Magnesium-containing bioactive polycrystalline silicate-based ceramics and glass-ceramics for...

Current Opinion in Solid State and Materials Science xxx (2014) xxx–xxx

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Current Opinion in Solid State and Materials Science

journal homepage: www.elsevier .com/locate /cossms

Magnesium-containing bioactive polycrystalline silicate-based ceramicsand glass-ceramics for biomedical applications

http://dx.doi.org/10.1016/j.cossms.2014.02.0041359-0286/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding author. Tel.: +49 9131 8528601; fax: +49 9131 8528602.E-mail address: [email protected] (A.R. Boccaccini).

1 Present address: Department of Biomaterials, Radboud University Medical Center,Philips van Leijdenlaan 25, 6525 EX Nijmegen, The Netherlands.

2 Present address: Max-Planck-Institute for Dynamics of Complex TechnicalSystems, Sandtorstr. 1, 39106 Magdeburg, Germany.

Please cite this article in press as: Diba M et al. Magnesium-containing bioactive polycrystalline silicate-based ceramics and glass-ceramics for biomapplications. Curr. Opin. Solid State Mater. Sci. (2014), http://dx.doi.org/10.1016/j.cossms.2014.02.004

Mani Diba 1, Ourania-Menti Goudouri, Felipe Tapia 2, Aldo R. Boccaccini ⇑Institute of Biomaterials, Department of Materials Science and Engineering, University of Erlangen-Nuremberg, 91058 Erlangen, Germany

a r t i c l e i n f o

Article history:Received 18 November 2013Revised 2 February 2014Accepted 17 February 2014Available online xxxx

Keywords:MagnesiumBioactive materialsSilicateCeramicsGlass-ceramics

a b s t r a c t

With improvement of orthopaedic technologies for bone replacement and regeneration, there is anincreasing need for materials with superior properties. Mg-containing silicate ceramics and glass-ceramicshave been shown to be bioactive and exhibit various advantages for biomedical applications. Thisreview paper is intended to summarize and discuss the most relevant studies carried out in the field ofMg-containing bioactive silicate ceramics and glass-ceramics.

� 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Degenerative and inflammatory conditions of bone, teeth andjoints are typical causes of chronic disability, while the numberof people suffering from musculoskeletal conditions has increasedby 25% over the past decade [1]. These problems affect several hun-dred millions of people worldwide, and a sharp increase of this fig-ure is expected due to the predicted doubling in the number ofpeople over 50 years old by the year 2020 [2,3]. Affordable mea-sures to treat musculoskeletal conditions are continuouslyrequired, with both replacement and regeneration of damagedhard tissues being suitable therapeutic alternatives.

Ceramic materials feature prominently among established bio-materials for the replacement or regeneration of the musculoskel-etal system and in dental applications [4–6]. For example, ceramicshave been widely used in the repair of skull bone defects [7,8] andmaxillofacial reconstruction [9], including alveolar ridge augmen-tation [10], periodontal pocket obliteration [11] as well as dental[12,13] and orthopaedic implants [14,15]. Bioceramics can be clas-sified in bioinert, bioactive and bioresorbable materials, based ontheir chemical surface reactivity [16,17]. Bioactive materials arematerials able to form a chemical bond with living tissues. In the

context of bone replacement materials, the bioactivity of a materialis commonly characterized by its ability to induce formation of anapatite layer on its surface upon immersion in biological fluids[18]. However, it is important to mention that some materials(e.g. b-tricalcium phosphate) have been reported not to be ableto form a surface apatite layer but still be able to bond to bone tis-sue in vivo [19]. Alternatively, bioceramics can be classified bytheir primary chemical composition [20]. Silicate and phosphateceramics and glass-ceramics are the two broadest categories andinclude, basically, all the bioactive and bioresorbable compositions.Oxide ceramics such as alumina and zirconia are among the bioin-ert ceramics and are mainly used as prosthetic devices in medicineand dentistry.

It is well-known that the composition of bioceramics is one ofthe key parameters affecting their properties and determiningtheir biocompatibility, bioactivity and biodegradability [6]. Fur-thermore, it has been shown that release of specific elements (ions)from inorganic, bioactive materials, e.g. bioactive silicate glassesand calcium phosphates, can induce positive effects in their sur-rounding biological environment [21,22].

The relevance of magnesium for biomedical applications isrelated to its fundamental role in cellular processes and humanmetabolism [23,24]. Magnesium is one of the most important -mineral elements in the human body, with approximately half ofthe total physiological magnesium stored in bony tissues [25]. -Several studies have indicated that divalent cations (e.g. Mg2+)have a key role in bone remodelling and skeletal development[26,27]. It is well known that magnesium is a co-factor for many

edical

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2 M. Diba et al. / Current Opinion in Solid State and Materials Science xxx (2014) xxx–xxx

enzymes and stabilizes the structures of RNA and DNA [24,28]. Fur-thermore, research studies have consistently demonstrated thatsurface modification of bioceramics with Mg2+ substantially affectsthe phenotype of osteogenic cells in vivo and in vitro [29,30].Osteoblast cells-biomaterials surface interactions are mainly med-iated by cell membrane adhesion receptors (integrins), in whichmagnesium plays an important role [28,31]. Mg2+ can bond to inte-grins a subunits [32,33], and extracellular changes of Mg2+ canmodulate the affinity of the cells to the biomaterial surface [34–36].

Various studies have investigated magnesium-containing poly-crystalline ceramics and glass-ceramics for biomedical applica-tions; which is the subject of the present review article [23–25,37–39]. It has been shown that these materials are promisingcandidates for applications in bone tissue regeneration and magne-sium-based bone fillers have been approved by the FDA (Osteo-crete; Bone Solutions, Inc., Dallas, TX), which confirms thepotential of Mg to be used as a component of bone healingmaterials.

Furthermore, Mg-containing bioactive ceramics and glass-ceramics are attractive from the point of view of their mechanicalproperties, bioactivity and biocompatibility [40–43]. These materi-als have been shown to promote cell adhesion, proliferation,spreading, and differentiation [39,44,45], which make them prom-ising candidates for tissue engineering applications. In addition, arecent study [46] has indicated the ability of a Mg-containing sili-cate bioceramic to induce angiogenesis in biological conditions.Although there is an increasing interest in these materials, therehas not been a comprehensive overview discussing their propertiesand potential for biomedical purposes.

The widespread use of magnesium-containing inorganic mate-rials in biomedical applications is at the centre of numerousresearch efforts and it represents an emerging area of biomaterialsresearch. The biomedical applications of Mg-containing bioactive(amorphous) glasses were reviewed in our recent article [47].The present review will focus on polycrystalline ceramics andglass-ceramics. In addition, several studies have investigated Mg-containing non-silicate systems, including calcium phosphate[48–50] and magnesium phosphate [51,52] systems. However, thisreview paper is intended to discuss the most relevant studies car-ried out for Mg-containing bioactive silicate ceramics and glass-ceramics.

2. Mg-containing bioactive silicate ceramics

Various Mg-containing silicate ceramics have been shown to bebioactive and attractive for biomedical applications. Table 1 pre-sents an overview of basic characteristics of these ceramics. More-over, Table 2 summarizes the relevant in vitro and in vivobiological properties reported in the literature.

Table 1Overview of Mg-containing bioactive silicate ceramics.

Name Crystal system Chemical formula Co

Mg

Akermanite Tetragonal Ca2MgSi2O7 8Bredigite Orthorhombic Ca7Mg(SiO4)4 3Diopside Monoclinic CaMgSi2O6 11Forsterite Orthorhombic Mg2SiO4 34Merwinite Monoclinic Ca3Mg(SiO4)2 7Monticellite Orthorhombic CaMgSiO4 15Proto-enstatite Orthorhombic MgSiO3 24

a Upper stability limit (decomposes to Ca2SiO4 and merwinite).b Decomposes to forsterite.

Please cite this article in press as: Diba M et al. Magnesium-containing bioactivapplications. Curr. Opin. Solid State Mater. Sci. (2014), http://dx.doi.org/10.101

2.1. Akermanite

2.1.1. SynthesisNormally, the naturally occurring akermanite is associated with

other minerals and accordingly is not pure. Therefore, sol–gel[37,53–58] and combustion synthesis [59] methods have beenused in order to synthesize pure akermanite ceramics for biomed-ical applications. In 2004, Wu and Chang [53] reported the chemi-cal synthesis of pure akermanite powders via sol–gel method. Itwas shown that calcification temperature is an important factorfor obtaining pure akermanite ceramics, and merwinite and diop-side impurities could be observed at calcification temperatureslower than 1300 �C. The optimum calcination temperature forthe sol–gel synthesis of pure akermanite powders was reportedto be 1300 �C, which leads to formation of porous agglomeratedakermanite particles with particle size about 5–40 lm containingpores of about 1–5 lm in size [53]. In 2011, Bhatkar and Bhatkar[59] synthesized akermanite powders using a combustion synthe-sis method. They claimed that this technique is easier than previ-ously used sol–gel methods [53] and it produces akermanitepowders with a narrower size distribution and average agglomer-ate particle sizes in the range 0.5–5 lm.

2.1.2. Sintering and mechanical propertiesThe sintering process of akermanite ceramics highly determines

the mechanical properties of the material in the bulk form. Table 3shows density and mechanical properties values of akermanitedisks prepared by uniaxial pressing (10 MPa) of sol–gel akermanitepowders and subsequent sintering using different schedules [55].It was reported that akermanite could not be obtained at 1400 �C[55].

It is important to consider that the mechanical properties ofceramics can be altered by producing porous structures. Thus,Wu et al. [57] showed an indirect relationship between the poros-ity and mechanical strength of akermanite porous structures. Itwas shown that with increase of porosity from 63.5% to 90.3%,the mechanical strength decreased from 1130 kPa to 530 kPa [57].

2.1.3. In vitro bioactivity and degradationTo the best of the authors’ knowledge, diopside ceramics were

the first type of Mg-containing silicate ceramics, which wasreported to be bioactive [60]. This finding motivated manyresearchers to study bioactivity and bioactivity related applica-tions of other Mg-containing silicate ceramics [53,61]. Variousstudies have evaluated the in vitro bioactivity and degradation ofakermanite ceramics [53–57]. Wu and Chang [53,55] revealedthe apatite-forming ability of sol–gel derived akermanite powdersand disks by immersion of the samples in simulated body fluid(SBF) solution. Due to the similar composition and differentstructure of akermanite compared to bioactive diopside [62], it

mposition (wt.%) Melting T (�C)

Ca Si O

.92 29.40 20.6 41.08 1454

.61 41.67 16.69 38.03 1372 ± 2a

.22 18.51 25.94 44.33 1391

.55 0 19.96 45.49 1890

.39 36.58 17.09 38.94 1450

.53 25.61 17.95 40.9 1454

.31 0 28.09 15.00 1557b

e polycrystalline silicate-based ceramics and glass-ceramics for biomedical6/j.cossms.2014.02.004


Table 2Summary of in vitro and in vivo biological studies of Mg-containing bioactive silicate ceramics.

Name In vitro In vivo

Cell type analyzeda Direct evaluationb results Indirect evaluationb results Animal model Results

Akermanite Adipose SC, MSCs, primaryosteoblasts, osteoblast CL,fibroblast CL, stromal cells,ligament cells, endothelialcells, and bacteria as E. coli

Akermanite samples enhancedproliferation and differentiation ofadipose SC and hPDLCs; goodspreading and proliferation ofosteoblastsc; hBMSCa showedgood spreading, differentmorphologies and enhancedproliferation, while ALPa activitywas twice higherc

Extracts enhanced proliferationof hBMSCa, primary osteoblastsand fibroblast cell line L929;enhanced ALPa activity, andenhanced expression ofosteocalcin, osteopontin andbone sialoprotein fromhBMSCsc; extracts enhancedosteoblastic differentiation ofhBMSCd. Osteogenesis andangiogenesis of hAECs wasstimulated. Also, it is proposedthat Ca, Mg, and Si present inakermanite extracts has apossitive effect ondifferentiation of hASCs viaextracellular signal-releasedkinases signalling pathway

New Zealandrabbits

It was observed a highermineral apposition rate of newbone formation in akermanitescaffoldsc that enhanced boneregeneration. Also,angiogenesis was faster inakermanite bioceramicimplants compared to thecontrolc

Bredigite Primary culture of PDLCs;hBMSCs; hAECs; fibrobastcell line L929; osteoblasts

Osteoblasts showed goodattachment and proliferation overbredigite surface after 6 h ofculture

Cementogenic proliferation anddifferentiation of primaryculture of PDLCsa werestimulated, possibly byactivating Wnt/b-cateninsignalling pathway; osteogenicdifferentiation of hBMSCsa andangiogenesis of hAECs weresuperior to those observed withakermanite and diopside;fibroblast L929 proliferationwas enhanced with extractconcentrations between 6.25and 25 mg/mL

There are noin vivo studiesfor bredigite

Diopside Ostegenic CL MC3T3-E1;fibroblast CL L-929; hAECs

Diopside discs allowed ALPactivity, differentiation andgrowth with mineralcrystallization of MC3T3-E1 cells,and allowed adhesion of thefibroblast CL L-929

