Findings and perspectives of β-Ti alloys with biomedical applications: Exploring beyond biomechanical and biofunctional behaviour

Findings and perspectives of β-Ti alloys with biomedical applications: Exploring beyond biomechanical and biofunctional behaviour Amir A. Elhadad a , L. Romero-Resendiz b , M.C. Rossi c , L.M. Rodríguez-Albelo d , Sheila Lascano e , Conrado R.M. Afonso c , A. Alcudia f , V. Amig´ o g , Y. Torres d,* a Department of Pre-Clinical Oral Health Sciences, College of Dental Medicine, QU-Health, Qatar University, Qatar b Facultad de Química, Departamento de Ingeniería Metalúrgica, Universidad Nacional Aut´ onoma de M´ exico, Mexico City, 04510, Mexico c Materials Engineering Department (DEMa), Universidade Federal de S˜ ao Carlos (UFSCar), 13.565-905, S˜ ao Carlos, SP, Brazil d Departamento de Ingeniería y Ciencia de los Materiales y del Transporte, Escuela Polit´ ecnica Superior, Universidad de Sevilla, 41011, Seville, Spain e Departamento de Ingeniería Mec´ anica, Universidad T´ ecnica Federico Santa María, Avda. Vicu˜ na Mackenna Poniente N 3939- San Joaquín, 8320000, Santiago, Chile f Departamento de Química Org´ anica y Farmac´ eutica, Facultad de Farmacia, Universidad de Sevilla, 41012, Seville, Spain g Instituto de Tecnología de Materiales, Universitat Polit` ecnica de Val` encia, Camino de Vera s/n, 46022, Valencia, Spain ARTICLE INFO Handling editor: P.Y. Chen Keywords: Beta-titanium alloys Novel manufacturing routes Surface modification treatments Bio-functional and tribomechanical behaviour Cellular and bacterial response ABSTRACT Early implant failure and bone resorption may occur in load-bearing conditions as a result of stress shielding brought on by a mismatch in the bone-Ti-implant modulus. A review with a novel multidisciplinary perspective is proposed in this work, which considers recent developments of β-Titanium alloys and new trends in novel microstructures, processing techniques, properties of dense and porous substrates, as well as the relationship between all these aspects and performance in service, in terms of improved its biomechanical and bio-functional balance. In addition to highlighting several modern and historical uses for Ti alloys, this review covers many cutting-edge novel β-Ti alloys and uses that promise to exceed historical standards. Also, it deepens through several important properties of these alloys, including toxicity of alloying elements, phase stability, thermomechanical processing, heat treatment, surface, and stress-induced modifications. The stiffness, hardness, fatigue and wear resistance, corrosion behaviour, biocompatibility, and manufacturing and surface modification effects on these parameters are also emphasized. In-vitro and in-vivo assays have been added to highlight important aspects of bioactivity and antibacterial behaviour, and future significant research areas are suggested along with new techniques to ensure the successful clinical application of β-Ti alloys. 1. Introduction In the US and Europe, 500,000 individuals undergo bone restoration operations each year [1]. In 2012, the market for biomaterials was at $94.1 billion USD; as of 2017, it is worth $134.3 billion USD [2]. Whereas the entire cost of expenses related to bone repair is expected to be US$ 17 billion annually, the annual global expenditures on bone fractures and orthopaedics are estimated to be 6.5 billion USD and 4.7 billion USD, respectively [3]. Several diseases or accidental bone fractures lead to the requirement to repair the specific damaged organs. Biological activity is unfortunately limited at the region of the damaged bone tissue, making implantation necessary. Once tissue is lost due to severe trauma or injury in an ageing population, it has a major clinical and socioeconomic impact on many people’s quality of life [4,5]. Additionally, the ageing of populations in developed nations and patients’ desire to retain their normal level of physical activity have also contributed to a remarkable rise in the usage of implants during the last few years [6]. Choosing the appropriate implant materials that can serve as a bridge and support and, having the proper mechanical and bone induction capabilities, is crucial [7]. Treatments for bone repair relies on autografts, allografts, xenografts, and synthetic scaffolds [8]. Due to their osteoinductive and osteoconductive capabilities, autografts are widely recommended; yet, they have significant downsides, including a limited availability and donor-site morbidity [9,10]. While successfully resolving the mentioned drawbacks, allografts and xenografts are nevertheless susceptible to immune rejection, and transmission of infection [9]. Use of synthetic biomaterials, such as metals, ceramics, and polymers, is another potential treatment approach [11]. However, * Corresponding author. E-mail address: [email protected] (Y. Torres). Contents lists available at ScienceDirect Journal of Materials Research and Technology journal homepage: www.elsevier.com/locate/jmrt https://doi.org/10.1016/j.jmrt.2024.09.248 Received 15 July 2024; Received in revised form 20 September 2024; Accepted 28 September 2024 Journal of Materials Research and Technology 33 (2024) 3550–3618 Available online 5 October 2024 2238-7854/© 2024 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC license ( http://creativecommons.org/licenses/bync/4.0/ ). due to a variety of issues, such as low support strength and the brittleness of inorganic ceramics, biocompatible metals have emerged as the ideal material for implants [12]. In addition, the criteria for specialized designs, such as load transmission, stress distribution, and implant articulation to facilitate movement in prosthetic knee joint applications, are likewise satisfied by metal-based implants [13]. Stainless steels, commercially pure titanium (CP–Ti), Ti–6Al–4V, and cobalt-chromium-molybdenum (Co–Cr–Mo), alloys are one of the most commonly employed metallic implants [5]. The Ti–6Al–4V alloy is an ideal choice for biomedical implants due to its biocompatibility, ease of manufacture, and superior corrosion resistance when compared to other alloys [5,14–16]. So far, one of the primary drawbacks of implants made of standard materials is the apparent mismatch between the mechanical characteristics of the implant and those of bone. The merely linear stress-strain attributes of metallic implants, which are considerably stiffer and do not match the plateau-like hysteretic activity of bones, causes implant loosening [17]. Accordingly, metallic implants have the downside of stress shielding due to their high elastic modulus, which might lead to bone resorption and eventually the early failure of the implant under load-bearing conditions [17]. Another major issue that could have an impact on the biological environment is the harmful consequences of the V and Al ions that are released by metallic implants when they are implanted in-vivo [18,19]. More research data indicate the possibility that vanadium (V) ions may alter the kinetics of an enzyme activity related to inflammatory response cells in human tissues. Aluminium (Al), on the other hand, raises the risk of Alzheimer’s disease [20]. Another major downside of the Ti–6Al–4V alloy is the implant-associated infection (IAI) continues to be a major clinical concern despite advances in the life expectancy of Ti alloy (Ti–6Al–4V) implants used in applications including joint and tooth replacement [21, 22]. The implant may get infected directly during surgery by contamination, indirectly by a contiguous route in which bacteria are transferred to the implant from the surgical site, an open traumatic wound, an infection at a nearby site, or through haematogenous seeding [23]. There has been a great deal of research in the field of super elastic β-Ti alloys driven primarily by the desire to develop an implant based on metallic materials that would combine super elasticity with a biocompatibility comparable to pure Ti. The β-Ti alloys have received a lot of attention over the recent decades due to the use of other biocompatible chemical elements that replace V and Al as a reasonable solution to obtain Ti alloys with a stiffness quite similar to that of the cortical bone tissue that they intend to replace as well as long-term implant rejection [24–26]. The usage of non-toxic, low elastic modulus β-Ti alloys is envisaged for the development of the next generation of implanted devices [27]. Such an alloy is completely biocompatible and possesses amazing mechanical, chemical, and electrochemical properties for the manufacture of bio implants. Despite numerous reviews of β-Ti alloys for biomedical applications in the literature, many of them omit or do not go into detail about specific aspects of the microstructure, manufacturing processes, and characteristics of this type of alloy, considering the real-world requirements in customer experience and the problems that negatively affect the clinical success of this type of metal alloys for replacing damaged bone tissues. For instance, among the most crucial measures affecting the strength of the bone-to-implant contact is surface roughness [28]. The implant’s surface roughness affords the bone a chance to grow and interlock with the implant. Cell distribution is affected by the geometry of the pores, their internal structure, and their pattern of connectivity. When the implant has a porous structure over its whole volume, the rate of osseointegration may be considerably higher [29]. The use of surface modifications, such as anodizing, grit-blasting, acid-etching, and resorbable blast media, has proven effective in improving bone regeneration around dense implants [30]. However, in individuals with poor bone density and osteoporosis, acid-etched and sandblasted Ti surfaces contributed to lower implant stability [31,32]. An implant with a porous structure will facilitate the attachment of biological bone to the implant [11]. Studies have demonstrated that porous titanium implants significantly enhance the rate of tissue ingrowth as well as the strength of the implant’s attachment to the bone [30]. Stress shielding brought on by a bone-Ti-implant modulus mismatch leads to poor osseointegration and bone formation [33]. Pore introduction is one approach for lowering the elastic modulus of pure bulk Ti to a value closer to that of human bone [33]. To reduce the impacts of stress shielding and boost porosity for bone tissue ingrowth, another strategy of implant surface treatment is the creation of porous structures [30]. Porous metal, also described as trabecular metal in the field of orthopaedic use, is a type of metal structure having a 3D network of interconnecting pores. When compared to dense metal, porous metal shows less rigidity and strength [12]. Its Young’s modulus is within or close to that of the characteristics of human bone (from 0.4 GPa of trabecular bone to 17.9 GPa of cortical bone) [34]. High friction coefficients improve the porous metal’s initial stability, and connected porosities and bone growth into the porous Ti implant provide long-term stability [12]. Porous Ti implants have larger surface areas compared to solid Ti implants, which enhances bone cell attachment and growth, allows bone ingrowth into the implant, and accelerates osseointegration [35]. Thus, porous Ti appeared to be useful choice for usage in biomedical field, such as prosthetics and scaffolds for tissue engineering. However, more than 50 GPa is the elastic modulus of β-Ti, exceeding the elastic modulus of human bone [36]. Due to the stiffness mismatch, the implants are subjected to most of the stress. The majority of the stress is placed on the implants as a result of the mismatch in stiffness. While going through the healing process, human bone regenerates and grows. As a result, the bone density around the implants may diminish over time due to a lack of load, which can lead to bone resorption, a stress shielding phenomenon, and slowed-down healing times [37]. Contrarily, a material with a low elastic modulus will deform substantially under stress and cease to have a supporting function [38]. Moreover, implanted Ti alloys are difficult to deeply bind to biological tissue, which increases the risk of implant fracture and tissue loss [5]. Consequently, research into the correlations between elastic modulus and chemical composition, hardness, and porosity is crucial. Various chemical compositions and microstructures of β-Ti alloys were designed to address the concern of stress shielding; some of them have been applied to the biomedical field. Even though there are many works reviewing β-Ti alloys for biomedical applications in the literature, large number them avoid mentioning or do not go into considerable detail regarding specific aspects of the microstructural features, manufacturing processes, and characteristics of this type of alloy, taking into consideration the practical needs in service and the issues that adversely impact the clinical success of this type of metal alloys replacing damaged bone tissues. In this context, the review on β-Ti alloys attached offers a multidisciplinary perspective, emphasizing the balance between biomechanical and biofunctional properties, which is a unique approach compared to previous literature. It covers recent developments in microstructures, processing techniques, and the service performance of both dense and porous β-Ti alloys. This comprehensive view, which includes in-vitro and in-vivo assays for bioactivity and antibacterial behaviour, and the discussion on the toxicity of alloying elements, phase stability, and surface modifications, sets it apart from existing reviews that may not delve into these aspects in detail. Additionally, the review suggests future research directions and new techniques for clinical application, highlighting a forward-looking perspective on β-Ti alloy development. The scope of this review is outlined schematically in Fig. 1. 2. Effect of alloying elements in β-Ti alloys This section describes the impact of different alloying elements on the crystal structure stability, toxicity, reactivity, and cost. The design of new and improved β-Ti alloys is under intense worldwide research. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3551 2.1. Selection and toxicity of alloying elements One of the primary concerns when using metallic implants in-vivo is the toxicity of the ions that are released [18,19]. As explained in the introduction, CP-Ti and Ti–6Al–4V are preferred compared to other metallic implants used in joint replacement since the 1940s [39]. Nevertheless, study findings and long-term clinical applications revealed that adding Al and V to a Ti–6Al–4V alloy would have significant negative physiological impacts [20]. Al has been associated with a variety of diseases, including brain damage, digestive issues, neurotoxicity, contact dermatitis, breast cancer, osteomalacia, anaemia, and encephalopathy [40]. Al can also damage the brain and cause encephalopathy, Parkinson’s disease, Alzheimer’s disease, neurotoxicity, and male infertility [41,42] In very large doses, aluminium (Al), changes the mineralization of the bone system, leading to pathological fractures, which can have a deleterious impact on bone metabolism [43]. Vanadium (V) is present in extremely small amounts in the human body [44]. Although it seems to not generate substantial toxic reactions when present in low concentrations (1–4%) in the Ti–6Al–4V alloy, it does exhibit some toxic phenomena that, over time, may have negative effects on human health [45,46]. Diarrhea, vomiting, weight loss, glomerulonephritis, pyelonephritis, seminiferous epithelium degradation, male Infertility, and cell viability are some of the severe symptoms [45,46]. The toxicity of Cobalt (Co) can affect several neurologically connected organ systems, (tinnitus, vision loss, vertigo), depression, lack of concentration, headaches, impaired reflexes, vibration/position sense, as well as limb paraesthesia [47,48]. Chronic exposure to Co was related to lung cancer, some respiratory system cancers, DNA damage in sperm, and the activation of DNA repair genes in the testes, fertility reductions in male and female individuals by lowering sperm production and ova quality, and signs of reduced reproductive success [48,49]. Molybdenum (Mo) at low levels might irritate the skin and eyes while ingesting excessive amounts can result in diarrhea, infertility, liver, kidney, and lung damage [50,51]. A significant amount of copper (Cu) exposure can result in tissue damage, haemolytic anaemia, gastrointestinal bleeding, liver and kidney failure, Wilson’s disease, Alzheimer’s disease, mortality, nausea, vomiting, headaches, lethargy, diarrhea, tachycardia, and other symptoms [52,53]. Given aforementioned, it is evident that there is a critical need to enhance the characteristics of current implant materials or to develop new ones. Findings suggest that metals like zirconium (Zr), titanium (Ti), niobium (Nb), tantalum (Ta), and tin (Sn) are safe for use as metallic implants and have good biocompatibility [42]. Several scientists have been working to develop implant-grade titanium alloys that only contain the aforementioned biocompatibility constituents, such as the Ti–Zr series [54],Ti–Nb series [55], Ti–Zr–Nb series [56], Ti–Nb–Ta–Zr series [57], Ti–Zr–Nb–Sn series [58], and others. However, the exploration of new alloying systems is necessary to maximize the synergy between non-toxicity, biocompatibility and mechanical performance. Thus, understanding the criteria for selecting alloying elements is crucial for future implant materials development. The chemical composition and stoichiometry are key factors in the toxicity of Ti alloys. The chemical composition determines the chemical species released from the implant material during service. Commonly, the elements are released in an ionic state. Numerous reports on developing biomaterials for dental implants include an ion release test [59–61]. This test measures the remaining ionic concentration in a corrosive media (commonly buffer solution or artificial saliva) after immersion of the implant material at specific times and temperatures [62]. The ion release concentration can be correlated to the recommended daily intake from local or international health authorities. An example of the above is the annual Dietary Reference Intakes (DRIs) from the National Academies of Sciences, Engineering, and Medicine in the USA [63]. The DRIs include recommended dietary allowances and adequate intakes of chemical elements that can be compared to the ion release concentrations by dental implant materials. An example of the above was the Ti–34Nb–6Sn, in which ion release concentrations of Nb (1.4–2.9 μ g cm −2 day −1 ) and Sn (0.89–1.45 μ g cm −2 day −1 ) were compared to the dietary recommended intake of 620 μ g and 200 mg/kg of food, respectively [60]. The ion release concentration can also be compared to reports of TC50, i.e., the concentration of certain elements that cause 50% cell death [64]. For example, a TC50 of 2310 μ M for ionic indium (In) has been reported [65]. A development of Ti–In alloys compared the indium released to the reported TC50 61 , concluding the non-toxicity due to indium content of those alloys. From the literature, the cytotoxicity of some elements can be ordered from the most to the less cytotoxic against L-929 cells as follows: Ag + , Zn 2+ , Au 3+ , Pt 4+ , Co 2+ , Cu 2+ , Ni 2+ , Pd 2+ , Mn 2+ , Nb 5+ , Mo 5+ , Cr 3+ , In 3+ , Sn 2+65 . The use of the TC50 for toxicity comparative purposes should consider the multiple parameters influencing in-vitro cell testing, such as parent cells kind, protein concentration in the culture media, initial cell concentration, and testing parameters like time, temperature, presence of CO 2 , pH, among others [65,66]. Fig. 2 presents an overview of how alloying elements can be harmful to human health. The cytotoxicity is also relevant for orthopedic and dental implant materials. Cytotoxicity refers to the cell adhesion and proliferation occurrence on the surface of the implant material [67]. Among the key parameters that influence cytotoxicity are surface roughness, chemical composition, and hydrophobicity [67]. The plate counting method is one of the most useful for counting the cell colonies remaining at the surface of the implant material after contact for specific times and temperatures [68]. Commonly, a temperature that is close to that of the human body is selected for this test, i.e., about 37 ◦C [68]. The influencing parameters for in-vitro cell testing described from the TC50 testing also influence cell adhesion and proliferation testing [67]. Moreover, the literature contains numerous examples of newly developed or improved implant materials with cytotoxicity analyses, and these approaches are further explained in section 8.3.2 [59]. 2.2. Physicochemical properties Electronegativity, defined as an atom’s capacity to attract valence electrons, plays an important role in determining the multiplicity of chemical bonding and physicochemical properties of the material [69]. Fig. 1. Overview of alloy processing, microstructure and performance evaluation that make up the scope of this review. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3552 Thus, electronegativity strongly influences the corrosion susceptibility and ion release of Ti alloys. The electronegativity of noble, low reactive elements such as platinum (Pt), gold (Au), palladium (Pd), etc., tends to be higher than elements with high chemical activity like Mg [70–77]. Furthermore, high electronegativity differences between phases at the same alloy [70–77] or different materials in contact [78–80] might encourage high electrochemical potential disparities (micro-galvanic or macro-galvanic couples) and boost corrosion processes (more details in section 8.2.) [81]. Considering that V triggers the breakout of passive surface films in V-bearing Ti alloys, the β-Ti phase is more prone to corrosion than the α -Ti phase [70]. However, the β-phase has been reported to be more corrosion resistant than the α -Ti phase when Nb, Mo, and Zr β-stabilizers are present [77]. Lanthanum (La) also creates micro-galvanic cells in Ti–Zr–Mo alloys, reducing corrosion resistance [75]. Other examples are Cu and Ag, which add antimicrobial properties to the Ti matrix, but also galvanic pairs when found as segregates or intermetallic phases [82]. Thus, the expected better corrosion performance of β-Ti with respect to α -Ti is based on the obtaining of a fully single-β microstructure without galvanic couplings [83]. Using pre-alloyed β-Ti powders or heat treatments to promote homogeneity and/or single-phase microstructures decrease the formation of micro-galvanic couples through the material [70,76,77]. The solubility of the alloying elements on the Ti matrix should also be considered to decrease the risk of immiscible particles or intermetallic phases. The formation kinetics of the surface passivation layer also plays a key factor in the kinetics of corrosion due to micro-galvanic couplings [84]. For last, the ratio between anodic (where oxidation occurs) and cathodic (where reduction occurs) also impacts the galvanic corrosion occurrence [80]. Thus, electronegativity between phases, including elements segregation in the microstructure, should also be considered for the chemical performance of biomedical implant materials. Fig. 3 compares the four key criteria—biocompatibility, electronegativity, autoignition temperature, and cost to be considered when choosing alloying elements. Fig. 2. Overview of the negative effect of alloying element in human health. Fig. 3. Comparison of main factors to consider for alloying elements selection; biocompatibility, electronegativity, autoignition temperature, and cost (variable depending on the size and purity of the powder). A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3553 2.3. Reactivity and fabrication safety The high surface contact area of powders increases the chemical reactivity of raw materials compared to their micrometric counterparts [85,86]. Thus, the management of elemental powders for additive manufacturing (AM) and powder metallurgy techniques requires higher safety measurements than the rods or chunks usually used for melting routes. The autoignition temperature, i.e., the temperature for a 50% thermal-induced phase conversion rate, is a useful parameter to consider during the methodology planning [87]. Fig. 3 includes the autoignition temperatures of different metallic powders. It also mentions the pyrophoric powders, i.e., materials with spontaneous ignition, to assist in the fabrication measurements consideration. Moreover, it should be considered that reducing the particle size might abruptly change the autoignition temperature of the elements. For instance, it has been reported a pyrophoric risk on nano-powder of Ti, which was not observable in the micrometric order [85]. Additionally, the oxidant media also plays a key role in the ignition temperature, which tends to be lower in the air (mostly N 2 and O 2 ) compared to CO 2 or N 2 [88]. Another useful parameter to identify the expected pyrophoric activity of metallic powders is the activation energy [85,89]. The activation energy is the required energy for undergoing a chemical transformation and it is closely related to the occurrence of combustion processes and their exothermic energy releases [85,89]. Typically, higher activation energies were associated with a higher self-ignition likelihood [85]. For example, Yavor et al. reported the activation energy of some elements according to their activation energies when reacting with water [90]. The activation energies of those elements can be ordered from the highest to the lowest as follows; Mg, Al, Mo, W, Ti, Zn, Zr, Cr, Sn, B, Ni, Mn, Cu, and Fe [90]. Finally, it should be remarked that the activation energy is higher in processes implying chemical bonding changes, such as combustion processes [91]. For practical purposes, an empirical explosibility index related to ignition sensitivity and explosion severity has been reported [88]. This index is a function of the ignition temperature and energy, the critical concentration of explosive material, and the explosion pressure. Specifically, the index of explosibility classifies some of the elements as follows; i) severe hazard: Al, Mg, Zr, and Ti, ii) strong hazard: C and Fe, iii) moderate hazard: B, Cr, Mn, and Sn, iv) weak hazard: Zn, V, Mo, Co, W, and Cu, and v) none hazard: Ni. Additionally, the mentioned report considered the effect of decreasing particle size on the increment of explosibility hazard; however, only micrometric particles were considered [88]. The severity of an explosion can be assessed by a fuel’s explosivity index (Ks), which is a function of the maximum pressure during combustion [86]. Finally, the concept of the explosion pentagon should be carefully considered. The pentagon includes the necessary conditions to join for the occurrence of an explosion, i.e., fuel, oxidant, ignition source, confinement, and dispersion of the fuel. A guideline for safety measurements during the handling of explosive powders can be found in the literature [86]. 2.4. Powder cost The development of production techniques that result in final netshape parts with no need for machining, such as additive manufacturing (AM) and powder metallurgy, have considerably decreased the manufacturing cost of metallic materials [92–94], including implant materials. The net-shape devices have been very useful considering that each patient has unique anatomy that frequently requires customized biomedical implants [94]. However, the cost of raw materials greatly impacts the final price, especially when using elements with low concentrations in nature or difficult to extract and refine. The price of biomedical implants might be unaffordable for a big percentage of the population. For example, the invoiced prices of some hip, knee, dental, and stents biomedical devices was reported from 249 to 2308 USD in Mississippi, USA [94]. Additionally, statistics from the same report mentioned demand for hip, knee, dental and stent implants of at least 154, 326, 342, and 560 in Mississippi during 2010 [94]. Contrastingly, the annual median individual income in 2010 in Mississippi was about 20,000 USD. Thus, an average patient in Mississippi required from 0.1 to 1.4 salary months to acquire a biomedical implant in 2010. Despite the cost of implants already including raw materials, manufacturing process, and supply chain 92, there are other complementary costs for patients, such as medical diagnoses, clinical procedures, hospital stays, and recovery [94]. In addition, the Mississippi state reported average hospital charges of 24,751 USD. Thus, the annual average income could be insufficient to cover biomedical implant expenses in Mississippi in 2010. Furthermore, in the case of infection, a cost higher than 50,000 USD is estimated to treat infected arthroplasty [95]. Astoundingly, the demand for biomedical implants is large enough to be in the USA’s top ten medical diagnoses. Moreover, previous numbers might be more critical for patients in low-income countries. Thus, the accessibility to biomedical implants is a worldwide burden. In this context, the authors of this review suggest considering the economic burden for the end users when designing new biomedical materials. To support this recommendation, Fig. 3 compares the cost of different elemental powders to be considered during new implant materials development. 