Elucidating the corrosion-related degradation mechanisms of a Ti-6Al-4V dental implant

Graphical abstract



The Ti-6Al-4V ( TAV ) alloy is commercially used as a dental implant material. This work seeks to elucidates the origins of degradation of Ti-6Al-4V (TAV) implant alloys that result in peri-implant bone loss. 
Methods: In this work, a combination of microstructure, surface, and solution analyses was utilized to study the corrosion mechanism of the TAV alloy in oral environments. The corrosion of TAV alloys in the F -enriched environment of a crevice was evaluated through nanoscale surface analysis. And, the findings were further rationalized via electrochemical means.


Our results suggest the bone loss was caused by crevice corrosion and the consequential release of by-products, and the crevice corrosion was potentially induced by the buildup of corrosive species such as fluorides, which are common additives in dental products. In turn, the corrosion properties of the TAV alloy were evaluated in fluoride enriched environments. Nanoscale analysis of corroded surfaces, carried out using vertical scanning interferometry (VSI) showed that the corrosion susceptibility of the constituent phases dictates the corrosion product species. In specific, the aluminum-rich α phase preferentially dissolves under potential-free conditions and promotes the formation of insoluble Al-Ti oxides. Notably, under conditions of applied potential, oxidative dissolution of the vanadium-rich β phase is favored, and the vanadium release is promoted.


These findings elucidate the origins of degradation of TAV-implants that result in the release of corrosion by-products into the local biological environment. More important, they offer guidelines for materials design and improvement to prevent this nature of degradation of dental implants.


Owing to the improved performance and the long-lasting durability, the use of dental implants has become a common option to treat tooth loss. Titanium (Ti)-based alloys such as Ti-6Al-4V (TAV: 90 mass % Ti, 6 mass % aluminum (Al), and 4 mass % vanadium (V)) are often used as dental implant materials because of their excellent mechanical performance and biocompatibility [ ]. However, the in vivo degradation of an osseointegrated (i.e., “attached to the bone”) implant can lead to saucerization and osteolysis, both of which are manifestations of bone mass loss [ ]. Assembled implant components are particularly susceptible to crevice corrosion [ ]. Following the infiltration of oral and body fluids, the presence of crevices promotes the accumulation of aggressive anions (e.g., F and Cl ) and localized acidification [ ]. Moreover, because fluoride- (F ) containing salts are common additives in dental products such as toothpaste, dental gels, and rinses, the buildup of F ions in crevices is inevitable. Therefore, assessing the crevice corrosion susceptibility of an implant alloy (e.g., Ti-6Al-4V) in F -enriched solutions is essential for predicting the implant’s durability. Although the Ti-6Al-4V alloy has been demonstrated as acutely susceptible to the presence of fluoride (F ) ions [ ], corrosion mechanisms in F -containing crevices remains unclear.

Following corrosion, the release of alloy elements and migration of corrosion-erosion debris are detrimental to the peri-implant biological system [ ]. In multi-phasic alloys, the release rate of the various corrosion by-products is often phase-dependent. Consequently, the composition and the speciation of the corrosion by-products rely on the corrosion susceptibility of the constituent phases. In TAV alloy, for example, distinct corrosion resistances of its two constituent phases, namely α and β phases, have been observed due to the partitioning of the alloying elements Al and V, respectively [ ]. However, the electrochemical basis of such phase-dependent degradation is not well understood. For instance, researchers have identified either the α phase [ ] or the β phase [ ] as being more susceptible to corrosion in various environments. Therefore, this study elaborates upon possible degradation processes of a failed commercial implant. The identified degradation mechanism is supported by electrochemical analyses (i.e., potentiodynamic (PD) polarization and electrochemical impedance spectroscopy (EIS)) of TAV alloy in simulated F -containing oral solutions. Further, corrosion susceptibilities of the α and β phases was resolved by analysis of surface evolutions, and solution composition. The findings provide insights into the cause-of-failure of the commercial implant and thereby suggest possible means for inhibiting corrosion and related degradation.


Sample preparation

Failure analysis was performed on a commercial implant assembly recovered from a patient who suffered from bone loss. The assembly, which consists of a commercially pure (CP) Ti base and a Ti-6Al-4V abutment, had remained in function for approximately ten years prior to removal. The bone loss was detected by a routine X-ray examination, signifying the early-stage peri-implantitis. The patient did not exhibit any other risk factors commonly known to cause the peri-implantitis (e.g., smoking, diabetes, or the use of bisphosphonates). Prior to failure analysis, the assembly was dissembled and ultrasonically cleaned in methanol and then deionized (DI) water (>18 MΩ).

