2018 Volume 59 Issue 1 Pages 1-13
The progress of metallic structural biomaterials, mainly titanium alloys, for implants with mainly focusing on mechanical biocompatibility is described. Mechanical biocompatibility includes not only Young's modulus but also broad sense of mechanical biocompatibilities such as balance of strength and elongation, fatigue endurance (fatigue strength) and fracture toughness. Specially, the present paper focuses on developments of high fatigue strength of (α + β)-type titanium alloys composed of non-toxic elements, low Young's modulus β-type titanium alloys composed of non-toxic and allergy-free elements, Young's modulus self-adjustable β-type titanium alloys composed of non-toxic elements, Ni-free β-type titanium alloys for biomedical applications. Ni-free stainless steels and Co-Cr-Mo alloys, cell viability of pure metals, and some very recent research and development topics are also briefly introduced in the present paper.
This Paper was Originally Published in Japanese in Materia Japan 56 (2017) 205–210. In order to introduce more recent topics on metallic biomaterials, sections 4, 5, 9, and 10 were newly added. According to adding new sections, Tables 2 and 3, and Figs. 1–6, 8, 9, 13, and 19–25 were newly added. The Refs. 1), 4–12), 14), 28–39), 51), and 57–72) were also newly added. Following these changing, the numbers of tables, figures, and references were newly changed.
Structural materials that show good mechanical biocompatibility (biomaterials) are required for use in implants that substitute for failed hard tissue, such as failed bone. These materials also need to show good stability under a high load. The implants fabricated from these biomaterials must exhibit bone functionality. Mechanical reliability is a very important attribute of structural materials used as bone substitute. From the viewpoint of mechanical reliability, metallic materials are superior to other materials because they exhibit an excellent balance of strength and ductility. They occupy almost of the structural materials for implants. In addition to mechanical reliability, structural materials must fulfill the safety requirements for use in a living body. Therefore, biomaterials must be composed of non-toxic and allergy-free elements, and no dissolution of the constituent elements must occur, that is, high corrosion resistance. In practice, stainless steels, cobalt (Co)-chromium (Cr)-molybdenum (Mo) alloys, and titanium (Ti) and its alloys have been used as representative implant materials. In the early stage of applications, mechanical reliability of these metallic biomaterials was regarded as the most important factor; hence, general structural metallic materials showing high corrosion resistance were applied in implants. Many of them are still used as structural biomaterials for implants. Subsequently, with the increased concern regarding the safety requirements, new alloys were designed for use as biomaterials to circumvent problems related to toxicity and allergic reactions. Furthermore, much attention has been paid to mechanical biocompatibility, with focus on Young's modulus, for the design of metallic biomaterials1).
During above mentioned research and development on metallic biomaterials with focusing on Ti and its alloys, research and development of Ni-free stainless steels2) and Co-Cr-Mo alloys3) for biomedical applications, biodegradable metallic biomaterials including magnesium (Mg) alloys4), cell viability of pure metals5,6), and biofunctional surface modifications7), etc. were progressed.
Very recently, the design of mechanically biocompatible metallic biomaterials, in which the user-friendliness during implantation of the implants is taken into consideration, has been proposed.
New trends of research and development of metallic biomaterials including zirconium (Zr) based alloys for biomedical applications8,9), additive manufacturing (AM) by selected laser melting (SLM)10) and electron beam melting (EBM)11) methods, and groove designing on the surface of the implant to enhance the bone growth with the preferential direction12) are also raised in very recent years.
The progress in designing and developments of metallic biomaterials are described in this paper, especially from the view point of mechanical biocompatibility. Young's modulus is commonly used as the main indicator of the mechanical biocompatibility of biomaterials; however, the suitability of mechanical characteristics (fatigue strength, balance of strength and ductility, and fracture toughness) to living tissue is also included into the broad sense of mechanical biocompatibility13) in this paper.
Some of very recent new topics including research and development of Zr based alloys for biomedical applications, AM for fabricating medical implants, and groove designing on the surface of the implant to enhance the bone growth with the preferential direction are also briefly described.
Ti and its alloys are superior in terms of biological biocompatibility, which is judged based on the chemical interaction of the material with living tissue and is evaluated by cyto-toxicity or animal tests14–16). Ti and its alloys are gaining widespread importance due to their high biocompatibility.
Young's moduli of SUS 316L stainless steel, which is the main stainless steel used for implants and Co-Cr-Mo alloys are approximately 180 and 210 GPa17), respectively. These values are much greater than that of cortical bone (hereafter, referred to as bone), approximately 10–30 GPa18). In contrast, Ti and Ti-6Al-4V ELI (Extra Low Interstitial), which find practical applications in implants, show Young's moduli of approximately 105–110 GPa18), which are much less than those of SUS 316L stainless steel and Co-Cr-Mo alloys, but relatively comparable to that of bone. The use of Ti and its alloys in implants can inhibit non-homogeneity of stress between the bone and the implant (referred to as stress shielding) and is advantageous for bone remodeling according to the stress (referred to as bone remodeling). Therefore, the mechanical biocompatibility of Ti and its alloys, considering Young's modulus, is superior to those of SUS 316L stainless steel and Co-Cr-Mo alloys. This concept is connected to the development of titanium alloys composed of non-toxic and allergy-free elements and showing a low Young's modulus close to that of bone.
