FIELD OF THE INVENTION

The present invention relates to a biocompatible ceramic material and in particular to a biocompatible calcium silicate-based material. In one embodiment, the invention has been developed for use in tissue regeneration, especially bone tissue. In other embodiments the invention has been developed as a suitable coating to improve the long-term stability of prior art implantable medical devices. In another embodiment the invention is suitable for use in drug delivery or for skeletal tissue regeneration. However, it will be appreciated that the invention is not limited to these particular fields of use.

BACKGROUND OF THE INVENTION The following discussion of the prior art is provided to place the invention in an appropriate technical context and enable the advantages of it to be more fully understood. It should be appreciated, however, that any discussion of the prior art throughout the specification should not be considered as an express or implied admission that such prior art is widely known or forms part of common general knowledge in the field. Bone, as a living tissue, has the ability to heal itself, however in some cases damage to the bone from whatever cause is too severe to allow natural healing to take place, and so a bone graft is required to stimulate regeneration. There are three main types of bone grafts: autografts, allografts and synthetic grafts. Significant research is being conducted in the field of synthetic grafts as bone substitutes since synthetic grafts can ameliorate many of the problems associated with autografts and allografts, such as limited supply, donor site pain, and immunogenicity issues.

In the case of advanced degenerative bone disease, joint replacement therapy remains the only treatment available for relieving the pain and suffering. However, the technologies available in this area of orthopaedics are far from satisfactory. For example, Australians require more than sixty-thousand hip and knee replacement operations annually, a rate that has been estimated to be increasing by some 10% per annum, and a staggering 25% of which are revisions of failed implants (Graves, S.E., et al., The Australian Orthopaedic Association National Joint Replacement Registry. Med. J. Aust., 2004 180(5 Suppl): S31-34). Further complications arise in situations where bone stock is compromised, or where initial implant stability is questionable ( e.g ., elderly patients, post-traumatic injuries, or in revision operations), in which cases short- and long-terms clinical results are typically inferior. The increases in life expectancy, and in the number of younger patients requiring implants, highlights the need for greater implant longevity and has driven biomedical research to develop novel micro-engineered surfaces to anchor the cementless prosthesis directly to the living bone through osteo-integration, thereby attempting to provide a stable interface strong enough to support life-long functional loading. It is clear that there is a serious problem with the longevity of current orthopaedic devices; a problem that is anticipated to only increase with the increasing demand from the aging population requiring such treatments. It is also clear that any improvement that could be made to increase the performance of these orthopaedic devices would be welcomed, not only by the orthopaedic community but also by the patients themselves.

3D scaffolds that promote the migration, proliferation and differentiation of bone and endothelial cells are becoming increasingly important in not only orthopaedic but also maxillofacial surgery. An ideal bone replacement material should support bone formation and vascularisation, show minimal fibrotic reaction and serve as a temporary biomaterial for bone remodelling. They should also degrade in a controlled fashion into non-toxic products that the body can metabolise or excrete via normal physiological mechanisms (Yaszemski, et al., Biomaterials , 1996 17: 175-185). Scaffolds need to be mechanically strong and matched with a similar modulus of elasticity to that of bone in order to prevent stress shielding as well as maintaining adequate toughness to prevent fatigue fracture under cyclic loading. At present, there are no successful strategies available for bone tissue regeneration and resurfacing arthritic joints with articular cartilage. The lack of cartilage reparative response creates a great demand for new modalities that promote tissue regeneration.

Over the last century, various ceramics have been investigated for the purpose of encouraging or stimulating bone growth and as scaffolds. For example, in the 1880s calcium sulfate (plaster of Paris) was utilised. However, calcium sulfate displays a relatively low bioactivity and a relatively high rate of degradation (Tay, et al., Orthop. Clin. North Am., 1999 30: 615-23). In the 1950s hydroxyapatite was utilised, but it suffers from a relatively low degradation rate and poor mechanical properties (Wiltfang J. et al J. Biomed. Mater. Res. 2002 63: 115-21). In the 1970s Bioglass ^® was developed. However, this material it is relatively hard to handle due to its inherent brittleness and has a relatively low bending strength (Cordioli G., Clin. Oral Implants Res. 2001 13: 655-665). In the 1990s calcium silicate ceramics began being used for stimulating bone growth. They are regarded as potential bioactive materials and their degradation products do not incite an inflammatory reaction. However, drawbacks exist with these materials that compromise their physical and biological properties including their a.) inability to combine the required mechanical properties with open porosity b.) poor mechanical strength making them unsuitable for load-bearing applications; and c.) poor chemical instability (high degradation rate) leading to a highly alkaline condition in the surrounding environment which is detrimental to cell viability and limits their long-term biological applications.

Whilst other more recent ceramics such as HAp, Bioverit ^® , Ceraverit ^® and other calcium silicates have been found to bond to living bone and meet wide clinical applications, i.e., good bioactivity, they cannot be used in highly loaded areas, such as the cortical bone found in, for example, legs, due to the relative brittleness of these materials. Thus, the materials possess good bioactivity, but lack full biodegradability after implantation and their mechanical strength is compromised (Hench L.L., J Am Ceram. Soc. 1998 81: 1705-28). They are too brittle and fracture frequently. For at least this reason such materials typically find their use limited to coatings on metallic implants. It is an object of the present invention to overcome or ameliorate at least one of the disadvantages of the above mentioned prior art, or to provide a useful alternative.

It is an object of an especially preferred form of the present invention to provide a biocompatible ceramic material that finds particular use as a replacement bone scaffold which shows good cell differentiation and bone formation at the tissue-implant interface, thereby avoiding or minimising interface fracture, and/or having similar mechanical properties as the replaced tissue, thereby avoiding or minimising issues such as stress shielding.

SUMMARY OF THE INVENTION The clinical success of implants depends on two main factors: (i) good cell differentiation and bone formation at the tissue-implant interface to avoid interface fracture, and (ii) matching the mechanical properties between the implant and the replaced tissue to avoid problems such as stress shielding. The present Applicant has developed an improved implantable ceramic material which substantially addresses these factors, as disclosed herein.

The present Applicant has previously developed a biocompatible ceramic named Baghdadite (see international PCT publication WO 2009/052583), which provides mechanical properties that are approximately equivalent to those of bone. A modified form of Baghdadite has now been generated that, in some cases, displays improved biochemical and mechanical properties over Baghdadite (as described in WO 2009/052583). The modified form of Baghdadite disclosed herein provides a bone graft substitute that has surprisingly excellent strength and is particularly suited for use in load bearing applications, and which reproduces or exceeds the bone regenerative and integrative capability of young adult bone grafts. In particular, the present invention relates to biocompatible ceramic materials that may be considered as Mg- and/or Fe- doped forms of Baghdadite. Doping as used herein refers to the incorporation of specific species of ions or atoms into a host lattice core structure to produce a hybrid material with new and useful properties. In the context of the invention described herein, the host lattice core structure is Baghdadite. As used herein, the dopant is Mg and/or Fe. Doping can influence the size and shape of the lattice structure, which is reflected in the transmission X-ray diffraction (XRD) pattern and lattice parameters. The XRD pattern of Mg- and/or Fe-doped forms of Baghdadite exhibit shifted peak lines of intensity compared with undoped Baghdadite. The a, b, c and b lattice parameters of Mg- and/or Fe-doped forms of Baghdadite shift relative to those of undoped Baghdadite. Doping may also influence other properties including modulus and strength. By way of overview, the results shown below indicate that, in some embodiments, the modulus and strength of a Baghdadite material doped with small amounts of magnesium (so-called “Mg-doped Baghdadite”) are up to around 17% and 35%, respectively, higher than “un- doped” Baghdadite disclosed in WO 2009/052583, while maintaining at least the same levels of bioactivity. Additionally, the cell proliferation data shows up to a surprising approximate 10% increase for “Mg-doped Baghdadite” compared to “un-doped” Baghdadite.

The present invention provides biocompatible ceramic materials that provide a bone graft substitute for defect repair that are able to fulfil the currently unmet medical need for bone graft substitutes that are strong enough for use in load bearing applications, and which reproduce or exceed the bone regenerative and integrative capability of young adult bone grafts. While the prior art shows that doping of certain elements in calcium silicates is useful, the effectiveness of adding these elements to Baghdadite is not obvious. Baghdadite itself is a Zr-modified calcium silicate, and adding further elements to this material has only now been surprisingly found to result in a synergetic effect between zirconia and the doped elements (Mg and/or Fe) that appears to lead to a completely different chemical structure, and bioactivity.

Moreover, as discussed in the previous literature, doping of elements results in either improvement of bioactivity, or mechanical properties, not both. Simultaneous improvement of these properties has been rarely reported and is not predictable. For example, Liu et al. (“Novel tricalcium silicate/magnesium phosphate composite bone cement having high compressive strength, in vitro bioactivity and cytocompatibility” Acta biomaterialia, 2015 21: 217-227) showed that a multiphasic tricalcium silicate/magnesium phosphate (C3S/MPC) material is stiffer and stronger than both tricalcium silicates and magnesium phosphates. However, this mechanical improvement comes at the expense of the bioactivity of the material at most concentrations of ionic extracts. Roohani et al. (“Design and fabrication of 3D printed scaffolds with a mechanical strength comparable to cortical bone to repair large bone defects” Scientific reports , 2016 6: 19468) modified Hardystonite and Gahnite by preparing a complex chemical structure that resulted in exceptional improvement in mechanical properties, but decreased osteo-conductivity (as observed from the fibrous tissue formation on the surface of implants).

Additionally, calcium phosphates are different to calcium silicates, and there is no reason to suspect that doping of elements into calcium phosphates to provide a beneficial effect will necessarily produce the same beneficial effect in calcium silicates. It is well known in the art that ceramics do not always necessarily display biocompatible properties. For example, it has been shown that doping Ti into calcium silicates increases the stability of the material, making it useful for coating applications where a reduced degradation rate is desired. However, this improvement is not necessarily accompanied by improved bio-integration. Furthermore, other ceramics have also been shown to be both poorly biocompatible and not able to readily support hydroxyapatite (HA) formation (see Salama et al, Ceram. Int. 2006 32: 357-364). Additionally, the addition of dopants do not always have expected effects on the final material properties, both in terms of mechanical properties, and biocompatibility, as discussed above. Accordingly, the Applicant asserts that the biocompatible and/or mechanical properties of the biocompatible ceramic materials of the present invention are novel and surprising, and the formation of an HA layer is not necessarily expected.