For corrosion analysis, CP-Ti (ASTM Grade 2, Titanium Processing Center) and Ti-6Al-4V alloy (ASTM Grade 5, Dynamet) rods were solution annealed using a vacuum tube furnace at 960 °C for two hours, slow-cooled (1 °C/min) to 700 °C, fast-cooled (40 °C/min) to 540 °C, and finally furnace-cooled to room temperature (around 23 °C). Such heat treatment was performed to obtain a coarse-grained microstructure [ ] that allows differentiation between the α-phase and the β-phase. The heat-treated rods were then sectioned to smaller specimens with dimensions of ϕ 3 mm × 7 mm (diameter × height) for CP-Ti and ϕ 4.5 mm × 7 mm for Ti-6Al-4V alloy. Thereafter, electrodes were fabricated by attaching a lead wire to the back of each specimen and then mounting the ensemble in epoxy. The exposed face of the specimen was progressively polished to a 0.05-μm finish using a colloidal silica slurry.

Dulbecco’s phosphate-buffered saline (DPBS, Thermo Fisher Scientific) was used to simulate a near-neutral pH body fluid, and the Fusayama-Meyer artificial saliva (AS, Pickering Labs), titrated to pH 5 and pH 3 using HCl, was used to represent acidic oral solutions. The solution compositions are provided in Table 1 . To assess the effect of F on corrosion, 100 mM NaF salt (ACS Reagent Grade) was added to all solutions. Immediately (within a few seconds) prior to the electrochemical analyses, the solutions were filtered using a cellulose filter (0.1 μm pore size) to remove any precipitates that may have formed.

Table 1
The compositions of solutions used to simulate oral environments.
NaCl KCl CaCl 2 MgCl 2 Na 2 HPO 4 NaH 2 PO 4 KH 2 PO 4 Na 2 S CO(NH 2 ) 2 (Urea) pH
DPBS 137.931 0.495 0.901 0.495 8.099 1.471 7.4
AS5 6.838 5.369 6.163 3.876 0.021 16.667 4.9 a
AS3 6.838 5.369 6.163 3.876 0.021 16.667 3.0 a

a pH adjusted by the addition of HCl.

Metallurgical characterization

The microstructure and phase compositions of Ti alloys were characterized by scanning electron microscopy (SEM, FEI NOVA 230 NanoSEM) and energy dispersive X-ray emission spectroscopy (EDS, Thermo Scientific UltraDry). The surface morphology was characterized using VSI (Zygo NewView 8200). To enable high-resolution imaging of micrometer-sized surface features, a 100× Mirau objective (numerical aperture, N.A. = 0.85) was used providing lateral ( x–y pixel size) and axial resolutions of 90 nm and 2 nm, respectively.

Electrochemical characterization

Electrochemical analysis was used to assess the corrosion behavior of the alloys. First, crevice corrosion was simulated on the heat treated Ti-6Al-4V electrode with a polished polytetrafluoroethylene (PTFE) crevice former covering half of the alloy surface. The ensemble was then potentiostatically held at 0.7 V Ag/AgCl in the DPBS NaF (DPBS + 0.1 M NaF) solution for 3 h to facilitate crevice corrosion. Such a high potential represents an aggressive oral environment with high oxidants concentrations. Secondly, potentiodynamic polarizations was conducted following ASTM G5-14 [ ], wherein first, the electrodes are cathodically cleaned and acclimated for one hour in the testing solutions, and then polarized from −0.25 V vs. OCP to 1.6 V Ag/AgCl at a rate of 0.167 mV/s. Finally, the immersion tests were coupled with electrochemical impedance spectroscopy (EIS) using a ±10 mV excitation voltage over the frequency range from 10 −3 (or 10 −2 for NaF-containing solutions) to 10 5 Hz. The acquired EIS results were examined using the Kramers-Kronig method [ ] and then analyzed by equivalent circuit modelling using ZsimpWin® (ver. 3.60).

A potentiostat (Princeton Applied Research VersaSTAT 4) equipped with a 20 mL microcell was used to perform all electrochemical measurements using the alloys as the working electrode, a platinum (Pt) wire as the counter electrode, and a Ag/AgCl reference electrode. The electrochemical analyses were conducted at room temperature (22 ± 1 °C) in 15 mL solutions. The compositions of solutions containing liberated corrosion products were determined using inductively coupled plasma-optical emission spectrometry (ICP-OES, PerkinElmer Avio 200).