Early metallic biomaterials contained toxic elements were derived from general structural materials as stated above. Therefore, attempts were made to designing metallic biomaterials composed of non-toxic elements. Firstly, the toxicity of the β-stabilizing element vanadium (V), which is a constituent element of Ti-6Al-4V ELI, was revealed19). Therefore, V in Ti-6Al-4V ELI was replaced by another β-stabilizing element, iron (Fe) or niobium (Nb), both of which are considered safer for living tissue than V. Ti-5Al-2.5Fe20,21) and Ti-6Al-7Nb22), which are also (α + β)-type titanium alloys, were introduced17). Therefore, it can be said that the development of titanium alloys, especially those intended for use in living bodies began with the introduction of these alloys. In addition, it has been reported that aluminum (Al) ions are neurotoxic and inhibit bone mineralization23). Therefore, (α + β)-type titanium alloys that do not contain V and Al, such as Ti-15mass%Zr (zirconium) and Ti-15mass% Sn (tin) system alloys, were introduced24).
Then the mechanical biocompatibility focusing on Young's modulus was considered. The concept that similar elastic deformation of the bone and implant is effective for inhibiting bone resorption (stress shielding) and can allow for efficient bone remodeling was considered. Specially, since Young's moduli of metallic biomaterials are greater than that of bone, researchers have focused on lowering these values.
Simultaneously, alloy design by taking into considering problems related to allergic reactions was initiated. Elements showing a high risk of allergic reactions were avoided in metallic biomaterials: the addition of Ni, which is a high-risk element from the viewpoint of allergic reactions25), has been strongly avoided. Therefore, the development of Ni-free metallic biomaterials was commenced. In the case of stainless steels, Ni was replaced with N, so that Ni-free stainless steel with a high N content was developed2). In the case of Co-Cr-Mo alloys, Ni- and carbon (C) free Co-Cr-Mo alloys were developed3). Ni in TiNi shape memory alloys were pointed out, and so Ni-free shape memory Ti alloys were developed26,27).
As stated above, developments of Ni-free stainless steels, Co-Cr-Mo alloys, and titanium alloys were carried out. Ni-free titanium alloys will be described later in the section of 8.2. Therefore, Ni-free stainless steels and Co-Cr-Mo alloys are introduced here.
Ni-free austenitic stainless steels, that is, high nitrogen stainless steels have been developed because of Ni allergy problem2). Deformation in wrought process of Ni-free stainless steel fabricated by ordinary fabricating process is not easy. The amount of its manganese (Mn) is also high. For solving these problems, new processing for making Ni-free stainless steel was developed. This new process fabricates the final products of Ni-free stainless steel without Mn28). This new process is shown in Fig. 128). Firstly, ferritic stainless steel is prepared and made into final shape of product (in this case, wire). Then, ferritic stainless steel product is heated in nitrogen (N) gas atmosphere at high temperature. N diffuses into ferritic stainless steel product and finally Ni-free austenitic stainless steel product can be fabricated. Since the diffusion depth of N is limited to be several mm from the surface of the product, the new process is applicable to make small products including wires with a small diameter in the present state. Ni-free stainless steels developed are listed in Table 329).
New process for making Ni- and Mn-free stainless steel.
| Stainless steel | C | Cr | Mn | Mo | Si | Ni | N | Cu |
|---|---|---|---|---|---|---|---|---|
| PANCEA P558 | 0.20 | 17.4 | 10.18 | 3.09 | 0.43 | ≤0.08 | 0.48 | - |
| Biodur®108 | 0.08 | 21 | 23 | 0.7 | 0.75 | ≤0.3 | 0.97 | 0.25 |
| X13CrMnMoN18-14-3 (P2000) | 0.13 | 18 | 14 | 3 | - | ≤0.05 | 0.75–1.0 | ≤0.3 |
| 24Cr-1N | - | 24 | - | - | - | - | 1.0 | - |
| 24Cr-2Mo-1N | - | 24 | - | 2 | - | - | 1.0 | - |
| BIOSSN4 | 0.043 | 17.9 | 15.3 | 2.02 | 0.02 | ≤0.2 | 0.46 | 0.66 |
Ni-free Co-Cr-Mo alloy has been already registered in ASTM standardization, but a small amount of Ni up to approximately 1 mass% is allowed to be contained, and a small amount of carbon (C) up to approximately 0.35 mass% is also allowed to be contained. It is effective to improve the mechanical properties of the Ni-free Co-Cr-Mo alloy by utilizing hot forging process, but follows difficulty applying hot forging because the Ni-free Co-Cr-Mo alloy exhibits high work hardening, which causes brittleness. On the other hand, C also leads to brittleness of Co-Cr-Mo alloy because the carbides form along with grain boundary. Therefore, the improvement of mechanical properties and wear resistance of Co-Cr-Mo with low Ni and C, whose chemical composition is Cr: 28, Mo: 6, Ni: ≺0.01, C: 0.03–0.07, Si: ≺0.01, O: 0.09–0.003, N: 0.0006–0.0009, Co: bal, was investigated3). The fabrication process is shown in Fig. 23). The forging process is carried out in γ region followed by rapid cooling because γ phase is ductile for easy deformation and the following rapid cooling avoids the precipitation of brittle ε phase. In this case, tensile gstrength and elongation of forged Ni- and C- free Co-Cr-Mo alloy (Forged samples I, II, and III; I, II, and III indicate the differences of the lots.) are much higher than those of as-cast Co-28Cr-5.4Mo-0.19Ni-0.24C, which is equivalent to the Co-Cr-Mo alloy registered in ASTM-F75 standardization. as shown in Fig. 33). The wear rate as a function of load of forged Ni- and C- free Co-Cr-Mo alloys (Forged samples I, II, and III) is shown in Fig. 43) with that of the cast Co-28Cr-5.4Mo-0.19Ni-0.24C annealed at 1503 K, which is equivalent to the Co-Cr-Mo registered in ASTM-F75 standardization and stellite 6B. The wear rate of forged Ni- and C-free Co-Cr-Mo alloy exhibits excellent wear resistance. The mechanical properties and wear resistance are greatly improved by grain refinement and deformation induced transformation of metastable γ phase to martensite.
