 10 , which enables G. sulfurreducens cells to generate the highest current density in bioelectrochemical systems 11 . Owing to large electron storage capacity, cytochromes also confer high supercapacitance to biofilms with low self-discharge and reversible charge/discharge 12 . Moreover, a network of purified nanowires can transfer electrons over distances of 10,000-times the size of a cell 9 . Therefore, G. sulfurreducens serves as an ideal model system for electrocatalysis, metal corrosion and production of fuels 13 , 14 . It was previously thought that conductive filaments on the surface of G. sulfurreducens are pili 15 and a network of pili confers conductivity to G. sulfurreducens biofilms 10 , 13 , 16 . However, structural, functional and subcellular localization studies revealed that nanowires on bacterial surface are composed of cytochromes 7 , 8 whereas pili remain inside the cell during EET and are required for the secretion of cytochrome nanowires to the bacterial surface 17 , 18 . Nanowires could be widespread and their photophysical properties could be physiologically important because many Geobacter -like metal-reducing bacteria form highly conductive biofilms 19 , 20 and are widely distributed at the surface of earth in shallow sediments that contain abundant sunlight and metal oxides 21 – 23 . The sediments are capable of transporting electrons over centimeters 24 and can convert incident light into electricity 25 . Illuminating visible light on G. sulfurreducens cells has been shown to improve their catalytic performance such as increases in metabolic electron transfer to metal oxides 26 or other semiconducting materials by over 8-fold compared to that observed under dark conditions 21 . Furthermore, light-induced bacterial electron transfer correlated well with the rates of microbial respiration and substrate consumption 26 . However, the underlying molecular and physical mechanism for this increased photocatalytic performance has remained unclear. In addition to light-induced whole-cell catalysis 21 , 26 , artificially expressing cytochrome OmcS in photosynthetic cyanobacteria, increased catalytic performance in diverse processes such as an increase in photocurrent by 9-fold 27 , increase in nitrogen fixation by 13-fold 28 , and improved photosynthesis due to 60% increase in biomass 29 compared to the wild-type cyanobacteria. These studies highlight the vital role of OmcS in light-driven biocatalysis. However, intrinsic photophysical properties of OmcS, which could account for these catalytic improvements, have not been investigated. Out of 111 cytochromes in G. sulfurreducens , OmcS is the only nanowire-forming cytochrome essential for EET to Fe(III) oxides abundant in subsurface 14 . Indeed, cytochromes abundant in subsurface during uranium bioremediation function similar to OmcS 30 . OmcS is also important for EET to electrodes during initial stages of biofilm growth 14 . OmcS is also required for interspecies electron transfer in Geobacter cocultures to conduct “electric syntrophy” 13 , 31 , 32 . This interspecies electron transfer via naturally conductive microbial consortia is important in diverse methanogenic and methane-consuming environments that affect global climate 33 – 35 . Photosynthetic bacterial species have also been shown to perform electric syntrophy with light-driven conversion of CO 2 to value-added chemical commodities 2 . However, the components and pathways responsible for such light-driven biocatalysis have not been identified and potential for photoactivity beyond photosynthetic microorganisms remains largely unknown. We hypothesized that cytochrome nanowires in the biofilms could be photoactive, enabling efficient electronic interface between microbes and electrodes. Here we show that living biofilms of Geobacter sulfurreducens use nanowires of cytochrome OmcS as intrinsic photoconductors. Surprisingly, nanowires show photoconductivity with ultrafast, sub-picosecond heme-to-heme electron transfer which could explain their influence on photocatalytic performance mentioned above. These rates are among the highest for excited-state electron transfer in biology 36 .",
"discussion": "Results and Discussion Photoconductivity in living biofilms made up of OmcS nanowire network To determine the role of OmcS nanowires in light-induced electron transfer, we used the genetically engineered G . sulfurreducens strain CL-1 because it overexpresses OmcS nanowires (Fig. 1b–d ) and forms highly conductive and cohesive biofilms that can be easily transferred to multiple surfaces 37 (Fig. 1a ). Upon laser photoexcitation (λ = 408 nm) which is specific to the Soret band of c-type hemes 4 , biofilm conductance remained ohmic and increased by 72 ± 21% (Fig. 1e, f ). These studies show that living G . sulfurreducens biofilms can serve as intrinsic photoconductors. As biofilm conductivity determines the bacterial rate of EET 11 , our results could explain the increased photocatalytic performance by G. sulfurreducens 21 , 26 . Rapid photoconductivity in purified nanowires reversible for hours To determine the origin of photoconductivity in biofilms, we purified nanowires from the CL-1 strain (Fig. 2a ). The ultraviolet-visible (UV-Vis) absorbance spectrum of nanowires showed a strong Soret band at 410 nm for air-oxidized nanowires (Fig. 2b ). Nanowires were fully oxidized under these conditions because addition of oxidant (ferricyanide) did not change the spectrum (Supplementary Fig. 1a ). We placed the nanowires on interdigitated gold electrodes and illuminated from the top (Laser Power = 100 mW/cm 2 ). Photoconductance of nanowire network initially increased more than 6-fold (Fig. 2c ), but the extent of conductance increase decreased over time, likely due to laser damage. Nanowires responded faster than 100 ms (Fig. 2c inset). The photoresponse persisted for hours but decreased over time (Fig. 2c ). Both the dark current and photocurrent were proportional to an applied voltage ranging from –0.2 to +0.2 V (Fig. 2d ), indicating an ohmic conduction behavior of nanowires similar to biofilms. Remarkably, nanowire networks, with and without laser excitation, showed a linear current-voltage response with an average conductance increase of 230 ± 28 % ( n = 7), which is higher than common perovskites 38 , 39 and porphyrin nanowires 40 (Fig. 2d, e ). Fig. 2 High photoconductivity in purified protein nanowires. a Heme staining gel of nanowires showing a single band of OmcS. b UV-Vis spectrum of oxidized (green) and reduced (red) nanowires. c Photocurrent response of nanowire network at 200 mV with the current decay of the off-state subtracted. Inset : Fast (<100 ms) photoresponse of nanowires. Axes are same as in Fig. 2c. d Current-voltage response of nanowire network and cytochrome c for comparison e Comparison of conductance of nanowire network with laser on or off. Values represent mean ± standard error of the mean (S.E.M) with individual data points shown as grey dots (n = 7 independent experiments). ** indicates p value = 0.003 using a paired two tail t -test. f Schematic of pc-AFM of individual nanowires. g Current-voltage response of an individual nanowire with a linear fit shown by a purple dashed line. h Comparison of conductance increase upon photoexcitation in individual nanowires. Values represent mean of all current-voltage curves measured on individual nanowires (number of curves ranges from 10 to 120 Supplemental Table 2 ). i Comparison of average conductance of individual nanowires with laser on or off. Values represent mean ± S.E.M. with individual data points shown as grey dots ( n = 15 independent experiments). ** indicates p value = 0.007 using a paired two tail t -test. Source data are provided as a Source Data file. Multiple control experiments confirmed that the observed photoconductivity is an intrinsic property of nanowires, owing to their polymerized cytochrome architecture. For example, the monomeric horse-heart cytochrome-c showed very low dark current and photocurrent as expected 3 , 4 when measured under identical conditions (Fig. 2d ). Upon addition of a chemical reductant sodium dithionite, the Soret band for reduced nanowires red-shifted to 420 nm as expected 4 (Fig. 2b ). These chemically reduced nanowires (λ Soret = 420 nm) did not show significant photoconductance upon excitation at λ = 405 nm, confirming that photoreduction of oxidized hemes are necessary for photoconductivity in nanowires at this excitation (Supplementary Fig. 1b ). Switching the electrode material from gold to tungsten also retained photoconductivity, confirming that the measured response is not an artifact of the electrode material (Supplementary Fig. 2 ). The ratio of laser-on/ laser-off (on/off) current of nanowires increased with increasing laser power, further demonstrating that the measured photoconductivity is solely due to laser excitation (Supplementary Fig. 3 ). All these experiments together confirm that the nanowires show intrinsic photoconductivity which can account for observed photoconductivity in living biofilms. The difference in photoconductivity between biofilms and purified nanowires is likely due to non-conductive materials such as cells and polysaccharides present in the biofilms. Individual nanowires show up to 100-fold photoconductivity increase To quantify the photoresponse of individual nanowires, we used photoconductive atomic force microscopy (pc-AFM) 41 (λ = 405 nm, Initial Laser Power = 3.20 kW/cm 2 , Fig. 2f ). Individual nanowires showed up to 100-fold increase in conductance upon photoexcitation (Fig. 2h–i , Supplementary Table 2 ). The differences in photoconductance are likely due to variation in the laser power caused due to experimental setup (see methods and Supplementary Fig. 11 for details). The difference in the photoconductivity between individual nanowires and nanowire network is likely due to inter-nanowire as well as nanowire-electrode contact resistance. Notably, the observed 10 to 100-fold increase in conductance for protein nanowires at relatively low bias (< 0.5 V) is substantially greater than that of synthetic porphyrins 42 that show only up to a 5-fold increase at very high bias of 12 V. These experiments on individual nanowires confirm that the observed photoconductivity response in networks of nanowires is due to nanowires alone and not because of an artifact of the measurement setup. Furthermore, the observed photoconductivity is not due to heating effects because all pc-AFM experiments were performed in a temperature-controlled environment thus inhibiting any substantial increase in temperature. Furthermore, the linearity and stability of our IV curves indicate that measured conductivity increase is not due to heating (Fig. 2g ). In addition, the conductivity of OmcS nanowires decreases upon heating 43 whereas we observed up to 100-fold increase in conductivity upon photoexcitation. fs-TA revels sub-picosecond charge separation in nanowire To understand the mechanism of photoconductivity in protein nanowires, we performed femtosecond transient absorption (fs-TA) spectroscopy by determining the electron dynamics upon photoexcitation on an ultrafast (~100 fs) time scale 5 , 44 (Fig. 3a ). The fs-TA tracks the UV-Vis spectral changes by changing the time delay Δτ between the femtosecond laser pump and the probe pulses and recording a differential absorbance spectrum (ΔA) at each time delay 44 (Fig. 3a ). This difference spectrum contains information on the dynamic processes occurring in the system such as excited state energy migration, electron or proton transfer processes and isomerization 44 . In contrast to the above studies of photoexcitation in the Soret band (Figs. 1 , 2 ), we performed fs-TA using excitation in the Q-band (λ = 545 nm) to avoid thermal damage, and to monitor changes in the region of the strongest absorption bands 44 . It is important to note that Soret and Q-band transitions arise from the same ground state making Q-band excitation a suitable proxy to monitor these processes 44 . Photoexcitation with λ = 530, 545 and 400 nm yielded similar dynamics, demonstrating a wide spectral range for photoconductivity (Supplementary Figs. 4 , 5 ). Neither buffer alone nor the blank substrate showed any response, measurements in solid and liquid state are similar, and the ET dynamics were independent of the laser intensity and power (Supplementary Figs. 6 , 7 , 8 ) indicating that observed dynamics are due to nanowires and not an artifact of the environment or the substrate. Fig. 3 Ultrafast (<100 fs) charge transfer between hemes in nanowires revealed by femtosecond transient absorption spectroscopy (fs-TA). a Schematic of fs-TA. A pump beam (λ = 545 nm) excites a nanowire sample and is followed by a probe beam after a time delay. The differential absorption between the initial and time-delayed spectra is detected and reported as optical density. b Averaged transient absorption data of nanowires ( n = 6 independent experiments) where colours represent the milli optical density (mOD). c Normalized change in differential absorption with wavelength at different delay times. Key wavelengths are marked as λ = 410 nm (green), λ = 424 nm (red), and λ = 367 nm (blue). d The experimental (solid) and simulated (dashed) spectra of oxidized, reduced, and singlet doubly-oxidized nanowires. Wavelength markers are same as in Fig. 3c. e Normalized change in differential absorption over delay time at key wavelengths. Time-markers are shown in the same colour as time traces in Fig. 3c. Traces in c and e represent mean of n = 6 independent experiments. Source data are provided as a Source Data file. Upon photoexcitation of protein nanowires, electrons are promoted from the ground state to the excited state, decreasing the ground state population. This decrease caused a negative signal in ΔA at 410 nm known as a ground state bleach 44 (Fig. 3b, c ). In addition, we observed a positive ΔA which is indicative of excited-state absorption at λ = 367 nm and λ = 424 nm after Δτ = 0.1 ps and 2 ps, respectively (Fig. 3c, d ). These absorptions are absent in the native, air-oxidized, unexcited nanowires (Fig. 3d ), indicating that the photoexcitation is causing these absorptions. In particular, the absorption at λ = 424 nm agrees well with the absorption of chemically reduced nanowires (Fig. 3d ), suggesting that upon photoexcitation, excited-state electron transfer is reducing the hemes in the nanowires and thus photoreduction contributes to nanowire photoconductivity. To understand the origin of different transient oxidation states, we determined the kinetics at the key wavelengths mentioned above using a sequential model 44 that yielded the first excitation timescale of 19 ± 23 fs (see methods). This timescale is faster than the instrument response function (100 ± 50 fs) and can, therefore, be treated as an instantaneous excitation on the timescale of the measurement (Fig. 4a ). Following this excitation, the charges were transferred between hemes with a decay time of 212 ± 27 fs. The corresponding spectra are a superposition of a ground state bleach and the appearance of a new feature around 367 nm, which can be attributed to the doubly oxidized hemes as per the spectral simulations (Fig. 3d ). Based on simulations (see below, Fig. 4d ), we conclude that the ultrafast charge transfer also results in the formation of a reduced heme in its excited state, which is spectroscopically dark. We also found a second decay with a time constant of 1.0 ± 0.1 ps that can be attributed to relaxation of the excited reduced heme that increases in absorption at the reduced Soret at λ = 424 nm. In addition, we found a third decay time of 7.9 ± 0.3 ps that can be attributed to the recombination to their initial state, including charge transfers back to the singly oxidized heme ground states. Fig. 4 Model for origin of photoconductivity in protein nanowires. a Simplified energy level diagram for hemes depicting the changes that occur upon photoexcitation in transient absorption and their respective decay times. b The dark current in the ground state arises due to propagation of a reduced state created by electron injection from the electrode. c The photocurrent is due to the laser excitation initiating an ultrafast charge transfer between hemes, creating newly reduced (red) and double oxidized hemes (blue). The photoreduction provides additional charge carriers and larger driving force for charge transfer, which therefore increases the current under bias. d Quantum dynamics simulations of ultrafast charge transfer between hemes in protein nanowires, forming a doubly oxidized heme and an excited state of a reduced heme. Computations suggest excited state electron transfer between hemes We further compared our experimental UV-Vis spectra of nanowires with time-dependent density functional theory calculations of hemes in the nanowire (Fig. 3d ). The maximum of the computed Soret band in the reduced heme (λ = 420 nm) is shifted by 9.5 nm to the red of the band maximum for the oxidized heme, in good agreement with the 10.5 nm shift observed experimentally (Fig. 3d ). These computational analyses further suggest that photoexcitation causes reduction of hemes in the nanowires. Our finding is consistent with prior studies of photoreduction of monomeric cytochromes mediated by the light-induced excited state of hemes even in the absence of external electron donors 45 , 46 . To evaluate the transient kinetics data obtained using the sequential model fitted to experimental data, we performed quantum dynamics simulations at the Extended Hückel level of theory 47 , 48 . We simulated the propagation of an electron wavepacket in the excited state from hemes in the nanowires, in a slip-stacked as well as in T-stacked orientation 7 . Our simulations suggest a ~100 fs timescale for photo-induced charge transfer between the slip-stacked pair of hemes (Fig. 4d ). This timescale agrees with the experimentally determined timescale for the excited state charge transfer (212 ± 27 fs). The survival probability for electron transfer in a slipped-stack heme pair remains low (<60%) for most energy levels, indicating a high probability for electron transfer to a nearby heme within 100 fs (Supplementary Fig. 9 ). Thus, the timescale for electron transfer between slipped-stack heme pairs remains similar for most energy levels. Hemes are likely electron source for observed photoreduction As no external electron donor was added, our results suggest that the additional electrons that reduce the heme are intrinsic to the nanowire itself. We further analyzed the possibility that surrounding protein causes the observed photoreduction of OmcS hemes. Several aromatic amino acids, including tryptophan and tyrosine, are within 5 Å of the hemes in the OmcS. Although excitation of either tryptophan or tyrosine is not possible at the wavelengths used in this study 45 , 46 , we considered a possibility that electron transfer can quench a photoexcited heme in a manner similar to flavins in a cryptochrome 49 . This quenching would reduce a heme and leave behind an amino acid radical. The most likely amino acid candidate for radical formation is tryptophan because its radicals have absorbance which would explain the 367 nm species 50 . While the formation of such radicals is possible, the signal strength in fs-TA measurements is determined by the molar extinction coefficients (ε) of the (transient) species. The molar extinction coefficient of the Soret band for OmcS is approximately 100 times larger than those of tryptophan radicals 50 , 51 . The ground state bleach represents all the photoexcited hemes in the OmcS nanowires and the species corresponding to λ = 367 nm and 424 nm have differential absorptions of ~20 and 10% of the total magnitude, respectively (Fig. 3e ). Therefore, the number of tryptophan radicals created from electron transfer needs to be larger than the number of excited hemes in OmcS if the radical species at λ = 367 nm arises from tryptophan. Such a possibility seems unlikely because only one radical can be created for every quenched excited heme. Thus, the observed spectra cannot be accounted for by amino acid radicals. We also evaluated the possibilities of other electron sources causing photoreduction. We found that multi-photon processes are absent in our experiments because the ET dynamics were independent of the laser intensity and power (Supplementary Fig. 8 ). The magnitude of photocurrent is also linear with increased power (Supplementary Fig. 3 ). Redox impurities also did not contribute to the measured spectra because of identical dynamics in solution and in solid-state (Supplementary Fig. 7 ). Photodegradation also did not change the electron transfer dynamics, only the magnitude of spectra by <10% over two hours. We therefore considered an alternative possibility that parallel-stacked hemes can serve as an electron donor and acceptor pair (Fig. 4 ). We hypothesized that the excited state charge transfer is occurring between two neighboring hemes with only one of the hemes being in the excited state. Such charge transfer would result in the appearance of a reduced heme and leave behind a doubly oxidized heme (Fig. 4 ). The computed UV-Vis spectrum of a doubly oxidized heme indeed showed an absorption maximum at λ = 365 nm which agrees with the experimentally observed species at λ = 367 nm. Our computed spectrum of a doubly oxidized heme thus recaptures the blue shift observed in the transient absorption experiment (Supplementary Fig. 10 ). The qualitative agreement between the computed and experimental spectra is independent of the spin state of doubly-oxidized species such as the singlet and triplet state. To identify the nature of doubly oxidized species, we performed an analysis of atomic spin populations. We found that the change in the spin populations occurs only on the ligands and not in the iron center. Therefore, our analysis suggests that doubly oxidized species are Fe 3+ + porphyrin radical which agrees with the observed spectra at 367 nm. These analyses further suggest that the doubly oxidized species are not Fe 4+ due to lack of change of spin density on the iron center upon additional oxidation of the heme in the Fe 3+ state (Supplementary Fig. 10 and Supplementary Table 1 ). To further evaluate the thermodynamic feasibility of radical heme species, we used the Rehm-Weller cycle. This analysis requires four energetic terms: (1) energy required to form radical heme species (based on iron-porphyrin systems 52 ) (1.7 V), (2) the ground state redox potential of OmcS (−212 mV) 51 , (3) the photon energy used to excite OmcS nanowires (λ = 545 nm = 2.3 eV), and (4) the vibrational energy difference between the ground and excited states, called the Coulomb stabilization energy associated with the intermediate radical ion pair 53 ( ω p ) ~60 meV. Therefore, the energetics of this process would be ΔG et = [1.7 eV–(−0.212 eV) + 0.06 eV]−2.3 eV = −0.4 eV. Thus, ΔG et < 0 for the formation of the radical heme species, making them energetically feasible. Our analysis is a lower estimate for the net energy available for the formation of the radical species. Therefore, in combination with our simulated analysis, our studies suggest that doubly oxidized species are Fe 3+ + porphyrin radical and nanowires are photoreduced by ultrafast light-induced heme-to-heme charge transfer. Proposed mechanism for ultrafast photoconductivity in OmcS nanowire Based on above results, we propose the following model for the origin of photoconductivity in OmcS nanowires (Fig. 4 ). This model is focused on the singlet states and not triplet states because these states are spectroscopically dark, and would be less pronounced due to their lower energies. As these nanowires transport charges through seamless stacking of hemes (Fig. 1c ), our prior experiments have shown that they can be treated as redox conductors, with the long-range charge transfer governed by a theoretically-predicted hopping mechanism with negligible carrier loss over micrometers 54 . All hemes in the nanowires are initially oxidized and in their ground state as confirmed by UV-Vis spectroscopy (Fig. 2b ). Upon applying a bias, electrons are injected from the electrode into the nanowire, creating a reduced state that travels through the nanowire (Fig. 4b ). The photoexcitation triggers an ultrafast charge transfer resulting in an additional reduced state that persists for picosecond timescale, without any applied bias, far away from the electrode (Fig. 4c ). This newly formed reduced state will have a mobility similar to the electrode-injected state as they both are present in the same nanowire with identical structure. Therefore, upon photoexcitation, the density of reduced states is increased, thus increasing the carrier density of the OmcS to generate photoconductivity in nanowires. The photoreduction observed in our fs-TA is consistent with this model. In addition to the higher carrier density due to photogenerated electrons, it is likely that the mobility of electrons increases upon photoexcitation due to increased driving force for charge transfer in the excited state of hemes 43 . Upon photoexcitation an electron is promoted from the ground state to an excited state. The ultrafast charge transfer between neighboring hemes creates a reduced-state heme in the excited state and a doubly oxidized heme (Fig. 4c, d ). The reduced-state heme can then relax from the excited to the ground state. Upon photoexcitation, the uniformly oxidized nanowire is thus partially reduced and partially double oxidized (Fig. 4c ). The generated doubly oxidized heme will alter the redox energies of the heme chain, with a more positive redox potential. We have previously found that the redox potential of OmcS hemes becomes substantially positive upon oxidation 43 . The OmcS nanowires transport charges via a hopping mechanism 54 —a process in which a charge (electron or hole) temporarily resides at a heme, changing its redox state. The driving force for charge transfer depends on the redox energies of the electron donating and accepting hemes. Therefore, the charge transfer rate is directly related to the mobility. For the fully oxidized (non-excited) state, this process initiates at the electrode surface where injected electrons hop to nanowire redox sites, creating locally reduced hemes. For the photoexcited state, this process is enhanced because transferring an electron to the double-oxidized species, and removing an electron from a reduced heme, are significantly more favorable in the illuminated nanowire than for the oxidized nanowire in the dark. The increased likelihood for charge transfer upon photoexcitation will then result in increased mobility. Furthermore, the initial ultrafast charge transfer between hemes increases the lifetime of the photogenerated state. Both the generation of a “new” mobile charge and the increase in its mobility will contribute to the observed increase in conductivity upon photoexcitation. In summary, we demonstrate, for the first-time, significant photoconductivity in a living system due to ultrafast light-induced charge transfer within protein nanowires. The surprising origin of photoconductivity in these natural systems lies in the higher carrier density and mobility upon photoexcitation. Although ultrafast electron transfer can occur in monomeric cytochromes, it typically requires incorporated dyes as photosensitizers and sacrificial electron donors 36 which can be toxic to cells 1 . In contrast, we find that the protein nanowires intrinsically exhibit robust and ultrafast charge transfer without any need for such site-selective labeling. Our studies thus establish OmcS nanowires as photoconductors intrinsic to cells with capability of ultrafast electron transfer, thus eliminating the need for foreign materials such as molecular dyes or inorganic nanoparticles that limit the catalytic performance 1 . Furthermore, our studies show that sub-ps charge transfer is possible in natural proteins in an excited state. Prior ultrafast electron transfer studies have reported the ground state rates of 15–90 ps in the closest-stacked hemes 36 . This difference is likely because excited-state rates are known to be faster due to higher energy and larger orbital delocalization compared to the ground-state rates 49 . Although many bacterial EET studies remain focused on electrons, protons play a very important role, not only in bacterial energy generation, but also in the electronic conductivity of proteins 55 . For example, through measurements of the intrinsic electron transfer rate, we previously found that both the energetics of a glutamine (proton acceptor) and its proximity to a neighboring tyrosine (proton donor), regulate the hole transport over micrometers in amyloids through a proton rocking mechanism 56 . Therefore, it is very important to couple electron/proton transfer to accelerate EET and for the development of electronically conductive protein-based biomaterials. The high surface area of these nanowires, combined with their biocompatibility and lack of toxicity, make them attractive candidates for an emerging field of light-driven whole-cell bioelectrocatalysis for a wide range of applications such as water splitting, chemical sensing and CO 2 fixation and production of chemicals, fuels and materials 57 . Our studies may also help establish the efficient and stable production of liquid fuels from sunlight using a liquid sunlight approach 5 . Future studies on nanowires with different heme stacking and protein environment 8 or substituting the metals from iron to zinc 58 or tin 59 could vary the interactions between the heme cofactors to alter the electronic and photophysical properties of nanowires for tuneable functionality 57 ."
