Abstract

1. Introduction

Bone tissue engineering (BTE) proposes innovative strategies to replace lost or injured bone tissue resulting from various causes, such as degenerative diseases, accidents, aging, or other specific medical conditions. Compared to traditional grafting techniques, which involve extracting bone material from the patient or a compatible donor, BTE offers significant advantages, being a less invasive approach and capable of promoting superior cell adhesion, differentiation, and proliferation¹,².

For this purpose, the combination of different materials plays a fundamental role in the search for effective solutions within BTE. A highly promising approach is bone cements, which should exhibit adequate mechanical strength, good biocompatibility, and radiopacity. Bone cements are basically composed of two phases: a solid and a liquid. Among the cements commonly used for bone-to-bone adhesion, acrylic bone cements (ABCs), calcium sulfate cements (CSCs), and calcium phosphate cements (CPCs) stand out. Of these, CPCs, the focus of the present study, are notable for their self-setting capacity, good osteogenic activity, and osteoconductivity when hydrated with water or another phosphoric acid solution³,⁴. The solid phase of CPCs involves the addition of ceramic materials based on calcium phosphates, also known as bioceramics, such as hydroxyapatite (HAp), tricalcium phosphate (TCP), and biphasic calcium phosphate (BCP), among others. These materials have demonstrated excellent results in bone reconstruction by mimicking the mineral composition of natural bone, creating a porous and crystalline structure that facilitates cellular infiltration and the formation of a biocompatible regenerative matrix⁵,⁶. Synthetic hydroxyapatite (HAp), for instance, stands out and has been widely used as a bone substitute due to its similarity to the mineral matrix of human bone tissues. It possesses key properties, such as good mechanical strength, appropriate color, low thermal and electrical conductivity, osteoconductivity, and a favorable cellular and vascular response without altering inflammation patterns. Moreover, it can react endothermically without the need for high temperatures (thus avoiding issues of cell degeneration), and it has a simple and cost-effective production process⁷,⁸.The liquid phase of CPCs is typically composed of aqueous ionic solutions which, when mixed with the solid phase, form a moldable and injectable paste (suitable for minimally invasive surgeries) that gradually hardens at the desired site. Additionally, CPCs tend to be resorbed due to their osteoclastic activity, promoting, at the same time, the formation of new bone tissue at the bone-implant interface. In other words, CPCs do not act as permanent bone substitutes but rather as temporary agents, gradually being replaced by new bone tissue⁹,¹⁰. Compared to conventional calcium phosphate ceramics in the form of granules or blocks, CPCs offer the advantage of creating a viscous paste that can be molded into complex bone defects or injected into the implant site, thus mitigating inflammatory responses⁹,¹¹,¹².

Polymethyl methacrylate (PMMA), a non-biodegradable synthetic polymer, has been one of the most widely used materials in the composition of bone cements due to its high chemical stability, low cost, and favorable biomedical properties¹³. However, PMMA has potential disadvantages; for instance, its elastic modulus differs significantly from that of porous bone, which can lead to repeated fractures of adjacent vertebrae or cement leakage. Furthermore, PMMA lacks biological activity, preventing its degradation or organic integration with bone, thus posing risks of inflammation and potential failure of the surgical procedure¹⁴. Methylmethacrylate (MMA), the monomer of PMMA, is considered a toxic substance that may remain in the cement due to incomplete polymerization, causing thermal and chemical necrosis in tissues and compromising blood circulation¹⁵,¹⁶. Therefore, the use of biodegradable polymers in the liquid phase of the cement not only provides temporary mechanical support for tissue regeneration, but also facilitates its gradual replacement by new bone tissue over time. These biopolymers, being natural and biodegradable, offer unique properties such as biocompatibility, renewability, easy availability, and cost-effectiveness¹⁷. Common examples of these natural polymers include glycerin, hyaluronic acid, lactic acid, chitosan, carboxymethylcellulose (CMC), alginate, and others. These biopolymers not only ensure good adhesion and rheological compliance of the cement, but also enhance injectability, improving fluidity and adaptation to the environment, thereby amplifying the results of manipulation and in situ application¹⁸. Glycerin, for example, acts as a plasticizer capable of modulating the properties of the cement, improving cohesion between the organic and inorganic phases, adjusting injectability, and increasing setting time. Glycerin is an excellent candidate as a solvent for use in non-aqueous alkaline cement pastes due to its biocompatibility and complete miscibility with water¹⁹. Carboxymethylcellulose (CMC), a biopolymer derived from cellulose, is widely used in the food industry as an edible coating, thickener, stabilizer, emulsifier, and gelling agent²⁰. The aqueous solution of this gum provides important properties for bone cement composition, such as high viscosity and pseudoplastic behavior, aiding in the manipulation and injectability of the paste.