Fabrication process of Ni- and C-free Co-Cr-Mo alloy.
Nominal stress-nominal strain curves of forged alloys I, II and III (Co-29Cr-6Mo), and as-cast Co-29Cr-6Mo-1Ni-0.35C (F75).
Specific wear rate, ws of forged alloys I, II and III (Co-29Cr-6Mo) vs. contact load, P in comparison with those of stellite 6B and Co-29Cr-6Mo-1Ni-0.35C (F75).
Research and development of Co-Cr-Mo alloys for biomedical applications are being further energetically carried out.
Data of cell viability of pure metals are useful to design metallic biomaterials composed of non-toxic elements. The widely known data of cell viability of pure metals are those reported by Kawahara et al.30). Further data of cell viability of pure metals are needed. Then, cell viability evaluations of various metal ion extract mediums5) or various metal salts6) using L929, MC3T3-E1 or V79 cells were carried out. For example, Fig. 56) shows compatibility of 21 kinds of metal salt with L929 cell. According to the results shown in Fig. 5, metal salts are grouped into low, intermediate and high toxic categories; low toxicity: FeSo4, FeCl3, SnCl4, TiCl4, ZrCl4, NbCl5, MoCl5, TaCl5, WCl6, and Al(NO3)3, intermediate toxicity: CuCl, CoCl2, CuCl2, SnCl2, MnCl2, NiCl2, PdCl2, ZnCl2, and Cr(NO3)3, and high toxicity: VCl3 and K2Cr2O75,6).
Compatibility of 21 kinds of metal salt with L929 cell.
Since magnesium (Mg) and its alloys show similar strength and Young's moduli as those of bone and corrosive property, and Mg is one of the essential elements, Mg and its alloys were studies as biodegradable metallic biomaterials with focusing on the stents applications although Mg is not so much safe element from the view point of cell viability. The commercial Mg alloys such as AZ31 (Mg-Al-Zn-Mn system alloy), AZ91 (Mg-Al-Zn system alloy), AM60 (Mg-Al-Mn system alloy), AZ80 (Mg-Al-Zn-Mn system alloy), WE43 (Mg-Y-RE-Zr system alloy), WE54 (Mg-Y-RE-Zr system alloy), and LAE442 (Mg-Li-Al-RE system alloy) were investigated as biodegradable Mg alloys31–35). However, these commercial alloys have possibilities to contain toxic elements. Therefore, biodegradable Mg alloys contains non-toxic elements were developed. Figure 631) shows relationship between yield stress and amount of dissolved Mg ion into simulated body fluid of newly developed biodegradable Mg alloys and commercial extruded Mg alloys. Biodegradable Mg alloys showing various strength and corrosion resistance levels were developed as shown in Fig. 6. As newly developed biodegradable Mg alloys, Mg-Ca alloy, Mg-Zn-Mn alloy, Mg-Nd-Zn-Zr alloy, Mg-Zn-Ca bulk metallic glass were developed36).
Relationship between yield stress and amount of dissolved Mg ion into simulated body fluid of newly developed biodegradable Mg alloys and commercial extruded Mg alloys.
Controlling of the dissolved rate (corrosion rate) is very important factor for biodegradable Mg alloys. Surface modifications using biopolymers or bioactive ceramics are investigating to control the corrosion rate of biodegradable Mg alloys37,38).
Biodegradable iron (Fe) and its alloys such as Fe-Mn system alloys were also developed34).
For Ti-15Zr-4Nb-4Ta-0.2Pd, a Ti-Zr system alloy, endurance, which is important factor as mechanical biocompatibility, namely fatigue strength was evaluated in addition to the biological biocompatibility39). The fatigue strength of Ti-6Al-4V was also evaluated for comparison. The fatigue strength of Ti-15Zr-4Nb-4Ta-0.2Pd (730 MPa) was reported to be greater than that of Ti-6Al-4V (685 MPa). Effort have been made to develop the fatigue strength of Ti-6Al-7Nb40). For example, the fatigue strengths of Ti-6Al-7Nb subjected to aging at 833 K for 4 h after solution treatment at 1273 K for 1 h followed by water quenching (WQ) or air cooling (AC) and Ti-6Al-4V subjected to aging at 813 K for 4 h after solution treatment at 1273 K for 1 h followed by AC were evaluated and compared with each other. As a result, the fatigue strength of Ti-6Al-7Nb subjected to aging at 833 K for 4 h after solution treatment at 1273 K for 1 h followed by AC was reported to be greater than that of Ti-6Al-4V subjected to aging at 813 K for 4 h after solution treatment at 1273 K for 1 h followed by AC as shown in Fig. 719,40).
