Nature produces load-bearing minerals possessing pores whose diameters range from nm up to mm. To control the pore size of our synthetic biominerals over a similar range of length scales, we combine different techniques. We tune the size of small pores, with diameters ranging from the 100s of nm up to the 100s of µm, with emulsion templates. By functionalizing and mineralizing the surface of emulsions, we create capsules with empty cores whose sizes correspond to those of the emulsion drops. These capsules can be used as building blocks of macroscopic porous materials.
Emulsion drops with diameters between 200 nm and 120 µm are prepared by mixing a fluoroether-based oil, HFE7500, with an aqueous solution containing calcium chloride (CaCl2) and poly(vinyl alcohol) (PVA). To prevent coalescence of the emulsions, we add an amphiphilic block copolymer surfactant whose hydrophilic block is end-functionalized with two pyrogallols to the oil phase, as schematically shown in Figure S1. These surfactants absorb at the emulsion surface where they are ionically crosslinked by calcium ions (Ca2+) contained in the aqueous phase. The oil/water mixture is tip sonicated for 6 minutes to result in polydisperse oil-in-water emulsions whose diameters are hundreds of nanometers. To enlarge the range of emulsion sizes we can produce, which directly translates into the range of pore sizes that can be introduced into our biomineral, we also fabricate emulsion drops by vortexing the oil/water mixture for 25 s. The resulting emulsions are again polydisperse with diameters ranging from 6 to 52 µm, with an average diameter of 22 µm and a standard deviation of 12 µm. Emulsions with a narrower size distribution are fabricated with microfluidic flow-focusing devices by injecting the inner oil phase containing the surfactant at a flow rate of 30 µl/min and the outer aqueous phase containing Ca2+ at a flow rate of 50 µl/min. The resulting emulsions display an average drop size of 112 µm and a standard deviation of 8 µm, as illustrated in the optical micrographs in Figure S2 and summarized with the diameter distributions in Figure S7.
To control the structure of porous materials on larger length scales, we combine emulsion templates with 3D printing. To enable 3D printing, the emulsions must be processed into a shear-thinning ink. To fulfill this requirement, we up-concentrate emulsions through sedimentation. Indeed, if the appropriate amount of PVA is added to the aqueous phase, the resulting ink is shear-thinning, as shown in Fig. 2a and detailed in Figure S3. The flow point, defined as the crossover of the storage (G’) and loss modulus (G’’) in oscillation amplitude sweep measurements, can be tuned with the Ca2+ concentration present in the surrounding aqueous media, as summarized in Fig. 2b. For example, inks containing 0.1 M Ca2+ possess an increased viscosity and lower flow point compared to those lacking any Ca2+ ions. We assign the difference in viscosity to the increase of inter-emulsion friction that comes from the rougher surfaces of capsules converted from surface functionalized emulsions and inter-emulsion crosslinking via Ca2+ ions24. The enhanced interactions among capsules composed of ionically crosslinked surfactants also condense the ink formulation and hence, decrease the strain at the flow point of the ink.
To ensure a good shape fidelity of the 3D printed parts, the ink must rapidly recover its storage modulus once the shear is released. To demonstrate this feature, we perform shear recovery tests on our samples by repetitively shearing them for 60 s at a strain of 1% followed by 60 s of shear at a strain of 100% while recording the storage and loss moduli. Indeed, the ink displays reversible transitions from solid-like to liquid-like states, as shown during the first 300 s in Fig. 2c.
The storage modulus of our viscoelastic ink is too low to ensure good shape retention of 3D printed parts. To overcome this limitation, we increase the storage modulus of our ink by printing it into an aqueous bath containing Na2CO3 to initiate the gelation of PVA25 and the mineralization of the capsule surfaces. Indeed, the storage modulus of the printed filaments increases to 2000 Pa upon insertion into the carbonate source, which is 400 times higher than that of the original ink, as shown in the blue shaded area in Fig. 2c. We assign the strong increase in storage modulus primarily to the precipitation of PVA, that is caused by its lower solubility in salt-rich aqueous solutions, according to the Hofmeister effect26.
To transform an emulsion-based ink into a solid porous structure, we must selectively induce mineralization at the surfaces of the capsules and in the surrounding aqueous solution. Our capsules present pyrogallol functionalities at their surfaces that point towards the aqueous phase, as schematically shown in Fig. 1a. Hence, the surfactants endow the capsule surface with a high affinity towards Ca2+ present in the aqueous phase owing to the Ca2+–pyrogallol interactions27. Previous reports demonstrated that PVA films28 or organic matrices containing metal-coordinate complexes29,30 template the growth of inorganic precipitates. We hence expect minerals to preferentially start to form at the capsule surface and thereafter within the polymer matrix, when capsules presenting high concentrations of Ca2+ ions at their surfaces are immersed in an aqueous solution containing a carbonate source. To test this expectation, we investigate the mineral-polymer interactions using Raman and Fourier-transform infrared (FTIR) spectroscopy. Indeed, we observe the peaks from surfactants, the PVA matrix as well as CaCO3 crystals in Raman spectra of samples that have been mineralized, as shown in Figure S4. The O-H stretch vibration from the phenolic hydroxyl groups in the FTIR spectrum shifts from 3300 − 3050 cm− 1 to 3400 − 3150 cm− 1 upon crosslinking with Ca2+ and mineralization, indicating the assumed pyrogallol-Ca2+ interactions.
