d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 Available online at www.sciencedirect.com ScienceDirect jo ur nal ho me pag e: www.int l .e lsev ierhea l th .com/ journa ls /dema Development of binary and ternary titanium alloys for dental implants Jairo M. Cordeiroa,b, Thamara Belinea,b, Ana Lúcia R. Ribeiroc,d, Elidiane C. Rangel e, Nilson C. da Cruze, Richard Landers f, Leonardo P. Faveranig, Luís Geraldo Vazh, Laiza M.G. Faish, Fabio B. Vicenteb,i, Carlos R. Grandinib,j, Mathew T. Mathewb,k,l, Cortino Sukotjob,l, Valentim A.R. Barãoa,b,∗ a University of Campinas (UNICAMP), Piracicaba Dental School, Department of Prosthodontics and Periodontology, Av. Limeira, 901, Piracicaba, São Paulo 13414-903, Brazil b Institute of Biomaterials, Tribocorrosion and Nanomedicine (IBTN), Brazil and USA c Faculdade de Ciências do Tocantins (FACIT), Rua D 25, Qd 11, Lt 10, Setor George Yunes, Araguaína, Tocantins 77818-650, Brazil d Faculdade de Ciências Humanas, Econômicas e da Saúde de Araguaína/Instituto Tocantinense Presidente Antônio Carlos (FAHESA/ITPAC), Av. Filadélfia, 568, Araguaína, Tocantins 77816-540, Brazil e Univ Estadual Paulista (UNESP), Engineering College, Laboratory of Technological Plasmas, Av. Três de Março, 511, Sorocaba, São Paulo 18087-180, Brazil f University of Campinas (UNICAMP), Institute of Physics Gleb Wataghin, Cidade Universitária Zeferino Vaz—Barão Geraldo, Campinas, São Paulo 13083-859, Brazil g Univ Estadual Paulista (UNESP), Aracatuba Dental School, Department of Surgery and Integrated Clinic, R. José Bonifácio, 1193, Aracatuba, São Paulo 16015-050, Brazil h Univ Estadual Paulista (UNESP), Araraquara Dental School, Department of Dental Materials and Prosthodontics, R. Humaitá, 1680, Araraquara, São Paulo 14801-903, Brazil i Universidade Paulista (UNIP), Av. Nossa Sra. de Fátima, 9-50, Bauru, São Paulo 17017-337, Brazil j Univ Estadual Paulista (UNESP), Laboratório de Anelasticidade e Biomateriais, Av. Eng. Luiz Edmundo Carrijo Coube, Bauru, São Paulo 17033-360, Brazil k University of Illinois College of Medicine at Rockford, Department of Biomedical Sciences, 1601 Parkview Avenue, Rockford, IL 61107, USA l University of Illinois at Chicago, College of Dentistry, Department of Restorative Dentistry, 801 S Paulina, Chicago, IL 60612, USA a r t i c l e i n f o a b s t r a c t Article history: Received 6 June 2017 Received in revised form Objective. The aim of this study was to develop binary and ternary titanium (Ti) alloys contain- ing zirconium (Zr) and niobium (Nb) and to characterize them in terms of microstructural, mechanical, chemical, electrochemical, and biological properties. ental alloys — (in wt%) Ti–5Zr, Ti–10Zr, Ti–35Nb–5Zr, and Ti–35Nb–10Zr 10 July 2017 Methods. The experim Accepted 13 July 2017 — were fabricated from pure metals. Commercially pure titanium (cpTi) and Ti–6Al–4V were used as controls. Microstructural analysis was performed by means of X-ray diffraction ∗ Corresponding author at: Av. Limeira, 901, Piracicaba, São Paulo 13414-903, Brazil. Fax: +55 19 2106 5218. E-mail address: vbarao@unicamp.br (V.A.R. Barão). http://dx.doi.org/10.1016/j.dental.2017.07.013 0109-5641/© 2017 The Academy of Dental Materials. Published by Elsevier Ltd. All rights reserved. dx.doi.org/10.1016/j.dental.2017.07.013 http://www.sciencedirect.com/science/journal/01095641 www.intl.elsevierhealth.com/journals/dema http://crossmark.crossref.org/dialog/?doi=10.1016/j.dental.2017.07.013&domain=pdf mailto:vbarao@unicamp.br dx.doi.org/10.1016/j.dental.2017.07.013 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 1245 Keywords: Alloys Titanium Zirconium Dental implant Corrosion and scanning electron microscopy. Vickers microhardness, elastic modulus, dispersive energy spectroscopy, X-ray excited photoelectron spectroscopy, atomic force microscopy, surface roughness, and surface free energy were evaluated. The electrochemical behavior analysis was conducted in a body fluid solution (pH 7.4). The albumin adsorption was mea- sured by the bicinchoninic acid method. Data were evaluated through one-way ANOVA and the Tukey test ( ̨ = 0.05). Results. The alloying elements proved to modify the alloy microstructure and to enhance the mechanical properties, improving the hardness and decreasing the elastic modulus of the binary and ternary alloys, respectively. Ti–Zr alloys displayed greater electrochemical stability relative to that of controls, presenting higher polarization resistance and lower capacitance. The experimental alloys were not detrimental to albumin adsorption. Significance. The experimental alloys are suitable options for dental implant manufactur- ing, particularly the binary system, which showed a better combination of mechanical and electrochemical properties without the presence of toxic elements. © 2017 The Academy of Dental Materials. Published by Elsevier Ltd. All rights reserved. 1 C m N c [ s t w c c o m b b V c c c a p a c t c t T r t a [ m s a g e . Introduction ommercially pure titanium (cpTi) has been widely used as the ain biomaterial for the manufacture of dental implants [1,2]. evertheless, like any other material used in physiological onditions, it is exposed to mechanical and biological factors 3] that may impair implant survival and long-term treatment uccess. In this context, alloys have been considered to be he treatment of choice [4], due to their improved properties, hich allow for the development of materials according to linical demands [5]. The Ti–6Al–4V alloy is widely used in the replacement of pTi in situations where high strength is required [6] because f its excellent mechanical performance [2]. However, this aterial has been shown to be biomechanically incompati- le owing to its higher elastic modulus compared with that of one. Further, Ti–Al–V has been associated with the release of into the blood and urine [7], initiation of the inflammatory ascade leading to osteolysis [8,9], neurotoxic effects, negative ell viability response, and, consequently, an undesirable out- ome for implant biocompatibility with ion release [10–12]. In ddition, Al has been shown to be present in brain tissue of atients with Alzheimer’s disease [13]. Metal ions and debris released from implant materials re strongly associated with implant corrosion tenden- ies in physiological conditions [14,15]. Besides affecting he implant’s biocompatibility, the corrosion phenomenon hanges the implant’s mechanical properties and affects he bone through the abrasion and wear regimes [16]. hus, implant materials must not only fulfill mechanical equirements but also offer appropriate biological and elec- rochemical properties. Experimental Ti alloys without the presence of Al and V re being processed and studied to achieve these properties 17]. Zirconium (Zr) and niobium (Nb) elements have attracted uch special attention [18]. Zr acts as a solid-solution trengthening component when alloyed with Ti [1,19]. Ti–Zr lloys present a predominantly �-crystalline structure, which uarantees increased mechanical resistance and excellent lectrochemical behavior [1,20]. In contrast, Nb is a ˇ-stabilizer that is added to Ti to create ̨ + ̌ and ̌ alloys, which have demonstrated more promising properties for biomedical use [21], such as an excellent combination of low elastic modu- lus and high tensile strength [22,23]. In addition, the Ti–Nb–Zr alloy has shown non-toxicity toward osteoblastic cells, no allergy-related problems, and excellent biocompatibility [18]. To extend the clinical application of implants, it is neces- sary to develop new alloys that are sufficiently strong, present low elastic modulus, and are stable in a physiological envi- ronment. As mentioned above, Ti–Zr and Ti–Nb–Zr alloys appear to be promising candidates for dental implant appli- cations. Extensive studies have been conducted with cpTi and Ti–6Al–4 V [24–28], but studies with Ti alloys containing Nb and Zr are limited. Thus, the aim of the current study was to characterize the microstructure and mechanical, chemi- cal, and electrochemical properties of binary and ternary Ti alloys containing Zr and Nb and to conduct a comparison with the materials that are widely used for dental implants: cpTi and Ti–6Al–4V alloy. The biological aspects of such alloys were investigated by means of a protein adsorption assay. 2. Materials and methods The experimental design of this study can be seen in Fig. 1. Two control groups were considered: cpTi and Ti–6Al–4V alloy discs (Mac-Master Carr, Elmhurst, IL, USA) 10 mm in diameter and 2 mm in thickness. These materials were chosen because they are widely used in the manufacture of dental implants. 2.1. Fabrication of experimental alloys The experimental alloys (in wt%) (Ti–5Zr, Ti–10Zr, Ti–35Nb–5Zr, and Ti–35Nb–10Zr) were melted from pure metals (Ti, Nb, and Zr presented degrees of purity equal or superior to 99.0%) (Sigma–Aldrich, St. Louis, MO, USA) in an arc-voltaic furnace with a water-cooled copper crucible under an argon atmo- sphere. The ingots were flipped and re-melted five times to ensure homogeneity of the samples [1,29,30]. The Ti–Nb–Zr ingots were encapsulated in quartz tubes, heat-treated at 1000 ◦C for 8 h, and furnace-cooled [29,30]. All ingots were dx.doi.org/10.1016/j.dental.2017.07.013 1246 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 of Fig. 1 – Schematic diagram heated to 1000 ◦C and hot-swaged to form bars ≈11 mm in diameter. Then, Ti–Nb–Zr was machined into discs (10 mm in diameter and 2 mm in thickness). After that, the Ti–Nb–Zr discs and the Ti–Zr ingots were heat-treated at 1000 ◦C for 1 h and air-cooled to improve the alloys’ mechanical behavior and to relieve tensions generated during the machining pro- cedure [30]. Ti–Zr ingots were also machined into discs with the abovementioned dimensions. All discs were polished with #320-, #400-, and #600-grit SiC abrasive papers (Carbimet 2, Buehler, Lake Bluff, IL, USA) in an automatic polisher (EcoMet 300 Pro with AutoMet 250; Buehler, Lake Bluff, IL, USA) for surface standardization. Then, samples were ultrasonically cleaned with deionized water (10 min) and 70% propanol (10 min) (Sigma–Aldrich, St. Louis, MO, USA) and dried with warm air [24]. 2.2. Phase characterization The microstructural phases of the alloys were determined by X-ray diffractometry (XRD). A diffractometer (XRD; Pana- lytical, X′Pert3 Powder, Almelo, The Netherlands) was used, with Cu–K� (� = 1.540598 Å) radiation and operating at 45 kV and 40 mA at a continuous speed of 0.02◦ per second and a scan range from 20◦ to 90◦. The microstructural analysis of the alloys was confirmed by scanning electron microscopy (SEM; JEOL JSM-6010LA, Peabody, MA, USA). For that, the samples were conventionally polished as described above and then polished to a mirror finish with diamond paste (MetaDi 9- micron, Buehler, Lake Bluff, IL, USA), a polishing cloth (TextMet Polishing Cloth, Buehler, Lake Bluff, IL, USA), and lubricant (MetaDi Fluid, Buehler, Lake Bluff, IL, USA). Finally, a col- loidal silica polishing suspension (MasterMed, Buehler, Lake Bluff, IL, USA) was used together with a ChemoMet I polish- ing cloth (Buehler, Lake Bluff, IL, USA) [24]. The samples were subsequently etched for 4–5 s with Kroll’s reagent (5% nitric acid, 10% hydrofluoric acid, and 85% water) (Sigma–Aldrich, St. Louis, MO, USA) [31]. 2.3. Mechanical properties The Vickers hardness was measured by means of an indenter (Shimadzu, HMV-2 Micro Hardness Tester, Shimadzu Corpora- tion, Kyoto, Japan) by the application of a 0.5 Kgf load for 15 s. The test was performed in five samples of each group at four randomly distributed points [32]. The mean was calculated for each sample and then for each group to obtain the hard- ness (expressed in Vickers hardness units—VHN). The elastic the experimental design. modulus was tested by means of a nano-indenter (TI 950 TriboIndente, Hysitron Inc., Eden Prairie, MN, USA) equipped with a diamond Berkovich-type indenter (100 nm in diameter). Indentations were performed in at least ten positions of each sample, with a trapezoidal load function with 2 mN of maxi- mum force. The loading, unloading, and dwell times were 5, 5, and 2 s, respectively. The results represent the average among the obtained values [33]. 