} | 8,779 |
31057985 | PMC6497424 | pmc | 3 | {
"abstract": "The\nbiological production of FDCA is of considerable value as a potential\nreplacement for petrochemical-derived monomers such as terephthalate,\nused in polyethylene terephthalate (PET) plastics. HmfF belongs to\nan uncharacterized branch of the prenylated flavin (prFMN) dependent\nUbiD family of reversible (de)carboxylases and is proposed to convert\n2,5-furandicarboxylic acid (FDCA) to furoic acid in vivo. We present\na detailed characterization of HmfF and demonstrate that HmfF can\ncatalyze furoic acid carboxylation at elevated CO 2 levels\nin vitro. We report the crystal structure of a thermophilic HmfF from Pelotomaculum thermopropionicum , revealing that the\nactive site located above the prFMN cofactor contains a furoic acid/FDCA\nbinding site composed of residues H296-R304-R331 specific to the HmfF\nbranch of UbiD enzymes. Variants of the latter are compromised in\nactivity, while H296N alters the substrate preference to pyrrole compounds.\nSolution studies and crystal structure determination of an engineered\ndimeric form of the enzyme revealed an unexpected key role for a UbiD\nfamily wide conserved Leu residue in activity. The structural insights\ninto substrate and cofactor binding provide a template for further\nexploitation of HmfF in the production of FDCA plastic precursors\nand improve our understanding of catalysis by members of the UbiD\nenzyme family.",
"discussion": "Results and Discussion Initial Identification,\nExpression, and Characterization of Thermostable FDCA (De)carboxylases It has previously been reported that the thermophilic bacterium Geobacillus kaustophilus HTA426 is capable of degrading\nfurfural. 17 A BLAST search of the G. kaustophilus genome 18 using the C. basilensis Hmf gene\ncluster suggested the presence of a similar Hmf gene cluster located\non plasmid pHTA426. Although there is no mention in the literature\nregarding the ability of G. kaustophilus to degrade HMF, a C. basilensis HmfF\nhomologue (WP_011229502) could be located on pHTA426, possessing 51%\nsequence identity and located adjacent to a HmfG/UbiX homologue. Active\nrecombinant G. kaustophilus HmfF was\nsuccessfully produced in E. coli when\nit was coexpressed with E. coli UbiX\n( Figure S2 ). However, while soluble recombinant G. kaustophilus HmfF could be produced, the protein\nhad a tendency to aggregate, hampering crystallogenesis and other\nbiophysical studies. Other thermophilic HmfF homologues were screened,\nwith the P. thermopropionicum HmfF\nenzyme being the most promising in terms of protein expression levels\nand stability. The purified recombinant HmfF enzymes (from both P. thermopropionicum and G. kaustophilus ) were capable of decarboxylating 2,5-furandicarboxylic acid to furoic\nacid in vitro ( Figure S2b ) but could not\nfurther decarboxylate furoic acid to furan. Expression and Detailed\nCharacterization of P. thermopropionicum HmfF Purified PtHmfF expressed in absence of E. coli UbiX coexpresssion was pale yellow and possessed\na UV–vis spectrum consistent with oxidized FMN binding. In\ncontrast, when it was coexpressed with UbiX, the purified recombinant\nprotein was pale pink, possessing a complex UV–vis spectrum\nwith three main features in addition to the protein peak at 280 nm\n( Figure 2 A). These\ninclude a feature at 390 nm, similar to that observed previously for\nthe model system A. niger Fdc1, 11 a peak at 450 nm (likely corresponding to the\npresence of a subpopulation bound to oxidized FMN rather than prFMN),\nand a broad peak centered around 550 nm. Similar spectral features\nat 550 nm were previously identified as corresponding to the semiquinone\nradical form of the prFMN cofactor. 11 , 13 , 19 Figure 2 HmfF spectral properties, in vitro reconstitution, and\noxygen dependence of activity. (A) UV–vis spectra obtained\nfor heterologous expressed P. thermopropionicum HmfF. Spectra are shown of the WT protein expressed on its own (orange\nline) or coexpressed with UbiX and purified either aerobically (purple)\nor anaerobically (green). Spectra were normalized on the A 280 peak. The inset shows the closeup of the cofactor-related\nspectral features present in the 300–800 nm region. (B) UV–vis\nspectra of single expressed “apo” P.\nthermopropionicum HmfF as isolated (orange), following\nreconstitution with in vitro synthesized prFMN under anaerobic conditions\n(red), and following exposure to air (blue). (C) Activity of reconstituted\nPtHmfF against aerobic or anaerobic substrate before and after exposure\nto air. Assays were performed against 900 μM FDCA at 25 ° C (error bars represent SEM, n = 3). P. thermopropionicum HmfF in Vitro Reconstitution Confirms Oxidative Maturation Is Required\nfor Activity While UbiX produces prFMN in a reduced state,\nthe cofactor must undergo oxidative maturation within UbiD to produce\nthe active prFMN iminium form. 8 − 11 To investigate the requirement\nfor oxidative maturation of the cofactor in HmfF, apo -enzyme was reconstituted in vitro as described previously for AroY\nand UbiD. 13 , 19 Single expressed HmfF lacking decarboxylation\nactivity was reconstituted under anaerobic conditions and revealed\nprominent features at 360 and 530 nm ( Figure 2 B). Exposure to air resulted in an enhancement\nof the spectral features at 360–380, 450, and 530 nm, a range\nof spectral features suggestive of a mixture of normal oxidized FMN,\nprFMN radical , and possibly prFMN iminium , similar\nto that observed in the as isolated coexpressed enzyme. Consistent\nwith this, the anaerobic reconstituted protein displayed low levels\nof decarboxylase activity when it was assayed under anaerobic conditions.\nHowever, the rate of enzymatic decarboxylation was 5-fold higher when\nthe protein was assayed under aerobic conditions ( Figure 2 C). Taken together, these data\nconfirm that, as with Fdc1 and AroY, HmfF requires oxidative maturation\nof prFMN for activity. Pt HmfF Is Light and Oxygen\nSensitive The activity of the as-isolated coexpressed Pt HmfF was found to rapidly decrease over time when it was\nstored on ice. The loss in activity appeared to be partially due to\nlight exposure, with the half-life of Pt HmfF increasing\nfrom 35 to ∼100 min when it was stored in the dark under aerobic\nconditions ( Figure S3 ). Similar observations\nwere made for the A. niger Fdc1 enzyme,\nwhere light exposure was found to induce a complex isomerization of\nthe cofactor leading to inactivation. 20 However, unlike Fdc1, protection from illumination was not sufficient\nto stabilize Pt HmfF activity. In contrast, Pt HmfF stored under anaerobic conditions did not appear\nto lose activity over the course of several hours, suggesting that\ninactivation was also the result of O 2 , as observed for\nAroY. 13 Subsequently, Pt HmfF was either purified anaerobically or purified aerobically and\nthen reconstituted in vitro and assayed under anaerobic conditions.\nThe Pt HmfF enzyme activity was found to have a pH\noptimum between 6 and 6.5, with a temperature maximum of ∼60\n°C ( Figure 3 ).\nHowever, from 55 °C and above the activity decreased rapidly\nover the course of a few minutes, indicating that the enzyme was being\ninactivated, making it difficult to obtain accurate initial rates.\nThermal denaturation of Pt HmfF monitored using CD\nspectroscopy revealed a melting temperature of ∼68 °C\n( Figure S4 ). Thus, all subsequent assays\nwere performed at 50 °C. An Arrhenius plot of the 25–50\n°C data points indicated an activation energy of 80.7 kJ mol –1 . At pH 6 and 50 °C, the apparent K m and k cat values for FDCA\nwere 49.4(±3.7) μM and 2.39(±0.05) s –1 , respectively ( Figure 3 C). The Pt HmfF enzyme was also found to have minor\nactivity with 2,5-pyrroledicarboxylic acid (PDCA). In contrast, no\ndecarboxylation could be detected for 2,3-furandicarboxylic acid,\n5-formyl-2-furoic acid, 5-hydroxymethyl-2-furoic acid, 5-nitro-2-furoic\nacid, 2,5-thiophenedicarboxylic acid, 2,6-pyridinedicarboxylic acid,\nterephthalic acid, isophthalic acid, or muconic acid. Figure 3 Pt HmfF\nenzyme activity. (A) Effect of pH on activity. (B) Effect of temperature\non activity. Inset: Arrhenius plot of data. (C) Steady-state kinetic\nparameters obtained for P. thermopropionicum HmfF against FDCA (blue) and PDCA (red) at 50 °C and pH 6.\nError bars represent SEM, n = 3. Pt HmfF Catalyzes H/D Exchange of a Small Range of\nHeteroaromatic Acids It has previously been shown that UbiD\nenzymes are capable of catalyzing deuterium exchange of substrates\nthat can undergo UbiD-mediated carboxylation. 20 , 21 1 H NMR showed that incubation of furoic acid with Pt HmfF in D 2 O resulted in depletion of the resonance\npeak at 7.6 ppm consistent with exchange of the proton in the 5-position\n(denoted H a ) with a deuteron ( Figure S5 ). This was further supported by a change in splitting of\nthe 6.5 ppm resonance (corresponding to H b ) from a doublet\nof doublets to a doublet resulting from the loss of coupling between\nH b and H a . Partial H/D exchange of the 5-position\nof pyrrole-2-carboxylate (∼30%, Figure S4B ) could also be observed under the conditions tested; however,\nno exchange of thiophene-2-carboxylate was detected ( Figure S5C ). These observations confirm that the level of\nH/D exchange follows the same trend as observed for the level of decarboxylation\nof the corresponding diacids. With this in mind, we used H/D exchange\nto assay Pt HmfF against substrates where the corresponding\ndiacids were not commercially available. The proton in the 5-position\nof 2-oxazolecarboxylic acid could only be readily exchanged for deuterium\nin the presence of enzyme ( Figure S5D ).\nIn contrast, no enzyme-dependent exchange could be observed for position\n2 of 5-oxazolecarboxylic acid ( Figure S5E ). Pt HmfF and Gk HmfF Catalyze\nFuroic Acid Carboxylation at Elevated [CO 2 ] The\nHmfF reverse reaction, carboxylation, has been demonstrated to occur\nin vivo for distinct UbiD members that function as dedicated carboxylases, 22 − 24 while those family members that act as decarboxylases under physiological\nconditions (such as AroY and Fdc1) can catalyze carboxylation in vitro\nat elevated levels of CO 2 . 11 , 13 , 14 To investigate the ability of HmfF enzymes to catalyze\nthe reverse reaction, carboxylation of furoic acid to produce FDCA,\npurified Pt HmfF and Gk HmfF enzymes\nwere incubated with 50 mM furoic acid and 1 M bicarbonate at 50 °C\novernight. HPLC analysis of the reaction mixtures revealed a peak\nwith retention time of 2.3 min that comigrates with an FDCA standard\n( Figure 4 A). Mass spectrometry\nconfirmed that this species had a mass of 154.99 Da, consistent with\nthe expected mass for FDCA. We sought to determine whether performing\nthe reaction under pressurized CO 2 could increase the amount\nof carboxylated product. Reaction mixtures containing 50 mM furoic\nacid were incubated overnight with HmfF at 50 °C. In the presence\nof 1 M KHCO 3 , there was no significant difference between\nreaction mixtures incubated under N 2 at atmospheric pressure\nor under CO 2 at 32 bar with ∼2 mM FDCA produced.\nIn the absence of bicarbonate, ∼150 μM FDCA was produced\nunder 32 bar CO 2 , whereas no FDCA was detectable under\nN 2 ( Figure 4 B). While HmfF presents an attractive route to the production of\n2,5-furandicarboxylic acid, a potential bioreplacement for polymer\nprecursors, yields remain low even under high [CO 2 ]. Given\nthe unfavorable equilibrium for the carboxylation reaction, future\nefforts aimed at increasing the yield for this reaction will likely\nrequire in situ conversion of FDCA. Figure 4 Pt HmfF-catalyzed carboxylation\nof furoic acid to FDCA. (A) HPLC chromatogram demonstrating enzymatic\nproduction of FDCA by carboxylation of furoic acid by P. thermopropionicum HmfF. Chromatograms of FDCA\n(red) and 50 mM furoic acid in 1 M KHCO 3 solution incubated\nin the absence (blue) and presence (purple) of the Pt HmfF enzyme. Mass spectrometry confirmed that the species that comigrated\nwith the FDCA standard also possessed a mass consistent with FDCA.\n(B) Furoic acid carboxylation under CO 2 pressure. Assays\nwith or without 1 M KHCO 3 were incubated overnight either\nunder N 2 at atmospheric pressure or under CO 2 at 32 bar. Error bars represent SEM, n = 3. Pt HmfF\nCrystal Structures Reveal FMN Binding Mode To aid crystallization,\nthe Pt HmfF was expressed without affinity tag. The\nbest crystals obtained belonged to the P 2 1 space group and diffracted to 2.7 Å. The procedure was repeated\nwith Se-Met-substituted enzyme, and the structure was solved using\nSe-Met SAD, revealing a Pt HmfF hexamer ( D 3 symmetry) in the asymmetric unit. Although the UV–vis\nspectra of the purified enzyme used for crystallization trials indicated\nthe presence of cofactor, no electron density corresponding to the\ncofactor could be detected in preliminary electron density maps. Final\nrefinement was done using data collected to 2.7 Å on Pt HmfF crystals soaked with FMN (as a stable analogue of\nthe prFMN cofactor) in the presence of K + and Mn 2+ , revealing clear electron density for both the FMN and the associated\nmetal ions in the prFMN binding site ( Figure S6 ). The Pt HmfF structure is similar to other UbiD\nfamily member structures with a Z score of 47 with\nthe bacterial protocatechuate decarboxylase AroY (rmsd 1.6 Å\nover 441 C-αs), 13 45 with the E. coli UbiD (rmsd 2.5 Å over 440 C-αs), 19 40 with the fungal cinnamic acid decarboxylase\nFdc1 (rmsd 3.1 Å over 440 C-αs) 11 and 38 with recently solved TtnD decarboxylase involved in polyketide\nbiosynthesis (rmsd 2.7 Å over 414 C-αs). 25 The Pt HmfF monomer consists of an N-terminal\nprFMN binding domain connected via an α-helical linker to the\noligomerization domain ( Figure 5 ). The C-terminus consists of an extended loop region with\nsome α-helical character that interacts with the prFMN binding\ndomain of an adjacent Pt HmfF monomer. An overlay\nof the six Pt HmfF monomers reveals that minor variation\noccurs in the respective positions of the N-terminal prFMN binding\ndomain and the oligomerization domain, suggestive of domain motion\nvia the hinge region connecting both domains ( Figure 5 B). As the active site (vide infra) is located\nat the interface between both domains, this could be relevant to catalysis.\nThe phosphate moiety of the bound FMN is coordinated by Mn 2+ and K + ions (the identity of these was derived from the\nfact they were added to the crystal and was not independently verified),\nwhile the isoalloxazine is positioned directly adjacent to the conserved\nE(D)-R-E ionic network of residues conserved in UbiD ( Figure 5 C). In the case of Pt HmfF, the active site is only partly occluded from solvent\nas a consequence of the relatively open conformation of the N-terminal\nprFMN binding domain; this is similar to what has been observed for\nthe canonical UbiD and AroY enzymes ( Figure S7 ). To achieve full occlusion from solvent, as is observed for the\nfungal Fdc1 enzyme, a hinge motion (akin to that observed by comparison\nof the various Pt HmfF monomers) leading to a closed\nconformation would be required. Figure 5 Pt HmfF crystal structure.\n(A) Pt HmfF hexamer (D3 symmetry) shown in two orientations,\nrepresented in cartoon depiction, with the prFMN binding domain in\nblue, the connecting helix in magenta, the hexamerization domain in\ngreen, and the C-terminal helix in red. The bound FMN is shown as\nyellow spheres. Arrows indicate the interfaces disrupted by mutagenesis\n(vide infra). (B) Overlay of the six Pt HmfF monomers\npresent in the asymmetric unit with the C-α traces depicted\nin ribbon using a color coding similar to (A). (C) Side-by-side comparison\nof the Pt HmfF active site with other structurally\ncharacterized UbiD family members. Key residues are show in atom color\nsticks, with carbons colored according to domain structure as used\nin (A). In the case of the Aspergillus niger Fdc1 enzyme, the α-fluorocinnamic acid complex is shown, with\nthe substrate shown in cyan carbons. In the case of the TtnD enzyme,\nthe loop containing residues E272–E277 is not ordered in the\nFMN-bound structure. The conserved E(D)-R-E motif is highlighted by\nthe use of red labels. Pt HmfF Active Site Contains a Furoic Acid Binding\nMotif All attempts to acquire a crystal structure of the Pt HmfF in complex with substrate through either soaking\nor cocrystallization failed, a possible consequence of the open configuration\nof the enzyme. Guided by the structure of the related Fdc1 in complex\nwith cinnamic acid substrates, 11 FDCA can\neasily be placed into the active site of Pt HmfF in\na similar position with respect to the prFMN cofactor. This positions\nthe substrate furan oxygen approximately within hydrogen-bonding distance\nof His296 and locates the distal substrate carboxylate adjacent to\nArg304 and Arg331. All three putative substrate binding residues are\nconserved in the HmfF branch of the UbiD family tree ( Figure S1 ). To support our hypothesis regarding\nthe role of H296, R304, and R331 in substrate binding, we made Pt HmfF H296N, R304A, and R331A variants. All variants possess\nUV–vis spectra similar to that of the WT with the exception\nof the H296N variant ( Figure 6 A). In the latter case, cofactor related features between\n300 and 400 nm are less intense, indicating lower cofactor content.