2.5. Phase-stabilization The stable room-temperature α -Ti phase corresponds to hexagonal close-packed (HCP) structure. However, the β-Ti phase (BCC structure) is more advantageous than the α -Ti phase in terms of lower elastic modulus, adequate strength, and biocompatibility [59,60]. Moreover, β-Ti alloys can be encouraged by tailoring the chemical composition of the alloys. The addition of β-stabilizers decreases the β-transus temperature, expands the β region in the phase diagram, and increases the β-phase stability. Thus, metastable β-microstructures will be formed if a critical concentration of β-stabilizers is met. The β and α stabilizer equivalencies, also called molybdenum (Mo eq ) and aluminium (Al eq ) equivalencies, allow estimating the content of β (Moeq wt.%)and α (Aleq wt.%) phases in the microstructure. Some authors define the Mo eq as the amount of different β-, α -, and neutral elements necessary to retain a fully β-phase quenched microstructure [96]. The Mo eq and Al eq can be calculated from Eqs. (1) and (2) [81]. Moeq wt.%= [Mo] + [Ta] 5+[Nb] 3.6+[W] 2.5+[V] 1.5+1.25 [Cr] + 1.25 [Ni] +1.7[Mn] + 1.7[Co] + 2.5[Fe] − [Al](1) Aleq wt.%= [Al] + [Zr] 5.9+[Sn] 3+[O+N] 0.1(2) where [x] refers to the alloying elements content in wt.%. Mo eq and Al eq consider the effect of the main β- and α stabilizers. Thus, β-transus temperature decreases with the increment of Mo eq . Different corrected coefficients have been introduced in Eqs. (1) and (2) [81,97]. Table 1 contains the MO eq calculations for various β-Ti alloys. The metastable β-phase can be transformed under the application of an energetic stimulus, whether thermal or plastic stress [108]. These phenomena are also known as stress-induced and strain-induced phase transformations, respectively. The first one is given when no plastic deformation is applied [108], e.g., during heat treatments. On the other hand, the strain-induced phase transformation occurs during the plastic strain regime [109]. The stress-induced transformation is the most studied in biomedical materials because they are subjected to lower loads than structural materials, which are more susceptible to strain-induced phase transformation. Moreover, the stress-induced phase transformation is affected by the alloying elements [110]. Thus, the β-Ti alloys can be obtained from heat treatments designed to retain the β phase (with cooling out of equilibrium) or by A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3554 tailoring the chemical composition [96]. Transformations from the metastable β phase might occur to the HCP Tiα ′ , orthorhombic Tiα ’’, or HCP Tiω phases during straining (tensile, compressive, or creep) or heat treatments. The formation of each phase is a function of the β stability, as is predicted by Mo eq . The β-stabilizers dependence for the β, Tiα , Tiα ′ , and Tiα ’’phase transformations in Ti alloys. One example of stress-induced phase transformation is the β-type Ti–Hf alloys, which transformed to mainly HCP- α ’ martensite after water quenching from solution treatment [111]. For more details, Kolli et al. reviewed the phase transformation (including the Tiω ) as a function of the Mo eq for several Ti alloys [96,112]. The critical concentrations of β-stabilizers, mainly Mo, Nb, Ta, V, W, Co, Cu, Cr, Fe, Mn, and Ni, to retain a fully β-phase microstructure in binary Ti-alloys can also be found in the literature [96]. Thus, chemical composition, temperature, and cooling rate have a significant impact on the phase transformation of Ti-alloys once a minimum critical strain is applied [113]. 2.6. Effect of Fe and Cr contents Segregation defects, particularly β-flecks, are critical phenomena impacting the microstructural integrity and mechanical properties of β-Ti alloys, especially those with higher contents of transition elements such as Fe and Cr. In addition, the heterogeneous distribution of these β-stabilizers during solidification can lead to the formation of localized, enriched regions known as β-flecks [114]. In alloys with higher Fe and Cr contents, the solidification range is extended, increasing the susceptibility to microsegregation. As the alloy solidifies, the solute elements (Fe and Cr) tend to segregate towards the interdendritic regions, which solidify later. Therefore, this segregation leads to the enrichment of β-stabilizing elements in these regions, promoting the formation of localized β-phase areas known as β-flecks. These flecks are typically harder and more brittle than the surrounding matrix, presenting as distinct, often darker, spots in the microstructure due to their different chemical composition and phase structure. The mechanism of β-fleck formation is influenced by several factors, including the cooling rate, alloy composition, and the presence of other alloying elements [114–116]. A slower cooling rate tends to exacerbate segregation by allowing more time for the β-stabilizing elements to migrate towards the interdendritic regions. Moreover, the overall alloy composition also plays a significant role; higher concentrations of Fe and Cr not only promote the stability of the β-phase but also increase the tendency for segregation. Furthermore, the interaction of Fe and Cr with other alloying elements can influence the kinetics of segregation and the resulting microstructure. To mitigate the formation of β-flecks, several strategies can be employed. These include optimizing the alloy composition to minimize the content of segregating elements, refining the microstructure through thermomechanical processing to disrupt the segregated regions, and employing rapid solidification techniques to reduce the time available for segregation. For example, additive manufacturing has been used to Table 1 Mo equivalency (MO eq ) for several β- or near β-Ti alloys. Alloy Fabrication technique MO EQ Reference Ti-13 Nb–13Zr Cast +Aging 3.6 [98] Ti–15Mo Cast +Annealing 15.0 [98] Ti–10Zr–5Ta–5Nb –2.4 [20] Ti–12Cr –15.0 [83] Ti–12Mo–5Ta –13.0 [20] Ti–12Mo–5Zr –12.0 [20] Ti–12Mo–6Zr–2Fe Aast +Annealing 17.0 [98] Ti–12Mo–6Zr–2Fe LPBF, chessscanned 17.0 [99] Ti–12Mo–6Zr–2Fe LPBF, simple-scanned 17.0 [99] Ti–12Nb–12Zr–3Ta SLM +SSol 3.9 [100] Ti–12Nb–12Zr–8Ta SLM +SSol 4.9 [100] Ti–12Nb–13Zr–4Ta SLM +SSol 4.1 [100] Ti–12Nb–13Zr–6Ta SLM +SSol 4.5 [100] Ti–13Nb–12Mo P&S +Laser surface modification with laser power of 1000 W 15.6 [101] Ti–13Nb–13Zr –3.6 [20] Ti–13Nb–13Zr SLM +SSol 3.6 [100] Ti–15Mo Press & Sintering −3.3 [102] Ti–15Mo Press & Sintering 15.0 [102] Ti–15Mo Press & Sintering 15.0 [102] Ti–15Mo Press & Sintering 15.0 [102] Ti–15Mo-0.1O Melting +Hot Forging + Hot Rolling 15.0 [103] Ti–15Mo-0.2O Melting +Hot Forging + Hot Rolling 15.0 [103] Ti–15Mo-0.3O Melting +Hot Forging + Hot Rolling 15.0 [103] Ti–15Mo-0.4O Melting +Hot Forging + Hot Rolling 15.0 [103] Ti–15Mo-0.5O Melting +Hot Forging + Hot Rolling 15.0 [103] Ti–15Mo–2.8Nb–0.2Si–0.26O (Annealed) 15.8 [83] Ti–15Mo-2.8Nb–3Al –12.8 [20] Ti–15Mo–3Nb Cast +Annealing 15.8 [98] Ti–15Mo–5Zr–3Al Press & Sintering 12.0 [104] Ti–15Mo–5Zr–3Al Press & Sintering 12.0 [104] Ti–15Mo–5Zr–3Al Press & Sintering 12.0 [104] Ti–15Mo–5Zr–3Al Press & Sintering 12.0 [104] Ti–15Mo–5Zr–3Al Press & Sintering 12.0 [104] Ti–15Mo–5Zr–3Al Cast 12.0 [105] Ti–16Nb-10Hf Cast +Aging 4.4 [98] Ti–19Nb–14Zr –5.3 [20] Ti–20Nb–10Mo P&S +Laser surface modification with laser power of 1000 W 15.6 [101] Ti–23Nb–0.7Ta–2Zr –6.5 [83] Ti–23Nb–0.7Ta–2Zr–1.2O –6.5 [83] Ti–24Nb–0.5N –6.7 [83] Ti–24Nb–0.5O –6.7 [83] Ti–24Nb–4Zr–7.9Sn –6.7 [20] Ti–24Nb–4Zr–8Sn Hot rolling 6.7 [106] Ti–24Nb–4Zr–8Sn Hot forging 6.7 [106] Ti–25Nb–2Mo–4Sn –8.9 [20] Ti–25Ta–25Nb –11.9 [20] Ti–27Nb–8Mo P&S +Laser surface modification with laser power of 1000 W 15.5 [101] Ti–28Nb–13Zr–0.5Fe –9.0 [20] Ti–28Nb–13Zr–2Fe –12.8 [20] Ti–29Nb–11Ta–5Zr –10.3 [20] Ti–29Nb–13Ta–2Sn –10.7 [20] Ti–29Nb–13Ta-4.5Zr –10.7 [20] Ti–29Nb–13Ta-4.6Sn –10.7 [20] Ti–29Nb–13Ta–4.6Zr (Aged) 10.7 [83] Ti–29Nb–13Ta–4Mo –14.7 [20] Ti–29Nb–13Ta–6Sn –10.7 [20] Ti–29Nb–13Ta-7.1Zr –10.7 [20] Ti–29Nb–6Ta–5Zr –9.3 [20] Ti–34Nb–6Sn P&S 100 MPa 9.4 [60] Ti–34Nb–6Sn P&S 200 MPa 9.4 [60] Ti–35.5Nb–7.3Zr–5.7Ta –11.0 [83] Ti–35Nb–4Sn –9.7 [20] Ti–35Nb–5.7Ta–7.2Zr –10.9 [20] Table 1 (continued) Alloy Fabrication technique MO EQ Reference Ti–35Nb–5ta–7Zr–0.4O –10.7 [20] Ti–35Nb–6Mo P&S +Laser surface modification with laser power of 1000 W 15.7 [101] Ti–35Nb–7Zr–5Ta –10.7 [20] TI–36Nb–2Ta–3Zr–0.3O –10.4 [83] TI-40Nb P&S 1-h sintering at 1200 ◦C 11.1 [107] TI–6Mn–4Mo –14.2 [83] Ti-7.5Mo–3Fe –15.0 [20] Ti–9Mn –15.3 [83] - Unknown thermomechanical route. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3555 reduceβ-flecks formation [117,118]. Additionally, post-solidification heat treatments can be used to homogenize the microstructure, dissolving β-flecks and redistributing the β-stabilizing elements more uniformly. 2.7. Compositional alloy design The design β-Ti alloys involves sophisticated strategies to achieve a desirable balance between high strength, ductility, and other mechanical properties, which are essential for biomedical as well as other applications like aerospace and high-performance engineering fields. Among these strategies, two prominent approaches in the compositional design of β-Ti alloys are the d-electron theory and the Bo-Md diagram. Each of these methods offers a framework for understanding and predicting the phase stability and mechanical properties of these alloys but also comes with certain limitations. 2.7.1. D-electron concept The d-electron concept, developed by Morinaga is based on the electronic structure theory [119]. It focuses on the role of d-electrons in determining the phase stability of titanium alloys. In this concept, the number of d-electrons (e/a ratio, where e represents the number of valence electrons and a represents the number of atoms) is a critical parameter. The theory posits that the e/a ratio influences the phase stability of the alloys, where a higher e/a ratio tends to stabilize the β phase, while a lower ratio favors the α phase. This approach allows alloy designers to predict the phase stability and, to some extent, the mechanical properties of β-Ti alloys by adjusting the alloying elements to achieve the desired e/a ratio. 2.7.2. Bo-Md diagram The Bo-Md diagram, named after the Bonding (Bo) and Metal dorbital energy level (Md) parameters, is another approach used in the design of β-Ti alloys. It offers a visual tool for understanding the effects of alloying elements on the stability of β-Ti alloys. This diagram plots the bond order (Bo), indicative of the strength of the metal-metal bonds, against the d-orbital energy level (Md), which represents the energy level of the d-orbitals of the metal atoms. The positioning of an alloy on this diagram can predict its phase stability, with specific regions corresponding to the stability of α , β, and other phases. By targeting a position within the β-stable region, alloy designers can tailor the composition of β-Ti alloys to achieve desired properties. For example, the Bo-Md diagram has been used to correlate the instability of the β-phase with lower Young’s modulus [120]. A detailed discussion of the d-electron theory and the use of Bo-Md diagrams can be found in the literature [121,122]. 2.7.3. Limitations The d-electron concept, while useful in predicting phase stability, does not account for all factors influencing the properties of alloys, such as microstructural effects and the influence of interstitial elements like oxygen and nitrogen. Moreover, it primarily focuses on electronic factors and overlooks other crucial aspects such as atomic size and lattice distortion. On the other hand, the Bo-Md diagram, on the other hand, provides a more visual and intuitive way of understanding alloy behavior. However, it simplifies complex alloy systems to two-dimensional representations, potentially overlooking the multidimensional nature of alloy properties. Additionally, this diagram also does not directly account for kinetic factors and the effects of processing conditions on the final alloy properties. In conclusion, while the d-electron concept and Bo-Md diagram offer valuable insights into the design of β-Ti alloys, their effectiveness is constrained by their inherent assumptions and the complexity of alloy systems. Future research and development in this field will likely integrate these theoretical frameworks with empirical data, advanced computational models, and considerations of processing-structure- property relationships, sustainability and health impact to develop the next generation of β-Ti alloys. 3. Powder production methods Powder metals are essential raw materials for neat-net-shape processing technologies such as powder metallurgy (PM), metal injection molding (MIM), hot isostatic pressing (HIP) and additive manufacturing (AM) [123]. In terms of how the powder is produced, chemical, physical, and mechanical methods can be distinguished [124], as is shown in Fig. 4. Most industrial processes combine the above types. Choosing the route, however, depends on the desired properties of the powder, along with its application-specific characteristics. Powders of Ti and its alloys have been processed commercially by chemical reduction, hydrogenation/dehydrogenation, gas atomization, plasma-rotating electrode, and plasma atomization [125]. In addition, low quantities of titanium powder can also be obtained by hydrolysis of melted titanium compounds (TiCl 4 , TiO 2 ), centrifugal atomization, and others. Here, we provide an overview of the commercial methods used to process titanium powder. 3.1. Chemical methods In general, typically, chemical methods can be categorized into two groups based on the processes involved: chemical reductions and chemical decompositions. In chemical reduction, a variety of chemical compounds are employed, including oxides, halides, or various metal salts. The following are some examples of such processes: 1) reducing tungsten oxide with hydrogen from a solid state; 2) reducing titanium tetrachloride with melted magnesium, the Kroll process [126]; or 3) reducing copper sulfate and iron from an aqueous solution. CP-Ti can be produced by chlorinating natural or synthetic rutile (TiO 2 ) [125,127]. A series of reduction processes are used to reduce titanium tetrachloride (TiCl4) into metallic titanium, including the Hunter process (reduction of sodium), Kroll process (reduction of magnesium), and the reduction of calcium hydride, which has been used in Russia for years [128]. Fig. 5 illustrates the chemical methods used to produce CP-Ti. 3.1.1. Hunter process Hunter process could be performed in a one-step or two-step procedure. A stoichiometric amount of sodium metal is employed in the one-step procedure to convert TiCl 4 directly to Ti metal in a single reaction container. In one-step procedure, TiCl 4 is directly converted to Ti metal in a single reaction container by the use of a stoichiometric amount of sodium metal (Eq. 3). A two-stage reaction, on the other hand, begins with TiCl4 being reduced to TiCl2 using half stoichiometric sodium, followed by the transfer of TiCl2 into another vessel to undergo a second reaction with sodium metal to obtain Ti metal (Eq. 4 and 5). The two-step operation offers the advantages of separating the sodium reduction process into two parts to reduce the amount of heat generated as well as to control the vapor pressure of sodium metal more effectively than a one-step process. Although, the Hunter process is attractive because the purity of the Ti produced, compared to the Kroll process. Even so, the Hunter method for commercial applications has been abandoned since 1993 due to the difficulty of effectively separating NaCl, which requires leaching in an aqueous solution and requires additional energy [125]. One-step operation: TiCl4(g) + 4Na(l)→Ti(s) + 4NaCl (l)(3) Two-step operation: TiCl4(g) + 2Na(l)→TiCl2(l,in NaCl) + 2NaCl (l)(4) TiCl2(l,in NaCl) + 2NaCl (l)→Ti(s) + 4NaCl(l)(5) A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3556 3.1.2. Kroll process The Magnesium reduction method or Kroll process is one most used via to obtain Ti from TiCl 4 . Due to the extensive heat treatment required to isolate pure Ti from Ti sponge obtained, this process is extremely energy intensive [129]. Kroll process occurs at 1073–1123K and its similar to the Hunter process but, magnesium is used instead NaCl to reduce TiCl 4 (Eq. 5) [125]. It is important to note that the reduction of Ti produces an exothermic reaction, releasing 412 kJ/mol, or even 686 kJ/mol if magnesium is added to the reaction in a gaseous state [129, 130]. Kroll process involves four stages [131], represented in Fig. 5: (1) chlorination, where TiO 2 (s) is chlorinated with carbon, obtaining TiCl 4 (Eq. 6); (2) reduction and separation, which involves the reduction of TiCl 4 with Mg metal to form Ti sponge (Eq. 7); (3) crushing and melting, where Ti the sponge is crushed and melted in vacuum arc or electron beam melting furnaces in order to reduce the impurities; and (4) the electrolysis process, magnesium chloride (MgCl 2 ) is electrolyzed to obtain Mg and Cl 2 gas (Eq. 8) [131]. TiO2(s) + 2Cl2(g) + C(s)→TiCl4(g) + CO2(g)(6) TiCl4(g) + 2Mg(l)→Ti(s) + 2MgCl2(l) (Reduction)(7) 2MgCl2(l)→Mg +Cl2(g) (electrolysis)(8) As a disadvantage of the Kroll process, the time required for processing is lengthy. In step 2, the sponge’s impurities are eliminated through vacuum distillation over 4–5 days, depending on the size of the sponge block. The high crush resistance of the sponge difficult its size reduction. In this case, prior to the crushing, Ti sponge is oxidized or hydrogenated to make it fragile. When the reduction and separation process end, the reactor is cooled down for 4 days, and opened. The sponge obtained is pressed, and then the reaction products are eliminated from its surface. Since, the quality of the obtained Ti sponge is not homogeneous (variable grades of Ti within sponge), thus, this is grinded and classified and the high-purity grade Ti is analyzed and post processed by vacuum arc re-melting to obtain CP-Ti [129]. More effective technology process is necessary to reduce cost of Ti extraction, to improve the cycle time of the process and to increase the quality of Ti obtained [130,132,133]. 3.1.3. Armstrong process The technique is a derivative of the Hunter process and produces Ti powder with oxygen levels below 1000 ppm and chlorine levels below 50 ppm. Using the reduction of TiCl 4 and other metal halides by sodium, this method produces CP-Ti and Ti–6Al–4V with a unique "coral-like" morphology [134,135] (Fig. 5). Enhanced particle size distribution and tap density are achieved by employing post-processing methods like ball milling. Due to their irregular shapes, porous surfaces, and ease of compression, the resultant powder is suitable for conventional powder metallurgy, producing compacts with high green resistance, even when cold uniaxial compaction is used [135]. 3.1.4. Hydride-dehydride process (HDH) In this process, Ti (powder, sponge, milled products, or scrap) is introduced into a batch furnace at pressures of 1 atm. and temperatures up to 800 ◦C, in a vacuum/hydrogen atmosphere. Then, the Ti sponge is milled, and the resultant powder is sieved to produce a fine powder from titanium hydride. After that, the classified powders are re-introduced into the batch furnace under a vacuum or argon atmosphere (dehydride process) in order to remove the hydrogen. After being magnetically screened, the powders are then acid washed to remove ferromagnetic impurities. Pure powders with irregular morphologies are produced. In the case of scrap being used as raw material, this process cannot be used for modifying the alloy composition [135]. We have already discussed the basic methods for obtaining CP-Ti Fig. 4. Manufacturing techniques for Tipowder and its alloys. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3557 metal powders via chemical methods. Some companies also use cuttingedge technology, such as atomization, to produce better quality powders. However, powders of Ti alloys are commonly obtained by blending and mechanical alloying [136]. During mechanical alloying, constituent powders are ground in a vessel containing grinding media [137]. The diffusion process over solid-state media produces homogeneous metal powders [137]. Next, we give an overview of the most common fabrication techniques for Ti alloys, which can be categorized into physical and mechanical methods. 3.2. Physical methods The most relevant physical methods to obtain Ti powders are the electrolytic and atomization methods. Fig. 6 shows an illustration of physical methods used to obtain Ti powders. The shape is more spherical in atomization process; in particular, plasma atomization and ultrasonic techniques produce more spherical particles. 3.2.1. Electrolytic method Electrolysis of Ti is an ionizable process for producing high purity Ti. It involves the preparation of an electrolyte, the electro-refining of the Ti, and the collection of the Ti alloy. The electrolyte is prepared by reacting TiCl 4 and high-purity Ti to obtain molten salts containing lowvalent Ti ions (Ti 2+ and Ti 3+ ) [138]. Thus, the salt contains Ti ions in solution, which are necessary to facilitate the hydrolysis process at a voltage below that required to decompose alkali and alkaline earth metals at the cathode [139]. During electro-refining, Ti sponge obtained by Kroll process acts as an anode, while high purity Ti plates serve as cathodes on which Ti ions are deposited. As an electrolyte, alkali and alkaline earth halides are utilized because of their eutectic composition, which can melt at temperatures as low as 350 ◦C. The final product collected is Ti of high purity. Table 2 shows possible reactions involved Fig. 5. Techniques for the production of Ti powders and its alloys by chemical methods. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3558 4. Manufacturing of porous Ti alloys Porous metal implants have been manufactured using different routes. On the one hand, the conventional techniques include conventional powder metallurgy [203], spark plasma sintering [204], metal injection molding (MIM) [205], pressureless sintering [206], low pressure sintering [60], the use of space holders [207], expansion of pressurized bubbles (injection of gas into the metal melt) [208], freeze casting techniques [209]. On the other hand, more sophisticated technologies referred as additive manufacturing (AM) technologies [210], which include three-dimensional printing (3DP) [211], electron beam melting (EBM) [212], selective laser melting (SLM) [213], direct metal deposit (DMD) [214], and selective laser sintering (SLS) [213], Fig. 12. 4.1. Powder metallurgy 4.1.1. Press and sinter Powder metallurgy has a variety of well-known advantages when it comes to the manufacturing of β-Ti alloys. The powder metallurgic approach offers a substantial financial advantage over the conventional route [203]. Due to the efficiency in manufacturing final shapes that it offers over other technologies, as well as the capability to blend a variety of alloys using both elemental powders mixing and pre-added powder mixing. The simplest process is press-and-sinter method, which allows pieces to be manufactured from an elementary mixture of the constituent particles of a certain alloy. Since Kroll first began sintering spiky Ti with low argon pressure in 1937, the conventional press-and-sinter method for Ti alloys has been evolving rapidly, along with new processes that have been incorporated into the processing of various materials, including the production of powders (Kroll, Amstrong, Hydrogen-assisted magnesium reduction HAMR, and electrochemical reduction) as in the densification of the same (Hot pressing HP, Hot isostatic pressing HIP, Field Assisted Sintering Technologies/Spark Plasma Sintering (FAST/SPS) or Electrical Resistance sintering ERS) [160]. Use of TiH 2 as a primary material that can dehydrate in a phase prior to the sintering cycle is one option for obtaining alloys while minimizing the impact of oxygen on the Ti. In this regard, Ti–Nb and Ti–Nb–Fe β-Ti alloys were developed from TiH 2 powders by Chirico et al. [215]. Furthermore, they studied the role that alloying elements play in the transformation of TiH 2 into β-Ti alloys [216]. Several alloys combining α and β phases of Ti have been developed in recent years for various industrial sectors. However, different methodologies have been used for powder technology. The majority of the work has been done on the CP-Ti or the Ti–6Al–4V alloy, both in the elemental mixture and the pre-alloying [217]. To address these last issues, it is necessary or preferable to turn to high-density processing (HIP or VHP: vacuum hot pressing). This final process was developed in 1980 and has primarily only been used in CP-Ti and Ti–6Al–4V, however it has been used to sinter both blending elements (BE) and prealloying (PA) powders [218]. Regarding β-Ti alloys, several alloys have been designed adjusting their molybdenum equivalent for various industries, such as aerospace, automotive, athletics, chemistry, or biomedicine [219]. One of the primary concerns with these alloys is their large concentration of alloying elements, which may make them unstable in the future, because β stability begins at a Mo eq of 25, with β-Ti alloys metastable for Mo eq ranging from 10 to 20, [220]. As a response, Bolzoni et al., developed low-cost alloys, including Fe as a alloying element as it is one of the elements with the greatest effect on β-phase stability [221]. Furthermore, the use of refractory metals as additives, such as molybdenum, niobium, or tantalium, causes diffusion problems in a solid state, allowing for a significant amount of alpha phase on the grain boundary that modifies not only the microstructure, but also affects powder metallurgy alloys in terms of their mechanical properties and corrosion resistance [222]. These mechanical characteristics are also influenced by the porosity obtained during the conventional press-and-sinter process [223,224] which mostly affects fatigue properties. However, one of the primary concerns is the high reactivity of Ti, which requires working in controlled atmospheres or in vacuum, yet despite this, certain oxygen contents (about 0.8–1.0 w%) are added, which reduces the plasticity of the alloys while increasing strength [225]. Luo et al. investigated the impact of high oxygen content on alloys produced by powder metallurgy technique such as α -Ti, α +β Ti, and finally β-Ti in the presence of a critical oxygen content that maintains a 6% elongation [226]. As a result, solubilization routes and thermal treatment methods play a significant role on final mechanical properties, particularly plasticity [227]. With an oxygen content of 0.34%, an alloy Ti–31Ta–12Nb–2V–8Zr obtained by cold isostatic pressing (CIP), vacuum sintering (VS), hot forging (HF), hot rolling (HR), solution treatment (ST), and cold work (CW) maintains a plasticity of 13% with a UTS of 1170 MPa, achieving a porosity close to 0%. In a similar processing conditions, Furuta et al. evaluate the influence of Niobium content in the alloys Ti-(30–36)Nb–2Ta–3Zr, with oxygen content of around 0.49%, resulting in tensile properties decreasing and elongation increasing [228]. The mechanism in which oxygen is presented also has a significant impact, as shown by M´ alek et al. that studied the impact of oxygen being in solution, homogeneous, or concentrated at the granular boundary [229]. This