Angiogenesis of hAECsa

cultured on ECMatrix™ wasstimulated with diopsideextractsc

White Japaneserabbits;Japanesemonkey;wistar rats

Diopside attached to rabbitbone, it was degraded andreplaced by new formed bone.Diopside implanted in the jawbone of rabbit and in themandible of a Japanese monkeyshowed uniform junction withnewly grown bone, withpresence of apatite crystals.Finally, in vivo osteogenesiswas reported with diopsidemicrospheres as bone fillingmaterial

Forsterite Osteoblast-like G292 CL;U20S-type humanosteoblast cells; primaryculture of osteoblasts

G292 cells and primary culture ofosteoblasts attached faster andspread over forsterite discs. Cellsseeded on composites of b-CaSiO3/forsterite proliferate faster thanon pure b-CaSiO3

Proliferation of G292 cells wasenhanced compared to mediumfree of extract; ionic productsstimulated U20S-type humanosteoblast cells proliferation

There are nostudiesexploring thein vivoperformance offorsteriteceramics

Merwinite Osteoblasts Osteoblasts can adhere and spreadover merwinite surface

Merwinite extracts promotedosteoblast cell growth

Rat femoraldefects

Osteogenesis in vivo is inducedand merwinite implantsshowed a higherbiodegradation at early and latestages of implantationcompared to HA bioceramics

Monticellite Osteoblasts Osteoblasts can adhere and spreadover its surface

Monticellite extracts promotedosteoblast cell growth

There are noin vivo studiesfor Monticellite

a Abbreviations: SC: stem cells; CL: cell line; MSCs: mesenchymal stem cells; ALP: alkaline phosphatase; hBMSC: human bone marrow stromal cells; hBMSCs: human bonemarrow MSCs; hPDLCs: human periodontal ligament cells; hAECs: human aortic endothelial cells; PDLCs: periodontal ligament cells.

b Direct evaluation consists on analysing cell performance by directly culturing cells over a sample, while indirect evaluation consists on exposing cells to extracts of thematerial at different dilutions.

c Beta tricalcium phosphate was used as control material.d Control consisted of akermanite extract-containing osteogenic reagents (L-ascorbic acid, glycerophosphate, and dexamethasone).

M. Diba et al. / Current Opinion in Solid State and Materials Science xxx (2014) xxx–xxx 3

was suggested [53] that compositional similarity has a more signif-icant effect on in vitro bioactivity than the structural crystallinity. Acomparison between bone-like apatite formation ability of akerma-nite and wollastonite ceramics in SBF was performed by Wu et al.


[54]. At the early stage of soaking (up to 5 days), there was no obvi-ous bone-like apatite-layer formation on the akermanite surfacecompared to wollastonite. However, after 20 days, a 130 lm thickapatite layer was formed on the akermanite surface which was



Table 3Sintering and mechanical properties of akermanite ceramics, sintered at different conditions (modified from Ref. [55]).

Sintering temperature (�C) 1000 1200 1300 1350 1370 1370 1370Sintering time (h) 2 2 2 2 2 4 6Relative density (%) 57.4 ± 0.2 63.4 ± 0.5 70.1 ± 0.7 77.1 ± 0.9 82.8 ± 0.4 85 ± 1 89.6 ± 3.3Line shrinkage (%) 3.0 ± 0.5 3.6 ± 0.2 5.1 ± 0.6 7.7 ± 0.3 10.0 ± 0.1 10.2 ± 0.1 10.6 ± 0.2Bending strength (MPa) 7 ± 1 42 ± 3 78 ± 7 134 ± 9 135 ± 2 153 ± 2 176.2 ± 9.8Fracture toughness (MPa m1/2) 0.23 ± 0.07 0.3 ± 0.1 0.46 ± 0.17 0.94 ± 0.15 1.47 ± 0.07 1.61 ± 0.04 1.83 ± 0.10Young’s modulus (GPa) 16.9 ± 0.2 18.6 ± 4.2 32.8 ± 5.3 28.5 ± 2.6 23 ± 2 21 ± 7 42.0 ± 5.4


similar to the 110 lm thick apatite layer formed on the wollastonitesurface. Since the only compositional difference between akerma-nite and wollastonite is the presence of Mg in the akermanitecomposition, the authors suggested that the lack of apatite formingability for akermanite ceramics at the early stage of soaking could bedue to an initial inhibition effect of the released Mg ions on apatitecrystallization and growth. However, with increase of soaking time,the faster release of Ca2+ compared to Mg2+, led to supersaturationof the Ca ions in the solution, and aided crystallization and growthof the bone-like apatite layer. Hence, after 20 days of soaking, thebone-like apatite-forming ability of the akermanite ceramics wassimilar to that of bioactive wollastonite ceramics. Moreover, theevaluation of ionic concentration and pH of the SBF solution duringthe soaking, suggested that the apatite formation mechanism forakermanite ceramics is similar to that of wollastonite ceramics[54]. In another study, immersion of 90% porous akermanite scaf-folds for 10 days in SBF led to formation of a hydroxyapatite(HAp) layer with a 1.62 Ca/P ratio, and HAp crystals of around100 nm in size [57]. Furthermore, degradability of the akermanitescaffolds was confirmed by 28 days immersion in Ringer’s solution,and by detecting the concentration of Ca, Mg, and Si released fromthe scaffolds, in addition to the scaffold weight loss, which increasedover the immersion time. A study on a drug delivery system basedon gentamicin sulphate-akermanite showed that incorporation of10 wt.% gentamicin sulphate in akermanite discs does not influencethe in vitro bioactivity of akermanite over 21 days of immersion inSBF [56].

2.1.4. In vitro biological properties of akermaniteAdipose stem cells [63], mesenchymal stem cells [39,64,65],

primary cultures of osteoblasts [54], osteoblast cell lines [66],fibroblast cell lines [54], stromal cells [67], ligament cells [68],endothelial cells [69], and bacteria as Escheichia coli [58] havebeen used to evaluate the influence of akermanite on cell attach-ment, spreading, proliferation, differentiation, as well as itsbacterial activity. Human adipose-derived stem cells (hASCs) havebeen studied [63] for their attachment, proliferation and osteo-genic differentiation over akermanite samples, and their behav-iour was compared with hASCs cultured on beta-tricalciumphosphate (b-TCP). It was found that the attachment and prolifer-ation of hASCs on akermanite ceramics were as good as those ofthe cells cultured on b-TCP, and their differentiation and prolifer-ation on akermanite ceramics were substantially higher after10 days of culture. Similarly, a study [39] using human bonemarrow-derived mesenchymal stromal cells (hBMSCs) showed asignificant enhanced proliferation of hBMSCs after 20 days of cul-ture with akermanite extract, when compared to cells exposed tob-TCP extract. The cell cultures were made with a 1/256 dilutionof a 200 mg/mL extract. The alkaline phosphatase (ALP) activitywas also enhanced in akermanite extracts. Furthermore, theexpression of alkaline phosphatase, osteopontin, osteocalcin, andbone sialoprotein was significantly enhanced in akermaniteextracts when compared to those cells cultured in b-TCP extract.In a similar investigation [54], osteoblasts isolated from calvariaof neonatal Sprague–Dawley rats, and the fibroblast cell line


L929 were used to perform indirect and direct evaluations ofakermanite ceramics. For the indirect evaluation, extracts ofakermanite powder with dilution ratios between 1 and 1/320,000 were prepared, and results showed that Ca, Si, and Mgions at a certain concentration range released from akermaniteceramics significantly stimulated osteoblast and L929 cell prolifer-ation. Osteoblast spread well on the surface of akermanite ceram-ics, and proliferated with increasing culture time. Sun et al. [67]investigated the proliferation, ALP activity, morphology, differen-tiation, and gene expression of human bone marrow stromal cells(hBMSC) cultured on akermanite bioactive ceramics. Discs ofb-TCP were used as a reference. By scanning electron microscopy(SEM) imaging taken 24 h after seeding it was clear that hBMSCspread well on both biomaterials, but with different morpholo-gies. The proliferation was more significant on akermanite thanon b-TCP. Furthermore, by using ALP staining, it was possible toobserve almost twice higher ALP activity on cells seeded onakermanite discs than those on b-TCP discs. Additionally, the cellswere cultured with akermanite extracts-containing osteogenicmedia as well as in akermanite extracts-containing medium with-out osteogenic reagents (L-ascorbic acid, glycerophosphate, anddexamethasone). It was shown that the osteoblastic differentia-tion of hBMSC was enhanced in both types of culture medium.These results suggest that akermanite is able to enhance cell pro-liferation, as well as to promote osteoblastic differentiation ofhBMSC in vitro by up-regulating osteogenic gene expression. Asimilar investigation was carried out by Xia et al. [68] but usinghuman periodontal ligament cells (hPDLCs). The results showedenhanced attachment, proliferation, and osteogenic differentiationwhen hPDLC’s were exposed to akermanite ceramics. In anotherrelevant investigation, Zhai et al. [69] have studied the effect ofionic extracts of akermanite ceramics on proliferation and geneexpression of human aortic endothelial cells (hAECs), as well asthe in vitro angiogenesis by using an in vitro angiogenesis assaykit (ECMatrix™). This study revealed that akermanite ceramicsare able to induce both osteogenesis and angiogenesis. In an effortto understand the underlying mechanism involved in the osteo-genic differentiation of hASCs exposed to akermanite ceramics,Gu et al. [44] used hASCs as a model for determining the rolesof mitogen-activated protein kinases (MAPKs) in the regulationof osteogenic differentiation of this cell line. It was found thatthe akermanite extract which contains Ca, Mg and Si ions withconcentrations of 2.36, 1.11 and 1.03 mM, respectively, has a posi-tive effect on osteogenic differentiation of hASCs via extracellularsignal-released kinases (ERK) signalling pathway. Furthermore,akermanite ceramics have been reported to induce antibacterialeffects [58]. It was shown that the bactericidal effect of akerma-nite on Escheichia coli (E. coli) (ATCC25922) is improved as the sizeof akermanite particles in solution is decreased [58]. The resultsshowed that akermanite particle concentrations between1 mg/mL and 100 mg/mL can lead to antibacterial activitiesbetween 20% and 80%, respectively. In general, the antibacterialactivity is improved with increasing Ca ions in solution, whichis directly correlated with increase of the particle concentrationand decrease of the particle size.




2.1.5. In vivo biological properties of akermaniteFew investigations have been carried out to evaluate the in vivo

properties of akermanite ceramics. Huang et al. [39] investigatedthe implantation of akermanite and b-TCP ceramics into the cavi-ties of New Zealand rabbits. After 16 weeks of implantation, mea-surements showed that the mineral apposition rate of new boneformation in implanted akermanite scaffolds was higher than thatin the b-TCP implants. Van Gieson’s picric-fuchsine stainingrevealed that new bone formation and material degradation weremuch more evident in the akermanite implantation than in theb-TCP after both 8 and 16 weeks. Akermanite implants enhancedbone regeneration because of their superior degradation rate andbiocompatibility. In a relevant investigation, Zhai et al. [69]implanted porous bioceramics into the cavities (6 mm in diameter)of New Zealand rabbits near the distal femur of the left and rightlegs. In this study, in vivo angiogenesis and new bone formationwere studied via Van Gieson’s picric-fuchsine staining. New boneformation was observed within the porous structure of theakermanite scaffolds at weeks 8 and 16. Under high magnification,Haversian system canals, composed of Haversian and Volkman’scanals were observed together with the new bone formed. Bloodvessels were found in much greater number in the akermanitebioceramic when compared to the b-TCP used as a control. Theseinteresting results revealed that akermanite bioceramic degrada-tion products have potential to promote vascularized bone forma-tion and angiogenesis was faster in the akermanite bioceramicimplant than in the b-TCP-bioceramic implant.

2.1.6. ApplicationsAkermanite ceramics have been used for development of bone

tissue engineering scaffolds [57], and drug delivery systems[56,71] as well as coatings on metallic implants [64]. Wu et al.[57] fabricated porous akermanite scaffolds using sol–gel derivedakermanite powders via the polymer sponge method (foam-rep-lica). It was shown that using the foam-replica method, it is possi-ble to fabricate mechanically stable akermanite scaffolds withinterconnected porosities in the range 63.5–90.3%, pore sizes inthe range 300–500 lm and about 100 lm thick pore walls. The cellculture study was carried out by using osteoblasts isolated fromcalvaria of neonatal Sprague–Dawley rats. It was demonstratedthat the ionic dissolution of porous akermanite scaffolds stimu-lated osteoblast proliferation. Moreover, good attachment, spread-ing and proliferation of the bone marrow stromal cells, obtainedfrom a newborn calf femur, on the akermanite scaffolds wereobserved. The results of the study [57] indicated that the highlyporous akermanite scaffolds are degradable, bioactive and cyto-compatible and therefore might be used for bone tissue engineer-ing applications.