Mineralization that happens within a polymer matrix typically increases the storage and loss moduli of the polymer matrix30. To assess if this is also the case for our structured materials, we perform frequency sweeps on our samples and quantify the plateau moduli as a function of the time that samples have been incubated in the mineralizing solution. Within the first hour of incubation, the stiffness of the 3D printed hydrated scaffolds increases with increasing incubation time before it reaches a plateau, as shown in Figure S5. This result indicates that the vast majority of minerals forms and grows within the first hour of incubation. To test if the mineral formation is limited by the amount of Ca2+ present within the system, we quantify the plateau storage modulus of scaffolds that have been mineralized for 1 h as a function of the Ca2+ concentration initially contained in the emulsion-based ink. Indeed, the plateau storage modulus of these samples increases with the concentration of Ca2+, as shown in Figure S6, indicating that the amount of Ca2+ in the system limits the degree of mineralization of our capsules.
To transform the mineralized structures into porous materials, we dry them in air. The fluoroether-based oil we use has a vapor pressure of 2.1 kPa at 25°C. This value is close to that of water, which possesses a vapor pressure of 3.2 kPa at 25°C, such that the solvent evaporates at room temperature. To quantify the sizes of the resulting pores, we visualize dried samples with scanning electron microscopy (SEM). Samples made from emulsions prepared by vortexing display pore sizes ranging from 4 to 45 µm with an average of 18 µm and a standard deviation of 11 µm, as shown in Fig. 3a and Movie M3. This pore size range is very similar to the diameter of the emulsion drop templates they have been made from, as shown in Figure S7. The slight decrease in pore sizes compared to the intial emulsion sizes can be assigned to the shrinkage of the structure caused by capillary forces. The good correlation between the pore size and the size of initial emulsions is in good agreement with previously reported results on the fabrication of porous ceramics from emulsion templates7,31. We test if this correlation also holds for larger emulsions, which are more difficult to stabilize against coalescence. Therefore, we analyze the pore sizes of samples prepared from monodisperse emulsion drop templates that have been prepared with microfluidic devices. These samples display a narrow pore size distribution with an average diameter of 87 µm and a standard deviation of 17 µm, as shown in Fig. 3b. To visualize the porous structure in 3D, we perform X-ray micro-computed tomography (µ-CT) on them. The vast majority of the pores are spherical, as shown in Fig. 3b (iv,v) and Figure S8b. This pore morphology suggests that the majoriy of the emulsions remains intact during the processing, thereby offering tight control over the size of pores at the micrometer range.
The best density-normalized mechanical performance is typically achieved in materials possessing a highly ordered hierarchical architecture32. However, processes that lead to structures with hierarchical porposities typically require high energy input, such as ice-templating33, pyrolysis5 or metal oxidation during sintering34. To test if our low energy process enables the formation of structures with hierarchical porosities whose mechanical properties outperform those possessing pores that are only in one size range, we mix two batches of emulsions possessing drops with different sizes: We mix one batch of emulsions produced through vortexing, whose diameters are in the tens of µm, with one batch produced by tip sonication, whose emulsion diameter is of order 100 nm. The emulsion mixture is up-concentrated, molded and mineralized before it is dried and visualized using SEM and X-ray µ-CT. Indeed, these samples contain pores ranging from 100s of nm up to 10s of µm, as exemplified in Fig. 3c.
An important parameter that determines the shape fidelity of 3D printed parts is their shrinkage upon drying. While shrinkage can be partially avoided by drying in a humid environment, this process prolongs the drying time. To shorten the processing time and prevent that the capillary forces collapse pores during drying, we strengthen our composite by enhancing the inter-particle links with additional minerals, thereby increasing the volume fraction of minerals. This can be achieved by increasing the initial Ca2+ concentration in the ink formulation to 1 M. The resulting samples are visualized using X-ray µ-CT. Indeed, these structures contain pores with diameters of order 10s of µm, as exemplified in the virtual tomography slice and 3D renderings in Figure S9a. However, the stability of the emulsions decreases if the Ca2+ concentration is increased above 1 M, and hence compromises the shape fidelity of the porous structure.
