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Bioceramics in the CaMgSi2O6–Li2O System: A Glass‐Ceramic Strategy for Excellent Mechanical Strength and Enhanced Bioactivity by Spontaneous Elemental Redistribution

Bioceramics in the CaMgSi2O6–Li2O System: A Glass‐Ceramic Strategy for Excellent Mechanical... IntroductionImplantable bioceramics bone tissue‐engineering materials that combine multiple modes of bioactivity to help facilitate natural bone regeneration are an emerging approach toward the treatment of significant bone defects. In recent years, a variety of biomaterials such as bioactive glasses,[1–3] glass‐ceramics,[4–6] crystalline calcium silicates (such as wollastonite [CaSiO3] and larnite [Ca2SiO4],[7,8] and calcium phosphate‐based bioceramics[9,10] have been the subject of numerous investigation toward bone tissue engineering. In particular, some researchers have proposed that CaSiO3 exhibits superior bioactivity and osseointegration properties to the widely studied hydroxyapatite (HAp), which motivates further investigation into CaSiO3‐based materials for bone tissue repair and regeneration.[11,12] The presence of bioavailable calcium (Ca) and silicon (Si) ions in tissue engineering materials can directly influence the quality of bone as Ca is the main constituent of biological apatite and Si was reported to have the ability of inducing osteogenesis and angiogenesis.[13,14] However, it has been proposed that the poor mechanical strength and rapid dissolution rate of CaSiO3 may hinder its biomedical application to be used as bulk implants and would be detrimental for cell growth because of a rapid increase in pH value in the surrounding environment during the first few weeks.[15,16]Magnesium‐containing bioceramics have gained interest recently due to the favorable combination of properties that is achievable in such systems in terms of biodegradation, bioactivity and mechanical strength.[17,18] As with Ca and Si, magnesium (Mg) is also an indispensable element for the human body, and bone regeneration in particular as Mg deficiency would stimulate the differentiation of osteoclasts and inhibit the production of osteoblasts, thus leading to osteoporosis.[19] Further, it has been proved that Mg‐containing calcium silicate‐based bioceramics can stimulate the apatite formation and induce a tight bond with bone tissues.[20,21] Lai et al.[22] suggested that an increase of new vessel development and blood perfusion after surgery was achieved by Mg‐incorporated PLGA/TCP porous scaffolds. Further, Mg‐containing thin films deposited on porous polymer scaffolds could significantly develop new bone formation and strengthen mechanical properties of newly formed bone,[23] indicating that Mg‐containing bioceramics could be used as a novel material for bone tissue engineering applications.Due to the various attractive attributes of calcium magnesium silicates as effective bioceramics, such materials from various silicate families have been selected as bone‐tissue engineering materials, including the sorosilicate akermanite (Ca2MgSi2O7),[24,25] orthosilicate monticellite (CaMgSiO4),[26,27] and the pyroxene diopside (CaMgSi2O6).[28–30] Among all the materials, diopside has superior mechanical properties,[31] and also possesses good cytocompatibility and surface apatite precipitation ability.[32,33] Consequently, diopside (CaMgSi2O6) was chosen as the primary bioceramic material in the present study. It is worth noting that the mechanical properties of CaMgSi2O6 was improved in comparison with that of CaSiO3, while the HAp mineralization in simulated body fluid (SBF) was compromised due to the incorporation of Mg.[34–37]Toward the repair of defects, the mechanical and biological properties of bioceramic implants should be compatible with native bone tissue and stimulate a favorable biological response.[38] One of the current techniques for addressing these issues is to introduce additional elements into main structures (doping[39–42] and/or composites[43,44]) to modify the crystallinity, lattice parameters, and morphology of the bioceramics, further affecting their mechanical and biological properties. Table 1 summarizes the previous works on the incorporation of additional elements in CaMgSi2O6 bioceramics.1TableSurvey of experimental works on introducing additional elements into CaMgSi2O6 bioceramicsElementsElement contentAnnealing treatmentEffectsReferenceSr/F2 mol%/1 mol%1200 °C/5 h All the doped scaffolds presented higher apatite‐forming ability Osteoblast‐like MG‐63 cells exhibited the highest compatibility to the Sr‐doped scaffolds Osteoblast‐like MG‐63 cells exhibited the lowest compatibility to the codoped scaffolds[45]Sr2 mol%1200 °C/3 h Sr‐doping improved the sinterability and apatite‐forming ability of the diopside‐based scaffolds Sr‐doping retarded the biodegradation of diopside‐based scaffolds Mesenchymal stem cells presented better adhesion and spreading with typical cell extensions on the Sr‐doped diopside‐based scaffolds[46]LiNaK2 mol%900 °C/2 h The crystallinity and lattice volume of diopside were changed with the most and least deviations for K and Na, respectively The dopants altered the in vitro bioactivity of diopside in the following ranking: K‐doped >  Li‐doped >  Na‐doped >  undoped The 2 mol% dopants improved the biocompatibility of diopside, where the most beneficial effect was found for Na and K K‐doping was the optimal doping for the bioactivity and cytocompatibility assessments[47]Ce0–100 mol%1000 °C/4 h Addition of 25 mol% Ce had the best biomineralization performance in vitro Less hydroxyapatite precipitates were found with further increasing Ce addition[48]Mo0–100 mol%1000 °C/4 h At a lower Mo content, the mixed phase materials showed higher hardness and slower biodegradation At a higher Mo content, mixed phase materials exhibited lower hardness and bioactivity[49]Lithium (Li) has been commonly used to treat bipolar disorder,[50,51] and the psychiatric patients who took Li medication have been reported to have enhanced bone mass and lower bone turnover states.[51] This may be explained by the mechanism that Li can stimulate the osteogenesis through activating Wnt and Hh signaling pathway.[52,53] Moreover, it was found that Li doping can improve the thermal stability of calcium phosphate.[54] Table 2 summarizes the previous works on the incorporation of Li element in different types of bioceramics. Padmanabhan et al.[55] prepared lithium‐substituted hydroxyapatite (Li‐HAp) nanoparticles by a sonochemical synthesis process. It was reported that the incorporation of Li into HAp caused the transformation of particles from a rod shape to a needle‐like structure due to the sonication effect. Further, Li‐HAp showed excellent antimicrobial activity and blood compatibility, and could be used as a drug delivery carrier owing to its improved cytocompatibility. Yuan et al.[56] synthesized Li‐doped calcium polyphosphate bioceramic scaffolds by a gravity sintering method. The study concluded that Li doping did not affect the crystallized phase but changed the surface morphology. Further, the Li‐doped samples were found to have better degradation properties and enhanced proliferation and differentiation of osteoblasts. Wang et al.[57] synthesized Li‐doped HAp by a precipitation method, where the scaffolds had better mechanical strength and Li doping was found to be beneficial to the proliferation of osteoblasts. Li‐doped β‐TCP was found to have two times higher compressive strength than initial β‐TCP.[58] The incorporation of nanoporous lithium doping magnesium silicate into calcium sulfate hemihydrate was found to enhance the degradability, biocompatibility, vascularization, and osteogenesis.[59] As summarized above, the incorporation of Li into calcium phosphate‐based bioceramics was expected as a viable way to enhance the mechanical and biological properties of bioceramics. However, few studies have been conducted on bioceramics in the diopside–lithium oxide system. Only one study synthesized Li‐doped diopside through a coprecipitation method, as shown in Table 1. The result showed that the substitution of 2 mol% Li for 1 mol% Mg has improved the in vitro bioactivity and biocompatibility.2TableSurvey of experimental works on introducing Li element into bioceramicsBioceramic typeLi amountMethodEffectsReferenceHAp0.1–0.5 mol%Sonochemical Li‐HAp particles were transformed from a rod‐like structure to needle‐like structure Reduction in crystallinity of HAp Excellent antimicrobial activity and blood compatibility Promoted cell proliferation[55]Calcium Polyphosphate (CPP)2 mol%Gravity sintering The crystallized phase was not affected Rougher and more uneven surface Better degradation properties Promoted proliferation and differentiation of osteoblasts Better osteoconductivity[56]HAp0.5–1 mol%Precipitation High mechanical strength Fast osteoblast‐mediated degradation rate Promoted osteoblast activity[57]β‐TCP1.9 mol%Precipitation and ball milling Compressive strength enhanced by two times Lowered degradation rate and cytotoxicity[58]Calcium Sulfate Hemihydrate–Precipitation and mixing Enhanced degradability in vitro and in vivo Prevention of acidification in the biological environment Improved new bone formation Better vascularization and osteogenic differentiation[59]HAp0.5–2 mol%Precipitation 1.5%Li‐HAp obtained optimal cell proliferation activity Enhanced new bone formation Better osteogenesis and angiogenesis potential[60]As the amounts of Li incorporated or substituted into diopside and other bioceramics were mostly limited to a maximum of 2 mol%, the present work attempts to investigate the effects of a wide range of Li content (25, 50, 75, and 100 mol%) on the diopside–lithium oxide (CaMgSi2O6–Li2O) system to determine the appropriate Li amount and thereafter optimize its mechanical and biological properties. The specimens were synthesized through a precipitation method, followed by sintering at 1000 °C for 4 h. The mineralogical, morphological, mechanical, and biological properties were investigated to comprehensively understand the factors and mechanisms affecting the properties of CaMgSi2O6–Li2O bioceramics system.Results and DiscussionCrystalline Phase IdentificationThe phase identification of the pure diopside and CaMgSi2O6–Li2O bioceramics was obtained by X‐ray diffraction (XRD), as demonstrated in Figure 1. In order to precisely investigate the crystallized phases in the bioceramics, the as‐sintered samples were ground into powder for random crystalline orientation and then measured by XRD, as shown in Figure 1a. It can be seen that a major phase of CaMgSi2O6, corresponding to ICSD code: 168107, was observed in the pure diopside. As the Li2O substitution increased to 25 mol%, the formation of secondary phases (CaSiO3, minor Li2Si2O5, and SiO2) was detected in accordance with ICSD code: 201537, 15414, and 81382, respectively. The CaSiO3 was identified as β‐CaSiO3 with triclinic structure which is generally formed below 1100 °C.[61] As the Li2O content increased to 50 and 75 mol%, the peak intensity of CaMgSi2O6 phase decreased along with the increment of peak intensity in CaSiO3, Li2Si2O5, and SiO2 phases, and two new secondary phases Li2SiO3 and Li2Ca2Si5O13 formed, matching with ICSD code: 100402 and 79379, respectively. With further increasing the Li2O content to 100 mol%, CaMgSi2O6 phase was completely diminished and replaced by CaSiO3 (major), Li2Si2O5, SiO2, Li2SiO3 (minor), and Li2Ca2Si5O13 (minor) phases as a result of the absence of Mg element. The details of the phase transition, peak intensities, and degree of crystallinity are clearly listed in Table 3. Following the substitution of Li2O for MgO, the amorphous phase started to emerge, and the amorphous content increased with the higher Li2O substitution. The XRD results can be further supported by the phase diagram of Li2O–SiO2 system, where Li2SiO3 and Li2Si2O5 would form at 936 to 1028 °C when high molar concentrations of SiO2 were added.[62]1FigureXRD patterns of pure diopside and CaMgSi2O6–Li2O bioceramics in the form of a) powder and b) pellet, and c) after SBF soaking for 28 days in the form of pellet.3TablePhase transition, peak intensities, and degree of crystallinity in pure diopside and CaMgSi2O6–Li2O bioceramicsSamplePhase transitionDegree of crystallinity [%]CrystallinityAmorphousPureCaMgSi2O6 only––Li0.25CaMgSi2O6 (↓) + CaSiO3 (↑) + minor Li2Si2O5 (↑)+ minor SiO2 (↑)95.544.46Li0.50CaMgSi2O6 (↓↓) + CaSiO3 (↑) + Li2Si2O5 (↑)+ SiO2 (↑) + minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)91.638.37Li0.75CaMgSi2O6 (↓↓) + major CaSiO3 (↑↑) + Li2Si2O5 (↑)+ SiO2 (↑) + minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)87.7412.26Li1.00major CaSiO3 (↑↑↑) + Li2Si2O5 (↑) + SiO2 (↑)+ minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)84.5615.44In order to study the crystallization behavior of the bioceramics in biological environment, the XRD patterns of the bioceramics pellets before and after SBF soaking for 28 days are compared in Figures 1b and 1c, respectively. The slight difference in peak intensity and peak position between the form of pellet and powder could be originated from the preferred orientation and the uneven sample surface. It can be seen that the basic compositions in the SBF‐soaked materials were practically the same as those in the as‐sintered samples, confirming that no significant phase transformation was presented in the matrices. Further, the peaks at 25.9° and 31.7–34.1°, representing the precipitate of HAp (Ca5(PO4)3OH, ICSD code: 34457), were observed in all the samples, where CaMgSi2O6–Li2O bioceramics had higher intensities of Ca5(PO4)3OH phase than pure diopside, indicating a superior HAp formation capacity of CaMgSi2O6–Li2O bioceramics. In addition, the peak intensity of CaSiO3 was lowered after SBF soaking, while the peak intensity of CaMgSi2O6 remained evident. The enhancement of HAp formation and the degraded CaSiO3 phase could be ascribed to the replacement of MgO by Li2O, as reported by Vallet‐Regí et al.