2.4. Surface characterization The chemical composition of the cpTi and Ti alloys (on the order of 1 �m3) was checked by energy-dispersive spectroscopy (EDS). X-ray photoelectron spectroscopy (XPS) analysis was used to verify the chemical composition and state of the outermost oxide layer by means of a spectrometer (Vacuum Scientific Workshop, VSW HA100) [34]. Atomic force microscopy (AFM) was used to observe the surface topography of the samples. Images of 50 × 50 �m were obtained by microscope (AFM; 5500 AFM/SPM, Agilent Tech- nologies, Chandler, AZ, USA) from two different areas in the non-contact mode (tapping). Gwyddion software (Gwyddion v 2.37; GNU General Public License; Czech Republic) was used for image processing [35]. The surface roughness parameters (average roughness, Ra; maximum height of the profile, Rt; average maximum height of the profile, Rz; and root mean square roughness, Rq) of the samples were evaluated by profilometry (Dektak 150-d; Veeco, Plainview, NY, USA). The roughness parameters were obtained with a cut-off of 0.25 mm at 0.05 mm/s over 12 s. Three mea- surements (right, center, and left of the sample) were obtained from five discs of each group and then averaged [32]. Surface free energy was analyzed with a goniometer (Ramé-Hart 100–00; Ramé-Hart Instrument Co., Succasunna, NJ, USA) and the sessile drop method. The water contact angle (polar component) and the diiodomethane contact angle (dispersive component) were calculated with Ramé-Hart DROPimage Standard software (Ramé-Hart Instrument Co., Succasunna, NJ, USA). The polar and dispersive components and the surface free energy were calculated [35]. 2.5. Electrochemical tests The corrosion assessment was carried out with a potentiostat (Interface 1000, Gamry Instruments,Warminster, PA, USA) and a standardized method of three-electrode cells as required by ASTM International (formerly the American Society for Test- dx.doi.org/10.1016/j.dental.2017.07.013 3 ( 2 0 1 7 ) 1244–1257 1247 i c g o t m w c ( H T c 6 a c a p o r l s a p i t E I c t a r p ( t S s v p I % % w a i t 2 T [ 1 u d e n t a l m a t e r i a l s 3 ng and Materials (ASTM)) (G61–86 and G31–72). A saturated alomel electrode (SCE) was used as the reference electrode, a raphite rod as the counterelectrode, and the exposed surface f the sample as the working electrode. The electrolyte solu- ion used was simulated body fluid (SBF) at 37 ± 1 ◦C (pH 7.4) to imic blood plasma. In total, a 10-mL quantity of electrolyte as used for each corrosion experiment [24]. The chemical omposition of the SBF (in kg/m3) was NaCl (12.0045), NaHCO3 0.5025), KCl (0.3360), K2HPO4 (0.2610), Na2SO4 (0.1065), 1 M Cl (60 mL), CaCl2.2H2O (0.5520), and MgCl2·H2O (0.4575) [36]. ris was used to achieve a pH = 7.4. The exposed area (in m2) of Ti materials was determined by AFM (cpTi = 1.07, Ti- Al–4V = 0.99, Ti–5Zr = 1.01, Ti–10Zr = 1.31, Ti–35Nb–5Zr = 1.06, nd Ti–35Nb–10Zr = 1.03). Electrochemical testing was conducted according to a spe- ific protocol [24,35]. A cathodic potential (−0.9 V vs. SCE) was pplied for 10 min to standardize the oxide layers of the sam- les. The open circuit potential was monitored for a period f 3600 s to assess the free corrosion potential of the mate- ial in the electrolyte solution. For evaluation of the passive ayer, electrochemical impedance spectroscopy (EIS) was mea- ured at a frequency of 100 kHz–5 mHz, with the AC curve at range of ±10 mV applied to the electrode at its corrosion otential. The values were used to determine the real (Z′) and maginary (Z′′) components of impedance, which were used o construct Nyquist, Bode (|Z|), and phase angle plots. The IS data were analyzed with Echem Analyst software (Gamry nstruments, Warminster, PA, USA). An equivalent electrical ircuit was fitted for quantification of the corrosion process in he passive/oxide film formation (polarization resistance, Rp, nd constant phase element, CPE). The samples were then polarized from −0.8 to 1.8 V (scan ate of 2 mV/s). Corrosion parameters such as Ecorr (corrosion otential), Icorr (corrosion current density), and Tafel slopes bc, ba) were obtained from the potentiodynamic polariza- ion curves by the Tafel extrapolation method (Echem Analyst oftware, Gamry Instruments, Warminster, PA, USA). The pas- ivation current density (Ipass) corresponds to the current alue of the passivation region of the polarization plot. The ercentage corrosion efficiencies with regard to the values of corr (Eq. (1)) and Rp (Eq. (2)) were calculated: CE = Icorr ∗ − Icorr Icorr ∗ × 100 (1) CE = Rp ∗ − Rp Rp ∗ × 100 (2) here Icorr* and Icorr are the corrosion current density of cpTi nd the Ti alloys, respectively, and Rp* and Rp are the polar- zation resistance of cpTi and the Ti alloys, respectively. The electrochemical tests were conducted at least five imes (n = 5) to ensure reliability and reproducibility. .6. Protein adsorption he protein adsorption assay followed a previous protocol 37]. Briefly, five samples of each group were incubated in 00 mg/mL of albumin (Sigma–Aldrich, St. Louis, MO, USA) nder horizontal stirring (7.85 rad/s) at 37 ◦C for 2 h. Sam- Fig. 2 – X-ray diffraction patterns of cpTi and Ti alloys. ples were washed in phosphate-buffered saline (PBS) (Gibco, Life Technologies, Gaithersburg, MD, USA) to remove non- adherent proteins, and the protein adsorption was measured by the bicinchoninic acid method (BCA Kit, Sigma–Aldrich, St. Louis, MO, USA). 2.7. Statistical analyses The normality of all response variables was tested by the Shapiro–Wilk method, and data were transformed when necessary. One-way ANOVA was used to test the statisti- cally significant differences among groups with regard to roughness, surface energy, hardness, elastic modulus, electro- chemical parameters (Rp, CPE, Ecorr, Icorr, and Ipass), and protein adsorption. Tukey’s HSD test was used as a post hoc technique for multiple comparisons. A mean difference significant at the 0.05 level was used for all tests (SPSS v. 20.0, SPSS Inc.). 