To address this limitation while still increasing the mineral fraction in the composites, we keep the Ca2+ concentration at 1 M and subject the 3D printed structures to repeated mineralization cycles. This is achieved by washing the 3D printed structures before immersing them in an aqueous solution containing 1 M Ca2+ for 2 h and again incubating them in an aqueous solution containing 1 M CO32− for 2 h, as schematically illustrated in Fig. 4a. We quantify the shrinkage of our composite by measuring the dimensions of samples as printed in the bath and after drying. Indeed, the shrinkage of the printed object linearly decreases from ~ 45% to ~ 16% with increasing number of mineralization cycles, as evidenced by direct measurements on a 20 mm\(\times\)20 mm\(\times\)3 mm printed lattice after two mineralization cycles, shown in Fig. 4c and Figure S10. This degree of shrinking is much lower than that typically observed during pyrolysis of polymer-derived ceramics, which is usually around 30%35.
We assign the measured reduction in shrinkage with increasing number of mineralization cycles the composite has been subjected to, to an increase in its mineral content. To test this hypothesis, we quantify the CaCO3 weight fraction of composites that have undergone up to three mineralization cycles using thermogravimetry (TGA). The CaCO3 content increases with mineralization cycles, with an average of 77 wt% in samples subjected to two mineralization cycles, as shown in Fig. 4b. To visualize the effect of the additional mineralization on the porosity of our samples, we image them with SEM and X-ray µ-CT. Samples that have not been subjected to any mineralization cycle possess many interconnected open pores, as shown in Fig. 4f,g. We assign the open porosity to a partial collapse of the structure during the evaporation of the solvent. By contrast, samples subjected to two mineralization cycles contain many spherical pores that are much less interconnected, as shown in Fig. 4h,i. We assign this finding to the newly formed CaCO3 that can only grow within the hydrophilic parts of the composites, which is the surrounding of the pores. As the mineral content increases, the interstitial spaces get increasingly mineralized, eventually leading to the formation of intact, closed, spherical mineral shells. Indeed, the mineral content can only be controllably increased by subjecting the samples to multiple mineralization cycles and not by increasing the initial Ca2+ concentration as a high Ca2+ concentration results in a collapse of these spherical structures, as shown in Figure S9. However, if we subject the composite to more than two mineralization cycles, we compromise its shape fidelity because crystals start to grow out of the 3D printed filaments. For example, filaments, formed by extruding the paste through a 515 µm diameter nozzle, initially have a diameter of approximately 360 µm, as shown in Fig. 4d. After three mineralization cycles, their diameter increases more than two-fold to 800 µm, as shown in Fig. 4e. Based on these results, we subject our samples to two mineralization cycles as this protocol results in the best shape fidelity while still reaching a high mineral content. These results nicely demonstrate the degree of control over the microstructure our process offers.
To assess the influence of the pore structure on the mechanical properties of our porous materials, we perform nanoindentation (NI) tests on them. The hardness of our polymeric scaffold increases up to 12-fold and its stiffness up to 4.5-fold upon mineralization, as shown in Fig. 5a. Indeed, samples subjected to two mineralization cycles, containing around 80 wt% CaCO3 display a hardness and stiffness similar to those of human trabecular bones, even though bulk calcite is weaker than hydroxyapatite36. We assign the excellent mechanical properties of our composites to the high CaCO3 content contained in them.
The compressive strength of composites is closely related to their density. To perform a fair comparison between our samples, we plot the strength measured in uniaxial compresion tests as a function of the density of the respective samples, as detailed in Fig. 5b. As expected, molded bulk samples possessing pores whose diameters span many orders of magnitudes, labeled as ‘hierarchical pores’, have lower densities than samples containing pore diameters that are all within a similar length scale, referred to ‘polydisperse pores’. Yet, the compressive strengths of the two types of samples are similar. We assign this finding to the stress distribution within porous composites: Composites possessing pore sizes that span many orders of magnitudes contain many small mineral struts that enable an efficient stress transfer and prevent extensive distortions and associated build-up of stress concentrations in the macro-struts34,37. Hence, they result in higher specific strength, defined as the compressive strength divided by the density of the material. Nature often uses this trick to build lightweight yet damage-resistant materials such as bamboo38, the skeleton of glass sponges1, avian bones39 and feathers40.
A comparison between our samples and lightweight, damage-resistant natural biominerals reveals that we still have a lower specific strength, as summarized in Table S1. To further reduce the density of our samples, we leverage the 3D printability of our ink to design structures that include mm sized pores. We 3D print cubes with diameters of 10 mm \(\times\) 10 mm \(\times\) 5 mm possessing a controllable porosity ranging from the 100 nm up to the mm length scale, as shown in Figure S11. Indeed, these structures are as light and strong as avian beaks, as summarized in Fig. 5b. Note that even though we did not use monodisperse emulsion drop templates to fabricate these hierarchical structures, our method offers the opportunity to do so. Thereby, it would enable the design of structures possessing ordered and well-defined pores or gradients in pore sizes that possess even higher strengths and stiffnesses, as can be found, for example, in cuttlebone41.