[37] that the decrease of MgO in CaO–MgO–SiO2–P2O5 bioceramics promoted the apatite formation rate and the vulnerability to aqueous solution because of better seeded HAp crystallization and lower CaO bonding energy than MgO.Morphology and Elemental DistributionThe scanning electron microscope (SEM) images (5000×) for sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets are shown in Figure 2. The surface morphology of pure diopside without Li2O addition was homogenous with a distribution of small pores, as shown in Figure 2a. With 25 mol% of Li2O substitution for MgO, a microstructure consisting of small grains embedded in a glass‐like substrate was clearly observed in Figure 2b. The presence of amorphous phase was also detected by XRD and degree of crystallinity analysis in Table 3. Moreover, with the increase of Li2O addition, the aggregation and grain growth phenomena became more apparent, leading to an increasingly porous structure, as shown in Figure 2c–e. The size of aggregated grains was getting larger from 3 to 12 µm in length and 2 to 5 µm in width. Li et al.[35] have reported that the incorporation of MgO into CaSiO3 gave rise to the formation of calcium–magnesium silicates, and therefore resulted in a compact surface and a decrease of total pore volume, and vice versa. However, in the present study, a fused and densified surface was observed in Li0.25 with an increased ratio of CaSiO3 due to the substitution of Li2O for MgO, which contradicted with what Li et al. has found. In this case, since the appearance of CaSiO3 should have increased the porosity, the densified surface of Li0.25 may be contributed by another secondary phase: Li2Si2O5. This has been supported by Wang et al. that the aggregated ZrO2 grains were embedded into the liquid matrix of Li2Si2O5 by a sol–gel synthesis at the sintering temperature of 1000 °C, leading to the removal of pores and the densification of ceramics.[63] This is so‐called “glass‐ceramics” composed of a Li2Si2O5 crystalline phase and an amorphous matrix, which is induced by a heterogeneous nucleation mechanism.[64–66] In the current study, the viscous amorphous matrix of Li2Si2O5 was able to fill the pores throughout the bioceramics and hence densified the materials. With the further increase of Li2O content, the porosification induced by the increasing CaSiO3 formation and grain growth would surpass the densification caused by Li2Si2O5, resulting in a porous material surface. A similar outcome was obtained by Li et al. that the incorporation of Li2Si2O5 into Ca0.7Sr0.3TiO3 enhanced the densification of materials, and the grain size increased with the elevating content of Li2Si2O5.[67]2FigureSEM images of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets: a) pure, b) Li0.25, c) Li0.50, d) Li0.75, and e) Li1.00.Elemental mapping based on Ca, Mg, and Si elements on the surface of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets is shown in Figure 3. Li element was excluded from the mapping analysis as the atomic order of Li is too small to be detected by field emission electron probe microanalyzer (FE‐EPMA). A homogeneous distribution of Ca, Mg, and Si can be observed in pure diopside, indicating the existence of CaMgSi2O6 single phase. With 25 mol% of Li2O substitution, the elemental distribution of Ca and Mg was concentrated in specific areas, suggesting the formation of CaSiO3 and the grain growth of CaMgSi2O6. As the substitution of Li2O increased to 50 and 75 mol%, the grain size of CaSiO3 gradually increased, while the CaMgSi2O6 phase decreased with the reduction of Mg, as confirmed by XRD (Figure 1) and SEM (Figure 2). When the Li2O content reached 100 mol%, the Mg signal has completely disappeared and the large grains of CaSiO3 partially embedded in the glass phase of Li2Si2O5 were clearly detected. As can be seen from CaMgSi2O6–Li2O bioceramics, it is noteworthy that Mg elements were mostly concentrated in the crystalline phase of CaMgSi2O6, while Ca elements were not only found in the crystalline phases of CaSiO3 and CaMgSi2O6, but heavily detected in the amorphous glass phase. This phenomenon may be explained by a mechanism that Mg2+ has higher field strength than Ca2+ and Li+, which increases the viscosity of the glass phase, leading to a reduced diffusion of ions and ionic complexes.[68] Because of the poor ion diffusion in the regions containing concentrated Mg2+ ions, Mg2+ may have a strong tendency to nucleate CaMgSi2O6 crystals with the localized Ca2+, Si4+, and O2− ions. By contrast, partial Ca2+ can be dispersed in the glass phase of Li2Si2O5 due to its high ionic diffusivity without the influence of Mg2+.3FigureElemental mapping analysis on the surface of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets.The SEM images (5000×) for pure diopside and CaMgSi2O6–Li2O bioceramics pellets after soaking in SBF solution for 28 days are shown in Figure 4. The presence of newly formed HAp layers can be clearly observed on all the samples. For pure diopside, it can be seen that the diameter of the spherical precipitates on the sample surface varied from 1 to 2 µm. Regardless of the initial matrix structure, the diameter of the spherical precipitates on all the CaMgSi2O6–Li2O bioceramics samples ranged from 5 to 12 µm, which was much larger than that on the pure diopside, indicating a superior HAp formation capacity of CaMgSi2O6–Li2O bioceramics, as confirmed by the XRD results in Figure 1c. The improved HAp formation could be attributed to the appearance of CaSiO3 and Li2Si2O5 caused by the Li2O addition, where CaSiO3 has been reported to have better apatite mineralization but more rapid dissolution rate than CaMgSi2O6.[34,35] Furthermore, the quantitative compositions of HAp formed on pure diopside and CaMgSi2O6–Li2O bioceramics were measured by the point mode of FE‐EPMA, as indicated by yellow arrows in Figure 4, and the results were tabulated in Table 4. This quantitative analysis mainly measured the concentrations of Ca, Mg, Si, O, and P as they are the major compositions of CaMgSi2O6‐based bioceramics and HAp precipitates. Other elements, such as undetectable Li, gold coating used for enhancing conductivity, or the elements adsorbed from SBF solution, were classified as “others.” It is clear that the HAp formed on CaMgSi2O6–Li2O bioceramics contained Ca, O, and P as its major components, corresponding to the composition of Ca5(PO4)3OH. In addition to Ca, O, and P, the HAp formed on the pure diopside had a noticeable amount of Mg and Si, indicating that it readsorbed the Mg and Si elements which were released from the matrix materials into SBF solution.4FigureSEM images of pure diopside and CaMgSi2O6–Li2O bioceramics pellets after soaking in SBF for 28 days: a) pure, b) Li0.25, c) Li0.50, d) Li0.75, and e) Li1.00.4TableComposition of HAp formed on pure diopside and CaMgSi2O6–Li2O bioceramics pellets (measured by the point mode of FE‐EPMA)SampleComposition of HAp [wt%]CaMgSiOPOtherspure17.14.110.523.46.138.8Li0.2527.20.50.011.114.746.5Li0.5028.40.30.311.811.347.9Li0.7529.30.30.611.413.245.2Li1.0027.10.30.816.112.343.4In particular, the elemental mapping analysis of Li0.25 bioceramics after soaking in SBF for 28 days was performed and shown in Figure 5. It can be seen that the soaked Li0.25 bioceramics was comprised of Ca, Mg, Si, O, and P ions. The elemental mapping implied the development of a uniform microstructure with regard to an appropriate spatial distribution of elements. High concentrations of Ca and P ions were found on the surface of HAp precipitates, while high concentrations of Mg and Si ions, and a medium concentration of Ca ions were found on the surface of the Li0.25 matrix, which were consistent with their chemical compositions.5FigureElemental mapping analysis of the Li0.25 bioceramics pellets after soaking in SBF for 28 days: a) SEI, b) Ca, c) Mg, d) Si, e) O, and f) P elements.Density and Porosity MeasurementTable 5 shows the bulk and apparent density of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets. Following the density results, the porosity given by open pores was calculated accordingly. As it can be seen, the apparent density (DA) of pure diopside is 3.19 g cm−3, which is very close to the theoretical density of diopside (CaMgSi2O6) at 3.26 g cm−3.[69] However, its bulk density (DB) is only 1.73 g cm−3, representing a high porosity of 45.8% due to the considerable difference of DA and DB. As 25 mol% of Li2O was incorporated into CaMgSi2O6, the bulk and apparent density of Li0.25 were slightly decreased to 2.94 and 2.95 g cm−3, respectively. The slight decrease in the density could be ascribed to the formation of CaSiO3, Li2Si2O5, and SiO2 which have theoretical densities of 2.91,[70] 2.5,[71] and 2.26[72] g cm−3, respectively, all of which are lower than the theoretical density of CaMgSi2O6 at 3.26 g cm−3. Further, Li0.25 had an extremely low porosity of 0.3%, indicating a dense structure which is in good agreement with the observation in SEM (Figure 2b). As the Li2O content increased to 50 mol%, the porosity of 0.7% of Li0.50 suggested a dense structure similar to Li0.25, while the bulk and apparent density dropped further to 2.67 and 2.69 g cm−3, respectively. This could be attributed to the decreased content of CaMgSi2O6 and the increased content of CaSiO3, Li2Si2O5, and SiO2, coupled with two newly formed secondary phases Li2SiO3 and Li2Ca2Si5O13 with theoretical densities of 2.53[73] and 2.75 g cm−3,[74] respectively. When 75 mol% Li2O was added, the apparent density of Li0.75 remained around 2.7 g cm−3, but the bulk density decreased to 2.14 g cm−3 with a porosity of 21.3%, indicating an increase in the connected open pores. When the Li2O substitution reached 100 mol%, the apparent density of Li1.00 remained at 2.7 g cm−3, but the bulk density significantly dropped to 1.59 g cm−3, leading to an increase in porosity to 41.1%, which is comparable to the high porosity of pure diopside. These phenomena revealed that the substitution of Li2O induced the formation of secondary phases with lower density (CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13), resulting in the reduction of the overall density. However, the addition of small amounts of Li2O decreased the porosity of CaMgSi2O6, leading to a fused and dense structure due to the presence of Li2Si2O5. As the amount of Li2O increased, the grain size and porosity increased owing to the massive formation of CaSiO3, following the same trend of structural changes observed in SEM (Figure 2b–e).5TableBulk density, apparent density, and porosity of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pelletsSampleBulk density [DB, g cm−3]Apparent density [DA, g cm−3]Porositya) [%]pure1.733.1945.8Li0.252.942.950.3Li0.502.672.690.7Li0.752.142.7221.3Li1.001.592.7041.1a)The porosity of tabulated samples is contributed by the open pores.Hardness MeasurementHardness results of sintered pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics before and after immersing in SBF solution are shown in Figure 6. Each sample was indented at three different locations, and then the averages and standard deviations were calculated. It was expected that the addition of Li2O would affect the hardness property through different combinations of phases (CaMgSi2O6, CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13) and structural change in the material.6FigureHardness results of sintered pure diopside and CaMgSi2O6–Li2O bioceramic: a) before and b) after SBF soaking for 28 days.As can be seen from Figure 6a, Li0.25 had the highest hardness of all the samples at about 700 Hv, which is about 38 times harder than the pure diopside. This result shows that the hardness of CaMgSi2O6–Li2O bioceramics decreased as the amount of Li2O increased, but the hardness was still generally much higher than that of the pure diopside. This indicates that the Li2O addition can significantly increase the mechanical strength of CaMgSi2O6. Based on Table 5 and Figure 6a, it was found that the hardness was directly proportional to the apparent density and inversely proportional to the porosity. The pure diopside had the highest apparent density, but the high porosity of 45.8% yielded a fragile structure and a low hardness. With the incorporation of Li2O, Li0.25 had a similar porosity but a higher apparent density compared to Li0.50, leading to a superior hardness of Li0.25 to that of Li0.50. As the increment of Li2O from 50 to 75 and 100 mol%, the apparent densities of Li0.50, Li0.75, and L1.00 remained at ≈2.7 g cm−3, while the porosity was increased from 0.7% to 21.3% and 41.1% with a remarkable drop in hardness from 521 down to 416 and 318 Hv. Based on the XRD results (Figure 1), the secondary phases (CaSiO3 and Li2Si2O5) occurred in Li0.25, suggesting that the highest hardness measured in Li0.25 was associated to the synergistic effect of CaMgSi2O6, CaSiO3, and Li2Si2O5. In addition, an amorphous phase of 4.46% was able to increase the bulk density of Li0.25 and reduce its porosity substantially. As elucidated in SEM and density measurement, the Li2O substitution varied the porosity and the density of Li0.25, inducing an aggregated and densified crystal structure as the reason of its outstanding hardness. With further increasing the Li2O quantity, though the amorphous phase increased to 8.37–15.44%, the Mg deficiency and the abundant CaSiO3 and Li2Si2O5 phase triggered progressive aggregation and grain growth, followed by a lowered density and a porous morphology, as shown in Figure 2c–e, resulting in the degraded hardness. Thus, it is expected that the hardness of CaMgSi2O6–Li2O bioceramics can be adjusted through the amount of Li2O substitution, depending on various implant applications. Moreover, all the samples were subjected to a decrease in hardness after SBF soaking for 28 days (shown in Figure 6b), which was due to the biodegradation behavior during soaking process. Among all the soaked samples, Li0.25 retained the highest hardness of 197 Hv, which is much higher than other samples’ hardness ranging from 4 to 12 Hv, revealing a strong resistance to biodegradation. In comparison with dense (>93%) TCP bioceramics synthesized by novel ultrafast high‐temperature sintering, the CaMgSi2O6–Li2O specimens have much higher hardness than dense TCP (maximum: 350 Hv).[75] In addition, compared to the most commonly used Ti‐based alloys for bone implant applications, the hardness of as‐sintered Li0.25 at 700 Hv is 1.7 times that of Ti–6Al‐4 V alloys at 416 Hv, while the hardness of soaked Li0.25 at 197 Hv is comparable to that of Ti–25Nb–3Zr–3Mo–2Sn alloys at 202 Hv.[76] The decreased hardness as a function of soaking time may allow Li0.25 to be more favorable for bone repair than Ti‐based alloys as the cortical bone has a relatively low hardness around 40–60 Hv.