The incorporation of magnetic materials into bone cement composition, whether in the form of magnetic particles (Fe, Co, Ni), oxides (FeO), ferrites (Fe₂O₃, Fe₃O₄, CoFe₂O₄, CuFe₂O₄, MgFe₂O₄), or magnetic glass-ceramics, has become an emerging practice to enable magnetic orientation and precise cell guidance in the regeneration process²¹-²⁶. Among these, magnetic nanoparticles (MNPs) stand out compared to conventional materials, particularly due to their electromagnetic properties, which have great potential for biomedical applications, including magnetic resonance imaging (MRI), cancer therapy, biosensors, and tissue engineering, among others²⁷,²⁸. The use of MNPs in BTE significantly enhances bone tissue regeneration, as they exhibit osteoinductive properties that promote cell growth, proliferation, and the expression of osteogenesis-related genes. In addition, they allow the creation of structures that are responsive to external stimuli, intelligent, and bioactive²⁹,³⁰. Studies show that the interaction between MNPs and cellular tissues stimulates cell growth due to their ability to reduce intracellular H₂O₂ through intrinsic peroxidase activity, accelerating cell cycle progression³¹. Moreover, MNPs can respond to the application of an external magnetic field, enabling directed delivery of cements and medications, uniformly distributing them throughout the biological environment. Additionally, they offer the potential for separating molecules and cells, inducing a series of changes in cellular behavior³²,³³.

Recently, several research efforts have indicated that ferrite MNPs can surprisingly accelerate cell proliferation, adhesion, and differentiation. Among the most studied systems, magnetite (Fe₃O₄) stands out, receiving significant attention due to its biocompatibility, non-toxicity, high rigidity and strength, as well as its large specific surface area and acceptable degradability²¹,³⁴-³⁷. Moreover, these Fe₃O₄ nanoparticles, as well as maghemite (γ-Fe₂O₃) nanoparticles, have already been approved by the U.S. Department of Health and Human Services – Food and Drug Administration (FDA) – for clinical use in the treatment of iron deficiency anemia and as contrast agents in MRI³⁸. Abo-zeid et al.³⁹ leveraged the FDA-approved status of Fe₃O₄ and γ-Fe₂O₃ nanoparticles to conduct a molecular docking study for the treatment and control of infections caused by COVID-19, and they observed a good interaction between ferrite ions and the protein molecules of the SARS-CoV-2 virus, leading to its inactivation. However, the high surface area of these MNPs leads to the formation of materials with superparamagnetic and/or paramagnetic characteristics, presenting greater challenges for their application under physiological conditions⁴⁰.

Another highlight can be given to cobalt ferrite (CoFe₂O₄) MNPs, which rank among the most important members of the magnetic spinel ferrite family due to their excellent properties, such as high chemical stability, moderate to high saturation magnetization (Ms = 3-80 emu.g⁻¹), high coercivity (Hc), high permeability, and strong magnetic anisotropy³⁵,⁴¹,⁴². CoFe₂O₄ nanoparticles below a critical size exhibit superparamagnetic properties at room temperature and have been successfully employed in hyperthermia treatments⁴²,⁴³. Beyond cobalt ferrite's promising physical and chemical properties, the presence of Co²⁺ ions throughout the nanoparticle structure offers therapeutic effects and antibacterial activity. Studies report that Co²⁺ ions can stimulate angiogenesis by mimicking hypoxia in tissue cells and promote osteogenesis by enhancing bone formation and the proliferation of osteoblastic cells, thus confirming their biocompatibility⁴⁴-⁴⁶. Additionally, research highlights cobalt's effectiveness in providing antibacterial and antimicrobial properties, making its application crucial in materials used for biomedical implants, as it effectively minimizes postoperative infections⁴¹,⁴⁷,⁴⁸.

In this context, the development of formulations combining hydroxyapatite (HAp) with hybrids of magnetite (Fe₃O₄@SiO₂) and/or cobalt ferrite (CoFe₂O₄@SiO₂), forming a magnetic calcium phosphate cement (MCPC), stands out as a promising avenue in biomedicine. This approach is supported by the non-toxic, bioactive, and mechanically robust properties of these materials, as well as their nanostructure, which allows for strong attraction to an external magnetic field. Hybrid MNPs can facilitate the development of carrier systems that can be injected intravenously, endovenously, or even grafted in situ, allowing them to be guided to specific locations in the body by applying a magnetic field⁴⁹,⁵⁰.

Building on this context, the aim of this work is to develop different formulations of MCPCs using magnetically activated hydroxyapatite (HAp) with hybrid nanoparticles (MNPs@SiO₂) of Fe₃O₄ and CoFe₂O₄ synthesized by combustion reaction, within an aqueous matrix of glycerin and carboxymethyl cellulose (CMC). In this study, the morphological, magnetic, mechanical, and biological properties of the MCPCs were investigated to evaluate their viability for biomedical applications in bone tissue repair.