\nProlonged incubation of large quantities of protein with substrate\nresulted in complete decarboxylation of FDCA by WT Pt HmfF and the R304A variant (assayed by HPLC). Under these conditions,\nH296N was able to decarboxylate ∼87% of the substrate, in comparison\nwith 30% using the R331A variant ( Figure 6 B). Using PDCA as a substrate, only the WT\nand H296N were able to perform 100% decarboxylation ( Figure 6 C). A continuous spectrophotometrically\nbased assay (using 1 mM substrate) revealed that all three variants\nwere severely compromised in activity in comparison to the wild type\nenzyme, with k cat values 30–400\nfold lower than the WT against FDCA ( Figure 6 D). Interestingly, the H296N variant displays\na preference for 2,5-pyrroledicarboxylic acid over FDCA and has a\nslightly higher activity for PDCA in comparison to the WT enzyme ( Figure 6 D). Michaelis–Menten\nkinetics revealed that both R304A and H296N variants were not saturated\nat 1 mM substrate (the maximum possible under the experimental conditions),\nwhile no reliable data could be obtained for R331A ( Figure 6 E). These data clearly indicate\nthe FDCA affinity has been compromised by substitutions at positions\nH296, R304, and R331. Figure 6 Characterization of Pt HmfF variants.\n(A) UV–vis spectra of Pt HmfF variants including\nWT (blue), H296N (green), R304A (magenta), R331A (orange), and L403A\n(red). Spectra were normalized on the A 280 peak. The inset shows a closeup of the cofactor-related spectral\nfeatures present in the 300–800 nm region. (B) Decarboxylation\nof 10 mM 2,5-furandicarboxylic acid (FDCA) to furoic acid after overnight\nincubation with Pt HmfF variants. (C) Decarboxylation\nof 10 mM 2,5-pyrroledicarboxylic acid (PDCA) to pyrrole-2-carboxylate\nafter overnight incubation with Pt HmfF variants.\n(D) Rate of decarboxylation of 1 mM FDCA (blue) or 1 mM PDCA (red)\nby Pt HmfF variants. (E) Steady-state kinetics of Pt HmfF variants. Error bars represent SEM, n = 3. A Dimeric Pt HmfF Variant Binds prFMN but Is Compromised for Activity The resolution of the hexameric Pt HmfF structures\nobtained is limited and is in sharp contrast to the atomic resolution\nroutinely achieved for the dimeric A. niger Fdc1. A structural alignment of Pt HmfF hexamer\nwith the related A. niger Fdc1 dimer\nstructure demonstrates that Pt HmfF A315, N348, F351,\nT355, A388, F393, V395, and M399 form key hydrophobic interactions\nbetween the individual Pt HmfF dimers. In Fdc1, the\nequivalent positions are D343, R382, D385, N389, P424, T429, F431,\nand R435, respectively: i.e. generally larger and/or charged residues.\nWe created a Pt HmfF variant by substituting for the\ncorresponding A. niger Fdc1 amino acids\n(i.e., A315N, N348R, F351D, T355N, A388P, F393T, V395F, and M399R)\nto disrupt dimer–dimer interactions. SEC-MALLS of the purified Pt HmfF dimer variant indicated a native mass of 110 kDa,\nbroadly consistent with the expected mass of a dimer. Similarly to\nthe WT protein, the purified dimer variant possesses a complex UV–vis\nspectrum with three main features in the 300–800 nm region\n( Figure S8 ). This suggests the presence\nof prFMN, in addition to minor populations of FMN and the radical\nprFMN. Despite the presence of prFMN, the dimer variant display weak\nactivity, and incubation of 10 mM FDCA against 20 μM Pt HmfF dimer mutant only resulted in 30% decarboxylation\nfollowing overnight incubation ( Figure 6 B). Crystal Structure of Dimeric Pt HmfF Suggests a Key Role for a Conserved Leu in Activity The P. thermopropionicum HmfF dimer\nvariant was crystallized and the structure solved to 2.3 Å using\nmolecular replacement with the WT Pt HmfF monomer.\nUnlike the wild type enzyme, the dimer variant crystals contain prFMN\nin the active site. Despite the extensive mutation of the WT dimer–dimer\ninterface, the Pt HmfF dimer variant is very similar\nin structure to an individual dimer module from the WT hexamer. The\nprFMN is bound in a similar position and configuration as the FMN\nin the Pt HmfF hexamer, with little difference in\nthe position of the majority of active site residues. A notable exception\nis Leu403, which is located on a loop region that is disordered in\nthe Pt HmfF dimer variant and therefore absent from\nthe active site ( Figure 7 ). The 398–410 region including Leu403 is disordered in both\nthe Pt HmfF dimer variant monomers, a likely consequence\nof the M399R mutation and/or the disruption of the WT dimer–dimer\ninterface. As a consequence, the Pt HmfF dimer variant\nactive site is exposed to the solvent. To confirm whether the absence\nof Leu from the dimer active site contributed to the low activity\nof the dimer mutant, a L403A Pt HmfF was created.\nWhile the UV–vis profiles of both WT and the L403A variant\nare comparable, the latter had a k cat value\n∼40-fold lower than that of the WT ( Figure 6 D,E). However, unlike the H296 and R304/331\nvariants, the L403A K m value for FDCA\nwas not significantly different from the WT, suggesting that L403\ndoes not contribute to substrate binding. The hydrophobic nature of\nthe carboxylic acid binding pocket has been implicated in the mechanism\nof other decarboxylases, 26 , 27 and it is plausible\nthat the highly conserved Leu403 fulfils a similar role. Furthermore,\nLeu403 is one of the residues most affected by the proposed domain\nmotion ( Figure 5 B and Figure S7 ) that occludes the Pt HmfF active site from solvent. Figure 7 Pt HmfF dimer variant\ncrystal structure. (A) Pt HmfF dimer variant shown\nin cartoon representation, with color coding as in Figure 5 a. The mutations disrupting\nthe hexamer formation interface are indicated by cyan spheres for\nthe corresponding Cα positions. (B) Overlay of the two Pt HmfF monomers present in the asymmetric unit with the\nCα traces depicted in ribbon using a color coding similar to\nthat in (A). In addition, a single monomer of the Pt HmfF hexamer structure is shown in gray. (C) Position of the Leu403\nregion (in red) at the prFMN-domain/multimerization domain interface\nthat is disordered in the Pt HmfF dimer structure.\nMutations at the hexamer formation interface are shown as cyan spheres,\nexcept for M399R, which is shown in red. (D) Overlay of the respective Pt HmfF hexamer active site (in complex with FMN, in gray)\nand the Pt HmfF dimer variant in complex with prFMN iminium . For comparison, the position of the α-fluorocinnamic\nacid substrate of A. niger Fdc1 with\nrespect the prFMN cofactor is shown (in cyan), as well as the corresponding\nposition of L439 (homologous to Pt HmfF L403). Proposed Mechanism for\nHmfF Previous studies for the Fdc1 enzyme have suggested\nthat the reversible decarboxylation occurs via a 1,3-dipolar cycloaddition\nbetween substrate and prFMN. In principle, a similar reaction scheme\ncan be proposed for any of the UbiD substrates. However, for those\nsubstrates where (de)carboxylation occurs directly on an aromatic\nring system, cycloaddition also requires dearomatization. An alternative\nproposal has been put forward for AroY, on the basis of formation\nof a quinoid intermediate that allows formation of a substrate–prFMN\nadduct. In view of the modest aromatic nature of the furan ring, an\nFDCA or furoic acid adduct with prFMN could be formed through either\ncycloaddition (Ib in Figure 8 ) or via formation of an oxonium ion (Ia in Figure 8 ) in the case of HmfF. Chemical\nprecedent exists for the 1,3-cycloaddition of furans to 1,3-dipoles. 28 It is unclear at present which route is preferred\nfor HmfF, and this will require further investigation. However, it\nis interesting to note that HmfF-catalyzed H/D exchange can be readily\nobserved for weakly aromatic heteroaromatic acids only at those positions\nthat are adjacent to a carbon, hinting at the possibility that species\nIb is indeed formed during the enzyme reaction. Furthermore, neither\ndecarboxylation of the more aromatic thiophenedicarboxylic acid nor\nH/D exchange of thiophene-2-carboxylate was observed. In fact, thiophene\ncompounds are not known to readily undergo cycloaddition reactions,\nin contrast to furan. While HmfF is able to catalyze furoic acid carboxylation\nunder ambient conditions, this requires a mechanism for increasing\n[CO 2 ]. Other decarboxylases have been found to catalyze\npyrrole carboxylation in supercriticial CO 2 , 29 and we intend to explore whether HmfF (natural\nor evolved variants) can be used under these conditions. Further studies\nwill also need to address cofactor stability and homogeneity to ensure\na robust biocatalyst for the carboxylation of furoic acid. Figure 8 Proposal for\nthe HmfF mechanism. The HmfF substrate is bound by polar interactions\nwith the HmfF specific R304/R331 and H296, in addition to the UbiD\nfamily conserved R152 (part of the UbiD Glu-Arg-Glu motif). Substrate\nbinding is possibly linked to domain motion, affecting the relative\nposition of L403 and R331. Formation of a covalent prFMN iminium –substrate adduct can occur either through nucleophilic attack,\nleading to species Ia, or through cycloaddition, leading to species\nIb. Decarboxylation of either species leads to intermediate II, following\nE260/CO 2 exchange; protonation of the substrate via E260\nleading to product release occurs through either intermediate IVa\nor IVb."
} | 6,651 |
39759078 | PMC11700627 | pmc | 4 | {
"abstract": "Summary Largely varied anti-icing performance among superhydrophobic surfaces remains perplexing and challenging. Herein, the issue is elucidated by exploring the roles of surface chemistry and surface topography in anti-icing. Three superhydrophobic surfaces, i.e., gecko-like, petal-like, and lotus-like surfaces, together with smooth hydrophobic and hydrophilic surfaces, are prepared and compared in ice nucleation temperature under both non-condensation and condensation conditions. As a result, in non-condensation condition, water droplet freezing is caused by interfacial heterogeneous nucleation, wherein both surface chemistry and surface topography contribute to deferring freezing, and the former is dominant. In condensation condition, the freezing strongly correlates to condensation frosting. Surface chemistry maintains as a strong deterrent, whereas surface topography has two competing effects on the freezing. The paper deepens the understanding of water freezing on superhydrophobic surfaces, unravels the correlation between superhydrophobicity and anti-icing, and provides design guidelines on application-oriented anti-icing surfaces.",
"introduction": "Introduction Icing, a ubiquitous phenomenon, however brings about devastating disasters to air and road traffic, malfunction of solar cells and wind turbines, and plummeted crop production. 1 To this end, much effort has been devoted to anti-icing, among which superhydrophobic strategy is a widely adopted recipe to lower ice nucleation temperature (INT) and prolong freezing delay time (FDT) in undercooling conditions. 2 , 3 , 4 Although the anti-icing performance of superhydrophobic surfaces has been intensively studied over the last decade, it is still in dispute, which can be reflected from a wide INT gap of over 10°C and two orders of magnitude deviation for FDT. 5 , 6 , 7 , 8 , 9 , 10 , 11 , 12 Apparently, the varied anti-icing performance was derived from different superhydrophobic genres. Superhydrophobic surfaces are realized via regulating both surface chemistry and surface topography. 13 , 14 Employment of low-surface-energy materials or coatings is necessary for superhydrophobicity. Designs of surface topographies owning multi-scale hierarchical roughness and trapped air pockets are also indispensable. Since the revelation of dual-scale nano-/micro-textures for lotus leaves, a plethora of biomimetic superhydrophobic surfaces have sprouted up, possessing distinctive topographies, e.g., lotus-like, petal-like, and gecko-like. 15 , 16 , 17 , 18 , 19 Generally, such characters of superhydrophobic surfaces benefit the anti-icing performance by lifting the energy barrier for ice nucleation and/or lowering the heat transfer between surfaces and water droplets. 20 , 21 , 22 , 23 , 24 Moreover, the varied anti-icing performance of superhydrophobic surfaces was tested in different conditions. The effects of droplet size and cooling rate can be well explained by classical nucleation theory (CNT). 25 In addition, relative humidity (RH) has a profound influence on the freezing of water droplets. On one hand, high RH probably destabilizes the trapped air pockets and increases surface-droplet contact in a direct manner. 26 , 27 , 28 On the other hand, the ambient water vapor may condense on surfaces to interfere the freezing of water droplets in an indirect manner. 29 , 30 Therefore, to find out the origin of anti-icing performance variation among superhydrophobic surfaces, the roles of surface chemistry and surface topography in anti-icing under different conditions should be traced and explicated, which though is absent to the best of our knowledge. Herein, five samples, i.e., smooth hydrophilic surface, smooth hydrophobic surface, gecko-like, petal-like, and lotus-like superhydrophobic surfaces, are prepared. And their anti-icing properties are compared in two test conditions, which are condensation and non-condensation conditions. Briefly, the respective INT of gecko-like, petal-like, and lotus-like superhydrophobic surfaces is −30.5, −30.5, and −30.4°C in non-condensation condition, which goes to −15.0, −12.8, and −16.1°C in condensation condition accordingly. The nearly same INT values in non-condensation condition deviate from each other in condensation condition, and the INT values in non-condensation condition are far superior to the counterparts in condensation condition. The reason is that the freezing mechanism of water droplets experiences a change when condensation occurs, and thus the roles of surface chemistry and surface topography vary, resulting in such distinct anti-icing performance.",
"discussion": "Discussion Herein, a huge INT gap up to 18°C was discerned among different kinds of superhydrophobic surfaces using different test environments. In non-condensation condition, the freezing of water droplets on surfaces is caused by heterogeneous nucleation at droplet-surface interface, which instead strongly correlates to condensation frosting in condensation condition. Owning to such distinct freezing mechanisms, the roles of surface chemistry and surface topography of superhydrophobic surfaces in anti-icing vary significantly. In non-condensation condition, both surface chemistry and surface topography contribute to anti-icing. The former raises the energy barrier for ice nucleation by employing low-surface-energy materials, and the latter lowers the nucleation kinetics by reducing the contact between droplets and surfaces. In condensation condition, the role of surface chemistry remains positive, which defers the frost propagation by forming dropwise condensation. But the role of surface topography becomes complex: on one hand, it speeds up the frost propagation due to the increasing condensate density; on the other hand, it slows down the frost propagation via the fast coalescence of condensates that is evident in lotus-like surface. As a result, low surface energy benefits the anti-icing in both non-condensation and condensation conditions. Nano-sized textures are promising in non-condensation condition, whereas frost-free/frost-delay textures enabling fast coalescence, jumping removal, 29 , 45 or large vapor pressure gradients 66 have great potential in condensation condition. Limitations of the study We have revealed two different freezing mechanisms of water droplets on superhydrophobic surfaces and thus the role variation of surface chemistry and surface topography in anti-icing. However, we set several preconditions: there was no heat transfer between surfaces and water droplets in non-condensation environment; there was no jumping removal of condensates in condensation environment; and all surfaces were placed horizontally. Therefore, future investigation is needed to explore more situations."
