Journal of the Australian Ceramic Society (2019) 55:187–200
https://doi.org/10.1007/s41779-018-0224-1
RESEARCH
In vitro bioactivity and corrosion resistance enhancement of Ti-6Al-4V
by highly ordered TiO2 nanotube arrays
M. Sarraf 1,2 & N. L. Sukiman 2 & A. R. Bushroa 1 & B. Nasiri-Tabrizi 3
W. J. Basirun 7
&
A. Dabbagh 4,5 & N. H. Abu Kasim 6 &
Received: 6 June 2017 / Revised: 2 November 2017 / Accepted: 6 June 2018 / Published online: 26 July 2018
# Australian Ceramic Society 2018
Abstract
In the present study, the structural features, corrosion behavior, and in vitro bioactivity of TiO2 nanotubular arrays coated on Ti–6Al–
4V (Ti64) alloy were investigated. For this reason, Ti64 plates were anodized in an ammonium fluoride electrolyte dissolved in a
90:10 ethylene glycol and water solvent mixture at room temperature under a constant potential of 60 V for 1 h. Subsequently, the
anodized specimens were annealed in an argon gas furnace at 500 and 700 °C for 1.5 h with a heating and cooling rate of 5 °C min−1.
From XRD analysis and Raman spectroscopy, a highly crystalline anatase phase with tetragonal symmetry was formed from the
thermally induced crystallization at 500 °C. Besides, the Ti 2p3/2 and Ti 2p1/2 binding energies showed the presence of the Ti4+
oxidation state. According to the in vitro bioassay, the modified surface proved its outstanding capability in enhancing the bioactivity, where a thick layer of bone-like apatite was formed on the annealed TiO2 nanotube surface. In addition, the corrosion
measurements indicated that the corrosion protection efficiency increased remarkably and reached 87% after annealing at 500 °C.
Keywords TiO2 nanotubes . Anodization . Ti–6Al–4V . Corrosion resistance . In vitro bioactivity
Introduction
The utilization of metallic implants for medical applications
can be traced back to the nineteenth century, leading up to the
* A. R. Bushroa
bushroa@um.edu.my
3
New Technologies Research Center, Amirkabir University of
Technology, Tehran, Iran
4
School of Medicine, Faculty of Health and Medical Sciences,
Taylor’s University, Subang Jaya, Malaysia
M. Sarraf
masoudsarraf@gmail.com
5
Department of Materials Science and Engineering, Sharif University
of Technology, Tehran, Iran
Centre of Advanced Manufacturing and Material Processing,
Department of Mechanical Engineering, Faculty of Engineering,
University of Malaya, Kuala Lumpur, Malaysia
6
Department of Restorative Dentistry, Faculty of Dentistry, University
of Malaya, Kuala Lumpur, Malaysia
7
Department of Chemistry, Faculty of Science, University of Malaya,
50603 Kuala Lumpur, Malaysia
* B. Nasiri-Tabrizi
bahman_nasiri@hotmail.com
1
2
era as the metal industry began to develop during the
Industrial Revolution [1, 2]. Recently, increasing research efforts in metallic biomaterials have been invested for applications in non-conventional reconstructive surgery of hard tis-
Centre of Advanced Materials, Department of Mechanical
Engineering, Faculty of Engineering, University of Malaya, Kuala
Lumpur, Malaysia
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J Aust Ceram Soc (2019) 55:187–200
sues/organs, such as the application of NiTi shape memory
alloys as vascular stents and the development of new
magnesium-based alloys for bone tissue engineering and regeneration [3, 4]. In addition, numerous attempts have been
made to improve the surface properties of common metallic
implants such as titanium-based implants [5, 6].
Titanium and its alloys are commonly utilized in biomedical applications, particularly as hard tissue implants due to
their desired features, such as good corrosion resistance, biocompatibility, fatigue strength, machinability, and formability
as well as relatively low modulus. Despite their outstanding
properties, the application of titanium-based implants is still
limited as they are unable meet all the clinical requirements.
Hence, in order to improve their biofunctionality, surface
modification is often performed [7]. In general, the instinctive
formation of TiO2 passive layer provides the desired inertness
and biocompatibility [8]. However, it has been reported that
this layer is easily damaged under wearing conditions which
results in the release of wear debris and unfavorable ions into
the biological media [9]. Therefore, the formation of metallic
oxide nanotubular surfaces such as TiO2 nanotubes on metallic implants to improve their biomedical effects is much desired. This surface modification for orthopedic implants is
usually performed by electrochemical anodization, where
self-organized nanotubular oxide configurations can easily
be created and controlled by changing the anodization conditions [10]. The formation of self-ordered TiO2 nanotubes can
be carried out by simple electrochemical anodization of a
titanium-based substrate under specific conditions [11–16].