2. Materials and Methods

2.1. Materials

To produce the magnetic bone cements, the following materials were used:

• Ferrites hybridized with a 3-aminopropyltrimetoxysilane [APTMS - H₂N(CH₂)₃Si(OCH₃)₃] silane agent, magnetite (Fe₃O₄@SiO₂) and cobalt ferrite (CoFe₂O₄@SiO₂), with particle sizes of 164.72 nm and 50.84 nm, respectively, supplied by the Ceramic Materials Synthesis Laboratory (LabSMaC) of UFCG-PB-Brazil, synthesized according to the methodology proposed by Araújo⁵¹;

• Hydroxyapatite (HAp), with a particle size of 25.66 nm, supplied by LabSMaC and synthesized by precipitation using a phosphorus/calcium ratio of 1.67, according to the methodology proposed by Saeri et al.⁵²;

• Anionic polymer carboxymethylcellulose (CMC) P.A. (-RO-O-O-RO-)n, where R = CH₂CO₂H, with a decomposition temperature of 250°C (NEON);

• Glycerin P.A. [C₃H₅(OH)₃], with a decomposition temperature of 290°C (BIOPACK);

• Dibasic sodium phosphate P.A. (Na₂HPO₄) as a setting accelerator agent (NEON).

2.2. Preparation of cement

The preparation of the cements followed a fixed solid-to-liquid ratio of 30:70, with the solid phase comprising silanized ferrites (Fe₃O₄@SiO₂ and CoFe₂O₄@SiO₂) and hydroxyapatite (HAp), forming the MNP@SiO₂:HAp hybrid, while the liquid phase was a mixture of carboxymethylcellulose (CMC), glycerin (G), and deionized water (H₂O), forming the hydrogel.

The hydrogel or liquid phase, representing 70% of the cement mass, was prepared by mixing a fixed formulation of 1.40% CMC, 13.72% G, and 54.88% H₂O for all cements. Initially, deionized water was preheated in a beaker on a hot plate with magnetic stirring at a temperature of 50 °C. Once this temperature was reached, glycerin was slowly added until a homogeneous mixture was formed. Next, the CMC was added in a powdered form, keeping the mixture under constant stirring for 2 hours at 50°C. By the end of this time, the hydrogel solution, the base for forming the cement, was obtained.

The solid phase of the cement, representing 30% of its mass, had variations in its MNP@SiO₂:HAp formulation with mass percentage ratios (%/%) of 9:21, 15:15, and 21:9, as shown in Table 1.

The MNP@SiO₂ and HAp were previously sieved through an ABNT #325 mesh (44 µm) and dried in an oven at 50 °C for 2 hours. The hydrogel solution, already prepared in a beaker, was subjected to continuous stirring using a mechanical stirrer (FISATOM model 7130). The MNP@SiO₂ and HAp samples were then gradually added in the appropriate proportions, with stirring maintained until a homogeneous cement mixture (MNP@SiO₂:HAp:Hydrogel) was achieved. Subsequently, 2% dibasic sodium phosphate (Na₂HPO₄), relative to the total mass of reagents, was added to initiate the cement setting reaction. Dibasic sodium phosphate acts as a setting accelerator, promoting ionic exchange between calcium ions (Ca²⁺) from HAp and phosphate groups (PO₄^3-), leading to the formation of a cohesive calcium phosphate matrix.

After the formulation, the cements were cast into plastic and metal molds in disk (12 x 6 mm), cylindrical (6 x 12 mm), and rectangular (30 × 6 × 0.6 mm) shapes. The specimens were then incubated at 37 °C for 48 hours to complete the crosslinking and setting process.

2.3. Characterization

The crystalline phases present in the samples were determined by X-ray diffraction (XRD) using a Bruker D2 Phaser X-ray diffractometer operating with a Cu-Kα radiation source, a voltage of 30 kV, a current of 10 mA, and scanning from 10° to 80°. The identification of the crystalline phases was performed using the DiffracPlus Suite Eva Software and the Joint Committee on Powder Diffraction Standards (JCPDS). The morphological characteristics of the samples were examined by scanning electron microscopy (SEM), using a Tescan Vega3 instrument.

The swelling ratio (SR) was determined using simulated body fluid (SBF) with an ionic concentration closely matching that of human blood plasma, prepared at a temperature of 36 ± 1°C. The SBF was obtained by dissolving high-purity reagents in deionized water and buffered with 0.05 mol/L tris-hydroxymethyl aminomethane solution and 0.1 mol/L HCl to achieve a pH of 7.4 ± 0.05. Table 2 outlines the proportions to obtain the ionic concentrations of the SBF in mol/L and provides a comparison between the SBF and blood plasma.

For SR determination, test specimens were fabricated using acrylic molds in a disk shape with a diameter of 12 mm and a height of 6 mm. Samples prepared in triplicate were then withdrawn from the incubator and weighed upon achieving mass stabilization. Subsequently, they were immersed in containers with SBF at 36 ± 1 °C, following the surface area to solution volume ratio of (0.1 cm²/cm³)⁵³. Fluid absorption (SBF) was quantified by weighing the samples at various time points: 0.5, 1, 2, 4, 8, and 24 hours until reaching swelling equilibrium, that is, when there was no significant variation in the mass of the samples. To remove excess fluid, at each weighing performed, the samples were promptly placed on filter paper (< 20 s), gently pressed, and then reweighed on an analytical balance. After reweighing, they were dried again in an incubator at 37 °C until mass stability was reached. The SR was obtained according to Equation 1, where m_s represents the swollen mass, and m_d signifies the final dry mass.