In another study [56], the drug release behaviour of akermaniteceramics was evaluated and compared to that for wollastoniteceramics. 10 wt.% of gentamicin sulphate was incorporated withinthe sol–gel derived akermanite powders, and the ceramic diskswere obtained by a two-step pressing of the powders. The drugrelease behaviour was studied by immersion of the disks in phos-phate buffered saline (PBS) solutions for 21 days. UV–VIS spectros-copy showed a lower drug release rate from akermanite comparedto that from wollastonite. It was shown that, for both materials,most of the drug is released within the first week of immersionwhich is followed by a significant decrease of the drug release rate.In another relevant study, it was shown [70] that incorporation ofakermanite powders within a poly (lactic-co-glycolic acid) (PLGA)matrix using a thermally induced phase separation (TIPS) methodcan produce composite microspheres with improved structural,physiochemical, mechanical, and biological properties in compari-son to neat PLGA beads. Composites of PLGA/Akermanite withweight ratios of 5:1 and 2:1 showed enhanced cell proliferation


when compared with pure PLGA beads. The ALP activity after7 days of exposition to the composite beads was substantiallyincreased when compared with the control. In addition, the pHof the surrounding medium could be efficiently controlled by add-ing akermanite to PLGA, thus having a positive effect on cell behav-iour. Therefore, it was concluded that incorporation of akermanitepowders neutralizes the biopolymer (PLGA) acidic products, con-trols the release of bovine serum albumin (BSA) from the compos-ite, induces the formation of an apatite-layer in SBF, and enhancesthe proliferation and alkaline phosphatase activity of bone marrowstromal cells.

Yi et al. [64] developed akermanite coatings on Ti alloys (Ti–6Al–4V) via plasma spraying. The coatings had thicknesses of�200–500 lm and bonding strength of �38.7–42.2 MPa. Theakermanite coatings were able to induce apatite mineralizationon their surface when immersed in SBF. After two days of SBFimmersion, apatite particles were observed on the surfaces andthey covered completely the coating surface after 6 days of immer-sion. Three regions could be observed in the cross section of theSBF immersed coatings which were akermanite, a Si-rich layerand an apatite layer. The Si-rich layer was suggested to play akey role in the apatite formation mechanism of akermanite ceram-ics. Finally, cell culture studies with rabbit bone marrow MSCsshowed that akermanite coatings supported cell differentiationand proliferation.

2.2. Bredigite

2.2.1. SynthesisSol–gel [61,71,72], combustion [73] and mechanical activation-

annealing [74,75] methods have been reported for preparation ofpure phase bredigite ceramic powders. Wu et al. [61,71,72] pre-pared bredigite ceramic powders by sol–gel process. It wasreported [71,72] that by calcination of the dry gel at 1150 �C for3 h pure bredigite powders with particle size in the range of 1–10 lm could be obtained. Synthesis of pure nanocrystalline bredig-ite powders via a combustion method was reported by Huang andChang [73]. It was reported that using a simple solution combus-tion method, pure nanocrystalline bredigite powders could be pre-pared at a relatively low temperature (650 �C), with particle sizesin the range 234–463 nm and containing spherical bredigite crys-tallites with sizes in the range 90–150 nm. Tavangarian and Emadi[74] synthesized nanostructured bredigite powders by mechanicalactivation (60 h milling) with subsequent annealing at 1200 �C for1 h. The prepared powders had a mean particle size around 779 nmand a mean crystallite size around 50 nm. Moreover, it was possi-ble to obtain bredigite nanopowders with altered particle and crys-tallite sizes by changing the milling time [75].

2.2.2. Sintering and mechanical propertiesNot many studies have focused on sintering and mechanical

properties of bredigite ceramics for biomedical applications. Inone of the few studies available, Wu et al. [61] prepared dense bre-digite ceramics by two-step pressing (uniaxially at 10 MPa and iso-statically at 200 MPa) of sol–gel derived bredigite powders andsintering at 1350 �C for 8 h. The sintered bredigite ceramics had arelative density of 94.2 ± 1.2%, bending strength of 156 ± 6 MPa,fracture toughness of 1.57 ± 0.12 MPa m1/2, and Young’s modulusof �43 GPa.

2.2.3. In vitro bioactivity and degradationWu et al. [61] evaluated the in vitro bioactivity of press-sintered

sol–gel derived bredigite discs by 20 days immersion in SBF. Afterimmersion for 7 days, lath-like nano-sized HAp crystals wereformed on the bredigite surface. By 10 days, the HAp crystallitesbecame more compact and were around 200–300 nm and




100 nm in length and diameter, respectively. By 20 days, the HApcrystallites became spherical with diameter in the range 80–100 nm. At last, the formed layer consisted of a 70 lm thick sil-ica-rich layer with a 20 lm thick HAp layer on it, which had theCa/P ratio of 1.65. By immersion of the pellets in SBF, the pH ofthe SBF solution increased during the first day and afterward itdecreased. The Mg and Si concentrations in SBF increased, P con-centration decreased, and Ca concentration increased in the first7 days and afterward decreased. These results suggested a HAp-layer formation mechanism for bredigite ceramics similar to thatindicated for wollastonite ceramics [61]. In another study by Wuand Chang [72], sol–gel derived bredigite powders were immersedfor 10 days in SBF which led to formation of aggregated HAp crys-tallites (50 nm in diameter) on the bredigite surfaces. Moreover,during immersion in SBF, Ca, Mg, and Si concentrations increased,while the concentration of P in the solution decreased. Thereviewed studies [61,72] suggested a comparable HAp-formationability of bredigite to bioactive glasses. Moreover, the apatite-forming ability of nanocrystaline bredigite powders prepared bya combustion method was evaluated by Huang and Chang [73].The nanopowders were immersed in SBF for 4 days and numerousuniform wormlike HAp crystallites (80–100 nm in length) with aCa/P ratio of 1.63 were formed on the bredigite powders, suggest-ing the superior in vitro HAp forming ability of the nanostructuredbredigite powders prepared by combustion method, over previ-ously reported [72] sol–gel derived coarse grained bredigitepowders.

2.2.4. In vitro biological propertiesThere have been only few studies regarding the biological prop-

erties of bredigite bioceramics. Recently, Zhou et al. [77] reportedthe cementogenic proliferation and differentiation of a primaryculture of periodontal ligament cells (PDLCs) after exposition tobredigite extracts at different concentration ranges. The possiblemechanism for the stimulatory effect of bredigite on PDLCs is theactivation of the Wnt/b-catenin signalling pathway. Furthermore,in a recent study of Zhai et al. [76] the effect of bredigite ceramicson the osteogenic differentiation of hBMSCs and the angiogenesisof HAECs were explored. The results of this study indicated thatbredigite showed the highest osteogenic and angiogenic potentialwhen compared to akermanite and diopside. In the previous dis-cussed study of Wu et al. [61], the effects of bredigite ceramicsextracts on the proliferation of the mouse fibroblasts cell lineL929, as well as the adhesion and proliferation of osteoblastsseeded on bredigite disks were investigated. The results fromexposing fibroblasts to bredigite ceramic extracts with concentra-tions of 6.25, 12.5 and 25 mg/mL indicated an enhancement in cellproliferation of nearly 2-fold when compared with the cells cul-tured without bredigite extracts, while cell proliferation was inhib-ited for bredigite extract concentrations between 100 and 200 mg/mL. Moreover, the osteoblast cells seeded on the bredigite surfaceshowed good attachment and proliferation after 6 h of cell culture.The same author reported that extracts of bredigite powders, at acertain concentration range, stimulated osteoblast proliferationin vitro [72]. These results showed that bredigite ceramics canstimulate osteoblastic cell growth at a certain concentration range,and exhibit good cytocompatibility which make them interestingbiomaterials for tissue engineering applications. Nevertheless, fur-ther studies are needed in order to evaluate the effects of bredigitebioceramics on different cells genotypes and lineages, as well astheir in vivo properties.

2.2.5. ApplicationsBredigite ceramics have been used for development of bone

tissue engineering scaffolds [71]. Wu et al. [71] developed porousbredigite scaffolds using sol–gel derived bredigite powders via a


polymer sponge method which included a sintering step at1350 �C for 3 h. The obtained scaffolds with interconnected poresof around 300–500 lm and pore walls of about 60–100 lm wereimmersed in SBF (pH 7.40, 37 �C) for 10 days. The immersion inSBF was used as a biomimetic process to mimic in vivo bone bio-mineralization and led to formation of a bone-like apatite layeron the scaffold struts, with apatite crystals of around 100 nm andCa/P ratio of around 1.61. The bredigite scaffolds with biomimeticapatite layer (BTAP) had porosities of about 90% and were soakedin Ringer’s solution (pH 7.40, 37 �C) in a shaking water bath upto 28 days, with the solution volume to scaffolds mass ratio of200 ml/g to evaluate their degradation. The results showed thatweight loss of the BTAP scaffolds increased with increasing soakingtime, and it was similar to that of similarly prepared b-TCP scaf-folds. Fig. 1 shows SEM and energy-dispersive X-ray spectroscopy(EDS) analysis of the bredigite scaffolds before and after immersionin SBF.

Compression tests of bredigite scaffolds before immersion inSBF, after immersion in SBF (BTAP), and of a similarly preparedb-TCP scaffold showed that compressive strength of the BTAP scaf-fold (101 kPa) was lower than that of the pure bredigite scaffold(233 kPa), however it was still higher than that of the b-TCP scaf-fold (50 kPa). A cell culture study of osteoblast-like cells in contactwith BTAP and b-TCP scaffolds showed that the osteoblast-likecells adhere and spread well on BTAP scaffolds. Moreover, the cellsseeded on BTAP scaffolds showed a higher ALP activity than thoseseeded on b-TCP scaffolds. This could be due to a stimulating effectoriginated from the Mg and Si ions release effect of the BTAP scaf-folds. In general, cells exposed to the BTAP scaffolds, compared tothose exposed to the b-TCP scaffolds, showed improved prolifera-tion rate and differentiation.

2.3. Diopside

2.3.1. SynthesisSolid-state [40], sol–gel [78] and coprecipitation [38] processes

have been used for production of diopside ceramics for biomedicalapplications. Nonami and Tsutsumi [40] prepared diopside pow-ders by mixing the appropriate ratio of CaCO3, MgO and SiO2, cal-cination at 1100 �C, and at last wet grinding. In another study byNonami et al. [78] sol–gel synthesis of diopside powders was car-ried out using metal alkoxides. It was reported that the sol–geldiopside powders prepared in this study had crystallization tem-perature of 840 �C. Iwata et al. [79] used metal alkoxide and metalsalts for sol–gel synthesis of diopside ceramics. The sol–gel diop-side powder prepared by this method had considerably lower crys-tallization temperature (751.4 �C) than the sol–gel diopsidepowder reported by Nonami et al. [78] that was prepared withoutusing metal salts. In another study by Iwata et al. [38] polycrystal-line diopside ceramics were synthesized by a coprecipitation pro-cess using a metal alkoxide and metal salts. The dried powderprepared by this method could be crystalized into diopside at845.5 �C. The crystallization process occurred smoothly withoutusing sintering additives [38]. Finally, successful synthesis ofwell-ordered mesoporous diopside ceramics has been achievedvia a sol–gel method using appropriate surfactants [80].

2.3.2. Sintering and mechanical propertiesMechanical properties of dense and porous diopside ceramics

have been studied. In two relevant studies [40,81], diopside ceram-ics were prepared by pressing diopside powders and firing thecompacts in air at 1300 �C for 2 h. The obtained samples had a den-sity of 3.20 g/cm3, bending strength of 300 MPa, fracture toughnessof 3.5 MPa m1/2, Young’s modulus of 170 GPa and Poison’s ratio of0.35. The mechanical properties were all higher than those forHAp sample with density 3.16 g/cm3. Wu et al. [82] studied the



Fig. 1. SEM micrographs of bredigite scaffolds according to Wu et al. [71] (a) before immersion in SBF, (b and c) after immersion in SBF and (d) the EDS analysis of the formedapatite layer (reproduced from Ref. [72] with permission of Springer).


mechanical properties of porous diopside structures (scaffolds),which will be discussed in Section 2.3.6.

