What if we could create custom bone implants that would trigger their own replacement with real bone? Jakus and colleagues have done just this with a promising biomaterial that can be 3D-printed into many shapes and easily deployed in the operating room. Made mainly of hydroxyapatite and either polycaprolactone or poly(lactic-co-glycolic acid), this “hyperelastic bone” can be 3D-printed at up to 275 cm 3 /hour, the authors report. It also promoted bone growth in vitro, in mice and rats, and in a case study of skull repair in a rhesus macaque. Its effectiveness, fast, easy synthesis, and ease of use in surgery set it apart from many of the materials now available for bone repair.

Despite substantial attention given to the development of osteoregenerative biomaterials, severe deficiencies remain in current products. These limitations include an inability to adequately, rapidly, and reproducibly regenerate new bone; high costs and limited manufacturing capacity; and lack of surgical ease of handling. To address these shortcomings, we generated a new, synthetic osteoregenerative biomaterial, hyperelastic “bone” (HB). HB, which is composed of 90 weight % (wt %) hydroxyapatite and 10 wt % polycaprolactone or poly(lactic-co-glycolic acid), could be rapidly three-dimensionally (3D) printed (up to 275 cm 3 /hour) from room temperature extruded liquid inks. The resulting 3D-printed HB exhibited elastic mechanical properties (~32 to 67% strain to failure, ~4 to 11 MPa elastic modulus), was highly absorbent (50% material porosity), supported cell viability and proliferation, and induced osteogenic differentiation of bone marrow–derived human mesenchymal stem cells cultured in vitro over 4 weeks without any osteo-inducing factors in the medium. We evaluated HB in vivo in a mouse subcutaneous implant model for material biocompatibility (7 and 35 days), in a rat posterolateral spinal fusion model for new bone formation (8 weeks), and in a large, non-human primate calvarial defect case study (4 weeks). HB did not elicit a negative immune response, became vascularized, quickly integrated with surrounding tissues, and rapidly ossified and supported new bone growth without the need for added biological factors.

It is for these reasons that the field has focused on CaP-polymer composites, including those produced via room temperature or hot-melt extrusion–based 3D printing ( 32 – 34 ) or other forms of AM ( 35 , 36 ). Nevertheless, these composites often still have suboptimal material, handling, and/or biological properties: Although composites with CaP, whether in hydrogel ( 37 ) or 3D-printed form ( 32 – 34 ), often have improved stiffness (elastic and compressive moduli) over pure polymers and increased mechanical elasticity or malleability over pure CaPs ( 5 ), the polymeric component often physically encapsulates the bioactive CaP particles, isolating them from the tissue and mitigating their therapeutic potential. In addition, many 3D-printed composites are fabricated with hot-melt fused deposition modeling or laser sintering techniques, which require temperatures greater than 100°C. This high-temperature processing precludes direct incorporation of biological molecules or factors ( 33 ) and is too slow for mass fabrication, with linear deposition rates not greater than 5 mm/s or volume deposition rates not greater than 1 mm 3 /s. Although these composites often do not undergo brittle fracture, their bioactivity is often limited, requiring surface modification with costly biomolecules or other factors ( 38 – 40 ), which also complicate regulatory approval and translatability. Here, we report a new synthetic osteoregenerative biomaterial, which we have called hyperelastic “bone” (HB), that avoids the technical, surgical, and manufacturing limitations of current bone graft materials.

Over the past two decades, much effort has been applied toward creating porous bioceramic and bioceramic composite constructs—from the traditional freeze-casting and gas-foaming methods that result in heterogeneously porous foams ( 13 – 20 ) to the emerging wide variety of additive manufacturing (AM) and three-dimensional (3D) printing technologies that can be used to create ordered porosity and user-defined constructs. These newer technologies have been used to make an array of CaP-based porous materials and structures, including some that are made of synthetic hydroxyapatite (HA), a highly bioactive ceramic ( 21 , 22 ). These have been fabricated with established ceramic AM methods ( 23 ), such as particulate slurry extrusion or powder bed–based inkjet binding followed by sintering ( 24 – 27 ), which require additional high-temperature processing steps and result in brittle final products unfriendly to surgical applications. This high-temperature processing is also incompatible with the direct incorporation of agents such as antibiotics ( 28 , 29 ), growth factors, and other advanced organic-based therapeutics ( 30 ), directly into the material, which can be used to enhance and hasten tissue integration and regeneration ( 31 ).