} | 1,691 |
40227506 | PMC11996751 | pmc | 5 | {
"abstract": "As an emerging memory device, memristor shows great potential in neuromorphic computing applications due to its advantage of low power consumption. This review paper focuses on the application of low-power-based memristors in various aspects. The concept and structure of memristor devices are introduced. The selection of functional materials for low-power memristors is discussed, including ion transport materials, phase change materials, magnetoresistive materials, and ferroelectric materials. Two common types of memristor arrays, 1T1R and 1S1R crossbar arrays are introduced, and physical diagrams of edge computing memristor chips are discussed in detail. Potential applications of low-power memristors in advanced multi-value storage, digital logic gates, and analogue neuromorphic computing are summarized. Furthermore, the future challenges and outlook of neuromorphic computing based on memristor are deeply discussed.",
"conclusion": "Conclusion and Perspectives Overall, memristors represent progress in neuromorphic computing architectures, bringing significant advantages with their inherent physical properties and operational characteristics. First, non-volatile resistance state allows them to store information without additional data transmission power consumption. Second, many memristors achieve stable switching characteristics at feature sizes below 10 nm, with great potential for expansion. In-memory computing eliminates the traditional von Neumann bottleneck and greatly reduces the energy consumption associated with data movement between independent processing and storage units. The adjustable multi-level storage state enables matrix multiplication and weight updates for neuromorphic computing. With excellent CMOS compatibility, memristors facilitate integration into existing semiconductor manufacturing workflows, while supporting new computing paradigms such as logic-in-memory and brain-inspired neuromorphic computing. Recent demonstrations of memristor-based neural networks have achieved remarkable energy efficiencies below 1 fJ per synaptic operation, which is orders of magnitude better than conventional digital implementations and biological computing. The development of new materials remains key to improving the performance of memristors. Researchers are exploring 2D materials such as graphene and transition metal dichalcogenides, which have unique electrical properties and atomic-level thickness. These materials can achieve more precise resistance modulation and lower power consumption. Research on metal oxides continues, focusing on designing defect states and interface properties to achieve better switching characteristics and reliability. Array structure optimization is to minimize sneak current and improve read/write margins. Advanced selectors can be developed, including volatile switch selectors and engineered tunnel barriers. At the same time, three-dimensional integration strategies are explored to increase storage density while maintaining low power consumption levels. For storage applications, researchers are developing more complex programming schemes and error correction methods. Research on new switching mechanisms such as phase change and magnetoresistance effects may produce hybrid devices that combine the advantages of different storage mechanisms. For the digital logic computing, future research focuses on optimizing device characteristics for logic operations, developing more efficient programming schemes, and creating new circuit topologies that exploit the unique properties of memristors. In neuromorphic computing, future research will focus on developing ultra-low-power devices and systems, achieving extremely low programming currents to achieve ultra-low-power-consumption pulse generation and transmission. In terms of training schemes, future development schemes need to take into account the non-ideality of the device and optimize the power-performance balance through approximate computing techniques and multi-device architectures. The development of multi-functional memristor is also advancing, which can perform synaptic and neural functions at the same time, thereby achieving more compact and efficient neuromorphic systems. In response to the challenges of neuromorphic applications, researchers are improving energy efficiency through innovative programming schemes and adaptive precision techniques. Future work will also implement online learning algorithms under low-power operation and explore the use of complementary memristor devices to simplify the weight update process. In addition, the integration of memristor neuromorphic systems with CMOS circuits is also being optimized, especially interface circuits operating at low voltages.",
"introduction": "Introduction Von Neumann architecture is the basic architecture of modern computers, proposed by mathematician John von Neumann in 1945. Its core idea is to store program instructions and data in the same memory block and process the data by reading and executing these instructions through a central processing unit (CPU). This architecture's primary benefit lies in its adaptability and malleability, allowing the computer to undertake various tasks by altering programs stored in its memory [ 1 ]. However, von Neumann structure has its inherent flaws, where data storage and computing share the same channel. Such working mode limits processing speed of computer, especially if it uses dynamic random access memory (DRAM) as its primary memory. DRAM access not only requires high energy consumption, but also requires periodic refreshing. During data processing, the processor has to run continuously even while waiting for data, leading to additional energy consumption. As a result, the so-called “energy wall” and “speed wall” are formed. As internet technology rapidly evolves, the demand for artificial intelligence is experiencing exponential growth. Artificial intelligence has achieved numerous breakthroughs in various domains, including image processing, natural language processing, and big data analysis [ 2 – 4 ]. The amount of data that need to be trained and processed are also increasing daily. To address this problem, complex hardware systems consisting of numerous CPUs and graphics processing units (GPUs) have been developed. As semiconductor technology is approaching its physical limits, Moore's law is also facing failure [ 5 , 6 ], and researchers must examine the constraints of von Neumann architecture through the lens of computer architecture and software algorithms. In this regard, researchers have proposed various approaches, such as the introduction of multi-level caches [ 7 ], the introduction of data streaming [ 8 ], and the proposal of in-memory computing. Among emerging technologies, in-memory computing, first conceptualized by W.H. Kautz in 1969 [ 9 ], seamlessly integrates computational functions within storage, drastically reducing the delay for data transfer. This integration further leads to reduced power consumption and improved efficiency and is hailed as the next-generation computer architecture poised to transcend the barriers of von Neumann architecture. In recent years, there has been a swift advancement in the development of novel non-volatile memory and in-memory computing technology. With high speed, low power consumption and high-density integration capability, memristor is becoming a research hotspot in in-memory computing fields. Inspired by human brain, memristors with weights updating functions are considered ideal for developing in-memory computing and artificial intelligence [ 10 ]. This paper summarizes the research progress of memristors in the field of in-memory computing and artificial intelligence from the perspective of power consumption, covering the aspects of the device structure, mechanism, and key performance parameters of memristors, as well as the introduction of memristor arrays. Then, the low-power functional materials applied in memristors are categorized and discussed. Afterward, the review focuses on discussion of reducing power consumption in several compelling application areas of memristors, especially in multi-bit memories, logic gates, and neuromorphic computing. By summarizing the principles of memristors applied therein, the low-power implementation mechanism is well analyzed. Furthermore, the existing research progress, future challenges and outlook are discussed in detail. Figure 1 shows the overview of this review article. Figure 2 a shows the von Neumann architecture diagrams mentioned above. Figure 2 b shows a schematic diagram of the “energy wall” and the “speed wall”. Fig. 1 Overview of memristors for low-power storage and computing: including devices, materials, artificial synapses and neurons, and neural networks. From the device level, resistive random access memory (RRAM), phase change random access memory (PCRAM), magnetoresistive random access memory (MRAM) and ferroelectric device are potential low-power neuromorphic computing electronics. From materials system level, ion transport materials, phase change materials, magnetoresistive materials and ferroelectric materials are main functional material layers for low-power memristors. These novel memristors could be used to act as artificial synapses and neurons for low-power neuromorphic computing, including artificial neural network (ANN), spiking neural network (SNN) and convolutional neural network (CNN) Fig. 2 a Schematic illustration of the segregation structure. b Schematic representation of the “energy wall” and “speed wall” facing the von Neumann structure. c Schematic diagram of RRAM device structure. d Schematic diagram of PCRAM device structure. e Schematic diagram of MRAM device structure. f Schematic diagram of ferroelectric device structure Memristor The concept of memristor was first proposed by Professor Chua in 1971, which was the fourth basic passive circuit element after resistance, capacitance, and inductance, filling the gap in the description of the relationship between electric charge and magnetic flux [ 11 ]. Its mathematical model is expressed as the ratio of magnetic flux to electric charge, \\documentclass[12pt]{minimal}\n\t\t\t\t\\usepackage{amsmath}\n\t\t\t\t\\usepackage{wasysym} \n\t\t\t\t\\usepackage{amsfonts} \n\t\t\t\t\\usepackage{amssymb} \n\t\t\t\t\\usepackage{amsbsy}\n\t\t\t\t\\usepackage{mathrsfs}\n\t\t\t\t\\usepackage{upgreek}\n\t\t\t\t\\setlength{\\oddsidemargin}{-69pt}\n\t\t\t\t\\begin{document}$$M = {\\text{d}}\\varphi /{\\text{d}}q$$\\end{document} M = d φ / d q , the resistance is determined by the magnetic flux. It is a nonlinear resistance element with memory characteristics. However, in actual physical systems, direct coupling of magnetic flux and charge is not easy to achieve, and ideal memristors remain more at the theoretical level. Although many devices do not strictly meet the definition of ideal memristors, they exhibit similar characteristics, especially the non-volatile characteristic and adjustability. The realization of generalized memristors is usually based on ion migration, the formation and breaking of conductive filaments (CFs), phase change or magnetic spin effects, etc. RRAM, PCRAM, MRAM and ferroelectric memristor have emerged. RRAM is one of the typical representatives of memristors. Its resistance state is determined by the distribution of oxygen vacancies or CFs inside the material. The resistance can be changed by voltage pulses and be retained after removing pulses. The structure is usually divided into electrodes and functional layers, presenting a sandwich structure of electrode-functional layer-electrode, as shown in Fig. 2 c. PCRAM uses phase change materials between crystalline and amorphous states to achieve resistance change. The material can be heated to different states under different current pulses, with low resistance in the crystalline state and high resistance in the amorphous state, thereby achieving data writing and storage. The PCRAM device structure is generally mushroom-shaped, with a wider top electrode, a narrower bottom electrode, and a layer of phase change material in the middle. The device structure is shown in Fig. 2 d. MRAM uses the non-volatile magnetic materials and spin electronics for storage. It stores data through a magnetic tunnel junction (MTJ), which consists of two layers of magnetic material and an insulating layer. One magnetic layer is fixed, and the magnetization direction of the other free layer can be changed by current. The resistance state of the MTJ represents the data, with low resistance corresponding to parallel magnetization and high resistance corresponding to antiparallel magnetization. The device structure is shown in Fig. 2 e. Different from early MRAM relying on magnetic field induced switching, spin-transfer torque (STT) technology directly changes the magnetization direction of the free layer through current, reducing power consumption and suitable for high-density storage. Spin-transfer torque random access memory (STT-RAM) is developed based on STT technology. Similarly, there is spin–orbit torque random access memory (SOT-RAM) that uses the spin–orbit torque (SOT) effect. Ferroelectric memristor uses the polarization characteristics of ferroelectric materials to regulate the resistance state of the device. Ferroelectric materials have reversible polarization direction. When an external electric field is applied, the polarization direction of ferroelectric materials can be flipped, thereby changing the barrier height or interface charge distribution. This change affects the tunneling behavior of the current and the conductivity characteristics and ultimately manifests as different resistance states, as shown in Fig. 2 f. Functional Materials According to the common memristor types, memristor functional layer materials can be divided into ion transport materials, phase change materials, magnetoresistive materials and ferroelectric materials, as shown in Fig. 3 . Ion transport materials are mainly targeted at RRAM. In recent years, research in this area has mainly focused on inorganic and organic materials, specifically oxides, perovskites, two-dimensional (2D) materials and organic materials. Inorganic oxides have excellent performance and mature preparation technology and are currently widely used, but traditional binary oxides still have problems such as large leakage current and large power consumption. By doping or constructing multi-layer oxide heterojunctions, the formation and dissolution of conductive filaments can be improved for low-power-consumption storage. Perovskites and two-dimensional materials have unique structures, so they have excellent ionic conductivity and low-voltage operation [ 12 , 13 , 14 , 15 ]. Organic materials are regarded as strong competitors for the next generation of memory due to their flexibility, adjustability and low-cost potential, especially in flexible devices [ 16 ]. Typical research performance reports are summarized in Table 1 . Phase change materials are mainly chalcogenide alloys, with Ge–Sb–Te (GST) as the core. Recent PCRAM devices are also based on GST for heterogeneous doping and proportion alloying. When evaluating the impact of phase change materials on the performance of PCRAM devices, crystallization temperature, thermal conductivity, etc. are key indicators [ 17 ]. Khan et al. introduced GeTe/Sb 2 Te 3 superlattice structure in PCRAM, reducing heat loss and power consumption by 25–30 times [ 18 ]. Yang et al. introduced a conductive bridge phase change mechanism into a heterogeneous Ge-Sb-O alloy, which achieved fJ-level energy consumption (43 fJ) [ 19 ]. These works provide evidence for low-power-consumption applications of PCRAM. Magnetoresistive materials with spin polarization characteristics are mainly used for MRAM, where the free layers and fixed layers are made of ferromagnetic materials. As a king of typical ferromagnetic material, CoFeB can form a good interface with the insulating layer and has a low magnetization reversal energy. MgO usually acts as an insulator in the magnetic tunnel junction and can achieve a high tunnel magnetoresistance ratio. Most applications require MTJ to have perpendicular magnetic anisotropy (PMA), that is, the magnetization direction of the material is more likely to be arranged in a direction perpendicular to the plane of the film. PMA is related to the interface effect, lattice structure and stress of the material. The general methods to improve PMA include stacking materials with strong spin–orbit coupling such as ruthenium, cobalt or platinum in the buffer layer, or using an external voltage to regulate the magnetic anisotropy of the magnetic material. STT-RAM has been partially commercialized, but due to high current requirements and material degradation, researchers introduce SOT-RAM to reduce power consumption and increase write speed through the spin–orbit torque effect. The most studied SOT materials are heavy metal materials and topological insulators with strong spin Hall effect or Rashba effect [ 20 ]. Heavy metal materials such as Ta, W and Pt are used for the SOT layer, which have a high spin Hall angle and can efficiently generate spin currents. The surface states of topological insulators (such as Bi 2 Se 3 , Bi 2 Te 3 ) have high spin polarization rates and can achieve efficient spin injection at low currents.\n Table 1 Summary of the characteristics of the four functional materials of RRAM related to device research Structure Thickness Operating voltage Programming power consumption Endurance Year of publication Inorganic oxides and heterojunctions ITO/Bi:SnO₂/TiN [ 21 ] 20 nm − 0.5 V/0.4 V The SET operating power is 16 µW 10⁷ 2020 Ag/SiO₂/Ta₂O₅/Pt [ 22 ] 6.5 nm 0.14 V to 0.24 V/− 0.06 V to − 0.14 V N/A > 1000 2020 Pd/BaTiO 3 :Nd 2 O 3 /La 0.67 Sr 0.33 MnO 3 (LSMO)/STO [ 23 ] BNO: 34 nm LSMO:12 nm − 1 V/2 V 0.45 fJ per synaptic event > 10 10 2024 Two-dimensional materials Au/h-BN/Ti [ 24 ] 5 nm − 0.5 V/0.5 V 1.2 pJ/pulse, 30 ns pulse width and 45 µA current > 6000 2023 Ti /h-BN/Au [ 25 ] ~ 2.3 nm 2.75 V < 2 pJ 600 2024 Pt/WSe 2 /Hf x Zr 1−x O 2 (HZO)/TiN [ 26 ] WSe 2 : ~ 0.7 nm HZO:10 nm − 1.2 V / 1.5 V N/A > 2000 2025 Au/CuInS2/Cu [ 27 ] N/A 0.6 V 10 nW 1000 2025 Perovskite materials Ag/CH 3 NH 3 PbI 3 /FTO [ 28 ] 350 nm − 0.2 V/0.2 V ~ 47 fJ μm −2 > 10 3 2020 Ag/BA 2 MA 5 Pb 6 I 19 /Pt [ 29 ] ~ 300 nm − 0.15 V/0.15 V ~ 150 μW, I cc = 1 mA > 5 × 10 6 2024 Organic Materials Al/Cu-doped pMSSQ/Al [ 30 ] ~ 80 nm < 0.9 V < 0.5 pJ per pulse 500 2017 Ag/PFC-73/ITO [ 31 ] 114 nm 0.86 V N/A 60 2023 ITO/PEDOT:PSS/D:A/PDINN/Ag [ 32 ] The light intensity used (ranging from 0.51 to 194.01 mW cm −2 ) 2023 Fig. 3 Schematic diagram of memristor classification of different functional materials, including ion transport, phase change, magnetoresistive and ferroelectric. Among them, ion transport materials include organic and inorganic types [ 33 ]. Copyright (2014) American Chemical Society [ 34 ]. Copyright (2019) Wiley‐VCH, phase change materials are mainly chalcogenide alloys [ 35 ]. Copyright (2022) The Authors [ 36 ]. Copyright (2020) The Authors, magnetoresistive materials mainly constitute MTJ [ 37 ]. Copyright (2023) Science China Press [ 38 ]. Copyright (2024) The Authors, and ferroelectric materials mainly have spontaneous polarization characteristics [ 39 ]. Copyright (2020) The Authors [ 40 ]. Copyright (2024) Wiley‐VCH Ferroelectric materials can achieve reversible polarization reversal under electric field, thereby regulating tunneling current or interface charge distribution and realizing resistance state storage. Classical ferroelectric materials include bismuth titanate (BTO) and barium strontium titanate (BST), which are widely used in ferroelectric tunneling junctions due to their high remanent polarization and low leakage current. Because of excellent complementary metal–oxide–semiconductor (CMOS) compatibility, hafnium oxide-based materials (such as doped HfO 2 ) have become a research hotspot in recent years, especially in low-power and high-density memories. Two-dimensional ferroelectrics is a kind of emerging ferroelectric materials, such as In 2 Se 3 and MoTe 2 , which have ultra-thin thicknesses and are suitable for high-density integration and flexible electronics. Figure 4 summarizes the power consumption of various memristors when completing synaptic operations. RRAM and ferroelectric memristors can reach a lower level than biological levels of 10 fJ. The reported lowest power consumption is 4.28 aJ of HfAlOx-based RRAM, indicating that RRAM exhibits great potential in low-power neuromorphic computing. Therefore, the following content will be expanded on low-power-consumption RRAM. Fig. 4 Power consumption of different low-power memristors when performing synaptic plasticity [ 40 – 58 ], where biological synaptic power consumption is ~ 10 fJ. The reported power consumption of novel memristors range from 5 nJ to 4.28 aJ, exhibiting great potential in neuromorphic computing Memristor Array Two typical structures of memristor array are the 1 transistor 1 resistor (1T1R) array and the crossbar array. As illustrated in Fig. 5 a, 1T1R arrays are active arrays where each memristor is connected in series with a transistor. The word lines connect to the gate electrode of transistor, and the source lines connect to the source of the transistor. The bit lines connect to the top electrode of the memristor, and the bottom electrode connects to the drain of the transistor. The cell area of a 1T1R array is typically 12F 2 (F is the minimum feature size). As illustrated in Fig. 5 b, crossbar arrays are passive arrays with 4F 2 , consisting of perpendicular word lines and bit lines that form a crossbar structure. Memristors are arranged at the cross-points, which is more suitable for integration than a 1T1R and has no quiescent power dissipation. However, crossbar structure is prone to latent path currents. The latent path currents will flow through the other path resistors, thus causing inaccurate readings in the calculations, as well as additional power losses. In contrast, the 1T1R array, with its larger cell area and better isolation of neighboring cells, has no risk of sneak currents, which has higher computational read accuracy. For the crossbar array, a common approach to solving this problem is increasing the I–V nonlinearity by connecting a selector in series with one end of each memristor cell. The selector can use either a diode a resistor (1D1R) for unipolar memristors or a two-terminal selector device for bipolar memristors (1S1R). The combined device effectively suppresses the leakage currents caused by the unipolar memristor’s reverse bias or bipolar memristor’s low bias, resulting in much lower currents [ 59 – 61 ]. In recent years, prototype chips based on memristor arrays have been widely developed. Figure 5 c–h shows recent studies of memristors arrays, which summarize the structures, the types, the sizes and the realized functions. Fig. 5 Physical diagram based on 1T1R and crossbar memristor arrays. a Schematic diagram of a basic 1T1R array [ 62 ]. Copyright (2023) The Authors. b Schematic diagram of a basic crossbar array [ 63 ]. Copyright (2019) The Authors. c 128 × 64 1T1R array for handwritten digit classification [ 64 ]. Copyright (2018) The Authors. d 32 × 32 1T1R reconfigurable memristor array for analog computing tasks [ 65 ]. Copyright (2022) The Authors. e 2K memristor chips and an FPGA board, which mainly uses memristor arrays to achieve high-precision medical image reconstruction [ 62 ]. Copyright (2023) The Authors. f Schematic diagram of 32 × 32 WO x memristor array realize temporal information processing and handwritten digit recognition [ 66 ]. Copyright (2017) The Authors. g SEM image of a 20 × 20 crossbar array, used for neuromorphic computing with each memristor acting as a synapse [ 67 ]. Copyright (2018) The Authors. h 12 × 12 crossbar memory array composed of self-selective van der Waals heterostructure memory cells [ 63 ]. Copyright (2019) The Authors"
} | 6,040 |
29430725 | null | s2 | 6 | {
"abstract": "Vast potential exists for the development of novel, engineered platforms that manipulate biology for the production of programmed advanced materials. Such systems would possess the autonomous, adaptive, and self-healing characteristics of living organisms, but would be engineered with the goal of assembling bulk materials with designer physicochemical or mechanical properties, across multiple length scales. Early efforts toward such engineered living materials (ELMs) are reviewed here, with an emphasis on engineered bacterial systems, living composite materials which integrate inorganic components, successful examples of large-scale implementation, and production methods. In addition, a conceptual exploration of the fundamental criteria of ELM technology and its future challenges is presented. Cradled within the rich intersection of synthetic biology and self-assembling materials, the development of ELM technologies allows the power of biology to be leveraged to grow complex structures and objects using a palette of bio-nanomaterials."