In a first aspect, the present invention provides a material comprising doped Baghdadite, wherein the doped Baghdadite is selected from Mg-doped Baghdadite, Fe- doped Baghdadite, and [Mg, Fe] -doped Baghdadite. Preferably the doped Baghdadite is selected from Mg-doped Baghdadite and Fe-doped Baghdadite. In one preferred embodiment, the doped Baghdadite is Mg-doped Baghdadite. In another preferred embodiment, the doped Baghdadite is Fe-doped Baghdadite. In one embodiment, the doped Baghdadite has a formula:

In one embodiment, the doped Baghdadite has a formula of Formula A, . In Formula A, 0 < x < 1, e.g., 0 < x < 0.5, or 0.1 < x < 0.5, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x < 0.9, or 0.5 < x < 1, or 0.75 < x < 1, e.g., x may be 0.01, 0.02, 0.03, 0.04,

0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, or 0.95. In some embodiments, x in Formula A may be any one of these values or selected from one or more of these values. In Formula A, any values of y and z may be taken such that y + z = x, e.g., y may be 0, 0.01, 0.02, 0.03, 0.04, 0.05,

0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, or 0.99, and similarly, z may be 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, or 0.99. In some embodiments, y and/or z in Formula A may be any one of these values or selected from one or more of these values. Preferably, either y = 0 or z = 0 in Formula A.

In one embodiment of Formula A, z = 0, thereby producing a material of Formula I: In Formula I, 0 < x < 1, e.g., 0 < x < 0.5, or 0.1 < x < 0.5, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x < 0.9, or 0.5 < x < 1, or 0.75 < x < 1, e.g., x may be 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1,

0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, or 0.95. In a preferred embodiment of Formula I, z = 0 and x = 0.05, 0.1, 0.2, 0.3 or 0.5, preferably z = 0 and x = 0.1, 0.2, 0.3 or 0.5. In a preferred embodiment of Formula I, z = 0 and x = 0.1. In another embodiment of Formula A, y = 0, thereby producing a material of Formula la: In Formula la, 0 < x < 1, e.g., 0 < x < 0.5, or 0.1 < x < 0.5, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x < 0.9, or 0.5 < x

< 1, or 0.75 < x < 1, e.g., x may be 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9,

0.95, or 0.99. In a preferred embodiment of Formula la, y = 0 and x = 0.1, 0.2, or 0.5.

In another embodiment, the doped Baghdadite has a formula of Formula B, Ca - x ₎ ([Mg(y ₎ Fe(z ₎ ]å(y,z)=x)ZrSi209 wherein 0 < x < 3. In Formula B, 0 < x < 3, e.g., 0 < x < 0.5, or 0.1 < x < 0.5, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x < 0.9, or 0.5 < x < 1, or 0.75 < x < 1, or 1 < x < 1.5, or 1 < x < 2, or 1.5 < x < 3, or 1.5 < x

< 2, or 1.75 < x < 2.5, or 2 < x < 3, or 2.5 < x < 3, or 1 < x < 2.5, or 0.5 < x < 3, e.g., x may be 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, 1.0, 1.25, 1.5, 1.75, 2.0, 2.25, 2.5, 2.75, 2.95, or 2.99. In some embodiments, x in Formula B may be any one of these values or selected from one or more of these values. In Formula B, any values of y and z may be taken such that y + z = x, e.g., y may be 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, 1.0, 1.25, 1.5, 1.75, 2.0, 2.25, 2.5, 2.75, 2.95, or 2.99 and similarly, z may be 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, 1.0, 1.25,

1.5, 1.75, 2.0, 2.25, 2.5, 2.75, 2.95, or 2.99. In some embodiments, y and/or z in Formula B may be any one of these values or selected from one or more of these values. Preferably, either y = 0 or z = 0 in Formula B.

In one embodiment of Formula B, z = 0, thereby producing a material of Formula II: In Formula II, 0 < x < 3, e.g., 0 < x < 0.5, or 0.1 < x < 0.5, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x < 0.9, or 0.5 < x < 1, or 0.75 < x < 1, or 1 < x < 1.5, or 1 < x < 2, or 1.5 < x < 3, or 1.5 < x < 2, or 1.75 < x < 2.5, or 2 < x < 3, or 2.5 < x < 3, or 1 < x < 2.5, or 0.5 < x < 3, e.g., x may be 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, 1.0, 1.25, 1.5, 1.75, 2.0, 2.25, 2.5, 2.75, 2.95, or 2.99.

In a preferred embodiment of Formula II, z = 0 and x = 0.05, 0.1, 0.2 or 0.5. In a preferred embodiment of Formula II, z = 0 and x = 0.1, 0.2 or 0.5. In a preferred embodiment of Formula II, z = 0 and x = 0.1 or 0.5. In another embodiment of Formula B, y = 0, thereby producing a material of Formula III: In Formula III, 0 < x < 3, e.g., 0 < x < 0.5, or 0.1 < x

< 0.5, or 0.05 < x < 0.2, or 0.25 < x < 0.75, or 0.4 < x < 0.6, or 0.1 < x < 0.75, or 0.5 < x

< 0.9, or 0.5 < x < 1, or 0.75 < x < 1, or 1 < x < 1.5, or 1 < x < 2, or 1.5 < x < 3, or 1.5 < x < 2, or 1.75 < x < 2.5, or 2 < x < 3, or 2.5 < x < 3, or 1 < x < 2.5, or 0.5 < x < 3, e.g., x may be 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, 0.5, 0.55, 0.6, 0.65, 0.7, 0.75, 0.8, 0.85, 0.9, 0.95, 1.0, 1.25, 1.5, 1.75, 2.0, 2.25, 2.5, 2.75, 2.95, or 2.99. In a preferred embodiment of Formula III, y = 0 and x = 0.05, x = 0.05, 0.1, 0.2 or 0.5. In a preferred embodiment of Formula III, y = 0 and x = 0.1 or 0.2.

In one embodiment, wherein the material is Mg-doped Baghdadite, the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source comprising the following 20 diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source having 20 diffraction angles selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. In some embodiments, the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source including at least one diffraction angle 20 plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 27.581, 29.539, 29.959, 31.167,

36.022, and 37.023 degrees. Preferably, the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source including at least one diffraction angle 20 selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. In one embodiment, wherein the material is Mg-doped Baghdadite of Formula I: , more preferably wherein x = 0.1, the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source having 20 diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source having 20 diffraction angles selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. Preferably, the material does not include Merwinite, un-doped Baghdadite or a combination thereof. More preferably, the material does not have a transmission X-ray diffraction pattern characteristic of Merwinite, un-doped Baghdadite or a combination thereof. Preferably, the material is monophasic.

In another embodiment, wherein the material is Mg-doped Baghdadite of Formula I: wherein x > 0.2, more preferably wherein x = 0.2, 0.3 or 0.5, the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source having 2θ diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a copper X-ray source having 20 diffraction angles selected from the group comprising: 27.581, 29.539, 29.959, 31.167, 36.022, and 37.023 degrees. The material may include Merwinite, more particularly the material may have a transmission X-ray diffraction pattern characteristic of Merwinite. Preferably, the material does not include un-doped Baghdadite, more preferably the material does not have a transmission X-ray diffraction pattern characteristic of un-doped Baghdadite. In one embodiment, wherein the material is Fe-doped Baghdadite, the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 20 diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 20 diffraction angles selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. In some embodiments, the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source including at least one diffraction angle 20 plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source including at least one diffraction angle 20 selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees.

In one embodiment, wherein the material is Fe-doped Baghdadite of Formula la: Ca ₃ Fe(y=x ₎ Zr(i-x ₎ Si20(9-x ₎ , wherein x < 0.2, more preferably wherein x = 0.05 or 0.1, the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 20 diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 2Q diffraction angles selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. Preferably, the material does not include calcium iron silicate, un-doped Baghdadite or a combination thereof. More preferably, the material does not have a transmission X-ray diffraction pattern characteristic of calcium iron silicate, un-doped Baghdadite or a combination thereof. Preferably, the material is monophasic.

In another embodiment, wherein the material is Fe-doped Baghdadite of Formula la: , wherein x > 0.2, more preferably wherein x = 0.2, the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 2θ diffraction angles plus or minus 0.01 - 0.03 degrees 20, selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. Preferably the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 20 diffraction angles selected from the group comprising: 12.541, 13.534, 13.998, 14.222, 22.414, 22.787, and 24.228 degrees. The material may include calcium iron silicate, more particularly the material may have a transmission X-ray diffraction pattern characteristic of calcium iron silicate. Preferably, the material does not include un-doped Baghdadite, more preferably the material does not have a transmission X-ray diffraction pattern characteristic of un- doped Baghdadite. In one embodiment, wherein the material is Mg-doped Baghdadite, the material may have a transmission X-ray diffraction pattern obtained from a copper source having 2Q diffraction angles of undoped Baghdadite, wherein the 2Q diffraction angles are shifted at least plus or minus 0.02 degrees 2Q, preferably 0.02 - 1 degrees. More preferably, the 2Q diffraction angles are shifted at least minus 0.02 degrees 2Q, preferably minus 0.02 - 1 degrees.

In one embodiment, wherein the material is Fe-doped Baghdadite, the material may have a transmission X-ray diffraction pattern obtained from a molybdenum X-ray source having 2Q diffraction angles of undoped Baghdadite, wherein the 2Q diffraction angles are shifted at least plus or minus 0.05 degrees 2Q, preferably 0.05 - 0.3 degrees. More preferably, the 2Q diffraction angles are shifted at least minus 0.05 degrees 2Q, preferably minus 0.05 - 0.3 degrees, wherein 0.05 - 0.3 includes 0.05, 0.06, 0.07, 0.08, 0.09, 0.10, 0.11, 0.12, 0.13, 0.14, 0.15, 0.16, 0.17, 0.18, 0.19, 0.20, 0.21, 0.22, 0.23, 0.24, 0.25, 0.26, 0.27, 0.28, 0.29 and 0.30. In one embodiment, wherein the material is Mg-doped Baghdadite the material may comprise at least one lattice parameter plus or minus 0.001 - 0.1 A selected from the group consisting of: a 7.354 A, b 10.180A, c 10.444 A, and b 90.820 A. Preferably, the material comprises one, two, three or four of the lattice parameters. In one embodiment, wherein the material is Mg-doped Baghdadite the material may comprise at least one lattice parameter of undoped Baghdadite, wherein the lattice parameter is shifted plus or minus between 0.001 - 0.05 A, preferably minus 0.001 - 0.05 A, more preferably minus 0.004 - 0.03 A. In a preferred embodiment, the material may comprise at least one lattice parameter of undoped Baghdadite selected from a, b, c, and b , wherein: a is shifted plus or minus 0.001 - 0.01 A, more preferably minus 0.001 - 0.01 A, wherein 0.001 - 0.01 A includes 0.001, 0.002, 0.003, 0.004, 0.005, 0.006, 0.007, 0.008, 0.009 and 0.01 A; b is shifted plus or minus 0.001 - 0.01 A, more preferably minus 0.001 - 0.01 A, wherein 0.001 - 0.01 A includes 0.001, 0.002, 0.003, 0.004, 0.005, 0.006, 0.007, 0.008, 0.009 and 0.01 A; c is shifted plus or minus 0.005 - 0.015 A, more preferably minus 0.005 - 0.015 A, wherein 0.005 - 0.015 A includes 0.005, 0.006, 0.007, 0.008, 0.009, 0.01, 0.011, 0.012, 0.013, 0.014 and 0.015 A; and b is shifted plus or minus 0.01 - 0.1 A, more preferably minus 0.01 - 0.1 A, wherein 0.01 - 0.1 A includes 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09 and 0.1

A.