[77] The deteriorated hardness of Li0.50, Li0.75, and Li1.00 after SBF soaking could be ascribed to the escalating CaSiO3 formation as CaSiO3 has been reported to have a higher dissolution rate than diopside.[34,35]Dissolution and Precipitation Behaviors in SBF SolutionThe dissolution and precipitation behaviors of pure diopside and CaMgSi2O6–Li2O bioceramics were evaluated by the weight loss of the samples as well as the variations of the pH values and ion concentrations of the SBF solution after soaking in SBF for certain periods. To detect the degradability of pure diopside and CaMgSi2O6–Li2O bioceramics, the variations of sample weight during the immersion in SBF solution were measured and shown in Figure 7. It is clear that pure diopside and Li0.25 had smaller percentage of weight loss between 0% and 2%, while bioceramics with the higher concentrations of Li2O substitution had higher percentage of weight loss around 4% after 28 days of immersion time.7FigureWeight loss of pure diopside and CaMgSi2O6–Li2O bioceramics after soaking in the SBF solution as a function of immersion period.The mass variation is related to the biodegradation of the sample (weight loss) and the HAp formation on the sample surface (weight gain). From the results of overall weight loss, it can be suggested that the biodegradation of all the samples was predominant over HAp formation (especially at higher concentrations of Li2O) as the sample weights were reduced. For bone tissue repair applications, bioceramics with suitable biological degradation are needed. Thus, Li0.25 sample was considered to be favorable for repairing bone defects due to its slowest degradation rate, which could be attributed to the compact structure, as observed in Figure 2b. The samples with Li2O addition more than 25 mol% revealed the accelerated biodegradation properties in accordance with the decreased hardness (Figure 6b), which could be resulted from the growing secondary phase of CaSiO3 with higher dissolution rate. A similar result was obtained by Du et al.[78] that the weight loss of pure CaSiO3 was more severe than that of Mg‐incorporated CaSiO3, which could be due to the stronger MgO bonding than the CaO bonding.[37] Compared to another CaSi‐based bioceramic: baghdadite (Ca3ZrSi2O9), where the network of ZrCa bonds stabilizes the crystal structure and enables baghdadite a 70% higher compressive strength than CaMgSi2O6, the biodegradation rates of CaMgSi2O6–Li2O bioceramics in the current work were much lower than that of baghdadite (weight loss: 4–10%).[79]The pH values in the SBF solution following immersion of bioceramics samples are shown in Figure 8a. The upward trend of pH values for all the samples during 28 days of soaking indicates a sustained ion exchange of Ca2+, Mg2+, and Li+ with H+ and H3O+ in the SBF solution, followed by the formation of amorphous silica‐rich colloid layer and the subsequent formation of HAp.[49,80] In the first 7 days, all the CaMgSi2O6–Li2O bioceramics, especially those with higher levels of Li2O addition, had significantly higher pH values than pure diopside. After 1 week of soaking, the rapidly rising trends in pH value for all the samples tended to level off, indicating that the ion exchange between the samples and the SBF solution gradually reached saturation. After 28 days of soaking, Li0.25 had a similar pH value to pure diopside at around 8.0, while the pH values of Li0.50, Li0.75, and Li1.00 were about 8.3, implying that more ions were exchanged with higher Li2O substitution. As mentioned above, the ion exchange between the cations of bioceramics (Ca2+, Mg2+, and Li+) and the cations of SBF (H+ and H3O+) is a prerequisite of superficial HAp formation. In other words, the higher pH values stood for the more frequent ion exchange and better silica‐rich layer formation for HAp deposition. This observation is consistent with XRD and SEM results that the increase of Li2O content would stimulate the generation of CaSiO3 and thus improve the HAp formation. Du et al.[78] obtained an analogous outcome that the pH value of soaking medium provided by pure CaSiO3 was higher than Mg‐incorporated CaSiO3. Porosity may have negligible effect on the pH value as the comparably high pH values were given by Li0.50, Li0.75, and Li1.00 with different porosities. Since a weak alkaline microenvironment would improve osteoblast production and hinder osteoclast differentiation,[81] the Li0.25 sample is considered to be more suitable for implant applications because of its outstanding hardness (Figure 6), low degradation rate (Figure 7), reasonable HAp formation, and minor influence on the environmental pH value.8Figurea) pH values and ion concentrations of b) Ca2+, c) Mg2+, d) Li+, e) Si4+, and f) P5+ of SBF solutions after soaking pure diopside and CaMgSi2O6–Li2O bioceramics as a function of immersion period.The changes of ion concentration in SBF solution during immersion are shown in Figure 8b–f. In this test, the ion concentration of SBF was dominated by three reactions: 1) ion exchange to form silica‐rich layer, followed by an increase in pH value, and an increase in the concentration of Ca2+, Mg2+, and Li+; 2) adsorption of ions by the silica‐rich layer to form HAp, followed by a decrease in the concentration of ions, especially in Ca2+ and P5+; 3) dissolution of bioceramics and HAp, followed by an increase in the concentration of each ion. Thus, it can be summarized that ion exchange and material dissolution led to an increase in ion concentration, while the formation of HAp caused a decrease in ion concentration.First, the interaction mechanism between pure diopside and SBF solution is discussed as the followings. Pure diopside mainly exchanged Ca2+ and Mg2+ for H+ and H3O+ in SBF, forming a silica‐rich layer on the surface. As can be seen in Figure 8c, the Mg2+ concentration of pure diopside kept increasing, representing a greater driving force for ion exchange of Mg2+ and material dissolution than for adsorption. In Figure 8b, the Ca2+ concentration of pure diopside at day 14 was slightly lower than that of fresh SBF, indicating that more Ca2+ ions were adsorbed by the silica‐rich layer than they were released, and HAp was being formed at this stage. By day 28, the Ca2+ concentration has risen, suggesting that the adsorption of Ca2+ has reached saturation; then the ion exchange and material dissolution have outpaced the adsorption. The change of P5+ concentration in Figure 8f also revealed that P5+ had not yet been fully adsorbed at day 14, confirming that HAp was still being formed by Ca2+ and P5+ at this stage. By day 28, P5+ ions were completely consumed, indicating the saturation of ion adsorption and the completion of HAp formation.Second, the interaction mechanism of Li0.25 is discussed. As can be seen from Figure 8b, the Ca2+ concentration of Li0.25 was slightly higher than that of fresh SBF at day 14 and then remained stable until day 28, suggesting a two‐way balance between the release and adsorption of Ca2+, with a continuous HAp formation. Based on Figure 8f, P5+ ions have not been completely consumed after 28 days, verifying an ongoing ion adsorption and HAp formation throughout the immersion. Figure 8c shows that the Mg2+ concentration of Li0.25 was slightly lower than that of fresh SBF at both day 14 and 28, indicating that the adsorption rate of Mg2+ was higher than the release rate. In Figure 8d, Li0.25 displayed a steady release of Li+ over 28 days of immersion.Li0.50, Li0.75, and Li1.00 were grouped into the same category for discussion. As can be seen from Figure 8b, the Ca2+ concentration of these three samples increased continuously from ≈100 to ≈265 ppm throughout the 28‐day immersion period, representing that the ion exchange and material dissolution were much more driven than ion adsorption and thus HAp growth had been completed. Figure 8f shows that the P5+ ions were fully consumed by day 14, further confirming a complete formation of HAp. Based on Figure 8c, these three samples were identical to Li0.25 in that the adsorption rate of Mg2+ was slightly higher than the release rate. Figure 8d illustrates that the Li+ concentration increased with the increase of soaking period and the amount of Li2O.Figure 8e shows the Si4+ concentrations of all samples, and it can be seen that all the samples continued to release Si4+. Huang et al.[80] stated that the dissolution of silicon ions was dependent on not only the chemical composition, but the crystal structure of silicates, especially the bonding types. Since the Si content of materials remained constant here, the concentration of Si4+ in SBF solution should be predominantly related to its incorporation in crystalline or amorphous structures. Among the structures present in our materials, CaMgSi2O6, CaSiO3, and minor phase Li2SiO3 are inosilicates, where two oxygen atoms of each silicate tetrahedron are shared with adjacent tetrahedra to form single chains of connected silicate tetrahedra. The Li2Si2O5 phase induced by the substitution of Li2O is a phyllosilicate, where sheet‐like silicates composed of six‐membered silica rings occur in a chair structure.[82] The SiO2 phase is a tectosilicate, where silicate tetrahedra are connected in the form of 3D framework. In addition, the minor secondary phase Li2Ca2Si5O13 is a complex silicate, containing four‐ and five‐membered rings of silicate tetrahedra to form double chains.[74] Therefore, it can be deduced that it is more difficult for Si4+ to be dissolved from Li2Si2O5, SiO2, and Li2Ca2Si5O13 than from CaMgSi2O6, CaSiO3, and Li2SiO3 because of the stronger SiO covalent bonds. This may have resulted in the Si4+ concentrations of CaMgSi2O6–Li2O bioceramics becoming lower than that of pure diopside at the end of immersion period. Further, the concentration of Si4+ in Li0.25 was significantly lower than in the other samples, which could be due to its dense structure.The change in the concentration of Mg2+ ions in the SBF solution is particularly worthy of closer scrutiny. As shown in Figure 8c, the amount of Mg2+ released by pure diopside throughout the soaking period was much higher than that released by CaMgSi2O6–Li2O bioceramics. The elemental analysis of precipitated HAp in Table 4 shows that the HAp formed on pure diopside contained 8–13 times more Mg2+ ions than the HAp formed on CaMgSi2O6–Li2O bioceramics, indicating that pure diopside adsorbed considerable Mg2+ from the surrounding environment when forming HAp. In contrast to pure diopside, mixed phase CaMgSi2O6–Li2O bioceramics adsorbed only a trace amount of Mg2+, corresponding to the slight decrease of Mg2+ concentration shown in Figure 8c. Moreover, after 28 days of immersion, Mg2+ ions remained undetected after the HAp had fully grown on the surface of CaMgSi2O6–Li2O bioceramics. The low levels of Mg2+ detected in both SBF and HAp of CaMgSi2O6–Li2O bioceramics implied that Mg2+ ions were neither released from the materials nor then readsorbed by the HAp. Thus, it can be inferred that the addition of Li2O suppressed the release of Mg2+, thereby reducing the adsorption of Mg2+ by HAp, which can be further confirmed by the elemental mapping in Figure 5, where high Mg signals were detected in the matrix of CaMgSi2O6–Li2O bioceramics instead of in the HAp. The reason for the restricted release of Mg2+ may be explained as the followings. As clarified in the elemental mapping in Figure 3, the amorphous phase of Li2Si2O5 started to form after the substitution of Li2O, and the viscosity increased in the region where Mg2+ ions were present, resulting in a decreased ion diffusion. Hence, Mg2+ began to react with the surrounding Ca2+, Si4+, and O2− to form large grains of CaMgSi2O6, which were more difficult to dissolve in SBF and release ions than the small grains of CaMgSi2O6. In addition, the embedment of CaMgSi2O6 in the amorphous phase of Li2Si2O5 further reduced the dissolution rate, leading to Mg2+ being trapped in the large grains of CaMgSi2O6.As an indicator of silica‐rich layer formation, the relationship between the pH value and ion concentration can be explored in depth here. According to Figure 8b–d, though the ion exchange of Mg2+ was hindered by the addition of Li2O, a large amount of Ca2+ and Li+ of CaMgSi2O6–Li2O bioceramics was exchanged for H+ and H3O+, leading to much higher pH values (≈8.3) than that of pure diopside (≈8.1). The high pH values of CaMgSi2O6–Li2O bioceramics demonstrated the formation of a large quantity of silica‐rich layer, laying the foundation for the following substantial HAp formation. However, the ion release of Ca2+ reached about 160 ppm, which was much higher than that of Li+ (11 to 33 ppm), coupled with the fact that Li+ is a monovalent ion and can exchange half as many H+ ions as Ca2+/Mg2+, indicating that Ca2+ was more dominant than Li+ in the formation of silica‐rich layer. It is believed that the high Ca2+ and Li+ concentrations were contributed by CaSiO3 and Li2Si2O5 phases, which were more likely released from the vulnerable amorphous matrix than from the solid crystals, as shown in Figure 3. Since Li0.25 had a dense structure and relatively small amount of CaSiO3, its overall ion exchange and material dissolution were lower than other CaMgSi2O6–Li2O bioceramics. Moreover, the pH value of Li0.25 was close to that of pure diopside, suggesting a comparable overall ion exchange and similar ability of silica‐rich layer formation.From the results of XRD (Figure 1c) and SEM (Figure 4) analyses, it can be observed that all the CaMgSi2O6–Li2O bioceramics were found to have a greater extent of HAp formation than pure diopside. It can be further confirmed that Li0.50, Li0.75, and Li1.00 had superior silica‐rich layer and subsequent HAp formation to pure diopside based on the variations of pH value (Figure 8a) and the consumption of P5+ ions (Figure 8f). In addition, the high Ca2+ levels found in the SBF of Li0.50, Li0.75, and Li1.00 may increase the supersaturating concentration of SBF, which would facilitate the apatite formation on the surface of the bioceramics.