Hydroxyapatite-Titanium Bulk Composites for Bone Tissue Engineering Applications: A Review Alok Kumar, 1 Krishanu Biswas, 2 and Bikramjit Basu1* 1 2
Laboratory for Biomaterials, Materials Research Centre, Indian Institute of Science, Bangalore-560012, India
Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kanpur-208016, India
Abstract The research work on bulk hydroxyapatite (HA)-based composites are driven by the need to develop biomaterials with better mechanical properties without compromising its bioactivity and biocompatibility properties. Despite several years of research, the mechanical properties of the HA-based composites still need to be enhanced to match the properties of natural cortical bone. In this regard, the scope of the present review on the HA-based bulk biomaterials is limited to the processing and the mechanical as well as biocompatibility properties for bone tissue engineering applications of a model system i.e. hydroxyapatite-titanium (HA-Ti) bulk composites. It will be discussed in this review how HA-Ti based bulk composites can be processed to have better fracture toughness and strength without compromising biocompatibility. The advantages of the functionally gradient materials (FGMs) to integrate the mechanical and biocompatibility properties is a promising approach in hard tissue engineering and has been emphasized here in reference to the limited literature reports. On the biomaterials fabrication aspect, the recent results are discussed to demonstrate that advanced manufacturing techniques, like spark plasma sintering can be adopted as a processing route to restrict the sintering reactions, while enhancing the mechanical properties. Various toughening mechanisms related to careful tailoring of microstructure are discussed. The in vitro cytocompatibilty, cell fate processes as well as in vivo biocompatibility results are also reviewed and the use of flow cytometry to quantify in vitro cell fate processes is stressed upon. Keywords: Hydroxyapatite; Titanium; Composite; Fracture toughness; Biocompatibility corresponding author: Email: [email protected], Tel.: +91 80 2293 3256; Fax: +91 80 2360 7316
This article has been accepted for publication and undergone full peer review but has not been through the copyediting, typesetting, pagination and proofreading process which may lead to differences between this version and the Version of Record. Please cite this article as an ‘Accepted Article’, doi: 10.1002/jbm.a.35198
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Contents: 1. Introduction 2. Physical and mechanical properties required for bone materials 3. Orthopedic grafting using hydroxyapatite-based composites 4. Processing of HA-Ti composites 4.1. Conventional sintering 4.2. Advanced sintering 5. Toughening mechanisms and toughness properties 6. Functionally graded materials 7. Biocompatibility properties of HA-Ti composites 7.1. In vitro bioactivity 7.2. In vitro cytocompatibility 7.3. In vivo biocompatility 8. Closure References
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1. Introduction: Healthy bone and joints (e.g. articulating, knee, and hip joints) are necessary for the structural stability and pain free movements. However, these healthy bones and joints can get damaged in some unfortunate cases due to accident / diseases and fail to perform their normal functions. The damaged bones can be replaced or repaired with allogeneic bones or allografts
1-3
. However, the possible risk of transmission of life threatening infections
with allogeneic bones as well as inferior osteoconductivity and poor mechanical properties of the freeze dried bones limit the application of such types of biomaterials
4-6
. In contrast, the designed biomaterials with bone-
mimicking properties can be used as an alternative to the allografts and autogafts. Biomaterials development has a long history. In 1940, Jean and Robert Judet of Paris performed the replacement arthoplasty to replace the femoral head with acrylic resin, which opens up the possibility of joint replacement 7. However, after an initial success, this implant failed due to its poor mechanical properties as well as inapt design. More than 10 decades later, when most of the issues related to joint replacement have been settled down, the interest in new biomaterials remain as fierce as ever. A biomaterial is a synthetic biocompatible material/device that has been intently designed to persuade a definite activity in the biological system, which can simulate the desired biological function, without generating a short and/or long term damage to the nearby or distant tissue as well as without being getting damaged in this procedure 8,9. Depending on the in vivo performance, different biomaterials can be subcategorized as, bioactive, bioinert, and bio-tolerant. Bioactive materials (CaP-based materials) can promote new bone growth under physiological conditions. In addition, the bony tissue will form around the bioactive implant and adhere strongly due to osteogenesis
10,11
. Bioinert materials (e.g. Ti, Ti alloys, Al2O3 etc.) exhibit very minimal interaction with the
surrounding tissue. Generally, a protective fibrous capsule of the connective tissue may form around the implant by the contact osteogenesis
12,13
. In case of the bio-tolerant materials (e.g.: stainless steel, PMMA etc.), a
capsule of connective tissue will form by distant osteogenesis, and this capsule does not cling to the implant 1315
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On the other hand, depending on the stability of a synthetic implant in the host environment, the implant materials can be categorized into biodegradable, bioresistive as well as bioresorbable. A biodegradable material degrades (or dissolves) gradually in the host environment. In contrast, the bioresistive biomaterials do not degrade under the effect of host environment, e.g. HA, which has very low solubility in alkaline solution
11
.
Bioresorbable materials are the CaP-based biomaterials (α and β-tricalcium posphates), which provide the therapeutic agent (e.g. Ca2+, PO43-) during controlled dissolution, essential for faster healing of bone injury. For many biomedical applications, a biomaterial should show the bioactivity, osteoconductivity, and osteoinductivity. Bioactivity is the property of implant material to make direct bonding with living bone
16,17
.
The bonding is normally mediated by an intervening (apatite) layer formed between the implant and the bone 18. On the other hand, osteoconductivity is defined as the ability of an implant material that can serve as a scaffold for new bone formation and therefore, promotes the osteoblast adhesion, proliferation, and differentiation
19,20
.
Osteoinductivity is the ability of a biomaterial to induce the new bone formation, implanted at non bone forming site, in vivo 21. It is well documented in case of bone tissue engineering applications that the success of an implant material depends on the biocompatibility, osseointegration property, physical/mechanical properties and antibacterial property etc.
22-25
. Metals have an advantage of good mechanical properties, like high fracture toughness and
flexural strength. But the corrosion as well as wear of the implant (results in leaching of metal ions) and stress shielding results in irritation, inflammation, and loosening of implants metal implants result in weak bonding with host bone, in vivo
13,15
26,27
. Additionally, bioinert nature of
. Although ceramics based biomaterials are
found suitable for hard tissue applications due to its biocompatibility, corrosion resistance, and compressive strength properties
28,29
, inherent brittleness of ceramics may result in the early stage failure of the implants
during load bearing applications
30
. A classic example of a brittle bioceramic is hydroxyapatite, which has
widely been researched due to its structural similarity with major inorganic component of bone composition 30.
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Bone can be considered as a biocomposite of HA and collagen and thus synthetic bulk composites made by these constituents can mimic the bone composition more closely. It is worthwhile to mention here that there are many inexpensive methods available that allow the hydroxyapatite synthesis in bulk quantity
31,32
. However,
due to weak mechanical properties as well as processing related challenges at high temperature (high temperature sintering leading to the denaturing of collagen), HA-collagen based composites can not be used for the load bearing applications. It is important to mention here that, to ensure better mechanical properties, HA needs to be sintered at above than 950 °C
33
. Although materials, like Co-Cr alloys, Ti and its alloys, and
stainless steels are extensively used in bone-tissue engineering applications, the expedition to search an ideal material for such application still needs extensive efforts 13. Despite good toughness, metals, like Co-Cr alloys, Ti and its alloys, and stainless steels do not considered reliable in complex physiological environment for longterm usages due to leaching of toxic ion (e.g. Cr, Ni) under physiological conditions as well as bioinert/biotolerant nature
12-15,34-36
. The application of these metallic implant materials can be extended by coating the
bioinert/bio-tolerant surfaces with bioactive materials, like calcium phosphate
37-39
. However, the poor
interfacial strength between base material and coating as well as fracture of coating durings the service period may result in serious clinical issues
40,41
. Another issue with metallic implants is the generation of metal ions
due to dissolution and/ or wear during the long service period main cause of local necrosis and/or inflammation
43
42
. The generation of nano-sized debris is the
. These wear debris may also be carried out by blood to
other organs, like heart, lever, and lung and this leads to the severe toxicity of these organs. In the context of load bearing implants, high elastic modulus of the implant with respect to the neighboring bone causes stress shielding, being responsible for the bone resorption near to the implant
44-48
. Therefore, the
elastic modulus of an implant ideally needs to be matched with bone, while efforts are to be made to enhance both fracture toughness and strength. In the above backdrops, the hydroxyapatite (HA) - based metal or ceramic reinforced composites can be used as a potential bulk implant material, where the second phase contributes specifically towards better mechanical properties. The choice of the second phase is limited as it needs to have good mechanical property as well as 5 John Wiley & Sons, Inc.
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biocompatibility. As compared to Al2O3, ZrO2, and stainless steels, titanium (Ti) is found to be more promising second phase material for orthopedic applications due to its excellent corrosion resistance property under physiological environment, low density (4.540 g/cc) as well as the lower elastic modulus (~110 GPa) than CoCr alloy (~230 GPa)/ stainless steel (~205 GPa) 22,48. The design of the HA-Ti composites underlies the fact that the Ti addition significantly improves the mechanical properties of the composites without degrading the cytocompatibility properties 49-51. In the present review, the processing of HA-Ti composites using conventional and advanced sintering route has been discussed. The improvement in mechanical properties due to Ti addition has been elaborated. The results of in vitro and in vivo experiments have also been provided to justify that the biocompatibility property is not compromised due to Ti addition to HA.