In one embodiment, wherein the material is Fe-doped Baghdadite the material may comprise at least one lattice parameter of undoped Baghdadite, wherein the lattice parameter is shifted plus or minus between 0.001 - 0.05 A, preferably minus 0.001 - 0.05 A, more preferably minus 0.004 - 0.03 A. In a preferred embodiment, the material may comprise at least one lattice parameter of undoped Baghdadite selected from a, b, c, and b , wherein: a is shifted plus or minus 0.001 - 0.01 A, more preferably minus 0.001 - 0.01 A, wherein 0.001 - 0.01 A includes 0.001, 0.002, 0.003, 0.004, 0.005, 0.006, 0.007, 0.008, 0.009 and 0.01 A; b is shifted plus or minus 0.001 - 0.01 A, more preferably minus 0.001 - 0.01 A, wherein 0.001 - 0.01 A includes 0.001, 0.002, 0.003, 0.004, 0.005, 0.006, 0.007, 0.008, 0.009 and 0.01 A; c is shifted plus or minus 0.005 - 0.015 A, more preferably minus 0.005 - 0.015 A, wherein 0.005 - 0.015 A includes 0.005, 0.006, 0.007, 0.008, 0.009, 0.01, 0.011, 0.012, 0.013, 0.014 and 0.015 A; and b is shifted plus or minus 0.01 - 0.1 A, more preferably minus 0.01 - 0.1 A, wherein 0.01 - 0.1 A includes 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09 and 0.1

A.

In one form, the material is a ceramic material.

In another form, the material is a biocompatible material.

Preferably, the material is a ceramic biocompatible material.

In one form, the material is monophasic. In another form, the material may include Merwinite, calcium iron silicate or a combination thereof.

Preferably, the material does not comprise undoped Baghdadite.

Baghdadite , Merwinite and Calcium Iron Silicate (Ca ₃ FeSi20 ₉ ) were prepared as reference samples, and 4 preferred forms of the invention were prepared according to the following formulae, i.e., Fe-doped and Mg-doped Baghdadite s:

In Formulae I, la, II and III, “x” represents the sum of the molar concentrations of Mg (y) and Fe (z) (i.e., y + z = x). In Formula I and la, Mg or Fe respectively substitute Zr, and the associated generic formula is given by Formula A. In Formulae II and III, Mg or Fe respectively substitute Ca, and the associated generic formula is given by Formula B.

The data on human primary osteoblast-like cell attachment and proliferation and on mechanical properties disclosed herein shows that substituting small amounts of Fe or Mg for Ca in a Baghdadite structure surprisingly results in an up to -15% improvement in modulus without affecting the strength or the bioactivity of the material, and substituting Zr in a Baghdadite structure with small amounts of Mg surprisingly results in an up to 17% and 35% increase in modulus and strength of the material, respectively, whilst maintaining the bioactivity of the material. The combined improvements in mechanical properties and cellular activity makes the novel biocompatible ceramic materials disclosed herein important for a range of biomedical applications. The skilled person will appreciate the term “biocompatible” as used herein defines a two-way response, i.e. the body’s response to the ceramic material and the material’s response to the body’s environment. The biocompatibility of a medical device refers to the ability of the device to perform its intended function, with the desired degree of incorporation in the host, without eliciting any undesirable local or systemic effects in that host.

Preferably the ceramic materials of the invention display biocompatibility when placed in physiological fluid. Preferably, the biocompatible ceramic material of Formula A (e.g., of Formula I or Formula la) or Formula B (e.g., of Formula II or Formula III) as described herein forms a hydroxyapatite layer upon exposure to bodily fluids. As the skilled person will appreciate, the formation of hydroxyapatite is widely recognised as strong evidence that the body accepts the material as sui generis and is a requirement for the implant to chemically bond with living bone and tissue.

Whilst in preferred embodiments the ceramic materials of the invention are substantially pure, in other embodiments the materials include impurities, which may be in significant quantities, for example impurities of other transition metals. As the skilled person will appreciate, some tolerance to impurities may be acceptable. In one aspect, the ceramic materials of the invention are intermixed with apatite or tricalcium phosphate crystals. The present disclosure is the first time that the ceramic materials of the invention have been synthetically prepared and their potential use as a biocompatible material has been explored. It has been found that, surprisingly, the ceramic materials of the invention display exceptional biocompatibility, and more particularly, are particularly suited for the regeneration of bone and other tissue. In one embodiment, the inventors contemplate that the biocompatible ceramic materials of the invention find particular utility in resurfacing arthritic joints to promote the growth of articular cartilage. In other embodiments, the ceramic materials of the invention are useful in the development of 3D scaffolds which promote migration, proliferation and differentiation of bone and endothelial cells, for example in orthopaedic and maxillofacial surgeries, and yet provide sufficient mechanical properties for load-bearing parts. The ceramic materials of the invention also support bone tissue regeneration/formation and vascularization, and yet also provide minimal fibrotic reactions. In one aspect, the present invention provides scaffolds for osteochondral defects. In yet other embodiments, the present invention provides biocompatible ceramic materials which are coatable on currently used orthopaedic and dental implants to provide enhanced long-term implant stability. In further embodiments, the ceramic materials of the invention are selectively coatable on currently used orthopaedic implants, for example on areas where wear is an issue.

As discussed previously, the development of bioglass, glass-ceramics, and bioceramics containing CaO and S1O2 for bone tissue regeneration has received great attention in the past 3 decades. The stimulatory effect of the Ca and Si containing ionic products released from materials on osteoblast proliferation, differentiation, and related gene expression, and mineralization have also been well documented (see for example Xynos I.D., et al in Ionic products of bioactive glass dissolution increase proliferation of human osteoblasts and induce insulin-like growth factor II mRNA expression and protein synthesis, Biochem. Biophy. Res. Commun. 2000; 276:461-465). CaSiO3 based materials are considered as potential bioactive materials for bone tissue regeneration and implant coatings due to their bioactivity. However, a major drawback of the CaSiC ceramics is their relatively high dissolution rate leading to a high pH value in the surrounding environment, (see for example Siriphannon P, et al “Formation of hydroxyapatite on CaSiC3 powders in simulated body fluid”, J Eur. Ceram. Soc. 2002 22: 511-520). Indeed, the bonding of CaSiO3 coatings to titanium substrate degrades with the increasing immersion time in simulated body fluid (SBF) due to the relatively fast dissolution rate of the coating, which limits further biological applications. It has been unexpectedly found that the chemical modification of Baghdadite with Mg- and/or Fe- provides a bioceramic with significantly improved properties compared to previously known calcium silicates and previously known bioceramic materials.

In particular, the biocompatible ceramic materials of the invention provide many of the advantages of the CaSiO3 materials but ameliorate many of its disadvantages. The ceramic materials of the invention display a relatively reduced dissolution profile, which is associated with a relatively reduced pH compared to CaSiC materials. Further, the ceramic materials of the invention exhibit excellent mechanical properties and allow attachment and proliferation of bone cells. In particular, the ceramic materials of the invention have been found to form a chemical bond with bone and to form an apatite layer. Furthermore, the ceramic materials of the invention display relatively reduced corrosion in biological environments.

According to a second aspect the present invention provides a method for preparing the biocompatible ceramic material of the first aspect above, the method comprising: • providing a sol of precursor materials;

• at least partially gelling the sol; and

• drying and sintering the at least partially gelled sol to thereby form the biocompatible ceramic material.

In one embodiment of the second aspect, there is provided a method for preparing the biocompatible ceramic material of the first aspect above, the method comprising: providing a sol of precursor materials selected from the group consisting of ; at least partially gelling the sol, and drying and sintering the at least partially gelled sol to thereby form the biocompatible ceramic material.

The ceramic materials according to the second aspect of the invention are sol-gel derived. However, it will be appreciated that, in other embodiments, any method of synthetic production of the ceramic materials described herein would fall within the purview of the present invention. For example, in another embodiment, S1O2, CaO, Fe20 ₃ , MgO and Zr02 may be melted at relatively high temperatures (for example see the methodology outlined in Mazerolles, L. et al. Aerospace Science and Technology, 2008 12: 499-505) and then cooled, and the resulting material pulverized. The resulting powder can then be formed and hot-pressed, as is well known in the art, for example (see Russias J. et al. Journal of the European Ceramic Society, 2007 27: 327-335).

According to a third aspect the present invention provides an implantable medical device comprising the biocompatible ceramic material of the first aspect of the invention. In one embodiment, the present invention provides an implantable medical device consisting of the biocompatible ceramic material of the first aspect of the invention.

The implantable medical device is preferably chosen from the group consisting of: a 3D implantable scaffold, an orthopaedic implant (e.g., for reconstructive surgery), a dental implant/prosthesis, a spine implant, an implant for craniofacial reconstruction or alveolar ridge augmentation, an implant for cartilage regeneration, an osteochondral defect implant, a strut, a stent and a stent-graft. However, it will be appreciated that there are many other devices which would be within the purview of the present invention. The skilled person will readily appreciate how to manufacture a medical device from the biocompatible material of the invention. For example, the inventors contemplate that the biocompatible material of the invention can be formed into a medical device in a similar methodology as outlined in the prior art, for example see Hench L.L. J. Am. Ceram. Soc. 1991 74: 1487-1510; and Zhao J. et al. Biomed. Mater. 2006 1: 188-92. The implantable medical device comprising a ceramic material as described herein may be suitable for permanent implant, or it may be substantially biodegradable in vivo.

An embodiment of the invention is a bone implant comprising the ceramic material of the invention. The implantable medical device comprising a ceramic material according to the first aspect of the invention may therefore be a bone implant.

An embodiment of the invention is a tooth filling implant comprising the ceramic material of the invention. The implantable medical device comprising a ceramic material according to the first aspect of the invention may therefore be a tooth filling implant.