Synthetic and naturally derived CaP-based ceramics have been commonly used to treat bone defects due to their inherent bioactivity ( 2 , 6 , 7 ). However, the same stiff mechanical characteristics of ceramics that are partially responsible for this osteogenic bioactivity make their surgical implementation challenging or inefficient. For example, porous ceramic constructs cannot be easily shaped and resized by surgeons on demand to better accommodate the defect site ( 8 ). Additionally, ceramic constructs are not amenable to minimally invasive surgical methods because they cannot be deformed without failing through fracture. Therefore, CaP-based products have different formulations—such as granules, malleable putties, or injectable cements created with plasticizers, monomers, and water—so that they can be packed into the surgical sites ( 9 , 10 ). Nevertheless, these approaches suffer from deficiencies. Packed granules or putties, as well as injected bone cements, are often washed away because of intraoperative bleeding within the defect site ( 10 ). Putties that harden upon placement often transfer significant heat as a result of the chemical curing process ( 11 ) and can damage surrounding tissues. The most important is the fact that CaP granules and putties have minimal interconnected porosity ( 10 , 12 ), which hinders host-implant integration, vascularization ( 12 ), and patient recovery while also increasing susceptibility to inflammation, infection, and revision surgeries.

The quest to discover an ideal bone graft material that is compatible with the extensive variety of osseous tissue–related medical indications has been an ongoing challenge. Although there are numerous clinical products used today as bone void fillers or temporary scaffolds ( 1 ) and an extensive body of literature that reports potential new bone-related biomaterials, these products and their surgical implementation suffer from substantial technical, surgical, and manufacturing/scaling shortcomings ( 2 – 4 ). First, an ideal bone graft material is safe and efficacious in vivo (that is, it should not elicit a strong immunoinflammatory host response or cause other undesirable biological responses while repairing and regenerating damaged or missing bone). Second, the material must be rapidly and easily deployed in the operating room by surgical teams ( 5 ). Last, to be translatable, it must be able to be produced at the relevant scales (construct size and number) and widely available at comparable or lower cost than existing clinical products, although this may be difficult to define and quantify at the research stages ( 5 ). Many bone-grafting products on the market and reported in the literature, often composed of calcium phosphates (CaP), polyesters, or composites of the two, excel at one or two of these criteria, but none excel at all three.

RESULTS

Examples of HB ink design, 3D printability, and versatility We synthesized HB particle–laden liquid 3D inks by combining ceramic powder (HA) and polycaprolactone (PCL) or poly(lactic-co-glycolic acid) (PLGA) (9:1 by weight) to produce hydroxyapatite-polycaprolactone (HAPCL) or hydroxyapatite-poly(lactic-co-glycolic acid) (HAPLGA), respectively, in a trisolvent mixture comprising excess dichloromethane (DCM; evaporant), 2-butoxyethanol (2-Bu; surfactant), and dibutyl phthalate (DBP; plasticizer). We recently illustrated that these new 3D ink systems containing biomedical elastomers can be used to rapidly 3D-print graphene inks into highly electrically conductive structures comprising many layers that exhibit strong neurogenic properties and excellent handling characteristics (41). We have also recently adapted this process to create 3D printable metal oxide (ceramic) and metallic inks, which could be thermally processed into an extensive variety of metals and alloys (42). We selected PLGA and PCL as the binders for HB inks because of their extensive use in medicine and tissue engineering (43–46). With these particular elastomers, we 3D-printed solid HB structures comprising many hundreds of layers (movie S1) from liquid inks (Fig. 1A, inset). The resulting structures did not require further postprinting processing (other than rinsing and sterilization) before use and exhibited mechanical and physical properties that permitted further manipulation. For example, a 3D-printed HAPLGA sheet could be rolled, folded, and cut (Fig. 1A and movie S2) to create architectures that might otherwise not be possible to 3D-print directly because of the large, unsupported overhangs. An example of how HB could be used surgically is illustrated in Fig. 1C, where HB cylinders of various sizes were 3D-printed (inset) and the correct size was selected. They were snugly slipped onto the terminal regions of soft tissue human tendon allograft, cut to size, and sutured to the graft. An augmented graft such as this could be used in arthroscopic procedures for replacing damaged anterior cruciate ligaments (ACLs), where it may accelerate and strengthen ligament-to-bone healing within bone tunnels after ACL reconstruction (47). Fig. 1. Versatility, scalability, and manipulation of 3D-printed HB. (A) Easy to synthesize volumes (~100 ml shown) of liquid-based HB inks (inset) can be 3D-printed into a variety of structures: 3D-printed 12 × 12–cm HAPLGA sheet comprising three layers, which can be manipulated in a variety of ways, including rolling, folding, and cutting. Origami methods may be used to create complex folded structures, whereas kirigami methods can produce complex structures from strategic folding and