The dynamic-mechanical analysis was performed on a multi-frequency DMA 2980 TA Instrument in the temperature range of 25°C to 200°C, at a heating rate of 5°C/min, under a frequency of 1Hz, and in a nitrogen atmosphere. The data were extracted only at temperatures of 25°C and 37°C, as this represents the temperature range corresponding to the handling conditions of the cements in the laboratory. The mechanical solicitation employed was tension/compression on rectangular specimens with approximate dimensions of 30 x 6 x 0.6 mm. For each sample, 10 analyses were performed. This analysis provided information regarding the storage modulus (E'), loss modulus (E'') in the form of heat during the material deformation, and the loss factor (tan δ = E''/E'), indicating the mechanical damping when the sample is subjected to periodic forces.

The compressive strength was measured using a Shimadzu Autograph AGX-100 KN universal compression testing machine, with a crosshead speed of 1 mm/min. According to the ASTM F451-08 standard⁵⁴, samples were made in the form of cylinders with a diameter of 6 ± 0.1 mm and a height of 12 ± 0.1 mm. A steel mold was used for this purpose. After demolding, the specimens were placed in a humid environment (100% relative humidity) for 24 hours and then immersed in SBF for different periods (1, 7, 14, 21, and 28 days) at a temperature of 36.0 ± 1.0 °C, with a liquid-to-solid volumetric ratio of 50/50. Ten samples were used for each test. The compressive strength is given by the ultimate compressive stress of the sample, which depends on the force acting on the effective cross-sectional area of the sample. It can be calculated by dividing the test result by the cross-sectional area of the sample in cm². The strength obtained is in Kgf/cm², and when divided by 10.1972, it gives the strength in MPa. The tensile strength and modulus of elasticity of the samples were also determined, based on Eq. 2 and Eq. 3, respectively:

where σ is the tensile strength (MPa), F is the force (N) required for fracture, A is the cross-sectional area of the specimens (cm²), E is the modulus of elasticity or Young's modulus (MPa), and ε is the strain.

The magnetic characterization was conducted using a Lake Shore model 7404 vibrating sample magnetometer (VSM), with a maximum applied magnetic field of 15,000 Oe at room temperature. Magnetic parameters, including saturation magnetization (Ms), remanent magnetization (Mr), and coercive field (Hc), were obtained from the hysteresis curve by analyzing the behavior of the curves near the origin of the Cartesian plane.

The cytotoxicity tests of the cements and HAp were performed in quintuplicate and followed ISO 10993-5⁵⁵ (Biological Evaluation of Medical Devices – part 5: Tests for in vitro cytotoxicity) and ISO 10993-12⁵⁶ (Biological Evaluation of Medical Devices - part 12: Sample preparation and reference materials). Femoral cells from 3-month-old male Fisher rats were used for the conduction of this test. Primary osteoblast cultures from femoral cells were placed in flasks and incubated at 37 °C with 5% CO₂ (Ultrasafe HF 212 UV Incubator). The culture medium was changed every three days, and the progression of the culture was assessed by inverted phase microscopy (Carl Zeiss Microscope - Axiovert 40C). The reference materials used were HDPE (negative control - non-cytotoxic) and latex (positive control - cytotoxic).

3. Results

Figure 1 shows the X-ray diffractograms of the magnetic cements obtained. Analyzing Figure 1a, which refers to the cements based on Fe₃O₄@SiO₂:HAp, the presence of peaks corresponding to the magnetite (Fe₃O₄), hematite (α-Fe₂O₃), and hydroxyapatite (Ca₅HO₁₃P₃) crystalline phases can be observed, matching well with the JCPDS standard cards n° 88-0315, n° 33-0664 and n° 90-01233, respectively. Following the same reasoning, in Figure 1b, which pertains to the cements based on CoFe₂O₄@SiO₂:HAp, it was possible to observe the presence of the crystalline phases of cobalt ferrite (CoFe₂O₄), hematite (α-Fe₂O₃), and hydroxyapatite (Ca₅HO₁₃P₃), with their peaks also matching well with the JCPDS standard cards n° 22-1086, n° 33-0664, and n° 90-01233, respectively. The presence of these peaks is also in accordance with studies reported in the literature⁵¹,⁵⁷,⁵⁸. As expected, as the quantity of ferrite in the cement formulation increased (Table 1), the presence of Fe₃O₄ peaks (Figure 1a) and CoFe₂O₄ peaks (Figure 1b) became more intense and pronounced to the detriment of hydroxyapatite phase. The amorphous characteristic, also evident in the cements' diffractograms, is likely attributed to the presence of the polymeric fraction responsible for hydrogel formation, as reported in the literature⁵⁷,⁵⁹.

Figure 1
XRD patterns for the cements: (a) MC1, MC2, MC3 and (b) CoC1, CoC2, CoC3. JCPDS standard cards are included.