Ti-6Al-7Nb and Ti-6Al-4V ELI subjected to each heat treatment: AC and WQ indicate air cooling and water quenching, respectively after solution treatment.
Hydrogenation and dehydrogenation process, which is referred to as thermochemical processing (TCP), is effective to improve fatigue strength of (α + β)-type titanium alloys because their microstructures are refined significantly by TCP 41). The basic TCP is schematically shown in Fig. 841). Figure 941) shows the high-cycle fatigue strength of TCP treated and un-treated Ti-6Al-4V and Ti-5Al-2.5Fe for biomedical applications. The TCP treated Ti-6Al-4V and Ti-5Al-2.5Fe exhibits highly improved fatigue strength compared with untreated ones indicated by Ti-6AI-4V as-received, Ti-6A1-4V as-β transformed, and Ti-5Al-2.5Fe as-received.
Schematic drawings of thermochemical processing (TCP).
Comparison of high cycle fatigue strength of as-received, as-β transformed and TCP treated Ti-6Al-4V, and as-received and TCP treated Ti-5Al-2.5Fe.
The mechanical compatibility (mechanical biocompatibility) between the bone (living tissue) and the implant can be given as the biocompatibility looking from both living body and implant side. Namely, mechanical biocompatibility refers to the similarity in mechanical properties between the bone and the implant. The main factors to be considered are prevention of failure or abnormalities, and homogeneous stress transfer. The latter is particularly regarded as an important factor. This leads to a similar Young's modulus between the bone and the implant, as already stated above. Therefore, lowering Young's modulus is imperative for designing metallic biomaterials. Ti alloys are advantageous because their Young's moduli are lower than those of stainless steels (SUS 316L) and Co-Cr-Mo alloys, as also stated above. Ti alloys are attracting much attention because they exhibit excellent corrosion resistance, specific strength, and good balance between the strength and the ductility in addition to excellent biocompatibility. Ti alloys are roughly grouped into α-, (α + β)- and β-types according to their constituent phases. The crystal structures of α- and β-type Ti alloys are hexagonal close-packed (hcp) and body centered cubic (bcc) structures, respectively. Since the packing ratio of atoms is less in the β-type Ti alloys than in the α-type Ti alloys, Young's moduli of β-type Ti alloys are expected to be lower than those of α- and (α + β)-type Ti alloys. Therefore, efforts have been devoted to developments of β-type Ti alloys composed of non-toxic and allergy-free elements with low Young's moduli for biomedical applications. Many such β-type Ti alloys have been developed to date, and some of them have been standardized in ASTM, ISO, or JIS.
Since β-typeTi alloys with a low Young's modulus generally contain a large amount of high-cost elements such as Nb, Ta, Mo, and Zr, alternative alloys containing low cost elements such as Fe, Mn or Cr are being developed. The representative low Young's modulus β-type Ti alloys developed to date are listed in Table 142).
| β-type titanium alloys | ASTM standard | ISO standard | JIS standard |
|---|---|---|---|
| Ti-13Nb-13Zr | ASTMF F1713 | - - - | - - - |
| Ti-12Mo-6Zr-2Fe (TMZF) | ASTM F1813 | - - - | - - - |
| Ti-12Mo-5Zr-5Sn | - - - | - - - | - - - |
| Ti-15Mo | ASTM F2066 | - - - | - - - |
| Ti-16Nb-10Hf (Tiadyne 1610) | - - - | - - - | - - - |
| Ti-15Mo-2.8Nb-0.2Si | - - - | - - - | - - - |
| Ti-15Mo-5Zr-3Al | - - - | ISO 5832-14 | JIS T 7401-6 |
| Ti-30Ta | - - - | - - - | - - - |
| Ti-45Nb | - - - | - - - | - - - |
| Ti-35Zr-10Nb | - - - | - - - | - - - |
| Ti-35Nb-7Zr-5Ta (TNZT: TiOsteum) | ASTM Task Force 4.12.23 | - - - | - - - |
| Ti-29Nb-13Ta-4.6Zr (TNTZ) | - - - | - - - | - - - |
| Ti-35Nb-4Sn | - - - | - - - | - - - |
| Ti-11.5M0-6Zr-4.5Sn | ASTM F 1713 | - - - | |
| Ti-50Ta | - - - | - - - | - - - |
| Ti-8Fe-8Ta | - - - | - - - | - - - |
| Ti-8Fe-8Ta-4Zr | - - - | - - - | |
| Ti-35Nb-2Ta-3Zr | - - - | - - - | - - - |
| Ti-22.5Nb-0.7Zr-2Ta | - - - | - - - | - - - |
| Ti-23Nb-0.7Ta-2.0Zr-1.2O (Gum Metal) | - - - | - - - | - - - |
| Ti-28Nb-13Zr-0.5Fe (TNZF) | - - - | - - - | - - - |
| Ti-24Nb-4Zr-7.9Sn (Ti2448) | - - - | - - - | - - - |
| Ti-7.5Mo | - - - | - - - | - - - |
| Ti-12Mo-3Nb | - - - | - - - | - - - |
| Ti-12Mo-5Ta | - - - | - - - | - - - |
| Ti-12Cr | - - - | - - - | - - - |
| Ti-30Zr-7Mo | - - - | - - - | - - - |
| Ti-30Zr-3Mo-3Cr | - - - | - - - | - - - |
| Ti-5Fe-3Nb-3Zr | - - - | - - - | - - - |
Young's moduli of representative α-, (α + β)- and β-type Ti alloys are shown in Fig. 1043,44). Young's moduli of β-type Ti alloys for biomedical applications are apparently lower than those of α- and (α + β)-type Ti alloys. Although there is a fairly large difference among Young's moduli according to the measurement methods, the corresponding values of β-type Ti alloys are approximately 60–80 GPa under solutionized conditions giving a single β phase. The Young's modulus values can be reduced further by severe cold working after solution treatments45). These trends are seen in poly-crystallineβ-type Ti alloys. However, in the case of a single crystal, the lowest Young's modulus is, for example, 35 GPa in the <100> direction for Ti-29Nb-13Ta-4.6Zr (referred to as TNTZ) as shown in Fig. 1146). This value is comparable to that of bone, which is approximately 10–30 GPa. Therefore, artificial bone made of single crystal β-type Ti alloys with low Young's modulus is highly desired for biomedical applications.