[83] While Li0.25 had the similar capacity of silica‐rich layer formation as pure diopside, the high remaining P5+ ions in SBF implied that Li0.25 should have had the worse HAp formation than pure diopside as fewer P5+ ions were adsorbed to form HAp, which is inconsistent with the results shown in XRD and SEM. This inconsistency could be due to the fact that the Mg2+ and Si4+ released from pure diopside into SBF have been readsorbed by the HAp formed on the surface of pure diopside and hence reduced the crystallization of HAp. Hafezi‐Ardakani et al.[84] have stated that Mg2+ ions may be able to enter the forming HAp nuclei and further inhibited the HAp evolution as Mg2+ ions are not readily accommodated in the HAp structure. In the present study, all the CaMgSi2O6–Li2O bioceramics blocked the release of Mg2+, and only trace amounts of Mg2+ were adsorbed by HAp, as shown in Table 4, which significantly improved the HAp growth. Therefore, Li0.25 eliminated the negative effect of Mg2+ on HAp growth by means of low‐level Li2O addition, allowing for much more efficient HAp formation with a lower P5+ consumption than pure diopside. It can be expected that extending the soaking time for Li0.25 would allow the complete utilization of P5+ ions and the further increase in the amount of HAp precipitate. Moreover, Li0.25 had the lowest overall ion release among all the samples due to its dense structure, which is in good agreement with the result of weight loss.Cell Viability3‐(4,5‐dimethylthiazol‐2‐yl)‐2,5‐diphenyl tetrazolium bromide (MTT) assay is an effective method to assess the cytotoxicity of materials by measuring the concentration of mitochondrial dehydrogenase, which shows cell proliferation and hence reflects the biocompatibility of materials. In the current work, cell viability was studied with pure diopside and CaMgSi2O6–Li2O bioceramics samples by MTT test. Their optical densities (24 h) with incubated D1 cells were compared to the positive control (with sodium dodecyl sulfate [SDS] addition) and negative controls (with only cells and with Al2O3 addition), as shown in Figure 9. From the result, it can be seen that cells were more viable in the negative controls and pure diopside than in the positive control and CaMgSi2O6–Li2O bioceramics samples. The cell viability percentage ranging from 60.7% to 70.5% with CaMgSi2O6–Li2O bioceramics samples suggested that the presence of high‐dose Li ions had a strong effect on the proliferation of mouse bone marrow cells (D1). Among all the Li2O‐incorporated samples, Li0.25 possessed a relatively acceptable cell viability at 70.5%, indicating that low doses of Li ion may be considered to minimize the cytotoxicity of CaMgSi2O6–Li2O bioceramics.9FigureCell viability of D1 cells with pure diopside and CaMgSi2O6–Li2O bioceramics.ConclusionPure diopside and CaMgSi2O6–Li2O bioceramics were prepared by precipitation method, followed by sintering at 1000 °C for 4 h. Alongside the primary CaMgSi2O6 phase, secondary phases of CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13 formed with the increasing replacement of MgO by Li2O in the material stoichiometry, with implications for the mechanical strength and bioactivity of the material. Based on SEM images and elemental analysis, low‐level Li2O substitution induced a densified microstructure consisting of CaMgSi2O6 and CaSiO3 grains embedded in Li2Si2O5 amorphous matrix, leading to a significant enhancement of hardness. However, the hardness decreased for higher Li2O levels due to the porosification induced by the formation of CaSiO3 and significant grain growth. The elemental redistribution was observed in CaMgSi2O6–Li2O bioceramics where Mg was mostly concentrated in the CaMgSi2O6 grains, while Ca was found both in the crystalline and amorphous phases. After 28 days of SBF soaking, the mass and the hardness of all the samples decreased with the increment of Li2O due to abundant CaSiO3 formation with relatively weak CaO bonding and high dissolution rate. Furthermore, highly improved HAp (Ca5(PO4)3OH) formation was found after SBF soaking on all the CaMgSi2O6–Li2O bioceramics compared to pure diopside, which can be ascribed to the limited release of Mg2+ from the large CaMgSi2O6 grains embedded in the Li2Si2O5 amorphous phase, which prevented Mg2+ entering HAp nuclei and thus promoted HAp crystallization. In this study, Li0.25 is more favorable for implant applications due to its excellent hardness (as‐sintered: 700 Hv; SBF‐soaked: 197 Hv), low degradation rate (≈0.6 wt%), and minor impact on the environmental pH value. In addition, Li0.25 eliminated the negative effect of Mg2+ on HAp growth using low‐level Li2O substitution, permitting much more efficient HAp formation with a lower P5+ consumption than pure diopside and hence resulting in a superior bioactivity. A more detailed composition adjustment (such as low concentration of ion substitution) can be considered in the future work to optimize the biocompatibility of CaMgSi2O6–Li2O bioceramics as the presence of high‐dose Li+ suppressed the proliferation of mouse bone marrow cells (D1).Experimental SectionSamples PreparationDiopside and mixed phase CaMgSi2O6–Li2O bioceramics were synthesized by a precipitation method based on a procedure described in detail elsewhere.[48,49,85] For pure diopside, 0.01 mol of calcium chloride (CaCl2, 95%, Fisher Scientific, UK), 0.01 mol of magnesium chloride hexahydrate (MgCl2·6H2O, 95%, EMSURE, Germany), and 0.02 mol of tetraethyl orthosilicate (TEOS, 98%, Seedchem, Australia) as reactants were dissolved in 200 mL of ethanol (C2H5OH, 95%) to produce 0.01 mol of CaMgSi2O6 stoichiometry solution. The solution was magnetically stirred at 80 °C for 2 h. Then ammonium hydroxide (NH4OH, 28%, Nihon shiyaku reagent, Japan) was added into the solution to produce white precipitates. After stirring overnight, the solvent had evaporated, and the remaining white precipitates were dried in an oven at 80 °C for 24 h. The dried precipitates were ground into fine powder, and then calcined at 600 °C for 3 h to remove impurities. Compact pellets (Ø13 × H1.8 mm) were prepared by compressing the calcined CaMgSi2O6 powder under 75 MPa for 30 s. Then the white compact pellets were sintered at 1000 °C for 4 h in the air to obtain the final CaMgSi2O6 products.To synthesize the mixed phase CaMgSi2O6–Li2O bioceramics, a similar procedure as for pure diopside was carried out with the variation that a wide range of Li2O amounts (25, 50, 75, and 100 mol%) were introduced by the corresponding substitution of lithium chloride (99%, Acros Organics, Germany) for magnesium chloride (MgCl2). The CaMgSi2O6–Li2O bioceramics with various Li2O amounts (0–100 mol%) were denoted as pure, Li0.25, Li0.50, Li0.75, and Li1.00.Materials CharacterizationThe mineralogy of the pure diopside and mixed phase bioceramics was analyzed by a powder X‐ray diffractometer (D2 Phaser, Bruker, USA) with Cu Kα radiation (λ = 1.54184 Å) at 30 kV and 10 mA with diffraction angles (2θ): 10–70°, step size: 0.01° 2θ and step speed: 0.5 s/step. The phase identification and degree‐of‐crystallinity analysis based on the diffraction patterns were processed by Match! software (Crystal Impact, Germany). The surface morphology and chemical composition of the samples were observed using a FE‐EPMA (JXA‐8530F, JEOL, Japan) which consisted of a SEM and a wavelength dispersive spectrometer. FE‐EPMA was operated at 15 kV with an analytical depth of ≈700 nm, where mapping mode was used to analyze the elemental distribution and point mode was used to precisely determine the elemental concentration. The hardness was measured using a micro Vickers hardness tester (FM‐110, Future‐Tech, Japan) with a force of 1 kg over a loading time of 10 s under room temperature.The bulk density and apparent density of the sintered pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics were measured using Archimedes technique in deionized water. The density measurement was performed as the following. First, the samples were dried in an oven at 110 °C for 24 h to remove free water, then naturally cooled down to room temperature in an electronic moisture‐proof box at a constant zero humidity for 24 h. The weight of the dried samples (denoted as m1) was measured by a precision analytical balance (AS 220.R2 PLUS, Radwag, Poland) with 1 mg of accuracy. Second, the samples were saturated with deionized water under a vacuum condition for 2 h. The saturated samples were then suspended in deionized water and the immersed mass of the samples (denoted as m2) was recorded. Third, the samples were taken out from deionized water and the remaining water on the sample surface was carefully removed. The weight of the saturated samples in air was measured and denoted as m3. According to Archimedes principle, bulk density was the ratio of the material mass without free water to the macroscopic material volume, which was occupied by solid material, open pores, and closed pores. The calculation of bulk density was expressed by the following equation.1Bulk density(DB)=(m1×ρ0)/(m3−m2)\[\begin{array}{*{20}{c}}{{\rm{Bulk}}\;{\rm{density}}\left( {{D_{\rm{B}}}} \right) = \left( {m1 \times {\rho _0}} \right)/\left( {m3 - m2} \right)}\end{array}\]Apparent density was the ratio of the material mass to the material volume including solid material and closed pores, which was calculated as the equation described below.2Apparent density(DA)=(m1 × ρ0)/(m1−m2)\[\begin{array}{*{20}{c}}{{\rm{Apparent}}\;{\rm{density}}\left( {{D_{\rm{A}}}} \right) = \left( {m1\: \times \:{\rho _0}} \right){\rm{/}}\left( {m1 - m2} \right)}\end{array}\]where ρ0 was the density of deionized water at 25 °C. Based on the bulk density and apparent density, the porosity given by the open pores was calculated as the following.3Porosity(%)=[1−(DB/DA)]× 100%\[\begin{array}{*{20}{c}}{{\rm{Porosity}}\left( \\end{array}\]In Vitro Biodegradation AnalysisIn vitro biodegradation tests were performed by soaking pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics pellets in a SBF solution at 37 °C. The SBF solution had a similar ion concentration with human blood plasma, as proposed by Kokubo et al.[86] For the evaluation of degradation, each pellet was immersed in 40 mL SBF solution with pH 7.4 at 37 °C for 7, 14, 21, and 28 days. After the set soaking time, each pellet was washed by deionized water and then completely dried, and the final weight loss was measured and expressed in percentage (wt%) as below:4WL(%)=[(W0−WF)/W0]×100%\[\begin{array}{*{20}{c}}{{W_{\rm{L}}}\left( \\end{array}\]where WL was the weight loss of pellet after SBF soaking, W0 was the initial weight of pellet, and WF was the final weight of pellet after soaking.The pH values of the SBF solution before and after immersion were measured by a pH meter (pH 510, Eutech, Singapore) and the ion concentrations of Ca2+, Mg2+, Li+, Si4+, and P5+ in the SBF solution were analyzed using an inductively coupled plasma‐mass spectrometer (ELEMENT XR, Thermo fisher scientific, USA). The initial concentrations of Ca2+, Mg2+, Li+, Si4+, and P5+ ions in fresh SBF were 92.6, 43.9, 0, 0, and 40 ppm, respectively.Cell Viability AssayThe cell viability was evaluated using the liquid extracts of bioceramic samples and MTT (98%, Alfa Aesar, USA) assay. Dulbecco's modified Eagle's medium (DMEM, Gibco, USA) containing 10% fetal bovine serum (Sigma‐Aldrich, USA) and 0.5% penicillin–streptomycin (Sigma‐Aldrich, USA) was used in the cell culture and the extraction of sample substances. MTT solution was prepared at the concentration of 5 mg mL−1 by dissolving MTT in sterile and filtered phosphate buffered saline (PBS). The samples for preparing liquid extracts were pure diopside and mixed phase bioceramics with four levels of Li2O substitution (25–100 mol%). A group with the addition of SDS (99%, J.T Baker, USA) was prepared as positive control. A group containing only cells and a group with the addition of aluminum oxide (Al2O3, 99.5%, Sigma‐Aldrich, USA) were prepared as negative controls.First of all, mouse bone marrow stromal D1 cells were seeded at a density of 2 × 105 cells/well in 96‐well plates and cultured in 100 µL of DMEM at 37 °C in an incubator under 5% CO2 atmosphere for 24 h. Second, liquid extracts of all the test groups were prepared in DMEM at a concentration of 0.1 g mL−1 (except SDS group at 0.2 mg mL−1) at 37 °C for 24 h, followed by a filtration of extracts with a 0.25 µm filter. After 24 h of incubation, the initial DMEM was removed and then 100 µL of the filtered extract was added to each well. The D1 cells were cultured in the extracts in the incubator for 24 h for the following cytotoxicity test. After 24 h of cell culturing in the presence of extracts, the osteoblast cell viability was evaluated by MTT assay. The extracts were removed and 100 µL of MTT solution (5 mg mL−1 in PBS) was added to the cultures in each well, and the cells were incubated at 37 °C in 5% CO2 for 2 h. Then the MTT solution was removed and 100 µL of dimethyl sulfoxide (99.5%, Sigma‐Aldrich, USA) was added to each well to dissolve purple crystals of formazan. The optical density of each well at 570 nm of wavelength was measured by an ELISA microplate reader (Model 550, Bio‐Rad, USA). The cell viability was calculated using:5Cell viability(%)=(ODs/ODc) × 100%\[\begin{array}{*{20}{c}}{{\rm{Cell}}\;{\rm{viability}}\left( \\end{array}\]where ODs was the optical density values of sample groups and ODc was the optical density value of control group containing only cells. The quantitative data were expressed as means ± standard deviation of six repeats per experiment.AcknowledgementsThe authors acknowledge financial support from the Ministry of Science and Technology of Taiwan under contract MOST 111‐2636‐E‐110‐008, the DAAD's Program for Project‐related Personal Exchange (PPP) (Project ID 57390207), the Field Emission Electron Probe Microanalyzer equipment provided by Joint Center for High Valued Instruments at National Sun Yat‐Sen University, and the Inductively Coupled Plasma‐Mass Spectrometer equipment provided by Core Facility Center of National Cheng Kung University, Taiwan. The authors thank National Center for High‐performance Computing (NCHC) for providing computational and storage resources.Conflict of InterestThe authors declare no conflict of interest.Data Availability StatementThe data that support the findings of this study are available from the corresponding author upon reasonable request.S. Kargozar, M. Montazerian, E. Fiume, F. 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Bioceramics in the CaMgSi2O6–Li2O System: A Glass‐Ceramic Strategy for Excellent Mechanical Strength and Enhanced Bioactivity by Spontaneous Elemental Redistribution