oxygen content is proportional to the mean particle size of the powder, increasing as the size decreases, and can range from 0.1% for particle sizes of 100–120 μ m to 0.6% for particle sizes less than 25 μ m for irregular HDH powder of Ti–6Al–4V [230] whereas the oxygen concentration of powder atomized by the same gas remained between 0.07 and 0.10% [199]. Accordingly, mechanical properties are determined by the parameters of the powder metallurgy press-and-sinter process, the most influential of which are temperature and sintering time. Many authors have observed that increasing the sinter temperature increases the properties slightly in all instances, since increasing the size of the grain decreases the porosity and changes the morphology of the pores [231]. However, there is a significant concern, when large amounts of alloying elements are added, microstructures of β phase with a large amount of α phase in the grain border that extends into the grain interior are formed. This could affect mechanical properties and weaken the grains in the same way that precipitated ω phase is observed in transmission microscopy. As a result, some researchers have returned to mechanical alloying, which allows for more chemical homogeneity and microstructural change in alloys. Amig´ o Mata et al., demonstrated that the effects of adding low Fe and Cr contents in a Ti–15Nb base alloy allow for close to 80% phase in its microstructure for 3% contents that provide a Mo eq of close to 12%, however, the mechanical properties of flexion of 700 MPa for these conditions are significantly reduced when P&S of elemental powder blends is performed compared to the 220 MPa obtained with mechanical alloying powders [232]. Fig. 13 shows the microstructure of Ti–Nb alloys obtained by P/M press & sinter at 1250 ◦C. They obtain two alloys with similar Mo eq , Ti–40Nb and Ti–25Nb–5Fe, through P/M. Mohan et al. used the same conditions to study the impact of Fe addition in Ti–Mo alloys with Mo contents of 12 and 15%, with similar outcomes [233]. Elshalakany et al. developed quaternary alloys Ti–15Mo-6-Zr-(1–4) Cr, which were obtained by P&S of blend element powders and sintered at 1250 ◦C after a 3-h dwell time, and obtained compressive properties and bending strength between 1200 and 1600 MPa, with elongations of less than 5%. Plasticity is reduced primarily by the porosity obtained around 4–5% and the stabilization of the β phase up to 98.9% when Cr is added to the base alloy Ti–15Mo–6Zr [234]. It is nevertheless feasible to obtain a chemically homogeneous alloy with 15% Mo, which does not occur when the addition is more, as in Ti–Nb alloys, as concluded by Caha et al. [235]. The first exhibits a lack of diffusion of larger Nb particles, even when the sintering cycle is extended to 4 h, whereas the alloying with Fe exhibit a high porosity due to the different velocities of diffusion of Fe to Ti, which results in porosity similar to that of the Fe particles, and which appears to decrease when the cycle is extended to 4 h, as the porosity decreases to 1% [235]. These alloys, which were tested for their A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3565 resistance to corrosion and tribocorrosion, demonstrated excellent resistance, allowing them to be used as biomedical alloys. Mohan et al. investigated the same effect of Fe addition on alloy porosity [236]. Despite a Fe addition of up to 4 w% in the Ti–Mo–Zr–Fe alloys with this addition, the β phase is completely stabilized; however, the bending strength increases when the temperature rises from 1250 to 1300 ◦C for additions of 1 and 2% Fe, whereas for higher proportions, the properties decrease [236]. They sinter elemental blend powders and mechanical alloying powders, obtaining lower mechanical properties ranging from 1000 MPa to 400–600 MPa, because the ultimate porosity in mechanical alloys is higher and the pores have a larger size. The microstructure obtained in these mechanical alloys is bimodal in all cases, with better distribution and lower porosity as the compacting pressure increases from 600 MPa to 900 MPa. In these conditions, plasticity is low, which is one of the limitations of using this technology to create products that may be used in the biomedical field, which requires an elongation minimum of 8%. However, there are other properties that should be considered for its application in this field, as well as mechanical resistance and corrosion resistance, as analyzed by Çaha et al., particularly when TiN coatings or TiN particle additions are performed in the alloys, as performed by Çaha et al. and Chirico et al. with the evaluation of the tribocorrosion behavior of the Ti–40Nb P/M [237,238]. This superficial modification is used not only to increase resistance to degradation, but also to significantly modify the surface properties that allow for greater interaction with organic tissues. To that purpose, various laser modification techniques have been used, yielding a reduction in elastic modulus in some cases due to microstructural transformations caused by rapid heating and cooling. Tendero et al. investigated the microstructural changes observed in P/M Ti–Nb–Mo alloys, observing a slight increase in bending strength but, more importantly, an increase in overall length without reaching the previously mentioned limits [101]. Rossi et al. studied the impact of different alloying elements on the mechanical characteristics of the Ti–35Nb–10Ta and Ti–30Nb–4Sn alloys, obtaining, according to process parameters, a slight increase in hardness in both alloys and a slight decrease in elastic modulus in the Ti–30Nb–4Sn alloy [239]. On this family of Ti–Nb–Sn alloys, biological studies by Batibay et al. (2023) revealed antibacterial properties against Staphylococcus aureus (S. aureus) and were not cytotoxic to healthy cells, irrespective of their compositions [240]. Despite improvements in mechanical behavior of P/M alloys following laser treatment, particularly in material plasticity, the decrease of fatigue properties with respect to forging products remains an issue. On the other hand, Yang et al. (2024) analyze the corrosion behavior and antibacterial properties of Ti–15Mo-xCu alloys with copper contents from 3 to 15%, using low-pressure techniques to obtain functionally graded porous Ti–Mo–Cu alloys. By forming the Ti 2 Cu phase mainly at the grain boundary, they obtain a micro-galvanic cell effect which slightly decreases their corrosion resistance compared to Ti–15Mo alloy. However, the antibacterial properties increased with Cu content [241]. As a result, for greater β-Ti alloys’s consolidation and elemental solubility, P&S Hot Isostatic Pressing (HIP) processes are introduced along with P&S processes, as demonstrated by Lario et al. (2021), who assessed the evolution of mechanical characteristics of Ti–35Nb–2Sn after HIP (1200 ◦C 2 h) after 3 h of sintering at 1350 ◦C and with temperature controls of 100 and 500 ◦C/min. Post-treatments can reduce porosity while increasing grain size from 19 to 57 μ m. As a result, there is no lasting effect, and there is even a slight decrease in length and an increase in elastic modulus, especially when the material’s cooling is slow [242]. However, this alloy results in a 10% increase in bending strength in the fast cycle and a 30% increase in bending strength in the slow cycle. In both cases, there was an increase in length that reached 4.2% and a decrease in elastic modulus from 79.8 GPa as-sintered to 76.8 GPa as-HIPed [243]. These HIP treatments maintain the excellent corrosion resistance demonstrated by the alloy Ti–35Nb–7Zr–5Ta and extend the passivation potential from 0.1V to 3V for the rapid cooling condition that exhibits the best mechanical and plastic properties. Other authors, however, refer to high-density processes such as HP, HIP, FAST/SPS, or ERS, which will be discussed further. However, these processes require the use of pre-alloyed atomized powders by various processes or mechanical alloying, as discussed in subsequent chapters. Guo et al., used direct hot pressing to obtain nearβ-Ti 5553 alloys from machining swarf. Thermal aging treatments cause the formation of a content in phase that reaches values of 36.5% with acicular and globular morphology [244]. When this occurred, the UTS for HT-600 ◦ C-4h + 700 ◦ C-0.5h is 1400 MPa with a 2.5% deformation and 2500 MPa in three-point bending strength for a 1.6 mm deflection. Properties that are comparable to, and even greater to, those obtained from hot forging + solution treatment, but with a shorter elongation. Other authors use two-step pressureless sintering consisting of cold isostatic press of the powder mixture for sintering, then sintering at different temperatures. Sun et al. (2024) use a 250 MPa cold isostatic press with pressureless sintering at 1250 ◦C of Ti–6Al–4V, obtaining a densification of 99% and a grain size around 90 μ m. These two-step sintered samples show a maximum tensile strength of 968 MPa, yield strength of 871 MPa and elongation of 16.1%, thus maintaining ductility [245]. Kuang et al. (2023) obtain highly densified Ti–6Al–4V alloy by pressureless sintering and hot extrusion, after cold isostatic press. After extrusion they obtain an ultimate tensile strength of the rod of 1213 MPa and the elongation increases up to 14 % [246]. Chen et al. (2023) also perform two-step pressureless sintering on TiC-reinforced Ti–6Al–4V, by a first step of mechanical alloying of Ti CP and 60Al40V master alloy. This results in a composite with ultrastrong hierarchical microstructure [247]. Furthermore, Guo et al. (2024) use a hot forging post-processing after two-step pressureless sintering of Ti–6Cr–5Mo–5V–4Al for high performance, using as a first step CIP with a pressure of 200 MPa. After forging and solution treatment and ageing, the sample exhibited a high maximum tensile strength of 1211 MPa and an elongation of 13.8 %, considering very interesting properties for this PM Ti6554 alloy [248]. Similarly, Franco et al. Using a similar two-step pressureless sintering Fig. 13. Microstructure of Ti–Nb alloys obtained by P/M press & sinter at 1250 ◦C: a) Ti–15Nb, b) Ti–35Nb [233]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3566 process, they obtain the Ti–35Nb–7Zr–5Ta alloy with a final density of 98%, although the oxygen content increases to 0.96%. This high oxygen content allows obtaining a fatigue resistance higher than 88 MPa after subsequent heat treatment, together with the formation of coarse α -phase precipitates at the grain boundaries [249]. In contrast, Cai et al. (2023) obtained Ti–20Al–19Nb (at.%) dense alloy by TiH 2 -assisted pressureless sintering; although its densification remains problematic. They found that TiH 2 participated in the reverse eutectoid transformation, α -Ti +δ-TiH 2 → β-Ti, improving the density and mechanical properties [250]. 4.1.2. Metal injection molding These mechanical characteristics are also found in P&S products, but to a lesser extent when techniques such as metal injection molding are used (MIM). Specifically, the technical basis and properties, as well as the many steps to obtain MIM pieces, are described in various German Works [205]. For instance, in terms of β-Ti alloy sintering, Kafkas and Ebel sintered Ti–24Nb–4Zr–8Sn dog-bone specimens at various temperatures and times, yielding elastic moduli ranging from 53 to 61 GPa and yield strengths of 650 MPa for 0.5% oxygen contents [251]. Similarly, zhao et al. investigated the sintering behavior and mechanical characteristics of a metal injection molded Ti–16Nb alloy, requiring a temperature of 1500 ◦C to minimize porosity to 5%, whereas Nb diffusion is nearly complete at 1300 ◦C. However, due to the use of an organic binder, TiC precipitates appeared. Furthermore, this same effect of TiC precipitation is described by Xu et al., who observed the formation of particles in the alloy Ti–10Zr–20Nb with O content less than 0.28% and carbon content less than 0.04% [252]. Moreover, despite having a porosity of 5–6% and grain sizes greater than 300 μ m, analyzing the mechanical properties according to the distribution of these particles, since they decrease the elongation although it is still greater than 4%, maintaining the elastic limit between 700 and 750 MPa. Eventually, the same authors (Xu et al.) investigated the impact of yttrium (0.3–1 w %) addition on the same Ti–10Zr–20Nb alloy by MIM processing. The alloys are sintered at 1500 ◦C for 4 h. Despite the control over grain size, which drops to 170 μ m for the addition of 1% Y, mechanical properties decrease to values of yield strength of 612 MPa, resulting in a decrease in ductility of 4.0%–3.3%, as does the elastic modulus, which ranges from 71 GPa to 61 GPa. Santos et al., also obtained a low-cost β -type Ti-(8–17 w%) Mn using MIM, reporting phase stabilization and observing phase formation in TEM. However, the obtained plastic properties are inferior to the Ti–6Al–4V alloy, whereas the Ti–Mn hipo-eutectoids exhibit greater resistance to the Ti–6Al–4V alloy [253]. Furthermore, these same authors produced alloys with low Mn (5–6 w%) content with the addition of Mo (3-4w%) contents manufactured by MIM and sintered at 1100 ◦C for 8 h. Notably, it continues to work around a Mo eq of 15, achieving tensile strengths of 1000 MPa and elongations of 2–5%, owing mostly to the hardness and fragility that ω phase provides as detected by the TEM [254]. Moreover, this transition from press & sinter to MIM is justified despite the fact that MIM does not achieve the same tolerances as axial pressing, but the improvement of mechanical properties and, more importantly, fatigue resistance is very important in these alloys, even though there are still some limitations in the amount of additives available for alloying development. In addition, Zhao et al., investigated the possibility of Ti–Nb mixture with Nb content ranging from 10% to 22% via MIM, evaluating porosities as high as 6% in different alloys but with a TiC content of 2% in Ti–22Nb. These properties and microstructures are compared to HIP-treated alloys in which the porosity has been completely closed. Even with this treatment, mechanical properties only slightly increase, from 649 MPa for as-sintered Ti–22Nb to 687 MPa for HIP Ti–22Nb, and elongation shows a slight decrease, from 1.43% for as-sintered Ti–22Nb to 1.30% for HIP Ti–22Nb, as well as an increase in elastic modulus from 70 to 75 GPa following HIP treatment [255]. Similarly, Xu et al., evaluated the alloy Ti–20Nb–10Zr generated by employing an elemental powder mixture and the MIM process to get samples for tensile and fatigue testing with porosities of 2.7%. (average pore diameter 7.3 μ m), considering two different levels of material purity based on vacuum conditions used during sintering. In their fatigue tests, they obtained a greater number of life cycles than Ti–6Al–4V or Ti–6Al–7Nb, which are also processed and tested under the same conditions, as well as a certain amortiguation in their behavior, which may indicate the fatigue limit. Additionally, the study of two lots with different levels of purity found that for 10 7 cycles of life, there is no noticeable difference between them, and that only when stress applied is more than 780 MPa the sintered more clean allow for more than 10 6 cycles of life [256]. 4.1.3. Porous alloys One technique to minimize the weight and rigidity of materials is to obtain them with a given porosity and employ one of the disadvantages of powder metallurgy, their residual porosity, to lower and manage the elastic and plastic properties. In addition to greatly increasing the specific contact surface of great significance in several applications. Consequently, powder metallurgy is one of the technologies suitable for the production of such products, and the development of porous Ti materials has included techniques such as sintering without the use of microspheres or the use of various spacers to control the quantity and size of the pores. Moreover, porosity has also been generated in the form of channels with extremely precise dimensions using trapped gas and freeze casting techniques. Finally, the incorporation of additive manufacturing techniques in the production of porous Ti allows for the design of pores and the production of porosity with various structures, whether through the use of laser, electron, or electric field [29,257,258]. i. Pressureless sintering The sintering of microspheres without pressure has been used to obtain porous coatings of Ti on Ti alloys, as proposed by Pilliar et al., who proposed a new technology using porous coating to improve bone tissue ingrowth while simultaneously improving fixation while the bulk material provides mechanical support. As described by Pilliar et al., a new technology using porous coating to increase bone tissue ingrowth while improving fixation whereas the bulk material offers mechanical support has been obtained by sintering microspheres without pressure [206]. Moreover, the impact of these porous coatings on the fatigue resistance of Ti and Co–Cr implants was obviously evident, and this is one of the most concerning features in porous alloys. Since, sintering occurs at high temperatures and changes the microstructure of the base alloy, particularly if it has been subjected to thermal treatments, leading in granular engrossment and lamellar microstructures whether the alloy is Ti–6Al–4V or Ti–6Al–7Nb. For example, Reig et al. studied the sintering of Ti–6Al–4V alloy by microspheres [259], and they investigate the effect of temperature and sintering time, which increases from 1300 to 1400 ◦C and from 2 to 8 h, bending strength increases from 140 to 300 MPa for atomized particles ranging from 180 to 240 μ m. However, if the microspheres are larger in size, between 425 and 600 μ m, bending strength between 60 and 100 MPa are obtained, which are insufficient for the majority of applications. Fig. 14 shows the union between the microspheres of Ti–6Al–4V, with a large size 450–600 μ m, pressureless sintered at 1400 ◦C and the microstructure of these sintered microspheres at 1300 ◦C for 8 h. Importantly, this particle size affects the size and number of pores that can be obtained by pressureless sintering of microspheres since the size of the pores is one-fourth of the size of the microspheres while the porosity only reaches 32% regardless of particle size [260]. To address the mechanical behavior of these materials, the authors use the Zero-Order Reaction Rate Model and Transition State Theory, which allows them to predict bending strength based on particle size and the application of different sintering temperatures and times [261]. However, because to the difficulty and high expense of obtaining Ti beta atomized powder, there are no reports describing the use of these pressureless sintering techniques in these areas. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3567 However, due to the difficulty and high cost of obtaining atomized pre-alloyed Ti beta powders, there are no documents on the application of pressureless sintering techniques for pressureless sintering of microspheres in these alloys. Nevertheless, two-step processes with pressureless sintering after cold isostatic compaction or mechanical alloying have been used on both Ti–6Al–4V [245–247] and Ti-6554 Alloy (Guo et al., 2024), although in these processes, densities around 98% and non-porous have been obtained [248]. Zhang et al., in 2017, used pressureless spark plasma sintering to obtain Ti–Mo foams with nano and sub-micro porous. They used a solid-state dealloying method in the preparation of foam from (TiMo)1-xCux master alloys. The dealloying is performed inside Mg powders where the Cu atoms of the master alloy are selectively dissolved. As a result, this process yields pore sizes of less than 50 nm, with nano-porosities of 55–75% [262]. ii. Low pressure sintering While conventional P/M uses high compacting pressures to provide high densification pieces, in contrast, the use of low compacting pressures, which are lower than usual, allows for greater porosity in green pieces and therefore in the sintered material. However, porosity is distributed uniformly within the finished piece. The amount of porosity depends upon the compacting pressure, particle size and shape, chemical composition, and sintering temperature and time. Rossi et al., investigated the mechanical behavior of Ti–34Nb–6Sn alloys compacted at 100 and 200 MPa and sintered at 900 ◦C for 2 h [60]. Additionally, they assessed the corrosion resistance using electrochemical tests and ion release. The higher Nb content generates a lack of diffusion, since the temperature has been insufficient to achieve complete chemical homogeneity, resulting in a porosity of 26% for 100 MPa compacting and 22% for 200 MPa compacting. In this manner, the elastic compression limit increases from 245 MPa for 100 MPa compacting to 320 MPa for 200 MPa compacting. Thus, the use of low powder compacting pressures, together with the use of different particle sizes, allows for the production of green pieces with very low densification and, as a result, a large number of porous in the sintered material [263]. Due to its ease of use, this technique has gained popularity in the production of materials with gradient porosity, which will be determined by the many layers with varying sizes [264]. Nevertheless, the pore size will be varied and difficult to control, making the final mechanical characteristics highly dependent on the largest size of defect obtained. Similarly, Torres et al., obtained and characterized cylinders with gradient porosity with CP-Ti with a particle size between 10 and 70 μ m of 25 μ m, alternating layers with different density and therefore different porosity in such a way that they achieve a gradual longitudinal porosity with a distribution of porosities from 12 to 45 %. These materials have an elastic modulus of 25 GPa and they conclude that an optimization of the process can produce materials with a good balance of properties that can reduce the effect of stress shielding in their application as biomedical prostheses and thus avoid bone resorption [265]. iii. Space-holder technique To obtain a larger porosity and a greater interconnection between the pores, different spacers have been used that consist of solid materials that can be removed at low temperature without excessive contaminating the Ti powder. Prior to extraction of the matrix and manipulation to eliminate the spacer thermally, the powder-spacer mixture can be compressed with the required pressure. For example, urea powder (carbamide) has been used, which can be removed at 200 ◦C with little CP-Ti contamination, because the oxygen content increases from 3000 to 5000 ppm [207]. However, when the porosity increases, the tensile strength decreases rapidly, from 760 MPa for 7% porosity to 68 MPa for 60% porosity and finally to 6 MPa for 77% porosity, which is also dependent on the distribution of the size of the particles used as spacers [207]. Amig´ o et al., employed ammonium bicarbonate (NH 4 HCO 3 ), which removed at 80 ◦C with the formation of CO 2 and H 3 N, resulting in very low contamination. Fig. 15 shows CP-Ti foams with two different spacers carbamide and NH 4 HCO 3 with particle size between 250 and 500 μ m, compacted at 300 MPa and 200 MPa respectively and vacuum sintered at 1200 ◦C. Spacer sizes ranging from 250-500 μ m to 500–1000 μ m are used, with compacting pressures of 100 and 200 MPa. In this context, bending strength drops up to 500 MPa with 25% porosity and remains between 100 and 200 MPa with porosities ranging from 40 to 60%, depending on the size of the spacer used and the compacting pressure [266]. Similarly, Torres et al., investigated a similar material by adding 50% NH 4 HCO 3 CP-Ti HDH, resulting in a 43% interconnected porosity [267]. Moreover, a compressive pressure of 200 MPa results in an elastic limit of 220 MPa. They evaluated the cytotoxicity of various porous samples and discover that none of them were toxic [267]. In addition, this same group developed Ti pieces with radial porosity gradients for use in orthopedic implants. In this scenario, CP-Ti combined with NaCl of 180 μ m diameter, was used, after removing NaCl with water at 40 ◦C for 24 h, Ti samples were sintered at 1250 ◦C for 2 h [268]. In Fig. 16, an example of porosity obtained is shown with details of pores roughness. They employed different combinations of porosities ranging from 40 to 70% to get modulus comparable to cortical bone with a porosity of roughly 50% with experimental compression resistances ranging from 40 to 200 MPa, which are quite comparable to cortical resistance [269]. Magnesium in the form of grains is another spacer used since its melting temperature is low and its boiling temperature is high, which reduces dramatically when sintering in high vacuum [270]. They mixed CP-Ti with Mg grains (150–850 μ m) in proportions of 50, 60, and 70% before compacting in circular matrices at 400 ◦C and 600 MPa uniaxial pressure. Removing Mg with HCl and ethanol for 24 h, and then using ethanol as a buffer solution to avoid the corrosive and damaging effects Fig. 14. a) the union between the microspheres of Ti–6Al–4V, with an average size of 450 μ m, pressureless sintered at 1400 ◦C. b) Microstructure of these sintered microspheres at 1300 ◦C for 8 h [260]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3568 of hydrochloric acid on Ti powder. Then, they assessed the distribution, volume, and size of the rather homogeneous pores, as well as the interior roughness caused by the hydrochloric acid attack. They have compressive properties of 100 MPa for a porosity of 60%, but close to 300 MPa for a porosity of 50%, despite the fact that these values are obtained by applying the load in a direction parallel to the plastic Fig. 15. CP-Ti foams with two different spacers (particle size 250–500 μ m) compacted at 200 MPa and vacuum sintered at 1200 ◦C: a) Foam fracture obtained with 55 vol% with carbamide b) Microstructure of a 45 vol% foam using NH 4 HCO 3 as spacer [266]. Fig. 16. CP-Ti foam obtained with 40 vol % NaCl as spacer a) general appearance of the tensile fracture of the foam. b) Detail of the roughness at the bottom of the pores [269]. Fig. 17. Porosity produced by the addition of 10% Mg in the Ti–34Nb–6Sn alloy Press to 200 MPa and different sinter temperatures: A) Blend elements sinter to 900 ◦C. B) Mechanical alloying sinter at 900 ◦C. C) Porosity due to the elimination of Mg particles in mechanical alloying sinter at 1100 ◦C. D) Detail the sinter microstructure at 1100 ◦C in the inter-porous zones [271]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3569 deformation obtained by Mg grains when compacted at 400 ◦C. When the application to the deformed pores is normal, these values decrease to 100 MPa and 80 MPa, respectively. Rossi et al. studied the impact of low Mg content on the properties and densification of Ti sintered at temperatures of 900 ◦C and 1100 ◦C for 2 h after compacting to a pressure of 200 MPa (Fig. 17). However, these compacts have a transverse tensile strength of 85 MPa, which ensures the ability to obtain a complex near-shape when Mg is removed, which is complicated when NaCl is used as a spacer [271]. Due to a lack of diffusion between the elemental powder used, notably Ti and Nb, less than 80% of the β phase is kept in all circumstances; therefore, the quantity of α phase retained corresponds to Ti that has not interacted to generate the β phase with Nb. Using the impulse excitation method, they observed that the elastic modulus ranged from 31 GPa to 49 GPa depending on the sintering temperature, with high hardness, and that final oxygen levels of up to 1.3% by weight were detected. They also assessed both the physical and biological properties of Ti–34Nb–6Sn/Mg, which was produced by mechanical alloying of hydrated elemental powder for 60 min. The green parts were sintered for 2 h at 700 ◦C and 800 ◦C, yielding porosities of 38 and 28% at the two sintering temperatures, respectively, and hardness of 92 and 120 HV. After 72 h, biochemical analyses of cell adhesion revealed that cells had proliferated both on the surface and within the pores, indicating that they are perfectly suitable as biomaterials [272]. iv. Expansion of pressurized bubbles The trapped gas expansion technique is carried out in two hightemperature stages. In the first stage, the powder is placed in a container that is compressed with argon gas to a specific pressure, typically 0.33 MPa, before being subjected to