2. Physical and mechanical properties required for bone materials: Apart from the high dissolution rate under physiological environment, inferior corrosion resistance and poor tribological properties as well as inferior bactericidal property, an implant may fail during the service period due to, (a) mismatch in elastic modulus between the implant and the surrounding bone and (b) poor fracture toughness 22. Therefore, we will focus on these aspects in the following sections. A lower load experienced by the host bone can result in stress shielding due to comparatively higher elastic modulus of the implant than bone
52
. This leads to the resorption of bone near the implant due to remodeling
followed by loosening of the implant
44-47
. The problem of stress shielding in a particular application can be
minimized using natural grafts, or choosing an appropriately designed material. For example, three dimensional scaffolds of biocompatible materials having gradient porosity with controlled shape, size and distribution is a promising solution in such application 53,54. As compared to Al2O3 (390 GPa) and ZrO2 (205 GPa), elastic modulus of HA (80-110 GPa) was found to be closer to the cortical bone (14-20 GPa)
22,55-61
. However, the fracture toughness of HA (0.7-1.2 MPa.m1/2) is
lower than the even lower limit of the cortical bone (2-12 MPa.m1/2)
55-58
. As mentioned earlier, the poor
fracture toughness restricts the use of HA in the load bearing applications. In contrast, Ti with exceptionally 6 John Wiley & Sons, Inc.
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high fracture toughness (60 MPa.m1/2) is found to have lower elastic modulus (117 GPa) than Al2O3 and ZrO2. Therefore, the HA-Ti system can be considered as a model system. Apart from biocompatibility properties, the presence of the second phase can affect other properties, like fracture toughness, strength, and elastic modulus. Since, we are more concerned about the fracture toughness of bioceramics materials, a detailed description on the effect of second phase on fracture toughness of HA has been provided in section 5.
3. Orthopedic grafting using hydroxyapatite-based composites: The natural bone is a composite of hydroxyapatite and collagen. Therefore, a biomaterial with HA and collagen (e.g. HA-collagen composites) can provide the suitable surface for higher cell adhesion, proliferation and osseointegration 62. However, due to poor mechanical properties (e.g. fracture toughness), the use of HA-based composites reinforced with polymer has been restricted only for low load-bearing applications
63
. As will be
discussed later, this has been the driving force for the research on HA-based biocomposites. From the perspective of material development, the processing window must always be chosen on the basis of base material. For load bearing application, the base material (e.g. HA) has to be densified up to maximum extent, which is only possible if the material can be sintered at temperature ≥ 950 oC
33,64
. Nevertheless, at this
temperature, most of the polymers are thermally unstable and high temperature sintering leads to the complete removal/denaturing of most of the polymers, like collagen
65,66
. In contrast, innovative processing routes, like
freeze casting may allow the fabrication of HA-polymer composites with porous structure 67,68. To develope a bone replacement material, the choice of the second phase depends on the application as well as anticipated functionality. For example, BaTiO3, CaTiO3, Fe3O4, and ZnO or Ag can be added to HA as a second phase to introduce the piezoelectric, electrical conductivity, magnetic, and bactericidal properties, respectively 69-73
. However, brittle nature and the load bearing incapability of the ceramics may limit the broader application
of such types of ceramic-ceramic composites; some examples include HA-20HAwhisker, HA-20ZrO2, and HA(Al2O3 coated) ZrO2 with fracture toughness of 1.4 MPa.m1/2, 1.5 MPa.m1/2, and 3 MPa.m1/2, respectively 74-76. Interestingly, a glaring example of HA-based ceramic composite with fracture toughness, higher or equal to that 7 John Wiley & Sons, Inc.
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of the cortical bone is still unavailable (see figure 1). In contrast, the incorporation of metallic reinforcement as second phase is deemed to improve the fracture toughness of the HA-based composites. However, the sintering reactions between HA and metal phase are of major concern and can affect the biocompatibility properties in vitro and in vivo. Although stainless steel and Co-Cr alloys are competitively better than ceramics counterparts in terms of fracture toughness, the high elastic modulus of metals however may lead to stress shielding during in vivo performance 25. In contrast, Ti was found to be most suitable for biomedical applications due to excellent biocompatibility, corrosion resistance property and stability in physiological conditions
77
. Thus, in order to
incorporate both load bearing capabilities without fracture under peak loading as well as osseointegration property, the HA-Ti composites can be used as a bulk biomaterials. Over the various composites (HA-BaTiO3, HA-CaTiO3, HA-Al2O3, HA-ZrO2, HA-ZnO, HA-Ag), the selection of HA-Ti composites as a model system in the present review can be rationalized on the basis of following aspects: (a) low density of HA (3.16 g/cc) 63 and Ti (4.54 g/cc); (b) high fracture toughness of metallic Ti (60 MPa.m1/2)
22
; (c) bioactivity of HA and bio-
inertness of Ti; (d) corrosion resistant properties of Ti; (e) limited dissolution of HA under physiological condition 78.
4.
Processing of HA-Ti composites:
HA-Ti composites can be made either by conventional pressureless sintering or advanced sintering, such as spark plasma sintering (SPS). SPS has the advantage of higher heating rate and shorter soaking time over the conventional sintering. The advanced sintering techniques have many other advantages as compared to the conventional sintering and have been described in section 4.2. Traditionally, the HA-Ti composites were prepared by the homogeneous mixing of the HA and Ti powders using a ball-mill, followed by green compaction followed by pressureless sintering at relatively higher temperatures (≥ 1100 °C) and soaking time (≥ 1 hour) 64. Recently, the FAST (field assisted sintering technique) has been utilized to prepare highly densified samples from the powder mixture within a short time 33,79,80.
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4.1.
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Conventional sintering:
In this subsection, the literature reports on conventional sintering of HA-Ti are briefly discussed. Fahami et al. 81
used CaHPO4, CaO, and Ti powders as precursors to prepare the HA-Ti composites by the mechanochemical
synthesis of a mixture of HA and Ti according to the following reaction:
6CaHPO4 + 4CaO + Ti → Ca10 ( PO4 )6 (OH ) 2 + Ti + 2 H 2O
-------(1)
However, mechanochemical synthesis leads to formation of HA with reduced crystallinity. It is important to note that the lower crystallinity of the HA powder results in the increased bioresorption and solubility under physiological conditions
82
. Therefore, further attempt has been made by Fahami et al.
81
to increase the
crystallinity from 13 to 69 % by annealing the ball-milled powder at 650 oC for 2 hours, synthesized after reaction (1). However, this results in the oxidation of Ti and decomposition of HA to β-TCP (beta- tri-calcium phosphate) phase. In a different approach, Chu et al.
83
prepared the HA-20 wt.% Ti composite by mixing HA
and Ti powders in a ball mill, followed by hot pressing at 900 oC and 1000 oC. Importantly, the sintering at 1000 oC leads to the formation of α-TCP and Ca4O(PO4)2 phases. The difference in the linear thermal expansion coefficient (α) of HA (17.3×10-6 oC-1)
84
and Ti (8.4×10-6 oC-1)
85
phases results in inferior sinterability in case
of HA-20 wt.% Ti composite with ~67% relative density, achieved when sintered at 1200 oC 83. However, the same sintering conditions result in ~88% relative density in monolithic HA. Importantly, the addition of Ti to HA leads to an increase in fracture toughness from 0.7 MPa.m1/2 to 1.0 MPa.m1/2. Although Ti addition to HA was intended to increase the fracture toughness significantly, the marginal improvement in the fracture toughness can be attributed to the poor sinter density of the composite. In an earlier work, Nath et al. 64 prepared the HA-Ti composites with varying amount of Ti (10-40 wt.%) using pressureless sintering. It has been reported that the sintering of different compositions at 1000-1400 oC for 2 hours leads to the extensive sintering reactions and therefore the formation of TiO2, CaO, TCP and CaTiO3 (Table 1). The thermodynamic calculations show the higher probability of oxidation of Ti during conventional sintering due to higher sintering temperature and
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longer holding time 64. Such calculations also indicate the possibility of a reaction between HA and TiO2 above 750oC. Considering the thermodynamical aspect, Yang et al.
86
prepared the HA-10 wt.% Ti and HA-20 wt.% Ti
composities by sintering at 1100 oC in vacuum to minimize the oxidation of Ti. The FT-IR and XRD analysis of the HA-Ti samples show the decomposition of HA to α-TCP, TTCP (tetracalcium phosphate) as well as Ca2Ti2O5. As the sintering temperature further increases, HA is reported to decompose to HA with deficiency of hydroxyl group in the following way 87,88: o
vacuum , T >800 C Ca10 ( PO4 )6 (OH ) 2 → Ca10 ( PO4 )6 (OH ) 2− 2 x Ox + xH 2O ↑ -------(2)
The water vapor produced from the reaction (2) reacts with Ti to form TiO2 [see eq. (3)], which can finally react with HA to form β-TCP and CaTi2O5 [see eq. (4)].
Ti + 2 H 2O → TiO2 + 2 H 2 -------(3)
Ca10 ( PO4 )6 (OH ) 2 + 2TiO2 → 3Ca3 ( PO4 )2 + CaTi2O5 + H 2O ----(4) Thus, it is not impossible, but extremely hard to retain Ti in the sintered product using conventional sintering routes. Alternatively, the HA can form solid solution with metal oxides, carbonates and halides
89
. The Ca2+
ions can also be substituted with Ti4+, Na+, K+, Sr2+, Ba2+, Pb2+ and Y3+ 90-93, while OH- ion can be replaced with CO32-, F- and Cl- ions 94-96. Even PO43- ions can be replaced with AsO43-, and VO43- 97. In order to investigate the influence of the Ti ions on the stability of the structure of HA, Ribeiro et al.