An embodiment of the invention is a biocement comprising the ceramic material of the invention. The implantable medical device comprising a ceramic material according to the first aspect of the invention may therefore be a biocement. In other embodiments, the doped Baghdadite of the invention may be formed into a surgical device or as a coating on a surgical device. For example, Ti-6A1-4V, a titanium alloy, is well established as one of the primary biomaterials for orthopaedic implants because of its excellent biocompatibility, low toxicity, high chemical stability, low rate of corrosion and favourable mechanical properties. However, Ti-6A1-4V has a crucial drawback, namely poor wear resistance. Adhesive and abrasive wear at the bone- implant interface and articulating surfaces generates debris. This debris - small particles and shards of metal that detach from the implant surface - enter the surrounding tissue and migrate into spaces between the bone and implant where they induce inflammation and associated bone destruction, leading to aseptic loosening. This jeopardises the stability of the prosthesis, leading to the premature failure of the device, as well as pain and disability in patients (Haynes, D.R., T.N. Crotti, and H. Zreiqat, “Regulation of osteoclast activity in peri-implant tissues”, Biomaterials , 2004 25: 4877-4885). As a result, global failure rates of orthopaedic implants, mainly hip and knee replacements, are unacceptably high. The success of orthopaedic implants depends on strong anchorage of the device material in bone tissue. Various biomaterials modifications have been applied in an attempt to enhance bone formation, but to date none forms a stable interface with the strength required to support functional loading for the lifetime of the patient. Ideally, the implant should also interact with the host tissue, recruiting and even promoting differentiation of osteogenic cells, rather than acting as a passive stage for the performance of any itinerant cells. An important factor in selecting orthopaedic implant material, therefore, is identifying the correct chemistry to support or stimulate an appropriate host response. Frequently implant materials are not preferentially compatible with bone cells responsible for bone formation; rather, they promote the formation of undesirable soft connective tissue by other cells such as fibroblasts. Considerable effort has gone into developing surface treatments and coatings to improve host tissue - implant integration. Although these approaches have had some success, they have been shown to have slow rates of osseointegration and poor mechanical anchorage in challenging clinical cases, such as those associated with large bone loss and poor bone quality (Sporer, S.M. and W.G. Paprosky, “Biologic fixation and bone ingrowth” Orthop Clin North Am, 2005 36: 105-111, vii). During the last two decades, various surface modification methods have been proposed to improve bone conductivity or bioactivity of Ti-6A1-4V by coating it with ceramic. The aim has been to enhance osseo-integration and thereby interlock the implant with the surrounding skeletal tissue, providing a stable interface strong enough to support life-long functional loading. The coating should prevent corrosion of the underlying substrate in a biological environment; create a barrier against the release of the toxic metal debris into the body (Sun, L., et al., “Material fundamentals and clinical performance of plasma- sprayed hydroxyapatite coatings: a review”, J Biomed Mater Res, 2001 58: 570-592); and combine the mechanical properties of the metal with the bioactivity of the ceramic. One such approach is to coat Ti-6A1-4V with bioactive ceramics such as HA and calcium silicate ceramics (CaSi0 ₃ ) (for example see Harle, J., et al., “Initial responses of human osteoblasts to sol-gel modified titanium with hydroxyapatite and titania composition”, Acta Biomater, 20062: 547-556). HA has been used to coat hip-joint endoprostheses for the enhancement of long-term fixation in femoral bone (Ha, S.W., et al., “Chemical and morphological changes of vacuum-plasma-sprayed hydroxyapatite coatings during immersion in simulated physiological solutions”, J Am Ceram. Soc. 1998 81: 81-88). These have been shown to improve the stability of the Ti-6A1-4V implant, the interface strength, the bone mineralization, and the bone ingrowth rate (Soballe, K., et al., “Gap healing enhanced by hydroxyapatite coating in dogs”, Clin. Orthop. Relat. Res., 1991 272: 300-307). It is contemplated that the biomaterial/bioceramic materials of the present invention, which have improved properties compared to these prior art coatings, provide a coated implant having improved service life and excellent osseointegration. In one embodiment the porosity of the medical device comprising the ceramic of the invention is between about 10% and about 80%, e.g., between about 20% and about 80%, or between about 20% and about 70%, or between about 20% and about 60%, or between about 20% and about 55%, or between about 25% and about 80%, or between about 25% and about 70%, or between 25% and 60%, or between about 25% and about 55%, or between about 30% and about 80%, or between about 30% and about 70%, or between about 30% and about 55%, or between about 30% and about 40%, or between about 20% to about 30%, or between about 10% and 40%, or between about 25% and 50%, or between about 40% and 75%, or between about 50% and 80%. However, it will be appreciated that the device could be configured to have lower or greater porosity according to the intended use, and any porosity range would be within the purview of the present invention. For example, porosities of 10, 15, 20, 25, 30, 35, 40, 45, 50, 52, 55, 60, 65, 70, 75 or 80% are possible. In some embodiments, porosities may be any one of these values, selected from one or more of these values or between any two of these values.

In one embodiment, the pore size of the device is between about 20 to about 500 micron, e.g., between about 75 and about 200 μm , or between about 20 and about 100 μm , or between about 50 and about 250 μm , or between about 100 and about 300 μm , or between about 150 and about 400 μm , or between about 300 and about 500 μm , or between about 350 and about 450 μm , or between about 400 and about 500 μm , or between about 75 and about 400 μm . However, it will be appreciated that the device could be configured to have lower or greater pore size according to the intended or desired use, and any pore size would be within the purview of the present invention. For example, pore sizes of 20, 30, 40, 50, 60, 70, 80, 90, 100, 110, 120, 130, 140, 150, 160, 170, 180, 190, 200, 210, 220, 230, 240, 250, 260, 270, 280, 290, 300, 310, 320, 330,

340, 350, 360, 370, 380, 390, 400, 410, 420, 430, 440, 450, 460, 470, 480, 490, or 500 micron are possible. As the skilled person will appreciate, the porosity of ceramics can be adjusted by controlling the content and size of porogens.

The compressive strength of the ceramics of the invention are preferably between 5 MPa and 250 MPa, such as between 5 MPa and 200 MPa, or between 5 MPa and 150

MPa, or between 10 MPa and 150 MPa, or between 20 MPa and 150 MPa, or between 30 MPa and 150 MPa, or between 40 MPa and 150 MPa, or between 50 MPa and 150 MPa, or between 60 MPa and 150 MPa, or between 70 MPa and 150 MPa, or between 90 MPa and 150 MPa, or between 40 MPa and 120 MPa, or between 50 MPa and 100 MPa, or between 60 MPa and 80 MPa. The compressive strength of the ceramics of the invention are preferably between 100 and 250 MPa, such as between 100 and 150 MPa, or between 150 and 200 MPa, or between 125 and 225 MPa, or between 200 and 250 MPa, or between 175 and 250 Mpa, e.g., may be 110, 120, 130, 140, 150, 160, 170, 180, 190, 200, 210, 220, 230, 240 or 250 MPa. The compressive strength of implantable medical devices comprising a ceramic of the invention may be between 5 MPa and 250 MPa, such as between 5 MPa and 200 MPa, or between 5 MPa and 150 MPa, or between 10 MPa and 150 MPa, or between 20 MPa and 150 MPa, or between 30 MPa and 150 MPa, or between 40 MPa and 150 MPa, or between 50 MPa and 150 MPa, or between 60 MPa and 150 MPa, or between 70 MPa and 150 MPa, or between 90 MPa and 150 MPa, or between 40 MPa and 120 MPa, or between 50 MPa and 100 MPa, or between 60 MPa and 80 MPa. The compressive strength of implantable medical devices comprising a ceramic of the invention may be between 100 and 250 MPa, such as between 100 and 150 MPa, or between 150 and 200 MPa, or between 125 and 225 MPa, or between 200 and 250 MPa, or between 175 and 250 Mpa, e.g., may be 110, 120, 130, 140, 150, 160, 170, 180, 190, 200, 210, 220, 230, 240 or 250 MPa. The modulus of the ceramics of the invention may be between 0.5 and 20 GPa, e.g. between 0.5 and 15 GPa, or between 0.5 and 10 GPa, or between 0.5 and 15 GPa, or between 2 and 20 GPa, or between 4 and 18 GPa, or between 6 and 16 GPa, or between 8 and 14 GPa, or between

9 and 12 GPa. The modulus of the ceramics of the invention may be between 10 and 20 GPa, e.g., between 10 and 15 GPa, or between 12 and 18 GPa, or between 15 and 20 GPa, or between 17.5 and 20 GPa, such as 10, 11, 12, 13, 14, 15, 16, 17, 18, 19, or 20 GPa. The modulus of implantable medical devices comprising a ceramic of the invention may be between 0.5 and 20 GPa, e.g. between 0.5 and 15 GPa, or between 0.5 and 10

GPa, or between 0.5 and 15 GPa, or between 2 and 20 GPa, or between 4 and 18 GPa, or between 6 and 16 GPa, or between 8 and 14 GPa, or between 9 and 12 GPa. The modulus of implantable medical devices comprising a ceramic of the invention may be between 10 and 20 GPa, e.g., between 10 and 15 GPa, or between 12 and 18 GPa, or between 15 and 20 GPa, or between 17.5 and 20 GPa, such as 10, 11, 12, 13, 14, 15, 16, 17, 18, 19, or 20 GPa.

These mechanical properties are ideal for scaffolds to be placed in load-bearing applications as the strength of the natural bone is within this range. In one embodiment, the implantable medical device comprising a ceramic of the invention features a rotated cubic unit cell. In a rotated cubic design the cubic lattice is rotated in three dimensions so that the strut that is aligned with axis x (vector [1 00]) in a cubic design is located along the vector [1 1 1] in the rotated design. In one embodiment, the implantable medical device comprising a ceramic of the invention comprises one or more apertures, preferably the ceramic may comprise a series of apertures formed in a regular pattern or design. Apertures may be incorporated into the ceramic after it has been formed, for example, through drilling or tunnelling, or may be introduced during formation, for example in a sterolithographic process or by forming the ceramic in a mould. The aperatures may adopt any suitable shape, such as cylindrical, box-like, and so on. The aperatures may be including blind or unblind. In some embodiments, the ceramic comprising the one or more apertures adopts a shape corresponding to a negative replication of a strut-based design. Incoporation of the one or more apertures may be advantageous to reduce the weight of the ceramic and allow increased fluid penetration which may promote its function as a scaffold for bone or cartilage growth.

Implantable devices according to the present invention have many properties that make them suitable for use as implants, including high mechanical strength, resistance to fatigue, corrosion resistance, and biocompatibility. The implants may be implanted in animals, non-limiting examples of which include reptiles, birds, and mammals, with humans being particularly preferred.

The devices of this invention may be implanted into a body in different ways, including, but not limited to subcutaneous implantation, implantation at the surface of the skin, and implantation in the oral cavity. In one embodiment, the medical device comprising a ceramic material of the present invention may be coated with at least one resorbable polymer material, non- limiting examples of which include polyglycolides, polydioxanones, polyhydroxyalkanoates, polylactides, alginates, collagens, chitosans, polyalkylene oxalate, polyanhydrides, poly(glycolide-co-trimethylene carbonate), polyesteramides, or polydepsipeptides etc. Alternatively, the coating material may comprise at least one healing promoter such as a thrombosis inhibitor, a fibrinolytic agent, a vasodilator substance, an anti-inflammatory agent, a cell proliferation inhibitor, or an inhibitor of matrix elaboration or expression. Examples of such substances are discussed in U.S. Patent No. 6,162, 537. The present invention also contemplates using a polymer coating, ( e.g . a resorbable polymer) in conjunction with a healing promoter to coat the implantable medical device, for example according to the reference [Wu C. Acta Biomateilia, 2008 4: 343-353].

As described herein, the present invention provides a method for producing an implantable medical device comprising: transferring a biocompatible ceramic material of the invention onto a substrate thereby forming said implantable medical device.

It will be appreciated that there are a number of methods of transferring the ceramic materials of the invention onto a supporting surface or substrate, and any of these methods fall within the purview of the present invention. For example, in one embodiment, the ceramic material is plasma spray coated. As is well known in the art, this method essentially comprises the steps of spraying molten or heat softened material onto a surface to provide the coating. The material, in the form of powder, is injected into a high temperature plasma flame, where it is rapidly heated and accelerated to a high velocity. The hot material impacts on the substrate surface and rapidly cools thereby forming a coating (see for example Liu X. Biomedicine & Pharmacotherapy 2008 62: 526-529). The coatings have a dense structure with a thickness of about 50 pm.