Young's moduli of representative α-type, (α+β)-type, and β-type titanium alloys for biomedical applications.
Young's modulus of single-crystal Ti-29Nb-13Ta-4.6Zr (TNTZ) in directions between [100] and [110].
The effectiveness of the low Young's moduli of implants for inhibiting stress shielding, namely bone resorption and facilitating bone remodeling was proved by evaluating bone fracture healing and bone remodeling situations, where intramedullary rods and bone fracture fixation plates were implanted into fracture model made in tibia of Japanese white rabbits47,48). For example, in the study on bone remodeling, utilizing bone plates made of SUS 316L stainless steel, Ti-6Al-4V ELI, and TNTZ, healing conditions were observed using X-ray photographs recorded at regular intervals over a period of up to 48 weeks after implantation48). Following this, both tibiae were extracted along with the bone plate, and the bone formation was externally observed. For all the materials, callus formation was observed at 2 weeks, which became distinct at 3 weeks and then bone union was obtained at 4 weeks after implantation, and the fracture line was barely discernible at 8 weeks after implantation. Moreover, there was no trace of the experimental fracture at 16 weeks to 20 weeks after implantation. However, bone atrophy (thinning of the bone) was observed under the bone plate, and the time at which this occurred varied between the different materials. According to the X-ray images taken from 4 weeks after implantation to 18 weeks after implantation for each plate, for SUS 316L stainless steel, the thinning of the bone was first observed at 7 weeks after implantation, and the bone almost disappeared at 12 weeks after implantation. For Ti-6Al-4V ELI, the thinning was first observed at 7 weeks after implantation, and the bone almost disappeared at 14 weeks after implantation. Finally, for TNTZ, the corresponding times were 10 weeks and 18 weeks after implantation. Furthermore, an increase in the diameter of the tibia and the double wall structure in the intramedullary bone tissue was observed only for the case of the bone plate made of TNTZ, as shown in Fig. 1248). In this figure, the inner wall bone structure represents the original bone, i.e., the remaining old bone, whereas the outer wall bone structure is the newly-formed part. This bone remodeling is the direct result of using a bone plate with a low Young's modulus. Therefore, in the implant made of the β-type Ti alloy with low Young's modulus for biomedical applications, stress transfer between the bone and the implant becomes homogeneous and good bone remodeling is achieved although this has been proved only at the stage of animal testing.
CMRs (contact micro-radiographs) of cross sections of fracture models implanted with and without bone plates made of TNTZ at middle and distal level at 48 weeks after implantation : (a) cross section of fracture model, (b) parts of □ of (a), namely high magnification CMR of branched parts of bones formed outer and inner sides of tibiae, and (c) cross sections of unimplanted tibiae.
For shape memory alloys, the practical applications of TiNi alloys are developing. TiNi alloys have been applied as guide wires of catheters and orthodontic wires. However, Ni is a high-risk element from the viewpoint of allergic reactions as stated above. Therefore, many Ni-free shape memory β-type Ti alloys for biomedical applications have been developed. The representative Ni-free β type titanium alloys for biomedical applications are listed in Table 249). They are roughly grouped into Ti-Nb, Ti-Mo, Ti-Ta, and Ti-Cr system alloys. These show a low Young's modulus.