89
incubated HA powder (pre-
heated at 1000 oC) in a solution containing Ti4+ ions. During incubation, Ti4+ ions get adsorbed on the HA surface and as a result, the Ca2+ ions get released into the solution. In this case, Ca2+ ions are replaced with Ti4+ ions according to the following reaction:
Ca10 (PO4 )6 (OH)2 (solid) + (n / 2)Ti4+ → (Ca10−nTi(n / 2))(PO4 )6 (OH)2 + nCa2+ + (n / 2) where, symbol
...
.....(5)
is a vacancy
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The above reactions result in a reduction in the lattice parameter of Ca-deficient HA due to smaller ionic radius of Ti4+ (0.68 Å) than Ca2+ (0.99 Å). The sintering under a protective atmosphere (e.g. argon) is another promising approach towards retaining the metallic Ti in the sintered product. In order to see the effect of protective environment on the sintering reaction in the HA-Ti composites, Ning et al. 98 conventionally sintered the HA-50Ti composite at 1200 oC for 30 min in argon atmosphere. However, the sintering leads to the formation of Ti2O, CaTiO3, CaO and TiP-like phases along with α-Ti. In summary, conventional sintering of HA-Ti composites even at very low temperature (1000 °C) leads to the extensive dehydroxylation of HA as well as the reaction between HA and Ti. The formation of TCP phases will result in higher dissolution of the implant materials and therefore, decreased mechanical strength during the service period, in vivo. As mentioned earlier, it is extremely difficult to retain the metallic Ti in the composite during the conventional sintering. The sintering in protective atmosphere of Ar was also not found to be helpful in retarding the oxidation reactions in conventional sintering due to the longer sintering duration (≥ 1 hour) and higher holding temperature (≥ 1000 °C). In the following, we shall discuss the applicability of advanced sintering techniques (e.g. SPS) to prepare the composites with predominant retention of metallic Ti.
4.2.
Advanced sintering
Among various advanced sintering techniques, field activated sintering (FAST) is widely investigated to resolve the issues related to the conventional sintering, discussed earlier. For example, Nakahira et al.
79
reported the pulse electric current sintering (PECS) of HA-20 vol.%Ti and HA-25 vol.%Ti composites at 800 C and 1000 oC under 30 MPa in a protective argon atmosphere. Although PECS suppresses the oxidation of Ti at 900 oC, the formation of Ti2O was found in samples sintered at 1000 oC. Mondal et al.
99
used the spark
plasma sintering technique to sinter β-TCP and Ti powders at 1200 oC temperature under 50 MPa pressure. Despite the retention of β-TCP and Ti, this results in the formation of CaTiO3. 11 John Wiley & Sons, Inc.
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In a recent work, the present authors have reported the results of extensive study on spark plasma sintering and phase evolution of HA-xTi (x = 5, 10, 20 wt%) composites
33
. The processing conditions were intelligently
tailored to achieve high toughness with the minimum sintering reaction between the HA and Ti. The XRD analysis confirms the predominant presence of HA and Ti. However, the noticeable diffraction peaks corresponding to β-TCP, CaO, and CaTi4(PO4)6 were also recorded. FT-IR results show the reduction in the absorption peak intensity corresponding to OH- stretching (3574 cm-1) as well as OH- liberation (636 cm-1), confirming the dehydroxylation of HA at a minute level. The EPMA study indicates the presence of a thin diffusion layer around the Ti particles with a composition different than that of Ti and HA phase. XRD, EPMA and FT-IR findings were supported by the detailed TEM investigations. The analysis of the selected area electron diffraction (SAED) patterns from different phases confirms the formation of CaTi4(PO4)6 at the interface of HA and Ti phases, due to the interfacial reaction between β-TCP and Ti
33
. The β-TCP has
been formed due to local dehydroxylation of HA. In summary, spark plasma sintering at relatively lower processing temperature with a shorter holding time than conventional pressureless sintering can enable maximum retention of HA and Ti with minimal sintering reactions.
5. Toughening mechanisms and toughness properties Fracture in brittle solids is related to the propagation of cracks under the tensile loading. Importantly, the crack propagates if and only if the stress at crack front is higher than a critical value. In brittle solids, the resistance against the crack propagation can be enhanced by introducing the mechanisms that reduce the driving force for the crack propagation, such as dissipating the additional energy associated with crack tip leading to the blunting of the crack tip 100. In polycrystalline ceramic materials, the microstructure dependent fracture toughness can be correlated with the dissipation of the crack tip energy and thereby with the reduction in the driving force for the 12 John Wiley & Sons, Inc.
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crack propagation
101
. It is generally recognized that weak interfaces, such as grain boundaries and/or matrix-
reinforcement boundary are the preferred crack deflection paths. Moreover, the presence of the second phase disturbs the crack propagation by forcing the crack either to go around it or through it. The extent to which the it is able to dissipate the energy depends on the morphology / distribution and mechanical properties of the second phase in ceramic matrix composite 102. Although the presence of second phase may improve the fracture toughness of composites, the chemical reaction at the interface during processing (e.g. sintering) always offers the challenge for the materials scientists to optimize the toughness. On the basis of microstructural characteristics of the second phase; i.e., nature, shape and size, the toughened composite materials can be categorized as 103: (a) particulate reinforced composite, (b) whisker reinforced composite, and (c) fiber reinforced composite. In the context of the current review paper, the toughening in metallic particulate reinforced ceramic composites is attributed to the shielding effect to the crack front propagation due to ductile metal reinforcements. In such case, the crack resistance energy of the matrix can be given by 101,
Rο = (1 − V f )2γ B .........(6) where, γ B is the intrinsic cohesive energy of the matrix material and V f is the volume fraction of the second (reinforcement) phase. Thus, both the matrix as well as the reinforcement material contributes towards the fracture toughness and hence, enhanced fracture toughness can be achieved by intelligently designing the composites having mixed shielding mechanisms. In case of the polycrystalline monolithic ceramic materials, the frontal zone shielding results due to the interaction of the crack front with dislocation cloud and/or microcrack cloud. The activation of the dislocations results in generation of dislocation cloud and this restricts the crack front propagation by applying a stress field on the crack tip
101
. The stress field generated around the dislocation is proportional to the shear modulus (G)
and can be given as 104, 13 John Wiley & Sons, Inc.
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σ∝
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Gb ..........(7) r
where, G, b, and r are the shear modulus of the material, Burger vector, and distance from the dislocation, respectively. On the other hand, the microcrack cloud can be activated in the stress field of primary cracks at the relatively weaker zone, like grain boundaries and interphase flaws. Unlike the combination of dislocation as well as microcrack clouds towards the improvement in fracture toughness in brittle ceramics, the phase transformation in zirconia (ZrO2) can also contribute to the fracture toughness improvement in ZrO2 reinforced ceramic matrix
101
. In such case, the partially stabilized ZrO2 in
metastable tetragonal state transforms to stable monoclinic phase at the crack tip stress field transformation can be controlled by using additives, like MgO, CaO, Y2O3 and CeO2
100
100
. This
. In addition, the grain
bridging of crack in monolithic polycrystalline materials results in the toughness improvement due to the dissipation of crack energy as a result of grain sliding and/or fracture. For example, Swanson et al. 105 reported the evolution of grain bridging in pure polycrystalline alumina. In such case, the transgranular fracture may occur in large grains along with primary crack due to high frictional stresses at sliding grain boundary surfaces. In 1983, Faber and Evans studied the effect of geometrical aspect of the second phase on the crack deflection and therefore, effect on the fracture toughness 106,107. The strength of the interfacial boundary plays an important role in deciding whether a crack will pass though second phase or get deflected. In a case when the interface is strong enough to deflect the crack, it penetrates through the second phase. This may result in pinning of the advancing crack front at the interaction site. In case of ductile metal phase, this results in the dissipation of crack tip energy due to crack bridging. In the context of biomaterials, the benefits of the second phase addition to a ceramic matrix are twofold: (1) second phase improves the fracture toughness dramatically due to frontal wake and/or bridge interface shielding 33
and (2) second phase leads to the modification in surface energy
108
, which can be used as a tool to
manipulate the cell-material interaction. 14 John Wiley & Sons, Inc.
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Although the reinforcement of ceramic matrix with fiber or whisker results in improvement in fracture toughness, the release of nanosized and needle shaped whisker/fibrous can lead to the severe toxicity to the cells caused by lipid layer damage and genotoxicity
43,109
. Therefore, the fiber/whisker toughened ceramic
composites are not considered for clinical applications, although related toughening mechanisms can be very well adapted to develop novel HA-based high toughness composites. In parallel, the reinforcement of ceramic matrix with ductile metal particles results in dramatic improvement in mechanical properties. The crack bridging at metal particles results in the dissipation of energy associated with crack tip
33,101
. The fracture
energies of the pulsed electric current sintered (PECSed) HA-Ti composites, measured by Nakahira et al.