Results and discussion

Failure analysis of the commercial implant assembly

The degradation of the commercial implant assembly was examined using microscopy. The exterior surfaces of both the CP-Ti base and the TAV abutment show no evident discoloration and abrasion damage ( Fig. 1 a). But, the surface of the TAV abutment inside the crevice (see Fig. 1 ) is pronouncedly deteriorated. This crevice was embedded in the gingival tissue adjoining to the bone, and was accessible to both body and oral fluids. The observed deterioration suggests that the accumulation of aggressive ions within the crevice promoted corrosion and resulted in the release of corrosion by-products. A large number of insoluble products (see dark particles in Fig. 1 c) can be found near the crevice rim. These Al-Ti-oxide particles feature a composition of 80 atom % Al-oxide, 20 atom % Ti-oxide. The scratches that are adjacent to the oxides ( Fig. 1 c) are indicative of mechanical wear caused by the hard oxide particles.

Fig. 1
(a) A photograph of a commercial implant which failed because of bone loss, showing the location of the crevice between the CP-Ti base and the TAV abutment, (b) An optical image along the crevice, showing a severely deteriorated TAV surface, and (c) A SEM image of the crevice, showing corrosion products identified as Al-Ti-oxides.

Because the interior surfaces along the crevice were altered by mechanical wear, the microstructural influences on crevice corrosion cannot be verified in the commercial implant. Therefore, to better understand these effects, crevice corrosion was reproduced by electrochemically reacting the solution-annealed TAV alloy in a DPBS NaF solution. The surface analysis of corroded TAV using VSI showed that crevice corrosion resulted in a morphology that varied with the distance from the crevice rim (i.e., the crevice depth) ( Fig. 2 ). The morphologies and compositions of the two constituent phases, α and β, were identified (see Table 2 ) [ ]. Two distinct regions characterized by the preferential corrosion of either phase are evident. In Region 1, the surface recessed significantly (up to 1 μm) wherein the β phase corroded preferentially relative to the α phase, whereas in Region 2, surface recession was minor (less than 100 nm), although the α phase had corroded preferentially. The transition between the two regions is gradual. In addition, the island-like particles, which are located primarily in Region 1, were identified as Al-Ti-oxides having a composition (50 atom % Al-oxide, 50 atom % Ti-oxide) that is nominally similar to that observed in the commercial implant (see Fig. 1 c).

Fig. 2
VSI images of the solution-annealed TAV alloy wherein crevice corrosion is electrochemically stimulated in DPBS NaF solution to better understand the origin and nature of degradation that was observed in a realistic dental implant (shown in Fig. 1 ).

Table 2
The compositions of the constituent phases in the TAV alloy, in units of atom %. a
Phase Ti Al V
α phase 86.0 ± 0.5 13.2 ± 0.6 0.8 ± 0.2
β phase 82.0 ± 1.7 5.2 ± 1.1 12.8 ± 0.8

a The compositions reported are the average and standard deviation obtained from discrete measurements at five different locations for each phase.

EDS analysis ( Table 2 ) reveals that the α phase is enriched in Al and depleted in V relative to the β phase. The distinct recession rates of α and β phases implies that there are mismatches in the corrosion resistance of these phases resulted from alloying elements. However, this does not fully explain the evolution of corrosion morphology vis-à-vis the crevice depth. Since local environments as given by localized electrochemical potentials and the solution composition may determine the corrosion susceptibilities of α and β phases, we evaluated the electrochemical corrosion of the TAV alloy under both potentiodynamic and “potential-free” conditions.

Alloy corrosion under potentiodynamic conditions

Potentiodynamic (PD) polarization curves were obtained for solution-annealed TAV and CP-Ti alloy surfaces immersed in simulated oral solutions (see Fig. 3 ). In NaF-free solutions, the PD curves exhibited spontaneous passivation with passivation currents (I pass ) of around 5 × 10 −7 A/cm 2 for both alloys. Except for the somewhat lower corrosion potentials (E corr ), the PD curves obtained in the DPBS NaF solution were largely similar to that obtained in the absence of NaF. The addition of NaF in acidified artificial saliva (AS5 + 0.1 M NaF (AS5 NaF) and AS3 + 0.1 M NaF (AS3 NaF)) had a significant effect on the PD curves, as shown in Fig. 3 . Both the TAV and CP-Ti alloys exhibited evident passivity breakdown, particularly in AS3 NaF, as signaled by the decrease in corrosion potentials and the increase in corrosion currents. Additionally, the AS3 NaF curves exhibit an “active-passive-transpassive” transition at the potential range from −1.4 to 0 V Ag/AgCl . This suggests that the oxidation states of the alloying elements vary under different applied potentials.