| Alloy system | Shape memory alloy |
|---|---|
| Ti-Nb system | Ti-Nb, Ti-Nb-O, Ti-Nb-Sn, Ti-Nb-Al, Ti-22Nb-(0.5-2.0)O (at%), Ti-Nb-Zr, Ti-Nb-Zr-Ta, Ti-Nb-Zr-Ta-O, Ti-Nb-Ta-Zr-N, Ti-Nb-Mo, Ti-22Nb-6Ta(at%), Ti-Nb-Au, Ti-Nb-Pt, Ti-Nb-Ta, Ti-Nb-Pd |
| Ti-Mo system | Ti-Mo-Ga, Ti-Mo-Ge, Ti-Mo-Sn, Ti-Mo-Ag, Ti-5Mo-(2-5)Ag (mol%), Ti-5Mo-(1-3)Sn (mol%), Ti-Sc-Mo |
| Ti-Ta system | Ti-50Ta (mass%), Ti-50Ta-4Sn (mass%), Ti-50Ta-10Zr (mass%) |
| Others | Ti-7Cr-(1.5, 3.0, 4.5)Al (mass%) |
| Super elastic alloy: alloy shows only elasticity | |
| Ti3(Ta + Nb + V) + (Zr, Hf) + O (at%), Ti-29Nb-13Ta-4.6Zr (mass%) |
Development of Ti alloys has focused on the functionality with a low Young's modulus, which is advantageous for patients because a low Young's modulus inhibits stress shielding, between the bone and the implant, as previously described. The development of Ti alloys for biomedical applications that are also advantageous for surgeons has recently progressed. For example, in operations for scoliosis disease, the rods of the spinal fixation devices undergo bending when they are manually handled by surgeons within the small space inside a patient's body, to achieve in-situ spine contouring50). The contoured shape of the rod should be maintained. Therefore, reverse bending, known as springback, of the bent rod should be prevented. For implant rods, the degree of springback should be low to facilitate their handling during operations. The degree of springback is thought to depend on both the strength and Young's modulus of the implant rod. If two implant rods with the same strength, but differing Young's moduli, are used, the implant rod with the lower Young's modulus will exhibit greater springback, as schematically shown in Fig. 1351). Therefore, about the patients' needs, low Young's modulus is required to prevent stress shielding, whereas to facilitate surgery, high Young's modulus is required to prevent springback. To satisfy these conflicting demands simultaneously, it should be possible to increase the Young's modulus value of the bent parts of the rod via deformation at room temperature, while allowing Young's modulus of the remainder of the rod to remain unchanged at a low value50). The metallic rods used in spinal-fixation devices are required to have a low Young's modulus, good biocompatibility, and low springback. Accordingly, it is necessary to develop novel Ti alloys that offer good biocompatibility and an adjustable Young's modulus. To accomplish this, it should be possible to increase the local Young's modulus to a high value at certain parts of the device via deformation at room temperature, while allowing the Young's modulus of the remainder of the device to remain unchanged at a lower value50). Figure 1450) schematically shows the concept of self-adjustment of Young's modulus in an implant rod, which has been proposed as a solution for the aforementioned problem. In the case of certain metastable β-type titanium alloys, non-equilibrium phases such as the α′, α″, and ω phases appear in the β matrix during deformation52–55). Among these non-equilibrium phases, the ω phase has a higher Young's modulus than the β phase. Since the precipitation of the α′, α″, and ω phases is significantly related to the stability of the β phase, the appropriate combination of Ti and its alloying element is important to control the stability of the β phase. Ultimately, the β-type Ti-Cr system alloys for biomedical applications, in which the ω phase is induced, are being developed50,53).
Relationship between Young's modulus and springback.
Schematic explanation of concept of spinal fixation rod with self-adjustable Young's modulus.
It is important to determine appropriate chemical compositions of Ti-Cr system alloys, in which the ω phase is most effectively induced by deformation and the degree of increase in Young's moduli is the greatest.
Figure 1555) shows the transmission electron microscope (TEM) diffraction patters of Ti-(10-14) Cr subjected to solution treatment (ST) (Ti-(10-14)-ST). The electron diffraction patterns of Ti-10Cr-STshow weak extra spots, in addition to the spots derived from the β phase, suggesting that a small amount of the athermal ω phase is formed in Ti-10Cr-ST during water quenching (Fig. 15 (a)). For Ti-11Cr-ST, the intensities of the ω reflections decrease, accompanied by a change to circular diffuse streaks. As shown in Fig. 15 (b)–(e), as the Cr content continues to increase, the circular diffuse streaks are weakened. It is well known that the intensity of the ω reflection is related to the amount of the ω phase. The amount of the athermal ω phase in the designed alloys is found to be dependent on the stability of the β phase. Specifically, as the Cr content increased, the β phase became more stable, so that the formation of the athermal ω phase during water quenching is suppressed. Therefore, the amount of the athermal ω phase in Ti-11Cr-ST is lower than that in Ti-10Cr-ST, and the athermal ω phase disappears almost completely in Ti-12Cr-ST, Ti-13Cr-ST, and Ti-14Cr-ST. In order to achieve the largest amount of the ω phase induced by deformation, the amount of the athermal ω phase should be as small as possible. Eventually, the Cr content at which the circular diffuse streaks related to the ω phase almost disappeared is found to be 12 mass%.
TEM diffraction patters of Ti-(10-14)Cr subjected to solution treatment (ST).
In contrast, Fig. 1655) shows the TEM diffraction patters of Ti-(10-14) Cr subjected to cold rolling (CR) with a reduction ratio of 10% as a simulation of bending deformation after ST (Ti-(10-14) Cr-CR). In contrast to the case of Ti-(10-14) Cr-ST, after CR, the ω reflections in Ti-10Cr-CR (Fig. 16 (a′)) are much sharper than those in Ti-10Cr-ST (Fig. 15 (a)), indicating that the amount of the ω phase increases in response to cold rolling. Furthermore, the intensities of the ω reflections are higher in Ti-11Cr-CR and Ti-12Cr-CR (Fig. 16 (b′) and (c′)) when compared to those in Ti-11Cr-ST and Ti-12Cr-ST (Fig. 15 (b) and (c)). These findings confirm that the deformation-induced ω phase transformation occurs in the Ti-10Cr, Ti-11Cr, and Ti-12Cr alloys during cold rolling. The amount of the ω phase in Ti-10Cr-CR, Ti-11Cr-CR, and Ti-12Cr-CR is too small to be detected by X-ray diffraction (XRD) analysis, but the deformation-induced ω phase is evident in TEM observation. Specifically, as the Cr content increases, the intensities of the ω reflections decreased. In addition, only diffuse streaks are observed in Ti-13Cr-CR and Ti-14Cr-CR (Fig. 16 (d′) and (e′)), suggesting that the amount of the ω phase does not change during cold rolling. These findings demonstrate that deformation-induced ω phase transformation does not occur in Ti-13Cr and Ti-14Cr alloys during cold rolling and that it is dependent on the stability of the β phase. In lower stability alloys such as Ti-10Cr, Ti-11Cr, and Ti-12Cr, deformation-induced ω phase transformation may occur, while with increasing Cr content, the β phase can become more stable, thereby suppressing the deformation-induced ω phase transformation.