3. Results 3.1. Alloy microstructure The XRD patterns of the cpTi and Ti alloys are shown in Fig. 2. Only peaks corresponding to the ̨ phase were observed for cpTi and Ti–Zr alloys. The Ti–6Al–4V and Ti–35Nb–5Zr alloys showed a two-phase ( ̨ + ˇ) structure. Intermetallic peaks of Al and TiAl3 were detectable for Ti–6Al–4V. In contrast, the Ti–35Nb–10Zr alloy exhibited only a ̌ microstructure. SEM micrographs were obtained after samples were etched with Kroll’s solution to confirm the microstructure of the alloys (Fig. 3). Different magnifications are demonstrated among the groups to show the most representative image of the microstructure. Only the grain boundary of the equiaxed ˛ structure was observed for cpTi and Ti–5Zr. Although Ti–10Zr has a composition similar to that of Ti–5Zr, its microstruc- ture was the typical Widmanstätten pattern, containing a fine needle-like structure (acicular ˛) with multiple orientations. Ti–Nb–Zr alloys showed a homogenous structure with unde- fined contours, which made it difficult for their microstructure to be characterized by micrographs. dx.doi.org/10.1016/j.dental.2017.07.013 1248 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 Fig. 3 – SEM micrographs of cpTi and Ti alloys. Arrows indicate the grain boundaries. The asterisk represents the ̌ phase of the Ti–6Al–4V alloy. 3.2. Hardness and elastic modulus The hardness and elastic modulus data of cpTi and Ti alloys are shown in Fig. 4. The cpTi hardness was statistically sig- nificantly lower than those of all alloys (p < 0.05). It can be seen that Ti–Zr alloys had the highest hardness (416–434 VHN), exceeding even that of the Ti–6Al–4V alloy (354 VHN). As expected, Ti–6Al–4V presented the highest elastic modulus, while Ti-Nb-Zr alloys showed statistically significantly lower values (p < 0.05). 3.3. Alloy and oxide layer composition The main elements that compose the alloys were deter- mined by EDS and can be observed in the distribution map in Fig. 5. Semi-quantitative analysis of each element (wt%) is also described. As might be expected, Ti was present in all groups. In addition to the specific elements of each alloy, C and O were detected in all materials evaluated. A homogeneous distribution of elements can be noted without the presence of aggregations or segregations, indicating that the thermome- chanical processing produced a highly homogeneous sample. Fig. 6 shows the detailed XPS spectrum with the expected compounds for the binding energies detected. The abovemen- tioned alloying elements determined by EDS assessment were also identified as signals by XPS. It can be observed that all bands (Ti2p, O1s, Al2p, Zr3d, Nb3d, and C1s) displayed similar spectral shapes among the groups, with a prevalence of strong peaks of Ti2p, O1s, and C1s (not shown). Regarding the Ti2p, Zr3d, and Nb3d spectral lines, each oxidation state exhibited two peaks (doublets). The calculated relative concentrations of each chemical state in the XPS analysis are described in Table 1. 3.4. Topography, roughness, and surface free energy AFM analysis was performed to extend the information regarding the characterization of the alloys’ surface topog- raphy. The two- and three-dimensional AFM images are shown in Fig. 7. Similar surfaces were noted among materials, showing longitudinal grooves characteristic of the polishing process. A slight difference was observed for Ti–10Zr and Ti–35Nb–10Zr alloys. For the three-dimensional images, these alloys exhibited a decrease in grooves and demonstrated bright regions. Comparative analysis among the surfaces of the alloys took into consideration a standard scale for the Z- axis. The uppermost regions are represented by lighter shades, while the deeper regions are represented by darker shades. In addition to the qualitative analysis performed by AFM, the surface was quantitatively characterized by profilometry and goniometry. The roughness and surface free energy are shown in Fig. 8. All alloys exhibited Ra values lower than those of cpTi (p < 0.05). Ti–Al–V and Ti–10Zr were the only alloys whose roughness was on the nanometer scale, at 90 and 100 nm, respectively. The Rq and Rz results followed a trend similar to that of Ra. Although there was no statistical sig- nificance, it can be observed from the Rt measure that cpTi had deeper valleys and higher peaks, as exhibited by AFM. In contrast, the surface energy values were close to each other among the evaluated materials. Alloys with higher Zr con- centration presented statistically significantly lower surface energy than did cpTi (p < 0.05). No statistically significant dif- ference was found among the experimental groups. 3.5. Electrochemical assays The electrochemical behavior of the experimental alloys in comparison with those of cpTi and Ti–6Al–4V was investigated dx.doi.org/10.1016/j.dental.2017.07.013 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 1249 Fig. 4 – (a) Hardness (n = 5) and (b) elastic modulus (E) (n = 1) of cpTi and Ti alloys. Different letters indicate statistically significant differences among the groups (p < 0.05, Tukey’s HSD test). Fig. 5 – Chemical mapping by EDS with element concentrations (wt%) for cpTi and Ti alloys. Table 1 – Relative element concentrations (at%) of the oxide layers of cpTi and its alloys by XPS. Element spectral line Groups cpTi (at%) Ti–6Al–4V (at%) Ti–5Zr (at%) Ti–10Zr (at%) Ti–35Nb–5Zr (at%) Ti–35Nb–10Zr (at%) Ti2p 11.4 7.0 14.0 8.6 4.3 6.5 O1s 32.4 25.5 36.7 35.4 19.6 30.3 C1s 56.2 65.0 48.7 54.8 74.1 59.6 b n 3 N a V3s + Al2p – 2.5 – Zr3d – – 0.6 Nb3d – – – y electrochemical impedance spectroscopy and a potentiody- amic polarization test. .5.1. Electrochemical impedance spectroscopy (EIS) yquist and Bode plots are shown in Fig. 9. Ti–Zr and Ti–6Al–4V lloys presented higher diameters of the semicircles in the – – – 1.2 0.3 0.9 – 1.7 2.7 Nyquist diagram. Regarding the Bode plot, the same alloys pre- sented higher phase angles and increased impedance values at a low frequency range. The impedance results were modeled