Figure 2 illustrates the SEM micrographs obtained from the cements under study. In general, it can be observed that all micrographs show agglomerates of particles with different sizes, shapes and textures, typical of multiphase materials, as is the case with cements. However, it is also possible to observe that as the amount of ferrite in the cement formulations increased, as shown in Table 1, the presence of ferrite particles became more evident in the micrographs of samples MC3 (Figure 2c) and CoC3 (Figure 2f), referring to the cements 21%Fe₃O₄@SiO₂:9%HAp and 21%CoFe₂O₄@SiO₂:9%HAp, respectively. The presence of ferrite is distinguished by the appearance of plates with slightly hexagonal shapes, larger and smaller than 1 µm, a denser texture, and larger pores.

Figure 2
Scanning electron microscopy of the cements: (a) MC1, (b) MC2, (c) MC3, (d) CoC1, (e) CoC2, e (f) CoC3.

On the other hand, in the micrographs of the samples with the lowest percentage of ferrite, samples MC1 (Figure 2a) and CoC1 (Figure 2d), referring to the 9%Fe₃O₄@SiO₂:21%HAp and 9%CoFe₂O₄@SiO₂:21%HAp cements, respectively, it is possible to observe the presence of agglomerates of small particles, interconnected by apparently weak bonds (Van der Waals), and having a friable appearance. This behavior is due to the greater presence of hydroxyapatite in the cement formulation, contributing to a morphology with more micropores. Similar results were found by Alves et al.⁴⁰ and Barbosa⁶⁰, who identified similar behavior in the morphological analysis of cements.

The presence of pores in the cements is crucial for biomedical applications, as it enhances the adhesion between newly formed bone tissues. According to the literature, the porosity of HAp facilitates cell diffusion responsible for bone tissue deposition, improving bio-integration and the mechanical stability of the implant⁶¹,⁶². However, it is important to consider that the application of porous materials is limited to regions of the skeleton not subjected to mechanical stress or used as a bone cavity filler, due to their inherently low mechanical resistance.

Figure 3 shows the swelling ratio (SR) of the cements in the SBF solution at different time intervals at 37°C. The degree of swelling influences the bone cement’s ability to expand and contract similarly to surrounding bone tissue. An appropriate degree of swelling provides greater drug entrapment in biomaterials, promoting better interaction between biological entities and fluids, thus providing the necessary space for blood circulation⁶³. Furthermore, it minimizes the formation of gaps and voids between the biomaterial and bone, preventing complications such as infections and improving the stability of fixation and support of the structure as a whole⁶⁴,⁶⁵.

Figure 3
Swelling ratio of the cements: (a) MC1, MC2, MC3, and (b) CoC1, CoC2, CoC3.

According to the results in Figure 3, in general, it can be observed that the cements with a higher percentage of HAp, specifically MC1 and CoC1, demonstrated a greater SR than the other compositions. Conversely, the cements with a higher percentage of ferrite, specifically the MC3 and CoC3 samples, exhibited a lower SR compared to the other compositions. This result is possibly related to the porosity rate present in cements, which, in turn, depends on the number and size of the pores.

As evidenced in the micrographs in Figure 2, the compositions with the highest amount of HAp are also the cements that have the smallest particle sizes due to the microstructure of hydroxyapatite. Therefore, they are responsible for the largest surface area, contributing effectively to a greater volume of inter- and intraparticle pores, thus allowing greater infiltration of the SBF solution into the cement. On the other hand, ferrites, having larger and denser particle sizes, also exhibit lower intraparticle porosity, making the penetration of the SBF solution more difficult. However, when comparing the microstructures of silanized cobalt and magnetite ferrites (MNPs@SiO₂), which have particle sizes of 50.84 nm and 164.72 nm, respectively, it can be concluded that magnetite particles, being larger, also have larger interparticle pores than those of cobalt ferrite. Furthermore, regarding functionalization with the silane agent, it can be said that there is a greater probability of small pores in the microstructures of these MNPs being sealed by the presence of this agent on their surface. Therefore, when comparing cements with cobalt ferrite and magnetite MNPs, it is observed that cements with magnetite (Figure 3a) showed greater degrees of swelling compared to cements with cobalt ferrite (Figure 3b). However, all compositions showed increasing SR as a function of immersion time in the SBF solution. SR values ranging from 64% to 91% were observed among cements containing magnetite, specifically for cements MC3 and MC1, and from 58% to 87% among cements containing cobalt ferrite, specifically for cements CoC3 and CoC1, at immersion times of 0.5h and 24h, respectively.

The swelling behavior of the cements under study was evaluated over a period of 24 h. Research indicates that, after this period, the degree of swelling tends to reach an equilibrium state⁶⁶,⁶⁷. Sanmugam et al.⁶⁵, in turn, while evaluating the swelling behavior of nanohybrid scaffolds at 37°C in a phosphate-buffered saline (PBS) solution, observed that the swelling rate increases drastically in the initial hours and reaches its maximum value and equilibrium in less than 24 h, specifically shortly after the 15 h interval. According to Fernandes⁶⁸, swelling reaches its maximum value in one day, after which there is a gradual decrease across all samples.