2.3.3. In vitro bioactivity and degradationTo the author’s knowledge, diopside ceramics were the first

type of Mg-containing silicate ceramics reported to be bioactive[60]. For example, Miake et al. [62] investigated the in vitro HApcrystals formation mechanism on diopside surfaces by immersionof granular diopside (0.4–0.6 mm in size) for 5 days in SBF. Afterimmersion, many needle- or platelet-like crystals were observedon the diopside surface that were about 25 Å thick, 100–200 Åwide, 100–300 nm long and had a Ca/P molar ratio of 1.38 ± 0.10,which suggested to be Ca-deficient non-stoichiometric HAp. Theauthors hypothesized that at early stage of crystal formation, epi-taxial crystal growth of Octa-Calcium Phosphate (OCP) or b-TCPmay occur on the diopside surface in SBF, which is then followedby a phase transition into HAp. Nonami and Tsutsumi [40]immersed press-sintered diopside pellets in aqueous lactic acidand physiological salt solutions for 2 days. The results indicatedthat the weight loss for the sample immersed in the lactic acidsolution was 2.8%, while for the sample immersed in the physio-logical salt solution the weight loss was 0.05%. These values werereported to be lower than those for press-sintered HAp samples.Iwata et al. [79] evaluated the apatite-forming ability of sol–gelderived diopside ceramics by immersion of diopside pellets inSBF for up to 7 days. After 3 days of immersion, fine apatite parti-cles were formed on the diopside surfaces, which grew to massiveparticles after 7 days of immersion and formed a bone-like apatitelayer on the diopside surface. Analysis of the SBF solutions showedthat the Mg2+ ions concentration slightly increased and almostremained constant during the immersion periods, suggesting lowpartial dissolution of Mg2+ ions from the diopside pellets. Inanother study, Iwata et al. [38] investigated the in vitro bioactivityof diopside pellets prepared by press-sintering coprecipitationderived powders. It was reported that after 3 days immersion ofthe samples in SBF, leaf-like particles uniformly formed on thediopside surfaces and continued to grow during the 7 days ofimmersion. Furthermore, analysis of the SBF solutions showed thatthe magnesium content almost remained constant during theimmersion period, which suggested that the highest content ofMg2+ remained in the diopside pellets. Accordingly, the authors


suggested that Mg2+ ions do not play a role for apatite formationof diopside ceramics. A schematic illustration of the apatite-form-ing mechanism on the surface of diopside in SBF is shown in Fig. 2,according to Ref. [38].

Another relevant study regarding the in vitro bioactivity ofdiopside ceramics was performed by De Aza et al. [83]. Diopsidepellets were immersed in simulated human parotid saliva (HPS)with addition of several thymol crystals to avoid bacteria forma-tion, up to 1 month. After 1 day of immersion, spherical HAp-likeparticles were formed on the diopside surface which transformedto a complete layer of globular HAp-like particles after 7 days ofimmersion. By longer immersion times, a continuous compactlayer on the diopside surface was formed. The results indicated afast HAp formation rate at the first week of immersion, whichdecreased during the following weeks due to depletion of P ionsin the HPS solution and due to the negative effect of the formedHAp layer on the ion exchange process between the underlyingdiopside and the HPS solution. A high pH condition (9.8) wasreported on the diopside/HPS interface which was caused by theionic exchange at the interface. Furthermore, a non-continuous sil-ica-rich interlayer was present between the diopside and the HAplayer due to partially dissolution of the amorphous silica phase inthe high pH environment, prior to the formation of the HA-likephase. The results indicated that the formed HA-like layer on diop-side had an ultrastructure similar to that of natural cortical boneand also dentine.

2.3.4. In vitro biological propertiesKobayashi et al. [60] reported the cytocompatibility of diopside

with the osteogenic cell line MC3T3-E1, by measuring ALP activity,protein content, crystals secretion, and by cell morphology evalua-tion. The results showed that ALP activities of MC3T3-E1 cells onthe diopside disks were similar to those of the control, which con-sisted on thermanol coverslips. The cell differentiation and growthwere not inhibited on diopside disks and mineral crystallizationoccurred in MC3T3-E1 cells seeded on diopside. Nonami and Tsuts-umi [40] seeded the fibroblast cell line L-929 (ATC, CCL1 NCTCclone 929 strain) on diopside discs and used a polystyrene sampleas the control. No significant differences in cell adhesion betweendiopside and the control were observed, meaning that diopsidewas not cytotoxic. Finally, the work of Zhai et al. [76] showed that



Fig. 2. Schematic illustration of the apatite-forming mechanism on the surface of diopside in SBF (reproduced from Ref. [38] with permission of Elsevier).


diopside stimulated angiogenesis of human aortic endothelial cellscultured on the ECMatrix™ compared to b-tricalcium phosphatebioceramics.

2.3.5. In vivo biological propertiesKobayashi et al. [60] carried out the histological evaluation and

pulling test of diopside using apatite wollastonite-containingglass-ceramic (AWGC) as the control. Prepared blocks of diopsideand AWGC were implanted in forty tibial metaphyses of white Jap-anese rabbits. No fibrous tissue formation or foreign body reactionwas observed at the interface of bone and AWGC after 12 and20 weeks of implantation. Diopside-bone attachment occurredthrough the formation of a Ca–P layer, and the diopside surfacewas degraded and replaced by new bone. Nonami and Tsutsumi[40] implanted rectangular diopside specimens into holes formedin the lower edge of the jaw bone of adult male rabbits. In parallel,a diopside dental root sample was implanted into a cavity in themandible of a Japanese monkey and left for six months. The resultsof both cases showed that diopside forms a uniform junction withnewly grown bone, in which growth of apatite crystals wasobserved. In addition, continuity between the diopside latticeand that of the newly formed apatite crystals was also observedon the interface. In a different work, Luo et al. [84] reported thein vivo osteogenesis of bioactive diopside microspheres as a bonefilling material. These microspheres were implanted in the femurof Wistar rats and it was observed that new bone tissue formedbridges between spheres after 2 weeks of implantation. Further-more, higher amount of bone was found within defects filled withdiopside compared to b-TCP spheres after 2 and 4 weeks of implan-tation. Finally, although the actual effect that the presence of Mghas on the biological properties of the material has not been clearlyreported, these few studies showed that diopside was capable ofinducing osteogenesis in vivo and therefore could be a promisingbioactive material for bone filling, artificial bones, and dental roots.

2.3.6. ApplicationsDiopside ceramics have been used for production of bone tissue

engineering scaffolds [82] and coatings on biomedical implants[85,86], and have been suggested for surgery hemostasis [80] andin vivo imaging [87] applications. Wu et al. [82] prepared porousdiopside scaffolds using coprecipitation-derived diopside powdersvia the polymer sponge method. The fabricated scaffolds werehighly porous (80%) and had interconnected pores with poresizes around 300 lm. With decreasing the scaffold porosity from90% to 75%, the compressive modulus and strength increased from10 ± 3.3 MPa and 200 ± 20 kPa, to 68 ± 20 MPa and 1360 ± 370 kPa,respectively. The compressive modulus of the diopside scaffoldswas maintained after 14 days immersion in SBF, but themechanical strength decreased about 30%. The degradation studies


indicated that diopside scaffolds retained a sustained Si release andhad 2% weight loss after 28 days immersion in SBF. Human osteo-blastic-like cells (HOBs) were cultured on diopside scaffolds withan initial cell concentration of 2 � 104 cells per scaffold and incu-bated up to 7 days. Attachment, morphology, mithocondrial activ-ity and ALP activity of the cells indicated that diopside has acomparable cytocompatibility to HAp scaffolds. Fig. 3 shows HOBcells on the surface of diopside scaffolds [82].

Diopside nanopowder has been used as reinforcing agent (0–40 wt.%) in silk fibrion polymer–matrix porous scaffolds producedvia a freeze-drying method [88]. Addition of diopside nanopawderto the matrix could improve the scaffolds mechanical propertiesand hydrophilicity as well as proliferation of mouse preosteoblastcells cultured with the scaffolds extracts. These composite scaf-folds might be used as pre-implant biomaterials for maxillofacialbone regeneration [88].

Xue et al. [85,86] developed diopside coatings onto Ti–6Al–4Vsubstrates via an atmospheric plasma spraying technique. Thediopside coatings had a thermal expansion coefficient (TEC) ofaround 8.41 � 10�6/�C (20–600 �C), that was closer than the TECof HAp (15.20 � 10�6/�C (20–600 �C)) to the TEC of Ti-alloy(9.40 � 10�6/�C). The immersion of the coated samples in SBF for15 days indicated formation of a HAp layer exhibiting a Ca/P ratioof 1.62 and 100 nm average size of HAp crystals. Investigation ofthe mechanical properties showed that the diopside coatings hada bonding strength of �32.5 MPa (higher than typical HAp coat-ings), and Young’s modulus of 38.56 GPa (close to that of corticalbone). Moreover, adult human spine osteoblast cells cultureddirectly onto the diopside coating revealed good attachment,spreading, and proliferation of the cells, confirming the cytocom-patibility of diopside [86]. Another study [80] showed that sol–gel derived well-ordered mesoporous diopside ceramics not onlyexhibit excellent bioactivity but also, due to their high surface area,can absorb a large amount of water showing high hemostatic activ-ity. Therefore, mesoporous diopside can also be applied for surgeryhemostasis. Finally, MnII-doped diopside exhibits red long-lastingphosphorescence and can be used for in vivo imaging of small ani-mals [87].

2.4. Forsterite

2.4.1. SynthesisSolid-state [89–94], sol–gel [41,95–97] and mechanochemical

[98] methods have been reported for synthesis of forsterite ceram-ics for biomedical applications. A common difficulty for productionof pure forsterite ceramics is formation of enstatite (MgSiO3) and/or periclase (MgO) impurities. Ni et al. [41] used a sol–geltechnique to synthesize agglomerated forsterite powders withparticle sizes in the range 5–50 lm. It was shown that calcination



Fig. 3. Morphology of HOB cells on the surface of diopside scaffolds (a and b) at 3 days, and (c and d) 7 days of cell culture (reproduced from Ref. [82] with permission ofElsevier).


temperature of 1200 �C was required to obtain stoichiometric for-sterite. Significant research has been carried out at Isfahan Univer-sity of Technology (Iran), for fabrication of forsterite ceramics forbiomedical purposes. For example, Fathi and Kharaziha [89] inves-tigated a simple solid state method for preparation of pure nano-crystalline forsterite powders. It was shown that purenanocrystalline forsterite could be successfully synthesized by10 h of mechanical activation (MA) followed by heat treatment at1200 �C. Furthermore, with decreasing the heat treatment temper-ature from 1200 �C to 900 �C, the crystalline size of forsteritedecreased from �57 nm to �21 nm. In general, MA could induceformation of nanocrystalline forsterite within particles smallerthan 1000 nm. Similar studies [90–92] led to production of nano-crystalline forsterite powders with particle sizes smaller than1 lm and crystalline size in the range 28–60 nm. In another studyby Fathi and Kharaziha [98] pure nanocrystalline forsterite pow-ders were synthesized via a mechanochemical method in the pres-ence of fluorine ion by 5 h of MA and subsequent heat treatment at900 �C for 1 h. It was shown that the presence of fluorine ionsdecreased the forsterite formation rate via combustion reaction,and by using a heat treatment process the fluorine ions could beremoved via hydrolysis phenomena. The synthesized forsteritenanopowders had particle sizes <100 nm and containednanocrystalline grains of around 30 nm in size. In an attempt[95] to produce forsterite nanoparticles, bioactive forsterite nano-powder with particle sizes in the range 25–45 nm was synthesizedvia a sol–gel method. Sanosh et al. [96] synthesizedforsterite nanopowders with particles size in the range 5–90 nm(average size �27 nm) which contained forsterite crystals ofaround 22 ± 7 nm in size, via a sol–gel method. Thenucleation and growth mechanism of the forsterite nanoparticlesduring the sol–gel process are schematically illustrated in Fig. 4[96].

In a recent study [97] three different temperatures (800, 900,and 1000 �C) were used for the heat treatment of sol–gel derivedforsterite. It was reported that forsterite was the main crystallinephase of the materials heat treated at 900 and 1000 �C. However,periclase (MgO) phases were present in all the samples. The


samples heat treated at 800 and 900 �C contained forsterite nano-crystals smaller than 60 nm, but the sample heat treated at 1000 �Ccontained micrometre-sized forsterite grains.

Ramesh et al. [94] synthesized forsterite nanocrystalline pow-ders via a solid-state method. The powders were synthesized viaMA (ultrasonication) and ball milling of MgO and talc for differentdurations. Moreover, the effect of a heat treatment step on theproperties of the powders was studied. It was reported that homo-geneity of the mixtures of the initial precursors play a key role inphase stability of the forsterite powders and higher amplitude ofultrasonication and prolonged ball milling could decrease the pres-ence of secondary phases (MgO and enstatite) while improving thecrystallinity of the powders. Heat treatment could significantlyavoid the secondary phases but even by heat treatment at hightemperatures (1500 �C) some secondary phases were observed. Inaddition, higher temperatures led to forsterite nanocrystalsgrowth.

In related studies, Lee et al. [93] reported the synthesis offorsterite powders by a solid-state reaction followed by 2 h heattreatment processing of powders at 1200 �C. Interestingly,Durdu and Usta [99] reported that using a micro oxidationmethod in aqueous solutions containing sodium silicate led toproduction of coatings containing forsterite phases ontomagnesium surfaces. This approach might be useful for develop-ment of forsterite coatings onto magnesium based metallicimplants controlling the implant degradation and inducing abioactive surface. Moreover, forsterite and forsterite containinginorganic composites can be synthesized by heat treatment ofpreceramic polymers [100] which can be useful for production ofcomplex structures.

2.4.2. Sintering and mechanical propertiesIn a study [41], sol–gel derived forsterite powders (5–50 lm

particle size) were uniaxially pressed at 10 MPa, cold isostaticallypressed at 200 MPa, and sintered at 1350–1550 �C. Table 4 showsthe mechanical properties of dense forsterite ceramics obtainedby different sintering conditions.