TEM diffraction patters of Ti-(10-14)Cr subjected to 10% cold rolling (CR) after solution treatment (ST).
Figure 1755) shows Young's moduli of Ti-(10-14) Cr alloys subjected to solution treatment and cold rolling, where cold rolling at a reduction ratio of 10% was carried out to simulate deformation. The degree of increase in Young's modulus is the highest at a Cr content of 12 mass%. Therefore, the chemical composition that exhibits the lowest spring back, is expected to be Ti-12Cr.
Young's moduli of Ti-(10-14)Cr alloys subjected to solution treatment (ST) and 10% cold rolling (CR) after ST.
Figure 1855) shows comparison profiles of the ratio of springback per unit stress as a function of the applied strain for TNTZ, Ti-12Cr, and Ti-6Al-4V ELI (Ti64 ELI). The ratios of springback per unit stress of all the alloys show a similar trend; initially decreasing significantly and then remaining approximately stable with increasing applied strain. The springback of Ti-12Cr is significantly lowered as compared with that of TNTZ, and is similar to that of Ti-6Al-4V ELI.
Ratio of springback per unit load as a function of applied strain for Ti-12Cr, Ti-6Al-4V (Ti64) ELI, and TNTZ.
To achieve more increase in Young's modulus than that of Ti-12Cr, which has an a thermal ω-phase, O is added to the alloy because O suppresses the formation of the athermal ω-phase. Finally，Ti-11Cr-0.2O, which exhibited a higher degree of increase in Young's modulus than that of Ti-12Cr has been develop56).
Ti and its alloys show the greatest biocompatibility among metallic materials for biomedical applications. However, they are grouped into bioinert materials as well as ceramics like alumina, zirconia, etc., judging from the pattern of osteogenesis15) However, their biofunctionalities are poor. Therefore, surface modifications using bioactive ceramics such as calcium phosphate (Ca-P) and hydroxyapatite (HAp: Ca(PO4)3OH)7), or blood compatible polymers such as polyethylene glycol (PEG)57) and segment polyurethane58) are carried out on Ti and its alloys to achieve high level biocompatibility with living tissue.
The biocompatible surface modifications are achieved, in general, using dry process and wet process. There are various dry and wet processes59).
Dry processes are, for example, plasma spray method, ion plating, RF magnetron sputtering, pulse laser deposition method, ion beam dynamic mixing method, super plastic joining method, etc., where HAp is formed directly on Ti alloy surface, and calcium ion implantation, calcium ion mixing method etc., where HAp are formed indirectly on Ti alloy surface7,59) Wet processes are, for example, electrochemical treatment etc., which are direct HAp forming methods, and alkali treatment, etc., which are indirect HAp forming methods7,59).
There is another interesting method60) where the powder of calcium phosphate invert glass mixed with distilled water is coated on the surface of the Ti alloy followed by heating at around 1073 K, and then phosphate calcium type ceramics precipitate.
Figure 1961) shows the progress in research and practical stage of surface treatment and modification for achieving bone formation function in dental implants as an example. The research stage of surface treatment and modification is progressed up to the 4th stage generation. The progress into the 5th stage generation is desired. The practical stage is progress up to the middle of the 3rd stage, and its further progress is required.
Progress in research and practical stage of surface treatment and modification for achieving bone formation function in dental implants as an example.
The magnetic susceptibilities of metallic biomaterials such as Zr-1Mo8), Ti-6Al-4V, CP-Ti and Ti-6Al-7Nb, Co-Cr-Mo alloy, and SUS 316L stainless steel are 0.98 × 10−6 cm3·g−1, 3.5 × 10−662), 3.0 × 10−662), 7.7 × 10−662), and 3380 × 10−6 cm3·g−163), respectively. Therefore, the magnetic susceptibilities of zirconium (Zr) alloys are lower than those of Ti and its alloys, Co-Cr-Mo alloy and SUS 316L.
Magnetic resonance imaging (MRI) has become a powerful diagnostic tool in orthopedics. MRI diagnostic is inhibited by the presence of metallic implants in the body because they are magnetized in the intense magnetic field of the MRI instrument. That may cause image artifact leading to prevent exact diagnosis64). Therefore, Zr alloy implants are advantageous to prevent the artifact of the MRI image.
Therefore, Research and development of Zr alloys for biomedical applications are nowadays energetically carried out. Zr-Nb9,65), Zr-Mo8), Zr-Ti66), Zr-Cu67) alloys, etc. for biomedical applications have been reported.