The pH is crucial for the rheological properties of bone cement and can impact graft size and human health⁶⁹. It controls the concentrations of calcium (Ca) and phosphorus (P) in the cement solution, influencing the setting reaction rates. Factors such as the chemical composition of the cement, component proportions, particle size, use of accelerators or retarders, liquid-to-solid ratio, and temperature affect pH variation⁷⁰,⁷¹. Table 3 presents the pH values of the studied cements.

It can be observed that right after curing, the cements exhibited an alkaline nature, with pH values ranging from 8.0 to 9.0, suggesting a strong presence of basic compounds in the cement. It is also noted that the higher the amount of HAp in the cement composition (represented by MC1 and CoC1), the more basic its behavior, as there is a significant presence of hydroxyl ions (OH^-) in its structure. The decrease in the pH values of the cements after 1 day of immersion in SBF solution is possibly associated with the release of acidic groups or protons (H⁺), due to the dissolution or hydrolysis reactions of some components in the cement, such as carboxymethylcellulose (CMC). Although HAp is generally alkaline in aqueous medium, especially in the presence of aggressive ions from SBF, an initial demineralization of HAp may have occurred, releasing phosphate ions (PO₄^3-) and calcium (Ca²⁺) into the solution, resulting in the formation of acids, such as phosphoric acid (H₃PO₄), and thus causing a reduction in pH. The pH values after 1 day of immersion in SBF ranged from 4.5 to 4.9.

After 7 days of immersion in SBF, the pH of all cement formulations stabilized, presenting values in the range of 7.0–7.3, possibly due to the neutralization and precipitation of calcium phosphate phases³⁴,⁷²,⁷³, as well as the continuous removal of acidic ions through the frequent exchange of the SBF solution (every 2 days). In other words, the increase in pH likely occurred due to reactions of H₃PO₄ with calcium ions present in the SBF solution, once again producing HAp, since this is the least soluble calcium phosphate composition at a pH of 4.2 or higher, which stabilizes at a pH level close to 6.8⁷³. Therefore, since the cements present pH levels within the range of 6.5 to 8.0⁷⁰,⁷⁴, they are suitable for implantation as bone grafts, demonstrating that the additives used do not compromise the final pH of the system, as it stabilizes at values appropriate for clinical use.

Table 4 and Figure 4 present the initial and final setting times of the cements whose formulations have a fixed solid-to-liquid ratio of 30:70, with the composition of the liquid phase constant and variations only in the solid phase regarding the HAp:MNPs@SiO₂ ratio, as described in Table 1.

Figure 4
Setting times of the cements: (a) MC1, MC2, MC3, and (b) CoC1, CoC2, CoC3.

A slight reduction in the setting time of the cements was observed as part of the HAp was replaced by the addition of MNPs. It was also noted that the shortest setting times were observed in cements containing Fe₃O₄ nanoparticles, with values of 13.17, 12.53, and 12.03 minutes for cements MC1, MC2, and MC3, respectively. Meanwhile, for cements containing CoFe₂O₄ (CoC1, CoC2, and CoC3), the values were 25.81, 25.55, and 24.96 minutes, respectively.

Additionally, it was observed that the cements containing Fe₃O₄ also exhibited the shortest initial and final setting times, with values decreasing as the Fe₃O₄ percentage increased, ranging from 8.22 min to 6.24 min for the initial setting time, and from 21.39 min to 18.27 min for the final setting time. In the case of the cements with CoFe₂O₄, these values ranged from 11.15 min to 9.23 min for the initial setting time, and from 36.96 min to 34.19 min for the final setting time. This behavior is possibly associated with the surface area of these MNPs, which, in turn, directly depends on the size of their particles. Since the average particle size of CoFe₂O₄@SiO₂ (50.84 nm) was significantly smaller than that of Fe₃O₄@SiO₂ (164.72 nm), the cements containing CoFe₂O₄ tend to present a larger surface area, thus favoring greater swelling, lower viscosity, and consequently, a longer setting time. According to Ginebra et al.⁷⁵, the ideal setting time required for surgical procedures is between 10 and 15 minutes. Other studies⁷⁶-⁷⁹ report a slightly longer interval as the suitable setting time, with values within the range of 10 to 20 minutes. Therefore, the cement produced with Fe₃O₄ is the most suitable for surgical procedures. However, the other cements produced with CoFe₂O₄ extended the setting time by only ~5–6 seconds, which may facilitate application by providing additional working time.

Figure 5 presents the results of the storage modulus (E’), loss modulus (E’’), and loss factor (tan δ) obtained through dynamic mechanical analysis at temperatures of 25°C and 37°C, for the cements under study.

Figure 5
Dynamic-mechanical analysis of the cements: (a) MC1, MC2, MC3, and (b) CoC1, CoC2, CoC3.