} | 262 |
24119078 | PMC3815454 | pmc | 7 | {
"abstract": "Spider dragline silk is considered to be the toughest biopolymer on Earth due to an extraordinary combination of strength and elasticity. Moreover, silks are biocompatible and biodegradable protein-based materials. Recent advances in genetic engineering make it possible to produce recombinant silks in heterologous hosts, opening up opportunities for large-scale production of recombinant silks for various biomedical and material science applications. We review the current strategies to produce recombinant spider silks.",
"introduction": "Introduction Spider silks have been a focus of research for almost two decades due to their outstanding mechanical and biophysical properties. Spider silks are remarkable natural polymers that consist of three domains: a repetitive middle core domain that dominates the protein chain, and non-repetitive N-terminal and C-terminal domains. The large core domain is organized in a block copolymer-like arrangement, in which two basic sequences, crystalline [poly(A) or poly(GA)] and less crystalline (GGX or GPGXX) polypeptides alternate. At least seven different types of silk proteins are known for one orb-weaver species of spider (Lewis, 2006a ). Silks differ in primary sequence, physical properties and functions ( Hu et al ., 2006 ). For example, dragline silks used to build frames, radii and lifelines are known for outstanding mechanical properties including strength, toughness and elasticity ( Gosline et al ., 1984 ). On an equal weight basis, spider silk has a higher toughness than steel and Kevlar ( Vepari and Kaplan, 2007 ; Heim et al ., 2009 ). Flageliform silk found in capture spirals has extensibility of up to 500%. Minor ampullate silk, which is found in auxiliary spirals of the orb-web and in prey wrapping, possesses high toughness and strength almost similar to major ampullate silks, but does not supercontract in water. Figure 1 depicts the location and structural elements of MaSp, MiSp and Flag silks. Figure 1 A. An adult female orb weaver spider Nephila clavipes and her web. B. Schematic overview of N. clavipes web composed of three different spider silk proteins and their structures. The coloured boxes indicate the structural motifs in silk proteins. An empty box marked ‘?’ indicates that the secondary structure of the ‘spacer’ region is unknown. Note: MaSp1 or MaSp2: major ampullate spidroin 1 or 2; MiSp1 and 2: minor ampullate spidroin1 and 2; Flag: flagelliform protein. The photo was taken by Olena and Artem Tokarev in the Florida Keys. Finally, there are other silk types such as aciniform, pyriform, aggregate and tubuliform (egg case) with unusual primary structure, composition and properties. Diverse and unique biomechanical properties together with biocompatibility and a slow rate of degradation make spider silks excellent candidates as biomaterials for tissue engineering, guided tissue repair and drug delivery, for cosmetic products (e.g. nail and hair strengthener, skin care products), and industrial materials (e.g. nanowires, nanofibres, surface coatings). Recent advances in genetic engineering have provided a route to produce various types of recombinant spider silks ( Prince et al ., 1995 ; Fahnestock and Bedzyk, 1997 ; Rabotyagova et al ., 2009 ; Xu et al ., 2007 ). However, production of spider silk proteins at a larger scale remains challenging. Moreover, recombinant silk threads do not recapitulate the full potential of native fibres in terms of mechanical properties. Different heterologous host systems have been investigated to develop suitable production systems. In this review, we discuss recent advances in the production of recombinant spider silks in heterologous host systems with the main focus on microbial production. In particular, we focus on dragline silks. Current cloning strategies, expression systems and purification strategies will be discussed to help researchers to engineer customized synthetic spider silk-like proteins for various needs, including biomaterials and material science applications. Structure of silk proteins Spider silks are fascinating polymers, as is the spinning process that members of Araneidae family use to make these exceptional materials. Spiders use complex spinning to rapidly transform water soluble, high molecular weight, silk proteins into solid fibres at ambient temperature and pressure, giving rise to an environmentally safe, biodegradable and high performance material ( Asakura et al ., 2007 ; Lewicka et al ., 2012 ; Teulé et al ., 2012a ). The details on anatomy and physiology of the spider spinning apparatus ( N. clavipes ) can be found elsewhere (Knight and Vollrath, 2001 ; 2002 ; Eisoldt et al ., 2011 ; Rising et al ., 2011 ). In order to understand the challenges and needs associated with biotechnological production of recombinant spider silks, primary protein motifs, composition and secondary structural elements must be discussed. As mentioned earlier, one spider is capable of producing up to seven different types of silks with varying mechanical properties. In spite of different mechanical and physiological properties, the majority of spider silks share a common primary structural pattern comprised of a large central core of repetitive protein domains flanked by non-repetitive N- and C-terminal domains. The most investigated silk is dragline silk, which shows a remarkable combination of strength and elasticity. The golden orb-weaver spider, N. clavipes , produces dragline silk in the major ampullate gland ( Knight and Vollrath, 2001 ). Dragline silk is the protein complex composed of major ampullate dragline silk protein 1 (MaSp1) and major ampullate dragline silk protein 2 (MaSp2). Both silks are approximately 3500 amino acid long. MaSp1 can be found in the fibre core and the periphery, whereas MaSp2 forms clusters in certain core areas. The large central domains of MaSp1 and MaSp2 are organized in block copolymer-like arrangements, in which two basic sequences, crystalline [poly(A) or poly(GA)] and less crystalline (GGX or GPGXX) polypeptides alternate in core domain. The main difference between MaSp1 and MaSp2 is the presence of proline (P) residues accounting for 15% of the total amino acid content in MaSp2 ( Hu et al ., 2006 ), whereas MaSp1 is proline-free. By calculating the number of proline residues in N. clavipes dragline silk, it is possible to estimate the presence of the two proteins in fibres; 81% MaSp1 and 19% MaSp2 ( Brooks et al ., 2005 ). Different spiders have different ratios of MaSp1 and MaSp2. For example, a dragline silk fibre from the orb weaver Argiope aurantia contains 41% MaSp1 and 59% MaSp2 ( Huemmerich et al ., 2004 ). Such changes in the ratios of major ampullate silks can dictate the performance of the silk fibre ( Vollrath and Knight, 1999 ). Specific secondary structures have been assigned to poly(A)/(GA), GGX and GPGXX motifs including β-sheet, 3 10 -helix and β-spiral respectively ( Humenik et al ., 2011 ). The primary sequence, composition and secondary structural elements of the repetitive core domain are responsible for mechanical properties of spider silks; whereas, non-repetitive N- and C-terminal domains are essential for the storage of liquid silk dope in a lumen and fibre formation in a spinning duct ( Ittah et al ., 2006 ). The primary amino acid sequence, composition and secondary structural elements of other silk types are reviewed elsewhere (Lewis, 2006b ; Humenik et al ., 2011 ). Production of recombinant silk proteins Spiders cannot be farmed, in contrast to silkworms, due to their aggressive behaviour and territorial nature ( Kluge et al ., 2008 ). Collecting silk from webs is a time-consuming task. It took 8 years to make a golden spider silk cape from 1.2 million golden orb webs ( Chung et al ., 2012 ). Therefore, biotechnological production of recombinant spider silks is the only practicable solution to harvest silks on a larger scale and to meet growing needs of medicine and biotechnology. A variety of heterologous host systems have been explored to produce different types of recombinant silks ( Table 1 and Table 2 ). Recombinant partial spidroins as well as engineered silks have been cloned and expressed in bacteria ( Escherichia coli ), yeast ( Pichia pastoris ), insects (silkworm larvae), plants (tobacco, soybean, potato, Arabidopsis), mammalian cell lines (BHT/hamster) and transgenic animals (mice, goats). Table 1 Summary of recombinantly expressed spider silks in bacteria and yeasts. Spider silk origin, number of monomers, molecular weight, cloning and expression plasmids as well as restriction enzymes and purification strategies used to produce recombinant silks are shown Type Host Origin Protein Number of monomers MW (KDa) Cloning plasmid RE Expression plasmid Purification Strategy Yield (mg L −1 ) References Bacteria E. coli Nephila clavipes 16, 32, 64, 96 55, 100, 193, 285 pET30a(+) Nhe I/ Spe I pET30a(+) Ammonium sulphate Xia et al ., 2010 6; 15 16, 39.5 pETNX PA 96.8 mg L −1 ; 200 mg L −1 Dams-Kozlowska et al ., 2012 MaSp 1 8; 16 65–163 pBR322 derived Pst I pFP202 (pET9a + pET11a) 300 mg L −1 Fahnestock and Irwin, 1997 8; 16 65–163 pFP202, pFP204, or pFP207 IMAC Fahnestock and Irwin, 1997 16, 24 46, 70 pBSSKII+ AvrII, Nhe1 pET19k N A −1 An et al ., 2011 poly(A) and GGX 1, 2, 3, 6 10, 18 pET30a(+) Spe I/Nde I pET30a(+) 25 mg ml −1 Rabotyagova et al ., 2009 E. coli Nephila clavipes 8; 16 65–163 pBR322 derived Pst I pFP202, pFP204, or pFP207 300 mg L −1 Fahnestock and Irwin, 1997 8, 16, 32 31, 58, 112 pBBSK Sca/Xma/BspEI pET19b 10 mg g −1 Lewis et al ., 1996 Argiope aurantia MaSp2 16 63 IMAC Brooks et al ., 2008 12 71 pET30a(+) N A −1 pET30a(+) N A −1 Brooks et al ., 2008 8 67 Brooks et al ., 2008 E. coli Nephila clavipes Masp1/Masp2 24/16 62/47 pBSSKII+ Xma1/Sca1/BspE1 pET19K IMAC 120 mg L −1 An et al ., 2011 1x-18x 15, 23, 32, 41 pUC18 Spe I/Nde I pQE-9 15, 7, 3, 2 mg L −1 Prince et al ., 1995 E. coli Argiope trifasciata AcSp1 2, 3, 4 19, 38, 51.7, 76.1 pET32 BamHI/ BsgI/ BseRI pET32 IMAC 80 mg L −1 ; 22 mg L −1 Nephila antipodiana TuSp1 11 190 Xma1/Pvu1/BspE1 40 mg L −1 Xu et al ., 2007 Salmonella Araneus diadematus ADF1 1x-3x 30–56 pJ2 HindIII/XbaI pTRC99a_Cm SEC N A −1 Widmaier et al ., 2009 ADF2 1x-3x Widmaier et al ., 2009 ADF3 1x-3x Widmaier et al ., 2009 Yeast Pichia Pastori Nephila clavipes Masp 1 8, 16 65 pBR322 derived Pst I pFP684 Ammonium sulphate 663 mg L −1 Fahnestock and Bedzyk, 1997 Table 2 Summary of recombinantly expressed spider silks in insects, plants and mammalians. Spider silk origin, number of monomers, molecular weight, cloning and expression plasmids as well as restriction enzymes and purification strategies used to produce recombinant silks are shown Type Host Origin Protein Number of monomers MW (KDa) Cloning plasmid RE Expression plasmid Purification strategy Yield (mg L −1 ) References Insects B. mori Nephila clavipes Masp1 2 83 pSLfa1180fa Spe1/Nde1 pBac[3xP3-DsRedaf] IMAC N A −1 Wen et al ., 2010 4 70 pSL1180 pFastBacHT-C 6 mg/larva Zhang et al ., 2008 Masp1+Flag multiple 75–130 pBSSKII+ and pSLfa1180fa Sca1/Xma1/BspE1 pBAC[3xP3-DsRedaf] N A −1 Teulé et al ., 2012b Plants Nicotiana tobaccum Solaum tubercum Masp1 multiple 12.9–99.8 pUC19 NgoMIV/HindIII pRTRA7/3 (NH 4 ) 2 SO 4 10–50% saturation 0.5-2 % total protein Scheller et al ., 2001 Nicotiana benthamiana Nephila clavipes Flag (intein) 4; 10 47, 72, 100, 250 pRTRA15 splicing events pCB301-Kan IMAC 1.8 mg/50 g leaft material 0.34%; 0.03% in leaves, 1.2%; 0.78% in seeds Hauptmann et al ., 2013 Arabdopsis thaliana Glycine max Masp1 8, 16 64, 127 pBSSK+ BglII/BamH1 Cong' + Pha3' ammonium sulphate 1% in somatic embryos Barr et al ., 2004 Mammalians Trangenic mice Nephila clavipes MaSP1 6 31–66 pGEM-5zf BamHI/NcoI pBC1 centrifugation 11.7 mg L −1 Xu et al ., 2007 COS-1 cells Euprosthenops sp. 25, 22 pER1-14 BamH1/EcoRV pSecTag2/Hygro A N A −1 N A −1 Grip et al ., 2006 Baby hamster kidney Nephila clavipes Masp1/ Masp2 N A −1 59, 106/ 59 pBSSK+ ApaI/SapI CMV promoter ammonium sulfate Baby hamster kidney Araneus diadematus ADF3 63, 60, 110, 140 pSecTag-C MscI/PvuII 25–50 mg L −1 Lazaris et al ., 2002 Unicellular organisms as heterologous host systems Unicellular organisms, such as bacteria and yeast, have been investigated as host systems for recombinant silks. A gram-negative, rod-shaped bacterium E. coli is a well-established host for industrial scale production of proteins. Therefore, the majority of recombinant spider silks have been produced in E. coli (Lewis et al ., 2011 ; Fahnestock and Irwin, 1997 ; Wang et al ., 2006 ; Rabotyagova et al ., 2009 ; Rabotyagova et al ., 2010 ; An et al ., 2011 ; An et al ., 2012 ; Teulé et al ., 2012a ). E. coli is easy to manipulate, has a short generation time, is relatively low cost and can be scaled up for larger amounts protein production. The recombinant DNA approach enables the production of recombinant spider silks with programmed sequences, secondary structures, architectures and precise molecular weight ( Rabotyagova et al ., 2011 ). There are four main steps in the process: (i) design and assembly of synthetic silk-like genes into genetic ‘cassettes’, (ii) insertion of this segment into a DNA vector, (iii) transformation of this recombinant DNA molecule into a host cell and (iv) expression and purification of the selected clones. Figure 2 summarizes the recombinant DNA approach used to prepare silk-like proteins. Figure 2 Recombinant DNA approach used to prepare silk-like proteins. The monomeric silk-like gene sequences can be synthesized as short single-stranded oligonucleotides (up to 100 bp) by commercial oligonucleotide synthesis or used directly as polymerase chain reaction products from cDNA libraries. Large repetitive sequences can be constructed by using concatemerization, step-by-step directional approach and recursive ligation ( Fig. 3 ). Concatemerization is a useful method when a library of genes of different sizes is desired but has limitations in the preparation of genes with specific sizes ( Meyer and Chilkoti, 2002 ). To overcome limitations of concatemerization, recursive directional ligation or a step-by-step ligation is employed ( Meyer and Chilkoti, 2002 ; Wright and Conticello, 2002 ). Recursive directional ligation allows for facile modularity, where control over the size of the genetic cassettes is achieved. Moreover, recursive directional ligation eliminates the restriction sites at the junctions between monomeric genetic cassettes without interrupting key gene sequences with additional base pairs that makes it different from the step-by-step ligation approach ( Higashiya et al ., 2007 ). Figure 3 Gene multimerization approaches. Note: RE Site stands for a restriction enzyme site. For example, we have employed step-by-step directional ligation to produce various partial recombinant spider silks as well as engineered silk-like proteins based on the sequences of dragline silk originated from N. clavipes ( Prince et al ., 1995 ; Wang et al ., 2006 ; Huang et al ., 2007 ; Rabotyagova et al ., 2009 ; 2010 ; Mieszawska et al ., 2010 ; Gomes et al ., 2011 ; Numata et al ., 2012 ). As one example, spider silk block copolymers were generated in E. coli (Rabotyagova et al ., 2009 ; 2010 ). In the first cloning step, a commercially available pET30a(+) vector (Novagen, San Diego, CA, USA) was modified with an adaptor sequence, carrying NheI and SpeI restriction sites. The adaptor was inserted into XhoI and NcoI sites of a pET30a(+) to generate pET30L. The coding sequences of two spider silk-like monomers A (hydrophobic block) and B (hydrophilic) were designed to carry SpeI and NheI restriction sites at the ends of the sequences. This allowed ligation of the domains into a pET30L vector. By using a step-by-step directional ligation approach, direct control over the assembly of monomeric genes into complex sequences was achieved. Six different constructs were cloned and transformed into the bacterial host for expression. An N-terminal His-tag was used for protein purification by immobilized metal affinity chromatography ( Rabotyagova et al ., 2009 ). Another genetic engineered strategy has been proposed by Lewis Laboratory to assemble long repetitive spider silk genes ( Teule et al ., 2009 ). This cloning strategy employs a one-step head-to-tail ligation that can produce large inserts in precise manner ( Lewis et al ., 1996 ; Brooks et al ., 2008 ; Teule et al ., 2009 ; Teulé et al ., 2012a ). The spider silk synthetic genes were optimized for codon usage in E. coli and were cloned into a plasmid vector pBluescriptII SK(+) (Stratagene). Each silk module was carrying compatible XmaI and BspEI restriction sites at the ends on the coding sequences. The vector also contained a unique restriction site ( ScaI ) in the ampicillin resistance gene. By simultaneously performing two double digestion reactions ScaI – XmaI and ScaI – BspEI two fragments each containing a copy of a silk monomer gene were obtained. The fragments were ligated together using T4 ligase resulting in the doubling of the size of silk genes and restoring the ampicillin resistance of the plasmid ( Fig. 4 ). Several round of cloning were performed to obtain repetitive sequences of a desired size. Next, the multimeric synthetic genes were subcloned into an expression pET19b vector using NdeI and BamHI restriction sites. Since the expression vector was carrying NdeI and BamHI sites, the liberated inserts were cloned in-frame with pET19b. Similar to pET30L, silk genes in pET19b are under control of the T7 promoter and require the addition of isopropyl-β-D-1-thiogalactopyranoside to initiate protein expression. The expressed proteins can be purified by immobilized metal affinity chromatography (IMAC) due to the presence of an N-terminal His-tag. Several recombinant spider silk proteins from different species were produced using this genetic engineering strategy including silks from N. clavipes ( Teule et al ., 2009 ) Argiope aurantia ( Brooks et al ., 2008 ). Recombinant spider silk proteins from Nephylengys cruentata , Parawixia bistriata and Avicularia juruensis were produced employing this cloning strategy ( Leopoldo et al ., 2007 ) (US patent 20 100 311 645). Figure 4 summarizes the strategy. Figure 4 Cloning strategy used by the Lewis group to engineer long repetitive spider silk sequences (in green). A. Cloning of a silk monomer into the vector pBluescript II SK+. B. The resulting plasmid is double digested and fragments containing silk monomers are ligated again to produce longer sequences. C. The synthetic spider silk multimer is ligated into pET19b expression vector. Note: Restriction digestion sites are indicated by star. Adapted from reference (Teule et al ., 2009 ). A three module cloning strategy based on the sequences of ADF-3 and ADF-4 was developed by Scheibel research group ( Huemmerich et al ., 2004 ), designed so that multiple modules can be combined. Moreover, additional coding sequences such as N- or C-terminal domains can be added if needed. The purification protocol is based on heat resistance of silk proteins followed by an ammonium sulphate precipitation that is different from Ni-NTA IMAC. Different purification strategies have been employed recently to optimize small and large-scale production of recombinant silks. Most of the spider silk proteins are produced with an N- or C-terminal His-tags to make purification simple and produce enough amounts of the protein. However, the presence of this tag can affect protein secondary structure and interfere with the process of spider silk fibre formation. Dams-Kozlowska et al . (2012 ) proposed two strategies to purify spider silks from lysates without the use of a His-tag. These protocols are based on thermal treatment and organic acid resistance of silk proteins and do not require the presence of the His-tag. After purification, silk proteins based on MaSp1 gene sequence were formed into films that subsequently were used to grow murine fibroblast cell culture. The results demonstrated that silk films were non-toxic to the cells ( Dams-Kozlowska et al ., 2012 ). Because of the highly repetitive core sequence of spider silk genes, frequent homologous recombination, deletions, transcription errors, translation pauses, accumulation in inclusion bodies and low yields were observed during the production of recombinant silks in E. coli . Moreover, when the protein size was increased from 43 kDa to higher (the size of native spidroins is between 300 and 350 kDa), protein yields decreased dramatically. Codon optimization for the specific host expression system helped maximize the translation of the foreign gene transcripts and thus, improved protein yields ( Fahnestock and Bedzyk, 1997 , Lewis, 2006b ). It was also suggested