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10.1002/admi.202202491
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Abstract

IntroductionImplantable bioceramics bone tissue‐engineering materials that combine multiple modes of bioactivity to help facilitate natural bone regeneration are an emerging approach toward the treatment of significant bone defects. In recent years, a variety of biomaterials such as bioactive glasses,[1–3] glass‐ceramics,[4–6] crystalline calcium silicates (such as wollastonite [CaSiO3] and larnite [Ca2SiO4],[7,8] and calcium phosphate‐based bioceramics[9,10] have been the subject of numerous investigation toward bone tissue engineering. In particular, some researchers have proposed that CaSiO3 exhibits superior bioactivity and osseointegration properties to the widely studied hydroxyapatite (HAp), which motivates further investigation into CaSiO3‐based materials for bone tissue repair and regeneration.[11,12] The presence of bioavailable calcium (Ca) and silicon (Si) ions in tissue engineering materials can directly influence the quality of bone as Ca is the main constituent of biological apatite and Si was reported to have the ability of inducing osteogenesis and angiogenesis.[13,14] However, it has been proposed that the poor mechanical strength and rapid dissolution rate of CaSiO3 may hinder its biomedical application to be used as bulk implants and would be detrimental for cell growth because of a rapid increase in pH value in the surrounding environment during the first few weeks.[15,16]Magnesium‐containing bioceramics have gained interest recently due to the favorable combination of properties that is achievable in such systems in terms of biodegradation, bioactivity and mechanical strength.[17,18] As with Ca and Si, magnesium (Mg) is also an indispensable element for the human body, and bone regeneration in particular as Mg deficiency would stimulate the differentiation of osteoclasts and inhibit the production of osteoblasts, thus leading to osteoporosis.[19] Further, it has been proved that Mg‐containing calcium silicate‐based bioceramics can stimulate the apatite formation and induce a tight bond with bone tissues.[20,21] Lai et al.[22] suggested that an increase of new vessel development and blood perfusion after surgery was achieved by Mg‐incorporated PLGA/TCP porous scaffolds. Further, Mg‐containing thin films deposited on porous polymer scaffolds could significantly develop new bone formation and strengthen mechanical properties of newly formed bone,[23] indicating that Mg‐containing bioceramics could be used as a novel material for bone tissue engineering applications.Due to the various attractive attributes of calcium magnesium silicates as effective bioceramics, such materials from various silicate families have been selected as bone‐tissue engineering materials, including the sorosilicate akermanite (Ca2MgSi2O7),[24,25] orthosilicate monticellite (CaMgSiO4),[26,27] and the pyroxene diopside (CaMgSi2O6).[28–30] Among all the materials, diopside has superior mechanical properties,[31] and also possesses good cytocompatibility and surface apatite precipitation ability.[32,33] Consequently, diopside (CaMgSi2O6) was chosen as the primary bioceramic material in the present study. It is worth noting that the mechanical properties of CaMgSi2O6 was improved in comparison with that of CaSiO3, while the HAp mineralization in simulated body fluid (SBF) was compromised due to the incorporation of Mg.[34–37]Toward the repair of defects, the mechanical and biological properties of bioceramic implants should be compatible with native bone tissue and stimulate a favorable biological response.[38] One of the current techniques for addressing these issues is to introduce additional elements into main structures (doping[39–42] and/or composites[43,44]) to modify the crystallinity, lattice parameters, and morphology of the bioceramics, further affecting their mechanical and biological properties. Table 1 summarizes the previous works on the incorporation of additional elements in CaMgSi2O6 bioceramics.1TableSurvey of experimental works on introducing additional elements into CaMgSi2O6 bioceramicsElementsElement contentAnnealing treatmentEffectsReferenceSr/F2 mol%/1 mol%1200 °C/5 h All the doped scaffolds presented higher apatite‐forming ability Osteoblast‐like MG‐63 cells exhibited the highest compatibility to the Sr‐doped scaffolds Osteoblast‐like MG‐63 cells exhibited the lowest compatibility to the codoped scaffolds[45]Sr2 mol%1200 °C/3 h Sr‐doping improved the sinterability and apatite‐forming ability of the diopside‐based scaffolds Sr‐doping retarded the biodegradation of diopside‐based scaffolds Mesenchymal stem cells presented better adhesion and spreading with typical cell extensions on the Sr‐doped diopside‐based scaffolds[46]LiNaK2 mol%900 °C/2 h The crystallinity and lattice volume of diopside were changed with the most and least deviations for K and Na, respectively The dopants altered the in vitro bioactivity of diopside in the following ranking: K‐doped >  Li‐doped >  Na‐doped >  undoped The 2 mol% dopants improved the biocompatibility of diopside, where the most beneficial effect was found for Na and K K‐doping was the optimal doping for the bioactivity and cytocompatibility assessments[47]Ce0–100 mol%1000 °C/4 h Addition of 25 mol% Ce had the best biomineralization performance in vitro Less hydroxyapatite precipitates were found with further increasing Ce addition[48]Mo0–100 mol%1000 °C/4 h At a lower Mo content, the mixed phase materials showed higher hardness and slower biodegradation At a higher Mo content, mixed phase materials exhibited lower hardness and bioactivity[49]Lithium (Li) has been commonly used to treat bipolar disorder,[50,51] and the psychiatric patients who took Li medication have been reported to have enhanced bone mass and lower bone turnover states.[51] This may be explained by the mechanism that Li can stimulate the osteogenesis through activating Wnt and Hh signaling pathway.[52,53] Moreover, it was found that Li doping can improve the thermal stability of calcium phosphate.[54] Table 2 summarizes the previous works on the incorporation of Li element in different types of bioceramics. Padmanabhan et al.[55] prepared lithium‐substituted hydroxyapatite (Li‐HAp) nanoparticles by a sonochemical synthesis process. It was reported that the incorporation of Li into HAp caused the transformation of particles from a rod shape to a needle‐like structure due to the sonication effect. Further, Li‐HAp showed excellent antimicrobial activity and blood compatibility, and could be used as a drug delivery carrier owing to its improved cytocompatibility. Yuan et al.[56] synthesized Li‐doped calcium polyphosphate bioceramic scaffolds by a gravity sintering method. The study concluded that Li doping did not affect the crystallized phase but changed the surface morphology. Further, the Li‐doped samples were found to have better degradation properties and enhanced proliferation and differentiation of osteoblasts. Wang et al.[57] synthesized Li‐doped HAp by a precipitation method, where the scaffolds had better mechanical strength and Li doping was found to be beneficial to the proliferation of osteoblasts. Li‐doped β‐TCP was found to have two times higher compressive strength than initial β‐TCP.[58] The incorporation of nanoporous lithium doping magnesium silicate into calcium sulfate hemihydrate was found to enhance the degradability, biocompatibility, vascularization, and osteogenesis.[59] As summarized above, the incorporation of Li into calcium phosphate‐based bioceramics was expected as a viable way to enhance the mechanical and biological properties of bioceramics. However, few studies have been conducted on bioceramics in the diopside–lithium oxide system. Only one study synthesized Li‐doped diopside through a coprecipitation method, as shown in Table 1. The result showed that the substitution of 2 mol% Li for 1 mol% Mg has improved the in vitro bioactivity and biocompatibility.2TableSurvey of experimental works on introducing Li element into bioceramicsBioceramic typeLi amountMethodEffectsReferenceHAp0.1–0.5 mol%Sonochemical Li‐HAp particles were transformed from a rod‐like structure to needle‐like structure Reduction in crystallinity of HAp Excellent antimicrobial activity and blood compatibility Promoted cell proliferation[55]Calcium Polyphosphate (CPP)2 mol%Gravity sintering The crystallized phase was not affected Rougher and more uneven surface Better degradation properties Promoted proliferation and differentiation of osteoblasts Better osteoconductivity[56]HAp0.5–1 mol%Precipitation High mechanical strength Fast osteoblast‐mediated degradation rate Promoted osteoblast activity[57]β‐TCP1.9 mol%Precipitation and ball milling Compressive strength enhanced by two times Lowered degradation rate and cytotoxicity[58]Calcium Sulfate Hemihydrate–Precipitation and mixing Enhanced degradability in vitro and in vivo Prevention of acidification in the biological environment Improved new bone formation Better vascularization and osteogenic differentiation[59]HAp0.5–2 mol%Precipitation 1.5%Li‐HAp obtained optimal cell proliferation activity Enhanced new bone formation Better osteogenesis and angiogenesis potential[60]As the amounts of Li incorporated or substituted into diopside and other bioceramics were mostly limited to a maximum of 2 mol%, the present work attempts to investigate the effects of a wide range of Li content (25, 50, 75, and 100 mol%) on the diopside–lithium oxide (CaMgSi2O6–Li2O) system to determine the appropriate Li amount and thereafter optimize its mechanical and biological properties. The specimens were synthesized through a precipitation method, followed by sintering at 1000 °C for 4 h. The mineralogical, morphological, mechanical, and biological properties were investigated to comprehensively understand the factors and mechanisms affecting the properties of CaMgSi2O6–Li2O bioceramics system.Results and DiscussionCrystalline Phase IdentificationThe phase identification of the pure diopside and CaMgSi2O6–Li2O bioceramics was obtained by X‐ray diffraction (XRD), as demonstrated in Figure 1. In order to precisely investigate the crystallized phases in the bioceramics, the as‐sintered samples were ground into powder for random crystalline orientation and then measured by XRD, as shown in Figure 1a. It can be seen that a major phase of CaMgSi2O6, corresponding to ICSD code: 168107, was observed in the pure diopside. As the Li2O substitution increased to 25 mol%, the formation of secondary phases (CaSiO3, minor Li2Si2O5, and SiO2) was detected in accordance with ICSD code: 201537, 15414, and 81382, respectively. The CaSiO3 was identified as β‐CaSiO3 with triclinic structure which is generally formed below 1100 °C.[61] As the Li2O content increased to 50 and 75 mol%, the peak intensity of CaMgSi2O6 phase decreased along with the increment of peak intensity in CaSiO3, Li2Si2O5, and SiO2 phases, and two new secondary phases Li2SiO3 and Li2Ca2Si5O13 formed, matching with ICSD code: 100402 and 79379, respectively. With further increasing the Li2O content to 100 mol%, CaMgSi2O6 phase was completely diminished and replaced by CaSiO3 (major), Li2Si2O5, SiO2, Li2SiO3 (minor), and Li2Ca2Si5O13 (minor) phases as a result of the absence of Mg element. The details of the phase transition, peak intensities, and degree of crystallinity are clearly listed in Table 3. Following the substitution of Li2O for MgO, the amorphous phase started to emerge, and the amorphous content increased with the higher Li2O substitution. The XRD results can be further supported by the phase diagram of Li2O–SiO2 system, where Li2SiO3 and Li2Si2O5 would form at 936 to 1028 °C when high molar concentrations of SiO2 were added.[62]1FigureXRD patterns of pure diopside and CaMgSi2O6–Li2O bioceramics in the form of a) powder and b) pellet, and c) after SBF soaking for 28 days in the form of pellet.3TablePhase transition, peak intensities, and degree of crystallinity in pure diopside and CaMgSi2O6–Li2O bioceramicsSamplePhase transitionDegree of crystallinity [%]CrystallinityAmorphousPureCaMgSi2O6 only––Li0.25CaMgSi2O6 (↓) + CaSiO3 (↑) + minor Li2Si2O5 (↑)+ minor SiO2 (↑)95.544.46Li0.50CaMgSi2O6 (↓↓) + CaSiO3 (↑) + Li2Si2O5 (↑)+ SiO2 (↑) + minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)91.638.37Li0.75CaMgSi2O6 (↓↓) + major CaSiO3 (↑↑) + Li2Si2O5 (↑)+ SiO2 (↑) + minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)87.7412.26Li1.00major CaSiO3 (↑↑↑) + Li2Si2O5 (↑) + SiO2 (↑)+ minor Li2SiO3 (↑) + minor Li2Ca2Si5O13 (↑)84.5615.44In order to study the crystallization behavior of the bioceramics in biological environment, the XRD patterns of the bioceramics pellets before and after SBF soaking for 28 days are compared in Figures 1b and 1c, respectively. The slight difference in peak intensity and peak position between the form of pellet and powder could be originated from the preferred orientation and the uneven sample surface. It can be seen that the basic compositions in the SBF‐soaked materials were practically the same as those in the as‐sintered samples, confirming that no significant phase transformation was presented in the matrices. Further, the peaks at 25.9° and 31.7–34.1°, representing the precipitate of HAp (Ca5(PO4)3OH, ICSD code: 34457), were observed in all the samples, where CaMgSi2O6–Li2O bioceramics had higher intensities of Ca5(PO4)3OH phase than pure diopside, indicating a superior HAp formation capacity of CaMgSi2O6–Li2O bioceramics. In addition, the peak intensity of CaSiO3 was lowered after SBF soaking, while the peak intensity of CaMgSi2O6 remained evident. The enhancement of HAp formation and the degraded CaSiO3 phase could be ascribed to the replacement of MgO by Li2O, as reported by Vallet‐Regí et al.