HIP, which is normally 100 MPa and temperatures ranging from 900 to 1100 ◦C for one or 2 h. HIP-ed component is designed for long-term gas expansion treatment at temperatures between 1100 and 1250 ◦C, generating plastic deformation at high temperatures [208]. The most well-known Kearns technique is controlled Ar expansion in metal compacts, as the group was the first to obtain Ti–6Al–4V fumes using this approach [273]. Furthermore, the original Ti–6Al–4V alloying method has been used effectively to produce CP-Ti and Ti–Ni powders [274,275]. Davis et al. used CP-Ti 130 μ m packed to a 70% relative density in a steel container, backfilled with 0.33 MPa argon gas, and soldered. The material was densified by HIP for 125 min at 890 ◦C and 100 MPa, then expanded in 4-min cycles between 830 and 980 ◦C for a total experiment time of 30 min to 24 h. After 5 h, the porosity is 38%, with ultrasonic elastic moduli of 39 GPa and yield strength to compression of 120 MPa compared the expected 160 MPa [276]. Similarly, Spoerke et al., obtained a porosity of 22% with pores with widths of 200 μ m after subjecting the HIP-ed material to 960 ◦C for 5.25 h to obtain an argon gas expansion. They then used a rotating bioreactor to analyze the behavior of pre-osteoblasts in-vitro, and the results were remarkable [277]. Dunand offered an in-depth review of Ti foam fabrication, distinguishing two procedures based on pressurized bubble growth within the section [263]. The first creep expansion technique incorporates the argon expansion process developed by Kearns et al., as well as the Low-Density Core Approach (LDC), which Martin et al., developed from the Kearns approach to enable the manufacture of Ti–6Al–4V sandwich structures [278]. Boeing (USA), a leading aerospace company, has used this technology in structural components made of Ti–6Al–4V and porous Ti alloys [279], but no papers have been found where this technique is used with β-Ti alloys. The Second Superplastic Expansion Processes, as established by Dunand et al., when superplasticity is applied in the liquid state of CP-Ti and Ti–6Al–4V, are characterized by a high rate of deformation [280]. This super-plasticity generates more porosities than creep expansion which is induced by heat cycles near the CP-Ti and Ti–6Al–4V transformation, [281,282]. Additionally, Murray and Dunand investigated the influence of this thermal history on titania superplastic expansion due to argon expansion on the creation of pores. They explored chemistry and microstructure in the first half and model the chemistry in the second half [283,284]. Meanwhile, oppenheimer and Dunand used Ti–6Al–4V spherical powder with different sizes of 75, 115, and 165 μ m that are encapsulated in mild steel canisters and backfilled with argon gas at 0.33 MPa of pressure before densifying by HIP at 100 MPa and 950 ◦C for 4 h. After the steel layer is removed, the foam is baked in a high-vacuum oven with an isothermal treatment at 1030 ◦C. They achieve up to 30% porosity while attaining compression limits of 500 MPa and elasticity modulus of 55 GPa by expanding pressurized argon trapped in micro-size pores during the initial stage of HIP consolidation [285]. In another study, Nugroho et al., employed this process to create Ti–29Nb–13Ta-4.5Zr emulsions, with the initial stage consisting of mixing the powders in a steel container. After filling the container with 0.68 MPa argon gas, it was subjected to a HIP at 1100 ◦C for 1 h under 100 MPa in an argon atmosphere. After that, the piece was subjected to 10 h of gas expansion at two temperatures, 110 and 1225 ◦C. This resulted in porosities ranging from 15% at 1100 ◦C to 27% at 1225 ◦C. The resistance to corrosion increases with temperature, however electrochemical corrosion demonstrates a stronger resistance to corrosion at lower temperatures, despite the notable pitting effect [286]. However, no new work has been found on porous β-Ti alloys obtained by superplastic expansion processes. v. Freeze casting The freeze casting technique, which was developed to manufacture ceramic products with aligned porosity, has also been applied to Ti alloys [287]. The freeze casting technique is used to prepare a solution containing the various elements that will form the alloy. Liquid N 2 is used to cool the slurry at a rate of less than −10 ◦C/min. Powder normally settles and forms canal structures with the solvent. The ice is removed from the frozen green substrate by sublimating it at low temperatures and pressure for several hours. Chino and Dunand used an aqueous slurry of Ti powders with particle sizes less than 45 μ m that was solidified directly, yielding a consistent performance in ice dendrites with interdendritic areas containing a high Ti content. Following freeze-drying to remove ice dendrites and Ti sinterization, foams with 57–67% aligned holes and a length of several millimeters are obtained. Similarly, Mao and others, in a similar way achieved similar results for TiH 2 contents of 20, 25, and 30% by obtaining porosities ranging from 40 to 50% using elastic moduli ranging from 25 to 32 GPa [287]. Koh et al., used the freeze casting technique to produce Ti foams with varied porosity and TiH 2 /camphene suspensions with changing hydration levels. After freeze drying, the material was exposed to a 3 h vacuum treatment at 1300 ◦C. This process produces layers with varying porosities of 35, 50, and 75%, with highly good union between them, depending on the TiH 2 amount used. Furthermore, they discovered that as the period of freezing increases, so does the number of particles suspended, generating pores ranging in size from 100 to 270 μ m after 7 days [209]. Du et al., developed porous Ti scaffolds using freeze casting technology for use as structural cores of sandwiches, heat exchangers, or filters, with porosities ranging from 15% to 73% that provide compression resistances in the parallel to freezing direction of 60–86 MPa, while compression elasticity varies from 0.8 GPa to 11.2 GPa, significantly lower than those reported by Chang et al. [288]. Dunand and Li produced Ti foams using a camphene solution, finding that reducing the size of the pores improves the foam’s mechanical properties [289]. Chang et al., assessed the impact of sintering temperature, sintering freeze dry Ti scaffolds from 800 to 1100 ◦C and varying the temperature every 50 ◦C, obtaining porosities ranging from 61% for 800 ◦C to 20% for 1100 ◦C that provide elastic modules ranging from 2.04 to 7.16 GPa and compression resistance ranging from 100 to more than 1100 MPa even though the pieces are tested until fracture A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3570 [290]. Jenei et al., prepared Ti foams with relative densities of 0.49, 0.37, and 0.32, assessing mechanical performance in relation to porosity and grain size after sintering treatments at 850 and 1000 ◦C for 4–6 h, depending on the density and grain size chosen [291]. In contrast, Jenei et al., on the other hand, use freeze casting to produce a Ti–5W foam with a porosity of 32–33% and pore diameters of 15–17 μ m, resulting in an elastic modulus of 23 GPa. The addition of W to Ti raises the compression resistance from 196 MPa to 235 MPa [292]. Choi et al., examined the same alloy by producing an elemental suspension of Ti and W (about 50%) in water with 2.8 w% Poly (vinyl alcohol). By presenting an acicular microstructure Widmast¨ aten α /β with W separated in the β phase, they obtained a 2%W solution in Ti. Ti–5W foam has a compressive yield strength of 323 MPa, which is roughly 20% larger than Ti foam, which they are testing at 256 MPa, which is significantly higher than the values reported by Jenei et al. [293]. Weaver et al., investigated the relationship between the structure and the freeze-cast processing of Ti–6Al–4V alloy using a mixture of low molecular weight chitosan and sodium carboxymethyl cellulose in water with 1% acetic acid as a filler, increasing suspension viscosity and decreasing the effect of premature powder sedimentation, which added in concentrations of 20 and 30%. Porosities of 65 and 34% are obtained. The elastic modulus is 7–29 GPa, with compression tensions of 80–400 MPa for 20% deformations, which are slightly higher than for camphene-prepared foams [294]. Li et al. obtained a volume shrinkage of 51 to 47% by combining 15–30% Ti–6Al–4V powder in an aqueous solution with 0.2 vol% Xanthan gum as a surfactant, achieving a porosity gradient of more than 20% from the foot to the head and peak values of 55–74%. As a result, when sintered at 1200 ◦C for 1 h, their elastic compression modulus ranges from 2 to 12 GPa, with compression resistance ranging between 40 and 300 MPa, and they later endure with the addition of SiC fiber to the suspension [295]. Additionally, they noted that the addition of SiC (2, 4.5, and 10 vol%) increases compressive strength to 400 MPa, but larger amounts of SiC fiber generate small amounts of Ti 5 Si 3 and TiC precipitates, which can reduce compressive strength and modulus to values less than those of Ti without the use of force [296,297]. Similar properties to those reported by Chen et al., control porosity, porosity size, and porosity distribution by using gelatin as a binder and polyvinylpyrrolidone as a dispersant. They obtained compression resistances ranging from 57 to 142 MPa and elastic moduli ranging from 2.8 to 5.1 GPa [298]. The addition of zirconium acetate to the solution influences the composition of the pore structure and mechanical properties of porous Ti. As a result, the structure of the pores can be customized to resemble honeycomb structures [298,299]. Trueba et al., analyzed the influence of the freeze-casting process parameters (diameter, mold material and temperature gradient) on the porosity and mechanical behavior of cylinders with elongated porosity obtained using a 5 w% of polyethylene glycol in binder demineralized H 2 O. They confirm the formation of a higher gradient porosity in the top, as the resistance to compression and the elastic modulus decrease with the higher porosity [300]. However, the freeze casting technique has not been applied so far to obtain porous β-Ti alloys. 4.2. Spark plasma sintering (FAST/SPS) Nowadays, there is a growing demand in using rapid sintering techniques such as Spark plasma sintering also called Field Assisted Sintering Technologies (FAST/SPS) to manufacture various forms of Ti alloys with a fine-grained microstructure attain outstanding mechanical characteristics [301–303]. In FAST/SPS, a compact powder is subjected to high pressure when sintering by applying an electric current passing thought the pressing die to the powders, which improves densification and lowers the sintering temperature and holding time. The grain growth rate was also revealed to be lowered by this technique, producing grains with outstanding mechanical features [304]. The advantages of this approach over ordinary PM include the quick sintering in SPS and the outcomes of a smooth size process. When applying the SPS technique to manufacture porous structures, there are two common approaches. The first approach involves partially sintering loose particles or fibre at low temperature. The SPS involves loading Ti powders into a graphite die, which is then subjected to a uniaxial force during the sintering. Spark discharges in the spaces between the particles are used to heat the material. This discharge refines and excites the particle surface, and the electric current produces a joule heating between the particles as a consequence and plasma is generated as a result [305]. Regarding both conductive and nonconductive powder, densification is accomplished in a brief period of time during the sintering cycle owing to the electric current delivery, which also decreases the rate of grain growth [305]. Due to its rapid heating rate, SPS is used to produce nanocrystalline materials by minimizing the grain size [306]. A low voltage is used in the SPS system, which leads it to produce high electric current. The process uses graphite tools that allow to apply high cooling rates varying from 150 ◦C/min to 400 ◦C/min and auxiliary gas cooling can produce temperatures of roughly 2400 ◦C [307]. SPS employs loads around 50 and 250 KN while running under vacuum, thereby accelerates the densification process [308]. According to studies, the SPS can heat up fast (at a rate of 1000 ◦C/min), allowing for shorter material manufacturing cycles [309]. The flow of electric current between the powder particles affords the SPS’s mechanism a quite special quality. Plasma is produced when the particle surfaces are heated. The temperature refines and energies the particle surface, which leads the impurities on the powder particles to evaporate. The heated particles later melt and merged to form a neck. By directing an electrical current through the particles in the neck during the second stage, joule heat is generated. The diffusion of the particles in the neck is increased by joule heat, promoting their growth. Eventually, when the heated material has turned soft under pressure, plastic deformation occurs. Furthermore, the compact is obtained due to densification [306,307]. Since there are no intermediate ingredients used in this approach, it has the advantage of minimizing time, energy, and expense while reducing the chance of contamination. 4.3. Additive manufacturing (AM) Typically, powder metallurgy, foaming, and casting technologies are used for manufacturing titanium alloys [310,311]. Nevertheless, these conventional methods need additional time, energy, and material resources due to their various processing stages [312]. Furthermore, there are a lot of issues with these approaches. For instance, titanium’s low elastic modulus, rapid oxidation rate, and limited heat conductivity make it difficult to deform over a high temperature [312]. Oxygen and titanium will react throughout the machining process to create a tough oxide coating on the surface. Tool wear will also be accelerated because excessive heat accumulation will result from the machining process’s heat generation not dissipating rapidly enough [313]. With the aid of moulds, large-scale components may be effectively formed by casting. Nevertheless, specific tools are needed for titanium alloys, which are quickly oxidized at high temperatures. Therefore, components with various forms are typically in demand for unique moulds. As a result, the cost of production will rise dramatically because of these extra tools and processes. Additionally, molten titanium alloys have very high viscosities as compared to other metals, which results in poor filling fluidity properties. When they solidify, they are susceptible to flaws like porosity [314]. As a result, cast Ti parts frequently fall short of performance standards. Likewise casting, forging similarly necessitates the use of expensive specialist equipment. Furthermore, forging large-scale components with intricate forms is challenging. Titanium alloys made using the aforementioned traditional techniques are therefore more expensive than many other metal alloys due to their difficult extraction process, difficult melting and manufacturing process, and mechanical imperfections [314]. Conversely, additive manufacturing features easy processing steps A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3571 and inexpensive equipment. Without the need for moulds or specialized tools, it may immediately create components with almost net shapes. Therefore, AM may be a more affordable option than casting and forging. The manufacturing of items with intricate forms requires the use of the porous structure approach. The advancement of additive manufacturing (AM) technology in the last few decades have accelerated and enhanced implant innovation [9,315,316]. Lacking the need for a separate processing program, the AM technology is capable of producing the porous implants instantly. The apparent benefits of these implants include their light weight and low Young’s modulus, that can encourage the formation of new bone cells. The ability of AM to precisely machine the pore structure of complex unit structures in comparison to traditional techniques has raised serious concerns. The term additive manufacturing (AM) describes a wide range of methods that convert 3D digital drawings into real, working components in a manner similar to how an office printer prints 2D digital data onto paper. "A process of joining materials to make objects from 3D model data, usually layer upon layer, as opposed to subtractive manufacturing methodologies," is how ASTM F2792 describes additive manufacturing [317]. The process begins with slicing a 3D solid computer assisted design (CAD) model into layers. Then, digital data is generated according to the geometry of each layer [316]. The selected layer-by-layer combining of unit materials—usually sheets, wires, and powders—into final (near) net-shaped structures is subsequently accomplished using this digital information as the tool path. The enhanced versatility of AM technology is provided by this working concept [9]. By carefully applying material just where it is required, complex structures that are not achievable by machining, such interior cooling channels, may be easily made. Instantaneous local manufacturing on a global scale is made feasible by the digital models’ freedom and worldwide shareability. Furthermore, AM technology eliminates the need for fixtures, moulds, and tooling, which significantly cuts down on the lengthy lead times associated with the design and manufacture of moulds and/or tools [9]. Another key component of AM technology is near-net shaping, which drastically reduces waste and material consumption [316]. The latest innovation in the realm of metal fabrication is additive manufacturing (AM), frequently referred to as three-dimensional printing (3D printing) [318]. The near-net-shape forming capability of AM is one of its most important material-related benefits [319]. Due to the fact that AM has been recognized as a promising technology offering significant geometric design flexibility [320], a great deal of effort has gone into developing an appropriate AM method for producing titanium alloy components with enhanced mechanical qualities. Metallic items with complicated geometries may be created in a single step without requiring time-consuming and costly machining processes by layer-by-layer assembly of the component based on the digital model. 3D printing, is a technology which produces products from 3D model data via melting a layer of powder with an electron or laser beam under computer control to create intricate component structures [316]. Aerospace and biomedicine have shown a great deal of interest in additive manufacturing because of its ability to precisely, quickly, and efficiently regulate the intricate interior pore structure and form of the components that are manufactured. In recent years, different alloys have been incorporated for additive manufacturing techniques by obtaining different pre-alloyed powders of suitable quality and price. Cheng et al. (2024) analyze the manufacturability and mechanical properties of Ti–35Nb–7Zr–5Ta porous titanium alloys produced by laser powder-bed fusion, requiring lower energy than for dense materials. These alloys were topologically optimized considering the compressive strength of cortical bone, achieving minimal defects in the material [321]. 4.3.1. The foundations of metal AM techniques Despite beginning with organic materials, AM technology has a unique niche in the production of metallic materials because it can fulfil performance criteria and improve manufacturing efficiency, sustainability, and expenditures as well as energy savings. Various AM methods have been devised to leverage these significant advantages. A number of these processes have been particularly created for the manufacture of metallic components, or they have been modified from polymer-based procedures. These including sheet lamination (SL), directed energy deposition (DED), material extrusion (ME), material jetting (MJ), binder jetting (BJ), and VAT photopolymerization (VP). For Ti-based materials used in medical applications, powder bed fusion (PBF) and DED are the most used AM methods [318]. PBF and DED laser-based technologies operate in an inert atmosphere, whereas electron beam technologies operate in a vacuum [322]. Selective laser melting and electron beam melting are the two most common additive manufacturing processes [316]. 4.3.2. Powder bed fusion (PBF) The basis of PBF technologies is the deposition of a layer of metal powder on the build plate. The laser/e-beam scanning of the powder layer is driven by the section drawing (slice of the 3D object model), which causes the powders to melt partially or completely. Each model slice goes through the process once again until the product is produced. The bonding between the various layers is controlled by the powders’ melting and re-solidification [323]. Selective Laser Melting (SLM) and Electron Beam Melting (EBM), the two basic PBF procedures, use laser or electron beams as heating sources, respectively, for powder fusion. SLM technology, which employs a laser that melts metal powder, is the technology most metal 3D printer’s use. i. Selective Laser Melting (SLM) The initial report on the selective laser melting system was published in 1995 in Germany by the Fraunhofer Institute ILT in Aachen. SLM is an AM approach that accounts for more than 80% of the market for metal 3D printers [324]. SLM is unique from other powder bed fusion processes, such as selective laser sintering (SLS) and direct metal laser sintering (DMLS), in that it completely melts metal rather than sintering it [324]. Prior to using an SLM printer, a 3D model needs to be translated into a format [Stereolithography (STL)] files that is compatible with it [325]. Using so-called slicing software, the model is then divided into layers, with each cross-section of the model representing a layer that will be printed by the SLM printer [326]. The proper metal powder is added to the powder stock through a hopper or automated cartridge once the model has been placed into the printer [326]. To protect the metal from oxidation while subjected to the intense laser beam, argon gas is fed into the sealed manufacturing chamber [327]. These lasers are usually focused high-power ytterbium fibre lasers, (Ytterbium fibre/Nd:YAG, λ ≈1 μ m and power <200 W) which are hot enough to completely melt the metal powder but also produce significant temperature differences. The sealed chamber is heated to reduce temperature fluctuations between the laser and powder in order to take these variances into account [327]. A computer-aided-designed (CAD) 3D model that has been sliced into slices of a specific thickness serves as the toolpath for the AM equipment, which is used to fabricate the product [328]. In a building chamber, a substrate plate is covered with a thin layer of metal powder to start the construction process. Once the powder has been placed, a high-power laser is employed to melt and fuse specific areas in accordance with the processed data [328]. After the laser scanning is done, a new layer of powder is added on top, the building platform is lowered, and the beam scans a second layer [328]. This procedure is then repeated for other powder layers until all necessary components have been fully assembled. After the laser scanning task is performed, the construction chamber is cleared of loose powders, and the product can either be manually or electrically separated from the substrate plate [329]. SLM is therefore a viable technology for the net-shape production of parts employing a wide range of high-quality engineered metals that are difficult to build using other traditional processing techniques. With the A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3572 assistance of this technology, devices with fully interconnected porosity needed to mimic the mechanical modulus and structure of cancellous bone could be manufactured. Table 4 classifies the many defect types that frequently influence the nodes and struts that comprise the microstructure of metallic implants produced through additive manufacturing in terms of their sources, impacted components, and effects. In contrast, Table 5 displays a classification of techniques for reducing AM defects. ii. Electron Beam Melting (EBM) Due to its simplicity of use in producing flawless metallic components, EBM has lately gained interest in the industrial industries. EBM, one of the metal AM technologies, has apparently been capable of fabricating complex products with no defects, especially when manufacturing porous structures with full density, specified external shapes, and interior designs based on Ti materials [330,331]. EBM is an established technology for producing metallic parts via a layer-by-layer approach as compared to other AM techniques [332]. In the production process, a quick-moving high-energy beam of electrons is used as a heating element [332]. It sweeps the powder bed and immediately melts powder, producing fully dense metallic components. Electromagnetic lenses are employed to guide the electron beam, ensuring that it is inertia-free. As a result, scanning velocities on the powder bed can increase to 10 5 m s −1 213 . Considering Ti and aluminium metals have a high reactivity in contact with oxygen and nitrogen, the EBM manufacturing procedure is carried out in a vacuum chamber (10 −4 –10 −5 mbar), which is particularly crucial for these metals [212]. Layer by layer, the sample is constructed with each layer passing through four procedures: Preheating, melting, dropping the construction plate, and 5) deposition a fresh layer of powder [212]. The samples are cooled to room temperature inside the equipment chamber after completion of the final layer [212]. 4.3.3. Directed energy deposition (DED) While employing directed energy deposition (DED), a substance is introduced into the laser or electron beam’s nozzles in order to melt and deposit it in accordance with the 3D design on the substrate. In the case of lasers or e-beams, the material is fed from a coaxial nozzle in the form of metallic granules or metallic wires. DED techniques allow greater compositional and design flexibility by utilizing different feeding nozzles within a single building project. Furthermore, with this approach, new components can be constructed as well as damaged components can be repaired [333]. In comparison to DED technologies (2000–4000 μ m of beam size and 300–1000 μ m for Laser Metal Deposition), PBF techniques are distinguished by a smaller beam size (50–500 μ m for SLM and 140–250 μ m for EBM) and smaller layer thickness (15–150 μ m for SLM and 50–200 μ m for EBM) [333]. Due to these advantages, PBF technologies enable an improved surface finish, but their application requires longer processing times. Thus, PBF is more recommended for small parts requiring high levels of precision and complexity while DED is better suited to large parts with a coarser finish [323]. Significant research has recently been conducted on metal additive manufacturing (AM) for a variety of alloy systems, particularly Ti alloys [11,334]. The majority of the studies on titanium alloys has gone into improving the processing conditions and powder feed stock, especially for Ti–6Al–4V (wt.%) as well as a few additional commercial alloys. The literature presently in publication has limited data on AM fabrication of metastable beta Ti alloys; however, a few recent investigations employing selective laser melting (SLM) and electron beam melting (e-beam) have been published [11,257,312,335,336]. The majority of this study has been devoted to enhancing processing parameters or assessing the associations between processing, structure, and properties of Ti–24Nb–4Zr–8Sn (wt.