According to an embodiment of the invention, an implantable drug delivery device is provided comprising a ceramic material of the invention. It will be appreciated that the drug delivery device can deliver any drug and the can be shaped to suit the particular application. For example, see Krajewski et al J. Mater. Sci.: Mater. In Med. 2006 12: 763-771.

According to a fourth aspect, the present invention provides a method for improving the long-term stability of an implantable medical device, the method comprising: coating the implantable medical device with the biocompatible ceramic material according to the first aspect of the invention.

Preferably the method further includes coating the implantable medical device or the biocompatible ceramic material with a biocompatible polymer. Accordingly, preferably the coating includes a biocompatible polymer, which in one embodiment is PLGA. In one aspect the implantable medical device is a biphasic scaffold for an osteochondral defect.

According to a fifth aspect, the present invention provides a method for regenerating or resurfacing tissue, the method comprising: contacting the regenerating or resurfacing tissue with a quantity of a ceramic material according to a first aspect of the invention for a sufficient period to at least partially effect said regeneration or resurfacing.

According to a sixth aspect, the present invention provides a method for forming osseous tissue on an orthopaedic defect, the method comprising: contacting said defect with the ceramic material according to the first aspect of the invention. The present inventors contemplate that the defect could be contacted with, for example, a cementing paste comprising the ceramics described herein and cured or allowed to set. The presence of a biocompatible ceramic of the invention would act to stimulate the formation of the osseous tissue on the orthopaedic defect. As described herein, the present invention provides a method for treating orthopaedic conditions comprising, contacting a patient in need of such treatment with an effective regenerating amount of biocompatible composition comprising a ceramic material of the invention.

As described herein, the present invention provides a kit for regenerating or resurfacing tissue, comprising a biocompatible ceramic material of the invention and a therapeutic agent which stimulates and accelerates tissue regeneration. Such therapeutic agents are well known the art.

In one embodiment, preferably the biocompatible ceramic material of the invention is a fully synthetic bone graft substitute. Due to its interconnected pores, the material serves as an ideal osteoconductive scaffold and supports the formation of new host bone. As highlighted above, the advantages of the new material can be summarised as one or more of the following:

• optimized porosity,

• enhanced bone ingrowth and vascularization, · avoids potential problems common for grafting methods,

• is formable to almost any shape to suit the application,

• easy to use,

• combines with autologous bone marrow or blood, and

• displays accelerated and enhanced osteointegration. The uses of the present invention include one or more of the following:

• for bone void fillings or augmentation in zones requiring cancellous rather than cortical bone, • for the filling of bone defects after trauma, reconstruction, or correction in non load or load-bearing indications,

• for trauma and orthopaedics - filling of voids caused by cysts or osteotomies, filling of defects arising from impacted fractures, refilling of cancellous bone harvesting sites, arthrodesis and non-unions,

• for spine surgery - postero-lateral fusion, interbody fusion (as cage-filling material), vertebrectomies (as filling material of the vertebral implants), refilling of bone graft-harvesting sites, and

• for cranio -maxillofacial surgery - reconstruction of mandibular defects and sinus lifts.

According to a seventh aspect, the present invention provides a method of preparing the biocompatible ceramic material of the first aspect above, the method comprising:

• providing a resin comprising ceramic powder;

• depositing layers of resin comprising ceramic powder using a stereolithography or 3D printing device;

• crosslinking the layers and sintering to thereby form the biocompatible ceramic material.

In one embodiment of the seventh aspect, the biocompatible ceramic material is formed with a rotated cubic unit cell. In a rotated cubic design the cubic lattice is rotated in three dimensions so that the strut that is aligned with axis x (vector [1 00]) in a cubic design is located along the vector [1 1 1] in the rotated design.

In one embodiment of the seventh aspect, the resin is photocurable.

In one embodiment of the seventh aspect, the resin further comprises a dispersant, such as Tween 20.

In one embodiment of the seventh aspect, the sintering temperature is at least about 1200 °C, 1250 °C, 1300 °C, 1350 °C, 1400 °C, 1450 °C or 1480 °C. In one embodiment, the sintering temperature is not more than about 2000 °C, 1900 °C, 1800 °C, 1700 °C, 1600 °C, 1550 °C or 1500 °C. The sintering termperature may be between any of these temperatures, for example, from about 1200 °C to about 2000 °C or from about 1350 °C to about 1550 °C. In one embodiment of the seventh aspect, the sintering temperature is greater than about 1400 °C. In one embodiment of the seventh aspect, the sintering temperature is about 1400 °C. In one embodiment of the seventh aspect, the sintering temperature is about 1480 °C.

Methods of preparing a ceramic material using a stereolithographic or 3D printing device are described in PCT/AU2020/050975. The methods described in PCT/AU2020/050975 may also be applied to the biocompatible ceramic material of the present invention. The entire disclosure of PCT/AU2020/050975 is incorporated herein by reference.

BRIEF DESCRIPTION OF THE DRAWINGS Embodiments of the invention will now be described, by way of example only, with reference to the accompanying drawings in which:

Figure la. shows an XRD of Baghdadite ( and Mg-doped Baghdadite .

Figure lb. shows a zoom of the 25 - 38 degree range of Figure la. Figure lc. shows the value shift of XRD peaks of the spectra in Figures la and lb (dashed lines).

Figure 2. shows an XRD of Baghdadite , Calcium Iron Silicate , and Fe-doped Baghdadite where x = 0.05, 0.1, or 0.2). Figure 3 shows the cell proliferation and mechanical tests results for (a) cell proliferation, (b) modulus, and (c) strength of the material.

Figure 4 shows the cell proliferation and mechanical tests results for (a) cell proliferation, (b) modulus, and (c) strength of the material. Figure 5 shows the cell proliferation and mechanical tests results for (a) cell proliferation, (b) modulus, and (c) strength of the material.

Figure 6 results for compressive stiffness, compressive strength and Runx2 expression for 13 compositions of one of the following formulae: Figure 7 shows the mechanical and physical properties of Mg doped Baghdadite bioceramics.: (a) the effect of doping Mg in Baghdadite on the flexural modulus and strength along with three-point bending inset showing the dimensions of the specimens, (b) representative scanning electron microscopy (SEM) images of i) fracture surfaces at varying Mg doping, along with the binarized images showing the void content (quantitatively compared in (c), and ii) the grain size for different compositions (c) The void content, density, and grain size as function of molar concentration of Mg (x), (d) correlation between grain size, density, and flexural strength. The size of the markers is indicative of the magnitude of the strength (plotting the modulus would yield the same trends). Shaded regions in (a) and (c) show 2 standard deviations. N shows the number of samples for each data point scale bars are 60 and 3 pm in i) and ii) in panel (b).

Figure 8 shows the osteogenic response of human adipose-derived stromal cells seeded on Mg doped bioceramics of varying Mg doping. Nuclear translocation of the osteogenic factor RUNX2 expression in nucleus relative to cytoplasm is shown for different compositions and against the reference expansion and osteogenic media. Shaded regions in Figure 8a show 2 standard deviation. The scale bar in Figure 8b is 100 pm.

Figure 9 shows stereolithography (SLA) of bioceramic scaffolds. Figure 9a shows the printing set up. Figure 9b shows the optical and SEM images of the scaffolds printed for in-vivo evaluation (scale bar 2mm).

Figure 10 shows the mechanics of scaffolds pre- and post-implantation. Figure 10a shows the stress strain curves. Figure 10b shows the in-situ optical images and different time spots during loading. Figure 10c shows the comparison between the stiffness, strength and energy absorption of the scaffolds with two materials pre- and post-implantation.

Figure 11 shows the architecture and mechanics of the optimized material. Figure 11a shows scaffolds with cylindrical holes. Figure lib shows the stress-strain curves. Figure 11c shows an assessment between between stress and strain.

DEFINITIONS

In describing and claiming the present invention, the following terminology will be used in accordance with the definitions set out below. It is also to be understood that the terminology used herein is for the purpose of describing particular embodiments of the invention only and is not intended to be limiting. Unless defined otherwise, all technical and scientific terms used herein have the same meaning as commonly understood by one having ordinary skill in the art to which the invention pertains.

The recitation of a numerical range using endpoints includes all numbers subsumed within that range (e.g., 1 to 5 includes 1, 1.5, 2, 2.75, 3, 3.80, 4, 5, etc.).

The terms “preferred” and “preferably” refer to embodiments of the invention that may afford certain benefits, under certain circumstances. However, other embodiments may also be preferred, under the same or other circumstances. Furthermore, the recitation of one or more preferred embodiments does not imply that other embodiments are not useful, and is not intended to exclude other embodiments from the scope of the invention.

As used herein, an “implant” refers to an article or device that is placed entirely or partially into an animal, for example by a surgical procedure. Unless the context clearly requires otherwise, throughout the description and the claims, the words ‘comprise’, ‘comprising’, and the like are to be construed in an inclusive sense as opposed to an exclusive or exhaustive sense; that is to say, in the sense of “including, but not limited to”.

Other than in the operating examples, or where otherwise indicated, all numbers expressing quantities of ingredients or reaction conditions used herein are to be understood as modified in all instances by the term “about”. The examples are not intended to limit the scope of the invention. In what follows, or where otherwise indicated, “%” will mean “weight %”, “ratio” will mean “weight ratio” and “parts” will mean “weight parts”. Notwithstanding that the numerical ranges and parameters setting forth the broad scope of the invention are approximations, the numerical values set forth in the specific examples are reported as precisely as possible. Any numerical value, however, inherently contains certain errors necessarily resulting from the standard deviations found in their respective testing measurements.

PREFERRED EMBODIMENT OF THE INVENTION

Preferred embodiments of the present invention will be described in the following. Material preparation

Baghdadite was initially prepared as described in WO 2009/052583, which is incorporated herein by reference. Fe-doped and Mg-doped Baghdadite with four different formulations were then prepared: In these formulations, “x” represents the molar concentrations of either Mg or Fe. In two of these materials, Mg and Fe substitute Zr, in the other two materials Mg and Fe substitute Ca. a) this material constitutes of the same elements as the material (c) but the difference is that Mg substitutes Zr instead of Ca. The values of x = 0 (i.e., Baghdadite), 0.05, 0.1, 0.2, 0.3, 0.5, 1 (i.e. Merwinite) were explored. b) this material constitutes of the same elements as material (d) but the difference is that Fe substitutes Zr instead of Ca. The values of x = 0 (i.e., Baghdadite), 0.05, 0.1, 0.2, were explored. c) in this material, Mg was used to substitute Ca. The ceramic precursor used for Mg was MgO (342793, Sigma Aldrich, Australia). The values of x = 0 (i.e., Baghdadite), 0.05, 0.1 and 0.2 were explored. d) this material was prepared by mixing stochiometric amounts of calcium oxide CaO (248568, Sigma Aldrich, Australia), zirconium oxide Zr02 (230693, Sigma Aldrich, Australia), Iron(III) oxide Fe20 ₃ (310050, Sigma Aldrich, Australia), and silicon oxide SiO2 (S5631, Sigma Aldrich, Australia). Different values of x = 0 (i.e., Baghdadite), 0.05, 0.1, 0.2 were explored.