with an equivalent electrical circuit and are described in Table 2. A simple elec- trical circuit consisting of Rsol (polarization resistance of the dx.doi.org/10.1016/j.dental.2017.07.013 1250 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 Fig. 6 – Detailed XPS spectra of cpTi and Ti alloys. Fig. 7 – Two- and three-dimensional AFM images of cpTi and Ti alloys. Fig. 8 – (a) Roughness parameters (n = 5) and (b) surface energy (n = 5) of cpTi and Ti alloys. Different letters indicate statistically significant differences among the groups (p < 0.05, Tukey’s HSD test). dx.doi.org/10.1016/j.dental.2017.07.013 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 1251 Fig. 9 – Nyquist (a) and Bode (b) diagrams for cpTi and Ti alloys in SBF. The electrical equivalent circuit is shown in the Nyquist diagram. Symbols represent experimental data, and solid lines represent fitted data. Table 2 – Means and (standard deviations) of electrical parameters obtained from the equivalent circuit models for all groups. Group Rp (M� cm2) Q (��−1 sn cm−2) � �2 × 10−3 CE (%)a cpTi 2.21 (0.51)a 15.48 (1.38)a 0.92 (0.01) 3.18 (0.72) – Ti–6Al–4V 5.56 (1.00)b 15.34 (1.78)a 0.91 (0.01) 2.17 (0.37) 151.58 Ti–5Zr 9.45 (4.48)b 11.36 (1.14)ab 0.93 (0.01) 1.89 (1.00) 327.60 Ti–10Zr 7.35 (4.54)b 10.83 (0.83)b 0.95 (0.01) 1.37 (0.56) 232.57 Ti–35Nb–5Zr 1.09 (0.26)c 24.31 (6.61)c 0.92 (0.01) 4.72 (1.64) −50.67 Ti–35Nb–10Zr 1.26 (0.39)ac 25.75 (7.60)c 0.92 (0.01) 4.36 (2.56) −42.98 he gro ative e p p s a e d t t I t c F T Different letters indicate statistically significant differences among t a Positive values of CE mean an improvement in efficiency, while neg lectrolyte), R1 (polarization resistance), and Q1 (constant hase element, CPE) was used (Fig. 9), suggesting a single com- act film on the surface. The fitting chi-square evaluation (�2) howed high quality (on the order of 10−3), indicating excellent greement between the measured and simulated values. The electrical parameters showed that Ti–Zr and Ti–Al–V xhibited statistically significantly higher resistance than id the other groups (p < 0.05). Ti–35Nb–5Zr demonstrated he worst result, showing a corrosion resistance efficiency hat was about 50% lower in comparison with that of cpTi. n contrast, Ti–5Zr seemed to have more than three times he efficiency of cpTi. Ti–Zr alloys also presented lower apacitance values, while the Ti–Nb–Zr alloys showed slightly ig. 10 – Potentiodynamic polarization curves for cpTi and i alloys in SBF. ups (p < 0.05, Tukey’s HSD test). values represent a decrease in efficiency with respect to resistance. unfavorable values of electrical parameters when compared with those of cpTi. 3.5.2. Potentiodynamic polarization curves The potentiodynamic polarization curves indicated the natu- ral formation of passive oxide layers over all of the sample surfaces (Fig. 10). Ti–Zr alloys presented the most passive feature when compared with the other groups, confirmed by the electronic parameters obtained through Tafel slopes (Table 3). Ti–Zr alloys had a much lower value of corrosion current density (p < 0.05) when compared with the control groups and Ti–Nb–Zr alloys. In addition, only Ti–Zr alloys had enhanced corrosion efficiency relative to the Icorr parameter when compared with the cpTi. Ti–Nb–Zr alloys demonstrated statistically significantly higher current density in the passive region (p < 0.05). 3.6. Albumin adsorption We measured the adsorption of albumin for all groups to understand how blood protein interacts with the material sur- face (Fig. 11). There was no statistically significant difference among the groups (p = 0.184, one-way ANOVA). However, a slight increase in albumin adsorption could clearly be seen for the alloys with higher amounts of Zr. 4. Discussion 4.1. Microstructural and mechanical properties The analysis of the microstructure is important, since several vital characteristics of alloys (i.e., mechanical and corrosion dx.doi.org/10.1016/j.dental.2017.07.013 1252 d e n t a l m a t e r i a l s 3 3 ( 2 0 1 7 ) 1244–1257 Table 3 – Mean and (standard deviation) values of electrochemical parameters obtained from the potentiodynamic polarization curves. Group Ecorr (V vs. SCE) Icorr (nA cm−2) ba (mV dec−1) −bc (mV dec−1) Ipass (�A cm−2) CE (%)a cpTi −0.50 (0.02)a 13.55 (3.48)ab 5.74 (0.03) 1.62 (0.02) 6.43 (0.52)a – Ti–6Al–4V −0.50 (0.06)a 15.64 (3.75)ad 7.11 (0.06) 1.77 (0.02) 5.97 (0.56)a 15.42 Ti–5Zr −0.51 (0.08)a 8.92 (2.27)bc 7.16 (0.14) 1.57 (0.04) 6.21 (0.51)a –34.16 Ti–10Zr −0.50 (0.03)a 7.17 (1.15)c 6.07 (0.04) 1.67 (0.02) 5.83 (0.39)a –47.30 Ti–35Nb–5Zr −0.51 (0.03)a 25.36 (10.83)d 5.95 (0.02) 1.55 (0.01) 9.36 (1.88)b 87.15 Ti–35Nb–10Zr −0.45 (0.04)a 21.92 (3.96)ad 5.20 (0.07) 1.64 (0.01) 9.44 (2.20)b 61.77 Different letters indicate statistically significant differences among the groups (p < 0.05, Tukey’s HSD test). a Positive values of CE mean a decrease in efficiency, while negative value current density. Fig. 11 – Albumin adsorption for cpTi and Ti alloys. trations [6,23,46,48]. The combination of high strength and low properties) are strongly influenced by microstructural features [38]. The presence of a single phase for the binary alloys showed that Zr, acting as a neutral stabilizer, did not signif- icantly change the crystalline structure of the material [1]. A similar result was described in the literature with Ti–Zr alloys, in which no secondary phase was detected [39]. Further, the Ti–Zr binary system exhibited a completely solid solution for both the high-temperature � phase and the low-temperature ̨ phase [40]. Regarding the Ti–6Al–4V alloy, the two-phase structure is due to the presence of Al (˛-stabilizer) and V (ˇ-stabilizer) [41]. It exhibited the ˛2 tetragonal structure (Ti3Al) as the main phase as a result of the ˇ↔˛↔˛2 transformations. In the micrographs, the presence of the ̌ phase (lighter region) dispersed in the contours of the � matrix (darker region) can clearly be seen [42]. Although the Ti–35Nb–5Zr alloy also pre- sented two phases ( ̨ + ˇ), its microstructure can be considered to