%) alloys, often referred to as Ti2448 alloy [257,258,335,337]. Due to its low modulus of elasticity and absence of components considered to be hazardous to humans, this alloy was designed for applications involving biomedical implants. Even though the Ti2448 alloy has a low modulus of elasticity that is about between 42 and 50 GPa, it is nevertheless higher than cortical bone and might potentially lead to stress shielding from an implant. A significant facet of this investigation is the creation of scaffolds with a high strength to elastic modulus. Scaffolds have the potential to further reduce the elastic modulus, becoming close to that of cortical bone. Nevertheless, the presence of Sn results in poor tensile and fatigue attributes since it may evaporate and cause defects to occur, such as porosity. In addition, studies conducted on this alloy showed that SLM and EBM need distinct ways to optimize and provide different microstructures and characteristics. This example shows how challenging it is to modify current beta Ti alloys for use in metal additive manufacturing. For Ti–Nb alloys, it was also demonstrated how mechanical characteristics might be tuned based on phase volume fraction and composition [336]. There are still many unanswered questions concerning metastable beta Ti alloys produced by metal AM, and it is expected that this will be a focus of substantial future study. To better comprehend the relationship between processing, structure, and property in metal AM, more study must be undertaken. The existing lack of understanding regarding the microstructure and characteristics of AM-processed metastable beta Ti Table 4 Structural defects in metallic additively manufactured implants. Geometric type Dimensional Inaccuracy On the Strut Waviness Thickness Varying Crosssection shape Cross-section eccentricity Over/under Size Missing Strut On the Node Parasitic Node Missing Node Surface Roughness Raw Powder Particles on the Boundary of Strut Orientation based Stair-stepping effect Upward/downward facing surface Microstructural Type Porosity Feed Stock Pores exist in Feed stock powder Pores created due to Feed stock oxidation Processing Induced Pores Metallurgical Pores LOF and incomplete melting Keyhole Material Property Imperfection On Strut On Node Table 5 Defect minimization methods. Process parameter optimization Scanning Speed Lase power Hatching distance Layer thickness Powder morphology Powder size Powder size distribution Post-processing Material Subtraction Machining Polishing Chemical etching Laser machining Shot peening Electro polishing Coating/ treatments Heat treatment HIP (Hot Isotactic Pressing) Anodizing Design Strategies and defect Simulations Thickness Compensation Strategy Design Constraints Overhang elimination A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3573 alloys is impeding their implementation. AM causes fast cycle heating and cooling, which can lead to bulk flaws or non-equilibrium microstructures that may lack identical characteristics as traditional wrought alloys with the same composition. To have a good enough knowledge of how an alloy behaves in an application and during processing, extensive experimental characterisation is needed. 4.4. Additive manufacturing of beta titanium alloys In recent years Additive Manufacturing (AM) has becoming as an innovative processing technology of materials. At the beginning, most of its main features were developed for polymeric materials due to the diversity, versatility, and low working temperatures [83,338]. However, ceramic, and metallic materials have been processing using AM despite the technological difficulties involved in forming these materials. Main studies of AM metallic materials are focused on preparation of metallic alloys powders, processing parameters of selected 3D manufacture technology, and final forming of the designed piece, which will be the determining parameter for the appropriate selection of AM technique, due to the high sintering or fusion temperatures of metals. Nevertheless, the exponential development of metallic biomaterials for implantology, and the need to design devices with increasingly complicated and difficult to access topologies, drive the imperative use of additive manufacturing techniques as the best choice for this purpose. Titanium and its alloys continue to be the metals alloys par excellence for biomedical uses; therefore, many studies have been carried out to use commercially pure Ti and Ti–6Al–4V in different additive manufacturing technologies [83,339]. However, beta titanium alloys have been rarely used due to the high melting points of the stabilizing elements that, depending on the manufacturing technique used, will induce the formation of undesirable phases, and with it, incompatible mechanical properties for use in implantology [340]. Despite that, in 2011, Zhang et al. used Selective Laser Melting (SLM) to produce beta Ti–24Nb–4Zr–8Sn components with near full density (>99%) for the manufacture of a sample acetabular cup [106]. Shortly after, Liu et al. have used same beta Ti–24Nb–4Zr–8Sn (Ti2448) alloy to prepare scaffolds by Selective Laser Melting (SLM) technique and exploring processing parameters and final structures characterization as pores distribution, geometrical accuracy, and the mechanical properties of the produced components [312]. Furthermore, they have employed Electron Bean Melting (EBM) for the fabrication of porous scaffolds with nominal porosities of 67.9%, 72.5%, 75.0%, 77.4%, 79.5 and 91.2%, assessing mechanical properties of EBM substrates, including super-elastic properties, Young’s modulus, compressive strength, and fatigue properties [337]. Additionally, same authors fabricated by EBM, Ti–24Nb–4Zr–8Sn substrates with 70 % porosity for deep understanding of processing-microstructure-property relationships of manufactured scaffolds, showing improving features compared to similar scaffolds fabricated with Ti–6Al–4V alloy [258]. In 2019, Ti–25Nb–3Zr–3Mo–2Sn was employed for the fabrication of substrates by Selective Laser Melting (SLM), achieving high strength and large tensile ductility [341]. More recently, Schaal et al. have employed Laser Powder Bed Fusion (L-PDF) fabricated beta Ti–22Zr–9Nb–2Sn samples to study mechanical properties via cyclic tensile tests that revealed an exceptionally high strain-hardening of prepared scaffolds [342]. In 2019, Hafeez et al. have used beta Ti–35Nb–2Ta–3Zr alloy for fabrication by Selective Laser Sintering (SLS) of tensile specimens for cyclic loading-unloading tensile testing, and microstructure evolution [343]. Afterwards with similar chemical composition, Yang et al. used beta Ti alloy Ti–35Nb–5Ta–7Zr (TNTZ) for the manufacture of scaffolds by Selective Electron Beam Melting (SEBM) reaching relative densities of >99.5%, and adding 2% of Si to improve wear resistance [344]. Moreover, biomedical Ti–30Nb–5Ta–8Zr (TNTZ) alloy have been studied by Luo et al., exploring mechanical properties, biocompatibility, and proteomics analyses of Selective Laser Melting (SLM) fabricated lattice structures [57]. Ti–35Nb–5Ta–7Zr scaffold was manufactured using Laser Power Bed Fusion (LPBF) which allowed simultaneous introduction of dislocations and twins into the alloy structure avoiding the strength–ductility expected relationship [345]. The group of Ackers et al. showed the potentiality of load-bearing implants fabricated by Power Bed Fusion (PBF-LB/M) of Ti-4.5Ta–4Fe- 7.5Nb–6Zr (TTFNZ) alloy, studying processing parameters, mechanical and corrosion properties, as well as PVD coat-ability [346–348]. The use of Wire Arc Additive Manufacturing (WAAM) technique for the preparation of Ti–3Al–8V–6Cr–4Mo–4Zr (Beta-C or ASTM Grade 19) samples with high stability, high strength, and high quality of metastable β-titanium alloys free of segregation defects, known as β-flecks [117,349]. On the other hand, the use of Directed Energy Deposition (DED) technique for the fabrication of Ti–1Al–8V–5Fe scaffolds allow to tune the microstructure with adjustment of process parameters, to have required mechanical properties without any other processing steps [350]. Additionally, the additive manufacturing technique of Selective Laser Melting (SLM) was employed for the preparation of dense samples of Ti-X (X =Cr, Nb, Mo, Ta) with well-developed crystallographic texture, from mixture of pure metallic powders [351]. 5. Structural and microstructural characterization After manufacturing, and regardless of their fabrication route, the achieved microstructural characteristics must be evaluated and correlated with properties relevant to specific applications. Therefore, describing the processing-microstructure-properties relationship is crucial for designing, selecting, and improving materials, including β-Ti alloys. Among the most commonly used microstructural characterization techniques are X-ray Diffraction (XRD), Differential Scanning Calorimetry (DSC), Scanning Electron Microscopy (SEM), Electron Backscatter Diffraction (EBSD), Focused Ion Beam (FIB), Transmission Electron Microscopy (TEM), Neutron and Synchrotron Diffraction, and Atom Probe Tomography (APT). Thus, structural and microstructural characterization can be applied to β-Ti alloys produced by different fabrication routes (casting, powder metallurgy, additive manufacturing, etc.), including those where surface modification is applied to improve their biocompatibility. Fig. 18 illustrates the main techniques for structural and microstructural characterization of β-Ti alloys. 5.1. X-ray diffraction (XRD) In order to characterize Ti alloys, X-ray diffraction (XRD) analysis is a crucial and fundamental tool. It allows for the identification and differentiation of potential phases that could form in alloys with a variety of compositions and β-stabilizing elements, including equilibrium α and β phases, and metastable ones: martensite α ′ , α ′′ and nanoscale ω phases [352]. When using conventional XRD patterns for the analysis and characterization of Ti–35Nb–7Zr–5Ta and Ti–35Nb–7Ta alloys (TNZT) following aging heat treatments at various temperatures, the use of a broad Cu-K α radiation (λ =1.54 Å) without a monochromator can result in the superposition of peaks for β and β ′ separated phases or even limit the detection of metastable nanometric ω phases when their fraction precipitated is less than 5% volume fraction reference. In the study by Ferrandini et al. [352], it can be observed that the XRD patterns mostly composed of the β -Ti phase do not significantly differ from one another when the Ti–35Nb–7Zr–5Ta alloy was aged at 200, 300, and 400 ◦C for 4 h after solution treatment at 1000 ◦C for 35 min and water quenching [352] (Fig. 19). However, when the Vickers hardness variation was evaluated for the Ti–35Nb–7Zr–5Ta alloy as a function of aging temperature (Fig. 20), it was observed that the hardness varies by 17% from the TNZT after solution (287 HV), rising to 336 HV after aging at 400 ◦C for 4 h. This raised the question of how the hardness was rising without a visible precipitation of metastable omega ( ω ) or equilibrium α phase. In order to detect phases in small volumetric fractions, such as the nanoscale metastable ω phase, one effective technique is to acquire X-ray A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3574 extent for the β phase. 5.5. Transmission electron microscopy TEM The transmission electron microscopy (TEM) technique is very important in the materials science and engineering, and in the research field of β-Ti alloys it is even more important, due to nanoscale omega phase precipitation, which occurrence in β-Ti alloys can damage mechanical properties due to fragile character of ω phase in such alloys microstructure. Besides that, TEM characterization can identify α phase in micron and nanoscale, α ′ martensite phase (hcp), α ״ martensite phase (orthorhombic), and more complex phase transformation, such as β and β ′ phase separation by spinodal decomposition, which occurs in nanoscale and just can be identified with a combination of XRD and TEM Fig. 29. SEM of morphology variations in the TiO 2 nanotubular structure as a function of anodization time at 20 V during (a) 6, (b) 12, (c) 24, and (d) 48 h [366]. Fig. 30. Representative (a) TEM bright field (BF) image of TiO 2 NTs in the as-anodized condition at 6 h/20 V, with NTs morphology in the amorphous state and (b) the corresponding selected area diffraction (SAD) pattern typical of amorphous structure and (c) multicomponent (Ti,Fe,Nb,Sn)O 2 composition obtained from the EDS spectrum [366]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3581 analyses, as showed previously in the XRD topic of this review and in the work of Afonso et al. [369]. The phase separation by spinodal decomposition in β and β ′ phases is displayed in Fig. 35for Ti–35Nb–7Zr–5Ta alloy aged at 400 C for 4 h. Fig. 35 a illustrates TEM image in bright field (BF) mode and the absence of α phase precipitation along the grain boundary of β-Ti grains. Fig. 35 b presents in detail the interior of β-Ti grain at the [1 1 1] β zone axis revealing diffraction contrast between the β (dark) and β ′ (bright) separated phases. Even in conventional TEM image, is already possible to have a good indicative of phase separation by spinodal decomposition in β-Ti alloys [369]. Increasing the magnification and reaching high-resolution transmission electron microscopy HRTEM images, as is observed in Fig. 36 for the Ti–35Nb–7Zr–5Ta alloy aged at 400 C for 4 h, is possible to detect different crystallographic phases. Fig. 36 b shows a HRTEM image along the [1 1 1] β zone axis. The respective Fourier transform (FT) reveals splitting of the (1 0 1) β reflections (white arrows). A schematic representation of the Fourier space map of the [1 1 1] β zone axis obtained from simulation using the JEMS© software includes separated phases with cubic (bcc) and tetragonal (BCT) structures [369]. In a recent work of Afonso et al. [356], Fig. 37, the conversion of porous powder metallurgy (PM) sample of β-type Ti35Nb10Ta and Ti35Nb10Ta3Fe alloys of tens of microns grain size to tens of nanometers grain size in 100% dense (no porosity) in just one step of severe plastic deformation (SPD) step through HPT (high pressure torsion). The TEM analysis combined with ASTAR crystallographic orientation mapping through ASTAR technique was a fundamental combination of characterization tools in order to determine nanoscale grain size of β-Ti alloys after HPT processing, because just TEM convention BF image can not detect neither determine clearly the grain size (and respective grain boundaries). In this way, ASTAR was used in order to determine as a more accurate way the nanoscale grain size, which is one of the main applications of ASTAR technique [356]. As can be seen in the TEM image Fig. 38 and respective SAD pattern for the PM sample of β Ti35Nb10Ta and Ti35Nb10Ta3Fe alloys, the occurrence of nanoscale metastable ω phase can be detected clearly in the right zone axis, showing an orientation relationship [1 1 0] β-Ti //[1 1 2 0] ω between the β-Ti matrix and metastable ω phase for both alloys. Regarding β Ti alloys after HPT processing with nanoscale grain size, the SAD pattern confirms typical ring-shaped patterns of nanocrystalline materials and in agreement with electron diffraction results, the ASTAR technique complementary determines nanoscale grain size of HPT sample. The TEM analysis is important as well, to characterize in nanoscale and also atomic scale, the structure and nature of TiO 2 nanotubes (NTs) formed, amorphous or (nano) crystalline. In the work of Chaves et al. [366], TEM images, Fig. 39, is showed in bright field (BF) mode of TiO 2 –based NTs formed in the anodization of Ti–Fe–Nb–Sn alloy anodized to 6 h/20 V and annealing at (a) 400 ◦C (b) 450 ◦C (c) 500 ◦C [366]. Besides that, corresponding high-resolution transmission electron microscopy HRTEM images (atomic scale) and respective selected area diffraction (SAD) patterns in each annealing condition were fundamental in order to define amorphous halo (characteristic of amorphous phase structure formation) and diffraction spots of anatase–rutile phases in the crystallized sample after heat treatment. It is worth to mention, Fig. 31. TEM analysis of MAO-treated Ti–15Zr–5Mo alloy at 400 V: a) Interface film-substrate; b) SAED patterns of the distinct regions of the film; c) HRTEM of the nanocrystalline region of the film; d) chemical ratio of the elements along the film [364]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3582 that just very few works in the literature show the TEM analysis of TiO 2 NTs, and the sample preparation can be much easier, just spread the nanotubes from the anodized sample over the TEM carbon coated (holey carbon type) copper grid with the aid of a metal blade. 5.6. Neutron and Synchrotron Diffraction Neutron and synchrotron diffraction techniques are powerful tools for investigating the structural features of materials. In particular, they provide insights into atomic-scale structures, defects, residual stress, as well as physical, chemical and mechanical behaviors of alloys, surpassing the capabilities of more conventional techniques [370–372]. Neutron diffraction is based on the interaction between neutrons and the atomic nuclei of the studied materials. Specifically, when a neutron beam interacts with a material, the neutrons are scattered by the nuclei in a pattern that depends on the crystal structure [370,371]. This means that, this scattering pattern can be analyzed to determine the crystallographic texture, lattice parameters, phase content, lattice strain, and atomic positions [370,373]. Neutrons, being uncharged particles, have high penetration depths, making neutron diffraction particularly useful for studying bulk properties and structural features of materials [371]. However, neutron diffraction may not be an appropriate selection for all β-Ti alloys. This is because, the very negative bound coherent scattering length of Ti (−3.438 fm) combined with the positive bound coherent scattering lengths of all chemical elements with atomic numbers higher than 27, may produce insufficient reflections for the scattering pattern to be analyzed. This is because the scattering cross section, i.e., the effective area interaction between the atomic nucleus and the passing neutron, is a function of the scattering length of the alloying elements [371].Weak nucleus-neutron interaction means small scattering signals. Examples of β-stabilizers with very positive scattering lengths are Fe (9.45 fm), Ni (10.3 fm), Mo (6.715 fm), Cu (7.718 fm), Nb (7.054 fm), Ta (6.91 fm), Hf (7.7 fm), etc. [370,374]. Consequently, there are limited reports of neutron diffraction on β-Ti alloys, such as Ti–10V–2Fe–3Al, Ti–12Mo, and Ti–15Mo [375–377], mainly alloyed with a limited quantity of elements with positive bound coherent scattering lengths or with other negative bound coherent scattering lengths elements (such as V of −0.3824 fm). Therefore, the rest of this section will be focused on synchrotron diffraction. Synchrotron diffraction utilizes highly energetic X-rays generated by synchrotron light sources. Specifically, these X-rays are produced when charged particles, typically electrons, are accelerated to nearly the speed of light and forced to travel in circular paths [372]. As a result, the resulting X-rays can be tuned to specific energies by a monochromator that selects the desired wavelength. Furthermore, synchrotron X-rays interact primarily with the electron cloud surrounding atoms, providing detailed information about the electron density and atomic arrangements in the material with a high spatial resolution. In addition, this technique is especially useful for studying surface and near-surface phenomena and can detect subtle changes in electron density, such as those associated with defects, dislocations, and phase transformations [378]. Moreover, the hutches, i.e., radiation shielding, do not require a high vacuum, allowing material testing under different environments [372]. Consequently, synchrotron diffraction, with its high resolution and rapid data acquisition (faster than neutron diffraction), is well-suited for studying fine-scale features, including thin films, nanomaterials, and interfaces [372,378]. Additionally, it can also be used to study dynamic processes such as phase transformations and deformation mechanisms in real time. However, its limited penetration depth and sensitivity to surface effects can be a disadvantage when studying bulk Fig. 32. ASTAR analysis of Ti–15Zr–15Mo alloy: a) virtual bright field (VBF) micrograph, b) Index image, c) Reliability, d) phases (red: β-Ti matrix, green: anatase TiO 2 , blue: rutile TiO 2 ), f) mixed image of phases and index [367]. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3583 properties. In particular, in-situ studies are one of the most powerful applications of synchrotron diffraction, allowing to observe real-time structural evolution under varying conditions. Several in-situ studies have focused on the β-to- α ’’ transformation of Ti–24Nb-0.5 N and Ti–24Nb-0.5O [379], Ti–17Nb–1Fe and Ti–17Nb–2Fe [380], Ti–17Nb [381], Ti–24Nb–4Zr–8Sn [382], Ti–30Zr–10Nb [383], Ti–12Mo [384], and Ti–5Al–5Mo–5V–3Cr β-Ti alloys [385], sometimes correlating them to the superelastic properties given by the strain-induced martensitic transformation [379,382], the high plasticity with ultra-low strain hardening [383], or with ω -phase generated during aging [380]. Furthermore, the dynamic and static mechanisms of microstructural restoration during hot working and heat treatments have been studied in Ti–6Al–4V alloy [386]. Furthermore, synchrotron X-ray computer tomography can provide imaging of bone/implant contact through their different X-ray absorption [387], as well as assist in fractography analyses of β-Ti alloys [388]. The latter one found a preferred preferential cracking starting at the β phase [388]. In addition, in-situ synchrotron diffraction has been employed to investigate the deformation mechanisms and texture evolution of β-Ti alloys under mechanical loading. These studies help in understanding the alloy’s behavior at the microstructural level, including the role of dislocations, twinning, and phase transformation in enhancing mechanical properties like strength and ductility. Future research may focus on real-time studies of β-Ti alloys in simulated physiological environments or with rigs simulating the stress state received by different human body parts. 5.7. Atom Probe Tomography (APT) Atom probe tomography is a nanoscale characterization technique that allows for the 3D mapping of chemical composition at the atomic scale. APT operates based on the principles of field evaporation, where atoms are ionized and then projected onto a detector to reconstruct a 3D image of the sample [389–391]. Specifically, the high electric field applied to a sharp needle-shaped specimen (tip radius <100 nm) causes surface atoms to evaporate as ions, which are then accelerated toward a detector [389]. By measuring the time of flight (TOF) of these ions, their mass-to-charge ratio can be determined, allowing the chemical identity and distribution of individual atoms to be reconstructed in 3D. Moreover, APT is particularly advantageous for detecting and analyzing trace elements, precipitates, and solute clusters that are often critical to the performance of these alloys. APT is sensible to chemical elements in the range of parts per million, including light elements such as hydrogen, carbon or lithium [391]. However, among the limitations of APT is the requirement for specimens to be in a needle-like shape, which can be difficult to prepare [389]. Additionally, the analysis is typically confined to a small volume of material, which might not be representative of the entire sample. In the context of biomedical Ti alloys, APT was first used to study a bone/titanium implant interface providing knowledge on the Caenriched chemical distribution required for osseointegration [392]. Furthermore, 4D (3D plus chemical composition) APT has also been employed to study the distribution of other elements and compounds like Ti, TiN, C, Mg, P, Na, and TiO at the bone/Ti-implant interface [393]. Additionally, APT studies have also been useful in determining that Fig. 33. STEM-ADF micrographs for the (a) Ti 80 Nb 15 Mg 5 , (b) Ti 77 Nb 15 Mg 8 , (c) Ti 69 Nb 15 Mg 16 , and (d) Ti 60 Nb 15 Mg 25 coatings showing their general microstructures [368]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3584 high laser pulse energies are not appropriate for studying nanoscale precipitation in Ti-alloys because they encourage the formation of complex molecular ions, cause spectral peak overlapping, and complicate data analyses [394]. Therefore, smaller laser pulse energies are recommended to study precipitation of, for example, ω phase in β-Ti alloys. This work also elucidated that aging time increases the Al and O ( α -stabilizers) enrichment at the ω -precipitate/β-matrix interface, decreasing the α nucleation barrier and acting as α nucleation site in the Ti–5Al–5Mo–5V–3Cr wt% (Ti-5553) β-Ti alloy [394]. Moreover, the formation of ω -phase embryos from compositional and structural instabilities β-Ti alloys was observed by APT and TEM, showing that β-to- ω phase transformation involves displacive and diffusional mechanisms [395]. Other works on β-Ti alloys have exploited APT to describe the chemical composition differences or evolution during thermomechanical processing to correlate them to different mechanical, chemical, or physical properties. For example, the chemical composition during precipitation evolution has been studied in an aged Ti-5553 β-Ti alloy, providing information on the varying chemical distribution among the isosurface, precipitate core, and matrix [396]. Similarly, the solute distribution at α -precipitates and β-matrix was also studied and correlated to the high strength of a hierarchical Ti–1Al–8V–5Fe (Ti185) alloy [397]. In addition, the microstructural evolution of biomedical Ti–24Nb–4Zr–8Sn (Ti-2448) alloy during aging has also been correlated to its elastic properties, which were negatively impacted by the dissolution of Nb-enriched domains [398]. Furthermore, a low-cost and Fig. 34. ACOM maps obtained by ASTAR for the (a) Ti 80 Nb 15 Mg 5 , (b) Ti 77 Nb 15 Mg 8 , (c) Ti 69 Nb 15 Mg 16 , and (d) Ti 60 Nb 15 Mg 25 coatings showing inverse pole figure (IPF) coloring for the z-axis [368]. Fig. 35. Ti–35Nb–7Zr–5Ta alloy aged at 400 ◦C for 4 h. (a) TEM bright field (BF) images showing the absence of a phase precipitation along the grain boundary. (b) Detail of the grain interior at the [1 1 1] β zone axis revealing diffraction contrast between the β (dark) and β ′ (bright) phases [369]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3585 feasible route for promoting the simultaneous enhancement of strengthening and ductility at β-Ti alloys has been proposed based on APT and synchrotron analyses, which is based on the combination of oxygen as an interstitial element and grain refinement [399]. Looking forward, the application of APT, complemented with other electron- or X-ray-based microscopies and spectroscopies, to the study of β-Ti alloys is expected to grow, particularly as advances in specimen preparation and data analysis continue to improve the accuracy and efficiency of the technique. A step forward has been made in Ti-alloys by Xe-based plasma focus ion beam (PFIB) to prepare APT specimens Fig. 36. Ti–35Nb–7Zr–5Ta alloy aged at 400 ◦C for 4 h. (a) HRTEM image of the region shown in Fig. 34b along the [1 1 1] β zone axis. (b) The Fourier transform (FT) reveals splitting of the (1 0 1) reflections (white arrows). (c) Schematic representation of the Fourier space map of the [1 1 1] β zone axis, considering separated phases β ′ with cubic (bcc) and β ′ with tetragonal (bct) structure obtained through simulation using JEMS© software [369]. Fig. 37. TEM micrographs in (a) bright field (BF) mode showing the ω -Ti nanoprecipitates dispersed in β-Ti grain in the P/M samples of the Ti35Nb10Ta alloy and (c) for the Ti35Nb10Ta3Fe alloy, both sintered at 1250 ◦C, together with the respective (b) and (d) SAD patterns of the β-Ti region with an orientation relationship [1 1 0] β-Ti //[1 1 2 0] ω between the β-Ti matrix and metastable ω phase for both alloys. ACOM map of the Ti35Nb10Ta alloy sintered at 1250 ◦C showing (e) virtual bright field (VBF) of the α +β region and (f) Phase map combined with a Virtual-BF image of the α precipitates (green) dispersed through the β-Ti (red) matrix [356]. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3586 without lift-out, reducing possible influence at the preparation process stage [400]. 