cutting. (B) Full-scale, anatomically correct parts, such as a human mandible, comprising >250 layers, can be designed, 3D-printed from HAPLGA, and washed to rapidly produce a ready-to-implant object. Final image shows 3D-printed mandible next to an adult cadaveric human mandible. (C) Photograph series illustrating that custom-sized HAPLGA sleeves can be snuggly stretched around, cut, and sutured to a soft tissue, such as human cadaveric tendon, facilitating arthroscopic ACL repair and replacement surgery. (D) Independently 3D-printed HAPLGA miniature-scale versions of a human skull, skull cap, mandible, and upper thoracic seamlessly fused together to create highly complex structures by using HB ink applied to points of contact. (E) Black light–illuminated optical photographs of the outside and internal cross sections of HAPLGA fiber with (top) and without (bottom) incorporated recombinant green fluorescent protein (rGFP). We also rapidly 3D-printed HB inks into anatomically scaled, patient-specific grafts, such as an adult human mandible (Fig. 1B). We achieved linear print speeds as high as 15 cm/s (the hardware limits of the instrument), with no drying time required before handling the completed object. This linear deposition rate equates to volume deposition rates as high as 275 cm3/hour, from a single nozzle, or total object(s) volume of 550 cm3/hour when the architecture was 50% porous. We were able to 3D-print, wash, sterilize, and prepare for use this mandible in less than 3 hours. We custom-made HB into complex scaffold designs such as porous long-bone sleeves, which could both stabilize a long-bone fracture and promote regeneration across the defect length (fig. S1A). Additionally, unlike many other 3D printable material systems, HB inks could be used as self-adhesives, allowing independently 3D-printed objects made of the same or similar materials to be seamlessly fused together. Individually 3D-printed components were merged to form highly complex geometries, which would be impossible to 3D-print as one monolithic object (Fig. 1D). Furthermore, we used HB inks as flexible coatings on other implantable materials, such as metallic screws (fig. S1B); this application is similar to adding bioceramic coatings to these screws, a process shown to improve tissue integration (31). Last, the ability to synthesize and 3D-print HB inks under ambient conditions with no need for further sintering or chemical cementation allowed incorporation of biological factors and molecules, such as proteins (Fig. 1E), peptides, genes, and antibiotics (fig. S1C) (48, 49), which may enhance tissue regeneration and reduce infection.

HB microstructure To better understand the unique characteristics of HB, we compared fibers created with different ink formulations and printing methods (hot-melt extrusion versus solvent-based room temperature extrusion) (Table 1). We examined the microstructures of (i) HA/PCL (1:1 by weight) hot-melt–extruded fiber; (ii) HA/PCL (1:1 by weight) room temperature, trisolvent-based printed fiber; (iii) HA/PCL (9:1 by weight ) room temperature printed fiber using only DCM as the solvent; and (iv) HA/PCL (9:1 by weight) room temperature, trisolvent-based printed fiber (HAPCL) (Fig. 2, A to D). Hot-melt fibers, which have been used in tissue engineering research (32, 50), are characterized by slow extrusion rates (0.5 to 5 mm/s) and dense polymer matrices encompassing the HA particles (Fig. 2A and fig. S2). This is in contrast to the solvent-based, room temperature 3D-printed fiber microstructures, which are rough and have a nano- and microporous architecture (Fig. 2, B to D). Furthermore, there are differences in the elastomer matrix morphology: The elastomer in the trisolvent-based printed fibers forms a smooth continuous matrix joining the particles (Fig. 2, D and G), whereas the single-solvent fibers are characterized by a fine web-like network of elastomer joining adjacent HA particles (Fig. 2, C and F). Table 1. Summary of formulations, material preparation, 3D printing process, and additional characteristics of the HA-polymer composite systems discussed throughout this work. The two materials, HAPCL and HAPLGA, are collectively referred to as HB. 3DP, 3D printing. View this table: Fig. 2. Microstructural characteristics of HB and related 3D-printed systems. (A to D) Scanning electron microscopy (SEM) micrographs of representative fibers produced by HA/PCL (1:1 by mass) hot-melt (A), HA/PCL (1:1) room temperature solvent mixture (B), HA/PCL (9:1) room temperature with DCM only (C), and HA/PCL (9:1) with a trisolvent (HAPCL) (D). (E) Schematic representation of proposed HA and elastomer distribution within fibers with single- or graded-solvent mixtures, as a function of time after extrusion. Higher-magnification SEM micrographs of DCM solvent only (F) and HAPCL microstructures (G). Details regarding material compositions and preparations can be found in Table 1. We hypothesize that the microstructural differences between the single- and trisolvent fibers are a result of the interplay between solvent evaporation and polymer condensation. In both cases, inks begin as particles homogeneously dispersed within elastomer-rich solutions (Fig. 2E). Immediately upon extrusion (t = 0), local particle density increases as a result of shearing forces (51). Once exposed to air, most of the DCM rapidly evaporates (t 1 ). In the single-solvent system, composed only of DCM, this results in rapid precipitation of all dissolved elastomer homogeneously throughout the fiber volume, resulting in a crater-web microstructure (Fig. 2F), which is characteristic of extreme “solvent popping” commonly observed in paints and other coatings when the suspending solvent evaporates too quickly. This leaves solid surface films that trap remaining solvent beneath, which eventually vaporizes and forcefully emerges through the solidified surface in the form of popping bubbles. The resulting thin elastomeric webbing presents concentrated regions of high stress and results in brittle structures that are unable to absorb significant loads without failing. In the trisolvent system, however, the two additional low-volatility solvents slow elastomer precipitation, permitting it to preferentially coat particles (t 1 ). Enough elastomer precipitates onto the particles to form robust interparticle bridges, whereas the inability for the spherical HA particles to densely pack results in interparticle pores (Fig. 2G). Over the course of several minutes (t 2 ), the remaining minority solvents evaporate, slowly precipitating the last of the solubilized material onto previously precipitated elastomer (52). The retention of the two minor solvents after initial DCM evaporation also immediately enables adjacent fibers and layers to fuse during 3D printing. This likely mitigates interlayer delamination and results in monolithic objects that can be handled immediately after being 3D-printed. HB mechanical properties. Although not as elastic as their pure polymer counterparts (fig. S3B), both HAPLGA and HAPCL retain a high degree of elasticity, capable of undergoing 36.1 ± 4.3% and 61.2 ± 6.4% strain and having similar tensile elastic moduli of 4.3 ± 0.4 MPa and 10.3 ± 1.3 MPa, respectively. Hot-melt–extruded and DCM-only solvent-based materials were too brittle to be accurately tested under tension. Additionally, DCM-only solvent-based inks were exceptionally difficult to 3D-print into high-fidelity multilayered structures because the extruded fibers dried too quickly and did not adhere well to previously deposited material, making multilayer 3D printing particularly difficult. The microstructural characteristics of HB permitted fibers to undergo various modes of deformation while being able to recover almost completely upon unloading (Fig. 3A and fig. S4). In all instances, macroscopic deformation resulted from the elastomer matrix straining under loads, pulling and pushing embedded particles along with it. Porosity within the fibers enabled rigid particles to translate while limiting direct, incompressible interactions with each other. Upon compressive loading, excess pore space was eliminated as particles flowed with the straining elastomer to fill the open volume. Tensile loads were carried almost entirely by the elastomer, and under extreme strains, temporary separation between the elastomer and particle surfaces occurred (fig. S4E). However, because the HA particles were physically encapsulated within the elastomer and not covalently bound to it, these interfacial tensile voids were not permanent and disappeared upon unloading. The elastomer produced antiparallel restoring forces upon unloading, which manifested itself as a macroscopically observable elastic response (large, recoverable deformation), with the HB returning to near-net shape over many cycles (fig. S4, D and G). For porous HB constructs that were 3D-printed into defined architectures, the previously defined compression, tension, and bending deformation modes were combined to impart elastic properties throughout the entire construct. Although the geometry and porosity of the 3D-printed object affected the ultimate mechanical behavior, simple 90° cylinders (printed fibers oriented perpendicular to adjacent layers) could be cyclically compressed up to 40% strain and rapidly returned to near-original form immediately after each cycle (Fig. 3A), with full recovery occurring over the course of minutes (fig. S4H). This behavior was not limited to quasi-static loading but is also evident under dynamic loading, such as a hammer impact; 3D-printed HB constructs, despite being composed of 90 weight % (wt %) ceramic, did not shatter, catastrophically fail, or permanently deform under high-impact loads (unlike hot-melt printed samples) but, rather, rebounded to their original form (movie S3). Fig. 3. HB mechanical properties. (A) Photograph series showing the compression and recovery of a 1-cm-diameter 3D-printed HAPLGA cylinder over a single compression cycle. (B) Digital representation of average adult human femur and corresponding femoral midshaft section longitudinal and axial views. Axial (C) and longitudinal (D) views of 3D-printed HB femoral midshaft construct using digital file shown in (B). (E) Longitudinal compressive loading profile of HB femoral midshaft (D) and corresponding photographs at indicated percent strain points. Plastic deformation of HB femoral midshaft begins at 2 (10.3% strain) and proceeds to buckle and barrel (3 and 4). Cyclic compression loading profile (10 cycles) of HB femoral midshaft loaded in axial direction (C) in strain domain (F) and time domain (G). (H) Photograph series of a single axial compression cycle displayed in (F) and (G) and the corresponding percent strain. Additional characterization of HB mechanical properties can be found in the Supplementary Materials. Although there are numerous bone-grafting indications that would not require HB to be under direct and substantial mechanical loads, such as those related to the craniofacial, torso, upper spinal, and upper extremity regions, it is nonetheless important to investigate the loading limits of HB in an