be near ˇ, due to the large amount of the ̌ phase stabiliz- ing element [41]. The presence of the � phase in this alloy can be related to the thermomechanical process performed. It has been reported that hot working and solution treatment at temperatures above the ̌ transus led to the transforma- tion of part of the ̌ phase to the ̨ phase [38]. The samples in this study were air-cooled, which has been demonstrated to be a process that favors the precipitation of the ̨ phase in the ̌ matrix [18]. In contrast, the fact that the Ti–35Nb–10Zr alloy contained only the ̌ phase can be justified by the higher s represent an improvement in efficiency with respect to corrosion Nb amount and the increased Zr concentration, which sta- bilized the ̌ phase and inhibited its transformation to the ˛ phase [43]. High concentrations of Nb have also been revealed to decrease the grain size of Ti–xNb–3Zr–2Ta alloys [44]. It is postulated that such a grain decrease was very significant, so it was difficult to perform the microstructural analysis by SEM. Nevertheless, regions with light and dark colors representa- tive of solute-rich and solute-poor areas, respectively, can be seen [45]. The structural investigation aids in the understanding of the results obtained in the mechanical tests, since both the hardness and the elastic modulus are intrinsically linked to the alloys’ microstructural phases. The addition of 5 and 10 wt% Zr to Ti was able to more than double the hardness of the alloy. The increase in hardness was mainly caused by the solid-solution strengthening of the ˛-phase [1,39]. Further- more, the air-cooling process that started from a temperature above the ̌ transus can contribute to this property. In this con- dition, there is the development of a grain refinement of the ˛ phase that was transformed from the ̌ phase, which can be the major contributor to hardness [18,38]. A slight decrease in hardness has been reported in the literature with increases in Zr concentration above 15 wt% [1,39] and was also verified in our study with a Zr content of 10 wt%. The hardness values of ternary Ti–Nb–Zr alloys with 5 and 10 wt% Zr were similar. It can be clearly observed that Nb decreased the alloy hardness to a value close to but sta- tistically significantly greater than that of cpTi. Similarly, a study found higher hardness of Ti–Nb–Zr alloys with greater concentrations of Zr [46]. The decrease of hardness when com- pared with that of Ti–Zr alloys demonstrates that the softening behavior of the ̌ phase (main phase) seems to have a greater effect on the Ti–35Nb–5Zr alloy than the solid-solution hard- ening induced by the ̨ phase. Regarding the elastic modulus, Ti–Nb–Zr alloys deserve attention for presenting the lowest values, closer to those of bone (≈10–30 GPa), due to the higher amount of the ̌ phase. The ̌ phase has a structure (body-centered cubic) that exhibits a lower bonding force among the atoms, which ensures a reduced elastic modulus [1,47] and favorable plastic deforma- tion capacity of the alloy [45]. This result is in agreement with those of other studies that have found a low elastic modulus in ̌ ternary alloys containing Zr and Nb with different concen- elastic modulus is expected to minimize the “stress shielding” effect [23,49], which is advantageous in terms of bone atrophy dx.doi.org/10.1016/j.dental.2017.07.013 3 ( 2 i i b d e p t a t w u fi s p l b o 4 S t I l b f t t g d t u p s s [ f t h p T h r [ p Z N c o p i t p [ o o e d e n t a l m a t e r i a l s 3 nhibition and therefore leads to much better bone remodel- ng [19] due to improved communication between implant and one tissue through effective load transfer and uniform stress istribution [49]. In contrast, the studied binary alloys showed an increased lastic modulus, even with a Zr content of less than 15 wt%, as reviously mentioned by Correa et al. [1]. The increased elas- ic modulus could be related to the heat treatment performed fter swaging, which led to the alloy hardening. However, he value for Ti–5Zr was significantly lower than that for the idely used Ti–6Al–4V alloy, which had a higher elastic mod- lus (139 GPa) due to its hardness behavior. In general, these ndings confirm that the mechanical properties are structure- ensitive. Therefore, they can be influenced by phase state, recipitate size and distribution, grain and subgrain size, dis- ocation density, and other factors [50]. Indeed, the mechanical ehavior of Ti–Nb–Zr alloys proved to be especially dependent n the quantity of the ̌ or ̨ phase. .2. Surface properties urface characteristics such as composition, wettability, and opography influence cell adhesion and proliferation [51,52]. n addition, the biomaterial surface is in direct contact with iving tissues, with an immediate effect on the biocompati- ility and corrosion process [52]. In this work, the proportions ound of the alloying elements showed a small difference from he nominal composition. However, the Zr and Nb concentra- ions were close to stoichiometric values [17], demonstrating reater incorporation during the melting process. This small ifference can be attributed to the excluded vibrational con- ribution to the theoretical formation energy and mixing of nwanted interstitial elements (mainly oxygen) during alloy roduction [53]. The composition of the oxide layer, assessed by XPS, is hown in Fig. 6. In general, all groups presented a native pas- ive film on a surface formed mainly of TiO2 protective oxide 16,54,55]. On the Ti–Al–V alloy, only Ti and Al oxides were ormed. Normally, V oxides (V2O3 and V2O5) can be observed in he oxide layer in extremely small quantities [42], which may inder their identification. It can be observed that Zr oxides are resent in much lower concentrations than Ti and Nb oxides. his is related to the fact that the Ti4+ ions display mobility igher than that of the Zr4+ ions, leading to a significantly educed amount of ZrO2 oxide at the outermost surface layer 25]. A previous study stated that the uppermost part of the assive layer of Ti–Nb–Zr alloys consists of