that depletion of tRNA pools upon protein expression resulted in transcription and translation errors ( Rosenberg et al ., 1993 ). Recently, Xia et al . (2010 ) employed a metabolic engineered strategy to enhance the production of recombinant spider silks. The authors reported production of full length (284.9 kDa) recombinant N. clavipes dragline silk proteins that were rich in glycine (43–45%). Production of these silk proteins was enhanced by the use of the metabolically engineered expression host within which the glycyl-tRNA pool was elevated. The fibres spun with the native-sized recombinant spider silk protein showed tenacity, elongation and Young's modulus of 508 MPa, 15% and 21 GPa, respectively, comparable to those of native spider dragline silk ( Xia et al ., 2010 ). Through extensive proteomic analysis, serine hydroxymethyltransferase (GlyA) and β-subunit of glycly-tRNA synthetase (GlyS) were found to be upregulated to meet the high cellular demand for glycly-tRNA when expressing glycine-rich silk proteins. Increased glycine biosynthetic flux by overexpressing glycyl-tRNA synthetase elevated the total tRNAGly pool and resulted in enhanced production of high molecular weight recombinant spider silks. Recently, large spider recombinant egg case silk protein from Nephila antipodiana , 378 kDa, was engineered using E. coli , where gene multimers were chemically linked by cysteine disulfide bonds. The recombinant silk sequence consisted of two silk proteins: tubuliform spidroin 1 (TuSp1) and C-terminal domain of MisP1. Non-repetitive C-terminal domain of MiSp1 was chosen due to its higher water solubility and stability compared with the C-terminal domain of TuSp1. A disulfide linkage between two C-terminal domains was formed by introducing a point mutation (S76 to S76C). This link allowed the formation of a hybrid DNA construct that was expressed in E. coli (DE3). The recombinant protein was expressed in E. coli . Moreover, the artificial fibres spun from this protein showed higher tensile strength and Young' modulus than natural egg case protein ( Lin et al ., 2013 ). The highly repetitive silk gene arrangement and the unusual mRNA secondary structure result in inefficient translation that limits the size of the silks produced in E. coli . To minimize the presence of truncated silk proteins and allow the extracellular secretion of silks, the mythylotropic yeast P. pastoris has been used. Fahnestock and Bedzyk (1997 ) produced N. clavipes spider dragline silks in yeast P. pastoris . Synthetic genes were expressed at high levels under control of the methanol-inducible AOX1 promoter. Transformants containing multiple gene copies produced elevated levels of silk protein. Results demonstrated that P. pastoris can be used to successfully produced produce long repetitive proteins ( Fahnestock and Bedzyk, 1997 ). Spider silks from Araneus diadematus (ADF-1, 2 and 3) have also been expressed using the type III secretion system of a gram-negative, non-spore-forming, enterobacterium Salmonella . The authors reported yield values range from 90 to 410 nmol L −1 h −1 that is similar to 10 mg L −1 h −1 for a protein the size of ADF-2. The results demonstrated the feasibility to use Salmonella for the large-scale spider silk production ( Widmaier et al ., 2009 ). Mammalian cell lines, such as bovine mammary epithelial alveolar and baby hamster kidney cells, were used to express MaSp1 and MaSp2 ( Lazaris et al ., 2002 ). The cells expressed recombinant proteins; however, as size of silk gene increased, the yield decreased dramatically due to inability of mammalian cells to cope with large repetitive sequences. Several factors have attributed to the decreased yields including, but not limited to, inefficient transcription, insufficient secretion, low copy numbers and translational limitations. The produced silk proteins were spun into fibres, and their mechanical properties were tested. It was noted that those recombinant silks that were produced without a His-tag demonstrated better mechanical properties compared with fibres made of silk proteins with a His-tag (i.e. fibres were brittle). Similar problems (i.e. transcription and translation limitations) have been reported when green monkey kidney fibroblast-like cell lines (COS-1) were used to express a 636-base pair gene fragment of MaSp1 from the African spider Euprosthenops sp . ( Grip et al ., 2006 ). Table 1 summarizes genetic engineering approaches, cloning strategies, and production yields of recombinant silk proteins produced in unicellular heterologous host systems. Multicellular organisms as heterologous host systems Due to the low production rate and instability (i.e. frequent homologous recombination, deletions, transcription errors, translation pauses) of spider silk repetitive genes in unicellular organisms, multicellular organisms such as insects, plants and mammals have been studied for production of recombinant spider silk proteins. Silkworms ( B. mori ) can be farmed and produce cocoons containing large quantities of silkworm silk known as fibroin ( Vepari and Kaplan, 2007 ; Hu and Kaplan, 2011 ). Moreover, to produce a solid thread, silkworms employ a spinning process that is similar to that used by spiders to make dragline silk. The presence of a natural silk production system in silkworms makes them excellent candidates to investigate as heterologous hosts for spider silk production. There have been several reports of the transfer of silk genes from spiders to silkworms ( Motohashi et al ., 2005 ; Zhang et al ., 2011 ; Teulé et al ., 2012b ). Baculovirus-based expression systems have been used to introduce silk genes into a heterologous host. Baculovirus infects silkworms and allows for production of large quantities of heterologous proteins in a short period of time ( Motohashi et al ., 2005 ). Using this expression system, MaSp1 from N. clavipies linked with an enhanced green fluorescent protein (EGFP) fusion protein was cloned and expressed in the B. mori cell line (BmN) and larvae ( Zhang et al ., 2008 ). The authors reported successful production of a recombinant EGFP-MaSp1 fusion protein in both systems. In the silkworm larvae, a total of 6 mg of fusion protein was expressed, whereas in the BmN cells, 5% of the cell total protein was occupied by this recombinant silk. The major limitations of this expression system were low solubility of silk proteins and inability to assemble spider silk fibres. It was shown that more than 60% of the fusion proteins formed aggregates via self-assembly. To overcome solubility issues, MaSp1 C-terminal domain is to be incorporated due to its role to prevent aggregate formation. To produce fibres, germline-transgenic silkworms ( B. mori ) were produced by injecting silkworm eggs with a piggyBac transformation vector carrying MaSp1 sequence ( Wen et al ., 2010 ). The insects were capable of spinning fibres and forming cocoons containing recombinant spider silk. However, the mechanical properties of the fibres were lower than dragline MaSp1 silk due to the low ratio of MaSp1 in the total silk protein. In a recent effort to develop tough fibres, transgenic silkworms encoding chimeric silkworm/spider silk proteins were produced using piggyBac vectors (Teulé et al ., 2012b ). The vector, used previously by the Tamada group ( Kojima et al ., 2007 ) included the B. mori fibroin heavy chain promoter and enhancer, a genetic sequencing encoding a 78 kDa synthetic spider silk protein, and an EGFP tag. Strong EGFP signals were observed by fluorescence ( Fig. 5 ). The composite fibres were tougher than the parental silkworm silk fibres and as tough as native dragline spider silk fibres. Figure 5 Expression of the chimeric silkworm/spider silk/EGFP protein in (A) cocoons, (B and C) silk glands and (D) silk fibres from spider 6-GFP silkworms. Reproduced with permission from (Teulé et al ., 2012b ). These results demonstrate that silkworms can be engineered to generate composite silk fibres containing stably integrated spider silk protein sequences, which significantly improved overall mechanical properties. Transgenic plants have also been investigated as heterologous host systems to produce recombinant spider silks. Advances in genetic engineering technology and transformation methods make it possible to produce non-plant proteins in plants ( Yang et al ., 2005 ; Rech et al ., 2008 ). Moreover, one plant offers several different expression systems, such as seeds, leaves, tubers and roots with potential for organelle-specific accumulation of recombinant proteins ( Scheller and Conrad, 2005 ). Stable transgenic tobacco and potato lines were engineered to express MaSp1 genes from N. clavipes ranging from 420 to 3600 bp ( Scheller et al ., 2001 ). Recombinant spider silk proteins were found in the endoplasmic reticulum (ER) of tobacco and potato leaves at the accumulation of 2% of total soluble protein. Moreover, the production levels were independent of the size of silk genes. Purification was performed using high temperature treatment followed by acidification and ammonium sulphate precipitation. Additionally, recombinant MaSp1-like proteins were also produced in the leaves and seeds of Arabidopsis (small flowering plants related to cabbage) as well as in somatic soybean embryos ( Barr et al ., 2004 ). The expression of recombinant silks was driven by the 35S promoter in leaves and the β-conglycinin α' subunit promoter in seeds and somatic soybean embryos. The results demonstrated that recombinant spider silk proteins had higher accumulation levels in seeds than in the leaves. Recently, a native-sized FLAG protein from N. clavipes was cloned and expressed in the ER of tobacco plant ( Nicotiana benthamiana ) leaf cells using an intein-based posttranslational protein fusion technology ( Hauptmann et al ., 2013 ). This method avoids the need for highly repetitive transgenes resulting in a higher genetic and transcriptional stability. Additional details on production of fibrous proteins in plants can be found elsewhere ( Scheller and Conrad, 2005 ). Transgenic production of recombinant silk proteins in mammary glands and secretion of them into milk has been investigated in mice and goats ( Williams, 2003 ; Xu et al ., 2007 ). In case of transgenic mice production, MaSp1 and MaSp2 synthetic genes (40 and 55 kDa) were synthesized and cloned into the pBC1 expression vector (Invitrogen, Carlsbad, CA, USA) together with a goat β-casein signal sequence. The chimeric gene construct was microinjected into pronuclei of fertilized eggs of Kunming white mice ( Xu et al ., 2007 ). Southern blot analysis was used to identify mice containing transgene construct as well as a copy number of transgene. The expression of dragline silk in milk was confirmed by Northern blot followed by Western blot analysis. The results revealed that transgenic mice were capable of expressing recombinant silk proteins in their milk. Genetically engineered (transgenic) goats capable of expressing spider silk proteins based on the sequences of MaSp1 and MaSp 2 were produced by Nexia Biotechnologies, and later by the Lewis group ( Lazaris et al ., 2002 ; Service, 2002 ). Silk protein expression was controlled by the β-casein promoter and was expressed in the milk of transgenic goats. Silk proteins were observed only in mammary tissues as confirmed by Western blot ( Steinkraus et al ., 2012 ). Maximum yields observed for the recombinant silk production in transgenic animals were low (11.7 mg l −1 ) when compared with bacterial expression ( Table 1 and Table 2 ). Today, the large-scale production of recombinant silk proteins from transgenic animals is relatively expensive and challenging in terms of animal breeding."
} | 8,217 |
30944413 | null | s2 | 8 | {
"abstract": "Quorum sensing is a process of bacterial cell-to-cell chemical communication that relies on the production, detection and response to extracellular signalling molecules called autoinducers. Quorum sensing allows groups of bacteria to synchronously alter behaviour in response to changes in the population density and species composition of the vicinal community. Quorum-sensing-mediated communication is now understood to be the norm in the bacterial world. Elegant research has defined quorum-sensing components and their interactions, for the most part, under ideal and highly controlled conditions. Indeed, these seminal studies laid the foundations for the field. In this Review, we highlight new findings concerning how bacteria deploy quorum sensing in realistic scenarios that mimic nature. We focus on how quorums are detected and how quorum sensing controls group behaviours in complex and dynamically changing environments such as multi-species bacterial communities, in the presence of flow, in 3D non-uniform biofilms and in hosts during infection."
} | 265 |
36505976 | PMC9720699 | pmc | 9 | {
"abstract": "Abstract Spider silk is the toughest fiber found in nature, and bulk production of artificial spider silk that matches its mechanical properties remains elusive. Development of miniature spider silk proteins (mini‐spidroins) has made large‐scale fiber production economically feasible, but the fibers’ mechanical properties are inferior to native silk. The spider silk fiber's tensile strength is conferred by poly‐alanine stretches that are zipped together by tight side chain packing in β‐sheet crystals. Spidroins are secreted so they must be void of long stretches of hydrophobic residues, since such segments get inserted into the endoplasmic reticulum membrane. At the same time, hydrophobic residues have high β‐strand propensity and can mediate tight inter‐β‐sheet interactions, features that are attractive for generation of strong artificial silks. Protein production in prokaryotes can circumvent biological laws that spiders, being eukaryotic organisms, must obey, and the authors thus design mini‐spidroins that are predicted to more avidly form stronger β‐sheets than the wildtype protein. Biomimetic spinning of the engineered mini‐spidroins indeed results in fibers with increased tensile strength and two fiber types display toughness equal to native dragline silks. Bioreactor expression and purification result in a protein yield of ≈9 g L −1 which is in line with requirements for economically feasible bulk scale production.",
"conclusion": "3 Conclusion Using biological principles, we employed protein engineering to design mini‐spidroins with predicted increased β‐sheet propensities and increased inter‐β‐sheet binding strengths. Prokaryotic expression, protein purification, and biomimetic fiber spinning resulted in four different types of fibers with significantly improved tensile strength compared to the original mini‐spidroin. Using this strategy, we successfully produced the first biomimetic fibers with toughness values matching those of native dragline silk fibers. Finally, we show that these fibers can be produced at very high yields in bioreactors, vouching for feasible large‐scale production.",
"introduction": "1 Introduction Spider silk is nature's high‐performance fiber. Its unique combination of high tensile strength and extensibility results in an unsurpassed toughness which makes it very attractive for many industrial applications. [ \n \n 1 \n , \n 2 \n , \n 3 \n \n ] Due to limited availability of the natural material, large scale production must involve the expression of the silk proteins (spidroins) in heterologous hosts. [ \n \n 4 \n \n ] \n Spiders have up to seven different types of silk glands in which the spidroins are being produced, stored, and processed. [ \n \n 5 \n \n ] The major ampullate gland makes the strongest silk, which is used in the dragline and for making the framework of the web. [ \n \n 6 \n , \n 7 \n , \n 8 \n , \n 9 \n , \n 10 \n , \n 11 \n \n ] The spidroins are synthesized by epithelial cells lining the major ampullate gland and are stored in the gland lumen as a highly concentrated dope. [ \n \n 9 \n , \n 12 \n , \n 13 \n \n ] Changes in the microenvironment along the gland, [ \n \n 14 \n \n ] for example, ion exchange, drop in pH from 8.0 to at least 5.7, [ \n \n 15 \n \n ] increased shear forces, [ \n \n 16 \n \n ] and dehydration [ \n \n 7 \n \n ] lead to conformational transitions of the spidroins and fiber formation. [ \n \n 15 \n , \n 17 \n , \n 18 \n , \n 19 \n , \n 20 \n \n ] \n Spidroins are composed of an N‐terminal domain (NT), [ \n \n 21 \n \n ] a repetitive region that often is extensive [ \n \n 22 \n \n ] and a C‐terminal domain (CT). [ \n \n 18 \n \n ] The terminal domains are crucial for solubility of the spidroins during storage and regulate the assembly of the spidroins into a solid fiber. [ \n \n 17 \n , \n 18 \n , \n 19 \n , \n 20 \n , \n 23 \n \n ] The repetitive region of most major ampullate spidroins (MaSps) contain up to 100 tandem repeats of poly‐Ala blocks and Gly‐rich motifs. [ \n \n 22 \n , \n 24 \n \n ] In the soluble dope, the spidroins are mostly in random coil and helical conformations, [ \n \n 25 \n , \n 26 \n , \n 27 \n , \n 28 \n , \n 29 \n \n ] whereas the solid silk fiber contains nanosized crystals made up by stacked antiparallel β‐sheets embedded in amorphous structures. [ \n \n 30 \n , \n 31 \n , \n 32 \n , \n 33 \n , \n 34 \n \n ] This heterogeneous structure of the silk fiber is important as the β‐sheet crystals confer the strength while the amorphous structures confer the extensibility to the fiber. [ \n \n 10 \n , \n 35 \n , \n 36 \n \n ] The amorphous matrix, containing β‐turns and ordered structures with conformational similarities to collagen and poly‐proline helices, are dominated by the glycine‐rich regions. The β‐sheets, formed by the poly‐Ala blocks, orient with the β‐strands parallel to the fiber axis, [ \n \n 37 \n , \n 38 \n , \n 39 \n , \n 40 \n \n ] and the Ala side chain of a given β‐strand fill the space close to an α‐carbon in a neighboring β‐stand, analogous to a tightly packed steric zipper. [ \n \n 41 \n , \n 42 \n , \n 43 \n \n ] \n There are two main strategies for producing artificial silk fibers; one being expression of insoluble spidroins with subsequent solubilization and fiber processing using organic solvents, [ \n \n 44 \n , \n 45 \n , \n 46 \n , \n 47 \n , \n 48 \n , \n 49 \n \n ] and another being a biomimetic approach involving only aqueous solutions throughout the purification and spinning procedures and in which the molecular mechanisms and triggers for fiber formation are replicated. [ \n \n 50 \n , \n 51 \n , \n 52 \n , \n 53 \n \n ] The first approach enables expression of large spidroins that can be spun into fibers with high tensile strength, but the protein yields are far from what is required for industrial production. [ \n \n 54 \n , \n 55 \n \n ] Using the second approach, mini‐spidroins composed of an NT, a short repeat region consisting of two poly‐Ala/Gly‐rich blocks and a CT, have been developed. Such mini‐spidroins are extremely water‐soluble and can be spun into fibers using biomimetic spinning set‐ups. [ \n \n 51 \n , \n 52 \n , \n 53 \n , \n 56 \n \n ] Moreover, one of these mini‐spidroins, NT2RepCT, can be produced at a yield of 14.5 g L −1 in bioreactor cultivations which vouch for economically feasible bulk production. [ \n \n 55 \n , \n 56 \n \n ] Fibers spun from NT2RepCT are superior compared to previously published as‐spun fibers, but still, the fibers only reach about 15% of the native silk fiber's tensile strength. [ \n \n 1 \n , \n 51 \n \n ] NMR spectroscopy revealed that the mini‐spidroin's two poly‐Ala blocks are in an α‐helical conformation in the soluble state and convert to β‐sheet conformation in the as‐spun wet fiber, as expected. However, upon drying the fiber, the poly‐Ala blocks are transitioning back to α‐helical conformation, [ \n \n 57 \n \n ] which could lead to the inferior mechanical properties of dried NT2RepCT fibers compared to the native silk fiber. We therefore hypothesize that the mechanical properties of recombinant fibers could be improved by increasing the β‐strand propensity and inter‐β‐sheet interactions of the poly‐Ala blocks, [ \n \n 58 \n \n ] as it has been suggested by replacing the poly‐alanines with amyloidogenic sequences. [ \n \n 59 \n \n ] \n Notably, Ala residues have a low propensity to form β‐strands, whereas more hydrophobic residues like Val, Cys, Ile, and Phe show a higher β‐strand propensity, [ \n \n 60 \n \n ] and thus could be considered better candidates for forming stable β‐sheets in the silk fiber. However, being secretory proteins, the spidroins need to pass through the translocon when produced by the gland epithelium. [ \n \n 61 \n \n ] If the nascent polypeptide chain contains segments that are rich in Val, Ile, Cys, or Phe the translocon will mediate insertion into the endoplasmic reticulum membrane, [ \n \n 62 \n , \n 63 \n \n ] and thus any spidroin segment rich in these amino acid residues would be trapped in the cell. In fact, Ala is the most hydrophobic residue that allows passage through the translocon, which suggests that the spidroins have evolved to optimize hydrophobicity in their β‐sheet forming segments to the extent possible for a secretory protein. [ \n \n 58 \n , \n 60 \n \n ] Intracellular expression in prokaryotes will bypass the restrictions imposed by the secretory pathway that native spidroins must adhere to since translation and accumulation of the target protein takes place in the cytosol. These fundamental biological principles led us to use rational design and protein engineering to generate mini‐spidroins that potentially can be produced at high yields in prokaryotic hosts and be used to generate stronger biomimetic artificial spider silk fibers ( Figure \n 1 A,B ). The Zipper database [ \n \n 64 \n \n ] was used to screen a large panel of mini‐spidroins with designed modifications of the poly‐Ala blocks and candidates with low Rosetta energies were chosen for heterologous expression. Soluble target proteins were identified, characterized biochemically, and spun into fibers using a biomimetic spinning device. The mechanical performance of the fibers reveals that engineering of the repeat domain of mini‐spidroins is possible and can result in fibers with increased tensile strength. Figure 1 Schematic representation of the designed constructs. A) NT2RepCT (A 15 ‐A 14 ) is composed of an N‐terminal domain (NT, red; PDB: 4FBS), a repeat region with two poly‐Ala blocks (green and yellow), and a C‐terminal domain (CT, blue, PDB 3LR2). Both subunits of the soluble NT2RepCT dimer are shown (one is shaded). B) Protein sequence alignment of the repetitive region from A 15 ‐A 14 and engineered constructs thereof. Note that all constructs contain NT, a repeat part, and CT. Substitutions in the poly‐Ala blocks are indicated in orange.",