According to the results in Figure 5, it is observed that the cements with a higher amount of HAp, specifically MC1 and CoC1, exhibited a higher storage modulus (E'), with values of 11.83 MPa and 10.92 MPa at 25°C, and 10.37 MPa and 9.63 MPa at 37°C, respectively, indicating greater elasticity of the material, i.e., a higher capacity to store energy during a load and unload cycle⁸⁰,⁸¹. It is noted that as the amount of MNPs@SiO₂ increased in the cement composition, the presence of organic chains from the molecular structure of the silane agent likely led to changes in the molecular interactions within the matrix, influencing properties such as cohesion and elasticity of the cement. Therefore, the cements with higher amounts of MNP@SiO₂ (MC2, MC3, CoC2, and CoC3) are responsible for the lower values of storage modulus (E') and loss modulus (E''), with values ranging from 4.47-1.72 MPa and 0.97-0.41 MPa, respectively. The loss modulus (E''), representing the amount of energy dissipated as heat during material deformation⁸², confirmed the less rigid nature of MC2, MC3, CoC2, and CoC3 cements, requiring less energy for chain mobility.

It was also observed that increasing the temperature from 25°C to 37°C led to reductions of 12.34% and 11.81% in the E' modulus values, and 51.24% and 17.18% in the E'' modulus, for the MC1 and CoC1 cements, respectively. This behavior can be associated with the breaking of secondary bonds and the disruption of the polymer chains in the liquid phase of the cements, as well as the structure of the silane agent, as previously described.

Since the loss factor (tan δ) represents the mechanical damping of the cements when subjected to periodic forces, it is observed that the cements analyzed at 25°C and 37°C showed values within the range of 0.21-0.55. However, it was noted that the cements analyzed at 37°C exhibited a lower loss factor, indicating a slightly reduced viscous behavior compared to those analyzed at 25°C, which showed more viscous behavior with greater conversion to mechanical energy. This is likely due to the fact that at lower temperatures, the cement is more glassy and thus more rigid, requiring more time for relaxation and mobility of the polymer chains used in the cement composition⁸³,⁸⁴.

Figure 6 presents the mechanical resistance to compression of the cements soaked in the SBF solution for different times at 37°C. Generally, similar behaviors were observed for the cements containing both Fe₃O₄@SiO₂ and CoFe₂O₄@SiO₂. It was detected that the cements exhibited higher compressive strength after 14 days of immersion in SBF, reaching maximum values of 28.91 MPa and 30.06 MPa for the MC1 and CoC1 cements, respectively. After 14 days of immersion in SBF, a gradual reduction in the strength of the cements was observed; however, they maintained a compressive strength in the range of 22-25 MPa.

Figure 6
Mechanical resistance to compression of the cements: (a) MC1, MC2, MC3, and (b) CoC1, CoC2, CoC3.

It was also observed that transitioning from a 1-day to a 7-day immersion in SBF led to significant increases in compressive strength, ranging from 24% to 38%, reaching maximum values of 26.84 MPa and 28.11 MPa for the MC1 and CoC1 cements, respectively. This likely occurred because the cements had not yet completed their curing process within the first 24 hours. A similar behavior was also noted by other authors who evaluated the compressive strength of bone cements⁸⁵,⁸⁶.

Thus, as in the present study, Rabiee, Moztarzadeh, and Solati-Hashjin⁸⁷ also observed an increase in the compressive strength of hydroxyapatite cements with 0-4% of the accelerator NaH₂PO₄:2H₂O in SBF solution during the first 14 days, followed by a slow reduction after this period. According to the authors, the increase in the strength of the cements during the first two weeks can be attributed to new crystallization of the calcium phosphate phase and the growth of crystals within the cement particles, while the gradual reduction in strength during the following two weeks may be due to the slow hydrolysis of DCPA and TCP and the gradual transition to HAp⁸⁸.

Figure 7 illustrates the mechanical tensile strength of the cement soaked in the SBF solution for different times at 37°C. The highest tensile strength values were obtained by the cements MC3 and CoC2 after a 7-day immersion period in SBF, reaching maximum values of 14.52 MPa and 15.31 MPa, respectively. After this period, a slight reduction in the tensile strength of the cements was observed, but it remained in the range of 10.38 to 13.53 MPa. This slight decrease may be related to the saturation of the cement pores with the SBF solution and the leaching of organic components, such as glycerin and CMC, as these materials may undergo degradation or displacement after long periods in aqueous solution, reducing the internal cohesion of the cement matrix.

Figure 7
Mechanical tensile strength of the cements: (a) MC1, MC2 and MC3, and (b) CoC1, CoC2 and CoC3.

Figure 8 presents the results of the elastic modulus (E) of the cements. A similar behavior was observed, as there was an increase in the E of the cements from 1 day to 7 days of immersion in SBF, reaching maximum values of 1893.04 MPa and 1864.42 MPa with the cements MC3 and CoC3, respectively. After the 7-day immersion period in SBF, the cements containing magnetite showed a slight reduction in the elastic modulus, but maintained a level in the range of 1534 to 1754 MPa between the 14 and 28 days of immersion in SBF. The cements with cobalt ferrite, on the other hand, exhibited a gradual reduction in E as the immersion time in SBF increased, reaching the lowest values in the range of 1352 to 1405 MPa after 28 days of immersion.

Figure 8
Elastic modulus of the cements: (a) MC1, MC2 and MC3, and (b) CoC1, CoC2 and CoC3.