In all these materials, wet ball milling was used to mix, homogenize and refine the ceramic pre-cursor particles. 20wt% particles were mixed with 60 wt% zirconia balls and 20 wt% ethanol. The zirconia ball had three different diameters of 20, 3 and 1 mm with ratios of 80 wt%, 15 wt%, and 5 wt% respectively. The mixture was ball-milled for 5 hours at 200 rpm. Typically, the homogenized particle size was D20 = 0.11 ± 0.02 pm, D50 = 0.19 ± 0.03 pm and D80= 1.28 ± 0.02 pm, measured using laser scattering particle size analysis (LA-960 HORIBA, Japan). The refined ceramic slurry was then dried at room temperature for one day. The dried material was manually ground with a mortar and pestle, which was then filtered through a 200 pm sieve to remove the large clusters. XRD and mechanical tests

Cylinder- shaped green bodies with two different sizes of (i) diameter d=5mm and height h= 10 mm and (ii) diameter d = 10mm and height h = 1 mm, were prepared using uniaxial dry pressing at approximately 2 MPa. The samples were kept at this pressure for 2 min. The samples were heated to up to (and kept for 1 hour) at 620 °C at 2 °C/min to remove organic residues and water. The temperature was then increased to 1400°C at 2 °C/min, at which temperature the samples were kept for 3 hours to be fully cured. Batch (i) were used for mechanical compression tests depicted in Figure 3 and batch (ii) were used for cell proliferation studies depicted in Figures 4 and 5. For compression tests, the top and bottom surfaces of the samples were polished and made parallel to each other. The tests were performed at 5 pm/s using an Instron axial system (MA, US). For each chemistry, the mechanical tests were performed on 5 samples (N=5), and N=3 for cell proliferation tests. X-ray powder diffraction (XRD) was used to study the phases and crystal structure of the modified baghdadite. For Mg-doped Baghdadite a copper X-ray source was used. For Fe-doped Baghdadite a molybdenum X-ray source was used to minimise the reflections/noise observed using a copper X-ray source in presence of iron.

Cell study. Human primary osteoblast- like cells (HQBs) isolation Permission to use discarded human tissue was granted by the Human Ethics

Committee of Western Sydney Children Hospital and informed consent was obtained. Human trabecular bone was chopped into 1 mm ³ pieces and washed several times in phosphate buffered saline (PBS), followed by digestion for 90min at 37°C with 0.02% (w/v) trypsin (Sigma-Aldrich, USA) in PBS. Digested cells were cultured in complete media containing a-Minimal Essential Medium (a-MEM, Gibco Laboratories, USA), supplemented with 10% (v/v) fetal calf serum (FCS, Gibco Laboratories, USA), 30 mg/ml penicillin, 100 mg/ml streptomycin (Gibco Laboratories, USA) and ImM 1- ascorbic acid phosphate magnesium salt (Wako Pure Chemicals, Osaka, Japan). The cells were cultured at 37°C with 5% C02, and the medium refreshed every three days until confluence when cells were passaged. All HOBs used in the experiments were at passage 2.

HOBs attachment and proliferation assays

HOBS were cultured on different discs placed in 24-well culture plate at an initial density of 2.5 x 10 ⁴ cells/cm ² and incubated in a-MEM culture medium supplemented with 10% FCS maintained at 37°C in a humidified atmosphere of 95% air and 5% CO2. At 5 and 24 hours, the discs were removed from the culture wells, rinsed with PBS (pH = 7.4), fixed with 1.25% glutaraldehyde, 4% paraformaldehyde, and 4% sucrose in PBS for 1 hour. The fixative was removed by washing with buffer containing 4% (w/v) sucrose in PBS and post fixed in 1% osmium tetroxide in PBS followed by sequentially dehydrated in graded ethanol (70, 90, 95, and 100%). The specimens were dried in hexamethyldisilizane (HMDS) for 3 minutes before coating with gold for SEM analysis. The morphological characteristics of the attached cells on the ceramic discs were determined using SEM.

The CellTiter 96® Aqueous One Solution Reagent contains a tetrazolium compound [3-(4,5-dimethylthiazol-2-yl)-5-(3-carboxymethoxyphenyl)-2-( 4-sulfophenyl) -2H-tetrazolium, inner salt; MTS] was used for assessing HOB proliferation after 1 and 7 days of culture in a-MEM complete culture medium. One hundred micro litres of the reacted reagent from each well was transferred to a 96-well plate, and the absorbance was recorded using a microplate reader at a wavelength of 490 nm.

Alternative preparation

Pallets were prepared using uniaxial dry pressing under - 0.15 GPa for 2 min (Hydraulic press). They were heated to (and kept 1 hour at) 600 °C at 2 °C/min to remove organic residues and water. The temperature was then increased to 1400 °C at 2 °C/min, at which temperature the samples were kept for 3 hours for the reaction between the precursors to be completed and the sample to be fully sintered.

Making beams for bending tests

Cylindrical disks (dimeter: 25mm) from the six compositions (x values) were prepared using uniaxial dry pressing under — 0.15 GPa for 2 min. After sintering at 1400 °C (using the same heat profile as described in the alternative preparation), the samples underwent - 22 % shrinkage. Beams were then cut out of the disks using a precision diamond saw (Accutom-50, Struers), which were then tested in 3 point bending. Results are depicted in Figure 6.

Cell studies

For cell screening test, disk- shaped samples with diameter d = 10mm and height h = 1 mmm were prepared following the protocol described for making beams for bending tests. The surface of the samples were then polished with a fine grade (1200) sand paper for 5 min. Human adipose-derived stromal cells (hADSCs) (Life Technology), were cultured in expansion media: MesenPRO RS™ Basal Medium (Invitrogen) with the supplement of 2 mM 1-glutamine and MesenPRO RS Growth Supplement (Life Technology). hADSCs at passage 3 were used for studies and cultured in growth media. For RUNX2 assay, hADSCs were treated with mitomycin (10 pg mL ^-1 ) (Cayman Chemical, 11435) for 2h, to inhibit proliferation, 24 h after seeding, after which media was replaced with an expansion media. For RUNX2 translocation studies, cells were seeded at 3,000 cells cm ² . For positive controls, an osteogenic media (Gibco, A1007001) was used as noted.

Immunocytochemistry and quantification of RUNX2 localisation.

For immunostaining, all solutions, with the exception of those with dilute antibody and fluorophores, were syringe filtered through 0.22 pm membrane filters (Merck Millipore SLGP033RS). hADSCs were fixed at room temperature 4% PFA in xl PBS buffer for 10 min and then washed three times with PBS, followed by 12 min permeabilization at room temperature with 0.1 w v ^-1 % Triton X-100 in PBS. Samples were then incubated in a blocking buffer of 3% BSA, 3.75 mg mL ¹ glycine and 0.05% w v ¹ Tween 20 in PBS for 1 h at room temperature. Antibodies were diluted in 1% BSA in PBS and added overnight at 4 °C, anti-RUNX2-AF488 (1:250; Santa Cruz sc-390351). Nuclear and actin counterstains were performed simultaneously using Abeam iFluor conjugated phalloidin (abl76759), and Hoechst 33342 at 0.1 pg mL ^-1 (Sigma, 14533) dilute in xl PBS and incubated for 30 min at room temperature. Following counterstain incubation, samples were washed an additional x3 with PBS containing 0.05 w v ¹ % sodium azide. Microscopy was completed on a NiE Fluorescence Widefield, with x40/1.2NA H20 objective, using software triggered autofocus, triggered to find the plane maximum fluorescence intensity, and monochrome - DS-Qi2 camera. Nuclear images from Hoechst staining were used to create masks that define nuclear area, and cytoplasmic area masks were defined from flood-filled phalloidin stains, followed by calculation of the average fluorescent intensity of each area, using MATLAB with a custom script making use of the open microscopy Bio-Formats tool. Therein, the RUNX2 nuclear to cytoplasmic translocation ratio was determined as the ratio of the mean RUNX2 fluorescent saturation intensity of the nuclear area divided by the fluorescent saturation intensity in the non-nuclear cell cytoplasmic area. Results are depicted in Figure 6.

Preparing the ceramic resins for printing

The sintered pallets were first ground using ball milling to make the ceramic particles. balls, 70 gr of the 3mm balls and 15gr of the 1mm balls). This mixture was ground for 10 hours at 150 rpm using a high energy ball mill (Retsch, PM 400, Dusseldorf, Germany) to achieve particles size distribution with Dio, D50 and D90 of 0.24, 9.52,

72.51 pm respectively. The particle size was measured by using laser diffraction analysis (LA-960 HORIBA, Japan). After drying in air, 65 wt% of the ceramic particles were mixed with 17.5% photosensitive polymer (clear resin V4, Formlabs, USA) and 17.5% dispersant (Twin 20, Sigma Aldrich) using the same ball mill as above but only large balls (1 per 20 gr of particles) were used to mix the resin for 30 min at 100 rpm. The particle size measurement after mixing showed that the particle size was not affected by the mixing procedure.

Silanization of glass slides as substrate for 3D printing

Silanization was performed following a protocol similar to that reported in Mirkhalaf,

M. and F. Barthelat, Journal of the mechanical behavior of biomedical materials, 2016. 56: p. 23-33. 1 mm thick glass slides (Corning, NY, US) were first cleaned using ultrasonication for 5 min in an ethanol bath. 3 vol% of a silane (3- (trimethoxysilyl)propyl acrylate, Sigma Aldrich) was prepared in a methanol (95vol%)- acetic acid (5vol%) solvent. The glass slides were incubated in the solution for 45 min and then rinsed with methanal and ethanol respectively before getting dried with a positive air flow.

Polishing scaffolds for mechanical tests

A polishing holder that contained holes with the same diameter as that of the scaffolds (4 mm) but with different depths (5.8mm and 5.5mm) were machined from stainless steel. Initially the samples were put in the deeper hole and were polished until a flat surface on one side was achieved. Then the sample was rotated and put it inside the shorter hole to polish the other surface. The text fixtures were made of stainless steel.