79
using single edge double notched beam method, show a substantial increment with the Ti content. The fracture energies of 9 J/m2 and 20 J/m2 were reported for HA-20 vol.% Ti and HA-25 vol.% Ti samples, respectively. The reason for the improvement in fracture toughness was crack deflection and crack bridging by ductile Ti particles, which are retained in the HA matrix. Similarly, substantial improvement in the fracture toughness of the spark plasma sintered HA-Ti composites were reported by the present authors (see Figure 1) 33. An addition of 10 wt.% Ti to HA increases the fracture toughness (KIC) from ~3.4 MPa.m1/2 to 4.7 MPa.m1/2. Further addition of Ti (20 wt.%) to HA does not lead to any further improvement of the fracture toughness (4.3 MPa.m1/2). The increased fracture toughness of HA and HA-Ti composites can be explained on the basis of crack deflection and crack bridging. In the context of toughness enhancement, the synergistic interaction of multiple toughening mechanisms can be explored. The presence of Ti particles provides the significant contribution toward the fracture toughness by crack bridging. Fracture toughness can also be further enhanced with the addition of ZrO2 particles. The transformation toughening due to ZrO2 phase along with crack bridging due to Ti phase can therefore provide the synergistic effect in order to improve the fracture toughness beyond the 4.7 MPa.m1/2 (fracture toughness of HA-10 wt. % Ti composite). Added to the linear thermal expansion coefficient of ZrO2 (9.4×10-6 oC-1)110 is found to be in between HA (17.3×10-6 oC-1) and Ti (8.4×10-6 oC-1). This helps in minimizing the difference in thermal expansion coefficient between HA and Ti and therefore, does not lead to the interfacial debonding in sintered 15 John Wiley & Sons, Inc.
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composites. It is worthwhile to mention that the most reliable measure of toughness properties of HA-Ti has been obtained only using long crack toughness measurement techniques, like SEVNB etc. Several aspects need to be discussed to explain the significant increase in the fracture toughness of the HA-Ti composites. The most important aspect is that the crack bridging due to lath shaped Ti phase as well as crack wake debonding at HA and Ti interface being responsible for enhancement of the fracture toughness (~ 4.7 MPa.m1/2) in the HA-10 wt. % Ti composite. However, no further improvement in fracture toughness of HA20Ti composite was noted due to non-uniform distribution as well as lower volume fraction of Ti, resulted due to cold welding of Ti during ball-milling. This can be confirmed using theoretical models proposed by Ashby 111
as well as Evans 112. The detailed analysis of fracture toughness of particulate (Ti) reinforced ceramic (HA)
has been reported elsewhere 33. According to these models, the particle size and the distribution of the metallic phase near the crack tip (which is directly proportional to the volume fraction of metallic phase in the matrix) has a strong influence on the fracture toughness (see Figure 2a-c). The details are reported elsewhere
111,112
.
Although the shape of metal particles has an effect on fracture toughness, the average number of particles near the crack tip is reported to the deciding factor in the ceramic-metal (e.g. HA-Ti) composites (see Figure 2b-c). The unification of Ti in HA-20Ti composite results in the reduction in volume fraction Ti in HA matrix and therefore reduces the distribution of Ti particles near the crack tip. Thus, fracture toughness of HA-20 wt.% Ti is found to be lower than HA-10 wt.% Ti composite. In future, cryo-milling of the HA-Ti composition may be found as a good alternative to conventional ball-milling in this particular case. The cryo-milling is expected to produce well distributed Ti particles in the starting powder mixture. This will allow good dispersion of Ti in the composites with the concentration of Ti more than 20 wt% causing further improvement in fracture toughness. It may also be possible to reduce the initial particle size considerably by cryo-milling. This will lead to the HATi composites with even higher fracture toughness. Therefore, designing the composites with different shape and size of Ti particles may be useful in improving the fracture toughness. In future, cryo-milling parameters need to be optimized and similar SPS experiments can be performed to prove this point.
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In the last two decades, the fractal modeling approach has also been used to a modest extent to predict the fracture toughness of the ceramic composites
113,114
. Xu and co-workers successfully used the Mandelbrot’s
modified fractal model to calculate the fracture toughness (KIC) of alumina with equiaxed grains as well as reinforced with elongated grains 115. According to this model, the microscopic (characterized by irregular crack propagation path) fracture toughness can be given as,
K IC = (2 Eγ s )1/2 ( 1r )
( d f )/2
..................(8)
where, E is the elastic modulus in tension; γs the interfacial energy equals to (Ebo/π2), where, E and bo are the elastic modulus and atomic dimension, respectively 101. In equation 8, df is the fractal dimension, which defined as, df =
log( N ) .................(9) 1 log r
where, r is the ratio that describe the segments generated in self-similar object. A self-similar object (fractal curve) can be broken into arbitrary small elements with each small element being a replica of the entire object; N is calculated from fractal curve and equal to the number of patterns generated for a ratio ‘r’ 116,117. Recently, Rishabh et al. 118 have also used the Mandelbrot’s modified model to estimate the fracture toughness of monolithic alumina with equiaxed grains. The fracture toughness of alumina reinforced with elongated alumina grains as well as with single wall carbon nanotubes was calculated. It has been found that the Mandelbrot’s modified model can be used to estimate the fracture toughness of ceramic composites instead of using Vickers indentation method. The fractal modeling approach in case of bioceramics is not being extensively utilized to predict the toughness properties and this is one area, which requires attention in future. Since the fractal methods are based on the localized interaction of crack with energy dissipating objects, therefore these methods underestimate the fracture toughness of ceramic materials, which are governed by Weibull statistics of brittle failure. Nevertheless, the fractal modeling approach can be useful to select the second phase reinforcement from the toughness enhancement perspective.
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6. Functionally graded materials: The functionally graded materials (FGMs) are also drawing the attention in biomedical engineering due to their exceptionally good mechanical and biological properties
51,119
. In last one decade, functionally FGMs are
regarded as a novel concept to perceive graded functionality that otherwise cannot be realized using conventional homogeneous materials. Therefore, in the context of integrating both biocompatibility and the mechanical properties, we now briefly discuss the relevance for the development of the functionally graded materials. As mentioned earlier that the bioinert nature of Ti does not allow the direct contact with bone tissue and, therefore, bonding is mediated by the fibrous tissue formation. In addition, the flaking off the HA coating from the metallic substrate (e.g. Ti) surface leads to the revision surgery, when HA-coated Ti is used in orthopedic applications. Therefore, graded distribution of HA and Ti in HA-Ti based FGM bulk composites can resolve the issue of debonding at the interface. In this way, FGM provides a better solution to the HA coated Ti and/or other metal implants. Despite such anticipated advantages, the development of FGM as a biomaterial is attempted only to a very modest extent with only few works reported so far. Chenglin et al. 119 reported the design of two different types of HA-Ti based FGMs, sintered by hot pressing. In the first design, Ti/HA content varies from the outer layer to the center layer (Figure 3a), whereas Ti concentration varies from one end to other end in the second design (Figure 3b). In the first design, the presence of HA as an outer layers provides the favorable surface for the osseointegration. However, the central layers with higher Ti improve the strength of the FGM. In the second design, one surface is bioactive due to the presence of 40 wt.% HA with the other surface being highly bioinert (monolithic Ti). In such design, HA-rich layer would enhance the osseointegration; while bioinert Ti surface promotes the fibrous tissue growth
120
. In
addition, the presence of layers with higher HA content (> 50 wt.%) will promote faster osseointegration and therefore, faster bone healing. As far as the mechanical properties are concerned, a gradual change in the fracture toughness has been reported due to the presence of various layers with different HA and Ti content, as shown in figure 3c.
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In particular, HA-Ti based FGM with the composition varying from Ti-0 vol.% HA to Ti-60 vol.% HA gives the closest approximation to mimic the mechanical properties of the bone. It is important to mention that the fracture toughness of the cortical bone varies over a range of 2-12 MPa.m1/2. Figure 3c shows a drastic increase in fracture toughness for Ti-20 vol. % HA and a high toughness of ~29 MPa.m1/2 was measured for pure Ti. The presence of multiple layers with different elastic modulus will allow the distribution of load to different layer and the most of the load will be transferred to Ti-rich layer. Although the presence of a higher HA content layer in FGM structure can provide the excellent osseointegration and low stress shielding effect, the presence of higher Ti content layer at the other end may result in poor osseointegration and higher stress shielding effect. Importantly, these issues have not been addressed and discussed in the existing literature 51,119,121,122. In order to overcome these issues, the present authors proposed the concentric type FGM with various cylindrical layers of HA-Ti compositions with the outermost layer being HA-rich. Such design with HA-rich outermost layer is expected to result in excellent osseointegration with the host bone. In contrast, the presence of inner (core) layer with higher Ti content than outer layers provides excellent fracture toughness and load bearing capability. Also, the outer layer with lower elastic modulus (due to lower in Ti content) is expected to mitigate the stress shielding near the host bone. The mechanical properties of an FGM depend on the processing as well as composition of different layers. Although the specific advantages for FGM design can be readily anticipated, the processing of such FGM structure is a challenging task. The difference in thermal expansion coefficient in each layer due to the compositional difference can result in the development of residual stresses in the interfacial region, leading to the interlayer cracking in the as-sintered FGM. Also, different layers in an FGM structure would have different shrinkage kinetics at any given sintering temperature. This will augment the residual stresses in an FGM structure. Both these effects however, can be minimized with the careful selection of the process window in terms of temperature – time – heating rate in SPS route. Some preliminary experiments in the authors’ research group revealed that the sintering of FGM structure with progressively higher Ti content with the outermost layer being monolithic HA using SPS route leads to cracking. However, SPS of HA-Ti based concentric FGM with 19 John Wiley & Sons, Inc.