The mechanical properties of CaCO3-based composites depend on the CaCO3 structure, which can be tuned with appropriate additives present during the mineral formation. To assess the influence of additives on the CaCO3 polymorph that forms within our porous scaffolds, we add magnesium chloride (MgCl2) to the solution used to initiate the mineralization of the capsule templates and perform X-ray diffraction (XRD) on the resulting samples. In the absence of any additives, we obtain calcite, as indicated by the characteristic rhombohedral morphology of the crystals in Fig. 6a,b. This result is confirmed by the synchrotron XRD trace for molded bulk samples in Fig. 6e. By contrast, if mineralized in the presence of Mg2+, we primarily obtain aragonite, as indicated with XRD in Fig. 6e and the nanograin morphology seen on SEM images in Fig. 6c,d. We also get extra phases of magnesite or calcite when Mg2+ is not uniformly distributed in the ink formulation while trona comes from the printing bath solution when samples are not completely rinsed. If we mineralize our samples in the presence of an organic additive, polyacrylic acid (PAA), we obtain a mixture of calcite and vaterite, as demonstrated in the XRD in Fig. 6e. Traditional XRD confirms these results from molded samples as seen in Figure S12. This result is consistent with results obtained by forming these minerals in bulk49,50, suggesting that knowhow on polymorphic control of CaCO3 formation in bulk aqueous solutions can be transferred to our process.
2D scanning synchrotron XRD allows us to determine the spatial distribution of different crystal phases in printed lattice samples, as illustrated in Figure S13. In accordance with the results on bulk samples and our XRD results, a mixture of calcite and vaterite is obtained throughout the scaffold, if the mineralization is conducted in the presence of PAA, as shown in Fig. 6f,g. In addition, we detect crystalline PVA from the polymer matrix and NaCl as a side product of the mineralization, as shown in Figure S14.
To assess the influence of additives on the mechanical properties, we measure the compressive strength of composites containing different CaCO3 polymorphs. Samples formed in the presence of Mg2+ have a 48% higher average strength compared to those formed in the absence of any additives, as shown in Figure S15. We assign the higher strength of the aragonite-based composites to their comparatively high inorganic content and the higher stiffness of aragonite compared to that of calcite and vaterite51. By contrast, samples formed in the presence of PAA have a 14% lower average strength compared to samples without additives.
To demonstrate the versatility of our approach, we form different types of porous minerals. Pyrogallol can complex a broad range of metal ions52–54, including Ca2+ and Fe3+. To leverage this ability, we ionically crosslink the pyrogallol-functionalized surfactants located at the emulsion surface with Fe3+ by exchanging CaCl2 contained within the initial aqueous solution with FeCl3 before we initiate the formation of Fe(OH)3 at the emulsion surface. The resulting capsules are processed into inks that are subsequently printed into an aqueous solution containing Na2CO3, using the protocols established for the formation of CaCO3. The resulting Fe(OH)3 composite possesses a microstructure that is similar to that observed for CaCO3, as shown in Figure S16 and Movie M2. We confirm the structure of the mineral phase using XRD, as shown in Figure S17. To further demonstrate the versatility of our process, we replace the carbonate source in the mineralizing solution with a phosphate source. We immerse pyrogallol functionalized capsules that have been crosslinked with Ca2+ into this mineralizing solution containing ammonium dihydrogen phosphate before the capsules are processed into inks that are 3D printed into macroscopic grids. The resulting mineral is mainly composed of brushite, as verified with XRD and shown in Figure S17.
To illustrate the potential of our material, we produce an artificial toucan bird beak. To achieve this goal, we cast our ink composed of upconcentrated capsules possessing a thin CaCO3 shell into a polymer-based beak mold. The de-molded sample has a macroscopic shape and microscopic structure that closely resembles that of beaks of toucan birds, as shown in Fig. 7a. Our in-situ mineralization approach offers the additional advantage that it enables firm connections between adjacent layers, such as filaments that have been sequentially deposited during a 3D printing process or two separately fabricated grids. For example, two individually printed grids are joined into an integral free-standing structure by bringing the two grids in contact before they are subjected to an additional mineralization cycle. The dried object is free-standing, as shown in Fig. 7b, indicating that the structures are well connected. Indeed, the distribution of calcium at the interfaces is homogeneous, as revealed by Energy-dispersive X-ray (EDX) mapping in Fig. 7c. These results highlight the potential of our approach to 3D print macroscopic porous biominerals that thanks to the firm inter-layer connections, possess a high strength even if printed as multi-layer structures, as demonstrated by the SEM micrographs in Figure S18.