Fig. 4. Schematic diagram showing the nucleation and growth mechanism of forsterite nanoparticles during a sol–gel process (reproduced from Ref. [96] by permission fromElsevier).


As shown in the Table 4, sintering of forsterite powders at1450 �C for 8 h led to optimum mechanical properties. Fathi andKharaziha [101] used a two-step sintering (TSS) process to controlaccelerated grain growth of sintered solid-state process derivedforsterite nanopowders, in order to produce dense nanocrystallineforsterite ceramics with improved properties. TSS was performedon forsterite nanopowder green bodies at T1 = 1100–1300 �C andT2 = 750 and 850 �C. The results indicated that when using TSS atT1 = 1300 �C (for 6 min) and T2 = 750 �C (for 15 h), dense nanocrys-talline forsterite ceramics (�98.5% theoretical density) with grainsizes of around 60–75 nm could be obtained. The nanostructuredforsterite ceramics had a fracture toughness of 3.61 ± 0.1 MPa m1/

2 and microhardness of 940 ± 10 HV, which were higher than thevalues for typical hydroxyapatite and coarse grain forsteriteceramics. In another study by Kharaziha and Fathi [42], a similarTSS process (T1 = 900–1300 �C, T2 = 750 and 850 �C) was appliedon sol–gel derived forsterite nanopowders with particle sizes ofaround 25–45 nm. Using this method, nanostructured forsteriteceramics could be obtained, which exhibited hardness of 1102HV and fracture toughness of 4.3 MPa m1/2; these values werehigher than those for the previous TSS prepared nanostructuredforsterite ceramics [101]. This superior mechanical propertiescould be due to the smaller nanopowders prepared by the sol–gel method, compared to forsterite nanopowders prepared bymechanical activation and heat treatment processes, that couldinduce improved sinterability and lead to smaller grain size andlower residual pores in the sintered forsterite specimens. In a studyby Lee et al. [93], dense forsterite materials were prepared byuniaxially pressing (2.5–3.0 MPa) of forsterite powders, obtainedfrom a solid-state method, followed by cold isostatic pressing(200 MPa) and sintering for 3 h at different temperatures (1200–1500 �C). The results revealed that the optimum mechanical prop-erties could be achieved by sintering at 1400 �C which led to rela-tive density of 88.3%, Vickers microhardness of 7.11 GPa, andfracture toughness of 4.88 MPa m1/2. Moreover, the mechanicalproperties of porous nanostructured forsterite ceramics (scaffolds)prepared by a gel-casting method were studied by Ghomi et al.

Table 4Mechanical properties of forsterite ceramics obtained by different sintering conditions [41

Sintering T (�C) Sintering time (h) Relative density (%) Shrin

1350 6 82.6 ± 1.0 7.0 ±1450 3 86.6 ± 0.8 8.2 ±1450 6 87.7 ± 1.0 9.0 ±1450 8 92.5 ± 0.4 9.2 ±1550 6 91.4 ± 0.6 10.1 ±


[102]. Table 5 shows the crystallite size, apparent density, totalporosity, compressive strength, and elastic modulus of porous for-sterite structures sintered at different temperatures.

In this study, bioactive porous nano-forsterite structures withcompressive strength of around 2.43 MPa and elastic modulus ofaround 182 MPa could be achieved; these values are close to thelower limit of the mechanical properties of spongy bone [102].

Ramesh et al. [94] studied the effect of synthesis method andsintering procedure on microstructure and mechanical propertiesof forsterite sintered bodies. Forsterite nanocrystalline powderswere obtained via a solid-state reaction method. Heat treated ornon-heat treated forsterite nanocrystalline powders were firstuniaxially pressed (2.5–3.0 MPa) and then compacted via a coldisostatic press at 200 MPa. Finally, the forsterite green bodies weresintered at 1200 �C, 1300 �C, 1400 �C and 1500 �C. For the heattreated samples, the elastic modulus was directly proportional tothe forsterite grain size which was not significantly affected bythe sintering temperature. For the non-heat treated samples, theelastic modulus increased with increasing the sintering tempera-ture until 1300 �C, but further increase of the sintering tempera-ture led to a reduction of the elastic modulus through a largeincrease of the ceramic grain size.

2.4.3. In vitro bioactivity and degradationPrevious studies have suggested forsterite as a bioceramic for

bone repair applications [41,103]. However, coarse grained forste-rite ceramics were not able to induce an apatite-layer on their sur-faces in SBF and their degradation rate was extremely low [103].These results, in addition to the reports by Webster et al. [104–106] regarding the improved properties of nanophase ceramics,motivated Kharaziha and Fathi [95] to develop forsterite nanopow-ders and to investigate their biological properties. The in vitro bio-activity of the forsterite nanopowders was evaluated by theirimmersion in SBF up to 28 days. After 14 days of immersion, tinyspherical Ca–P particles with crystallite sizes of around400–500 nm were observed on the surfaces of forsterite nanopow-ders, which were transformed to clusters of agglomerated HAp

].

kage (%) Bending strength (MPa) Fracture toughness (MPa m1/2)

0.2 150 ± 8 1.8 ± 0.40.2 152 ± 8 2.1 ± 0.00.2 181 ± 9 2.3 ± 0.10.1 203 ± 8 2.4 ± 0.00.4 145 ± 8 1.6 ± 0.2



Table 5Crystallite size, physical and mechanical properties of porous forsterite ceramics (scaffolds) sintered at different temperatures [102].

Sintering T (�C) Crystallite size (nm) Apparent density (g/cm3) Total porosity (%) Compressive strength (MPa) Elastic modulus (MPa)

900 26 (±2) 0.45 (±0.04) 86 (±1) 2.06 (±0.09) 145 (±9)1000 28 (±1) 0.48 (±0.08) 85 (±1) 2.19 (±0.06) 165 (±12)1100 31 (±1) 0.55 (±0.05) 83 (±2) 2.31 (±0.07) 171 (±21)1200 35 (±2) 0.61 (±0.06) 81 (±1) 2.43 (±0.11) 182 (±19)


particles and covered the nanopowder surfaces after 28 days ofimmersion in SBF. Based on ionic concentration measurements inthe SBF solution used, the authors proposed that by immersionof forsterite nanopowder in SBF, initially Mg2+ ions from the nano-powders are exchanged with H+ ions present in SBF, pH increasesand negatively charged functional silanol groups form on the for-sterite surface. Then the Ca2+ and PO3

4� ions are absorbed on thesurface and formation of a HAp layer occurs. Therefore, forsteritenanopowder, unlike large particle-sized forsterite, was shown tobe bioactive, which is likely due to their high surface area obtainedfrom nanostructure processing. The bioactivity of forsteritenanopowders was also confirmed in a recent study [97] reportingformation of apatite on forsterite nanopowders by 7 days ofimmersion in SBF.

Moreover, immersion of forsterite pellets obtained by uniaxiallypressing of forsterite nanopowders with different crystalline size(20 nm and 31 nm) in SBF and Ringer’s solution, revealed thatthe forsterite ceramics with lower degree of crystallinity showhigher bioactivity and degradation rate, in SBF and Ringer’s solu-tion, respectively [107,108].

2.4.4. In vitro biological propertiesThe few cell culture studies that have been carried out for eval-

uation of in vitro biological properties of forsterite ceramics aresummarized in this section. Kharaziha and Fathi [42] investigatedthe cytocompatibility of nanostructured forsterite specimens viaindirect and direct cell cultures. A dilution of forsterite nanopow-der extract in contact with osteoblast-like G292 cells was used.Similarly, nanostructured discs in direct contact with osteoblast-like G292 cells were also analyzed. In the indirect assay, cell prolif-eration in presence of forsterite extracts was significantlyenhanced when compared with the negative control (medium freeof extract). Direct cell seeding on nanostructured discs showedgood attachment and spreading of the cells. These results suggestthat nanostructured forsterite bioceramics possess suitable cyto-compatibility. A similar result was reported by Naghiu et al. [97],who showed that forsterite powder extracts enhanced prolifera-tion of U20S-type human osteoblast cells. In a related investiga-tion, Ni et al. [41] evaluated the cytocompatibility of forsteritediscs using direct cell culture of osteoblast cells isolated from cal-varia of neonatal Sprague–Dawley rats. It was observed that thecells seeded on forsterite attached and spread before cells seededon the controls, and the cell proliferation on forsterite bioceramicswas significantly enhanced. In a different study, Ni and Chang[109] reported that the proliferation of osteoblast-like cellsisolated from calvaria of neonatal Sprague–Dawley rats wasstimulated by ionic products of forsterite bioceramic with Mgconcentrations in the range 1.04–0.204 mmol/L. In a related inves-tigation, Ni et al. [103] synthesized composites of b-CaSiO3/forste-rite and tested their cytocompatibility with osteoblast-like cellsisolated from calvaria of neo-natal Sprague–Dawley rats. Thematerial resulted in cell spreading, intimate cell-material contactand the formation of a flattened cell sheet that completely coveredthe surface of the material. Moreover, composites containing 50and 30 wt.% of forsterite were shown to enhance cell proliferationcompared to control samples of pure b-CaSiO3 and forsterite.


These few investigations confirmed that forsterite bioceramicscan improve the attachment, spreading and proliferation of osteo-blasts, which make them interesting bioceramics for bone tissueengineering. Further studies are required to explore the interactionof forsterite with other cell lineages, as well as its capacity toinduce cell differentiation in vitro and angiogenesis in vivo.

2.4.5. ApplicationsForsterite ceramics have been used for development of bone tis-

sue engineering scaffolds [102,110,111] and coatings on biomedi-cal implants [112]. Diba et al. [110,113] fabricated porousforsterite/polycaprolactone (PCL) nanocomposite scaffold via a sol-vent casting/salt leaching technique. Different amounts (0–50 wt.%) of sol–gel derived forsterite nanopowders (25–45 nm par-ticle sizes) were incorporated in a biodegradable polymer–matrix(PCL), and sodium chloride particles with particle sizes of 250–297 lm were used as porogen to obtain a porous structure by saltleaching. The fabricated nanocomposite scaffolds had porosities ofabout 90–92.5%, interconnected pores in the range 100–300 lm,and also many micropores (1–10 lm) were present on the scaffoldwalls. With increasing forsterite nanopowder content from 0 to50 wt.% the porosity and average pore size of the scaffoldsdecreased from 92.6% to 90.94% and from 193 to 98 lm, respec-tively, and pore interconnectivity also decreased. The nanocom-posite scaffolds with more than 40 wt.% forsterite nanopowdercontent lacked a homogeneous dispersion of forsterite nanopow-der and sufficient pore interconnectivity. Moreover, by increasingthe forsterite nanopowder content from 0 to 30 wt.% the compres-sive strength and elastic modulus of the scaffolds improved from0.0024 MPa to 0.3 MPa and from 3.1 MPa to 6.9 MPa, respectively.However, additional increase of forsterite nanopowder contentover 30 wt.% resulted in a decrease of the mechanical propertiesof the scaffolds. In vitro bioactivity of nanocomposite scaffoldswas studied by immersion of samples in SBF up to 28 days. Theresults showed the pure PCL scaffolds were not bioactive, but allthe samples containing forsterite were able to form a Ca–P layeron their surface by immersion in SBF. Interestingly, with increaseof forsterite nanopowder content in the composite scaffolds from10 wt.% to 50 wt.%, Ca:P ratio of the formed layer decreased from1.58 to 1.26 and the surface morphology changed from sphericalparticles (size of around 5–10 lm) with needle like crystallites toa smooth scaly structure with a characteristic dimension of 2 lm.A calcium-deficient hydroxyapatite (CDHA) layer was formed onthe 10 wt.% forsterite nanopowder containing scaffold, whichcould be formed due to incorporation of Mg2+ ions in the Ca2+

ion sites during the layer formation in SBF. Furthermore, an OCPlayer was formed on the 50 wt.% forsterite nanopowder containingscaffold. It was concluded that increasing the amount of forsteritenanopowder in the scaffolds compensates the low wettability ofPCL, improves the scaffold contact/interaction with SBF, andincreases the activity and tendency for transformation of theformed Ca–P layer structure from HAp to OCP. Immersion of thescaffolds in PBS for up to 30 days showed that the degradation rateof the pure PCL scaffold was increased by the incorporation of for-sterite nanopowder. Moreover, analysis of SBF and PBS solutionsshowed that forsterite nanopowder can act as a buffering agent




by releasing alkaline ions (e.g. Mg2+) that compensate the acidifi-cation effect of the products released by the polymer (PCL)degradation. Furthermore, the in vitro cytotoxicity of the nano-composite scaffolds was evaluated using the osteoblastic cell lineSaOS-2 [113]. The results indicated that the nanocompositescaffolds were not cytotoxic. In addition, cell proliferation wasgradually enhanced by forsterite nanopowder containing scaffoldswhen compared with the pure PCL scaffold. Analysis of thenanocomposite conditioned medium revealed an increasingrelease of Mg and Si ions from the scaffolds during immersion,which might stimulate cell proliferation. In addition, cell attach-ment and spreading on the nanocomposite was considerableenhanced when compared to the pure PCL scaffold.