Fig. 3
Representative potentiodynamic curves for: (a) TAV and (b) CP-Ti measured in DPBS, AS5, and AS3 solutions ( Table 1 ) both with (solid lines) and without (dashed lines) 0.1 M NaF as an additive to the solutions.

After polarization to 1.6 V Ag/AgCl only a negligible change in the surface morphology was observed for surfaces reacted with the DPBS NaF solution, whereas more evident changes were observed on those exposed to AS3 NaF and AS5 NaF solutions ( Fig. 4 ). Both solutions resulted in a morphology resembling Region 1 in Fig. 2 . Specifically, a phase-dependent degradation process was observed wherein the β phases became more recessed than the α phase by about 200 nm. SEM images revealed the persistence of nanostructured α lamellae within the corroded β region (see the “transformed β region” in Fig. 4 d). The residual V content in the corroded β regions is < 1 atom %, suggesting a preferential release of V relative to Ti and Al at anodic potentials (i.e., more oxidizing conditions).

Fig. 4
Surface topography as visualized using vertical scanning interferometry for: (a) pristine TAV alloy, (b) the same region following polarization to 1.6 V Ag/AgCl in AS3 NaF. (c) and (d) SEM images of the TAV surface after polarization to 1.6 V Ag/AgCl in AS5 NaF.

Addition of equimolar amounts of NaF to the different solutions (DPBS, AS3, AS5) resulted in distinct corrosion behaviors, indicating that the F concentration is not the sole factor governing the corrosion response. To better rationalize the alloys’ electrochemical response to the fluoride addition, the speciation of ions in solutions was examined using PHREEQC with the minteq.v4 database [ ]. The calculated F activities, {F } were similar within 3 mM across all NaF-containing solutions, as noted in Table 3 . Significant corrosion of the alloys occurred only in acidified NaF solutions, suggesting that {F } and {H + } contribute collectively to the passivity breakdown. Hence, the aggressiveness of the solution can be represented by the HF 0 activity, {HF 0 }, in line with previous observations [ , ]. The DPBS NaF solution features negligible {HF 0 }, consistent with its insignificant impact on the alloys’ passivation under anodic potentials. Notably, HF 0 can also be produced inside the crevice via a localized acidification process [ , ], resulting in a severe degradation of the TAV alloy in an otherwise neutral solution (Section 3.1 ).

Table 3
The calculated ionic speciation in the solutions and the measured and calculated pHs for the solutions containing 0.1 M NaF.
Solution Precipitates F concentration (mM) F activity * 10 −3 HF 0 activity
* 10 −3
pH (simulated) pH (measured)
DPBS NaF CaF 2 98.2 63.0 0.003 7.47 7.03
AS5 NaF CaF 2 87.7 64.8 0.216 5.65 5.37
AS3 NaF CaF 2 87.7 62.0 1.233 4.88 4.89

Alloy corrosion under “potential-free” conditions

“Potential-free” tests were also conducted on the TAV alloy in a AS3 NaF solution, which has the highest HF° activity (1.2 mM), and in a fluoride-free AS3 solution. As shown in Fig. 5 , The alloy’s open circuit potential (OCP) gradually increased from −0.4 V Ag/AgCl to −0.2 V Ag/AgCl during immersion in AS3 solution indicating the development of a robust passive film. In contrast, following immersion in AS3 NaF solutions, the alloy’s OCP dropped to about −1.2 V Ag/AgCl (OCP1 in Fig. 5 a) signifying active corrosion. The OCP increased gradually thereafter until after 20 h, when a sudden rise to around −0.7 V Ag/AgCl (OCP2 in Fig. 5 a) was observed followed by fluctuations in OCP. The elevation of OCP in AS3 NaF was concurrent with the decrease in {HF°} resulting from cathodic hydrogen evolution. In turn, the solution pH increased from 4.89 (as prepared AS3 NaF) to a critical value of 5.09 after 20 h of immersion, just before the OCP increased sharply. At the critical pH (or equivalently, critical {HF°}), the surface evolved from one undergoing active corrosion to one that is somewhat passivated (i.e., in an “active-passive” transition).