anatomically scaled scenario, such as those experienced in the adult femur. To do so, we 3D-printed a 4-cm-long (135 layers) section of an anatomically correct HB midfemoral construct from a digital file (Fig. 3B) with a 600-μm nozzle. The construct was given a hollow shaft and ~25% porous cortex (Fig. 3, C and D), emulating the natural interior and cortical regions of natural femoral bone, and mechanically compressed in both the axial and longitudinal directions. Under uniaxial longitudinal loading, the femoral HB construct was capable of supporting about 650 N (~150 pounds) before the onset of plastic deformation at 10% strain (Fig. 3E, 2), at which point the construct begins to buckle (Fig. 3E, 3), barrel (Fig. 3E, 4), and fracture (movie S4). Despite the fact that the solid volume fraction was only ~25% of the total HB construct functional volume (~75% from hollow interior and porous cortical walls), the 3D-printed HB femoral midshaft can support about two to three times greater mechanical loads than existing resorbable osteoconductive bone cements and press-fit osteochondral autograft (53) and approaches the stiffness and load-bearing capacity of human cortical bone (54). In addition to being both mechanically strong and stiff in the longitudinal loading direction, the HB femoral construct remained compliant and elastic under axial compressive loading (Fig. 3, F to H), capable of rapidly undergoing numerous compression cycles (up to 25% strain) without permanent deformation (movie S4). These highly anisotropic properties, which can be influenced by the interior pore design, illustrate that anatomically scaled HB constructs can sustain anatomical loads in one direction while remaining elastic and compliant in the orthogonal loading directions, permitting the constructs to be deformed by hand (fig. S5 and movie S4), an ideal characteristic for fitting 3D-printed bone graft into an osseous defect (5). These results indicate that HB, in addition to being translationally relevant to craniofacial and other nondirect load-bearing indications, may have potential applications for treating defects that experience substantial mechanical loading. Like any other implant or graft surgically positioned into a defect site, however, time must be allotted to permit the region to heal, graft to integrate, and, in the case of HB, to ossify, before substantial mechanical motion and loading. HB physical properties. Comparison of the measured dry and theoretical solid densities of HAPCL and HAPLGA fibers (the weight for a given volume if all space was filled: 90 wt % HA and 10 wt % polymer) revealed that HAPCL and HAPLGA fibers are about 50% porous (Fig. 4A and Supplementary Materials and Methods) (55). Once saturated with water, the density increases by 0.5 g/cm3, indicating that the porosity is open and accessible. This accessible porosity is a vital characteristic for promoting nutrient diffusion, cell migration and viability, and tissue integration. These microstructural characteristics also impart hydrophilicity and high liquid absorbency (Fig. 4B) to the HB, which is ideal for enhanced cell, bioactive factors, and nutrient infiltration. In contrast to other 3D-printed HA composite systems, which do not exhibit significant absorbance of fluids, the surfaces of HAPCL and HAPLGA are dominated by exposed HA (Fig. 4, C and D), as indicated by the dark red coloring after Alizarin Red S calcium-specific staining, a property distinct from that of HA-containing hot-melt printed scaffolds (Fig. 4E), trisolvent (1:1 HA/PLGA by weight) printed scaffolds (Fig. 4F), and even human allograft bone (Fig. 4G), which has less CaP at the surface. Fig. 4. Physical properties of HB. (A) Dry and wet densities of HAPCL and natural bone. Asterisk denotes value from literature (53). Upward arrow indicates that the wet density is expected to increase in vivo as water is replaced with proteins and tissues. (B) Time series contact angle profile of water on solid HAPCL surface (top) and small volume of red-colored water being injected into the end of a complex 3D-printed HAPLGA object (bottom). (C to G) Alizarin Red S–stained and Alizarin Red S–unstained (insets) photographs of HAPCL (C), HAPLGA (D), 1:1 HA/PLGA hot-melt (E), and 1:1 HA/PLGA room temperature solvent scaffolds (F) and human cadaveric bone (G). (H) Thermogravimetric profile of as-printed, water-washed, and 70% ethanol–washed HAPLGA scaffolds. Expected evaporation or decomposition temperature ranges for individual components indicated. Note that the PLGA decomposition temperature range and the DBP boiling temperature overlap. Thermogravimetric analysis (TGA) (Fig. 4H) revealed that HB constructs contained as much as 15 to 20 wt % residual solvents immediately after 3D printing. These solvents are biologically toxic and must be removed before applying the HB to any biological system. Rinses with deionized water are effective at removing DCM, but not the remaining two solvents. However, a 20- to 30-min rinse in 70% ethanol not only removes all residual solvents but also has a sterilizing effect (56), although this is not a clinically approved sterilization process. Gamma irradiation is a U.S. Food and Drug Administration–approved process commonly used to sterilize polyester and CaP products before clinical use and is likely the best translational sterilization approach for HB. The TGA results validate that final, rinsed, and dried constructs, although 50% porous, contain 90 wt % HA, which is the only component that does not decompose during the TGA process (horizontal dashed line in Fig. 4H).