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an outermost layer as HA-xTi (x
Laboratory for Biomaterials, Materials Research Centre, Indian Institute of Science, Bangalore-560012, India
Department of Materials Science and Engineering, Indian Institute of Technology Kanpur, Kanpur-208016, India
Abstract The research work on bulk hydroxyapatite (HA)-based composites are driven by the need to develop biomaterials with better mechanical properties without compromising its bioactivity and biocompatibility properties. Despite several years of research, the mechanical properties of the HA-based composites still need to be enhanced to match the properties of natural cortical bone. In this regard, the scope of the present review on the HA-based bulk biomaterials is limited to the processing and the mechanical as well as biocompatibility properties for bone tissue engineering applications of a model system i.e. hydroxyapatite-titanium (HA-Ti) bulk composites. It will be discussed in this review how HA-Ti based bulk composites can be processed to have better fracture toughness and strength without compromising biocompatibility. The advantages of the functionally gradient materials (FGMs) to integrate the mechanical and biocompatibility properties is a promising approach in hard tissue engineering and has been emphasized here in reference to the limited literature reports. On the biomaterials fabrication aspect, the recent results are discussed to demonstrate that advanced manufacturing techniques, like spark plasma sintering can be adopted as a processing route to restrict the sintering reactions, while enhancing the mechanical properties. Various toughening mechanisms related to careful tailoring of microstructure are discussed. The in vitro cytocompatibilty, cell fate processes as well as in vivo biocompatibility results are also reviewed and the use of flow cytometry to quantify in vitro cell fate processes is stressed upon. Keywords: Hydroxyapatite; Titanium; Composite; Fracture toughness; Biocompatibility corresponding author: Email: [email protected], Tel.: +91 80 2293 3256; Fax: +91 80 2360 7316
This article has been accepted for publication and undergone full peer review but has not been through the copyediting, typesetting, pagination and proofreading process which may lead to differences between this version and the Version of Record. Please cite this article as an ‘Accepted Article’, doi: 10.1002/jbm.a.35198
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Contents: 1. Introduction 2. Physical and mechanical properties required for bone materials 3. Orthopedic grafting using hydroxyapatite-based composites 4. Processing of HA-Ti composites 4.1. Conventional sintering 4.2. Advanced sintering 5. Toughening mechanisms and toughness properties 6. Functionally graded materials 7. Biocompatibility properties of HA-Ti composites 7.1. In vitro bioactivity 7.2. In vitro cytocompatibility 7.3. In vivo biocompatility 8. Closure References
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1. Introduction: Healthy bone and joints (e.g. articulating, knee, and hip joints) are necessary for the structural stability and pain free movements. However, these healthy bones and joints can get damaged in some unfortunate cases due to accident / diseases and fail to perform their normal functions. The damaged bones can be replaced or repaired with allogeneic bones or allografts
1-3
. However, the possible risk of transmission of life threatening infections
with allogeneic bones as well as inferior osteoconductivity and poor mechanical properties of the freeze dried bones limit the application of such types of biomaterials
4-6
. In contrast, the designed biomaterials with bone-
mimicking properties can be used as an alternative to the allografts and autogafts. Biomaterials development has a long history. In 1940, Jean and Robert Judet of Paris performed the replacement arthoplasty to replace the femoral head with acrylic resin, which opens up the possibility of joint replacement 7. However, after an initial success, this implant failed due to its poor mechanical properties as well as inapt design. More than 10 decades later, when most of the issues related to joint replacement have been settled down, the interest in new biomaterials remain as fierce as ever. A biomaterial is a synthetic biocompatible material/device that has been intently designed to persuade a definite activity in the biological system, which can simulate the desired biological function, without generating a short and/or long term damage to the nearby or distant tissue as well as without being getting damaged in this procedure 8,9. Depending on the in vivo performance, different biomaterials can be subcategorized as, bioactive, bioinert, and bio-tolerant. Bioactive materials (CaP-based materials) can promote new bone growth under physiological conditions. In addition, the bony tissue will form around the bioactive implant and adhere strongly due to osteogenesis
10,11
. Bioinert materials (e.g. Ti, Ti alloys, Al2O3 etc.) exhibit very minimal interaction with the
surrounding tissue. Generally, a protective fibrous capsule of the connective tissue may form around the implant by the contact osteogenesis
12,13
. In case of the bio-tolerant materials (e.g.: stainless steel, PMMA etc.), a
capsule of connective tissue will form by distant osteogenesis, and this capsule does not cling to the implant 1315
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On the other hand, depending on the stability of a synthetic implant in the host environment, the implant materials can be categorized into biodegradable, bioresistive as well as bioresorbable. A biodegradable material degrades (or dissolves) gradually in the host environment. In contrast, the bioresistive biomaterials do not degrade under the effect of host environment, e.g. HA, which has very low solubility in alkaline solution
11
.
Bioresorbable materials are the CaP-based biomaterials (α and β-tricalcium posphates), which provide the therapeutic agent (e.g. Ca2+, PO43-) during controlled dissolution, essential for faster healing of bone injury. For many biomedical applications, a biomaterial should show the bioactivity, osteoconductivity, and osteoinductivity. Bioactivity is the property of implant material to make direct bonding with living bone
16,17
.
The bonding is normally mediated by an intervening (apatite) layer formed between the implant and the bone 18. On the other hand, osteoconductivity is defined as the ability of an implant material that can serve as a scaffold for new bone formation and therefore, promotes the osteoblast adhesion, proliferation, and differentiation
19,20
.
Osteoinductivity is the ability of a biomaterial to induce the new bone formation, implanted at non bone forming site, in vivo 21. It is well documented in case of bone tissue engineering applications that the success of an implant material depends on the biocompatibility, osseointegration property, physical/mechanical properties and antibacterial property etc.
22-25
. Metals have an advantage of good mechanical properties, like high fracture toughness and
flexural strength. But the corrosion as well as wear of the implant (results in leaching of metal ions) and stress shielding results in irritation, inflammation, and loosening of implants metal implants result in weak bonding with host bone, in vivo
13,15
26,27
. Additionally, bioinert nature of
. Although ceramics based biomaterials are
found suitable for hard tissue applications due to its biocompatibility, corrosion resistance, and compressive strength properties
28,29
, inherent brittleness of ceramics may result in the early stage failure of the implants
during load bearing applications
30
. A classic example of a brittle bioceramic is hydroxyapatite, which has
widely been researched due to its structural similarity with major inorganic component of bone composition 30.
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Bone can be considered as a biocomposite of HA and collagen and thus synthetic bulk composites made by these constituents can mimic the bone composition more closely. It is worthwhile to mention here that there are many inexpensive methods available that allow the hydroxyapatite synthesis in bulk quantity
31,32
. However,
due to weak mechanical properties as well as processing related challenges at high temperature (high temperature sintering leading to the denaturing of collagen), HA-collagen based composites can not be used for the load bearing applications. It is important to mention here that, to ensure better mechanical properties, HA needs to be sintered at above than 950 °C
33
. Although materials, like Co-Cr alloys, Ti and its alloys, and
stainless steels are extensively used in bone-tissue engineering applications, the expedition to search an ideal material for such application still needs extensive efforts 13. Despite good toughness, metals, like Co-Cr alloys, Ti and its alloys, and stainless steels do not considered reliable in complex physiological environment for longterm usages due to leaching of toxic ion (e.g. Cr, Ni) under physiological conditions as well as bioinert/biotolerant nature
12-15,34-36
. The application of these metallic implant materials can be extended by coating the
bioinert/bio-tolerant surfaces with bioactive materials, like calcium phosphate
37-39
. However, the poor
interfacial strength between base material and coating as well as fracture of coating durings the service period may result in serious clinical issues
40,41
. Another issue with metallic implants is the generation of metal ions
due to dissolution and/ or wear during the long service period main cause of local necrosis and/or inflammation
43
42
. The generation of nano-sized debris is the
. These wear debris may also be carried out by blood to
other organs, like heart, lever, and lung and this leads to the severe toxicity of these organs. In the context of load bearing implants, high elastic modulus of the implant with respect to the neighboring bone causes stress shielding, being responsible for the bone resorption near to the implant
44-48
. Therefore, the
elastic modulus of an implant ideally needs to be matched with bone, while efforts are to be made to enhance both fracture toughness and strength. In the above backdrops, the hydroxyapatite (HA) - based metal or ceramic reinforced composites can be used as a potential bulk implant material, where the second phase contributes specifically towards better mechanical properties. The choice of the second phase is limited as it needs to have good mechanical property as well as 5 John Wiley & Sons, Inc.
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biocompatibility. As compared to Al2O3, ZrO2, and stainless steels, titanium (Ti) is found to be more promising second phase material for orthopedic applications due to its excellent corrosion resistance property under physiological environment, low density (4.540 g/cc) as well as the lower elastic modulus (~110 GPa) than CoCr alloy (~230 GPa)/ stainless steel (~205 GPa) 22,48. The design of the HA-Ti composites underlies the fact that the Ti addition significantly improves the mechanical properties of the composites without degrading the cytocompatibility properties 49-51. In the present review, the processing of HA-Ti composites using conventional and advanced sintering route has been discussed. The improvement in mechanical properties due to Ti addition has been elaborated. The results of in vitro and in vivo experiments have also been provided to justify that the biocompatibility property is not compromised due to Ti addition to HA.