Aligned nanofibrous forsterite-PCL composite membranes werefabricated via an electrospinning method [114]. Addition of10 wt.% forsterite nanopowder into a PCL matrix led to decreaseof fibre diameters from 872 ± 361 nm (for pure PCL fibres) to258 ± 159 nm. Elastic modulus and tensile strength of the mem-branes made of nanocomposite fibres were higher than those ofpure PCL membranes. However, further increase of forsterite con-tent produced agglomerations which led to decreased mechanicalproperties of membranes. Addition of forsterite improved themembranes degradation by increasing their hydrophilicity and pro-duced bioactive membranes able to form apatite on their surfaceupon immersion in SBF. Moreover, pre-osteoblast cells showed sig-nificantly higher attachment, proliferation and mineralization onthe nanocomposite membranes than on pure PCL membranes. Thisstudy suggested that these aligned nanofibrous composite scaffoldsmight be applicable for guided bone regeneration.

Ghomi et al. [102] fabricated porous forsterite scaffolds usingforsterite nanopowders (with crystallite size of 23–35 nm) via agel-casting technique followed by sintering at different tempera-tures (900–1200 �C). The scaffolds had interconnected sphericalpores with pore sizes between 50 and 200 lm, total porosities ofaround 81–86%, and open porosities of around 69–78%. Increaseof sintering temperature from 900 to 1200 �C decreased the poresize, total porosity (by 5%), and open porosity (by 9%) of the scaf-folds. Furthermore, in vitro bioactivity evaluation of the forsteritescaffolds in SBF showed that the forsterite scaffolds were able toform a HAp layer on their surfaces. (The mechanical properties ofthe scaffolds were discussed in Section 2.4.2.) In another study[111], bovine bone derived- HAp scaffolds were coated by forste-rite nanopowders via a dip coating method followed by a heattreatment process. The nanostructured forsterite coated HAp scaf-folds had porosities around 83%, pore sizes about 740 lm, and alsonano-sized pores were present within the forsterite coating. A uni-form nanostructured forsterite coating was formed on the HAp sur-faces without cracks or interfacial delamination. The forsteritegrain size within the coating increased by increasing the sinteringtemperature from �60 nm at 900 �C to �2 lm when sintered at1100 �C. It was shown that the forsterite coating had a significanteffect on the mechanical properties of scaffolds, and by applyingthe nanostructured forsterite coatings the compressive strengthof the porous HAp could be improved from 0.12 to 1.61 MPa. Xieet al. [112] applied sol–gel derived forsterite powder as aplasma-sprayed coating on a Ti–6Al–4V alloy substrate for hardtissue replacement applications. The results showed that the coat-ing was mainly composed of forsterite phase, in addition to a smallamount of glass and MgO phases. The formation of the glassy phasewas reported to be due to the quenching during the plasma-spray-ing process and the formation of the MgO phase could be partiallyoriginated from dissociation of the forsterite phase by the hightemperature plasma spraying process. The coating exhibited sur-face roughness of Ra = 7.86 ± 0.42 lm and bonding strength of41.5 ± 5.3 MPa, that was claimed to be much higher than for previ-ously reported plasma-sprayed HAp coatings. In vitro cytocompat-


ibility studies showed good attachment, spreading, proliferation,and improved cell differentiation of canine bone mesenchymalstem cells (MSCs) when cultured with forsterite coated scaffolds.Many polygonal-shaped cells were observed to have a close adhe-sion with the coating surface; they spread and assumed a morecompact shape with short cellular extensions. The proliferationrate was similar to that of the HAp coating control. A higherdifferentiation level was observed on cells seeded on the forsteritecoating, as the ALP activity after 21 days of culture remained highfor the forsterite coatings, while for HAp coatings it started todecrease.

2.5. Other Mg-containing silicate ceramics

Apart from the silicate materials described in previous subsec-tions, other Mg-containing silicate ceramics, such as merwinite(Ca3Mg(SiO4)2) and monticellite (CaMgSiO4), have also been stud-ied as potential bioactive materials for biomedical applications.In particular, merwinite and monticellite ceramics have recentlydrawn the attention of the biomedical research community, sincethey presented not only mechanical properties comparable tothose of cortical bone [115,116], but also an apatite forming abilityin vitro [115–120]. Furthermore, osteoblasts are reported to adhereand spread both on merwinite and monticellite surfaces [115,118],and the ionic products from both ceramics promoted significantlyosteoblast cell growth [115,116]. When compared to akermaniteceramics, it was observed that the mechanical properties wereimproved, from merwinite to akermanite and monticellite ceram-ics, with the increase of MgO content, while apatite forming abilityin SBF as well as cell proliferation decreased. Additionally, the mor-phological features and proliferation of cells on merwinite ceram-ics were more obvious than for cells seeded on the other twoceramics [121]. Additionally, merwinite is reported to form bioac-tive composites with b-TCP with enhanced mechanical propertiescompared to pure b-TCP ceramics [122]. Recently, the in vivo per-formance of merwinite and HAp bioceramics has been compared inrat femoral defect models [123,124]. Histomorphology and histo-morphometry evaluations indicated that merwinite bioceramicscould induce osteogenesis and showed a higher biodegradationat both early and late stages of implantation. Moreover, at the latestage of implantation, merwinite exhibited a higher rate of newbone formation compared to that of HAp bioceramics.

Finally, it is noteworthy to mention that montmorillonite (Na1/

3(Al5/3Mg1/3)Si4O10(OH)2) layered ceramics present high cationexchange capacity, non-toxicity, good biocompatibility and con-trolled release rate [125–128] and for this reason they have beenextensively studied for various biomedical [129,130] and pharma-ceutical applications [127,131]. However, to the best of theauthors’ knowledge, the in vitro apatite forming ability of theseceramic materials has not been addressed. Furthermore, clinoen-statite and enstatite ceramics (two polymorphic forms of MgSiO3)have been reported to present cell biocompatibility and enhancedmechanical properties and have been proposed as potential mate-rials for bone implants or bioceramic coatings. However, coarsegrained clinoenstatite was not able to form an apatite layer onits surface in SBF and no other study was found reporting the apa-tite forming ability of enstatite [132,133]. Nevertheless, one study[134] has reported that protoenstatite, another polymorphic formof MgSiO3, could precipitate apatite on its surface upon immersionin a pseudo body solution.

2.6. Comparison between silicate ceramics with different Mg contents

Few systematic studies have compared properties of various sil-icate ceramics with different magnesium contents. Wu et al. [37]compared in vitro bioactivity, degradation, and cytocompatibility




of diopside (11.22 wt.% Mg), akermanite (8.92 wt.% Mg), and bre-digite (3.61 wt.% Mg) ceramics. Coprecipitation derived diopsidepowders and sol–gel derived akermanite and bredigite powderswere used to prepare press-sintered ceramic discs. The apatiteforming ability of the ceramic disks was compared by soakingthem in SBF for 7 days. In vitro degradation of the ceramics wasevaluated by immersion of the disks in Tris–HCl buffer solutionsup to 28 days. The results of the study [37] revealed that withthe increase of Mg content from bredigite to diopside, the apa-tite-forming ability and degradation of the ceramics in SBF andTris–HCl solution, respectively, decreased. Moreover cell culturestudies, using osteoblast-like cells from rat calvaria were carriedout to evaluate the effects of ionic products from ceramic powderdissolution on cell proliferation, and the growth of cells seededdirectly onto ceramic disks. The results showed that ionic productsobtained at low concentrations from the three bioceramics (1.25and 12.5 mg/mL) could stimulate osteoblast proliferation. How-ever, at high extract concentrations, bredigite showed a more obvi-ous inhibitory effect on the cell proliferation compared to the twoother ceramics, which might be due to the elevated pH of the cellculture medium originated from high degradation rate of bredigite.It was suggested that Mg content in the bioceramics might affectthe ceramic degradation and apatite formation ability, which inturn influence the release of Si ions that can be responsible forthe stimulatory effect of osteoblasts. Furthermore, it was observedthat with increase of Mg content, i.e. from bredigite to diopside, thecell proliferation level was improved.

Chen et al. [121] studied the effect of MgO content on themechanical properties, in vitro bioactivity and cytocompatibilityof silicate bioceramics. Dense monticellite (25.76 wt.% MgO,1 Mg/Ca mol ratio), akermanite (14.78 wt.% MgO, 0.5 Mg/Ca molratio) and merwinite (12.26 wt.% MgO, 0.33 Mg/Ca mol ratio)ceramics were prepared by uniaxial pressing sol–gel derived pow-ders and by 6 h subsequent sintering at 1480, 1350 and 1400 �C,respectively. Table 6 shows the relative density, average crystallitesizes and mechanical properties of the three ceramics aftersintering.

As can be seen from Table 6 the three ceramics had similar rel-ative densities and average crystal sizes. However, with increase ofMg content and decrease of Mg/Ca molar ratio from monticellite tomerwinite, the mechanical properties improved. It was suggestedthat as the bond energy of Mg–O is higher than that of Ca–O bonds,the improvement of mechanical properties by increasing Mg con-tent could be due to the occupation of Ca atom positions by Mgatoms which might lead to more stable crystal structures. More-over, the different mechanical properties of these three ceramicsmight be related to their different crystal structures (Table 1).14 days immersion of the ceramics in SBF revealed that the apatiteformation rate and the size of the apatite crystals formed on theseceramics decrease with increasing Mg content and Mg/Ca molarratio. Spherical apatite crystals, 50–100 nm in diameter, wereformed on the monticellite ceramics, worm-like apatites with crys-tal size of around 100–300 nm in length and 50 nm in diameterwere formed on the akermanite ceramics, and worm-like apatiteswith crystal size of around 400–800 nm in length and 75 nm indiameter were formed on the merwinite ceramics. The decrease

Table 6Various properties of silicate ceramics with different magnesium contents [121].

Monticellite Akermanite Merwinite

Relative density (%) 92.7 ± 1.4 91.6 ± 3.2 91.3 ± 1.7Average crystal sizes (Å) 605 611 555Bending strength (MPa) 163.9 ± 3.6 141.8 ± 2.3 128.4 ± 4.7Fracture toughness (MPa m1/2) 1.65 ± 0.12 1.53 ± 0.10 1.57 ± 0.17Young’s modulus (GPa) 45.5 ± 4.1 56.2 ± 5.4 49.3 ± 2.3


of apatite forming ability by increasing Mg content in these ceram-ics was suggested to be due to the lower dissolution rate andrelated slower Ca ion release with higher Mg content, which arecaused by the stronger Mg–O bonds formed instead of Ca–O bonds.Furthermore, osteoblast cell-ceramic direct and indirect interac-tions were studied by seeding the cells directly on the ceramicwafers and by using the ceramic extracts, respectively. It wasfound that the Ca, Mg, and Si ions released from akermanite, mon-ticellite and merwinite ceramics present in the 1.25, 12.5 and25 mg/mL ceramic extracts could stimulate cell proliferation, asshown in Table 7.

Merwinite showed the highest stimulatory effect at 12.5 mg/mLafter 7 days of culture followed by akermanite and monticellite. Atextracts concentrations of 100 mg/mL no stimulatory effect wasobserved, whilst even merwinite showed an inhibitory effect. Fromthe analysis of the ion concentrations in the 100 mg/mL merwiniteextract, it was found that, compared with those of the other ceram-ics; the Mg concentration was at its lowest level while Si concen-tration was significantly higher, suggesting that low Mgconcentration and high Si concentration in solution may lead toinhibitory effects on cell proliferation. Cells seeded directly onceramic wafers attached well, and the morphological features weremuch more clear in merwinite, with no significant differencesbetween akermanite and monticellite. It was observed that cellproliferation was enhanced for the three bioceramics, with thehigher stimulatory effect for merwinite (lowest Mg content), fol-lowed by akermanite, and monticellite (highest Mg content). Huet al. [58] studied antibacterial activity of bredigite and akermaniteby preparing extracts of 1, 10, 50 and 100 mg/mL and using E. colias a model bacterium. The results showed better antibacterialactivity for bredigite (lowest Mg content) than akermanite. Thebactericidal percentage for akermanite increased from nearly 20%to 80% for 1 and 100 mg/mL extracts, respectively, while for bre-digite bactericidal percentage increased from 70% to 100% for 1and 10 mg/mL extracts, respectively.