Implantation The viability of the Ca ₃ Mgo.iZro.9Si208.9 composition as a bone substitute was assessed with an in vivo bone formation model using a murine hind limb muscle capsule transplant. Undoped Baghdadite (Ca3ZrSi2O9) was employed as a control. The study design used 6 rats randomly assigned to two groups (N = 6 per group) to receive bilateral implants. Test groups included (a) Baghdadite scaffold, and (b) Mg-doped baghdadite (Ca ₃ Mgo.iZro.9Si208.9) scaffolds. The animal study was approved by the local animal ethics committee (protocol K373). Animals were purchased from the Animal Resources Centre (Perth, Australia). Male Wistar rats were operated on when they were 12-14 weeks old with a mean weight of ~350g. Buprenorphine (0.1 mg. kg ^-1 ) was given subcutaneously one hour before surgery for pain relief. Surgical anesthesia was induced with intraperitoneal ketamine (75 mg. kg ¹ ) and xylazine (10 mg. kg ^-1 ) and maintained with inhaled isoflurane (2-3% with 1.5-2L oxygen) as required. The right and left legs of each animal were shaved and prepared with a povidone-iodine solution prior to surgery. A 1.5 cm lateral mid-thigh approach was made, and then, the vastus lateralis (the thigh muscle on top of the femur) was dissected with a 1 cm incision and created a potential space for implantation. The pre-prepared implants were inserted bilaterally, and the muscle tissue and skin were closed with 4-0 Vicryl (Ethicon). Baseline X-ray radiographs (Faxitron Bioptics, Tuscan, AZ) were taken at the time of surgery to confirm the orientation of the implant. Animals recovered from anesthesia on a heated pad and were given subcutaneous saline for hydration and Buprenorphine (0.1 mg. kg ^-1 ) for postoperative analgesia. Rats were given Enrofloxacin (~50 mg. kg ^-1 ) via their drinking water as prophylaxis against postoperative infection. Animals were monitored for any signs of declining overall health (loss of body weight, lethargy, pyrexia, poor coat condition, non-weight bearing, and inflammation of the surgical site), with the protocol specifying euthanasia in case of distress. At the study endpoint of 4 weeks postoperatively, rats were euthanized by carbon dioxide. The implants, including the surrounding muscle tissue, were harvested from the left and right legs. Results

XRD data analysis confirmed the doping of both Mg and Fe into Baghdadite (Figs. 1 and 2).

Undoped Baghdadite is characterised by an XRD pattern obtained from a copper X-ray source having the following diffraction angles 2Q: lines of strong intensity: 31.385; 31.075 and 29.940 degrees; and lines of medium intensity: 27.662; 36.045 and 36.997 degrees.

The phase peak positions of Mg-doped Baghdadite where x = 0.1 - 0.5), shifted towards lower 2-theta angles relative to undoped Baghdadite with increasing magnesium concentration. Mg-doped Baghdadite where x = 0.1) is characterised by an XRD pattern obtained from a copper X-ray source having the following diffraction angles 2Q : lines of strong intensity: 29.539, 29.959, and 31.167 degrees; and lines of medium intensity: 27.581, 36.022, and 37.023 degrees.

The shift in peak postion was accompanied by peak broadening, and changes in crystal lattice parameters (Table 1) due to the incorporation of magnesium into the crystal structure. Rietveld refinement was used to measure the lattice parameters of the Mg-doped Baghdadite.

Table 1 Lattice parameters of undoped Baghdadite , and Mg-doped Baghdadite

The results showed that all the lattice parameters, a, b, c and b decreased relative to undoped Baghdadite with increasing magnesium incorporation by 0.08%, 0.05%, 0.1% and 0.04% respectively in the Mg-doped Baghdadite where x = 0.1).

Undoped Baghdadite is characterised by an XRD pattern obtained from a molybdenum X-ray source having the following diffraction angles 2Q: lines of strong intensity: 12.607; 13.621, 14.169, and 14.307 degrees; and lines of medium intensity: 22.538; 22.966 and 24.441 degrees The phase peak positions of Fe-doped Baghdadite where x = 0.05 - 0.2), shifted towards lower 2-theta angles relative to pure Baghdadite with increasing iron concentration. The inventors surprisingly found that the doping process results in the formation of doped-Baghdadite, without the formation of undoped Baghdadite. At lower doping concentrations, where x < 0.2, monophasic doped-Baghdadite was observed. At high doping concentrations, where x > 0.2, doped Baghdadite and Merwinite (in the case of Mg-doped Baghdadite) or calcium iron silicate (in the case of Fe-doped Baghdadite) was observed. The shifts in the 2-theta values in the XRD profiles of doped Bagdhadite relative to undoped Baghdadite, indicate slight changes in the dimensions of the crystal unit cell, but still retains the crystal structure that defines the material ‘ Baghdadite’.

The mechanical and cell proliferation assessments show that substituting Ca with a small amount of Fe or Mg with x=0.05) results in 15% improvement in modulus without affecting the strength or the bioactivity of the material (see Fig. 3, and Fig. 4). On the other hand, substituting Zr with small amount of results in 17% and 35% increase in modulus and strength of the material respectively, and maintaining of the bioactivity of the material (see Fig. 5). The cell differentiation results show an increase (approx. 10%) in cell proliferation activity.

Bending studies

Given that bending is one of the most common mode of bone failure, the Mg- doped Baghdadite that showed highest potential to improve mechanics in the testing depicted in Figure 6, was tested in bending. 3-point bending was preferred over alternative bending tests such as “ring on ring” or “ball on three balls” tests, because an elastic modulus has to be assumed for the disks in these tests.

The respective stiffness and strength of Mg-doped Baghdadite (with x = 0.1) is 23% and 45% higher than Baghdadite (Fig. 7a). These properties did not significantly differ for increasing concentrations of Mg but decreased significantly for x = 1 (Merwinite). This result is consistent with previous reports showing increases to the material strength through doping of low concentration of elements to different ceramics, improvements that were attributed to crack deflection at the grain boundaries or at the interfaces between different phases.

SEM and density analysis

The fracture surfaces as well as polished and thermally etched (at 1350 °C for 15 min) surfaces of the samples was analysed using SEM (Zeiss, Germany) to search for the reasons behind improvements in mechanical properties (Fig. 7b). Examining the fracture surfaces revealed a significant difference in the void content. To quantify this, the void content as the relative surface area of the voids was approximated, which was measured on three different locations on the fracture surface of the samples using the binary images prepared with Image J (Fig. 7b: the black areas in the middle figures represent the voids). Baghdadite contained - 4.2% voids, while this decreased to -

0.03% in the Mg-doped Baghdadite with x = 0.1 (Fig 7c): > 10-fold decrease in the void content. The void content stayed low in materials with higher concentration of Mg, but increased to 2.2% in Merwinite (x = 1). The density of the samples was also measured using a pycnometer (N = 5, Fig. 2d). The void content and the density had opposite trends: the Mg-doped Baghdadite was denser than both Baghdadite and Merwinite.

The polished surfaces were analysed to characterize the size and shape of the materials’ grains. The 2D grain size (Fig. 7c) decreased with incorporation of Mg but increased significantly for Merwinite despite the same size of the homogenized precursors used for all the compositions. This is likely due to the effects of doping on decreasing in the mobility of the grain boundaries. Plotting the density, grain size, and strength against each other revealed interesting trends (Fig 7d): the density and the strength correlate positively with each other and negatively with grain size, showing that the smaller gains resulted from doping lead to lower void content and in turn higher density and mechanical properties.

Cell Screening (in-vitro analysis)

To assess the bioactivity of Mg-doped Baghdadite, in vitro testing was completed specifically quantifying the proliferative and osteogenic response of human adipose- derived stromal cells when seeded on the Mg doped Baghdadite bioceramics. The immunofluorescent expression of the RUNX2, an essential transcriptional regulator of osteogenesis necessary for commitment of multipotent mesenchymal cells to the osteoblast lineage, was reported. The endogenous subcellular translocation of RUNX2 to the nucleus was measured, as a ratio of the average fluorescent intensity within the nucleus as compared to the cytoplasm (RUNX2 nuc:cyto). Comparing RUNX2 expression across the differing levels of Mg-doping, it was observed that the maximal RUNX2 nuc:cyto expression was present for Mg-doped Baghdadite (with x = 0.1, p < 0.05), with decreases to RUNX2 nuc:cyto for further increases to the molar concentration of Mg-doping. Results are depicted in Figure 8a. Similar results have been observed during bone development wherein RUNX2 expression correlates with the increasing proliferative potential of immature osteoblasts populations. Further, for complete maturity of osteoblasts, RUNX2 must downregulate, a phenomenon that could explain the relatively lower expression of RUNX2 in osteogenic media conditions, wherein RUNX2 expression may be relatively lower, as the cells progress further towards osteogenic maturity. Interestingly, distinct changes to cell morphology can be seen comparing conditions for Mg-doping x = 1.0, where a rounded morphology was observed for all cells on this material. This rounded morphology correlated with low proliferative potential and poor osteogenic potential (Fig. 8b). Further, the flat-and-wide spread morphology of cells cultured in the osteogenic media conditions correspond with the well-known behavior of stromal cells during osteogenesis, wherein increasing cell spread area correlates with osteogenesis.

3D printing of bioceramic scaffolds

A stereolithography (SLA) 3D printing procedure was used to make porous scaffolds (with rotated cubic unit cell) made of Ca ₃ Mgo.iZro.9Si20s.9 and to compare its in-vivo performance with the scaffold with the same architecture but made with the control Baghdadite. SLA of ceramics consists of dividing a complex 3D architecture into a set of layers and depositing the layers sequentially through photochemical crosslinking of a photocurable ceramic resin, followed by sintering. The success of the process depends on several factors. Firstly, enough photopolymer must exist in the resin for successful photo-crosslinking of the layers. Secondly, the viscosity of the resin should be low enough so that the resin can flow during the printing process. And thirdly, the resin must contact enough particles with tailored size. Low concentation or large particles result in unsuccessful sintering. These requirements were satisfied by formulating a ceramic resin containing 65wt% ceramic powder refined to size D5o= 9.52 pm (using ball milling), 17.5 wt% commercial photopolymer (Form lab clear resin V4, MA, US), and 17.5 wt% dispersant (Tween 20, Sigma Aldrich). A desktop printer equipped with a resin mixer and elevated resin bath temperature was also used. This prevented particle sedimentation, or segregation of resin constituents and resulted in reduced resin viscosity due to higher bath temperature. The photosensitivity of the resulting resin and its relatively low viscosity of 4 Pa.s at 100 s ^-1 (Physica MCR 302 rheometer; Anton Paar, Austria) ensured successful printing (Fig. 9).

Scaffolds with diameter d = 4 mm, height h = 6 mm and with the same rotated cubic architecture were printed from both materials (Fig. 9b). The scaffolds had porosity = 52.4 ± 3.2 % and average pore size = 514 ± 24.3 % pm (Fig. 9b) calculated from the CT reconstruction of the samples. SEM microscopy was also used to measure pore size. The design of these similar scaffolds to achieve a certain average pore size and porosity is known in the art. The rotated cubic design is basically a cubic lattice rotated in three dimensions so that the strut that is aligned with axis x (vector [1 00]) in a cubic design is located along the vector [1 1 1] in the rotated design (Fig. 9b). This rotation strategy can be used to effectively decrease or increase the stiffness and strength of the scaffolds along certain directions. For example, the above rotation of the cubic architecture results in ~ 3.2x decrease in its stiffness and strength along the vertical direction (axis z, Fig. 8b). This strategy was used to reduce the stiffness and strength of the scaffolds so that the difference in the new bone formation in the two different materials (if any) would have a more pronounced effect on the mechanical properties post-implantation. The scaffolds with rotated cubic design were sintered following the steps used to sinter disks, with the only difference being that the material was kept at 600 °C for three hours (instead of one hour) to make sure that the organic photopolymer was completely degraded.