2. Physical and mechanical properties required for bone materials: Apart from the high dissolution rate under physiological environment, inferior corrosion resistance and poor tribological properties as well as inferior bactericidal property, an implant may fail during the service period due to, (a) mismatch in elastic modulus between the implant and the surrounding bone and (b) poor fracture toughness 22. Therefore, we will focus on these aspects in the following sections. A lower load experienced by the host bone can result in stress shielding due to comparatively higher elastic modulus of the implant than bone
52
. This leads to the resorption of bone near the implant due to remodeling
followed by loosening of the implant
44-47
. The problem of stress shielding in a particular application can be
minimized using natural grafts, or choosing an appropriately designed material. For example, three dimensional scaffolds of biocompatible materials having gradient porosity with controlled shape, size and distribution is a promising solution in such application 53,54. As compared to Al2O3 (390 GPa) and ZrO2 (205 GPa), elastic modulus of HA (80-110 GPa) was found to be closer to the cortical bone (14-20 GPa)
22,55-61
. However, the fracture toughness of HA (0.7-1.2 MPa.m1/2) is
lower than the even lower limit of the cortical bone (2-12 MPa.m1/2)
55-58
. As mentioned earlier, the poor
fracture toughness restricts the use of HA in the load bearing applications. In contrast, Ti with exceptionally 6 John Wiley & Sons, Inc.
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high fracture toughness (60 MPa.m1/2) is found to have lower elastic modulus (117 GPa) than Al2O3 and ZrO2. Therefore, the HA-Ti system can be considered as a model system. Apart from biocompatibility properties, the presence of the second phase can affect other properties, like fracture toughness, strength, and elastic modulus. Since, we are more concerned about the fracture toughness of bioceramics materials, a detailed description on the effect of second phase on fracture toughness of HA has been provided in section 5.
3. Orthopedic grafting using hydroxyapatite-based composites: The natural bone is a composite of hydroxyapatite and collagen. Therefore, a biomaterial with HA and collagen (e.g. HA-collagen composites) can provide the suitable surface for higher cell adhesion, proliferation and osseointegration 62. However, due to poor mechanical properties (e.g. fracture toughness), the use of HA-based composites reinforced with polymer has been restricted only for low load-bearing applications
63
. As will be
discussed later, this has been the driving force for the research on HA-based biocomposites. From the perspective of material development, the processing window must always be chosen on the basis of base material. For load bearing application, the base material (e.g. HA) has to be densified up to maximum extent, which is only possible if the material can be sintered at temperature ≥ 950 oC
33,64
. Nevertheless, at this
temperature, most of the polymers are thermally unstable and high temperature sintering leads to the complete removal/denaturing of most of the polymers, like collagen
65,66
. In contrast, innovative processing routes, like
freeze casting may allow the fabrication of HA-polymer composites with porous structure 67,68. To develope a bone replacement material, the choice of the second phase depends on the application as well as anticipated functionality. For example, BaTiO3, CaTiO3, Fe3O4, and ZnO or Ag can be added to HA as a second phase to introduce the piezoelectric, electrical conductivity, magnetic, and bactericidal properties, respectively 69-73
. However, brittle nature and the load bearing incapability of the ceramics may limit the broader application
of such types of ceramic-ceramic composites; some examples include HA-20HAwhisker, HA-20ZrO2, and HA(Al2O3 coated) ZrO2 with fracture toughness of 1.4 MPa.m1/2, 1.5 MPa.m1/2, and 3 MPa.m1/2, respectively 74-76. Interestingly, a glaring example of HA-based ceramic composite with fracture toughness, higher or equal to that 7 John Wiley & Sons, Inc.
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of the cortical bone is still unavailable (see figure 1). In contrast, the incorporation of metallic reinforcement as second phase is deemed to improve the fracture toughness of the HA-based composites. However, the sintering reactions between HA and metal phase are of major concern and can affect the biocompatibility properties in vitro and in vivo. Although stainless steel and Co-Cr alloys are competitively better than ceramics counterparts in terms of fracture toughness, the high elastic modulus of metals however may lead to stress shielding during in vivo performance 25. In contrast, Ti was found to be most suitable for biomedical applications due to excellent biocompatibility, corrosion resistance property and stability in physiological conditions
77
. Thus, in order to
incorporate both load bearing capabilities without fracture under peak loading as well as osseointegration property, the HA-Ti composites can be used as a bulk biomaterials. Over the various composites (HA-BaTiO3, HA-CaTiO3, HA-Al2O3, HA-ZrO2, HA-ZnO, HA-Ag), the selection of HA-Ti composites as a model system in the present review can be rationalized on the basis of following aspects: (a) low density of HA (3.16 g/cc) 63 and Ti (4.54 g/cc); (b) high fracture toughness of metallic Ti (60 MPa.m1/2)
22
; (c) bioactivity of HA and bio-
inertness of Ti; (d) corrosion resistant properties of Ti; (e) limited dissolution of HA under physiological condition 78.
4.
Processing of HA-Ti composites:
HA-Ti composites can be made either by conventional pressureless sintering or advanced sintering, such as spark plasma sintering (SPS). SPS has the advantage of higher heating rate and shorter soaking time over the conventional sintering. The advanced sintering techniques have many other advantages as compared to the conventional sintering and have been described in section 4.2. Traditionally, the HA-Ti composites were prepared by the homogeneous mixing of the HA and Ti powders using a ball-mill, followed by green compaction followed by pressureless sintering at relatively higher temperatures (≥ 1100 °C) and soaking time (≥ 1 hour) 64. Recently, the FAST (field assisted sintering technique) has been utilized to prepare highly densified samples from the powder mixture within a short time 33,79,80.
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4.1.
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Conventional sintering:
In this subsection, the literature reports on conventional sintering of HA-Ti are briefly discussed. Fahami et al. 81
used CaHPO4, CaO, and Ti powders as precursors to prepare the HA-Ti composites by the mechanochemical
synthesis of a mixture of HA and Ti according to the following reaction:
6CaHPO4 + 4CaO + Ti → Ca10 ( PO4 )6 (OH ) 2 + Ti + 2 H 2O
-------(1)
However, mechanochemical synthesis leads to formation of HA with reduced crystallinity. It is important to note that the lower crystallinity of the HA powder results in the increased bioresorption and solubility under physiological conditions
82
. Therefore, further attempt has been made by Fahami et al.
81
to increase the
crystallinity from 13 to 69 % by annealing the ball-milled powder at 650 oC for 2 hours, synthesized after reaction (1). However, this results in the oxidation of Ti and decomposition of HA to β-TCP (beta- tri-calcium phosphate) phase. In a different approach, Chu et al.
83
prepared the HA-20 wt.% Ti composite by mixing HA
and Ti powders in a ball mill, followed by hot pressing at 900 oC and 1000 oC. Importantly, the sintering at 1000 oC leads to the formation of α-TCP and Ca4O(PO4)2 phases. The difference in the linear thermal expansion coefficient (α) of HA (17.3×10-6 oC-1)
84
and Ti (8.4×10-6 oC-1)
85
phases results in inferior sinterability in case
of HA-20 wt.% Ti composite with ~67% relative density, achieved when sintered at 1200 oC 83. However, the same sintering conditions result in ~88% relative density in monolithic HA. Importantly, the addition of Ti to HA leads to an increase in fracture toughness from 0.7 MPa.m1/2 to 1.0 MPa.m1/2. Although Ti addition to HA was intended to increase the fracture toughness significantly, the marginal improvement in the fracture toughness can be attributed to the poor sinter density of the composite. In an earlier work, Nath et al. 64 prepared the HA-Ti composites with varying amount of Ti (10-40 wt.%) using pressureless sintering. It has been reported that the sintering of different compositions at 1000-1400 oC for 2 hours leads to the extensive sintering reactions and therefore the formation of TiO2, CaO, TCP and CaTiO3 (Table 1). The thermodynamic calculations show the higher probability of oxidation of Ti during conventional sintering due to higher sintering temperature and
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longer holding time 64. Such calculations also indicate the possibility of a reaction between HA and TiO2 above 750oC. Considering the thermodynamical aspect, Yang et al.
86
prepared the HA-10 wt.% Ti and HA-20 wt.% Ti
composities by sintering at 1100 oC in vacuum to minimize the oxidation of Ti. The FT-IR and XRD analysis of the HA-Ti samples show the decomposition of HA to α-TCP, TTCP (tetracalcium phosphate) as well as Ca2Ti2O5. As the sintering temperature further increases, HA is reported to decompose to HA with deficiency of hydroxyl group in the following way 87,88: o
vacuum , T >800 C Ca10 ( PO4 )6 (OH ) 2 → Ca10 ( PO4 )6 (OH ) 2− 2 x Ox + xH 2O ↑ -------(2)
The water vapor produced from the reaction (2) reacts with Ti to form TiO2 [see eq. (3)], which can finally react with HA to form β-TCP and CaTi2O5 [see eq. (4)].
Ti + 2 H 2O → TiO2 + 2 H 2 -------(3)
Ca10 ( PO4 )6 (OH ) 2 + 2TiO2 → 3Ca3 ( PO4 )2 + CaTi2O5 + H 2O ----(4) Thus, it is not impossible, but extremely hard to retain Ti in the sintered product using conventional sintering routes. Alternatively, the HA can form solid solution with metal oxides, carbonates and halides
89
. The Ca2+
ions can also be substituted with Ti4+, Na+, K+, Sr2+, Ba2+, Pb2+ and Y3+ 90-93, while OH- ion can be replaced with CO32-, F- and Cl- ions 94-96. Even PO43- ions can be replaced with AsO43-, and VO43- 97. In order to investigate the influence of the Ti ions on the stability of the structure of HA, Ribeiro et al.