Yamamoto et al. [134] compared apatite forming ability of dif-ferent bioceramics in the CaO–MgO–SiO2 system. Pseudowollas-tonite (CaSiO3), akermanite, diopside, protoenstatite (MgSiO3)and forsterite bioceramics were synthesized via a solid solutionmethod and thereafter sintered and immersed in pseudo bodysolution. The results showed that while pseudowollastonite, diop-side, and protoenstatite were able to precipitate HAp on their sur-faces, akermanite and forsterite did not show any HApprecipitation. Mg release into the solution correlated with theHAp formation results but Ca release did not show any correlationwith HAp formation. It was shown that HAp precipitation abilityincreased with decreasing Mg content. This correlation was sug-gested to be due to a disruptive effect of Mg on HAp crystal forma-tion. The non-sensible effect of Ca elution on HAp formation couldbe due to the low degree of Ca release in this study. Interestingly,the results showed that although akermanite did not contain thehighest amount of MgO compared to the other bioceramics usedin this study, it showed the highest degree of Mg elution. Neverthe-less, it was suggested that since the Mg elution might be beneficialfor in vivo performance of the bioceramics, diopside can be consid-ered as a preferred bioceramic since it can simultaneously releaseMg and form HAp on its surface. In another study [135] diopside,akermanite and pseudowollastonite bioceramics were immersedin calcium free pseudo body fluid and distilled water. Diopsideshowed the highest amount of calcium release compared to thetwo other bioceramics. All three bioceramics were able to formHAp on their surfaces when immersed in pseudo body fluid con-taining 25 mM calcium chloride. Moreover, these materials wereapplied as a paste on HAp pellet surfaces and were able to connectto the HAp pellets through formation of a connected HAp layerbetween the interfaces when immersed in biological fluids.



Table 7Concentrations range of Ca, Mg and Si ions in bioceramic extracts reported as stimulatory and non-stimulatory for cell proliferation [37,121].

Diopside Akermanite Bredigite Monticellite Merwinite

Stimulatory ion concentrationCa (mg/mL) 0.43–0.51 0.42–0.60 0.42–0.44 0.44–0.56 0.45–0.63Si (mg/mL) 0.02–0.24 0.02–0.35 0.08–0.37 0.02–0.35 0.04–0.70Mg (mg/mL) 0.32–0.58 0.32–0.88 0.29–0.33 0.33–0.94 0.32–0.79

Non-stimulatory ion concentrationCa (mg/mL) 0.95 0.78 and 1.07 0.60a 0.93 1.14a

Si (mg/mL) 1.89 2.04 and 1.75 6.98a 1.49 2.78a

Mg (mg/mL) 2.57 2.82 and 2.61 0.29a 2.83 2.27a

a Cell proliferation inhibitory concentrations reported in the literature.


The bioactivity mechanism of the Mg-containing silicate ceram-ics and the discussed effects of their Mg content on the HAp layerformation are similar to those suggested for Mg-containing bioac-tive silicate glasses, as presented in a previous review article [47].Studies have shown a relationship between degradation rate andbioactivity of the ceramics in the CaO–MgO–SiO2 system[37,38,121]. Although a faster degradation of these ceramics canlead to a faster release of Ca2+ ions which can facilitate the forma-tion of HAp layer on their surface [54], Ca does not seem to be anessential compositional element for them to be bioactive. Forste-rite, as a Ca free bioceramic, is a good example to show the impor-tance of degradation over the Ca content in the bioactivity processof these materials. Coarse grained forsterite does not degrade eas-ily and is not bioactive [103]; however, nanostructured forsterite isdegradable and bioactive [95]. The key role of degradation in thebioactivity process of these materials is correlated to the ionexchange rate at the ceramic/media interface which affects the for-mation of a silica-rich layer on the ceramics surface [38,95]. Never-theless, it is important to consider that a high degradation rate ofMg-containing ceramics can lead to high release of Mg ions intothe media which can limit the formation of the HAp layer on theceramic surface [134].

3. Mg-containing bioactive glass-ceramics

3.1. Structure, physical and mechanical properties

Mg-containing amorphous bioactive glasses have beenreviewed elsewhere [47]. This section will cover Mg-containingglass-ceramics. Mg-containing glass-ceramics commonly used forbiomedical applications are traditionally synthesized by annealingof a parent glass. The composition of this glass as well as theannealing schedule play an important role in the crystallizationprocedure. The most commonly used parent glasses are those ofthe composition CaO–MgO–SiO2 [137], while glass-ceramics basedon a SiO2–MgO–Al2O3 glass have attracted the interest of the scien-tific community, as alternative materials for biomedical applica-tions due to their good machinability [137,138], suitablebioactivity [139,140] and favourable mechanical properties[138,139]. A complete overview of reported synthesized Mg-con-taining glass-ceramics, their investigated properties and the keyeffects of Mg for each material, found in the present literaturesearch, is available as a table in the Supplementary Data section.

In the CaO–MgO–SiO2 system, the major crystallinephases formed after annealing are akermanite [70,137,141–162]and diopside [70,137,148,150–153,158,159,161–175], whilemerwinite [136,142–145,149,162], monticellite [164], whitlockite[136,174,175] and enstatite [174] have also been reported to becrystallized as primary and secondary phases by annealing of thesame system. Tulyaganov et al. [137] reported that the crystalliza-tion of a glass in the CaO–MgO–SiO2 system with B2O3, P2O5, Na2O


and CaF2 additives starts at relatively lower temperatures that thoseusually reported in these systems. Furthermore, it is well estab-lished that the addition of P2O5, Na2O and CaF2 plays an importantrole in lowering the melting point of glasses and in enhancing sinter-ability [137,169,176,177]. In addition, in cases of a high amount ofMgO content (>20 wt.%), the addition of Na2O and K2O has beenreported to promote the crystallization of potassium fluorrichterite(KNaCaMg5Si8O22F2), a chain silicate glass-ceramic that has beenproposed as a promising candidate for bone regeneration [178–181]. The formation of diopside and akermanite is reported to affectthe physical and mechanical properties of the glass-ceramic materi-als. Specifically, the density of the glass-ceramics increases withMgO content in the parent glass [182], while the formation of diop-side after annealing reduces the total porosity of the glass-ceramics[137], probably due to the significant difference between the densityof diopside in the glassy (2.75 g/cm3) and crystal (3.27 g/cm3) states[173]. Similarly, the addition of MgO in the parent glass is reportedto affect positively the compressive [182] and bending strength[137] and the microhardness [164] of the glass-ceramics, while Liuand Chou [146] claimed that the formation of akermanite crystalsalong with the crystallization of fibrous and dendritic wollastoniteimprove significantly the flexural strength and fracture toughnessof the glass-ceramic.

Concerning the second group of Mg-containing bioactive glass-ceramics that are based on a parent glass of the system SiO2–MgO–Al2O3, diopside and fluorophlogopite are the main crystallinephases formed after annealing. In particular, diopside is the onlycrystalline phase formed in the case of glass-ceramic materialsfrom parental glass containing less than 9 wt.% of MgO [183,184],while the crystallization of fluorophlogopite, which is responsiblefor the machinability of the glass-ceramic [139,185], is favouredwhen the parental glass contains more than 11.5 wt.% [184–187].However, the addition of amounts of less than 4.3 wt.% of MgOhas been reported to result in glasses with no Mg-based phases[188,189]. The nucleating agent used is reported to affect the kindof Mg-containing crystalline phase formed as well as the crystalmorphology [139,185,187,190]. Specifically, Dittmer and Russel[190] reported that the addition of ZnO and ZrO2 results in thecrystallization of spinel, while according to Molla and Basu [185]varying the F� content of the parent glass has a distinct effect onthe fluorophlogopite crystal morphology. In the latter case, the var-iation of F� content seems to affect the mechanical properties andthe hardness of the glass-ceramics, since they are highly connectedwith the morphology of the fluorophlogopite crystals [185]. Fur-thermore, addition of CaO and P2O5 to the base glass compositionleads to the formation of fluorophlogopite crystals together withapatite crystals [139,187], which results in the decrease of themachinability of the material [139]. However, the needle–likefluorapatite crystallized in the case of fluorophlogopite-basedglass-ceramics containing a high amount of CaO and P2O5 isreported to influence positively their Vickers’ hardness and




fracture toughness [187]. Finally, it is worth mentioning that thecrystallization behaviour of glass-ceramics originating from thecombination of the oxides SiO2, MgO and Al2O3 is drasticallyaffected by the fabrication method [191].

A third group of glass-ceramics containing MgO is the Apatite–Wollastonite (A–W) group introduced by Nakamura et al. in 1985[192], which are synthesized by annealing a parent glass of thecomposition CaO–P2O5–SiO2 with additions of MgO and other oxi-des [193–202]. In the case of these glass-ceramics as well as in thecase of some calcium phosphate glass-ceramics without silica(fourth group), MgO is added in the composition in small amounts(64.6 and 2.24 wt.%, respectively) as a stabilization agent, sinceMg2+ additions are reported to induce and stabilize the crystalliza-tion of phases containing high calcium content [203]. Just as in thecase of glass-ceramics originating from a SiO2–MgO–Al2O3 glass,the small amount of MgO results in the formation of no Mg-basedphases [48,49,203–209]. However, the increase of MgO content inthe CaO–MgO–P2O5–SiO2 composition provides the necessaryamount of MgO for the crystallization of diopside [184], andakermanite [210–212] instead of wollastonite [151,152], while inthe case of CaO–P2O5 composition only minor amounts of Na2-

Mg(PO3)4 have been reported to crystallize [203].

3.2. In vitro bioactivity and degradation

Concerning the bioactivity, it is well documented that thein vitro apatite forming ability of glass-ceramics cannot be directlyevaluated by their oxide composition, since there is no evidencewhether the various elements of the parent glass will remain inthe glassy matrix or they will be incorporated in the ceramic phase.In fact, the composition of the parent glass is shared between thecrystalline phases and a residual glassy matrix, the compositionof which is sometimes difficult to evaluate. In addition, it is likelythe surface reactivity of the glassy phase that is the main factordetermining the bioactivity of glass-ceramics [174]. However, theincrease in MgO content in glass-ceramics derived from SiO2–MgO–CaO glasses is reported to affect negatively their apatiteforming ability [164,170,182] which has been similarly reportedto occur in bioactive silicate glasses [47,213,214]. Indeed glassesand glass-ceramics of the same composition are reported by Oli-veira et al. [174] to attain different bioactive behaviour in vitro.Specifically, it was shown that the precipitated hydroxyapatitelayer was poorly attached to the parent glassy specimen, while itformed a strong bond to the glass-ceramic. It is worth mentioningthat although some studies have evaluated the ion release of theseglass-ceramic materials in SBF, to the authors’ knowledge no previ-ous work has focused on their in vitro degradability in other rele-vant solutions (e.g. PBS or DMEM).

The glass-ceramics of the group SiO2–MgO–Al2O3 present noapatite forming ability in vitro [139], while very few studies reportthe degradability of the materials. The introduction of CaO andP2O5 has been reported to induce the bioactive response in vitro[139,140], while their bioactivity is improved with the increasein F� content [185,187]. However, Molla and Basu [185] reportedno weight loss or gain within the accuracy of the digital balance(three digit resolution) even for bioactive samples.

Concerning glass-ceramics with low amount of MgO (A–W andcalcium phosphate glass-ceramics), MgO has not been reported toaffect their in vitro apatite forming ability. However, glass-ceram-ics of the A–W composition [215] with increased MgO content anddiopside and althausite as their main Mg-based crystalline phasespresented limited in vitro apatite forming ability [151,152]. In sup-port to these results, Ma et al. [210] reported that with increasingMgO concentration in glass-ceramics in the system CaO–MgO–P2O5–SiO2, the material degradability gradually decreased andthe formation of apatite was delayed. This is likely the result of


the depletion of the residual glass in Ca ions because of crystalliza-tion of diopside and apatite during heat treatment [194]. The crys-tallization of these two phases provides a chemically stabilizedmaterial, thus hindering Ca2+ exchange with H3O+ of SBF and sub-sequent formation of silanol groups, which explains the low bioac-tivity of the glass-ceramic [151]. However, the in vitro bioactivityof these glass-ceramics can be significantly improved by the chem-ical modification of their surface with HCl solution [151,152,206].The HCl attacks the silicate phases and the residual glass, resultingin the formation of a silica-gel layer and consequently leading tofaster mineralization kinetics compared to the untreated glass-ceramic material [151,152,206]. Concerning the in vitro degrada-tion of this family of glass-ceramic materials, they are reportedto be degradable up to the composition containing 5 wt.% MgO[208], while there are no reports regarding the degradation ofthe compositions with higher MgO content.

3.3. Biological properties

There are limited references dealing with the biologicalresponse of glass-ceramics in the system CaO–MgO–SiO2. How-ever, glass-ceramics of this system had not shown cytotoxic effects,at least within the investigated compositional range (4.6–12.8 wt.%of MgO), in the presence of osteoblasts [142,143,154] and fibro-blasts [158]. Furthermore, the glass-ceramics were reported notto hinder cell attachment and proliferation [142,143], which canbe probably attributed to the increased surface roughness of theglass-ceramics due to the crystallization of Mg-containing phases,including diopside, akermanite and merwinite, that favours celladhesion [154]. Additionally, the in vivo performance of anakermanite containing glass-ceramic has been evaluated byimplantation in New Zealand rabbits and the histological resultssuggested that the glass-ceramic allowed bone growth over its sur-face, by means of mesenchymal cells recruitment from the sur-rounding bone, and their differentiation into bone forming cells[153]. The bone covering area was significantly higher in rabbitswith the glass-ceramic implant when compared to control animalswithout the implant. In a different study, glass-ceramics of thegroup SiO2–MgO–Al2O3 were proved to be biocompatible, eventhough they exhibited no apatite forming ability in vitro [139].However, the addition of B2O3 had a significant influence on cyto-compatibility, cell adhesion and cell attachment properties, whilethe increase of F� content decreased cell viability [186]. Further-more, in vivo studies of Holland et al. [140] and Vogel et al. [139]revealed a direct bonding between a SiO2–MgO–Al2O3–P2O5

glass-ceramic and the spongy metaphyseal bone and the adjacentcortical bone of female guinea pigs.