TiO2, Nb2O5, and rO2, followed by an intermediate sub-layer containing Ti2O3, bO2, and NbO phases [29]. It can also be observed that the oncentration of Ti oxides decreases with the incorporation f alloying elements; moreover, the presence of Zr oxides is roportional to its concentration in the alloy. In all cases, several oxidation states were observed, includ- ng weaker contributions of metallic emission coming from he base material (Ti0, Zr0, Nb0, and Al0) caused due to a sam- ling depth for XPS similar to that of the oxide layer thickness 52,54,55]. These results indicate the formation of a native xide layer that was not fully oxidized [52]. In addition to the xides formed with the alloying elements, the O1s binding nergies were associated with the formation of three peaks 0 1 7 ) 1244–1257 1253 that are representative of O2−, OH ions, and adsorbed H2O [16]. In contrast, the C1s band is correlated to carbon contaminants, such as saturated hydrocarbons, C O, and C H bonds [55]. The profilometry analysis showed a higher surface rough- ness for cpTi, corroborating the results from a previous study [51]. The AFM assessment also showed that cpTi, Ti–5Zr, and Ti–35Nb–5Zr have greater variations in height, which can sug- gest increased roughness. In addition, Ti–6Al–4V and Ti–10Zr alloys demonstrated smoother surfaces in both analyses. The results may be justified by the material properties, such as hardness, which can in some way influence the polishing pro- cess. Another driving force toward the roughness parameter is the alloy grain size. The alloys demonstrated micro-features of grain size in the SEM images. Changes in the topogra- phy that lead to a rough surface enhance the differentiation of osteoblasts, while a smooth surface increases cell pro- liferation [65–69]. Surfaces that induce the maturation and proliferation of cells are crucial for improving and accelerating osseointegration [5]. In addition to roughness, biological properties are also influenced by surface free energy. Materials having a higher surface energy show greater wettability, which is important to achieve better cell proliferation and adhesion as well as to absorb more proteins on the surfaces of Ti alloy implants. An important observation of the wettability in this work is the low values of surface energy for the polar component (H2O) in Ti–10Zr, Ti–35Nb–5Zr, and Ti–35Nb–10Zr alloys. A less hydrophilic surface can be the main factor generating lower total surface energy values for these alloys. 4.3. Electrochemical behavior The native oxide layer on the Ti surface confers inherent pro- tective corrosion behavior on the implant [2,52] as long as the integrity of the film is maintained [56]. Corrosion resis- tance is correlated with the lifetime of the implant and the harmfulness of corrosion processes that occur in the body [57]. The improved corrosion resistance found for the binary alloys can be attributed to the fact that Zr is an anodic alloying element for Ti that acts to reduce the anodic activity directly [56]. A higher resistance suggests lower corrosion rates, since this parameter is considered a measure of the barrier effect of the passive film against charge transfer [29]. In addition, Ti–Zr and Ti–6Al–4V alloys showed the largest diameters on the Nyquist plot, indicating enhanced dielectric properties, namely, an oxide layer that is a better capacitor and is more protective [16,58,59]. Similarly, the Bode plot results corrobo- rate the higher protective properties of the alloy’s oxide film [16,25]. In contrast, the Ti–Nb–Zr alloy showed the worst behavior as a result of the less resistant passive film on the surface. It can be related to the fact that the � phase proved to be less resistant to corrosion than was the � phase [60]. This result dif- fers from those of previous studies of Ti–Nb–Zr alloys, in which similar or even nobler electrochemical behavior was found for the alloys in comparison with cpTi [16] and the Ti–6Al–4V alloy [29]. In these studies, the main reason for the superiority of the alloys was the modification of the passive TiO2 layer by the alloying elements that participate, along with their ZrO2 dx.doi.org/10.1016/j.dental.2017.07.013 s 3 3 1254 d e n t a l m a t e r i a l and Nb2O5 oxides, in the thickening and improved structural integrity of the alloy’s native passive film [16,61,62]. Although such oxides were detected via XPS, their presence was insuf- ficient to improve the corrosion properties of Ti–Nb–Zr alloys in this study. A study carried out by Cvijović-Alagić et al. [25] evalu- ated the corrosion and wear behavior of Ti–13Nb–13Zr and Ti–6Al–4V. A better combination of these properties was observed for Ti–6Al–4V. A possible reason for these findings is the greater hardness of the Ti–6Al–4V alloy, which favors the maintenance of a thicker and more firmly adhered oxide layer than does a softer material [25]. Thus, the greater corrosion protection for Ti–Zr alloys in this study can also be justified by their increased hardness, since the solid-solution strengthen- ing of the alloy may be responsible for the enhanced protection against oxidation [39]. Further, alloys with Zr produce signif- icantly more stable anodic oxides, improving the corrosion resistance of the biomaterial [63]. With regard to Ti–6Al–4V, the detected intermetallic compound TiAl3 can also influence the electrochemical behavior of the alloys by having good oxi- dation resistance [64]. Regarding the polarization curves, the shifts of the curves toward lower current densities (upper left area of the graph) [25] illustrate the improved behavior of Ti–Zr alloys. It can be observed that the current density increased gradually with increases in the potential from the corrosion potential, which can be attributed to the activation polarization [60]. The cur- rent density increased until a certain time, when it achieved a constant value (Ipass) of anodic polarization without any active–passive transition [41], indicating the absence of breaks and confirming the thickening of the oxide layer [60]. Confirming