Mechanics of scaffolds pre- and post-implantation

The Mg-doped Baghdadite and Baghdadite scaffolds were tested under compression at lpm/s (Instron, MA, US) under in-situ optical imaging

(Canon, Japan). The top and bottom surfaces of the scaffolds were polished to make sure that the surfaces are parallel and flat as described in the Methods section. The force was measured using a 1 kN load cell and the displacements were measured from the optical images (this was crucial in reporting the actual values of strains). The top and bottom surfaces of the test fixture that were in contact with samples were covered with a thin (50μm) layer of plastic to reduce the contact stresses at the surface and prevent the immature failure of the samples under contact stresses. Figure 10a shows the representative compression stress strain curves of the scaffolds (total number of samples N = 6). Figure 10b shows the in-situ optical images at different stages of loading. Pre implantation, the scaffolds showed the expected linear elastic response followed by catastrophic failure: several cracks propagated through the material resulting in shattering into pieces and complete loss of functionality at a small strain (1.2%, Fig.

10a). These trends were the same for both materials; however, Mg-doped Baghdadite showed 42%, and 34 % higher stiffness and strength respectively compared with Baghdadite (Fig. 10c). This difference was expected because the rotated cubic architecture is a bending dominated architecture, meaning that while the scaffolds were tested in compression, it is the bending stresses developed inside the material that dominate their deformation and failure. As a result, the properties of the scaffolds with the rotated cubic design followed the same trends observed in the bending of the monolithic beams.

Post-implantation, the behavior was completely transformed because of the new bone formation. Although the maximum stress was again achieved at the same strain (1.2%) where the struts started to break, the failure was not catastrophic: the materials showed gradual failure and a triangular stress-strain curve with a long tail typical of tough materials was observed. Both Baghdadite and Mg-doped Baghdadite lost only - 42% of their load bearing capability at 5% strain (~ 5x higher than the strain at failure for the scaffolds) (Fig. 10a). The in-situ images showed that the material was still in one piece and capable of resisting load (up to 40% of the maximum load) at very high strain of 8% (Fig 10b). As a result, the energy absorption of the Mg-doped scaffolds calculated from the area under the force-deflection curve improved 17-fold post-implantation. While this is impressive, the actual level of improvement in energy absorption might be even higher due to three reasons:

(1) for a brittle material like the ceramics before implantation, the area of the force-deflection curve overestimates the actual energy dissipation capability of the materials because a lot of the energy stored in the material is dissipated through dynamic effects (such as generation of sound) upon failure. The area under the curve for the material post-implantation however reflects the actual energy dissipation in the material as the material gradually fails and dynamic effects are not present.

(2) Previous studies showed that formaldehyde fixation (that was done on the scaffolds) has minimal effect on the quasi-static properties of bone. If the properties of the new bone are slightly degraded by the fixation process, the actual improvements could be even more significant than found here.

(3) The area under the curves were calculated up to the strain at which the stress was decreased to 40% of the maximum value (the tests were stopped at this point) but the scaffolds post-implantation were still resisting deformation at this point and capable of absorbing/dissipating more energy.

The high level of improvement in energy absorption shows the significance of the new bone formation on transformation of the behavior of these naturally brittle scaffolds. The stiffness and strength of these scaffolds were also improved by 24% and 38% respectively. Similar trends were observed for baghdadite scaffolds (which underwent the same fixation process as the Mg-doped ones); however, lower levels of improvement in the mechanical properties were achieved (Fig. 10a, c): 10%, 18% and 12x improvements were achieved respectively for stiffness, strength and energy absorption used the Mg-doped material. Optimisation of Mg-doped Baghdadite scaffolds

In the results discussed above, the stiffness and strength of the Mg-doped Baghdadite scaffolds were improved compared with Baghdadite. However, they were still lower than cortical bone (stiffness ~ 1.1 GPa and strength ~ 7.5 MPa, both 10 -15 times less than cortical bone). To improve the properties of the scaffolds and reach values closer to those of cortical bone making them even better adapted for load bearing applications, the following adaptations were made:

(a) changes in the architecture,

(b) reduction in the porosity; and

(c) optimization of the sintering parameters. An architecture based on a solid with cylindrical holes (i.e., a negative replication of a strut-based design, as depicted in Fig. 11a was used. Studies on the architectural effects have shown that the internal stresses within these scaffolds can be 4 times lower than the rotated cubic architecture resulting in 3-4 times improvement in strength (that would result in strength ~ 30 MPa). The porosity was decreased from ~ 52% to - 35% (27% difference). The strength of porous ceramics is known to increase logarithmically with a decrease in porosity. It was therefore hoped that this would result in a 27% or more improvement in strength by this reduction (that would result in strength of at least - 38 MPa). It was found that by making these two modifications, the strength of the scaffolds reached 63 MPa (Fig. 1 lb). The sintering parameters were also optimized by increasing the sintering temperature from 1400 °C to 1480 °C (Fig. lib).

To ensure the accuracy of the temperature in the furnace, the temperature was measured using an independent thermocouple (in addition to the in-built one), and consistent temperature readings were found. To make the material stable at this temperature, CaO were used crucibles. This was because the Portland cement crucible (used for the previous experiments) started to strongly adhere to samples at 1480 °C, resulting in sample melting at crucible contact points. XRD characterization of the material showed that the crystal structure of the material was not affected by sintering at a higher temperature. Sintering at 1480 °C increased the density of the material relative to the material sintered at 1400 °C. Density was measured using a pycnometer. A 45 % increase in the mechanical properties of the scaffolds were achieved (Fig. lib). After these improvements, the stiffness and strength of the Mg-doped Baghdadite scaffolds (with 35% porosity) reached those of cortical bone (Fig. 11). As expected, the stiffness and strength of Baghdadite scaffolds with the same architecture and processing parameters were lower than Mg-doped Baghdadite (Fig. 11c).

Although the invention has been described with reference to specific examples, it will be appreciated by those skilled in the art that the invention may be embodied in many other forms. In particular features of any one of the various described examples may be provided in any combination in any of the other described examples.

ITEMISED LIST OF EMBODIMENTS

Item 1. A material comprising doped Baghdadite, wherein the doped Baghdadite is selected from Mg-doped Baghdadite, Fe-doped Baghdadite, and [Mg, Fe]- doped Baghdadite.

Item 2. The material according to item 1, wherein the doped Baghdadite is Mg-doped Baghdadite.

Item 3. The material according to item 1, wherein the doped Baghdadite is Fe-doped Baghdadite.

Item 4. The material according to any one of items 1 to 3, wherein the doped Baghdadite has a formula wherein 0 < x < 1 [Formula A] or wherein 0 < x < 3 [Formula B].

Item 5. The material according to item 4, wherein in Formula A, z = 0.

Item 6. The material according to item 5, wherein x = 0.05, 0.1, 0.2, 0.3 or 0.5.

Item 7. The material according to item 5 or 6, wherein x = 0.1.

Item 8. The material according to item 4, wherein in Formula B, z = 0 or y = 0.

Item 9. The material according to item 8, wherein when z = 0, x = 0.05, 0.1, 0.2 or

0.5.

Item 10. The material according to item 8, wherein when y = 0, x = 0.05, 0.1, 0.2 or 0.5.

Item 11. The material according to any one of items 1 to 10, wherein the material is a biocompatible ceramic material.

Item 12. The material according to any one of items 1 to 11, wherein the material is monophasic.

Item 13. The material according to any one of items 1 to 12, wherein the material does not comprise undoped Baghdadite.

Item 14. The material according to any one of items 1 to 13, wherein the material forms a hydroxyapatite layer upon exposure to bodily fluids.

Item 15. A method of preparing the material according to any one of items 1 to 14, the method comprising: providing a sol of precursor materials; at least partially gelling the sol; and drying and sintering the at least partially gelled sol to thereby form the material. Item 16. The method according to item 15, wherein the precursor materials are selected from the group consisting of:

Item 17. An implantable medical device comprising the material according to any one of items 1 to 14.

Item 18. The implantable medical device according to item 17, wherein the implantable medical device is selected from the group consisting of: a 3D implantable scaffold, an orthopaedic implant, a dental implant/prosthesis, a spine implant, an implant for craniofacial reconstruction or alveolar ridge augmentation, an implant for cartilage regeneration, an osteochondral defect implant, a stmt, a stent and a stent-graft.

Item 19. The implantable medical device according to item 17 or item 18, wherein the implantable medical device is permanently implanted.

Item 20. The implantable medical device according to any one of items 17 to 19, wherein the implantable medical device is substantially biodegradable.

Item 21. The implantable medical device according to any one of items 17 to 20 comprising a porosity of between about 10 to about 80%.

Item 22. The implantable medical device according to item 21 having a pore size of between about 20 to about 500 micron.

Item 23. The implantable medical device according to any one of items 17 to 22, wherein the compressive strength of the implantable medical device is between 100 and 250 MPa.

Item 24. The implantable medical device according to any one of items 17 to 23, wherein the modulus of the implantable medical device is between 10 and 20 GPa.

Item 25. The implantable medical device according to any one of items 17 to 24, wherein the implantable medical device is coated with at least one resorbable polymer material selected from the group consisting of: a polyglycolide, a polydioxanone, a polyhydroxyalkanoate, a polylactide, an alginate, a collagen, a chitosan, a polyalkylene oxalate, a polyanhydride, a poly(glycolide-co- trimethylene carbonate), a polyesteramide, and a polydepsipeptide.

Item 26. The implantable medical device according to any one of items 17 to 25, wherein the implantable medical device is coated with at least one healing promoter selected from the group consisting of: a thrombosis inhibitor, a fibrinolytic agent, a vasodilator substance, an anti-inflammatory agent, a cell proliferation inhibitor, and an inhibitor of matrix elaboration or expression.

Item 27. The implantable medical device according to any one of items 17 to 26, wherein the implantable medical device is a bone implant.

Item 28. The implantable medical device according to any one of items 17 to 26, wherein the implantable medical device is a tooth filling implant.

Item 29. The implantable medical device according to any one of items 17 to 26, wherein the implantable medical device is a biocement.

Item 30. A method for improving the long-term stability of an implantable medical device, the method comprising: coating the implantable medical device with the material according to any one of items 1 to 14.

Item 31. The method according to item 30, wherein the method further comprises: coating the implantable medical device or the material with a biocompatible polymer.

Item 32. The method according to item 30 or item 31, wherein the implantable medical device is a biphasic scaffold for an osteochondral defect.

Item 33. A method for regenerating or resurfacing tissue, the method comprising: contacting the regenerating or resurfacing tissue with a quantity of the material according to any one of items 1 to 14 for a sufficient period to at least partially effect said regeneration or resurfacing.

Item 34. A method for forming osseous tissue on an orthopaedic defect, the method comprising: contacting said defect with the material according to any one of items 1 to 14.

Item 35. A method of preparing the material according to any one of items 1 to 14, the method comprising: providing a resin comprising ceramic powder; depositing layers of resin comprising ceramic powder using a stereolithography or 3D printing device; crosslinking the layers and sintering to thereby form the material.