In vitro human mesenchymal stem cell behavior on HB scaffolds To evaluate the regenerative potential of HB, we performed in vitro experiments to assess the ability of 3D-printed HAPLGA and HAPCL scaffolds (Fig. 5, A and B) to support human mesenchymal stem cell (hMSC) adhesion and proliferation and induce osteogenic differentiation and function in the absence of osteogenic growth factors. hMSCs seeded on both HAPLGA and HAPCL scaffolds proliferated over the course of 14 days (Fig. 5, C to E) to fill the entire scaffold volume, after which time (day 28) the cell number was maintained but did not significantly increase, presumably a result of lack of space within the construct. Thus, both HAPLGA and HAPCL support human stem cell viability and proliferation. Within HAPLGA scaffolds, alkaline phosphatase (ALP) activity, an early marker of osteogenesis (57), initially decreased during proliferation but increased substantially between days 14 and 28 (Fig. 5F), as expected for cells that proliferated and then underwent differentiation once proliferation plateaued (57). Analysis of the expression of pro-osteogenic genes further supports the notion that HB promotes differentiation. hMSCs cultured on HAPLGA showed significant up-regulation of the pro-osteogenic genes, collagen type I (~15-fold increase) (58), osteopontin (~20-fold increase) (59, 60), and osteocalcin (~350-fold increase) (61) by day 28 (Fig. 5G). This was accompanied by extensive extracellular matrix (ECM) and nanocrystalline HA synthesis and deposition (fig. S6, H and I). These crystals had an atomic Ca/P ratio of 1.69 (Fig. 5H), which is similar to that of natural human bone (1.65 to 1.69) (62) and different from that of the synthetic HA used to create HB, which was measured in the same scaffolds to be 1.59. This indicates that these HA nanocrystals are likely synthesized and deposited by the cells and that HB is inherently osteoinductive. Fig. 5. In vitro evaluation of 3D-printed HB scaffolds with MSCs. Photograph of 90° offset HAPCL scaffold (A) and the corresponding cross-sectional SEM micrograph highlighting the offset architecture between layers (B). (C and D) (Top) Top-down view, laser-scanning confocal reconstructions of live (green; calcein AM) and dead (red, ethidium homodimer-1) stains. (Bottom) Corresponding cross-sectional SEM micrographs of HAPCL samples 7 days (C) and 28 days (D) after seeding with hMSCs. (E) PicoGreen quantitation of DNA in HAPLGA and HAPCL scaffolds at indicated time points after seeding hMSCs. Values were normalized to average DNA measured on day 1 (n = 3 for all time points; error bars refer to SD). (F) ALP activity normalized to corresponding DNA content from HAPCL scaffolds at indicated time points (n = 3 for all time points; same samples as those used for DNA quantification; error bars refer to SD). (G) Gene expression levels of osteogenic-relevant transcripts in hMSCs grown on HAPCL scaffolds, normalized to sample-specific GAPDH (glyceraldehyde-3-phosphate dehydrogenase) values, followed by normalization to day 0 hMSC values (n = 3; error bars refer to SD). (H) Atomic Ca/P measured in the HA of the scaffold itself (Scaffold HA) and of the nanocrystals within the ECM formed after in vitro culture. Gray box, Ca/P range of natural HA. (E to G) *P < 0.05, over previous time point for the same group. Confidence intervals (P values; two-tailed, equal variance t tests) are as follows: for DNA quantification of HAPCL samples: days 1 and 7, 0.0016; days 7 and 14, 0.013; days 14 and 28, 0.152; for DNA quantification of HAPLGA samples: days 1 and 7, 0.0032; days 7 and 14, 0.043; days 14 and 28, 0.862; for ALP/ng DNA of HAPCL samples: days 1 and 7, 0.034; days 7 and 14, 0.0026; days 14 and 28, 0.0053; and for fold increase in gene expression of HAPCL samples: osteopontin: days 7 and 14, 0.00041; days 14 and 28, 0.00026; collagen I: days 7 and 14, 0.0033; days 14 and 28, 0.017; osteocalcin: days 7 and 14, 0.0049; days 14 and 28, 0.050. Additional in vitro related figures for hMSCs seeded onto 30° advancing angle HB scaffolds can be found in the Supplementary Materials.