the EIS results, the potentiodynamic parame- ters establish the superior corrosion resistance of Ti–Zr alloys. The lowest Icorr and Ipass values found for these alloys reflect low electrochemical activity and high corrosion resistance properties [16]. A higher Icorr value is representative of the degree of degradation of the alloy [59]. Lower current den- sity is preferable, since it indicates a more stable and resistant passive film [16,25]. Additionally, there was no statistically sig- nificant difference among groups with respect to corrosion potential (Ecorr), indicating that the corrosion processes below the passive films were similar [25]. A study carried out by Ribeiro et al. [29] with the same Ti–Nb–Zr alloys showed an increase of Ecorr and resistance parameters as a function of the immersion time in artificial saliva, representing decreased reactivity of the alloy and improved efficiency of the corro- sion protection of the passive films when in contact with the electrolyte over time. 4.4. Protein–surface interaction Concerning protein adsorption, the alloy chemical composi- tion and the small differences in roughness and surface free energy had no influence on albumin adsorption. The role of small variations of surface roughness (Ra <0.50 �m) in cellular responses and protein adsorption has not been well-defined [51]. However, a previous study found that albumin is adsorbed preferentially onto the smooth Ti–6Al–4V substratum, while the rough substratum binds a higher amount of total serum protein and fibronectin [65]. ( 2 0 1 7 ) 1244–1257 Composition can also influence protein adsorption. A study that evaluated the addition of boron to the Ti–Nb–Zr alloy showed that the chemical alteration provided by its incorpora- tion was responsible for a decrease in the protein adsorption and cell response [70]. In our study, Zr and Nb were not detrimental to protein adsorption. Indeed, it was possible to observe that higher concentrations of Zr led to an increase in this property. It is expected that hydrophilic nanotopography can facil- itate the penetration of nanoscale proteins. Several surface treatments have been investigated with the aim of modifying surface properties. An electrochemical anodization process was capable of creating a nanoporous surface that signifi- cantly enhanced albumin and fibronectin protein adsorption, improving the adhesion, migration, proliferation, and min- eralization of human bone marrow mesenchymal stem cells in the Ti–25Nb–25Zr alloy [66]. Further studies are needed to understand how changes in alloy surfaces can influence pro- tein adsorption. 4.5. Clinical implications, limitations, and future perspectives In this study, we attempted to perform a full evaluation of Ti alloys containing Zr and Nb as alternatives to cpTi and Ti–Al–V alloys. In an attempt to address all significant issues, it was observed that the experimental alloys are indeed viable options for the manufacture of dental implants. The enhanced mechanical properties of Ti–Zr and Ti–Nb–Zr alloys, such as higher microhardness and lower elastic modulus, may ensure that the dental implant has greater clinical applicability. For example, Ti–Zr alloys may be particularly suitable in situations that require greater strength, such as in rehabilitation involv- ing small-diameter implants. In these cases, cpTi is more susceptible to failure, and the use of Ti–Al–V is indicated. How- ever, as already mentioned, the use of the Ti–Al–V alloy has been avoided due to the toxicity related to the release of Al and V, but this is not a barrier in the case of the proposed non-toxic alloys. The native passive surface film of cpTi can be easily destroyed due to the intrinsic low wear resistance of Ti [6]. Thus, the improved corrosion behavior observed for Ti–Zr alloys could prevent passive film damage by mechanical solici- tations, which can provide long-term success of rehabilitation by decreasing both the probability of corrosion in physiologic environments and the probability that osseointegration will be harmed. In addition, the similar responses to protein adsorp- tion among the studied materials lead us to believe that none of the experimental alloys will negatively affect the initial adhesion of cells. Some limitations can be mentioned due to the intrinsic characteristics of the type of research, such as the validation of the results under in vivo physiological conditions. There- fore, a need for in vitro biocompatibility studies and in vivo randomized controlled trials that consolidate the efficacy and safety of alloys as alternatives to cpTi has been reported [5]. Future research should test the experimental alloys by cell cul- ture analysis and under in vivo conditions with the aim of validating their outstanding properties. Further, the influence of immersion time on the electrochemical stability of such dx.doi.org/10.1016/j.dental.2017.07.013 3 ( 2 a m p t t 5 A c • • • • A T F 2 a t f a i u D f r d e n t a l m a t e r i a l s 3 lloys is warranted for observation of the possible improve- ent in the Ti–Nb–Zr alloys’ behavior. It may be interesting to erform different heat and surface treatments in an attempt o decrease the elastic modulus of Ti–Zr alloys and to enhance he electrochemical properties of Ti–Nb–Zr alloys. . Conclusions ccording to the results found in this study, the following con- lusions can be drawn: Ti–Zr and Ti–Nb–Zr alloys exhibited an improved combina- tion of hardness and elastic modulus, respectively; Ti alloys presented a roughness and surface energy slightly lower than those of cpTi; Ti–Zr alloys exhibited superior electrochemical properties, presenting higher polarization resistance and lower values of Icorr, corrosion rate, and capacitance parameters; and the experimental alloys exhibited albumin adsorption sim- ilar to that of control groups. cknowledgments his work was supported by the São Paulo State Research oundation (FAPESP), Brazil (grant numbers 2013/08451-1, 014/26853-2, 2015/25562-7, and 2016/11470-6). 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