Concerning the in vivo biological response of A–W glass-ceram-ics, they are well accepted upon implantation in the tibia or femurof animal models, including rats, rabbits and German shepherds[192,208], while they permit new bone to be formed directly onthe implant surfaces. Furthermore, Abiraman et al. [207] demon-strated that sol–gel derived A–W glass-ceramics enhanced theneo-osteogenesis process, and the healing of the defect was com-plete within 6 weeks after implantation. Finally, it is worth men-tioning that Al2O3 doped A–W glass ceramics have been reportedto provoke the osteogenic differentiation of marrow stromal stemcells in vitro [204].

4. Mg-containing bioactive inorganic composites

Several bioactive ceramic/ceramic and ceramic/glass inorganiccomposites have been developed by mixing and press-sinteringbioceramic and glass powders. Due to the attractive mechanicalproperties of Mg-containing silicate ceramics, diopside or forsterite




powders have been used to improve the mechanical properties ofcrystalline calcium phosphate ceramics and silicate glasses. In1995 Nonami and Satoh [216] developed composites of HAp withelongated diopside grains via press-sintering of diopside and HAppowders. The results showed that the sample containing 60 vol.%of diopside and sintered at 1280 �C for 2 h exhibited the optimummechanical properties with 265 MPa bending strength,3.2 MPa m1/2 fracture toughness, and 141 GPa elastic modulus,which were 2 or 3 times higher than those of similar HAp samples.Furthermore, cell culture studies using osteoblastic (KUSA) andfibroblastic (CCL1) cells indicated the cytocompatibility of thesecomposites. Ni et al. [103] fabricated wollastonite–forsterite com-posites with 0–100 wt.% forsterite content. It was observed that thesintering process used (1200–1300 �C for 5 h) led to a reactionbetween the forsterite and wollastonite powders, and the obtainedcomposites contained a mixture of wollastonite, forsterite, diop-side, and akermanite phases. With increasing the initial forsteritecontent in the composites, their mechanical properties steadilyimproved (bending strength of 168.4 MPa and Young‘s modulusof 22.3 GPa were achieved for the sample with 70 wt.% forsterite),their dissolution rate in Tris–HCl buffer solution decreased, andtheir apatite formation in SBF also decreased. Cell culture studiesshowed that the samples with initial 30 and 50 wt.% of forsteriteinduced higher degree of osteoblastic cells proliferation. Carrode-guas et al. [217] studied the degradation and apatite-forming abil-ity of 40/60 wt.% b-TCP/diopside composite ceramics (sintered at1250 �C for 4 h) in SBF. The results of this study indicated that dur-ing immersion, the b-TCP grains had a higher degradation ratecompared to diopside grains, which led to the formation of a por-ous surface enriched with diopside that after 7 days of immersionwas covered by a CDHA layer (Ca/P = 1.62). Sainz et al. [218] usedstoichiometric eutectic composition of wollastonite (36.77 wt.%)and diopside (63.23 wt.%) powders to obtain a composite bioce-ramic via sintering at 1250–1350 �C for 4 h. Fig. 5 shows the phaseequilibrium diagram of the wollastonite–diopside system.

With sintering of the samples up to 1300 �C, porous structureswere obtained and the present phases corresponded to the initialpowders indicating the occurrence of a solid state sintering processin this temperature range. However, sintering at 1350 �C led to theformation of a dense sample containing stable low temperaturepolymorph wollastonite-2M (Ca0.83Mg0.17SiO3) and diopsidephases. Zhang et al. [219] produced Al2O3/diopside(0, 1, 20 wt.%)ceramic composites via hot-pressing of the powders at 1450 �Cfor 30 min. It was observed that addition of 1 wt.% diopsideimproved the mechanical properties; samples exhibited Vickershardness of 18.6 GPa, flexural strength of 427 MPa and fracture

Fig. 5. Phase equilibrium diagram of the wollastonite–diopside (CaO–MgO–SiO2)system (reproduced from Ref. [218] by permission from Elsevier).


toughness of 4.3 MPa m1/2. However, further increase of diopsidecontent (30 wt.%) led to inferior mechanical properties. After5 days of immersion in SBF, a HAp layer was formed on the surfaceof the samples containing diopside. On the contrary, pure Al2O3 didnot induce an HAp layer on its surface. In another relevant studyZhang et al. [220] introduced alumina and diopside in a HAp matrixvia uniaxial hot pressing of ceramic powders in order to obtain acomposite material with improved mechanical properties. It wasobserved that the composite containing 2 wt.% diopside and30 wt.% alumina had improved mechanical properties and alsowas able to form an apatite layer in SBF. In a similar study [221],addition of 10 wt.% diopside into a HAp matrix led to a bioactivecomposite with improved mechanical properties.

Fig. 6a shows a CDHA layer formed on a CaSiO3–CaMg(SiO3)2

composite surface after 3 weeks of immersion in SBF [218]. Fur-thermore, Fig. 6b shows the higher degradation of wollastonitephase surface compared to the diopside phase by immersion inSBF (Fig. 6b). In addition, it was reported that with increasing sin-tering temperature, the reactivity of the composite samples in SBFdecreased.

Yazdanpanah et al. [222] incorporated up to 30 wt.% of forste-rite nanopowder (10–50 nm) into a bioactive glassy matrix(64 mol.% SiO2, 31 mol.% CaO, and 5 mol.% P2O5) by pressing andsintering of mixed powders at 1000 �C for 2 h. X-ray diffraction(XRD) analysis of the prepared composites showed formation of apseudowollastonite phase after sintering. The results showed thatincreasing the forsterite nanopowder content improved the frac-ture toughness of the composites up to 0.22 MPa m1/2, butdecreased the elastic modulus and yield stress. A 14 days immer-sion period in SBF showed that all samples were able to induce aHAp layer on their surface and incorporation of forsterite nano-powder did not have a significant effect on the in vitro apatiteforming ability of the samples. Furthermore, it was reported thatwith increase of forsterite content, the Ca/P ratio of the formedapatite layer increased from 1.69 up to 3.21. However, these resultsare in contrast with other reports regarding the effect of Mg con-tent on the Ca/P ratio of the apatite layer formed on Mg-containingmaterials. It is expected that the released Mg ions incorporate intothe Ca sites of the apatite layer, and the Ca/P ratio decreases [47].

Few studies have developed glass-ceramic composites by press-sintering of the mixture of crystalline bioceramic and Mg-contain-ing amorphous glass powders. Pogrebenkov et al. [223] fabricatedapatite/glass-ceramic composites with 0–100 wt.% apatite content.The glass-ceramic powders were obtained by melting natural diop-side minerals at 1425 �C for 1 h, and the quenched glass-ceramicpowders contained 75% of diopside phase and 25% glassy phasein the 11 wt.% CaO–14 wt.% Al2O3–75 wt.% SiO2 system. Duringmelting, Al2O3 was added in order to reduce the crystallization ofdiopside glass. Different sintering schedules (950–1200 �C) of themixed powders produced different combination of diopside, HApand amorphous phases within the final composite. In general, withincreasing sintering temperature up to 1150 �C, the porosity of thesamples decreased and their bending strength improved. However,at 1150 �C signs of overburning were observed in some samplesand at 1200 �C fusion of phases in samples which contained 10–80 wt.% of HAp powder occured. The highest bending strengthbelonged to a pure glass sample and with increasing HAp contentthe bending strength decreased to 5–10 MPa. In a similar study,Ashuri et al. [224] press-sintered a mixture of sol–gel derivedMg-containing bioactive glass (64 SiO2–26 CaO–5 MgO–5 ZnO,mol.%) and HAp powders (5–30 wt.%) at 1100 �C for 1 h. Sinteringtemperatures higher than 800 �C led to formation of a b-TCP phasevia thermal decomposition of HA. Compression tests indicated thatwith increasing HAp content up to 20 wt.% the compressivestrength and work of fracture of the samples increased up toaround 18 MPa and 2.0 kJ m�2, respectively. However, further



Fig. 6. SEM micrographs of CaSiO3–CaMg(SiO3)2 sample sintered at 1350 �C after 3 weeks immersion in SBF showing (a) HAp particles formed on the surface (b) polishedcross-section with partial dissolution of the wollastonite region (reproduced from Ref. [218] by permission from Elsevier).

Fig. 7. Osteoblastic cells adhered and spread on a Mg-containing bioactive glass/CaCO3 composite surface (reproduced from Ref. [225] with permission from TransTech Publications Ltd.).


increase of HAp content had a negative effect on the mechanicalproperties. After 2 weeks immersion, the sample containing20 wt.% bioactive glass showed formation of an apatite layer onits surface and release of Si ions occured. However, due todegradation of present of the b-TCP phase during SBF immersion,the compressive strength decreased up to 65%. Furthermore, cellcultures studies using osteoblastic SaOS-2 cells confirmed thecytocompatibility of the composite. Similar results regarding themechanical properties of the composites had been previouslyreported by Lin et al. [225]. The authors reported that samplesprepared by mixing 20 wt.% Mg-containing melt-derived glass(44.4 CaO–34.2 SiO2–16.3 P2O5–4.6 MgO–0.5 CaF2, wt.%) andCaSiO3 powders sintered at 1100 �C for 5 h, showed favourablemechanical properties with values of 171.57 MPa for the bendingstrength and 51.52 GPa for the elastic modulus, which wasreported to be due to the positive effect of addition of bioactiveglass. In addition, osteoblastic cells showed enhanced proliferationand good adhesion (Fig. 7) on the surface of the composite sample,which could be due to the release of Mg and F ions from the glass[225].

It is worth mentioning that the sol–gel technique has beenreported as an effective method for development of hydroxyapa-tite–forsterite–bioactive glass [226–228] and fluorapatite–forste-rite [229] composite nanopowders and nanostructured coatings.Furthermore, gel-casting technique has been successfully utilizedfor development of hydroxyapatite–forsterite–bioactive glassfoams as scaffolds for bone tissue engineering applications [230].In general, the combination of different bioceramic phases pro-vides the opportunity to benefit from diverse properties of eachphase, and especially the superior mechanical properties of Mg-containing phases [231].

5. Conclusions, outlook and future work

Significant research has been performed on calcium phos-phates and Ca-containing silicates. However, these materials donot provide the entire requirements as an ideal material for bonetissue replacement or orthopedic applications. For example, theload bearing applications of these materials have been alwaysa challenging issue, and the lack of adequate vascularizationhas limited the applications of many of these bioceramics. Onthe other hand, crystalline and partially-crystalline Mg-contain-ing silicate bioceramics have been shown to be bioactive andexhibit various advantages for orthopedic applications and boneregeneration. The crystalline forms of these silicates show supe-rior mechanical properties compared to other typical bioceramicssuch as calcium phosphates. Furthermore, the release of Mg andSi ions from these materials in many cases was reported to havepositive effect on cell proliferation, differentiation and adhesion.


Moreover, interestingly, akermanite and bredigite, as well asdiopside in a smaller extent, have been proved to stimulateosteogenesis and angiogenesis in vitro and in vivo. In addition,the combination of these bioceramics with other types of inor-ganic and organic biomaterials has been useful for obtainingcomposite materials with optimized properties. In particular,the superior mechanical properties of Mg-containing silicatephases can be combined with the high bioactivity of bioactiveglasses or with the high degradability of biopolymers for specificapplications. The studies have shown that although some phaseswith higher Mg content, such as forsterite, are cytocompatibleand exhibit good mechanical properties, they are not able toinduce the formation of HAp on their surface and have extre-mely low degradation rate in biological fluids. However, thenanostructured form of forsterite is bioactive and has anenhanced degradation rate, which suggests that the nano-pro-cessing of other Mg-containing phases (e.g. clinoenstatite) mightbe useful for production of materials with superior properties forbiomedical applications. Finally, considering the reviewed studieson Mg-containing ceramics and glass-ceramics, these materialsseem to be promising candidates for development of variousproducts for biomedical applications, ranging from tissue engi-neering scaffolds to coatings for biomedical implants. Dependingon the application and its specific conditions, one of these sys-tems or their combinations can be selected to produce promisingbiomedical devices. For example, for load bearing applications,forsterite can be preferred due to its superior mechanicalproperties.




Acknowledgement

Dr. Ourania-Menti Goudouri acknowledges financial support bythe Marie Curie EU Fellowship ‘‘BIODENTTISSUE’’.

Appendix A. Supplementary material

Supplementary data associated with this article can be found, inthe online version, at http://dx.doi.org/10.1016/j.cossms.2014.02.004.v

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