HB in vivo biocompatibility Although both HAPLGA and HAPCL can be successfully 3D-printed into defined structures, HAPLGA was significantly easier to print than HAPCL, specifically in fabricating large (many dozen layers) objects as well as conforming to tight, organically shaped contours. HAPLGA also had superior mechanical properties and could be elastically strained upward of 60% (tensile) and 50% (compression); HAPCL, in contrast, could strain upward of 35% (tensile) (fig. S3B) and 50% (compression). For these reasons, we evaluated the in vivo biocompatibility and osteogenic function of HAPLGA; however, because PCL has been extensively used in the bone and tissue engineering field (43), it is a reasonable assumption that HAPCL would show similar biocompatibility in vivo, although long-term degradation times would likely vary from HAPLGA. To examine biocompatibility in vivo, we subcutaneously implanted HAPLGA scaffolds (that is, HB) in female BALB/c mice. HA/PLGA (1:1) hot-melt scaffolds were also implanted as a comparison, because they are similar to materials that have been previously used and evaluated as bone implants (32). After 7 days, tissue had already begun to infiltrate and vascularize throughout the HB scaffolds (fig. S7A). The hot-melt scaffold explants could not be histologically processed successfully, a result of their highly brittle nature and the dissolution of the PLGA (majority scaffold component) in histological process solvents, which caused the loss of integrated tissues and embedded HA particles during processing (Fig. 6F). After 35 days in vivo, the HB scaffolds were completely integrated with the surrounding host tissue (Fig. 6, A to C). SEM imaging of explanted scaffold tissues revealed that, in both material systems, the tissue formed intimate contact with the material within and throughout the scaffold volume by day 35. However, there was a distinct difference in the structure and texture of the tissue within HB (Fig. 6, D and E) compared to that within the hot-melt printed scaffolds (Fig. 6, G and H). Tissue surrounding HB more closely mimicked healthy ECM, with defined collagenous ECM (fig. S8) and blood vessels ranging from 2-μm single-cell capillaries to multihundred-micrometer vessels present throughout the scaffold (Fig. 6, B to D, and fig. S8). In contrast, the tissue within the 1:1 HA/PLGA hot-melt scaffolds was characteristically dense and relatively acellular compared to the HB counterparts (Fig. 6, G and H). SEM imaging also revealed a population of unhealthy blood cells (burr cells) inhabiting the tissue within the 1:1 HA/PLGA scaffolds (Fig. 6I), which may indicate a strong local fibrotic response. Additional staining with Alizarin Red S did not indicate any obvious mineralization within the integrated tissues by day 35 (fig. S7C); however, because there was no source of osteoprogenitor cells in this subcutaneous model, de novo mineralization was not expected. Fig. 6. Biocompatibility evaluation in vivo with a mouse subcutaneous implant. (A) Gross hematoxylin and eosin (H&E) histological image of day 35 HB explant cross section. Blue arrows, HB fiber cross sections. The densely packed HA particles within the HB stain dark purple to black. (B and C) H&E histological image of day 35 explant with HB fiber cross sections (dotted yellow circles), vessels (dashed yellow circles), and capillaries (solid yellow circles). (C) Higher-magnification section of (B). (D) SEM micrograph of day 35 explanted HB scaffold. Several vessels are indicated by dotted yellow circles; soft tissue completely fills the space between the HB fibers. (E) Higher-magnification SEM micrograph of (D) highlighting the structure of HB, new tissue, and the HB scaffold-tissue interface (dashed line). (F) Gross H&E histological image of day 35 hot-melt 1:1 HA/PLGA–explanted scaffold. Because of the solubility of PLGA in common histological solvents, and given that more than half the scaffold volume is composed of PLGA, the 1:1 HA/PLGA scaffold materials and incorporated tissues did not survive histological processing. (G) SEM micrograph of the cross section of day 35 hot-melt 1:1 HA/PLGA–explanted scaffold and tissue. No vessels are visible. (H) Increased magnification SEM micrograph of (G) illustrating the hot-melt HA/PLGA–tissue interface (dashed line). (I) Representative burr cell; these were found throughout the tissue within the hot-melt 1:1 HA/PLGA–explanted scaffolds. Additional related figures can be found in the Supplementary Materials.

HB efficacy in a rat spinal fusion model To assess the capacity of HAPLGA HB to induce bone regeneration in a preclinical in vivo model, we evaluated HB scaffolds with and without incorporated recombinant human bone morphogenic protein 2 (rhBMP-2) in a previously described rat posterolateral spinal fusion (PLF) model (63, 64). The PLF model was chosen because it replicates the spinal fusion procedure in humans, and it is widely accepted to assess and predict regenerative potential in the spinal fusion setting (65). Although the amount of osteoinductive stimulus necessary to achieve fusion is different in the rat than in the human (primarily because of the difference in anatomical size), it is useful because it is cost-effective and reproducible, with established positive and negative controls (66). HB scaffolds were implanted bilaterally at the L4 and L5 transverse processes, which had been decorticated immediately before. As a control, an equivalent amount of HA powder (42 mg, identical source and mass of HA that was used to create each HB scaffold) was implanted. A demineralized bone matrix (DBM) scaffold in sheet form (Bacterin) was also used as a control, because DBM scaffolds routinely elicit a 50 to 60% per-animal fusion rate in this model (64). Autografts promote spinal fusion in humans and, when used in higher-order animal models, elicit fusion rates of ~50%; autografts are not used as standard positive or comparative controls in the rat PLF