"discussion": "2 Results and Discussion Based on the β‐strand/α‐helix propensity ratios of amino acid residues as well as their hydrophobicity, Ile and Val were chosen to design 13 different constructs with substitutions in the poly‐Ala blocks of the original NT2RepCT sequence (referred to as A 15 ‐A 14 to reflect the composition of the two poly‐Ala blocks), (Figure 1 ). Additionally, the less hydrophobic residue Thr was used since it is branched at the β‐carbon and hence favors β‐strand conformation. [ \n \n 65 \n , \n 66 \n \n ] \n Figure 1B shows the amino acid sequences of the repetitive regions from A 15 ‐A 14 and engineered constructs with substitutions indicated (complete sequences can be found in Table S1 , Supporting Information). Substitutions were mainly introduced at every second position resulting in β‐strands with mutated side chains on the same side. Mutations were introduced in either both (e.g., (AV) 7 ‐(AV) 7 ) or only in one of the poly‐Ala blocks (e.g. (AV) 7 ‐A 14 ). The number of substitutions varied between 15 (e.g., V 15 ‐A 14 , in which all Ala are replaced by Val in the first poly‐Ala block) and 3 as in, for example, (A 3 V) 3 ‐(A 14 ), which contains Val substitution at every fourth position in the first poly‐Ala block. A few additional constructs were designed to analyze the impact of the position of the substituted residues, for example, (A 3 I) 3 ‐A 14 , A 15 ‐(A 3 I) 3 and IA 6 IA 6 I‐A 14 that all have three Ile substitutions but in different locations. The packing of β‐sheets in amyloid‐like fibrils involve steric zippers, [ \n \n 41 \n , \n 67 \n \n ] which are also found in spider silk β‐sheet crystals. [ \n \n 36 \n , \n 43 \n \n ] Steric zippers are formed by tightly bound β‐strands with high complementarity of the involved side chains. [ \n \n 41 \n , \n 67 \n \n ] The Zipper database predicts the stability and propensity of hexapeptides in a given amino acid sequence to form steric zippers by calculating the energies of the interstrand interactions. Rosetta energies equal or below −23 kcal mol −1 suggest a high propensity to form steric zippers. [ \n \n 64 \n \n ] \n \n Figure \n 2 A shows the Rosetta energies estimated for constructs A 15 ‐A 14 and (A 3 I) 3 ‐A 14 (corresponding profiles for all engineered mini‐spidroins are shown in Figure S1 , Supporting Information, and summarized in Table S2 , Supporting Information, and Figure 2B ). As expected, the hexapeptides in the poly‐Ala region of the A 15 ‐A 14 construct have low Rosetta energies (−24.6 kcal mol −1 ) and thus should be able to form steric zippers (Figure 2C ). All designed constructs contain at least one hexapeptide with a Rosetta energy lower than that of A 15 ‐A 14 (Table S2 , Supporting Information), ranging from −24.9 to −29.4 kcal mol −1 (for (AT) 7 ‐(AT) 7 and V 15 ‐A 14 , respectively). Generally, the effect on the Rosetta energies increased with an increasing number of hydrophobic replacements in the poly‐Ala region. Figure 2 Rosetta energy profiles of A) A 15 ‐A 14 and (A 3 I) 3 ‐A 14 (profiles for all designed proteins are found in Figure S1 and Table S2 , Supporting Information). Bars show Rosetta energies for moving hexapeptides (indicated at the first residue of each hexapeptide), red bars indicate Rosetta energies equal or below −23 kcal mol −1 (dashed line). Green bars indicate Rosetta energies above the threshold and are unlikely to form steric zippers ( https://services.mbi.ucla.edu/zipperdb/ ). [ \n \n 64 \n \n ] B) Bars indicate the Rosetta energy of the hexapeptide with the lowest predicted energy from A 15 ‐A 14 and the engineered mini‐spidroins (all hexapeptides are shown in Table S2 , Supporting Information). C) Hypothetical zipper structure of two β‐sheets composed of hexapeptides AAAAAA from A 15 ‐A 14 and AIAAI derived from (A 3 I) 3 ‐A 14 , respectively. Of the 15 designed proteins, seven were overexpressed and six were highly overexpressed in E. coli BL21 cells ( Table \n 1 \n and Figures S2 and S3 , Supporting Information). Constructs with Val substitutions had lower expression levels than corresponding constructs with Ile substitutions, but the number of substitution and the hydrophobicity did not have any general impact on expression levels (Figure S4 , Supporting Information). The (AT) 7 ‐(AT) 7 construct did not express well which could be due to that this repeat was designed to resemble a “CAT tail” which is known to lead to aggregation of the nascent polypeptide chain and to degradation by the proteasome. [ \n \n 68 \n \n ] \n Table 1 Summary of number of substitutions, expression levels, solubility after cell lysis, protein yield, and spinnability into fibers of the engineered proteins. Expression levels, solubility after cell lysis, and spinnability into fibers are rated from very high (+++), intermediate (++), low (+), and not at all (0). Rating of expression level and solubility after cell lysis were estimated by appearance of the target band on SDS‐PAGE (Figures S2 and S3 , Supporting Information). (−) indicates not tested. ( 1 ) indicates degradation during expression. (*) marks purification using gravity columns instead of FPLC Construct Number of substitutions Expression levels Solubility after cell lysis Average protein yield [mg L −1 culture] Spinnability into fibers 1. A 15 ‐A 14 \n 0 +++ +++ 250 +++ 2. (AT) 7 ‐(AT) 7 \n 14 + − − − 3. (A 3 T) 3 ‐(A 3 T) 3 \n 6 ++ +++ 58* +++ 4. (AV) 7 ‐(AV) 7 \n 14 +++ 0 − − 5. (AV) 7 ‐A 14 \n 7 +++ 0 − − 6. V 15 ‐A 14 \n 15 +/++ 1 \n 0 − − 7. (A 3 V) 3 ‐(A 3 V) 3 \n 6 +++ +++ 139* +++ 8. (A 3 V) 3 ‐A 14 \n 3 ++ +++ 216 +++ 9. (AI) 7 ‐(AI) 7 \n 14 + + 4* − 10. A 15 ‐(AI) 7 \n 7 + + − − 11. (AIA 2 ) 3 ‐(AIA 2 ) 3 \n 8 ++ + − − 12. (A 3 I) 3 ‐(A 3 I) 3 \n 6 +++ +++ 94* + 13. (A 3 I) 3 ‐A 14 \n 3 +++ +++ 207 +++ 14. A 15 ‐(A 3 I) 3 \n 3 +++ +++ 233 +++ 15. (A 2 I) 4 ‐A 14 \n 4 ++ +++ 243 +++ 16. IA 6 IA 6 I‐A 14 \n 3 ++ ++ 139 − John Wiley & Sons, Ltd. In addition to A 15 ‐A 14 , seven of the constructs were found mainly in the soluble fraction after cell lysis in 20 m m Tris‐HCl, and four constructs were in both the soluble and insoluble fraction (Table 1 and Figure S3 , Supporting Information). Increased hydrophobicity, number of substitutions, and lower Rosetta energies correlated with lower solubility after cell lysis (Figure S4 , Supporting Information). 9 of the 15 designed constructs plus the control A 15 ‐A 14 yielded sufficient soluble protein for purification. Nondenaturing immobilized metal affinity chromatography yielded between 4 and 243 mg of pure target protein per 1 L shake flask culture (average of 10 × 1 L cultures). Notably, six of the engineered mini‐spidroins gave very high yields (>100 mg L −1 Table 1 ). (AV) 7 ‐(AV) 7 , (AV) 7 ‐A 14 , and V 15 ‐A 14 expressed well but were insoluble after lysis, likely due to high hydrophobicity of the engineered segments. Expression and purification of the A 15 ‐(AI) 7 and (AIA 2 ) 3 ‐(AIA 2 ) 3 constructs did not result in enough soluble protein for further characterization. The constructs that showed intermediate to high expression levels but were insoluble after cell lysis were treated with 8 m urea but could not be solubilized to the extent needed for enabling purification of enough protein for fiber spinning (not shown). The position of the Ile replacements within one Ala block had an impact on the protein yield but whether these were located in the first or second poly‐Ala block did not matter. For example, (A 3 I) 3 ‐A 14 and A 15 ‐(A 3 I) 3 both have three Ile substitutions in the first and second poly‐Ala block, respectively, and show comparable yields. In contrast, (A 3 I) 3 ‐A 14 and IA 6 IA 6 I‐A 14 have the same number of Ile replacements in the first block, but their location differ as does the yield (207 vs 139 mg L −1 culture for (A 3 I) 3 ‐A 14 and IA 6 IA 6 I‐A 14 , respectively). Next, we investigated the secondary structure content and the thermal stability of the purified constructs by circular dichroism (CD) spectroscopy ( Figure \n 3 \n ). We found that all constructs had an overall α‐helical secondary structure (Figure 3A ) which indicates that the amino acid substitutions did not affect the secondary structure of the soluble proteins to any large extent. Heating to 90 °C led to a decreased signal for all constructs and concomitant transition to β‐sheet dominated secondary structures (Figure 3C ). The heat‐induced conformational changes were irreversible upon cooling of the samples (Figure 3D ). Melting curves for all constructs showed that the proteins unfolded around 46–50 °C, which is in line with reports on the isolated terminal domains, [ \n \n 15 \n \n ] and means that the substitutions in the repetitive region of the mini‐spidroins only had a minor effect on the thermal stability of the proteins (Figure 3B ). Figure 3 CD spectroscopy of purified engineered mini‐spidroins. A) Initial spectra at 20 °C and B) molar ellipticity measured at 222 nm from 20 to 90 °C was converted to fraction natively folded (%) and then normalized. CD spectroscopy of different constructs C) heated to 90 °C and D) after cooling to 20 °C. Out of the nine engineered mini‐spidroins that were successfully purified (excluding A 15 ‐A 14 ), eight could be concentrated to at least 200 mg mL −1 to generate spinning dopes, while (AI) 7 ‐(AI) 7 yielded too little protein (Table 1 ). The dopes made from the eight constructs were transferred to syringes and extruded through a thin glass capillary into a low pH aqueous buffer according to a previously described biomimetic spinning procedure. [ \n \n 50 \n , \n 51 \n \n ] Seven engineered mini‐spidroins could be spun into fibers, and only the IA 6 IA 6 I‐A 14 protein aggregated prematurely in the syringe. One of the mini‐spidroins, (A 3 I) 3 ‐(A 3 I) 3 , formed fibers that were too fragile to be retrieved. The reason for the poor integrity of the (A 3 I) 3 ‐(A 3 I) 3 fibers is not known but was not related to premature aggregation in the dope. The other six engineered fiber types, plus the A 15 ‐A 14 fibers, were successfully collected onto a motorized wheel at the end of the spinning bath ( Figure \n 4 A and Video S1 , Supporting Information). There was no difference in the appearance of the spun fibers (Figure 4B ) and the diameter of the different fiber types, determined by light microscopy, varied between 4 and 19 µm (Figure S5H and Table S3 , Supporting Information). The reason for the differences in diameter between the different fiber types is not known but is likely linked to differences in the properties of the proteins since the spinning conditions were kept constant. Figure 4 Mechanical properties of spinnable engineered mini‐spidroins in comparison with A 15 ‐A 14 . A) Photographs of the biomimetic spinning set‐up; a video of the spinning can be found in Video S1 , Supporting Information. B) Photographs of spun fibers. C) Strength, D) strain at break, E) toughness modulus, dashed line indicates toughness modulus of a native dragline silk, [ \n \n 10 \n \n ] and F) Young's modulus. Whiskers show standard deviation. * p < 0.05; ** p < 0.01; *** p < 0.001; **** p < 0.0001. Representative stress–strain graphs for all spinnable engineered mini‐spidroins are shown in Figure S5A–G , Supporting Information. The diameters of the fibers are shown in Figure S5H , Supporting Information. The values and corresponding standard deviations are shown in Table S3 , Supporting Information. The tensile strength of all fibers spun from engineered proteins increased significantly compared to A 15 ‐A 14 except for (A 3 V) 3 ‐A 14 and (A 2 I) 4 ‐A 14 (Figure 4 and Table S3 , Supporting Information). The two similar fiber types (A 3 I) 3 ‐A 14 and A 15 ‐(A 3 I) 3 displayed the highest increase in strength, the former reaching 131 MPa, which is almost three times higher than that of A 15 ‐A 14 (Figure 4C ). This indicates that rational protein engineering of the spidroin poly‐Ala blocks indeed can result in increased fiber tensile strength and stiffness. Unexpectedly, the introduced amino acid substitutions also had a high impact on the extensibility of the fiber, as the strain at break varied from 0.03 to 2.0 mm mm −1 (Figure 4D and Table S3 , Supporting Information). The two strongest fiber types ((A 3 I) 3 ‐A 14 and A 15 ‐(A 3 I) 3 ) displayed an exceptional increase in strain (to 1.6 and 2.0 mm mm −1 , respectively), while (A 3 V) 3 ‐A 14 , (A 2 I) 4 ‐A 14 fibers showed moderately increased strain (0.79 and 0.85 mm mm −1 , respectively) compared to A 15 ‐A 14 (0.45 mm/mm). (A 3 T) 3 ‐(A 3 T) 3 and (A 3 V) 3 ‐(A 3 V) 3 fibers were the least extensible (0.03 and 0.08 mm mm −1 , respectively). These two proteins contain substitutions in both poly‐Ala blocks and, possibly, the reason for the inferior strain of these fibers could be an increased propensity of the engineered segments to interact intra‐molecularly over forming intermolecular contacts. Apparently, the mechanical properties of artificial spider silk fibers can be significantly improved by introducing Ile in every fourth position in the first or second poly‐Ala block. These two mini‐spidroins, (A 3 I) 3 ‐A 14 and A 15 ‐(A 3 I) 3 , formed fibers with a toughness modulus that is comparable to native dragline silk (146 and 125 MJ m −3 , respectively, compared to 136 MJ m −3 for a native dragline silk from Argiope argentata ), (Figure 4E ). [ \n \n 10 \n \n ] Fibers formed by (A 3 V) 3 ‐A 14 and (A 2 I) 4 ‐A 14 also reached a significantly higher toughness modulus than A 15 ‐A 14 (50 and 37 MJ m −3 , respectively, compared to 18 MJ m −3 ). To investigate the link between fiber secondary structure content and mechanical properties, we used attenuated total reflection Fourier‐transform infrared (ATR‐FTIR) spectroscopy. The results, shown in Figure \n 5 \n and Figure S7 and Table S4 , Supporting Information, indicate that no large differences in secondary structure content between fibers were detected, but (A 3 V) 3 ‐A 14 , (A 3 I) 3 ‐A 14 and A 15 ‐(A 3 I) 3 had a slightly increased β‐sheet content, along with decreased α‐helix/random coil content compared to A 15 ‐A 14 fibers. However, the (A 3 V) 3 ‐(A 3 V) 3 and (A 2 I) 4 ‐A 14 fibers failed to show increased β‐sheet content compared to A 15 ‐A 14 fibers and we could detect no strong correlations between secondary structure content and mechanical properties of the fiber (Figure S6 , Supporting Information). Thus, ATR‐FTIR spectroscopy of the different fiber types did not detect any significant differences in secondary structure content. Therefore, we decided also to use solid‐state NMR spectroscopy to investigate the unmodified fibers (A 15 ‐A 14 ) and the best performing engineered fibers, (A 3 I) 3 ‐A 14 . As expected, more Ala residues were found in a β‐sheet conformation in (A 3 I) 3 ‐A 14 compared to A 15 ‐A 14 fibers ( Figure \n 6 \n ). Figure 5 FTIR spectroscopy of engineered fibers. Normalized and baseline‐subtracted absorbance spectrum in the amide I region of A) A 15 ‐A 14 , (A 3 V) 3 ‐(A 3 V) 3 , (A 3 V) 3 ‐A 14 , and (A 3 T) 3 ‐(A 3 T) 3 and B) A 15 ‐A 14 , (A 3 I) 3 ‐A 14 , A 15 ‐(A 3 I) 3 , and (A 2 I) 4 ‐A 14 . C) Percent secondary structure content determined by cofitting the absorbance spectrum and the second derivative. Horizontal line indicates β‐sheet content of A 15 ‐A 14 . Fits of absorbance spectra and second derivative of fibers spun are shown in Figure S7 , Supporting Information. Figure 6 Solid‐state NMR 13 C‐ 13 C correlation spectra (aliphatic region) of A 15 ‐A 14 (blue) and (A 3 I) 3 ‐A 14 (red) fibers. The Cα/Cβ correlations of Ala and Ile in α‐helical and β‐sheet conformation are indicated. The altered mechanical properties of the fibers made from the engineered spidroins indicate that intermolecular interactions in the spidroins are affected. In the native dragline silk fiber, pulling the fiber first results in reversible deformation of the amorphous regions up until the yielding point, after which the hydrogen bonds in the amorphous region break, resulting in softening of the material. [ \n \n 36 \n , \n 43 \n \n ] When the amorphous protein chains are extended, the load is transferred onto the β‐sheet crystals leading to a stiffening of the fiber. Upon further increased load, the β‐sheet crystals undergo stick‐slip deformation and the fiber breaks. [ \n \n 36 \n , \n 43 \n , \n 69 \n \n ] The increased tensile strength of the fibers made from engineered proteins suggests that our strategy to increase the β‐strand propensity and inter‐β‐sheet interactions indeed can result in stronger fibers, although some of the engineered fibers concomitantly displayed a decreased strain. Theoretically, increased β‐sheet formation and intermolecular interactions in the stacked β‐sheets could not only result in increased fiber strength, but also increased extensibility, since the amorphous region would be allowed to extend fully before the load is transferred to the crystalline region. In lack of poly‐Ala β‐sheet crystals, as in the A 15 ‐A 14 fibers, the intermolecular contacts may be too weak to allow a full extension of the amorphous protein chains before fiber failure. At the same time, it may be disadvantageous that all β‐sheets stack in crystals since only about 40% of the Ala residues in the native dragline silk are found in this conformation and the rest form less ordered β‐sheets. [ \n \n 70 \n \n ] In this study, introducing replacements in both poly‐Ala blocks resulted in fibers with dramatically reduced strain which suggest a suboptimal packing of the proteins in the fiber. Since the (A 3 I) 3 ‐A 14 fibers displayed superior mechanical properties, these fibers are attractive candidates for bulk‐scale production. Previously, A 15 ‐A 14 has been shown to express at very high levels (≈21 g L −1 ) in a bioreactor‐based E. coli fed‐batch culture. [ \n \n 56 \n \n ] Following the same protocol, the expression level of (A 3 I) 3 ‐A 14 amounted to 13 g L −1 and the final yield after purification using an automated purification protocol was 8.9 g L −1 (Figure S8 A,B, Supporting Information). To our knowledge, these yields are the second highest reported for any recombinant spidroin produced in E. coli and line with what is required for economically viable bulk production. [ \n \n 55 \n , \n 56 \n \n ] After purification, (A 3 I) 3 ‐A 14 was concentrated to 300 mg mL −1 and could easily be spun into fibers. Notably, 8.9 g recombinant silk protein is enough to produce an ≈18 km long fiber. When comparing (A 3 I) 3 ‐A 14 fibers produced from proteins recovered from bioreactor and shake flask fermentations, respectively, the former had slightly lower strength (Figure S9D , Supporting Information). However, the bioreactor produced (A 3 I) 3 ‐A 14 fibers still had a significantly higher tensile strength and strain compared to A 15 ‐A 14 fibers (Figure S9 , Supporting Information)."