[37] that the decrease of MgO in CaO–MgO–SiO2–P2O5 bioceramics promoted the apatite formation rate and the vulnerability to aqueous solution because of better seeded HAp crystallization and lower CaO bonding energy than MgO.Morphology and Elemental DistributionThe scanning electron microscope (SEM) images (5000×) for sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets are shown in Figure 2. The surface morphology of pure diopside without Li2O addition was homogenous with a distribution of small pores, as shown in Figure 2a. With 25 mol% of Li2O substitution for MgO, a microstructure consisting of small grains embedded in a glass‐like substrate was clearly observed in Figure 2b. The presence of amorphous phase was also detected by XRD and degree of crystallinity analysis in Table 3. Moreover, with the increase of Li2O addition, the aggregation and grain growth phenomena became more apparent, leading to an increasingly porous structure, as shown in Figure 2c–e. The size of aggregated grains was getting larger from 3 to 12 µm in length and 2 to 5 µm in width. Li et al.[35] have reported that the incorporation of MgO into CaSiO3 gave rise to the formation of calcium–magnesium silicates, and therefore resulted in a compact surface and a decrease of total pore volume, and vice versa. However, in the present study, a fused and densified surface was observed in Li0.25 with an increased ratio of CaSiO3 due to the substitution of Li2O for MgO, which contradicted with what Li et al. has found. In this case, since the appearance of CaSiO3 should have increased the porosity, the densified surface of Li0.25 may be contributed by another secondary phase: Li2Si2O5. This has been supported by Wang et al. that the aggregated ZrO2 grains were embedded into the liquid matrix of Li2Si2O5 by a sol–gel synthesis at the sintering temperature of 1000 °C, leading to the removal of pores and the densification of ceramics.[63] This is so‐called “glass‐ceramics” composed of a Li2Si2O5 crystalline phase and an amorphous matrix, which is induced by a heterogeneous nucleation mechanism.[64–66] In the current study, the viscous amorphous matrix of Li2Si2O5 was able to fill the pores throughout the bioceramics and hence densified the materials. With the further increase of Li2O content, the porosification induced by the increasing CaSiO3 formation and grain growth would surpass the densification caused by Li2Si2O5, resulting in a porous material surface. A similar outcome was obtained by Li et al. that the incorporation of Li2Si2O5 into Ca0.7Sr0.3TiO3 enhanced the densification of materials, and the grain size increased with the elevating content of Li2Si2O5.[67]2FigureSEM images of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets: a) pure, b) Li0.25, c) Li0.50, d) Li0.75, and e) Li1.00.Elemental mapping based on Ca, Mg, and Si elements on the surface of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets is shown in Figure 3. Li element was excluded from the mapping analysis as the atomic order of Li is too small to be detected by field emission electron probe microanalyzer (FE‐EPMA). A homogeneous distribution of Ca, Mg, and Si can be observed in pure diopside, indicating the existence of CaMgSi2O6 single phase. With 25 mol% of Li2O substitution, the elemental distribution of Ca and Mg was concentrated in specific areas, suggesting the formation of CaSiO3 and the grain growth of CaMgSi2O6. As the substitution of Li2O increased to 50 and 75 mol%, the grain size of CaSiO3 gradually increased, while the CaMgSi2O6 phase decreased with the reduction of Mg, as confirmed by XRD (Figure 1) and SEM (Figure 2). When the Li2O content reached 100 mol%, the Mg signal has completely disappeared and the large grains of CaSiO3 partially embedded in the glass phase of Li2Si2O5 were clearly detected. As can be seen from CaMgSi2O6–Li2O bioceramics, it is noteworthy that Mg elements were mostly concentrated in the crystalline phase of CaMgSi2O6, while Ca elements were not only found in the crystalline phases of CaSiO3 and CaMgSi2O6, but heavily detected in the amorphous glass phase. This phenomenon may be explained by a mechanism that Mg2+ has higher field strength than Ca2+ and Li+, which increases the viscosity of the glass phase, leading to a reduced diffusion of ions and ionic complexes.[68] Because of the poor ion diffusion in the regions containing concentrated Mg2+ ions, Mg2+ may have a strong tendency to nucleate CaMgSi2O6 crystals with the localized Ca2+, Si4+, and O2− ions. By contrast, partial Ca2+ can be dispersed in the glass phase of Li2Si2O5 due to its high ionic diffusivity without the influence of Mg2+.3FigureElemental mapping analysis on the surface of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets.The SEM images (5000×) for pure diopside and CaMgSi2O6–Li2O bioceramics pellets after soaking in SBF solution for 28 days are shown in Figure 4. The presence of newly formed HAp layers can be clearly observed on all the samples. For pure diopside, it can be seen that the diameter of the spherical precipitates on the sample surface varied from 1 to 2 µm. Regardless of the initial matrix structure, the diameter of the spherical precipitates on all the CaMgSi2O6–Li2O bioceramics samples ranged from 5 to 12 µm, which was much larger than that on the pure diopside, indicating a superior HAp formation capacity of CaMgSi2O6–Li2O bioceramics, as confirmed by the XRD results in Figure 1c. The improved HAp formation could be attributed to the appearance of CaSiO3 and Li2Si2O5 caused by the Li2O addition, where CaSiO3 has been reported to have better apatite mineralization but more rapid dissolution rate than CaMgSi2O6.[34,35] Furthermore, the quantitative compositions of HAp formed on pure diopside and CaMgSi2O6–Li2O bioceramics were measured by the point mode of FE‐EPMA, as indicated by yellow arrows in Figure 4, and the results were tabulated in Table 4. This quantitative analysis mainly measured the concentrations of Ca, Mg, Si, O, and P as they are the major compositions of CaMgSi2O6‐based bioceramics and HAp precipitates. Other elements, such as undetectable Li, gold coating used for enhancing conductivity, or the elements adsorbed from SBF solution, were classified as “others.” It is clear that the HAp formed on CaMgSi2O6–Li2O bioceramics contained Ca, O, and P as its major components, corresponding to the composition of Ca5(PO4)3OH. In addition to Ca, O, and P, the HAp formed on the pure diopside had a noticeable amount of Mg and Si, indicating that it readsorbed the Mg and Si elements which were released from the matrix materials into SBF solution.4FigureSEM images of pure diopside and CaMgSi2O6–Li2O bioceramics pellets after soaking in SBF for 28 days: a) pure, b) Li0.25, c) Li0.50, d) Li0.75, and e) Li1.00.4TableComposition of HAp formed on pure diopside and CaMgSi2O6–Li2O bioceramics pellets (measured by the point mode of FE‐EPMA)SampleComposition of HAp [wt%]CaMgSiOPOtherspure17.14.110.523.46.138.8Li0.2527.20.50.011.114.746.5Li0.5028.40.30.311.811.347.9Li0.7529.30.30.611.413.245.2Li1.0027.10.30.816.112.343.4In particular, the elemental mapping analysis of Li0.25 bioceramics after soaking in SBF for 28 days was performed and shown in Figure 5. It can be seen that the soaked Li0.25 bioceramics was comprised of Ca, Mg, Si, O, and P ions. The elemental mapping implied the development of a uniform microstructure with regard to an appropriate spatial distribution of elements. High concentrations of Ca and P ions were found on the surface of HAp precipitates, while high concentrations of Mg and Si ions, and a medium concentration of Ca ions were found on the surface of the Li0.25 matrix, which were consistent with their chemical compositions.5FigureElemental mapping analysis of the Li0.25 bioceramics pellets after soaking in SBF for 28 days: a) SEI, b) Ca, c) Mg, d) Si, e) O, and f) P elements.Density and Porosity MeasurementTable 5 shows the bulk and apparent density of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pellets. Following the density results, the porosity given by open pores was calculated accordingly. As it can be seen, the apparent density (DA) of pure diopside is 3.19 g cm−3, which is very close to the theoretical density of diopside (CaMgSi2O6) at 3.26 g cm−3.[69] However, its bulk density (DB) is only 1.73 g cm−3, representing a high porosity of 45.8% due to the considerable difference of DA and DB. As 25 mol% of Li2O was incorporated into CaMgSi2O6, the bulk and apparent density of Li0.25 were slightly decreased to 2.94 and 2.95 g cm−3, respectively. The slight decrease in the density could be ascribed to the formation of CaSiO3, Li2Si2O5, and SiO2 which have theoretical densities of 2.91,[70] 2.5,[71] and 2.26[72] g cm−3, respectively, all of which are lower than the theoretical density of CaMgSi2O6 at 3.26 g cm−3. Further, Li0.25 had an extremely low porosity of 0.3%, indicating a dense structure which is in good agreement with the observation in SEM (Figure 2b). As the Li2O content increased to 50 mol%, the porosity of 0.7% of Li0.50 suggested a dense structure similar to Li0.25, while the bulk and apparent density dropped further to 2.67 and 2.69 g cm−3, respectively. This could be attributed to the decreased content of CaMgSi2O6 and the increased content of CaSiO3, Li2Si2O5, and SiO2, coupled with two newly formed secondary phases Li2SiO3 and Li2Ca2Si5O13 with theoretical densities of 2.53[73] and 2.75 g cm−3,[74] respectively. When 75 mol% Li2O was added, the apparent density of Li0.75 remained around 2.7 g cm−3, but the bulk density decreased to 2.14 g cm−3 with a porosity of 21.3%, indicating an increase in the connected open pores. When the Li2O substitution reached 100 mol%, the apparent density of Li1.00 remained at 2.7 g cm−3, but the bulk density significantly dropped to 1.59 g cm−3, leading to an increase in porosity to 41.1%, which is comparable to the high porosity of pure diopside. These phenomena revealed that the substitution of Li2O induced the formation of secondary phases with lower density (CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13), resulting in the reduction of the overall density. However, the addition of small amounts of Li2O decreased the porosity of CaMgSi2O6, leading to a fused and dense structure due to the presence of Li2Si2O5. As the amount of Li2O increased, the grain size and porosity increased owing to the massive formation of CaSiO3, following the same trend of structural changes observed in SEM (Figure 2b–e).5TableBulk density, apparent density, and porosity of sintered pure diopside and CaMgSi2O6–Li2O bioceramics pelletsSampleBulk density [DB, g cm−3]Apparent density [DA, g cm−3]Porositya) [%]pure1.733.1945.8Li0.252.942.950.3Li0.502.672.690.7Li0.752.142.7221.3Li1.001.592.7041.1a)The porosity of tabulated samples is contributed by the open pores.Hardness MeasurementHardness results of sintered pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics before and after immersing in SBF solution are shown in Figure 6. Each sample was indented at three different locations, and then the averages and standard deviations were calculated. It was expected that the addition of Li2O would affect the hardness property through different combinations of phases (CaMgSi2O6, CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13) and structural change in the material.6FigureHardness results of sintered pure diopside and CaMgSi2O6–Li2O bioceramic: a) before and b) after SBF soaking for 28 days.As can be seen from Figure 6a, Li0.25 had the highest hardness of all the samples at about 700 Hv, which is about 38 times harder than the pure diopside. This result shows that the hardness of CaMgSi2O6–Li2O bioceramics decreased as the amount of Li2O increased, but the hardness was still generally much higher than that of the pure diopside. This indicates that the Li2O addition can significantly increase the mechanical strength of CaMgSi2O6. Based on Table 5 and Figure 6a, it was found that the hardness was directly proportional to the apparent density and inversely proportional to the porosity. The pure diopside had the highest apparent density, but the high porosity of 45.8% yielded a fragile structure and a low hardness. With the incorporation of Li2O, Li0.25 had a similar porosity but a higher apparent density compared to Li0.50, leading to a superior hardness of Li0.25 to that of Li0.50. As the increment of Li2O from 50 to 75 and 100 mol%, the apparent densities of Li0.50, Li0.75, and L1.00 remained at ≈2.7 g cm−3, while the porosity was increased from 0.7% to 21.3% and 41.1% with a remarkable drop in hardness from 521 down to 416 and 318 Hv. Based on the XRD results (Figure 1), the secondary phases (CaSiO3 and Li2Si2O5) occurred in Li0.25, suggesting that the highest hardness measured in Li0.25 was associated to the synergistic effect of CaMgSi2O6, CaSiO3, and Li2Si2O5. In addition, an amorphous phase of 4.46% was able to increase the bulk density of Li0.25 and reduce its porosity substantially. As elucidated in SEM and density measurement, the Li2O substitution varied the porosity and the density of Li0.25, inducing an aggregated and densified crystal structure as the reason of its outstanding hardness. With further increasing the Li2O quantity, though the amorphous phase increased to 8.37–15.44%, the Mg deficiency and the abundant CaSiO3 and Li2Si2O5 phase triggered progressive aggregation and grain growth, followed by a lowered density and a porous morphology, as shown in Figure 2c–e, resulting in the degraded hardness. Thus, it is expected that the hardness of CaMgSi2O6–Li2O bioceramics can be adjusted through the amount of Li2O substitution, depending on various implant applications. Moreover, all the samples were subjected to a decrease in hardness after SBF soaking for 28 days (shown in Figure 6b), which was due to the biodegradation behavior during soaking process. Among all the soaked samples, Li0.25 retained the highest hardness of 197 Hv, which is much higher than other samples’ hardness ranging from 4 to 12 Hv, revealing a strong resistance to biodegradation. In comparison with dense (>93%) TCP bioceramics synthesized by novel ultrafast high‐temperature sintering, the CaMgSi2O6–Li2O specimens have much higher hardness than dense TCP (maximum: 350 Hv).[75] In addition, compared to the most commonly used Ti‐based alloys for bone implant applications, the hardness of as‐sintered Li0.25 at 700 Hv is 1.7 times that of Ti–6Al‐4 V alloys at 416 Hv, while the hardness of soaked Li0.25 at 197 Hv is comparable to that of Ti–25Nb–3Zr–3Mo–2Sn alloys at 202 Hv.[76] The decreased hardness as a function of soaking time may allow Li0.25 to be more favorable for bone repair than Ti‐based alloys as the cortical bone has a relatively low hardness around 40–60 Hv.