6. Beta BCC high (HEA) and medium entropy alloys (MEA) in Ti–Nb–Zr–Ta–Mo system as multiprincipal alloys with low elastic modulus 6.1. High entropy or multiprincipal alloys for biomedical applications The concept of entropy of mixing (ΔS m ) classifies metallic alloys into three categories: low entropy (1 or 2 components), medium entropy (3 or 4 components), and high entropy (5 or more components) [401]. Such an approach is based on mixing together multiple elements with the same crystalline structure in relatively high (often equiatomic) contents. Briefly, the high entropy alloys (HEA) can be composed of at least 5 main elements, which can be equiatomic (20%at. for each element in a 5 components alloy) and varying percentage from 10 to 35% [402]. Authors suggested alternative names for high (HEA) or medium entropy alloys (MEA), such as multi-component alloys, compositionally complex alloys and multiprincipal-element alloys. By incorporating multiple principal components, we can increase the mixing entropy of the system ΔSm, with the aim of creating a single-phase alloy in solid solution (e.g. BCC in the case of β Ti alloys or composed by β-stabilizing elements: Ti, Nb, Hf, Mo, V, Zr, Ta and others), which reduces the number of phases predicted by the Gibbs rule, making the structure simpler and expected to present superior and optimized properties (effect of high entropy) [402]. In HEAs the atoms are bonded to different atoms that lead to chemical bonds among them asymmetric (atoms with different atomic size and crystalline structure) unlike conventional alloys, where the majority of atoms in the matrix form bonds with atoms of the same species. Such effect, results in a distortion in the crystalline structure of the alloy, promoting changes in its mechanical properties [402]. Since diffusion in HEAs is slower than in conventional or diluted alloys, there are few gaps for the substitutional diffusion process (supersaturated solid solutions), the nucleation and growth of others phases are delayed, resulting usually in one solid solution phase (BCC, FCC or HCP). On the other hand, the slower diffusion rate, can facilitate precipitates refinement, improves grain size control and increases the recrystallization temperature of the material resulting in grain refinement of the alloy [402,403]. Despite the fact that ΔS m is maximum when the elements are in equiatomic proportion, alloys with this composition generally have a high elastic modulus and, in the case of the Ti–Nb–Zr–Ta–Mo system, they have a high melting point (refractory alloys) since these alloys are based on Ta and Mo (when converting to weight%) due to the high atomic weight of these elements compared to the others. When just the optimization of mechanical properties of the alloy is the main goal the HEAs are very welcome, but in the case of metallic biomaterials for implant applications a lower elastic modulus (around E =60 GPa) is highly desirable in order to avoid stress shielding. Hereupon, there is an open window for the development of HEAs or MEAs alloys with a lower elastic modulus for biomedical application (BioHEA) and, therefore, equimassic BioHEAs can be an interesting alternative [404]. New equiatomic BCC HEA alloys have been reported [404], but just a few works in the literature reportthe elastic modulus E (GPa) obtained, probably because they are higher than that of pure titanium (E =110 GPa) and conventional Ti–6Al–4V alloy (E =100 GPa). These works focus, for the most part, on microstructure, corrosion resistance, tensile strength limit, improvement of ductility, application of empirical parameters (such as Mo eq , and B o /M d diagram), and, in their minority, on biological tests. The results are promising, since the properties are more suitable than those of commonly used alloys, but there is still plenty of room for further development. But there is a strong drawback for equiatomic alloys in Ti–Nb–Zr–Ta–Mo or similar systems which is the refractory character of such alloys and resulting in elastic modulus in the range E =116–153 GPa [358–364]. In order to circumvent it, one Fig. 38. TEM micrograph in (a) bright field (BF) mode of the nanocrystalline β Ti35Nb10Ta alloy after HPT processing and the respective (b) SAD ring pattern typical of nanostructured materials with aleatory orientation. ACOM analysis of the HPT sample of the Ti35Nb10Ta alloy with (c) virtual bright field (VBF) of the nanocrystalline α +β region and (d) IPF-z orientation image showing a general view of the nanostructure. Higher magnification ACOM analysis showing in detail (e) the virtual dark field (VDF) of the ω precipitates dispersed through the nanocrystalline β-Ti grains and (f) IPF-z orientation image showing smaller soft β nanograins and α precipitates as well [356]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3587 approach is to deal with medium entropy alloys (MEAs) or equimassic multiprincipal BCC alloys with 3 or more β-stabilizing alloying elements, looking for the optimal entropy range (1.1 <ΔS m <1.3R) Such uch equimassic ternary Ti–33Nb–33Zr, quaternary Ti–25Nb–25Zr–25Ta and quinary Ti–20Nb–20Zr–20Ta–20Mo (weight%) alloys show elastic modulus E =72–88 GPa and enhanced mechanical strength [368]. At present, the development of multicomponent, equimassic, stable β-Ti alloys with medium and high entropy of the systems Ti–33Nb–33Zr, Ti–25Nb–25Zr–25Ta, Ti–20Nb–20Zr–20Ta–20Mo (weight%), in addition to alloys from the Ti–40Nb-xZr system with high Zr fractions and based on electronic parameters Bo and Md seeking an optimal value of elastic modulus (between 40 and 60 GPa). The Ti–40Nb–40Zr alloy showed the lowest modulus of elasticity (E =42 GPa) in the rolled condition. The Ti–25Nb–25Zr–25Ta alloy showed better results in biological tests (adhesion and cell viability assay), and better combination of low modulus (E =72 GPa) and high mechanical strength, making it possible to reconcile these properties in β-Ti stable. Table 6 lists the chemical composition of β multiprincipal alloys as well as their elastic modulus, phases formed and processing conditions of some alloys for application with biomaterial. 7. Surface modification routes Various pre-treatment methods are required for Ti alloys in order to treat implant surfaces, prepare surfaces for coating processes, and modify their surfaces. These procedures have an impact on the physical, chemical, and electrochemical surface properties (roughness, wettability, and energy surface), which improves the biological environment’s response. Different kind of surface treatment or pre-treatments methods to create a functionalized surface have been investigated, such as: sandblasting, grinding, fine abrasive polishing, acid etching, electrochemical etching, anodization, and solvent degreasing [414]. Therefore, in this section, we’ll describe the methods adopted to improve the surface of various β-Ti alloys to enhance their biological Fig. 39. TEM bright field (BF) images of TiO 2 –based NTs formed from TFNS substrate anodized to 6 h/20 V and annealing at (a) 400 ◦C (b) 450 ◦C (c) 500 ◦C followed by corresponding HRTEM image (atomic scale) and respective SAD patterns in each annealing condition with amorphous halo and diffraction spots of anatase–rutile phases [366]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3588 properties and boost their corrosion resistance in biological environments. 7.1. Physical (texture and roughness) Independent of the implantation site, bone quantity, or bone quality, an ideal implant biomaterial should provide a surface that will not interfere with and may even promote the normal processes of bone healing. The procedures for modifying implant surfaces are shown in Fig. 40 and described in the section below. 7.1.1. Dual-acid etching and electrochemical treatment Acids are applied to metal surfaces during the acid etching process in order to not only clean them or disclose the grain and grain boundaries, but additionally to change how tough they are. This approach often makes use of a strong acid like hydrofluoric (HF), nitric (HNO 3 ), or sulfuric (H 2 SO 4 ), or a mixture of these acids. Many studies claimed that acid-etched surfaces had boosted bone production and cell adhesion, which improved the osseointegration mechanism [415–418]. The osseointegration process may be significantly impacted by changes in surface roughness and surface parameters such as surface charge, Table 6 Chemical composition of in atomic (at.%) and weight (wt.%) Beta Multiprincipal alloys or HEA/MEA and conventional alloys, elastic modulus, phases and condition of some alloys for application as biomaterial. Alloy Composition E (GPa) Phase(s) Condition ΔS m xR (J/K) Ref. Ti 34 Zr 52 Nb 14 (%at.) 76 βCo-sputtering 0.98R [401] TiZrNbTaMo (Equiatomic) – β Rapidly solidified (~2000K/s) 1.609R [404] TiZrNbTaMo (Equiatomic) 153 β +β ′ Copper mold cast 1.609R [405] TiZrNbTaHf (Equiatomic) 115 βCopper mold cast 1.609R [406] Ti 37.5 Nb 12.5 Zr 25 Hf 12.5 Ta 12.5 (%at.) 99 βCopper mold cast 1.494R [406] TiZrNbTaHf (Equiatomic) 80 βCopper mold cast 1.609R [407] TiZrNbTaHf (Equiatomic) 66 βCopper mold cast 1.609R [407] TiZrNbTaMo (Equiatomic) – β +β ′ Annealed 1.609R [408] TiZrNbTa (Equiatomic) – β Annealed 1.386R [408] Ti 2-X Zr 2-X Nb x Ta x Mo x (x ¼0.6) (%at.) – β Rapidly solidified (~2000K/s) 1.28R [409] Ti 35 Zr 27.5 Hf 27.5 Nb 5 Ta 5 (%at.) 79 β + α ″ Cold rolled and solution heat treatment. air cooled 1.38R [410] Ti–35Nb–2Ta–3Zr (%wt.) 49 β + α ″ Cold rolled (reduction of 99% transversal section) 1.10R [411] Ti 17.5 Zr 17 Nb 21.5 Ta 21.5 Mo 22.5 (%at.) 113 βHomogenized 1.60R [412] Ti 17.5 Zr 17 Nb 21.5 Ta 21.5 Mo 22.5 (%at.) 87 βAged at 500 ◦C/6 h 1.60R [412] Ti 29 Zr 12.5 Nb 28 Ta 29 Mn 1.5 (%at.) 76 βHomogenized 1.38R [412] Ti 29 Zr 12.5 Nb 28 Ta 29 Mn 1.5 (%at.) 80 βAged at 500 ◦C/6 h 1.38R [412] Ti–33Nb–33Zr (wt.%) 79 βAs-cast 1.04R [413] Ti–33Nb–33Zr (wt.%) 59 β + α ″ Hot rolled. air cooled 1.04R [413] Ti–40Nb–30Zr (wt.%) 83 βAs-cast 1.06R [413] Ti–40Nb–30Zr (wt.%) 65 β + α ″ Hot rolled. air cooled 1.06R [413] Ti–40Nb–40Zr (wt.%) 63 βAs-cast 1.10R [413] Ti–40Nb–40Zr (wt.%) 42 β + α ″ Hot rolled. air cooled 1.10R [413] Ti–25Nb–25Zr–25Ta (wt.%) 73 ВAs-cast 1.28R [413] Ti–25Nb–25Zr–25Ta (wt.%) 59 β + α ″ Hot rolled. air cooled 1.28R [413] Ti–20Nb–20Zr–20Ta–20Mo (wt.%) 88 ВAs-cast 1.52R [413] Stainless Steel 316L 215 γ (FCC) –[406] Co–Cr–Mo 252 FCC –[406] Ti-6AL-4V (wt%) 131 α +β – 0.39R [406] Fig. 40. Schematic techniques and process used for surface modification and treatment on β-Ti alloys. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3589 surface energy, and chemical composition caused by acid etching [415–418]. This strategy’s relevance also includes the homogeneous roughening it affords the substrate, independent of its sizes and shapes [419]. The durability of bone growth and resorption at the junction of bone implants is partly defined by the roughness of Ti [420]. According to Rama et al., an implant’s roughness may be improved by a nano-topography that permits bone ingrowth using acid etching [421]. According to a prior study, both the type and concentration of the acid employed affect the rate of etching [422]. The appropriateness of these acids for etching, meanwhile, was not verified since they needed additional testing, notably on the bone implant interface and torque elimination. Surface roughness increased once Ti samples were etched by H 2 SO 4 at various concentrations. Notably for biological purposes, concentrated H 2 SO 4 has been shown to be an efficient way to roughen surfaces [421]. In a manner analogous to acid etching, dual acid etching (DAE) is likewise capable of treating the surface by chemical or acid, whether in succession [423] or with the combination of both [424,425]. By using dual etching on a micro-rough surface, rapid osseointegration can be attained [426]. An assessment of surfaces that have been machined and those that have been treated with HF and HCl/H 2 SO 4 (DAE) has revealed that the acid-treated surface has superior osseointegration and is more resistant to reverse torque removal [426]. Yang et al. [427], implanted 15 implants into rabbits’ tibias in to investigate the surface roughened by the DAE. Amazingly, implants with roughened surfaces outperformed those with machined surfaces in terms of removal torque at 2, 4, and 8 weeks. A histomorphometric examination at the exact same time revealed that the development of peri implant bone occurred concurrently with a considerable increase in the bone-to-implant contact. Thus, the DAE can offer a surface that has a particular level of microroughness, promoting quick osseointegration [422]. Nonetheless, the procedure and acid type are very important factors in the acid etching treatment. An acid-etched Ti implant showed a surface topography that was comparable to that obtained from a sand-blasted large-grit acid etched (SLA) surface treatment, according to Juodzbalys et al. [428]. They observed that following etching with H 2 SO 4 and then HCl, the Ti sample displayed excellent surface roughness with 1–10 μ m micropits as opposed to an inferior surface microtexture by HCl and then H 2 SO 4 429 . In general, a chemical treatment is done to Ti alloys, specifically using alkali solutions or acids, not only to improve the surface roughness but also to alter the chemical makeup and to increase the surface energy or wettability [429]. It was revealed that the acid-treated surface of the implants demonstrated superior resistance to reverse torque removal and enhanced osseointegration as compared to the machined surface implants [426]. The density of the bone implant as well as the relative potential for osseointegration have been observed to be similarly affected by acid etching the zirconia implant’s surface [430]. Sadly, the usage of these acids also led to unintended consequences such porosities, which can range in size from 0.5 to 2 μ m [431, 432]. In some unspecified way, it is also thought that this procedure aids with the dental implant’s ability to support tissue ingrowth and cell surface contacts [432]. Reverse torque rotation was used to gauge the effectiveness of osseointegration or implant anchoring. More osseointegration was achieved as the torque rotation force value rose along with the bone-to-implant contact (BIC) [417]. By applying acid etching before coating, this value might be raised to 8.68 ±0.37 μ m [433]. However, the dual acid etching, which produced surfaces with roughness varying from 0.44 to 3.51 μ m, played a significant role in producing better surfaces [422,434,435]. Due to its high composition, amount, and concentration, DAE is typically superior to a single acid etching. Depending on the type, concentration, temperature, and duration of the acid, acid etching can affect dental implants differently. It is common to see an increase in surface roughness when the acid concentration increases [424]. Acid etching is among the most commonly utilized methods for modifying the surface of dental implants as it increases the surface’s bioactivity [429]. Implant bioactivity is influenced by surface chemistry, chemical composition, wettability, micro-topography, roughness, and overall increased interfacial surface area [436–438]. According to earlier studies the acid etched surface promotes greater cell attachment than implants without surface treatment [434,439,440]. Moreover, Park et al., showed that Ti surfaces subjected to dual acid etching result in an increase in fibrin and osteogenic cell adhesion [441]. For implants with acid etched surfaces compared to those with machined surfaces, Elias et al., observed a higher torque removal [442]. In the process of osseointegration, the surface topography and roughness of the implant have a significant influence on the first bone cellular response and the osseointegration process in general [422,443,444]. In order to form a fresh surface layer with a thicker coating of Ti oxide, an aggressive liquid is used on a Ti substrate to etch surface depressions. The goal of acid etching procedures is typically to change surface chemistry and micro-roughness and hence affect cell response [445]. The roughness of an implant is said to be improved by nano-topography that enables bone ingrowth by acid etching [420]. Etching is technically known as pickling, and it is typically used to remove oxide deposits to provide a smooth and even surface finish [446]. Chemical modification influences the chemical composition of the implant surface, resulting in unique interactions between molecules on the surface of cells that change not only those surface molecules but also the inner structure and function of cells [447]. These are carried out to eliminate oxide scales and contaminants in an effort to increase bioactivity and biocompatibility [448]. Chemical surface treatments have been used regularly on dental implant over the past ten years, among others. Comparing implants with Ti-plasma sprayed surfaces to those with sandblasted and acid-etched (SLA) Ti, it has been noted that the latter induces more osseous contact [449]. This is owing to the high surface roughness of SLA caused by grit blasting and acid etching, the latter of which increased the concentration of hydrogen beneath the surface and caused Ti hydride to develop [450]. By changing the treatment conditions, the SLA surface modification has also been improved [422]. These surface treatment techniques change the surface’s wettability and roughness, which changes the parameters of cell adhesion and proliferation. For dental implants, osseointegration at the bone-implant contact is a requirement to achieve and preserve their long-term stability. Understanding the biological mechanisms behind the interaction between osteoblasts and biomaterials is crucial in this context because osseointegration is the desired outcome. Titanium ’s surface can be modified for use in biological applications by etching it in concentrated acid, which seems to be an appealing technique. Studies in this field are, however, hardly reported in biomaterials science. A rougher Ti surface was created by etching in concentrated sulfuric acid than by using HCl, H 3 PO 4 , HF, or HNO 3 according to the prior study [449]. In the research done by Lario et al., Ti–6Al–4V ELI and Ti35Nb10- Ta1.5Fe (β-Ti alloy) were used to test the effects of the acid-etching procedure in a two-step acid treatment, samples were etched [451]. The outcomes showed that the two-step acid treatment altered the alloy’s morphology, increased its surface area, and modified its chemical composition. During the acid etching procedure, two distinct zones were observed in the Ti–35Nb–10Ta-1.5Fe alloy: an area with lower roughness parameters and other with greater mean roughness due to its lower chemical resistance. Acid-etching procedures are used to produce a microroughness topography on the implant surface [422,452,453]. The surface area and roughness of implants can be adjusted in order to enhance bone regeneration and accelerate short- and midterm healing times [454,455]. Surface topography can be changed by electrochemical procedures, plasma-sprayed coatings, grit blasting, acid etching, and other A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3590 Table 7 Orthopaedic alloys designed for use in orthopaedic implants and their mechanical characteristics (Elastic modulus (E), yield strength (YS), and ultimate tensile strength (UTS)). ALLOY DESIGNATION E (GPa) YS (MPa) UTS (MPa) REFERENCE Ti–17Nb 84 640 768 [562] Ti–27.5 Nb (at. %) 70 800 820 [563] Ti–35Nb (wt.%) 75 343 555 [564] Ti-40.5Nb 75.1 –447 [565] Ti–42Nb (wt.%) 60.51 674 683 [566] Ti–45Nb (wt.%) 63 –986 [567] Ti–16Nb–10Zr 70 485 520 [568] Ti–13Nb–13Zr 79–84 973–1037 [569] Ti–13Nb–13Zr 79 900 1030 [570] Ti–34Nb–25Zr 62 810 839 [565] Ti–28Nb–35.4Zr (wt.%) 63 611 633 [571] Ti–31Nb-1.0Fe 81 –477 [572] Ti–27Nb-1.5Fe 94 –496 [573] Ti–23Nb-2.0Fe 95 –604 [572] Ti–19Nb-2.5Fe 90 –672 [573] Ti–15Nb-3.0Fe 94 –695 [573] Ti–11Nb-3.5Fe 97 –715 [573] Ti-38.1Nb-2.1Mo 54.5 –663 [574] Ti-39.9Nb-3.2Mo 67 –535 [574] Ti-39.3Nb-6.3Mo 63.6 –621 [574] Ti-38.7Nb-9.2Mo 55.7 –634 [574] Ti–32Nb–2Sn (wt.%) 93 870 930 [575] Ti–32Nb–2Sn (wt.%) 82 900 1070 [575] Ti–37Nb–6Sn (wt.%) 66 –891 [576] Ti-30.8Nb-9.8Sn 61 –435 [577] Ti-25.4Nb-10.1Sn 62 –611 [577] Ti-21.2Nb-10.4Sn 75 –772 [577] Ti-16.6Nb-10.6Sn 80 –787 [577] Ti–25Nb–16Hf (wt.%) 42 709 –[578] Ti–29Nb–13Ta–6Sn 65 523 –[579] Ti–29Nb–13Ta–2Sn 48 469 –[579] Ti–29Nb–13Ta–4.6Sn 78 513 –[579] Ti–30Nb–2Ta–3Zr 57.1 –663 [580] Ti–25Nb–2Ta–3Zr 75.2 –864 [581] Ti–30Nb–10Ta–5Zr 58 558 689 [582] Ti–30Nb–10Ta–5Zr 66.9 804 [583] Ti–29Nb–13Ta–4.6Zr 60 800 1100 [584] Ti–29Nb–13Ta–4.6Zr >80 >1000 >1001 [585] Ti–29Nb–13Ta–4.6Zr 60 549 [586] Ti–29Nb–13Ta–4.6Zr 63 400 [579] Ti–29Nb–13Ta–4.6Zr (wt.%) ~ 80 ~ 1000 ~ 1100 [571] Ti–29Nb–13Ta–4.6Zr (wt.%) ~ 90 1200 –[587] Ti–29Nb–13Ta–4.6Zr (wt.%) ~ 70 ~ 900 ~ 1000 [588] Ti–29Nb–13Ta–4.6Zr (wt.%) 86 –1375 [589] Ti–29Nb–13Ta–4.6Zr (wt.%) 68 ~400 ~600 [590] Ti–29Nb–13Ta–4.6Zr (wt.%) 70 900 1000 [588] Ti–29Nb–13Ta–4.6Zr (wt.%) 50 490 740 [591] Ti–35.3Nb–5.1Ta–7.1Zr 55 547 597 [592] Ti–40Nb–2Ta–3Zr 57.8 –589 [580] Ti–35Nb–2Ta–3Zr 52 –431 [580] Ti–35Nb–2Ta–3Zr (wt.%) 50 –896 [411] Ti–35Nb–5Ta–7Zr (tnzt) 55 530 590 [570] Ti–24Nb–4Zr–7.9Sn 46 700 –[579] Ti–24Nb–4Zr–7.6Sn (wt.%) ~75 ~1000 –[593] Ti–24Nb–4Zr–7.9Sn (wt.%) 63.5 629 705 [594] Ti–24Nb–4Zr–8Sn (wt.%) ~ 57 –1050 [595] Ti–24Nb–4Zr–8Sn (wt.%) 53 563 665 [106] Ti–24Nb–4Zr–8Sn (wt.%) 56 –1150 [596] Ti–35Nb–7Zr–5Ta (wt.%) 93 –2120 [597] Ti–35Nb–7Zr–5Ta 55 823 [579] Ti–35Nb–3Zr–2Ta (wt.%) 59.3 389.9 765.4 [598] Ti–25Nb–3Zr–3Mo–2Sn 80 946 1220 [599] Ti–25Nb–3Zr–3Mo–2Sn (wt.%) 75 770 >800 [600] Ti–19Nb–1.5Mo–4Zr– 8Sn (wt.%) 55 823 973 [601] Ti–29Nb–13Ta–4.6Zr–0.7 (wt.%) 80 1050 1100 [602] Ti–32.5Nb–6.8Zr–2.7Sn (wt.%) 47 –883 [603] Ti–35Nb–5Ta–7Zr–0.4O 66 1010 –[604] Ti–35Nb–5Ta–7Zr-0.4 (TNZTO) 66 976 1010 [570] Ti–35Nb–9Zr–6Mo–4Sn (wt.%) 65 925 950 [605] Tiadyne 1610 81 736 851 [98] Ti65.5Nb22.3Zr4.6Ta1.6Fe6 (at. %) 52 2475 2650 [606] Ti–10Mo (wt.%) 47 690 731 [607] Ti–12Mo–3Nb 105 450 –[608] (continued on next page) A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3597 Table 7 (continued) ALLOY DESIGNATION E (GPa) YS (MPa) UTS (MPa) REFERENCE Ti–15Mo 78 544 874 [98] Ti–15Mo (aged) 78 448 [579] Ti–4.7Mo–4.5Fe (at. %) 124 –1390 [572] Ti–15Mo–3Nb-0.3O (21SRx) 82 1020 1020 [570] Ti–8Mo–4Nb–2Zr 35.4 798 –[609] Ti–8Mo–4Nb–5Zr 52 483 –[609] Ti–8Mo–5Nb–3Zr 69 947 –[609] Ti–8Mo–6Nb–4Zr 72 899 –[609] Ti–15Mo-2.8Nb–3Al 82 771 812 [570] Ti–15Mo-2.8Nb–3Al 100 1215 1300 [570] Ti–12Mo–6Zr–2Fe 74–85 1030–1100 –[610] Ti–15Mo–5Zr–3Al 75 838–1060 –[579] Ti–15Mo–5Zr–3Al 80 532 –[608] Ti–15Mo–5Zr–3Al 75 870–968 882–975 [570] Ti–15Mo–5Zr–3Al 88–113 1087–1284 1099–1312 [570] Ti–12Mo–6Zr–2Fe (TMZF) 74–85 1000–1060 1060–1100 [570] Ti-4.8Mo-5.9Sn-18.5Nb-9.1Zr 72 –519 [572] Ti-4.9Mo–6Sn-14.2Nb-9.3Zr 56 –701 [577] Ti-5.2Mo-6.5Sn–10Zr-10.2Nb 43 –577 [436] Ti-5.3Mo-6.5Sn-10.2Nb–10Zr 48 –715 [580] Ti-5.3Mo-6.5Sn-15.2Nb–5Zr 55 –692 [572] Ti–5Mo-6.2Sn-9.7Nb-9.5Zr 54 –715 [572] Ti–15Mo–6zr– (1–4) Cr (wt.%) 86 1160 –[611] (Ti–23.72Nb4.83Zr1.74Ta)100-xSi 37 1296 3263 [136] [(Mo0.5sn0.5)- (Ti13Zr1)] (Ti0.5Nb1.5) 56 566 701 [(Mo0.5Sn0.5) -(Ti13Zr0.5Nb0.5)] Nb1 55 533 692 [(Mo0.5Sn0.5) -(Ti13Zr1)] Nb1 48 478 715 [(Mo0.5Sn0.5) -(Ti13Zr1)] Nb2 72 499 715 [(Mo0.5Sn0.5)-Ti14] (Ti0.5Zr0.5Nb1) 59 275 762 [(Mo0.5Sn0.5)-Ti14] (Zr1Nb1) 54 235 715 Ti–30Zr–1Cr–5Mo 69.1 –760 [581] Ti–30Zr–2Cr–4Mo 64.2 –780 [581] Ti–30Zr–3Cr–3Mo 68.5 –800 [581] Ti–30Zr–4Cr 69.4 –700 [574] Ti–30Zr–5Cr 66.6 –900 [581] Ti–32Zr–30Nb 65 782 794 [565] Ti–35.4Zr–28Nb 64 729 755 [565] Ti–40.7Zr–24.8Nb 63 682 704 [565] Ti–50Ta 88 380 –[608] Ti–25Ta–25Nb (wt.%) 92 2142 2423 [612] Ti–50Ta (wt.%) 75.8 882 924 [613] Ti–5Al–2.5Fe 110 1020 –[579] Ti–5Al—2.5Fe 110 914 1033 [570] Ti-6-4 Eli 110 875 965 [98] Ti–6Al–4V 113 1100 –[608] Ti–6Al–4V 110 850–900 960–970 [570] Ti–6Al–7Nb 105 900 –[609] Ti–6Al–7Nb (protasul-100) 105 921 1024 [570] Ti–6Al–7Nb (soln.) 110 110–1050 –[609] CP-Ti 105 692 785 [570] Co–Cr–Mo 200–230 275–1585 600–1795 [570] CP-Ti (grade 1–4) 102.7 170–485 –[614] Stainless steel 316L 200 170–750 465–950 [570] Cp-Ti 109 607 758 [565] Ti Grade IV 104.1 485 550 [592] Porous titanium alloys Ti-20.5Nb-5.6Zr 47.1 % 3.4 –[276] [615] Ti-20.5Nb-5.6Zr 56.7 % 2.5 –[158] [615] Ti-20.5Nb-5.6Zr 65.2 % 1.5 –[74] [615] Ti–10Nb–10Zr 42% 21 368 –[616] Ti–10Nb–10Zr 50% 7.9 235 –[616] Ti–10Nb–10Zr 59% 5.6 137 –[616] Ti–10Nb–10Zr 69% 3.9 67 –[616] Ti–10Nb–10Zr 74% 1.6 27 –[616] Ti–35Nb 47.9% 7.7 ––[617] Ti–13Nb–13Zr 39.9% 11.7 ––[617] Ti–35Nb–7Zr–5Ta 42.5% 9.2 ––[617] Ti–35Nb–5Ta–7Zr 40% 0.8 ––[618] Ti–34Nb–6Sn 26% 14.1 245 590 [60] Ti–34Nb–6Sn 22% 18.8 318 660 [60] Ti-13-Nb-30Ta–2Mn 70.2 % 9.8 41.6 –[619] Ti–13Nb–30Ta–2Mn 69.2 % 10.1 52.7 –[619] Ti–13Nb–30Ta–4Mn 68.5 % 9.2 67.5 –[619] Ti–13Nb–30Ta–4Mn 68.7 % 8.3 42.3 –[619] Ti–13Nb–30Ta–6Mn 71.2 % 11.8 30.4 –[619] Ti–13Nb–30Ta–6Mn 70.8 % 9.8 43.3 –[619] (continued on next page) A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3598 four passes, with finely ground grains and homogenous microstructures. In addition, a low elastic modulus of around 59 GPa was attained [598]. To enhance the hardness of the TNTZ alloy, Kheradmandfard et al. [630] treated it using the ultrasonic nanocrystal surface modification (UNSM) technique. As a result, the surface hardness increased from 195 HV to 385 HV; nevertheless, as one moves away from the surface, the rate at which surface hardness increases diminishes. Both grain refinement as well as a working hardening phenomenon are attributed for the UNSM’s increase in hardness. Chauhan et al. [631] modified the surface microstructure of VT3-1 α -β-Ti alloy by laser surface heat treatment (LST). LST was performed using a variety of laser energies (100–250 W) and scanning velocities (150–500 mm/min). A laser-affected zone has been analyzed using Electron backscatter diffraction (EBSD), which has revealed a unique microstructure that changes with laser power and scanning speed. A nearly β-complete -containing microstructure was formed at the surface with a high hardness level of 750 HV at a lower scanning speed (150 mm/min) and higher laser power (200 W). Hardness is improved by obtaining dendritic microstructures and uniform elemental distributions. On the other hand, in their investigation about re-melting on Ti–35Nb–2Ta–3Zr, Zhang et al. [632] also employed a laser. Laser re-melting of the material resulted in a change from an equiaxed crystal to a needle-like structure, which caused an increase in the average microhardness value from 165 HV to 264 HV. Moreover, the ECAP technology can generate sizable homogeneous plastic stresses to improve mechanical behaviour and purify grains. Successful plasma electrolytic oxidation (PEO) coating development on the Ti–39Nb–6Zr alloy was achieved by Chen et al. [633], investigated the alloy’s electrochemical corrosion and wear behavior prior to as well as following modification in PBS solution. The findings demonstrated that the Ti–39Nb–6Zr alloy in the PBS solution greatly enhanced its corrosion and wear resistance following PEO surface treatment. 