Figure 209) shows Young's moduli of Zr-5Nb, -10Nb, -20Nb, and -30Nb. The lowest Young's modulus is obtained to be approximately 40 GPa at a Nb content of 20 mass%. This value is lower than those of low Young's modulus Ti alloys as mentioned above. Therefore, Zr alloys are advantageous from the view point of low magnetic susceptibilitiy and low Young's modulus for biomedical applications.
Young's moduli and Vickers hardness of Zr-5Nb, -10Nb, -20Nb, and -30Nb.
Nowadays, additive manufacturing (AM), namely 3D printing, wherein metal powders are accumulated layer by layer to make products without a mold. In the AM methods, laser or electron beam melting is used, which is called as selective laser melting (SLM)10) or electron beam melting (EBM)11).
The AM method of fabrication has attracted attention in many fields including fabrication of implants, especially porous structural implants. The Young's modulus of titanium and its alloys are easily controlled by making their porous body (Fig. 21)68) by AM methods. As an example of the animal test of Ti porous implant made by AM method (EBM method), lateral X-ray and micro-computed tomography (CT) images of porous Ti-6Al-4V without the bone graft made by AM method (EBM) and PEEK cages with the bone graft at 3 months and 6 months after implantation into the spine of sheep are shown in Fig. 2269). At 3 months after implantation (A1-3, B1-3), porous Ti-6Al-4V cages without the bone graft exhibit better osteointegration with surrounding bone tissue as compared with PEEK (Polyether-ether-ketone) cages with the bone graft, as shown by the reduced radiolucent region and smaller empty gap (white arrow head) around the edge of the cages. At 6 months after implantation (C1-3, D1-3), both cages achieve complete spine fusion. However, micro-CT shows that there are still small empty gaps around PEEK cages with the bone graft whereas porous Ti-6Al-4V cages without the bone graft are well integrated with surrounding bone tissue. The growth of the bone tissue into the porous Ti-6Al-4V cage is evident in the figure. The porous Ti-6Al-4V cage without the bone graft exhibits better osteointegration as compared with that of the bone grafted PEEK.
Young's modulus of porous Ti as a function of porosity.
Lateral X-ray and micro-CT images of porous Ti-6Al-4V fabricated AM method and PEEK cages at 3 months (A1-3, B1-3) and 6 months. (C1-3, D1-3) after implantation.
It has been recognized very recently that not only bone mineral density (BMD), but also bone orientation (biological apatite (BAp) crystal orientation) is important for the bone regeneration70).
Figure 2370) shows variations in the relative intensity ratio of the (002) diffraction peak to the (310) peak to evaluate the orientation degree of the BAp c-axis orientation with different directions A, B, and C for the ulna, skull bone, and mandible71,72). It can be concluded from this figure that the preferential orientation of the BAp c-axis corresponds to the in vivo stress distribution and BAp c-axis tends to orient along the principal stress direction in the original bones
Variations in the relative intensity ratio of the (002) diffraction peak to the (310) peak to evaluate the orientation degree of the BAp c-axis orientation with different directions A, B, and C for the ulna, skull bone, and mandible.
Figure 2470) shows degree of BAp orientation intensity ratio of (002)/(310) showing recovery of local BMD and BAp c-axis alignment in the regenerated ulna. The diffracted intensity ratio of (002)/(310) in the regenerated bone tissue increases significantly Significant increases were observed from 12 through 24 weeks in both local BMD and intensity ratio. There are no significant differences from baseline (intact bone) at 24 weeks in either local BMD or the intensity ratio. Thus, in regeneration of rabbit ulnas, where controlled release of rBMP-2 is carried out, recovery of the preferential BAp c-axis orientation tends to follow that of BMD.
Degree of BAp orientation intensity ratio of (002)/(310) showing recovery of local BMD and BAp c-axis alignment in the regenerated ulna.
Above results leads to groove design of the surface of the implant, which is favorable for the rapid bone fusion with implant as shown in Fig. 2512). A maximum principle stress is found to be distributed in 60° grooves by FEM: this alignment is favored for bone growth12).
3D FEA (finite element analysis) model of the beagle femur–artificial hip joint implant used in this study. (A) Two types of groove angle combinations on the surface of the proximal medial region ofthe femoral stem: (60°/0°/−60°) groove and (30°/0°/−30°) groove. (B) Definition of the angle, θ, between groove wall and the maximum principal stress direction in the groove. (C) Loading and boundary conditions used for FEA.
Mechanical biocompatibility is evaluated by taking into consideration both the biomaterials and the living tissue (bone). The similarity in Young's moduli between the biomaterial and the living tissues a measure of mechanical biocompatibility has been considered for designing metallic biomaterials, with focus on Ti alloys. The design of Ti alloys with self-adjusting Young's modulus for biomedical applications, which satisfy the demands of both patients and surgeons (i. e., a low Young's modulus for patients and a high Young's modulus for surgeons) has progressed. Imparting mechanical biocompatibility to biomaterials, which self-adjusts to match that of living tissue, is required in addition to controlling general mechanical properties. Research on the biofunctioal surface modification needs progress into the 5th generation stage. Nowadays, research and development of Zr based alloys for biomedical applications are getting much attentions. Additive manufacturing (AM) using SLM and EBM are aggressively applied to fabricate implants made of metallic biomaterials. Bone orientation design on the surface of the implant is effective for the rapid fusion of implant with bone.