Fig. 5
(a) OCP evolution of the TAV alloy immersed in AS3 and AS3 NaF solutions. EIS was conducted periodically during immersion, and the regressed charge-transfer resistances are plotted in (b) and (c).

In order to resolve the electrochemical differences between phases, EIS was carried out periodically over the course of immersion. The passivated TAV alloy in AS3 solution can be represented by a simple circuit. In this circuit, R s , which the denotes the solution resistance, is in series with the R O , the charge-transfer resistance, paralleled with a constant phase element, CPE O , which denotes the capacitance of the dielectric passive layer. The persistence of a single relaxation time constant indicates that the passive films formed on the α and β phases are similar and compatible. Moreover, the large magnitude of R O , and its increase over time suggest substantial, and increasing corrosion resistance of the passive layer ( Fig. 5 b). On the other hand, active corrosion of the alloy in AS3 NaF solution resulted in two distinct time constants, revealing the electrochemical mismatch between the two phases. The charge transfer interactions at the α and β surfaces are characterized by R α and CPE α , and R β and CPE β in a circuit as: R s [(CPE α R α )(CPE β R β )]. A decrease in {HF°} led to increases in both R α and R β (see Fig. 5 c) with significant increases occurring following the “active-passive” transition after 20 h of immersion.

The electrochemical mismatch identified by EIS was consistent with the surface morphology of the TAV alloy following 45 h of immersion in AS3 NaF solution. Interestingly, the surface features resembles that in the Region 2 of the crevice-corroded surface ( Fig. 2 ), which is characterized by the preferential dissolution of the α phase and the endurance of β phase ( Fig. 6 (a–d) ). The nano-lamellar α phase present within the transformed β regions was also preferentially corroded (see Fig. 6 d). Analysis of larger areas (200 μm × 200 μm) indicates that the unreacted (polished) surface featured a unimodal distribution with a full width at half maximum (FWHM) of 50 nm ( Fig. 6 e). On the other hand, immersion in AS3 NaF solution produced a bimodal distribution in which the α phase regions had recessed by about 700 nm relative to the β phase regions (see Fig. 6 f).

Fig. 6
Representative surface topography of: (a) as-polished TAV alloy, and (b) the same area following 45 h of immersion in the AS3 NaF solution. (c) and (d) are SEM images of the area shown in (b). (e) Shows the surface height distribution for a pristine surface prior to immersion. (f) Shows the surface height distribution of the same area following immersion. In (e–f), the lowest point of the as-polished surface is established as the “0/reference” location.

How the magnitude of the applied potential affects degradation of the biphasic alloy?

The reversal in the phase dependence of TAV corrosion under potentiodynamic and potential-free conditions suggests that the degradation kinetics are altered by anodic excitation. To further investigate how the magnitude of applied potential affects the rate of the phase-specific corrosion, two additional PD polarizations up to lower potentials of 0.55V Ag/AgCl and 0.65V Ag/AgCl were carried out in the AS3 NaF solution. Consistent with the transition region of the crevice-corroded surface ( Fig. 2 ), it is evident that the β regions persisted when the alloy was immersed but progressively recessed relative to the alpha regions when the alloy was polarized to increasing potentials ( Fig. 7 ). Further, the Al-Ti-oxides, which were absent on the potential-free surface formed within the α grains at anodic potentials.

Fig. 7
Representative surface topography images and line profiles of alloy surfaces corroded in the AS3 NaF solution after (from top to bottom) immersion for 45 h, and following polarization to 0.55V Ag/AgCl , 0.65V Ag/AgCl , and 1.6 V Ag/AgCl , respectively.

Such transitions in the phase-specific corrosion rates are likely due to changes in oxidation states of the alloy elements, signified by the “active-passive-transpassive” transition in the PD curve. Based on the Pourbaix diagrams for Ti, Al, and V [ ], Al 2 O 3 is the most stable form of Al over the entire range of potentials applied herein [ ], whereas Ti and V could demonstrate multiple oxidation reactions over the evaluated potential range ( Table 4 ) [ , ]. The equilibrium potential of these reactions, E’, was calculated from the Nernst equation given by:

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Apr 27, 2020 | Posted by in Dental Materials | Comments Off on Elucidating the corrosion-related degradation mechanisms of a Ti-6Al-4V dental implant
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