8.1.3. Fatigue Fatigue fracture initiation is typically observed to take place at stress concentration locations and manufacturing flaws in the metallic component, including holes, fillets, welds, notches, pits, and surface defects from machining procedures [634]. Fatigue could also initiate in microstructural imperfections, such as grain boundaries and non-metallic inclusions [635]. Cyclical loading is a known factor in the failure of orthopaedic implants over the course of their lifespans [636]. Through the phenomena of corrosion-fatigue, the corrosive physiological environment intensifies the damage done by fatigue. Micro-motions at implant interfaces may result in fretting fatigue, which in turn may lead to implant failure [637]. Biomedical Ti alloys must have fatigue strengths at least equal to and ideally higher than Ti–6Al–4V (700 MPa) [638]. It is crucial to first characterize and then appropriately enhance the fatigue resistance of Ti alloys. The alloy Ti–29Nb–13Ta-4.6Zr has been extensively studied for its fatigue properties [602,639]. In general, single β phase microstructures lack the necessary fatigue strength for load-bearing orthopaedic applications. In single β phase conditions, the fatigue strengths of alloys such Ti–29Nb–13Ta-4.6Zr [602], Ti–24Nb–4Zr– 8Sn [596], Ti–35Nb-2.5Sn [640], and Ti–35Nb [564] are, respectively, 320, 375, 225, and 275 MPa. These fatigue strength values are undoubtedly lower than those for Ti–6Al–4V [638]. Precipitation strengthening, plastic deformation, particle reinforcing, and adjustment of the O content are methods to boost the fatigue strength of Ti alloys. To enhance strength while keeping a lower modulus, Nakai et al. [641] regulated the precipitation of the phase during the aging of cold rolled Ti–29Nb–13Ta-4.6Zr. Yield strength increased from 700 to 900 MPa and fatigue strength increased concurrently from 350 to 600 MPa after 10.8 ks of aging at 300 ◦C [641]. The ductility of the material reduced to almost zero when aged for longer than 10.8 years, so fatigue strength tests were not conducted. The improvement in both tensile and fatigue strength was achieved while maintaining a modulus of less than 80 GPa. This was due to the texture of phase [641]. Achieving high fatigue strength requires both a boost in strength after precipitation and considerable ductility. A cold-rolled Ti–29Nb–13Ta-4.6Zr with the maximum strength was obtained by aging it at 325 ◦C. Ti–29Nb–13Ta-4.6Zr aged at 450 ◦C produced a steel with a lower degree of strengthening but acceptable ductility [585]. As a result, the fatigue strength after aging at 450 ◦C was substantially higher than the fatigue strength after aging at 325 ◦C [585]. Investigations have also been made into the impact of corrosive body environments on the fatigue life of Ti alloys [582,585,640]. The high cycle fatigue strength of Ti alloys is reported to be comparable in both air and corrosive environments that mimic the conditions of the human body, as described in a few reports [582,585,640]. It is most likely due to Ti alloys’ excellent corrosion resistance. In contrast to the fatigue response to an airy environment, the fatigue response to an environment with corrosive compounds is dependent on the processing history. After water quenching and furnace cooling, the fatigue strength of Table 7 (continued) ALLOY DESIGNATION E (GPa) YS (MPa) UTS (MPa) REFERENCE CP-Ti 39% 12.01 ––[617] Ti–6Al–4V 40.6% 12.8 ––[617] (−) means no available data. Fig. 44. A plot showing the modulus of elasticity, E (GPa) versus Mo equivalency (Moeq). Fig. 45. A plot showing the modulus of elasticity, E (GPa) versus Yield Strength, YS (MPa) [620–622]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3599 solution-treated Ti–35Nb alloy was similar in air and 0.9% NaCl (275 MPa), but lower in NaCl (150 MPa) than in air (225 MPa)) [564]. After furnace cooling, the fatigue strength was lower than it had been after water quenching, despite the initial strength being higher. In Fig. 46, the limit and sensitivity ( σ f /UTS) to fatigue of some titanium alloys is showed. As is expected, regardless of the type of alloy, test with R =0.1 has a higher fatigue life than test with R = − 1, and has lower fatigue sensitivity (higher values). In comparison to biphasic alloys, fatigue resistance is generally improved by ternary and quaternary titanium alloys, with oxygen control being a key component of the alloys. In addition, Ti alloys’ resistance to fatigue is decreased by the presence of pores, as shown in Fig. 47. Although pores are thought to be places where cracks form, they also function as a mechanism for making fractures more difficult to break (R-curve - "stop-hole"). Given that the implant’s resistance to wear, it will rise once the pores have been filled with the bone tissue that grows inward, therefore, the combination of a sufficient number of pores could be an intriguing solution for fatigue limit improvement. 8.1.4. Wear It is assumed that a material’s wear resistance will grow as hardness increases. The wear resistance of three β-Ti alloys with the compositions Ti–24Nb–4Zr–8Sn, Ti–29Nb–13Ta–7Zr, and Ti–35Nb–4Sn was studied by Chapala et al. [644]. Due to the alloy’s high-volume fraction of the high hardness phase, Ti–35Nb–4Sn had the best wear resistance. Likewise, due to the earlier alloy’s higher hardness, Ti–15Zr-7.5Mo has better wear resistance than Ti–15Zr–15Mo [645]. Wear occurs to implants under physiologically wet environments. Because of this, tribological characteristics are frequently assessed in simulated body fluid. A variety of studies have been conducted on tribo-corrosion, or pure wear behaviour in moist environments, while other studies evaluate the combined impact of wear and corrosion. Wet environments have the capability of speeding up as well as slowing down the wear rate of an alloy in a particular tribological system [646]. According to Yang et al. [547], in contrast to dry conditions, the wear rate of Ti–12Mo–6Zr–2Fe increased while that of Ti–6Al–4V decreased. In contrast, compared to dry conditions, the lubricating function of SBF decreased the wear of Ti–6Al–4V [547]. The wear rate of a Ti–29Nb–13Ta-4.6Zr disc against a ball of the same alloy was found to be higher in Ringer’s solution than in a dry environment in another study [646]. Ringer’s solution’s presence made it possible to get rid of worn particles that was acting as a lubricant and speeding up wear. On the other hand, because the debris generated in Ringer’s solution had no lubricating effect, the wear rate of a Ti–6Al–4V disc against a Ti–6Al–4V ball was lowered in comparison to air [646]. In comparison to Ti–6Al–4V, β-Ti alloys may have a reduced, equal, or greater wear rate. The tribological system has a significant influence on the wear response. Alumina or zirconia mating balls resulted in a higher and lower wear rate for Ti–29Nb–13Ta-4.6Zr, respectively, than Ti–6Al–4V [121]. Generally, β-Ti alloys’ wear performance can be compared to Ti–6Al–4V. It is noteworthy that there is no literature that discusses the wear of several new compositions of titanium alloys. The majority of the time, medicinal Ti alloys exhibit poor wear characteristics and are easily worn in the human body [647,648]. Large wear loss and high friction coefficients are regarded as negative aspects of tribological behaviour. Despite having a high corrosion resistance, the Ti alloys’ protective oxide layer may become degraded during an infection or when in touch with biomolecules or physiological fluids that contain low oxygen. Several β-Ti alloys do not show a consistent trend in their wear rates when compared to Ti–6Al–4V. The tribological system has a major impact on how well β-Ti alloys resist wear [649]. The high volume fraction of the high hardness phase in the alloy is what gives the material its greatest wear resistance [644]. Clearly, the wear performance of Ti-based alloys remains an unresolved research issue, which offers numerous opportunities for improving its performance under, wear. 8.2. Corrosion resistance Throughout their use in the human body, orthopaedic implants corrode. The physiological liquid which surrounds implants is a saline solution that contains Na + , Mg 2+ , Cl − , SO 4 2− and HCO 3 2− ions [650]. Several factors contribute to the ability of metallic implants to resist corrosion. One such factor is the formation of a passive film, a stable, compact, and continuous oxide surface film that protects the underlying bare metal surface from corrosion [651]. Nevertheless, the passive layer may become dissolved locally, particularly by chloride ions, resulting in pits that spread rapidly and pitting corrosion. Implant composition should be carefully engineered and, as a result, must be highly corrosion resistant, to minimize adverse body reactions. When metal ions are liberated from an implant, they can elicit allergic reactions, cause localised tissue reactions, and raise the concentration of metal ions in bodily fluids, which can lead to toxicity [652]. This ion release is influenced by both the alloy’s rate of corrosion and the solubility of the earliest corrosion products. To evaluate the long-term impacts of the porosity on the material’s behaviour in physiological fluids, it is crucial to consider the corrosion process. Within seconds to minutes following implantation, the biomaterial’s surface becomes covered in plasma/- tissue proteins. Monocytes attach to the protein-covered surface of the implant four to 8 h later and differentiate to macrophages, which generate a variety of growth factors, proinflammatory cytokines, and cytokines that promote fibrosis [29]. Additionally, damaging reactive oxygen species (ROS) are produced by surface-activated macrophages that might contribute to oxidative damage to the implant. Along with the locally elevated ROS content, neutrophils also produce lactic acid, Fig. 46. Limit and sensitivity ( σ f /UTS) to fatigue of C.P. titanium and their alloys. Note: estimated values (using Goodman Equation) [570,585,593,602, 639,640,642]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3600 which reduces the pH of the surrounding microenvironment [653]. All these factors make metals more prone to corrosion. The composition of the implant is the core problem regarding corrosion resistance [654, 655]. The question is how common alloying ingredients are employed in Ti alloys. According to Okazaki et al., adding Zr, Nb, Ta, and a tiny amount of Pd reduces the critical current density for passivation Ic. An upward trend in the Ic value was seen when Sn was employed in place of Zr in Ti alloys [656]. The obtained Ic values for Ti–Zr alloys were significantly lower than for Ti–6Al–4V ELI, although they were higher for Ti–Sn alloys. Whenever the final corrosion potential (E corr ) value drops along with the metal’s porosity, porous metals become susceptible to localised corrosion [29]. It was hypothesised that relatively small, isolated pores produced by denser compacts would favor the trapping of electrolyte species and the depletion of oxygen, both of which are essential for maintaining a stable titanium oxide layer [29]. Ti alloys are the good choice for orthopaedic applications owing to their excellent corrosion resistance. Ti is corrosion resistant in a wide range of environments due to the stable thin TiO 2 layer that is formed on its surface [121]. β-Ti alloys have a remarkable resistance to corrosion, as shown by their low corrosion rates, which vary from 10 −9 to 10 −7 A/cm 2 [657]. Comparing a β-Ti alloy’s corrosion resistance to Ti–6Al–4V or CP-Ti is a common way to evaluate its corrosion resistance. The corrosion resistance of β-Ti alloys may be superior to CP-Ti or Ti–6Al–4V [55], equal to Ref. [658], or worse to Ref. [659]. The rates of corrosion are influenced by several variables, including alloying elements, microstructure, and corrosive media. The relationship between alloying elements and corrosion resistance may be explained by their impact characteristics of the oxide layer in β-Ti alloys. The mechanism by which Nb improves the corrosion resistance of Ti–6Al–6Nb alloy is one such leading example [660]. The presence of lower valence Ti +3 ions cause oxygen vacancies in the oxide, which are diminished by the higher valence Nb +5 ions [660]. Thus, as the number of defects in the oxide layer decreases, corrosion resistance also increases [660]. The corrosion rate may also change due to phase transformation during thermomechanical treatment. After subjecting the Ti–32Nb–2Sn alloy to a surface mechanical attrition procedure, researchers noticed an increase in the corrosion rate because the microstructure’s heterogeneity had grown [661]. Additionally, the corrosion media significantly affects how quickly alloys corrode [662]. This brings attention to a crucial point when assessing the corrosion rate of β-Ti alloys. It is important to highlight that the temperature used for the corrosion testing in simulated environments is either room temperature or that of a human being. Furthermore, the implants will encounter a different environment in-vivo than they will in an artificial environment that is simulated with a variety of electrolytes, such as simulated body fluid (SBF), Ringer’s solution, phosphate buffer saline (PBS), Hank’s balanced salt solution (HBSS). As a result, in-vivo animal tests are necessary to accurately assess the corrosion resistance of recently developed Ti alloys. In this regard, the chemical composition of the alloy and the presence of pores have a significant impact on the corrosion resistance of Ti implants. Since allergic, toxic/cytotoxic, or carcinogenic species (such as Ni, Co, Cr, V, or Al) may be released into the body during corrosion processes, this is quite concerning in the case of metal implants and devices for in-vivo service [644]. In contrast to Ti–6Al–4V alloys, Ti–35Nb–7Zr–5Ta alloy does not contain any of these elements. Additionally, earlier research shows that Ti–Al–V and TNZT alloys display various levels of corrosion resistance in simulated fluids, including Hank’s solution [663], Ringer’s solution [664,665]. However, as they increase the surface area exposed to an aggressive environment, alloy porosity content, shape, and distribution also affect the corrosion properties of porous Ti and its alloys [666,667]. For instance, Xu et al. noticed that when Ti–10Mo alloy manufactured using powder metallurgy and the space holder technique was exposed to naturally aerated phosphate buffered saline solution, the corrosion rate increased exponentially from 1.6 g/m 2 day to 17.1 g/m 2 day when the porosity increased from 2.8 to 66.9% [667]. Additionally, crevice corrosion is driven on by high and interconnected porosities [668]. However, from a corrosion and protection perspective, an increase in porosity does not always result in a worsening of the implant’s corrosion behaviour. Contrarily, there is broad agreement that increased porosity can be quite advantageous since it dramatically improves the infiltration and adherence of various coatings, which, in turn, significantly improves their corrosion protection compared to their totally dense counterparts [669]. 8.2.1. Electrochemical behavior Corrosion is a degradation process involving multiple simultaneous electrochemical reactions between oxidized (anode) and reduced Fig. 47. Fatigue behavior of porous titanium alloys [337,643]. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3601 (cathode) regions placed in contact by a conductive medium (electrolyte) [670]. From section 2(Table 1), Ti matrix possesses higher electronegativity than other more active metals such as Mg and Al. From the above, CP-Ti has been widely applied due to its high corrosion resistance. However, the mechanical characteristics of CP-Ti are not suitable for the high-impact or cyclic load required for articulation implants. The requirement for mechanically resistant, low-elastic-modulus, and non-toxic alloy systems is an immediate worldwide priority to attend to the increasing demand for implant procedures [671,672]. As seen in section 2, the corrosion susceptibility of Ti alloys is strongly influenced by their chemical composition and the natural oxide passivation layer on Ti. A TiO 2 surface passive layer forms when Ti encounters the air, improving the corrosion resistance [70]. However, the surface layer can be broken by atomic diffusion of alloying elements, like V [70]. Chemical heterogeneities might also trigger the formation of micro-galvanic couplings, i.e., regions with high electrochemical potential disparities. Micro-galvanic couplings encourage the breakage of the passive film [82] and corrosion losses. Since the passive oxide film limits the ionic and electronic migration across the oxide/body fluid interface, its breakage allows corrosion on the implant. Besides chemical composition, the electrochemical performance is also determined by the microstructural features and the service environment, which are the focus of this section. 8.2.2. Microstructural effect in corrosion behavior Pitting, also known as localized corrosion, may lead to premature implant failure. Defects in the microstructure, such as interfaces (grain boundaries, multiphasic borders, different defects transitions, etc.), chemical heterogeneities (inclusions, intermetallic particles, segregates, etc), pores, among others, are preferred pitting nucleation sites in Ti alloys [673–675]. The preferred attack is based on the higher energy and cation mobility at defects compared to the homogeneous inner grain regions [676]. In the case of porous structures, the pores reduce the oxygen supply, allow Cl − adsorption, and trigger higher electronic donor density [673]. The previous conditions accelerate the dissolution of the oxide film in the porous materials. The negative impact of pores on the corrosion resistance of NiTi and Ti–39Nb alloys has been reported [677,678]. Regarding interfaces in metallic materials, intergranular corrosion occurs preferentially at high-energy grain boundaries (HAGBs) than at low-energy grain boundaries (LAGBs) [679]. However, HAGBs also promote the atomic diffusion to form thick surface passive layers [673]. Considering that twin are less energetic than HAGBs [680], it may be interesting to study the effect of high-dense twined microstructures on the corrosion resistance compared to that in high-dense HAGBs Ti alloys. These systematic studies remain a future opportunity for research. In general, corrosion resistance is associated to coarse (micrometric grained) alloys due to their lower defects (grain boundaries) density compared to nanostructured or ultrafine materials, i.e., coarse materials have less preferred sites for the corrosion [681]. However, low density of grain boundaries also decreases the sites for adherence of the surface passive layer compared to fine-grained systems [682–684]. The high density of grain boundaries and triple junctions (intersection of leastwise three different neighboring grains in nanostructured alloys results in better atomic diffusion and well-adhered passive layers [685]. Thus, nanostructured materials have more uniform and compact surface passive surface layers compared to those in coarse-grained alloys [684]. As a result, nano-grained materials have a low susceptibility to localized corrosion. One example is the wrought and EBM Ti–6Al–4V alloys, where the EBM showed better oxide film stability due to higher grain boundary density [686]. However, the high density of defects in fine-grained materials increases the surface energy and the chemical activity in non-passivating media [684].Therefore, grain size may have a dual effect on the electrochemical behavior of Ti alloys, being strongly dependent on the severity of the corrosive environment [683]. Nanostructured alloys may show low corrosion susceptibility in passivating environments, and increase it in non-passivating media [683]. Regarding multiphasic microstructures, the coexistence of different phases, such as HCP- α , BCC-β, HCP- α ′ , orthorhombicα ’’, or HCP- ω , encourage electrochemical potential disparities. The reason for such electrochemical disparities is the different chemical concentrations along the phases. It is expected that the β- and α -phases will have a higher concentration of β- and α -stabilizers (described in section 2), respectively. As a result, each phase will have different electrochemical potential and corrosion performance [687,688]. The galvanic activity might also trigger stress corrosion cracking (SCC) in Ti alloys, where the cathodic and anodic regions will be determined by the chemical composition of the present phases. As an example of the above, the β-Ti phase has been reported as more susceptible to corrosion than the α -Ti phase of V-bearing Ti alloys [70]. However, the β-phase was more electrochemically resistant than the α -Ti phase when Nb, Mo, and Zr β-stabilizers were present [77]. Besides multiphasic microstructures, galvanic couplings can also be formed by impurities on the microstructure. The volume fraction and distribution of the cathodic and anodic regions through the microstructures also impact the corrosion performance of the alloy. Two parameters were described to study this effect size; 1) a critical dimension (λ) of the electrochemical cell big enough to affect the distribution of current density on the system, and 2) the path length between anode and cathode (L). If the ratio λ/L is low, then a uniform current distribution might be expected through the system. Contrastingly, for large λ/L, the current density is not expected to be uniform, and the corrosion is controlled by the electrical resistance of the electrolyte [689]. Simulations on Al alloys have agreed with the previous phenomena [690]. From the above, if the anodic area is too small (below 0.5 μ m for Al-alloys [690]), the ionic transportation is restricted, and local corrosion stabilizes. For much larger anodic than cathodic areas, there is not enough driving force to activate the anodic dissolution. Besides, a small inter-cathodic distance encourages local acidification. However, the previous simulation was based on Al alloys. Different crystallographic anisotropy, chemical composition, and activation energies may be expected for Ti alloys. Thus, systematic modeling and experimental studies about the anode/cathode effect size in Ti alloys remain a future opportunity for researchers. 8.2.3. Effect of the environment The prosthetic materials are subjected to severe requirements to safeguard the safety of the patients. The shelf life and stability in the human environment are critical factors for their feasibility as biomedical materials. The body fluids can be a severely corrosive environment for many metallic materials. As a reference, serum and interstitial fluids can hold concentrations of Cl − of about 113–117 mEq l −1 692 . The body fluids are electrolytes rich in different ionic species, such as Na + , Ca 2+ , PO 4 3− , SO 4 2− , etc., that cause degradation on the implant surface. Furthermore, there is a high concentration of dissolved oxygen in the blood and intercellular spaces [691]. If the material is designed for dental applications, the ionic F − concentration also becomes a significant challenge. The aforementioned conditions constitute a severe corrosive media for implant materials. Thus, the implant material should be electrochemically stable under different in-vitro and in-vivo corrosive media, simulating the physiological body fluids or the oral environment. As body fluids, which are almost water, will surround the implant, depending on the hydrophobicity, the water molecules might be dissociated, chemisorb (undissociated), or physiosorbed. Thus, the interactions between water-oxide may be covalent, ionic, or Van der Waals’ types [691]. The hydrated oxide film surface is, therefore, an electrolyte medium. However, its electronic conductivity is inherent to its chemical bonding and structure. The TiO 2 is considered an n-type semiconductor. Cl − rich electrolytes are other common corrosive media used to test biomedical Ti alloys. The potential-pH diagrams, also known as Pourbaix diagrams, are useful to predict the passivation or corrosion A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3602 behavior of Ti when immersed in a corrosive medium as a function of the pH. Examples of Pourbaix diagrams for Ti when immersed in water and chloride solutions at 25 ◦C and 37 ◦C can be found in the literature [692, 693]. It should be mentioned that 37 ◦C corresponds to the average body core temperature [694] and is commonly used for simulating Ti alloys in body fluids environments. Examples of the above are the ion release assessments from Ti–In, Ti–34Nb–6Sn, and Ti–6Al–4V alloys at 37 ◦C in artificial saliva [59–61]. Ti alloys form a spontaneous protective oxide film of TiO 2 at specific pH and potentials that may vary with the corrosive media chemistry and the activated electrochemical reactions. As the electrochemical potential and available chemical species for electrochemical reactions vary for different Ti alloys, their Pourbaix diagrams may also vary. 8.2.4. Microbial corrosion Biofilms formed by bacteria may influence (accelerate or reduce) the implant degradation during service [695,696]. Biofilms produce corrosive metabolites, such as organic and inorganic acids [697–699], which assist on the formation of local oxygen/ion concentration cells. Those factors are responsible for the occurrence of microbiologically influenced corrosion (MIC), i.e., acceleration or inhibition of corrosion by microorganisms [697]. From the literature, the main pathogenic species influencing orthopedic implant infections are S. aureus (34%), S. epidermis (32%), Pseudomonas (8%), Enterococcus (5%), among others [95]. Raising or decreasing the corrosion rates by MIC depends on the type of pathogens involved and the characteristics of the media [700]. Surface roughness, chemical composition, microstructure, temperature, kind of pathogen, and nutrients content at the media influence the bacterial attachment and biofilm formation [697,701]. An example of affecting microstructural factors is porosity. Staphylococcus aureus was found to adhere and colonize the surface of porous Ti–6Al–4V alloy, changing the oxide layer composition due to the increment of metabolic products [702]. Besides, as explained in previous section 7.1., diverse defects may encourage pitting formation, as well as bacteria adherence. MIC may involve high risk of implant failure due to the generation of pitting. Pitting breaks the Ti-rich oxide passive layer, allowing the infiltration of ions, such as Cl − , F − , etc. Furthermore, the kinetics and severity of MIC are still unpredictable. Regardless of the good corrosion resistance of Ti alloys, pitting by MIC is so far inevitable and frequently leads to premature implant failure or revision surgeries. The risk of infection from 0.5 to 5% of the total joint replacements is estimated to occur due to microbes [95]. Tailoring chemical composition can assist microbiologically influenced corrosion (MIC) mitigation on Ti alloys. Different β- or nearly β-Ti alloys have developed antimicrobial properties due to adding antimicrobial metallic elements [532,533,526,703]. Many metallic elements are reported as promising antimicrobial materials. They are an alternative against the microbes’ multidrug resistance that has been encouraged by the extensive use of antibiotics [704]. Silver and copper are some of the most effective metals against multiple pathogens, including bacteria, viruses, fungi, and algae [705]. Both metals are efficient against several pathogens, including Influenza H1N1, Candida albicans, Salmonella enterica, SARS-CoV-2, Escherichia coli, and Staphylococcus aureus, among others [705–707]. However, Cu provides a more cost-efficient route than Ag. From the previous section 2.6., designing low-cost implant materials is a priority. However, the implant material should compromise efficient mechanical properties, low toxicity, biocompatibility, and good corrosion performance. Thus, attention should be paid to certain trade-offs such as antimicrobial properties-corrosion. An example of the above is the addition of Cu and Ag in a Ti matrix, which generated galvanic pairs due to segregates or intermetallic phases [82]. 8.3. In-vitro behavior of prosthesis based on β-Ti The in-vitro evaluation of different aspects that prosthesis usually tackle is a key issue to prevent early failure. In this section will be detailed most critical strategies to investigate new insights into mechanisms, impact of alloy composition with coating and non-coating approaches as depicted in Fig. 48. 8.3.1. Antibacterial properties Due to the host tissue’s inability to generate or sustain osseointegration, prosthetic joint bacterial infections and consequent aseptic loosening are among the most frequent complications following prosthesis or implant surgery [708–710]. Depending on the type of cell wall they possess, bacteria are either classified as Gram-positive or Gram-negative. The lack of a membrane in the peptidoglycan thick external layer of Gram-positive bacteria makes it extremely beneficial for antibiotic targeting. Gram-negative bacteria are wrapped by an outer cell membrane and cytoplasmic membrane, with a thinner layer of peptidoglycan (2–3 nm) in between them that is known as the periplasmic space. The oral cavity contains a complex microbial flora that is in homeostasis, however an imbalance carried on by surgery may imply the colonization of new pathogenic microorganisms, such as bacteria, that might easily infect this cavity. Gram-positive dental infections include Eubacterium nodatum, Filifactor alocis, Slackia exigua, Parvimonas micra, Staphylococcus aureus, Staphylococcus Warneri, and Streptococcus intermedius; the corresponding gram-negative, Bacteroides fragilis, Prevotella intermedia, Porphyromona gingivalis, Treponema denticola, Tannerella forsythia, Aggregatibacter actinomycetemcomitans, Campylobacter rectus, Treponema socranskii, and Pseudomonas aeruginosa are among examples. Moreover, S. aureus is a common bacterium in various prosthetic infections. E. aureus aeruginosa, E. coli, and S. E. Epidermidis S. faecalis C. lutea and albicans [711,712]. In this sense, following implant surgery, there is a phenomenon known as the "race for the surface," in which bacteria, macrophages, fibroblasts, and even platelets are drawn to the surface, especially when it is 50 nm or less away from the implant so they can adhere to it [713,714]. In a second step, the bacteria colonization spreads the formed and matured biofilm infection, affecting not only the implant but also the tissues after the bacteria proliferation [715]. In fact, the basis of bacterial reduction strategies is anti-adhesion to a metal surface, anti-colonization, anti-biofilm, or even anti-proliferative strategies. It is established that S. aureus can infect bone canaliculi, forming a bacterial reservoir that can result in chronic or recurrent infections like osteomyelitis [716]. The spread of bacterial Fig. 48. Research activities concerning antibacterial behavior aspects. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3603 biofilms can occur quite quickly; in one study that used a pig model, the bacterial biofilm had spread to the surroundings in just six days [715]. There are numerous established approaches for assessing the antibacterial activity of solid materials, the majority of which rely on one or more of the agar diffusion plate test, plate count method, or live/dead stain techniques [82]. For the agar diffusion method, a hot agar solution is poured onto sterilized Petri dishes, and then a suspension of bacteria is sprayed over the Petri dish. A test sample, a control sample, and a positive sample are each placed in contact with the agar separately. They are then incubated for 18–24 h at (37 ±1) ◦ C. The width of the inhibition zone, which is used to quantify antibacterial activity, ought to be up to 1 mm to qualify as having good antibacterial effects. In the plate-count method, a bacterial suspension (approx. 0.4 mL) is poured onto the sample, control sample, and the negative sample, and they are incubated at (35 ±1) ◦ C for 24 h. After the incubation, the active bacteria colonies are counted. The live/dead stain method evaluates bacterial adhesion, death, or biofilm formation. In this case, the bacterial suspension is dripped onto the samples, or the sample is immersed inside a bacterial suspension after incubation at 37 ◦C for 24 h. The diagnosis of a prosthesis infection is frequently difficult to accomplish since it may require a joint aspiration and analysis of leukocyte, D-lactate, or protein levels [717,718], and it is challenging to be treated because it frequently entails high costs and loss of quality of life [719]. Unfortunately, the idiosyncrasy of bone tissue exhibiting a peculiar pharmacokinetic environment and undesirable metabolic reactions, facilitated by the infection, usually offer difficult circumstances to use clinical antibiotic therapy [720,721]. Furthermore, antibiotic resistance, which occurs as a result of their use and leads to a change in which they lose their ability to kill or stop bacterial growth, is also based on a number of mechanisms, including enzymatic inactivation or a decrease in membrane permeability, and is unfortunately growing to be a significant public threat on a global scale. Hence, antibiotic intake should be controlled to avoid excessive or misuse to ensure the success of antimicrobials in treating infections. To preserve the functionality of implanted devices exhibiting optimal characteristics at the biomaterial-tissue interface and to ensure that implanted devices remain functional in situ, several innovations dealing with surface treatments, based on coated or non-coated methodologies, or biofilm microbiology, have been developed to tackle these issues. i. Coating methods: It is known that antibacterial prosthesis containing metallic materials inhibits regular mechanism of bacteria. For instance, to increase wear resistance as well as antibacterial properties, TiN coatings have been fabricated [722,723], although they have not been clinically evaluated a decrease of bacteria is produced caused from the incorporation of hydroxyapatite (HAp) using a low-cost plasma spraying [724–726]. Other approaches including sol-gel methods with antibacterial entities, such as silver nanoparticles, could provoke undesirable chemical structures or delamination in numerous cases, but fortunately on the other hand, it implements mechanical properties and decrease the possible adherence of S. aureus bacteria in a rate of 99.3%, after 24 h [727]. Also, methods such as PVD (Physical Vapor Deposition) [347, 348,728], easily suffers delamination, although they increase wear resistance, but the presence of silver in these cases guarantees the antimicrobial properties against methicillin-sensitive S. epidermidis [729]. Furthermore, bioactive glasses such as BGF18 has shown important anti-biofilm behavior at low concentrations (1.34 mg/mL) for P. aeruginosa and S. epidermidis, also promoting osseointegration mechanism [730]. Additionally, other approaches concerning antibiotic coatings have been developed, although resistance issues are always an important drawback, as already mentioned [718]. ii. Non-coating methods: Another promising area of research to induce antibacterial properties is the use of acid etching on the Ti surface. CP-Ti and Ti–6Al–4V, has been used in this sense, also implemented with silver ions to exert this action, but unfortunately residual ions are commonly introduced in surrounding tissues [731]. It is well established that smoother surfaces implement bacteria growth, so studies to induce irregularities are appropriated. Laser Treatments, also with CP-Ti and Ti–6Al–4V has been evaluated and this is considered a promising research area to combat infections. Electrochemical anodization has accomplished good results in-vivo, but unfortunately is associated with a lack of uniformity although it displayed bactericidal properties when the fabricated nanotubes were combined with gentamicin/chitosan against S. aureus [732]. Interestingly, metals and alloys containing Ti and other metals such as Cu or Ag has been described to provide adequate antibacterial behavior [733–739]. More recent findings use stainless steel, or include other alloys with cobalt, zinc, or magnesium to induce antibacterial properties and that fulfill mechanical properties, corrosion resistance and biocompatibility as well [82,740]. In general, β-Ti alloys do not exhibit antibacterial characteristics, and after surgery the implants usually suffer infection [626]. Therefore, it is critical and significant to endow the antibacterial property of β-Ti alloys via surface implementation. In this sense, laser-nitrided surfaces could induce significant antibacterial effect as shown by Chang et al., that employed an automated fiber laser system (Micro Laser Systems) with Ti–30Nb–4Sn alloy to reduce the S. aureus biofilm coverage ten times [741]. Similar surface fabrication procedure using Ti–35Nb–7Zr–6Ta (TNZT) was evaluated with S. aureus, to show a decrease of the bacteria percentage from 5% to less than 1%, and it seems that TiN in the outermost surface oxide could improve biocompatibility and antibacterial behavior [530]. The fabrication of Ti–13Nb–13Zr–5Cu alloy proportioned good antibacterial property (S. aureus>90%), due to the precipitation of the Ti2Cu phase [742]. Furthermore, Liu et al. prepared SLA–TiCu using a combined sandblasting and large-grits etching (SLA) methodology to increase antibacterial rate against S. aureus [743]. An alternative for surface Ti–Mo alloys to bear appropriated texture was used with an electron beam surface combined with a thermal treatment to enhanced antifouling against S. aureus, as it was described by Ferraris et al. [744]. Other alloys based on Ti–Mn–Cu alloys are potential candidates and have shown that the more content of Cu added, the less colonies of E. coli colonies are present around the implant. Furthermore, Ti5Cu and Ti–10Cu sintered alloys can be considered materials with antibacterial ability within the Food Safety Standard [529]. 8.3.2. Biocompatibility evaluation Ti alloys are simply made using non-toxic alloying ingredients to increase biocompatibility. However, it is crucial to evaluate the biocompatibility of new compositions and associated processing impacts [553]. The cytotoxicity of alloys is frequently assessed using the fibroblasts’ viability. Investigations have revealed that various β-Ti alloy compositions are similarly biocompatible compared to either CP-Ti or Ti–6Al–4V. A limited number of studies claim that β-Ti alloys perform better than CP-Ti or Ti–6Al–4V in in-vitro osteogenic differentiation [745,746]. Alteration in the passive layer’s characteristics and the emission of metal ions that promote osteogenic differentiation were both suggested as causes of the improvement [745,746]. The proposed alloys are assessed for their inflammatory response in addition to cell viability and differentiation. The inflammatory response of alloys is assessed by measuring the amounts of macrophage-released cytokines in-vitro. Low levels of inflammation are reported for β-Ti alloys, which are equivalent to CP-Ti and Ti–6Al–4V [747]. The adhesion of stem cells to severely deformed surfaces of aged Ti–32Nb–2Sn alloy, on the other hand, revealed a decrease, but proliferation and osteogenic differentiation were increased [661]. The increased focal adhesions of stem cells induced by roughness were attributed for the higher proliferation and differentiation. Generally, in-vitro experiments have proven that β-Ti alloys are biocompatible. It is crucial to understand that different A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3604 materials may have different biological responses when used in-vivo. The in-vivo reaction of only a few compositions of β-Ti alloys has been studied. The in-vivo bone tissue biocompatibility of β-Ti alloys is measured using variables including bone to tissue volume ratio, bone apposition ratio, bone implant contact, and new bone area. The majority of the time, the alloys are implanted in the femur (or tibia) of rabbits [748,749]. According to Prado et al. [617], Ti–35Nb–7Zr–5Ta had a higher bone to tissue volume ratio and more bone implant contact than CP-Ti. Similar to this, it was discovered that Ti–Mo based alloys had larger new bone areas at 26 weeks than Ti–6Al–4V [750]. Contrary to the alloys, Ti–6Al–4V showed a reduction in bone area with an increase in time from 12 to 26 weeks [750]. Due to their lower modulus than Ti–6Al–4V, the alloy did not experience bone loss [750]. However, several investigations [662,748,749] have revealed that CP-Ti/Ti–6Al–4V and β-Ti alloys exhibit comparable in-vivo responses. Overall, β-Ti alloys appear to meet the requirements established by CP-Ti or Ti–6Al–4V tested in-vivo. To determine the biocompatibility of β-Ti alloys, however, significant, and prolonged in-vivo testing is required. For instance, the release of metal ions after implantation and their potential harmful effects have not been studied. i. In-vitro assessments The studies concerning the biocompatibility describe all the appropriate biological requirements, host response and capacity of a prosthesis to be used as a medical device in a specific application. It is also one of the mandatory requirements for the clinical use of biomaterials to evaluate potential injurious and toxic effects of biological systems. Specifically, cell biocompatibility that compile numerous studies related to cell viability, apoptosis, differentiation, or migration. In the last decade, it has been described through a bionic perspective that the acquisition of similar human bone structure topography and morphology is critical to the potential biological activity of Ti implants, in terms of cell growth and osseointegration [751]. The micro-topography on the surface could be obtain via sandblasting, acid etching, anodizing, spin-coating, sputtering deposition, or introducing other active materials for an implementation [39,752,753]. Hence, Kheradmandfard et al., described a gradient nanostructure layer on the TNTZ surface bearing a microstructure composed of nanoflakes with a width of 60–200 nm, to demonstrate induced enhanced cell adhesion, osteoblast differentiation, excellent osseointegration using mesenchymal stem cells [630]. It was concluded that this micro-pattern favored excellent bone regeneration as well as biological activity. Recently Yao et al. have published that commercial purity Ti bearing an ultra-fine-grained structure, influences the increased osteoblast adhesion due to the increased number of grain boundary sites [754]. In the Ti alloys it has been widely described that the ions from some of the elements such as Cu or Ag could cause cytotoxicity [82]. If the alloy exerts no damage, it has been described that on the surfaces the cell should have a polygonal shape containing a numerous lamellipodia or filopodia [755]. Also, alloys could not only bear antibacterial atoms, but they could also contain Zn or Mg ion to promote stimulation for osteogenesis as well as viability, proliferation, and adhesion [756,757]. In general, β-Ti alloys made up of non-toxic elements (Nb, Ta, and Mo, Zr, Au, and Sn) are biocompatible devices and acceptable mechanical and chemical properties, with an elasticity comparable to cortical bone, with a superior cell response and osteo-compatibility, to be employed as implant [122]. The Ti–45Nb binary alloy with β-Ti phase exhibit good mechanical properties bearing convenient balance between low elastic modulus (64.3 GPa), high tensile strength (527 MPa) and high hardness (233.4 MPa), and a controlled chemical corrosion [662]. Also, it is known that both L929 and MG-63 cells presented good rates of attachment, proliferation, differentiation on the surface, influencing a good bone tissue compatibility, compared to that of pure Ti. In this sense, Chernozen et al. have described the preparation and comparison of hybrid composites based on β-alloy Ti−xNb and oxide nanotubes, to highlight that the best hMSC cell adhesion after 24 h was observed for the NTs, obtained at an anodization with an inner diameter of ~50 nm, also hMSC proliferation exhibit better results on flat substrates [758]. In conclusion, surface nanostructurization of NTs with a diameter of ~50 nm is more favorable than fabricated NTs with a larger inner diameter. Bai et al. have fabricated a Ti–45Nb binary alloy with β-Ti phase with appropriate mechanical and chemical performance (low elastic modulus, high tensile strength and high hardness, good corrosion resistance), that after in-vitro cytotoxicity evaluation, reflected that both L929 and MG-63 cells presented a convenient attachment, proliferation, differentiation, and what it has formed critical new bone and tissue formation (identical to pure Ti) [662]. With regard to Ti–25Nb–11Sn alloy disks, another example includes cell culture tests (L929 and MC3T3-E1) that showed good proliferation, similar relative cell growth ratios, and relative absorbance ratios when compared to Ti–6Al–4V alloy and CP-Ti, along with surface stability and appositional new bone formation around the rods, after histopathological examination [748]. This result suggests that the cytotoxicity is identical if compared to Ti–6Al–4V alloy or CP-Ti, already in the market. Moreover, new beta Ti–Mo–Nb alloy compositions lead to good biocompatibility in terms of adhesion, morphology, proliferation, no critical inflammatory response was described in RAW 264.7 macrophages and differentiation potential of MC3T3-E1 pre-osteoblasts if compared to CP-Ti [759]. Other combination for alloys such as Ti–25Nb–11Sn alloy has demonstrated that direct contact cell culture tests two types of cells exhibit good proliferation, compared to Ti–6Al–4V alloy and CP-Ti [748]. Besides, it was concluded that it is a bioinert material with bone tissue compatibility, good surface stability and has demonstrated appositional new bone formation in rods, one of the mail requirements for use in the orthopedic area. Also, Donaghy et al. described a fiber laser treatment for TNZT alloys that studies dealing with cell attachment, proliferation, and differentiation of MSCs promote significant decrease for bacteria adhesion [474]. In this case, the texture acquired is a critical feature that induce nano-topography to increase surface area that enables a more easily attachment of cells since this is related to the protein interactions, filaments, and tubules, which controls cell signaling and regulates cell functionality as well. The same author reported in 2019 that nitrided surfaces showed important antibacterial effect but offered no advantage to MSC response for beta TNZT alloys [760]. Ferraris et al., established that the procedure of electron beam technology applied to Ti15Mo alloys to obtain parallel grooves with 10 μ m showed that they can support gingival fibroblast adhesion and alignment with no bacterial contamination [744]. Hence, this is an example that highlight that without the use of a lot of active agents, often associated to cytotoxicity and regulatory problems, the biological behavior could be improved. In summary, to validate the clinical translation of prosthesis based on β-Ti, is clear that more research is required on these topics. ii. In-vivo assessments In the early to mid-1960s, Branemark and his colleagues performed in-vivo tests on dogs and rabbits using Ti implants, and they observed that the implants were perfectly implanted into the bone without inflaming it. This marked the beginning of titanium’s use in dental implants [761]. In-vivo studies are critical to implement prosthesis due to the limitations offered by the presence of multispecies such as pathogens, fibrinogen, fibronectin, collagen, plasma albumin, or immune system cell and diverse body fluids. However, in general to establish reliable in-vivo procedures is not an easy task, and it usually takes a long time, is also expensive and requires specialist manipulation. That could be presumably the reason why there are only few examples in the literature concerning these kinds of combined studies focused on the use of β-Ti alloys for example, in the case of the Ti–Cu alloy, an in-vivo study was carried out using rabbits [762]. Briefly, bacterium was implanted in the muscle and pathological observation, the white blood cell test and the plate-count method were analyzed to observe a general appearance. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3605 Also, Niinomi et al. reported that TNTZ alloy can promote bone remodeling, although it has a low elastic modulus that inhibit its practical use for prosthesis [763]. In this case, the X-ray photographs of intramedullary rods show a clear bone formation as well as the pictures of extracted bone plates fixed to tibiae. An alternative alloy, the Ti–15Mo–1Bi, although does not exhibit appropriate mechanical properties if compared to others Ti Mo alloys, however the addition of Bi with antibacterial behavior has been used to implement its use in clinical [750]. Besides, in-vivo experiments employing 6 mm ×32 mm metal pin implanted into the epicondyle of rabbit femur, has shown excellent new bone growth (the new bone measured 249%), and better persistence of the deposited surrounded tissue, if compared to Ti–6Al–4V after 26 weeks. In summary, although only few examples of in-vivo experiments have been described, no drawbacks have been found when using β-Ti alloys, so new and exciting data would probably be expected to be compiled in the near future in this area. 9. Conclusion and future topics Ti alloys is a metal with numerous advantageous properties, including light weight, substantial strength, and superior corrosion resistance. The strategy to obtain strength comparable to or even superior to Ti–6Al–4V while maintaining a low modulus is the main motivation behind β-Ti alloys development. The likelihood of cytotoxicity is diminished by using non-toxic constituent elements. The main motivation of ongoing research in the field of β -type Ti-based super elastic alloys is the need for a metallic implant material that unites super elasticity with biocompatibility comparable to that of pure Ti and Ti–6Al–4V alloys. It is crucial to evaluate how processing affects a broad range of material performance metrics, including biocompatibility, corrosion resistance, fatigue resistance, and tribological performance. Data on how processing influences the fatigue and tribological performance of β-Ti alloys is woefully missing in the literature. Furthermore, it is crucial to confirm the long-term biocompatibility of these metals utilizing ethical in-vivo tests using large animal models. Future advancements in the investigation of β-Ti alloys for orthopedics can be made in the following crucial sectors: 1. New medical devices should be made using material computing science, to mimic the evolution of microstructures during processing and their impact on mechanical characteristics, sophisticated approaches, such as phase modeling, can be applied. A phase-field model, for instance, can be used to represent the rapid growth that takes place during aging treatments. Then, mechanical properties might be predicted using the developed microstructure. In order to create new β-Ti alloy compositions, first principle calculations might be expanded further. 2. To improve thermomechanical processing so that alloys have improved fatigue and wear resistance, research should be supported. This could be accomplished by creating high throughput approaches for determining acceptable microstructures. 3. Investigating whether peptides comprising the amino acids arginine, glycine, and aspartic acid (RGD), which have been demonstrated to regulate osteoblast adhesion, would encourage osteoblast differentiation when attached to a β-Ti surface. 4. Long-term biocompatibility must be determined using large animal models. These investigations may emphasize on the quantity of leached metal ions and their cytotoxic effects, corrosion rate, production of wear debris and consequent inflammatory response, osseointegration, and stress shielding effect. The performance of these materials should be assessed using commercially available medical device simulators in addition to in-vivo tests. Additionally, exploring new chemical entities currently in development could be an intriguing aspect to consider as a supplement to antibacterial properties when incorporated into coatings as a preventative measure. Assuming ongoing research efforts are focused on eventually engineering high performance orthopedic devices, Ti alloys have the potential to become the next generation of materials for orthopedic applications. Author contributions Amir A. Elhadad: Writing – original draft, Writing – review & editing. L. Romero-Resendiz: Conceptualization, editing. M. C. Rossi: Writing – original draft, Writing – review & editing. L. M. Rodríguez-Albelo: Writing – original draft, Writing – review & editing. Sheila Lascano: Writing – original draft, Writing – review & editing. Conrado R M Afonso: Writing – original draft, Writing – review & editing. A. Alcudia: Writing – original draft, Writing – review & editing. V. Amig´ o: Writing – original draft, Writing – review & editing. Yadir Torres: Conceptualization, Methodology, Project administration, Funding acquisition. Funding This publication has been part of the following projects: - R+D+i projects PDC2022-133369-I00 and PID2022-137911OB-I00, financed by MCIN/AEI/10.13039/501100011033/ - CNPq Universal Project #422015/2018-0 (C.R.M.A.); - FAPESP (S˜ ao Paulo Research Foundation) for financial support through “Projeto Tem´ atico” # 2018/18293-8, post doc Grant # 2021/03865-9 (M.C.R.) Data availability No data was used for the research described in the article. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. References [1] Amini AR, Laurencin CT, Nukavarapu SP. Bone tissue engineering: recent advances and challenges. 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Amir Elhadad: Amir Elhadad previously worked as an Assistant Professor at the Faculty of Science, Al-Azhar University in Cairo, Egypt. He received his Ph.D. from Carlos III University of Madrid, Spain, with a focus on designing novel coatings for biomedical applications. Currently, he is a Research Associate at the College of Dental Medicine, Qatar University. Liliana Romero: Liliana Romero-Resendiz is currently an assistant professor at the National Autonomous University of Mexico, Mexico; senior Marie Curie research fellow at Bournemouth University, UK; and honorary research fellow at the University of Birmingham, UK. Her research interest includes the design and development of heterostructured alloys with improved strength-ductility trade-off for interdisciplinary applications, e.g., antimicrobial, structural, biomedical, etc. Mariana Correa Rossi: Mariana Correa Rossi is currently works as a postdoctoral fellow at Federal university of S˜ ao Carlos and develops studies of surface modifications of beta Ti alloys via electrochemical routes (anodization and oxidation by electrolytic plasma) in order to improve tribological and corrosive properties for biomedical applications. She has experience in manufacturing and characterization of porous and dense metallic materials. Also has experience using laser surface treatment to improve the microstructure and consequently mechanical and wear properties of the materials. She obtained two scholarship grants and at the moment is as researcher at Universitat Polit´ ecnica de Valencia. Luisa Marleny Rodríguez-Albelo: L. M. Rodriguez-Albelo is lecturer at Department of Pharmaceutical Organic Chemistry, Faculty of Pharmacy, at University of Seville. Her current research interests include fabrication of metallic, ceramic, and polymeric porous materials, by powder metallurgy and additive manufacture, for biomedical purposes. Furthermore, her skills arise from the study of coordination polymers (metal-organic frameworks, MOFs) as porous nanomaterials. Sheila Lascano: Sheila Lascano Farak has a PhD in Mechanical Engineering, and she is a member of the steering committee of the Postgraduate in Mechanical Engineering at UTFSM and director of the Multidisciplinary Research Group in Materials and Advanced Manufacturing Laboratory. Her research is focused on the development and characterization of new materials and the development of metallic biomaterials with controlled porosity, through powder metallurgy and FAST/SPS rapid sintering. Conrado R M Afonso: Conrado is an Associate Professor of Materials Science and Metallurgical Engineering at the Universidade Federal de S˜ ao Carlos (UFSCar) in Brazil and researcher at the Materials Science and Engineering Graduation Program (PPG-CEM). He is specialized in transmission electron microscopy, Beta Ti alloys for implant applications, laser additive manufacturing, surface modifications and characterization of advanced metallic materials. He is interested in nanostructured, complex and multiprincipal alloys and correlation of processing, microstructure, and mechanical properties. Ana Alcudia: Ana Alcudia’s research journey began after graduating in Chemical Sciences from the University of Seville in 1994, eventually leading to her current position as an Associate Professor in Organic and Medicinal Chemistry since 2011. She pursued her PhD in "Optimally Pure N-sulfinilimines: Application to the Synthesis of Aziridines and Aminoalcohols," earning outstanding qualifications and an Extraordinary Doctorate Award in 1999. Following this, she undertook a postdoctoral stay at Emory University, focusing on molybdenum complex synthesis and its applications to biological products. Her research took her to Johnson and Johnson Pharmaceuticals Laboratories and PharmaMar S.A. Later, she secured a postdoctoral research contract at CSIC. Vicente Amig´ o: Vicente is Professor of Materials Science and Metallurgical Engineering at the Universitat Polit` ecnica de Val` encia and researcher at the University Institute of Materials Technology. He is specialized in titanium alloys for biomedical applications and powder metallurgy of dense and porous alloys. He is interested in the microstructural, mechanical, chemical and biological characterisation of biomaterials. Yadir Torres: Dr. Torres is a Full Professor at Seville University, serving as the Deputy Director of Postgraduate Studies and R&D +i at the Polytechnic School, and leading the research line on Advanced Porous Applications. His primary research areas and expertise lie in the design and manufacturing of custom-made porous materials, surface modification, as well as the bio-functional and tribo-mechanical behavior of materials for biomedical applications. A.A. Elhadad et al. Journal of Materials Research and Technology 33 (2024) 3550–3618 3618