[77] The deteriorated hardness of Li0.50, Li0.75, and Li1.00 after SBF soaking could be ascribed to the escalating CaSiO3 formation as CaSiO3 has been reported to have a higher dissolution rate than diopside.[34,35]Dissolution and Precipitation Behaviors in SBF SolutionThe dissolution and precipitation behaviors of pure diopside and CaMgSi2O6–Li2O bioceramics were evaluated by the weight loss of the samples as well as the variations of the pH values and ion concentrations of the SBF solution after soaking in SBF for certain periods. To detect the degradability of pure diopside and CaMgSi2O6–Li2O bioceramics, the variations of sample weight during the immersion in SBF solution were measured and shown in Figure 7. It is clear that pure diopside and Li0.25 had smaller percentage of weight loss between 0% and 2%, while bioceramics with the higher concentrations of Li2O substitution had higher percentage of weight loss around 4% after 28 days of immersion time.7FigureWeight loss of pure diopside and CaMgSi2O6–Li2O bioceramics after soaking in the SBF solution as a function of immersion period.The mass variation is related to the biodegradation of the sample (weight loss) and the HAp formation on the sample surface (weight gain). From the results of overall weight loss, it can be suggested that the biodegradation of all the samples was predominant over HAp formation (especially at higher concentrations of Li2O) as the sample weights were reduced. For bone tissue repair applications, bioceramics with suitable biological degradation are needed. Thus, Li0.25 sample was considered to be favorable for repairing bone defects due to its slowest degradation rate, which could be attributed to the compact structure, as observed in Figure 2b. The samples with Li2O addition more than 25 mol% revealed the accelerated biodegradation properties in accordance with the decreased hardness (Figure 6b), which could be resulted from the growing secondary phase of CaSiO3 with higher dissolution rate. A similar result was obtained by Du et al.[78] that the weight loss of pure CaSiO3 was more severe than that of Mg‐incorporated CaSiO3, which could be due to the stronger MgO bonding than the CaO bonding.[37] Compared to another CaSi‐based bioceramic: baghdadite (Ca3ZrSi2O9), where the network of ZrCa bonds stabilizes the crystal structure and enables baghdadite a 70% higher compressive strength than CaMgSi2O6, the biodegradation rates of CaMgSi2O6–Li2O bioceramics in the current work were much lower than that of baghdadite (weight loss: 4–10%).[79]The pH values in the SBF solution following immersion of bioceramics samples are shown in Figure 8a. The upward trend of pH values for all the samples during 28 days of soaking indicates a sustained ion exchange of Ca2+, Mg2+, and Li+ with H+ and H3O+ in the SBF solution, followed by the formation of amorphous silica‐rich colloid layer and the subsequent formation of HAp.[49,80] In the first 7 days, all the CaMgSi2O6–Li2O bioceramics, especially those with higher levels of Li2O addition, had significantly higher pH values than pure diopside. After 1 week of soaking, the rapidly rising trends in pH value for all the samples tended to level off, indicating that the ion exchange between the samples and the SBF solution gradually reached saturation. After 28 days of soaking, Li0.25 had a similar pH value to pure diopside at around 8.0, while the pH values of Li0.50, Li0.75, and Li1.00 were about 8.3, implying that more ions were exchanged with higher Li2O substitution. As mentioned above, the ion exchange between the cations of bioceramics (Ca2+, Mg2+, and Li+) and the cations of SBF (H+ and H3O+) is a prerequisite of superficial HAp formation. In other words, the higher pH values stood for the more frequent ion exchange and better silica‐rich layer formation for HAp deposition. This observation is consistent with XRD and SEM results that the increase of Li2O content would stimulate the generation of CaSiO3 and thus improve the HAp formation. Du et al.[78] obtained an analogous outcome that the pH value of soaking medium provided by pure CaSiO3 was higher than Mg‐incorporated CaSiO3. Porosity may have negligible effect on the pH value as the comparably high pH values were given by Li0.50, Li0.75, and Li1.00 with different porosities. Since a weak alkaline microenvironment would improve osteoblast production and hinder osteoclast differentiation,[81] the Li0.25 sample is considered to be more suitable for implant applications because of its outstanding hardness (Figure 6), low degradation rate (Figure 7), reasonable HAp formation, and minor influence on the environmental pH value.8Figurea) pH values and ion concentrations of b) Ca2+, c) Mg2+, d) Li+, e) Si4+, and f) P5+ of SBF solutions after soaking pure diopside and CaMgSi2O6–Li2O bioceramics as a function of immersion period.The changes of ion concentration in SBF solution during immersion are shown in Figure 8b–f. In this test, the ion concentration of SBF was dominated by three reactions: 1) ion exchange to form silica‐rich layer, followed by an increase in pH value, and an increase in the concentration of Ca2+, Mg2+, and Li+; 2) adsorption of ions by the silica‐rich layer to form HAp, followed by a decrease in the concentration of ions, especially in Ca2+ and P5+; 3) dissolution of bioceramics and HAp, followed by an increase in the concentration of each ion. Thus, it can be summarized that ion exchange and material dissolution led to an increase in ion concentration, while the formation of HAp caused a decrease in ion concentration.First, the interaction mechanism between pure diopside and SBF solution is discussed as the followings. Pure diopside mainly exchanged Ca2+ and Mg2+ for H+ and H3O+ in SBF, forming a silica‐rich layer on the surface. As can be seen in Figure 8c, the Mg2+ concentration of pure diopside kept increasing, representing a greater driving force for ion exchange of Mg2+ and material dissolution than for adsorption. In Figure 8b, the Ca2+ concentration of pure diopside at day 14 was slightly lower than that of fresh SBF, indicating that more Ca2+ ions were adsorbed by the silica‐rich layer than they were released, and HAp was being formed at this stage. By day 28, the Ca2+ concentration has risen, suggesting that the adsorption of Ca2+ has reached saturation; then the ion exchange and material dissolution have outpaced the adsorption. The change of P5+ concentration in Figure 8f also revealed that P5+ had not yet been fully adsorbed at day 14, confirming that HAp was still being formed by Ca2+ and P5+ at this stage. By day 28, P5+ ions were completely consumed, indicating the saturation of ion adsorption and the completion of HAp formation.Second, the interaction mechanism of Li0.25 is discussed. As can be seen from Figure 8b, the Ca2+ concentration of Li0.25 was slightly higher than that of fresh SBF at day 14 and then remained stable until day 28, suggesting a two‐way balance between the release and adsorption of Ca2+, with a continuous HAp formation. Based on Figure 8f, P5+ ions have not been completely consumed after 28 days, verifying an ongoing ion adsorption and HAp formation throughout the immersion. Figure 8c shows that the Mg2+ concentration of Li0.25 was slightly lower than that of fresh SBF at both day 14 and 28, indicating that the adsorption rate of Mg2+ was higher than the release rate. In Figure 8d, Li0.25 displayed a steady release of Li+ over 28 days of immersion.Li0.50, Li0.75, and Li1.00 were grouped into the same category for discussion. As can be seen from Figure 8b, the Ca2+ concentration of these three samples increased continuously from ≈100 to ≈265 ppm throughout the 28‐day immersion period, representing that the ion exchange and material dissolution were much more driven than ion adsorption and thus HAp growth had been completed. Figure 8f shows that the P5+ ions were fully consumed by day 14, further confirming a complete formation of HAp. Based on Figure 8c, these three samples were identical to Li0.25 in that the adsorption rate of Mg2+ was slightly higher than the release rate. Figure 8d illustrates that the Li+ concentration increased with the increase of soaking period and the amount of Li2O.Figure 8e shows the Si4+ concentrations of all samples, and it can be seen that all the samples continued to release Si4+. Huang et al.[80] stated that the dissolution of silicon ions was dependent on not only the chemical composition, but the crystal structure of silicates, especially the bonding types. Since the Si content of materials remained constant here, the concentration of Si4+ in SBF solution should be predominantly related to its incorporation in crystalline or amorphous structures. Among the structures present in our materials, CaMgSi2O6, CaSiO3, and minor phase Li2SiO3 are inosilicates, where two oxygen atoms of each silicate tetrahedron are shared with adjacent tetrahedra to form single chains of connected silicate tetrahedra. The Li2Si2O5 phase induced by the substitution of Li2O is a phyllosilicate, where sheet‐like silicates composed of six‐membered silica rings occur in a chair structure.[82] The SiO2 phase is a tectosilicate, where silicate tetrahedra are connected in the form of 3D framework. In addition, the minor secondary phase Li2Ca2Si5O13 is a complex silicate, containing four‐ and five‐membered rings of silicate tetrahedra to form double chains.[74] Therefore, it can be deduced that it is more difficult for Si4+ to be dissolved from Li2Si2O5, SiO2, and Li2Ca2Si5O13 than from CaMgSi2O6, CaSiO3, and Li2SiO3 because of the stronger SiO covalent bonds. This may have resulted in the Si4+ concentrations of CaMgSi2O6–Li2O bioceramics becoming lower than that of pure diopside at the end of immersion period. Further, the concentration of Si4+ in Li0.25 was significantly lower than in the other samples, which could be due to its dense structure.The change in the concentration of Mg2+ ions in the SBF solution is particularly worthy of closer scrutiny. As shown in Figure 8c, the amount of Mg2+ released by pure diopside throughout the soaking period was much higher than that released by CaMgSi2O6–Li2O bioceramics. The elemental analysis of precipitated HAp in Table 4 shows that the HAp formed on pure diopside contained 8–13 times more Mg2+ ions than the HAp formed on CaMgSi2O6–Li2O bioceramics, indicating that pure diopside adsorbed considerable Mg2+ from the surrounding environment when forming HAp. In contrast to pure diopside, mixed phase CaMgSi2O6–Li2O bioceramics adsorbed only a trace amount of Mg2+, corresponding to the slight decrease of Mg2+ concentration shown in Figure 8c. Moreover, after 28 days of immersion, Mg2+ ions remained undetected after the HAp had fully grown on the surface of CaMgSi2O6–Li2O bioceramics. The low levels of Mg2+ detected in both SBF and HAp of CaMgSi2O6–Li2O bioceramics implied that Mg2+ ions were neither released from the materials nor then readsorbed by the HAp. Thus, it can be inferred that the addition of Li2O suppressed the release of Mg2+, thereby reducing the adsorption of Mg2+ by HAp, which can be further confirmed by the elemental mapping in Figure 5, where high Mg signals were detected in the matrix of CaMgSi2O6–Li2O bioceramics instead of in the HAp. The reason for the restricted release of Mg2+ may be explained as the followings. As clarified in the elemental mapping in Figure 3, the amorphous phase of Li2Si2O5 started to form after the substitution of Li2O, and the viscosity increased in the region where Mg2+ ions were present, resulting in a decreased ion diffusion. Hence, Mg2+ began to react with the surrounding Ca2+, Si4+, and O2− to form large grains of CaMgSi2O6, which were more difficult to dissolve in SBF and release ions than the small grains of CaMgSi2O6. In addition, the embedment of CaMgSi2O6 in the amorphous phase of Li2Si2O5 further reduced the dissolution rate, leading to Mg2+ being trapped in the large grains of CaMgSi2O6.As an indicator of silica‐rich layer formation, the relationship between the pH value and ion concentration can be explored in depth here. According to Figure 8b–d, though the ion exchange of Mg2+ was hindered by the addition of Li2O, a large amount of Ca2+ and Li+ of CaMgSi2O6–Li2O bioceramics was exchanged for H+ and H3O+, leading to much higher pH values (≈8.3) than that of pure diopside (≈8.1). The high pH values of CaMgSi2O6–Li2O bioceramics demonstrated the formation of a large quantity of silica‐rich layer, laying the foundation for the following substantial HAp formation. However, the ion release of Ca2+ reached about 160 ppm, which was much higher than that of Li+ (11 to 33 ppm), coupled with the fact that Li+ is a monovalent ion and can exchange half as many H+ ions as Ca2+/Mg2+, indicating that Ca2+ was more dominant than Li+ in the formation of silica‐rich layer. It is believed that the high Ca2+ and Li+ concentrations were contributed by CaSiO3 and Li2Si2O5 phases, which were more likely released from the vulnerable amorphous matrix than from the solid crystals, as shown in Figure 3. Since Li0.25 had a dense structure and relatively small amount of CaSiO3, its overall ion exchange and material dissolution were lower than other CaMgSi2O6–Li2O bioceramics. Moreover, the pH value of Li0.25 was close to that of pure diopside, suggesting a comparable overall ion exchange and similar ability of silica‐rich layer formation.From the results of XRD (Figure 1c) and SEM (Figure 4) analyses, it can be observed that all the CaMgSi2O6–Li2O bioceramics were found to have a greater extent of HAp formation than pure diopside. It can be further confirmed that Li0.50, Li0.75, and Li1.00 had superior silica‐rich layer and subsequent HAp formation to pure diopside based on the variations of pH value (Figure 8a) and the consumption of P5+ ions (Figure 8f). In addition, the high Ca2+ levels found in the SBF of Li0.50, Li0.75, and Li1.00 may increase the supersaturating concentration of SBF, which would facilitate the apatite formation on the surface of the bioceramics.[83] While Li0.25 had the similar capacity of silica‐rich layer formation as pure diopside, the high remaining P5+ ions in SBF implied that Li0.25 should have had the worse HAp formation than pure diopside as fewer P5+ ions were adsorbed to form HAp, which is inconsistent with the results shown in XRD and SEM. This inconsistency could be due to the fact that the Mg2+ and Si4+ released from pure diopside into SBF have been readsorbed by the HAp formed on the surface of pure diopside and hence reduced the crystallization of HAp. Hafezi‐Ardakani et al.