In general, the increase in the mechanical resistance of the cements can be explained by three main mechanisms: (i) the initial hardening, which likely occurred through a chelation reaction between the carboxymethylcellulose present in the cement formulations and the calcium from hydroxyapatite; (ii) the influence of the polymeric compositions, which enhance the internal cohesion of the matrix; and (iii) the progressive transformation of the cement components into hydroxyapatite, resulting in a more robust structure. The mechanical reduction of the cements after 14 days of immersion in SBF can be attributed to a combination of factors related to the saturation of the cement pores weakening internal cohesion, the leaching or degradation of organic components (glycerin and CMC) present in the matrix, or even the possible hydrolysis of residual phases of calcium phosphates.

Figure 9 shows the behavior of the magnetization (M) as a function of the applied magnetic field (H) through the hysteresis loop for the cements. In general, a typical behavior of ferrimagnetic materials can be observed for all samples, with the formation of a well-defined S-shaped hysteresis loop. However, the distinction between the two types of cement lies in the nature of the magnetic material used. Magnetite (Fe₃O₄), considered a soft magnetic material, is easily magnetized and demagnetized, while CoFe₂O₄ ferrite, a hard magnetic material, is characterized by residual magnetism with a wider hysteresis loop³⁵.

Figure 9
Magnetic hysteresis curves referring to de cements: (a) MC1, (b) MC2, (c) MC3, (d) CoC1, (e) CoC2 e (f) CoC3.

The magnetic parameters of the cements, such as saturation magnetization (Ms), remanent magnetization (Mr), coercivity (Hc), and remanence ratio (Mr/Ms) determined from the M x H curves, are described in Table 5. It was observed that as the amount of MNPs in the cement composition (MNPs@SiO₂:HAp:Hydrogel) increased, its magnetic response also increased, to the detriment of the non-magnetic influence inherent to both HAp and the hydrogel. In this sense, the highest Ms values were found in the CoC3 and MC3 cements, with values of 31.08 emu/g and 27.30 emu/g, respectively. An increase in Mr was also observed as the MNPs content in the cement increased, with values ranging from 2.29 emu/g to 4.29 emu/g in the cements with magnetite, and from 2.87 emu/g to 11.55 emu/g in the cements with cobalt ferrite. The coercivity (Hc) values were lower in the magnetite cements, ranging from 294 Oe to 347 Oe, confirming the typical behavior of a soft magnetic material, while the cobalt ferrite cements exhibited higher coercivity values, ranging from 1253 Oe to 1289 Oe, typical of a harder magnetic material. However, the low remanence ratio values (Mr/Ms), also known as the squareness ratio, in the range of 0.16 to 0.37, are associated with larger particles and domain-wall formation⁸⁹.

In Figure 10, the results of the cell viability of HAp and the studied cements, when exposed to femoral cells from 3-month-old rats, are shown. The cytotoxicity of a sample is determined by the percentage of cells that remain viable after the cell population is exposed to various concentrations of the test substance extract (analyzed sample). The results, therefore, confirm the non-toxicity of HAp and all the cement formulations studied, whether containing magnetite or cobalt ferrite MNPs, consistent with data previously reported in the literature³⁵,⁷³.

Figure 10
Cell viability of HAp and the studied MNPs@SiO₂:HAp:Hydrogel, as well as the corresponding positive (PC) and negative (NC) controls.

HAp demonstrated cell viability of 83%, while the cements exhibited values ranging from 107% to 128%, all significantly higher than the target threshold of 70%. It was also observed that as the MNP content in the cement increased, there was a slight reduction in cell viability. However, this reduction did not compromise the overall viability, which remained well above 107%.

4. Conclusion

The production of the Fe₃O₄ and CoFe₂O₄ bone cement formulations (MNPs@SiO₂:HAp:Hydrogel) was successfully achieved, with all cements displaying properties suitable for clinical application. The cements exhibited favorable physicochemical properties, including setting times ranging from 12.03 to 25.81 minutes, an alkaline pH (7.0–7.3), and enhanced swelling behavior, with reaction kinetics in SBF solution increasing up to 91% over 24 hours. Cements with higher HAp content showed greater swelling, while those with a higher percentage of MNPs (MC3 and CoC3) demonstrated superior mechanical performance, with compressive strengths of 28.91 MPa and 30.06 MPa, respectively, after 14 days in SBF. Additionally, the cements exhibited typical ferrimagnetic behavior, allowing for magnetic guidance, non-invasive tracking, and potential use in additional therapies. The formulations were non-toxic and biocompatible, promoting tissue integration and healing. Among these cements, CoC3 stood out due for its ease of pilot-scale production and lower cost, making it a promising candidate for vertebroplasty surgeries. Therefore, the MC3 and CoC3 formulations are recommended for further exploration in clinical applications, offering an effective balance of mechanical strength, swelling behavior, and magnetic response.

5. Acknowledgments

National Council for Scientific and Technological Development – CNPq, Coordination for the Improvement of Higher Education Personnel – CAPES and Foundation for Research Support of the State of Paraíba – FAPESQ for the financial support of scholarships.

6. References

Publication Dates

History