} | 7,180 |
38343096 | PMC10934265 | pmc | 11 | {
"abstract": "Despite the considerable\ninterest in the recombinant production\nof synthetic spider silk fibers that possess mechanical properties\nsimilar to those of native spider silks, such as the cost-effectiveness,\ntunability, and scalability realization, is still lacking. To address\nthis long-standing challenge, we have constructed an artificial spider\nsilk gene using Golden Gate assembly for the recombinant bacterial\nproduction of dragline-mimicking silk, incorporating all the essential\ncomponents: the N-terminal domain, a 33-residue-long major-ampullate-spidroin-inspired\nsegment repeated 16 times, and the C-terminal domain (N16C). This\ndesigned silk-like protein was successfully expressed in Escherichia coli , purified, and cast into films from\nformic acid. We produced uniformly 13 C– 15 N-labeled N16C films and employed solid-state magic-angle spinning\nnuclear magnetic resonance (NMR) for characterization. Thus, we could\ndemonstrate that our bioengineered silk-like protein self-assembles\ninto a film where, when hydrated, the solvent-exposed layer of the\nrigid, β-nanocrystalline polyalanine core undergoes a transition\nto an α-helical structure, gaining mobility to the extent that\nit fully dissolves in water and transforms into a highly dynamic random\ncoil. This hydration-induced behavior induces chain dynamics in the\nglycine-rich amorphous soft segments on the microsecond time scale,\ncontributing to the elasticity of the solid material. Our findings\nnot only reveal the presence of structurally and dynamically distinct\nsegments within the film’s superstructure but also highlight\nthe complexity of the self-organization responsible for the exceptional\nmechanical properties observed in proteins that mimic dragline silk.",
"conclusion": "Conclusions In this study, we have\ndemonstrated the successful cloning and\nexpression of a designed, synthetic spider dragline silk based on\nthe amino acid sequence of the dragline silk proteins of Nephila clavipes . The construct contained all the\ncharacteristic building blocks of natural spider silks, including\nthe nonrepetitive N- and CTDs, as well as the repetitive core domains\nrepeated 16 times (N16C), resulting in a final protein size of 68.1\nkDa. We achieved the straightforward directional assembly of the modules\nusing Golden Gate assembly, which can be easily utilized to further\nincrease the repetition size. Additionally, we developed an\nefficient recombinant protein production\nand purification strategy and characterized the structure and dynamics\nof the recombinantly produced, 13 C, 15 N-labeled\nspidroin mimic with solution- and solid-state NMR spectroscopy. By\nanalyzing the 1 H, 13 C, and 15 N chemical\nshifts, we evaluated the secondary structure of the core repetitive\ndomain of N16C dissolved in DMSO- d 6 . Furthermore,\nwe analyzed the structure and dynamics of hydrated N16C film cast\nfrom formic acid solution using proton-detected solid-state MAS methods,\nemploying both cross-polarization and J -coupling-based\nmagnetization transfers. Comparing the solution and solid-state chemical\nshifts, as well as the 15 N and 13 C relaxation\nrate constants measured in the solid state, we identified three structurally\nand dynamically distinct segments in the film superstructure. These\nsegments include a rigid, strongly hydrogen-bonded β-nanocrystalline\ncore surrounded by a solvent-exposed, dynamic α-helical shell\nin exchange with fully solubilized, flexible repeat units with random\ncoil characteristics. These findings highlight the complexity of the\nhierarchical organization responsible for the remarkable mechanical\nproperties of dragline-silk-mimicking proteins. By harnessing the\npower of recombinant silk production and advanced spectroscopic techniques,\nwe are one step closer to understanding the correlation between molecular\nstructure and mechanical response of silk-based high-performance materials.",
"introduction": "Introduction Spider dragline silks, or spidroins, are\nnatural fibers with exceptional\nproperties including their unique lightweight, impressive extensibility,\nhigh tensile strength, durability, and biocompatibility. 1 − 7 These unparalleled qualities make spidroins valuable materials for\nvarious industrial and biomedical applications. Throughout history,\nthey have been utilized in diverse fields such as clothing, fishing,\npainting, medicine, and weaponry. 4 , 8 Despite significant\nefforts in replicating the mechanical features of spidroins, artificial\nfibers often fall short in terms of versatility and overall performance\ncompared to those of natural spider silks. 9 − 11 However, recent\nadvancements in producing spidroins with previously unattainable molecular\nweights through heterologous bacterial expression 12 or utilizing transgenic silkworms 13 are bringing us closer to the commercialization of tailor-made bioengineered\nsilk fibers. Spider dragline silk has attracted considerable attention\ndue to its exceptionally high tensile strength, exceeding that of\nsteel, and toughness that is three times higher than that of Kevlar. 4 , 14 Most dragline silks are made from two major ampullate spidroin proteins\n[major ampullate spidroin protein 1 (MaSp1) and major ampullate spidroin\nprotein 2 (MaSp2)], which share a common structure, characterized\nby a low-complexity, highly repetitive core with up to a few hundred\nrepeats of 20–200 amino acids. This core is flanked by highly\nconserved, nonrepetitive α-helical N- and C-terminal domains\n(NTD and CTD). The repeat regions have a block copolymer structure\nin which hydrophilic “soft” glycine (Gly)-rich and hydrophobic\n“hard” alanine (Ala)-rich segments alternate. These\nsegments are composed of shorter consensus motifs, including poly(A),\npoly(GA), GGX, GSG, QQ, and GPGXX stretches, where X = Y (Tyr, tyrosine),\nL (Leu, leucine), or Q (Gln, glutamine). Solid-state nuclear\nmagnetic resonance (NMR) data suggested that\nthe Gly-rich segments predominantly adopt semiextended 3 1 -helices, while the GPGGX and GPGQQ elements of MaSp2 form elastin-like\ntype II β-turns. 15 − 24 Alternatively, these segments can be incorporated into β-sheet\nstructures, and thus, they are part of the hard segments. 19 , 20 , 24 , 25 The low density of hydrogen bonds in the soft Gly-rich regions grants\nextensibility to the dragline fiber, 7 , 26 while the\nsuccessive β-turns are associated with the supercontractive\nproperty of the silk fiber. 21 Upon sheer-induced\nstress, the majority of the alanine residues arrange into β-sheets\nand form β-nanocrystals, whereas the alanines in poly(GGA) adopt\n3 1 -helix structures. 15 , 18 , 24 , 25 In the β-nanocrystals,\nthe poly(A) β-sheets tightly interlock, resembling the steric\nzippers of amyloid crystals, thereby impeding water penetration between\nthe poly(A) sheets. 17 The relatively small\ncrystalline domains 27 − 31 embedded within the semiamorphous matrix act as intermolecular cross-links,\ncontributing to the high tensile strength of the spun fibers. 7 , 24 , 26 Additionally, the large number\nof repeated core units enhances interchain interactions and reduces\nchain-end defects, further augmenting the unprecedented tensile strength\nof the silk fibers. 7 , 12 , 13 , 32 The importance of the high repeat\nnumber in recombinant spider\nsilk proteins was recognized early on, 33 − 35 and attempts have been\nmade to increase the repeat number all the way up to 192. 12 , 32 Achieving such a feat required the implementation of extensive metabolic\nengineering and synthetic biology approaches, which have not yet been\nadapted for large-scale biotechnological production. In addition to\nthe repeat number, the terminal domains play fundamental roles in\ndetermining the properties of silk proteins, especially with respect\nto protein gland solubility and initiation of fiber assembly through\nsalt- and pH-dependent dimerization. 36 − 38 Previous studies have\nestablished a correlation between recombinantly produced silk’s\ntensile properties, such as the Young’s modulus, strength,\nand toughness of the fibers, with the presence of the CTD and the\nNTD. 9 , 39 For such reasons, the inclusion of the highly\nconserved globular terminal domains in the final spidroin product\nis necessary, especially when assembly and solubility properties are\nof interest. 36 Repetitive DNA sequences\nprovide challenges for standard cloning\ntechniques, for example, due to their inherent ability to recombine\nand hence instability. Typically, recombinant tandem-repeat DNA sequences\nof fibrous proteins, including mimics of spider silks (reviewed in\nrefs ( 5 , 11 , 40 – 42 )), collagen, 43 elastin, 44 , 45 keratin, 46 and resilin, 47 are constructed via stepwise concatenation, 48 , 49 recursive directional ligation, 12 , 44 , 45 or step-by-step directional approaches, 50 − 52 which involve several rounds of plasmid amplification, digestion,\nligation, and possible sequencing between repeat extensions. 52 − 54 In these stepwise techniques, at each oligomerization step, the\ncloning efficiency is greatly reduced by the possibility of obtaining\nempty vectors or self-ligated and circularized inserts. Aside from\nthe time and material costs associated with these traditional approaches,\na major drawback of a few of these methods is the scars that remain\nat the recombinant sites in the final constructs. These scars are\nthen translated into extraneous amino acids in the primary protein\nsequence, compromising the accuracy and potentially the structural\nintegrity of the expressed protein. Nevertheless, seamless stepwise\ncloning has been successfully applied in a few instances to produce\nrecombinant fibrous proteins. 44 , 55 − 59 To overcome these shortcomings, in this study, we used Golden\nGate\nassembly 60 to generate an expression construct\nfor a spidroin mimic, called N16C ( Figure 1 ), which includes both the N- and C-terminal\ndomains and 16 repeats of the repetitive core units. Golden Gate assembly\nrelies on type II restriction endonucleases, which cut double-stranded\nDNA outside their recognition sequence and leave a short, single-stranded,\nuser-defined overhang that guarantees ordered gene assembly of multiple\nconstituents. The restriction sites are eliminated during subcloning,\nwhich allows for simultaneous digestion and ligation in a one-pot\nreaction and facilitates the seamless assembly of gene constituents.\nPreviously, Golden Gate assembly has been successfully used to straightforwardly\nassemble repetitive DNA sequences encoding elastin-like proteins (ELPs). 59 Figure 1 Schematic representation and amino acid sequence of the\nN16C protein\nconstruct: NTD (blue), derived from the N. clavipes MaSp2, the repetitive central domain (S16, black), based on N. clavipes MaSp1 and MaSp2, and CTD (maroon) from N. clavipes MiSp1. Amino acids labeled with green\naccount for the His-tag used for purification purposes. High repeat numbers and the presence of terminal\ndomains\nare necessary\nbut are not sufficient requirements for reproducing the properties\nof natural silk fibers. The failure in natural silk reproduction arises\nfrom the incomplete understanding of the driving forces behind the\nprotein’s self-assembly across multiple spatial scales. Open\nquestions include the exact structures of the involved molecules and\nsequence motifs, the types of interactions and transformations they\nundergo, the kinetics and thermodynamics of their interactions, and\nthe dynamics that affect the local energetics and the macroscale mechanics\nof the silk fibers. Addressing these questions requires atomic-level\ninsights into the protein assembly both experimentally and computationally.\nSolid-state magic-angle spinning (MAS) NMR spectroscopy is arguably\nthe most suitable experimental technique to investigate complex heterogeneous\nsystems as it has been demonstrated for natural 17 − 21 , 23 − 25 , 61 − 65 and genetically engineered spider silks, 4 , 56 , 66 − 70 or for selectively isotope-enriched silk-mimicking\nmodel peptides as reviewed in refs ( 71 – 74 ). To fully exploit the advantages of recombinant spider silk\nproduction,\nwe expressed and purified a uniformly 13 C– 15 N-labeled version of our designed spidroin mimic N16C and measured\nits solid-state MAS NMR spectra to gain insights into the local structure,\nhydrogen-bonding, nanosecond and microsecond time-scale dynamics,\nand hydration-induced macroscopic organization of each amino acid\ntypes in the cast film. Fast MAS (55.55 kHz) and high magnetic fields\n(700 MHz) combined with proton detection significantly improved the\nsensitivity and the resolution of the acquired spectra, allowing for\nthe almost complete amino-acid-specific assignment of the repetitive\ncore, which then facilitated the analysis of site-specific dynamics\nacross multiple time scales. The uncovered correlation between the\natomic-level structure and hydration-induced local dynamics of the\ndesigned synthetic spidroins provides critical insight into the dynamic\norganization of natural silks and engineered silk-like proteins.",
"discussion": "Discussion Spidroins\nhave attracted significant attention in the fields of\nbiotechnology and material science due to their remarkable properties,\nincluding the ability to form strong, stable, and tough fibers. However,\ntheir expression in alternative hosts like E. coli presents challenges, primarily due to the inherent instability and\nrepetitive nature of the encoding DNA sequences, as well as their\nlimited solubility in the expression host. 93 Consequently, extensive efforts have been dedicated to improving\nprotein production and purification processes, optimizing cloning\nstrategies, and designing genetic circuits to regulate gene expression\nresources. Despite these endeavors, achieving a high yield in fibrous\nprotein production remains an unresolved issue. Encouragingly, we\nmanaged to successfully express a 68 kDa model spidroin protein, N16C,\ncomprising 16 repeat units, with a comparatively high yield in TB\nmedium, without the need for additional metabolic or genetic engineering\ninterventions. The Golden Gate assembly technique employed in constructing\nthe silk gene sequence in this study can be readily applied to generate\nDNA sequences encoding 32, 64, or even 128 repeat units, thereby enabling\nthe expression of N32C, N64C, or N128C spidroin mimics, respectively.\nHowever, it is important to note that producing such large repeat\nproteins will undoubtedly present several challenges associated with\nheterologous protein expression. An extra benefit of being able\nto express silk-like proteins in E. coli is that they can be easily produced in isotopically\nenriched forms for downstream NMR investigations. Solid-state NMR\nprovides unprecedented insights into the atomic-resolution structure,\ndynamics, organization, and interaction of semiamorphous semicrystalline\nsystems, like silk. The sensitivity enhancement achieved by uniform 13 C and 15 N labeling combined with proton detection\nand fast solid-state MAS NMR gave us access to a wide range of multidimensional\nNMR experiments that we customized to focus on specific aspects of\nthe material property. For example, we used two fundamentally different\nmagnetization transfers (cross-polarization and INEPT) to access either\nthe rigid core or the solubilized flexible segments of the hydrated\nN16C film, and we built these transfer steps into both assignment\nexperiments and spin relaxation measurements. The resulting chemical\nshifts and relaxation rate constants allowed for the extraction of\nsite-specific information on hydrogen bonding, secondary structure,\nand nanosecond-to-microsecond time scale dynamics of both rigid and\nsoluble states. Since Ala and Gly are the major constituents of the\nprotein sequence, we focused more on their analysis as they selectively\nreport on the properties of the hard and soft segments. Since\nin our designed spidroin sequence, only one type of Ala exists\nin the repetitive segment, the observed three different sets of Ala\nresonances must stem from the same poly(A) sequence in different chemical\nenvironments. Based on their chemical shifts and relaxation properties,\nwe assigned them to belong to Ala (i) inside the β-nanocrystals,\n(ii) at the water–protein interface, (iii) and in solution.\nIn the solid-state MAS spectra of native dragline silk fibers, Holland\net al. observed β-stranded as well as α-helical Ala and\nGly environments, but they associated the separate shifts with sequentially\ndifferent residues, e.g., with Ala that is part of the poly(A) segments\nor Ala in GGA elements flanking the poly(A) regions. 19 , 20 , 63 The structural and dynamic\ninformation on the Ala resonances suggested\na model for the repetitive region of the N16C film, where a densely\npacked β-nanocrystalline core is surrounded by a solvent-exposed\nlayer with α-helical conformation that is in dynamic exchange\nwith fully solubilized repetitive units of N16C, which are highly\nflexible and unstructured ( Figure 8 ). Figure 8 Schematic representation of the conformational states\nin the hydrated\nN16C film. The enlargement shows a structural model of the repetitive\nregion in which the poly(Ala) and some of the Gly-rich segments form\nβ-sheets and arrange into randomly oriented β-nanocrystals\n(blue), whereas the rest of the Gly-rich segments are unordered or\nadopt a 3 1 -helical structure. At the interface of the β-nanocrystals\nand amorphous segments, the poly(Ala) repeats form α-helices\n(red). Some protein chains are fully solubilized and adopt a random\ncoil conformation (green). The majority of Gly in the soft segment adopted\nan extended β-sheet\nconformation and became part of the β-nanocrystals, while a\nsmaller fraction was found in a different conformation that we tentatively\nassociated with the semiextended 3 1 -helical conformation.\nThe large fraction of Gly in β-sheet could be a consequence\nof casting the film from formic acid. Formic acid is known to initiate\nβ-sheet formation in regenerated Bombix mori silk fibroin and in recombinant spider silk films. 69 , 94 − 96 The freshly prepared dry N16C film was brittle but\nbecame elastic after isopropanol treatment and incubation in water\novernight. Stiffness and brittleness are associated with a high number\nof β-sheets and long-range order. 84 , 94 As water interacts\nwith the ordered β-crystalline structures, the crystallinity\ndecreases and the film becomes more elastic, as demonstrated on B. mori fibroin films. 94 , 97 In the hydrated\nN16C film, even though most Gly was incorporated into the rigid β-crystalline\ndomain, it actively contributed to the elasticity of the film as it\nshowed extensive μs time scale flexibility. Our findings align\nwith the results of other X-ray and NMR studies of spider silk fiber\nthat suggested a three-phase model, where the rigid β-crystalline\ncore and the elastic amorphous matrix are interconnected by a semiordered\nphase. 3 , 98"
} | 4,713 |
19229199 | null | s2 | 12 | {
"abstract": "The extreme strength and elasticity of spider silks originate from the modular nature of their repetitive proteins. To exploit such materials and mimic spider silks, comprehensive strategies to produce and spin recombinant fibrous proteins are necessary. This protocol describes silk gene design and cloning, protein expression in bacteria, recombinant protein purification and fiber formation. With an improved gene construction and cloning scheme, this technique is adaptable for the production of any repetitive fibrous proteins, and ensures the exact reproduction of native repeat sequences, analogs or chimeric versions. The proteins are solubilized in 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP) at 25-30% (wt/vol) for extrusion into fibers. This protocol, routinely used to spin single micrometer-size fibers from several recombinant silk-like proteins from different spider species, is a powerful tool to generate protein libraries with corresponding fibers for structure-function relationship investigations in protein-based biomaterials. This protocol may be completed in 40 d."
} | 271 |
End of preview. Expand in Data Studio
README.md exists but content is empty.
- Downloads last month
- 11