[84] have stated that Mg2+ ions may be able to enter the forming HAp nuclei and further inhibited the HAp evolution as Mg2+ ions are not readily accommodated in the HAp structure. In the present study, all the CaMgSi2O6–Li2O bioceramics blocked the release of Mg2+, and only trace amounts of Mg2+ were adsorbed by HAp, as shown in Table 4, which significantly improved the HAp growth. Therefore, Li0.25 eliminated the negative effect of Mg2+ on HAp growth by means of low‐level Li2O addition, allowing for much more efficient HAp formation with a lower P5+ consumption than pure diopside. It can be expected that extending the soaking time for Li0.25 would allow the complete utilization of P5+ ions and the further increase in the amount of HAp precipitate. Moreover, Li0.25 had the lowest overall ion release among all the samples due to its dense structure, which is in good agreement with the result of weight loss.Cell Viability3‐(4,5‐dimethylthiazol‐2‐yl)‐2,5‐diphenyl tetrazolium bromide (MTT) assay is an effective method to assess the cytotoxicity of materials by measuring the concentration of mitochondrial dehydrogenase, which shows cell proliferation and hence reflects the biocompatibility of materials. In the current work, cell viability was studied with pure diopside and CaMgSi2O6–Li2O bioceramics samples by MTT test. Their optical densities (24 h) with incubated D1 cells were compared to the positive control (with sodium dodecyl sulfate [SDS] addition) and negative controls (with only cells and with Al2O3 addition), as shown in Figure 9. From the result, it can be seen that cells were more viable in the negative controls and pure diopside than in the positive control and CaMgSi2O6–Li2O bioceramics samples. The cell viability percentage ranging from 60.7% to 70.5% with CaMgSi2O6–Li2O bioceramics samples suggested that the presence of high‐dose Li ions had a strong effect on the proliferation of mouse bone marrow cells (D1). Among all the Li2O‐incorporated samples, Li0.25 possessed a relatively acceptable cell viability at 70.5%, indicating that low doses of Li ion may be considered to minimize the cytotoxicity of CaMgSi2O6–Li2O bioceramics.9FigureCell viability of D1 cells with pure diopside and CaMgSi2O6–Li2O bioceramics.ConclusionPure diopside and CaMgSi2O6–Li2O bioceramics were prepared by precipitation method, followed by sintering at 1000 °C for 4 h. Alongside the primary CaMgSi2O6 phase, secondary phases of CaSiO3, Li2Si2O5, SiO2, Li2SiO3, and Li2Ca2Si5O13 formed with the increasing replacement of MgO by Li2O in the material stoichiometry, with implications for the mechanical strength and bioactivity of the material. Based on SEM images and elemental analysis, low‐level Li2O substitution induced a densified microstructure consisting of CaMgSi2O6 and CaSiO3 grains embedded in Li2Si2O5 amorphous matrix, leading to a significant enhancement of hardness. However, the hardness decreased for higher Li2O levels due to the porosification induced by the formation of CaSiO3 and significant grain growth. The elemental redistribution was observed in CaMgSi2O6–Li2O bioceramics where Mg was mostly concentrated in the CaMgSi2O6 grains, while Ca was found both in the crystalline and amorphous phases. After 28 days of SBF soaking, the mass and the hardness of all the samples decreased with the increment of Li2O due to abundant CaSiO3 formation with relatively weak CaO bonding and high dissolution rate. Furthermore, highly improved HAp (Ca5(PO4)3OH) formation was found after SBF soaking on all the CaMgSi2O6–Li2O bioceramics compared to pure diopside, which can be ascribed to the limited release of Mg2+ from the large CaMgSi2O6 grains embedded in the Li2Si2O5 amorphous phase, which prevented Mg2+ entering HAp nuclei and thus promoted HAp crystallization. In this study, Li0.25 is more favorable for implant applications due to its excellent hardness (as‐sintered: 700 Hv; SBF‐soaked: 197 Hv), low degradation rate (≈0.6 wt%), and minor impact on the environmental pH value. In addition, Li0.25 eliminated the negative effect of Mg2+ on HAp growth using low‐level Li2O substitution, permitting much more efficient HAp formation with a lower P5+ consumption than pure diopside and hence resulting in a superior bioactivity. A more detailed composition adjustment (such as low concentration of ion substitution) can be considered in the future work to optimize the biocompatibility of CaMgSi2O6–Li2O bioceramics as the presence of high‐dose Li+ suppressed the proliferation of mouse bone marrow cells (D1).Experimental SectionSamples PreparationDiopside and mixed phase CaMgSi2O6–Li2O bioceramics were synthesized by a precipitation method based on a procedure described in detail elsewhere.[48,49,85] For pure diopside, 0.01 mol of calcium chloride (CaCl2, 95%, Fisher Scientific, UK), 0.01 mol of magnesium chloride hexahydrate (MgCl2·6H2O, 95%, EMSURE, Germany), and 0.02 mol of tetraethyl orthosilicate (TEOS, 98%, Seedchem, Australia) as reactants were dissolved in 200 mL of ethanol (C2H5OH, 95%) to produce 0.01 mol of CaMgSi2O6 stoichiometry solution. The solution was magnetically stirred at 80 °C for 2 h. Then ammonium hydroxide (NH4OH, 28%, Nihon shiyaku reagent, Japan) was added into the solution to produce white precipitates. After stirring overnight, the solvent had evaporated, and the remaining white precipitates were dried in an oven at 80 °C for 24 h. The dried precipitates were ground into fine powder, and then calcined at 600 °C for 3 h to remove impurities. Compact pellets (Ø13 × H1.8 mm) were prepared by compressing the calcined CaMgSi2O6 powder under 75 MPa for 30 s. Then the white compact pellets were sintered at 1000 °C for 4 h in the air to obtain the final CaMgSi2O6 products.To synthesize the mixed phase CaMgSi2O6–Li2O bioceramics, a similar procedure as for pure diopside was carried out with the variation that a wide range of Li2O amounts (25, 50, 75, and 100 mol%) were introduced by the corresponding substitution of lithium chloride (99%, Acros Organics, Germany) for magnesium chloride (MgCl2). The CaMgSi2O6–Li2O bioceramics with various Li2O amounts (0–100 mol%) were denoted as pure, Li0.25, Li0.50, Li0.75, and Li1.00.Materials CharacterizationThe mineralogy of the pure diopside and mixed phase bioceramics was analyzed by a powder X‐ray diffractometer (D2 Phaser, Bruker, USA) with Cu Kα radiation (λ = 1.54184 Å) at 30 kV and 10 mA with diffraction angles (2θ): 10–70°, step size: 0.01° 2θ and step speed: 0.5 s/step. The phase identification and degree‐of‐crystallinity analysis based on the diffraction patterns were processed by Match! software (Crystal Impact, Germany). The surface morphology and chemical composition of the samples were observed using a FE‐EPMA (JXA‐8530F, JEOL, Japan) which consisted of a SEM and a wavelength dispersive spectrometer. FE‐EPMA was operated at 15 kV with an analytical depth of ≈700 nm, where mapping mode was used to analyze the elemental distribution and point mode was used to precisely determine the elemental concentration. The hardness was measured using a micro Vickers hardness tester (FM‐110, Future‐Tech, Japan) with a force of 1 kg over a loading time of 10 s under room temperature.The bulk density and apparent density of the sintered pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics were measured using Archimedes technique in deionized water. The density measurement was performed as the following. First, the samples were dried in an oven at 110 °C for 24 h to remove free water, then naturally cooled down to room temperature in an electronic moisture‐proof box at a constant zero humidity for 24 h. The weight of the dried samples (denoted as m1) was measured by a precision analytical balance (AS 220.R2 PLUS, Radwag, Poland) with 1 mg of accuracy. Second, the samples were saturated with deionized water under a vacuum condition for 2 h. The saturated samples were then suspended in deionized water and the immersed mass of the samples (denoted as m2) was recorded. Third, the samples were taken out from deionized water and the remaining water on the sample surface was carefully removed. The weight of the saturated samples in air was measured and denoted as m3. According to Archimedes principle, bulk density was the ratio of the material mass without free water to the macroscopic material volume, which was occupied by solid material, open pores, and closed pores. The calculation of bulk density was expressed by the following equation.1Bulk density(DB)=(m1×ρ0)/(m3−m2)\[\begin{array}{*{20}{c}}{{\rm{Bulk}}\;{\rm{density}}\left( {{D_{\rm{B}}}} \right) = \left( {m1 \times {\rho _0}} \right)/\left( {m3 - m2} \right)}\end{array}\]Apparent density was the ratio of the material mass to the material volume including solid material and closed pores, which was calculated as the equation described below.2Apparent density(DA)=(m1 × ρ0)/(m1−m2)\[\begin{array}{*{20}{c}}{{\rm{Apparent}}\;{\rm{density}}\left( {{D_{\rm{A}}}} \right) = \left( {m1\: \times \:{\rho _0}} \right){\rm{/}}\left( {m1 - m2} \right)}\end{array}\]where ρ0 was the density of deionized water at 25 °C. Based on the bulk density and apparent density, the porosity given by the open pores was calculated as the following.3Porosity(%)=[1−(DB/DA)]× 100%\[\begin{array}{*{20}{c}}{{\rm{Porosity}}\left( \\end{array}\]In Vitro Biodegradation AnalysisIn vitro biodegradation tests were performed by soaking pure diopside and mixed phase CaMgSi2O6–Li2O bioceramics pellets in a SBF solution at 37 °C. The SBF solution had a similar ion concentration with human blood plasma, as proposed by Kokubo et al.[86] For the evaluation of degradation, each pellet was immersed in 40 mL SBF solution with pH 7.4 at 37 °C for 7, 14, 21, and 28 days. After the set soaking time, each pellet was washed by deionized water and then completely dried, and the final weight loss was measured and expressed in percentage (wt%) as below:4WL(%)=[(W0−WF)/W0]×100%\[\begin{array}{*{20}{c}}{{W_{\rm{L}}}\left( \\end{array}\]where WL was the weight loss of pellet after SBF soaking, W0 was the initial weight of pellet, and WF was the final weight of pellet after soaking.The pH values of the SBF solution before and after immersion were measured by a pH meter (pH 510, Eutech, Singapore) and the ion concentrations of Ca2+, Mg2+, Li+, Si4+, and P5+ in the SBF solution were analyzed using an inductively coupled plasma‐mass spectrometer (ELEMENT XR, Thermo fisher scientific, USA). The initial concentrations of Ca2+, Mg2+, Li+, Si4+, and P5+ ions in fresh SBF were 92.6, 43.9, 0, 0, and 40 ppm, respectively.Cell Viability AssayThe cell viability was evaluated using the liquid extracts of bioceramic samples and MTT (98%, Alfa Aesar, USA) assay. Dulbecco's modified Eagle's medium (DMEM, Gibco, USA) containing 10% fetal bovine serum (Sigma‐Aldrich, USA) and 0.5% penicillin–streptomycin (Sigma‐Aldrich, USA) was used in the cell culture and the extraction of sample substances. MTT solution was prepared at the concentration of 5 mg mL−1 by dissolving MTT in sterile and filtered phosphate buffered saline (PBS). The samples for preparing liquid extracts were pure diopside and mixed phase bioceramics with four levels of Li2O substitution (25–100 mol%). A group with the addition of SDS (99%, J.T Baker, USA) was prepared as positive control. A group containing only cells and a group with the addition of aluminum oxide (Al2O3, 99.5%, Sigma‐Aldrich, USA) were prepared as negative controls.First of all, mouse bone marrow stromal D1 cells were seeded at a density of 2 × 105 cells/well in 96‐well plates and cultured in 100 µL of DMEM at 37 °C in an incubator under 5% CO2 atmosphere for 24 h. Second, liquid extracts of all the test groups were prepared in DMEM at a concentration of 0.1 g mL−1 (except SDS group at 0.2 mg mL−1) at 37 °C for 24 h, followed by a filtration of extracts with a 0.25 µm filter. After 24 h of incubation, the initial DMEM was removed and then 100 µL of the filtered extract was added to each well. The D1 cells were cultured in the extracts in the incubator for 24 h for the following cytotoxicity test. After 24 h of cell culturing in the presence of extracts, the osteoblast cell viability was evaluated by MTT assay. The extracts were removed and 100 µL of MTT solution (5 mg mL−1 in PBS) was added to the cultures in each well, and the cells were incubated at 37 °C in 5% CO2 for 2 h. Then the MTT solution was removed and 100 µL of dimethyl sulfoxide (99.5%, Sigma‐Aldrich, USA) was added to each well to dissolve purple crystals of formazan. The optical density of each well at 570 nm of wavelength was measured by an ELISA microplate reader (Model 550, Bio‐Rad, USA). The cell viability was calculated using:5Cell viability(%)=(ODs/ODc) × 100%\[\begin{array}{*{20}{c}}{{\rm{Cell}}\;{\rm{viability}}\left( \\end{array}\]where ODs was the optical density values of sample groups and ODc was the optical density value of control group containing only cells. The quantitative data were expressed as means ± standard deviation of six repeats per experiment.AcknowledgementsThe authors acknowledge financial support from the Ministry of Science and Technology of Taiwan under contract MOST 111‐2636‐E‐110‐008, the DAAD's Program for Project‐related Personal Exchange (PPP) (Project ID 57390207), the Field Emission Electron Probe Microanalyzer equipment provided by Joint Center for High Valued Instruments at National Sun Yat‐Sen University, and the Inductively Coupled Plasma‐Mass Spectrometer equipment provided by Core Facility Center of National Cheng Kung University, Taiwan. The authors thank National Center for High‐performance Computing (NCHC) for providing computational and storage resources.Conflict of InterestThe authors declare no conflict of interest.Data Availability StatementThe data that support the findings of this study are available from the corresponding author upon reasonable request.S. Kargozar, M. Montazerian, E. Fiume, F. 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Journal

Advanced Materials Interfaces – Wiley

Published: Apr 1, 2023

Keywords: bioactivities; diopside; lithium oxides; mechanical properties

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