89
incubated HA powder (pre-
heated at 1000 oC) in a solution containing Ti4+ ions. During incubation, Ti4+ ions get adsorbed on the HA surface and as a result, the Ca2+ ions get released into the solution. In this case, Ca2+ ions are replaced with Ti4+ ions according to the following reaction:
Ca10 (PO4 )6 (OH)2 (solid) + (n / 2)Ti4+ → (Ca10−nTi(n / 2))(PO4 )6 (OH)2 + nCa2+ + (n / 2) where, symbol
...
.....(5)
is a vacancy
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The above reactions result in a reduction in the lattice parameter of Ca-deficient HA due to smaller ionic radius of Ti4+ (0.68 Å) than Ca2+ (0.99 Å). The sintering under a protective atmosphere (e.g. argon) is another promising approach towards retaining the metallic Ti in the sintered product. In order to see the effect of protective environment on the sintering reaction in the HA-Ti composites, Ning et al. 98 conventionally sintered the HA-50Ti composite at 1200 oC for 30 min in argon atmosphere. However, the sintering leads to the formation of Ti2O, CaTiO3, CaO and TiP-like phases along with α-Ti. In summary, conventional sintering of HA-Ti composites even at very low temperature (1000 °C) leads to the extensive dehydroxylation of HA as well as the reaction between HA and Ti. The formation of TCP phases will result in higher dissolution of the implant materials and therefore, decreased mechanical strength during the service period, in vivo. As mentioned earlier, it is extremely difficult to retain the metallic Ti in the composite during the conventional sintering. The sintering in protective atmosphere of Ar was also not found to be helpful in retarding the oxidation reactions in conventional sintering due to the longer sintering duration (≥ 1 hour) and higher holding temperature (≥ 1000 °C). In the following, we shall discuss the applicability of advanced sintering techniques (e.g. SPS) to prepare the composites with predominant retention of metallic Ti.
4.2.
Advanced sintering
Among various advanced sintering techniques, field activated sintering (FAST) is widely investigated to resolve the issues related to the conventional sintering, discussed earlier. For example, Nakahira et al.
79
reported the pulse electric current sintering (PECS) of HA-20 vol.%Ti and HA-25 vol.%Ti composites at 800 C and 1000 oC under 30 MPa in a protective argon atmosphere. Although PECS suppresses the oxidation of Ti at 900 oC, the formation of Ti2O was found in samples sintered at 1000 oC. Mondal et al.
99
used the spark
plasma sintering technique to sinter β-TCP and Ti powders at 1200 oC temperature under 50 MPa pressure. Despite the retention of β-TCP and Ti, this results in the formation of CaTiO3. 11 John Wiley & Sons, Inc.
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In a recent work, the present authors have reported the results of extensive study on spark plasma sintering and phase evolution of HA-xTi (x = 5, 10, 20 wt%) composites
33
. The processing conditions were intelligently
tailored to achieve high toughness with the minimum sintering reaction between the HA and Ti. The XRD analysis confirms the predominant presence of HA and Ti. However, the noticeable diffraction peaks corresponding to β-TCP, CaO, and CaTi4(PO4)6 were also recorded. FT-IR results show the reduction in the absorption peak intensity corresponding to OH- stretching (3574 cm-1) as well as OH- liberation (636 cm-1), confirming the dehydroxylation of HA at a minute level. The EPMA study indicates the presence of a thin diffusion layer around the Ti particles with a composition different than that of Ti and HA phase. XRD, EPMA and FT-IR findings were supported by the detailed TEM investigations. The analysis of the selected area electron diffraction (SAED) patterns from different phases confirms the formation of CaTi4(PO4)6 at the interface of HA and Ti phases, due to the interfacial reaction between β-TCP and Ti
33
. The β-TCP has
been formed due to local dehydroxylation of HA. In summary, spark plasma sintering at relatively lower processing temperature with a shorter holding time than conventional pressureless sintering can enable maximum retention of HA and Ti with minimal sintering reactions.
5. Toughening mechanisms and toughness properties Fracture in brittle solids is related to the propagation of cracks under the tensile loading. Importantly, the crack propagates if and only if the stress at crack front is higher than a critical value. In brittle solids, the resistance against the crack propagation can be enhanced by introducing the mechanisms that reduce the driving force for the crack propagation, such as dissipating the additional energy associated with crack tip leading to the blunting of the crack tip 100. In polycrystalline ceramic materials, the microstructure dependent fracture toughness can be correlated with the dissipation of the crack tip energy and thereby with the reduction in the driving force for the 12 John Wiley & Sons, Inc.
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crack propagation
101
. It is generally recognized that weak interfaces, such as grain boundaries and/or matrix-
reinforcement boundary are the preferred crack deflection paths. Moreover, the presence of the second phase disturbs the crack propagation by forcing the crack either to go around it or through it. The extent to which the it is able to dissipate the energy depends on the morphology / distribution and mechanical properties of the second phase in ceramic matrix composite 102. Although the presence of second phase may improve the fracture toughness of composites, the chemical reaction at the interface during processing (e.g. sintering) always offers the challenge for the materials scientists to optimize the toughness. On the basis of microstructural characteristics of the second phase; i.e., nature, shape and size, the toughened composite materials can be categorized as 103: (a) particulate reinforced composite, (b) whisker reinforced composite, and (c) fiber reinforced composite. In the context of the current review paper, the toughening in metallic particulate reinforced ceramic composites is attributed to the shielding effect to the crack front propagation due to ductile metal reinforcements. In such case, the crack resistance energy of the matrix can be given by 101,
Rο = (1 − V f )2γ B .........(6) where, γ B is the intrinsic cohesive energy of the matrix material and V f is the volume fraction of the second (reinforcement) phase. Thus, both the matrix as well as the reinforcement material contributes towards the fracture toughness and hence, enhanced fracture toughness can be achieved by intelligently designing the composites having mixed shielding mechanisms. In case of the polycrystalline monolithic ceramic materials, the frontal zone shielding results due to the interaction of the crack front with dislocation cloud and/or microcrack cloud. The activation of the dislocations results in generation of dislocation cloud and this restricts the crack front propagation by applying a stress field on the crack tip
101
. The stress field generated around the dislocation is proportional to the shear modulus (G)
and can be given as 104, 13 John Wiley & Sons, Inc.
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σ∝
Page 14 of 40
Gb ..........(7) r
where, G, b, and r are the shear modulus of the material, Burger vector, and distance from the dislocation, respectively. On the other hand, the microcrack cloud can be activated in the stress field of primary cracks at the relatively weaker zone, like grain boundaries and interphase flaws. Unlike the combination of dislocation as well as microcrack clouds towards the improvement in fracture toughness in brittle ceramics, the phase transformation in zirconia (ZrO2) can also contribute to the fracture toughness improvement in ZrO2 reinforced ceramic matrix
101
. In such case, the partially stabilized ZrO2 in
metastable tetragonal state transforms to stable monoclinic phase at the crack tip stress field transformation can be controlled by using additives, like MgO, CaO, Y2O3 and CeO2
100
100
. This
. In addition, the grain
bridging of crack in monolithic polycrystalline materials results in the toughness improvement due to the dissipation of crack energy as a result of grain sliding and/or fracture. For example, Swanson et al. 105 reported the evolution of grain bridging in pure polycrystalline alumina. In such case, the transgranular fracture may occur in large grains along with primary crack due to high frictional stresses at sliding grain boundary surfaces. In 1983, Faber and Evans studied the effect of geometrical aspect of the second phase on the crack deflection and therefore, effect on the fracture toughness 106,107. The strength of the interfacial boundary plays an important role in deciding whether a crack will pass though second phase or get deflected. In a case when the interface is strong enough to deflect the crack, it penetrates through the second phase. This may result in pinning of the advancing crack front at the interaction site. In case of ductile metal phase, this results in the dissipation of crack tip energy due to crack bridging. In the context of biomaterials, the benefits of the second phase addition to a ceramic matrix are twofold: (1) second phase improves the fracture toughness dramatically due to frontal wake and/or bridge interface shielding 33
and (2) second phase leads to the modification in surface energy
108
, which can be used as a tool to
manipulate the cell-material interaction. 14 John Wiley & Sons, Inc.
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Although the reinforcement of ceramic matrix with fiber or whisker results in improvement in fracture toughness, the release of nanosized and needle shaped whisker/fibrous can lead to the severe toxicity to the cells caused by lipid layer damage and genotoxicity
43,109
. Therefore, the fiber/whisker toughened ceramic
composites are not considered for clinical applications, although related toughening mechanisms can be very well adapted to develop novel HA-based high toughness composites. In parallel, the reinforcement of ceramic matrix with ductile metal particles results in dramatic improvement in mechanical properties. The crack bridging at metal particles results in the dissipation of energy associated with crack tip
33,101
. The fracture
energies of the pulsed electric current sintered (PECSed) HA-Ti composites, measured by Nakahira et al.