model, however, because the limited volume of obtainable bone is insufficient to reproducibly promote successful fusion (67). Radiographic imaging of each spine shows placement of the scaffolds at the L4 and L5 transverse processes (Fig. 7A). Scaffold placement relative to spinal features can be also observed in the gross photograph of a spinal column cross section explanted after 4 weeks in vivo (Fig. 7B); after 4 weeks in vivo, the HB scaffolds became infiltrated by and integrated with the surrounding tissue. Fig. 7. Evaluation of HB in vivo for rat spinal fusion. (A) Representative radiograph illustrating bilateral placement of HAPLGA (HB) scaffolds across the L4 and L5 vertebral body transverse processes. (B) Photograph of the cross section of rat spinal column containing HB scaffolds 4 weeks after implantation. Placement of HB scaffolds on transverse processes (black arrows) is visible, as is significant tissue incorporation into the scaffolds. (C) Fusion scores of spinal segments with scaffolds explanted after 8 weeks. An established scoring system for the fusion score was used, whereby 0 = no bridging bone, 1 = unilateral bridging bone, and 2 = bilateral bridging bone. ACS, absorbable collagen sponge (historical control); DBM, demineralized bone matrix; HA, hydroxyapatite granules; HB, HAPLGA; HB + BMP, HAPLGA preloaded with 1.5 μg of rhBMP-2 before implantation. (D) Fusion rates were calculated as both percent sides fused (two sides per animal) as well as on a per-animal basis [n = 6 for ACS, DBM, HA, HB + BMP groups; n = 12 (two n = 6 replicates) for HB group]. For the latter, unilateral bridging bone was considered successful fusion (fusion score of ≥1.0). (E) Laboratory microCT–based quantification of new bone volume within and surrounding the HB scaffolds (with and without 1.5 μg of rhBMP-2 per scaffold). n = 6 animals for all groups. (F and G) Representative single-slice synchrotron microCT images of the cross section of HB scaffolds without (F) and with (G) rhBMP-2 added. Green rectangles, region enlarged at right; bright white, native bone; speckled black-white, HB scaffold. Error bars for (C) and (E) refer to SD. Confidence intervals (P values; two-tailed, equal variance t tests) for groups shown in (C) are as follows: HB/HB-BMP, 0.0036; DBM/HB-BMP, 0.0004; HB/DBM, 0.21. Confidence intervals (P values; two-tailed, equal variance t tests) for groups shown in (E) are as follows: per-scaffold HB/HB-BMP, 0.00003; per-animal HB/HB-BMP, 0.0003. *P < 0.05, between indicated groups or when compared with previous time point for the same material group. The double asterisk (**) indicates a value of 0 (no fusion or newbone formation wa observed). After 8 weeks in vivo, scaffolds and surrounding tissues were explanted and evaluated for fusion as well as for new bone formation via manual palpation and laboratory synchrotron–based micro–computed tomography (microCT), respectively. The mean fusion score of the HB group was significantly higher than that of a collagen scaffold historical control (1.2 versus 0) (63), as well as the HA powder control, which also had a mean fusion score of 0. Moreover, the mean fusion score of HB group was statistically equivalent to that achieved with a DBM control scaffold (1.2 versus 0.7; Fig. 7C). The fusion rate (as determined by the number of sides successfully fused, with two sides possible per animal) was also significantly higher in the HB group relative to the collagen historical negative control and HA powder groups (46% versus 0 and 0%) and was comparable to the DBM group (25% of sides successfully fused; Fig. 7D). These results indicate that the HB scaffold has intrinsic osteoinductive capacity sufficient to induce bone regeneration and successful spinal fusion in the rat without any added growth factors and shows similar osteogenic capacity not only to the DBM used here but also to other commercially available DBMs that have been evaluated in this animal model (64). We also found that HB can serve as an effective carrier for the delivery of rhBMP-2. When HB scaffolds (~45 mg) were preloaded with 1.5 μg of rhBMP-2 (0.003% loading by scaffold weight), the mean fusion score was significantly higher than that of the HB alone group (1.8 ± 0.4 versus 1.0 ± 0.7; Fig. 7C), and the fusion rate was similarly elevated (83% versus 46%; Fig. 7D). The host bone volume in the L4 and L5 transverse processes was quantified in four control animals outside of the study groups and averaged (267 ± 32 mm3) using laboratory microCT. The volume of new bone formed around the HB scaffolds was calculated by subtracting scaffold volume from total bone plus scaffold volume for each left and right implant region in each animal. This resulting value for total bone volume was evaluated against historically determined, natural mean bone volumes for rats of the same strain, age, and weight range. The new bone volume data illustrate that HB alone induces substantial new bone formation, and the rhBMP-2–carrying HB significantly induces more than HB alone (19.5 ± 6.3 mm3 versus 38.9 ± 11.4 mm3; Fig. 7E). High-resolution, synchrotron microCT analysis of HB samples without and with rhBMP-2 revealed that there was substantially more bone formation both around the exterior and throughout the interior of the HB + BMP scaffolds compared to the scaffolds without added rhBMP-2; however, new bone formation within the HB scaffolds without BMP was still apparent (Fig. 7, F and G, and movie S5). These results demonstrate that 3D-printed HB is osteogenic but may be optionally enhanced through addition and absorption (Fig. 4B) of established growth factors, such as BMPs (68–70). Not only is the material as efficacious as DBM, but it is more widely available and substantially cheaper and easier to process and fabricate into complex and patient-specific shapes relative to DBM, which is in limited supply, carries risks for disease transmission, and is subject to variability in efficacy due to the differences in allograft sources.