79
using single edge double notched beam method, show a substantial increment with the Ti content. The fracture energies of 9 J/m2 and 20 J/m2 were reported for HA-20 vol.% Ti and HA-25 vol.% Ti samples, respectively. The reason for the improvement in fracture toughness was crack deflection and crack bridging by ductile Ti particles, which are retained in the HA matrix. Similarly, substantial improvement in the fracture toughness of the spark plasma sintered HA-Ti composites were reported by the present authors (see Figure 1) 33. An addition of 10 wt.% Ti to HA increases the fracture toughness (KIC) from ~3.4 MPa.m1/2 to 4.7 MPa.m1/2. Further addition of Ti (20 wt.%) to HA does not lead to any further improvement of the fracture toughness (4.3 MPa.m1/2). The increased fracture toughness of HA and HA-Ti composites can be explained on the basis of crack deflection and crack bridging. In the context of toughness enhancement, the synergistic interaction of multiple toughening mechanisms can be explored. The presence of Ti particles provides the significant contribution toward the fracture toughness by crack bridging. Fracture toughness can also be further enhanced with the addition of ZrO2 particles. The transformation toughening due to ZrO2 phase along with crack bridging due to Ti phase can therefore provide the synergistic effect in order to improve the fracture toughness beyond the 4.7 MPa.m1/2 (fracture toughness of HA-10 wt. % Ti composite). Added to the linear thermal expansion coefficient of ZrO2 (9.4×10-6 oC-1)110 is found to be in between HA (17.3×10-6 oC-1) and Ti (8.4×10-6 oC-1). This helps in minimizing the difference in thermal expansion coefficient between HA and Ti and therefore, does not lead to the interfacial debonding in sintered 15 John Wiley & Sons, Inc.
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composites. It is worthwhile to mention that the most reliable measure of toughness properties of HA-Ti has been obtained only using long crack toughness measurement techniques, like SEVNB etc. Several aspects need to be discussed to explain the significant increase in the fracture toughness of the HA-Ti composites. The most important aspect is that the crack bridging due to lath shaped Ti phase as well as crack wake debonding at HA and Ti interface being responsible for enhancement of the fracture toughness (~ 4.7 MPa.m1/2) in the HA-10 wt. % Ti composite. However, no further improvement in fracture toughness of HA20Ti composite was noted due to non-uniform distribution as well as lower volume fraction of Ti, resulted due to cold welding of Ti during ball-milling. This can be confirmed using theoretical models proposed by Ashby 111
as well as Evans 112. The detailed analysis of fracture toughness of particulate (Ti) reinforced ceramic (HA)
has been reported elsewhere 33. According to these models, the particle size and the distribution of the metallic phase near the crack tip (which is directly proportional to the volume fraction of metallic phase in the matrix) has a strong influence on the fracture toughness (see Figure 2a-c). The details are reported elsewhere
111,112
.
Although the shape of metal particles has an effect on fracture toughness, the average number of particles near the crack tip is reported to the deciding factor in the ceramic-metal (e.g. HA-Ti) composites (see Figure 2b-c). The unification of Ti in HA-20Ti composite results in the reduction in volume fraction Ti in HA matrix and therefore reduces the distribution of Ti particles near the crack tip. Thus, fracture toughness of HA-20 wt.% Ti is found to be lower than HA-10 wt.% Ti composite. In future, cryo-milling of the HA-Ti composition may be found as a good alternative to conventional ball-milling in this particular case. The cryo-milling is expected to produce well distributed Ti particles in the starting powder mixture. This will allow good dispersion of Ti in the composites with the concentration of Ti more than 20 wt% causing further improvement in fracture toughness. It may also be possible to reduce the initial particle size considerably by cryo-milling. This will lead to the HATi composites with even higher fracture toughness. Therefore, designing the composites with different shape and size of Ti particles may be useful in improving the fracture toughness. In future, cryo-milling parameters need to be optimized and similar SPS experiments can be performed to prove this point.
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In the last two decades, the fractal modeling approach has also been used to a modest extent to predict the fracture toughness of the ceramic composites
113,114
. Xu and co-workers successfully used the Mandelbrot’s
modified fractal model to calculate the fracture toughness (KIC) of alumina with equiaxed grains as well as reinforced with elongated grains 115. According to this model, the microscopic (characterized by irregular crack propagation path) fracture toughness can be given as,
K IC = (2 Eγ s )1/2 ( 1r )
( d f )/2
..................(8)
where, E is the elastic modulus in tension; γs the interfacial energy equals to (Ebo/π2), where, E and bo are the elastic modulus and atomic dimension, respectively 101. In equation 8, df is the fractal dimension, which defined as, df =
log( N ) .................(9) 1 log r
where, r is the ratio that describe the segments generated in self-similar object. A self-similar object (fractal curve) can be broken into arbitrary small elements with each small element being a replica of the entire object; N is calculated from fractal curve and equal to the number of patterns generated for a ratio ‘r’ 116,117. Recently, Rishabh et al. 118 have also used the Mandelbrot’s modified model to estimate the fracture toughness of monolithic alumina with equiaxed grains. The fracture toughness of alumina reinforced with elongated alumina grains as well as with single wall carbon nanotubes was calculated. It has been found that the Mandelbrot’s modified model can be used to estimate the fracture toughness of ceramic composites instead of using Vickers indentation method. The fractal modeling approach in case of bioceramics is not being extensively utilized to predict the toughness properties and this is one area, which requires attention in future. Since the fractal methods are based on the localized interaction of crack with energy dissipating objects, therefore these methods underestimate the fracture toughness of ceramic materials, which are governed by Weibull statistics of brittle failure. Nevertheless, the fractal modeling approach can be useful to select the second phase reinforcement from the toughness enhancement perspective.
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6. Functionally graded materials: The functionally graded materials (FGMs) are also drawing the attention in biomedical engineering due to their exceptionally good mechanical and biological properties
51,119
. In last one decade, functionally FGMs are
regarded as a novel concept to perceive graded functionality that otherwise cannot be realized using conventional homogeneous materials. Therefore, in the context of integrating both biocompatibility and the mechanical properties, we now briefly discuss the relevance for the development of the functionally graded materials. As mentioned earlier that the bioinert nature of Ti does not allow the direct contact with bone tissue and, therefore, bonding is mediated by the fibrous tissue formation. In addition, the flaking off the HA coating from the metallic substrate (e.g. Ti) surface leads to the revision surgery, when HA-coated Ti is used in orthopedic applications. Therefore, graded distribution of HA and Ti in HA-Ti based FGM bulk composites can resolve the issue of debonding at the interface. In this way, FGM provides a better solution to the HA coated Ti and/or other metal implants. Despite such anticipated advantages, the development of FGM as a biomaterial is attempted only to a very modest extent with only few works reported so far. Chenglin et al. 119 reported the design of two different types of HA-Ti based FGMs, sintered by hot pressing. In the first design, Ti/HA content varies from the outer layer to the center layer (Figure 3a), whereas Ti concentration varies from one end to other end in the second design (Figure 3b). In the first design, the presence of HA as an outer layers provides the favorable surface for the osseointegration. However, the central layers with higher Ti improve the strength of the FGM. In the second design, one surface is bioactive due to the presence of 40 wt.% HA with the other surface being highly bioinert (monolithic Ti). In such design, HA-rich layer would enhance the osseointegration; while bioinert Ti surface promotes the fibrous tissue growth
120
. In
addition, the presence of layers with higher HA content (> 50 wt.%) will promote faster osseointegration and therefore, faster bone healing. As far as the mechanical properties are concerned, a gradual change in the fracture toughness has been reported due to the presence of various layers with different HA and Ti content, as shown in figure 3c.
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In particular, HA-Ti based FGM with the composition varying from Ti-0 vol.% HA to Ti-60 vol.% HA gives the closest approximation to mimic the mechanical properties of the bone. It is important to mention that the fracture toughness of the cortical bone varies over a range of 2-12 MPa.m1/2. Figure 3c shows a drastic increase in fracture toughness for Ti-20 vol. % HA and a high toughness of ~29 MPa.m1/2 was measured for pure Ti. The presence of multiple layers with different elastic modulus will allow the distribution of load to different layer and the most of the load will be transferred to Ti-rich layer. Although the presence of a higher HA content layer in FGM structure can provide the excellent osseointegration and low stress shielding effect, the presence of higher Ti content layer at the other end may result in poor osseointegration and higher stress shielding effect. Importantly, these issues have not been addressed and discussed in the existing literature 51,119,121,122. In order to overcome these issues, the present authors proposed the concentric type FGM with various cylindrical layers of HA-Ti compositions with the outermost layer being HA-rich. Such design with HA-rich outermost layer is expected to result in excellent osseointegration with the host bone. In contrast, the presence of inner (core) layer with higher Ti content than outer layers provides excellent fracture toughness and load bearing capability. Also, the outer layer with lower elastic modulus (due to lower in Ti content) is expected to mitigate the stress shielding near the host bone. The mechanical properties of an FGM depend on the processing as well as composition of different layers. Although the specific advantages for FGM design can be readily anticipated, the processing of such FGM structure is a challenging task. The difference in thermal expansion coefficient in each layer due to the compositional difference can result in the development of residual stresses in the interfacial region, leading to the interlayer cracking in the as-sintered FGM. Also, different layers in an FGM structure would have different shrinkage kinetics at any given sintering temperature. This will augment the residual stresses in an FGM structure. Both these effects however, can be minimized with the careful selection of the process window in terms of temperature – time – heating rate in SPS route. Some preliminary experiments in the authors’ research group revealed that the sintering of FGM structure with progressively higher Ti content with the outermost layer being monolithic HA using SPS route leads to cracking. However, SPS of HA-Ti based concentric FGM with 19 John Wiley & Sons, Inc.
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an outermost layer as HA-xTi (x