This process is influenced by the control parameters such as
anodization voltage, period, and concentration of acid. The
formation of TiO2 nanotubular arrays in aqueous-based solutions occur by two common steps: (i) hydrolysis of titanium
ions and (ii) chemical dissolution of oxide layer at the oxide/
electrolyte interface. Upon anodization, an oxide film is generated initially as a result of the interplay between the Ti ions
and anionic oxygen in the solution, followed by the homogeneous spreading across the metal surface. At the anode surface, the Ti4+ ions are liberated due to the oxidation of the
metal, while the dominant reaction at the cathode is the release
of hydrogen gas [17].
Ti þ 2H2 O⇌TiO2 þ 4Hþ þ 4e–
ð1Þ
4Hþ þ 4e– →2H2
ð2Þ
interfaces. Consequently, the oxide layer partially and sometimes completely dissolves in the presence of F− anions. The
TiO2 dissolution process occurs at the oxide/electrolyte interface which results in the formation of [TiF6]2− according to
Eq. (4) as follows:
TiO2 þ 6Hþ þ 6F− ⇌½TiF6 2− þ 2H2 O þ 2Hþ
ð4Þ
A schematic illustration of the one-pot anodization process
and the development of the nanotubular arrays under a steady
anodization voltage is shown in Fig. 1 [18]. From Fig. 1a, with
the anodization onset, a thin layer of oxide film develops on the
Ti64 surface. As shown in Fig. 1b, the formation of small pits in
the oxide layer occurs as a result of the localized dissolution of
the oxide, causing the barrier layer at the bottom of the pits to
become relatively thin which in turn increases the intensity of
the electric field across the remaining barrier layer. This effect
results in further pore growth as illustrated in Fig. 1c. It has been
proven that the pore entrance is not affected by the electric field
assisted dissolution. Therefore, it stays relatively narrow as the
electric field distribution in the curved bottom surface of the
pore causes pore widening, as well as the deepening of the pore.
The consequence of this process is a pore with a scallop configuration [19]. Given the high bonding energy of Ti–O
(323 kJ mol−1), here it is reasonable to presume that only pores
having thin walls can be developed due to the relatively low ion
mobility and relatively high chemical stability of the oxide layer
in the electrolyte. Thus, the un-treated metallic surface can initially exist between the pores. When the pores grow deeper, the
electric field in these protruded metallic areas increases, thus
enhancing the field assisted oxide growth and oxide dissolution.
Accordingly, the general reaction can be introduced by Eq.
(3) as follows:
Ti þ 2H2 O⇌TiO2 þ 2H2
ð3Þ
In addition to the anodic oxidation of Ti for the preparation
of TiO2 nanotubes which was described above, the development of the nanotubular arrays in F− electrolyte promotes the
chemical dissolution of TiO 2 at the oxides/electrolyte
Fig. 1 A schematic illustration of the one-pot anodization process and the
development of the nanotubular arrays under a steady anodization voltage: a oxide layer formation, b pit creation, c pit growth, d oxidation and
field assisted dissolution of the metallic region between the pores, and e
fully developed nanotubular configurations f with a corresponding top
view. Redrawn from [18]
J Aust Ceram Soc (2019) 55:187–200
Accordingly, the simultaneous formation of well-defined pores
and inter-pore voids also occurs (Fig. 1d). After this, an equilibrium is established between the void growth and the tube
enlargement. The length of nanotubes increases until the rate
of electrochemical etching is equal to the rate of chemical dissolution on the surface of the nanotube. Beyond this point, the
length of nanotubes will be independent from the anodization
time, as determined for a given electrolyte concentration and
anodization potential.
Given that a quick ingrowth of orthopedic implants in the
bone is required and the apatite formation is imperative for
osseointegration, a vital aspect is to rapidly stimulate bonelike apatite formation from body fluid [20]. There are several
studies that have focused on the improvement of apatite formability of metallic implants using various chemical and physical treatments [21, 22]. Therefore, it is imperative to examine
the formation and growth of self-organized TiO2 nanotubular
surfaces, in view of the hydroxyapatite-induced behavior for
biocompatibility approach. The beneficial effect of these
nanotubular structures is optimal for the embedding of precursors during apatite formation, which also promotes the apatite
nucleation and accelerates its growth. Oh et al. [23] reported
the formation of a very thin layer (~ 25 nm) of nanoscale
hydroxyapatite phase on TiO2 nanotubes after immersion in
SBF for a week. In this regard, Tsuchiya et al. [24] provided a
systematic assessment of the apatite formation on TiO2
nanotubular arrays for the first time. They found that to form
a uniform and thick apatite layer, it is highly desirable to
fabricate nanotubular arrays with larger opening diameters
and deeper depth for calcium and phosphorous nucleation
and growth where a minimum opening diameter of 15 nm is
needed to form the apatite layer by immersion of the TiO2
nanotubes in simulated body fluid (SBF). According to the
literature, a minimum of 14 days is required to achieve a
bone-like apatite layer more than 1 μm thickness on TiO2
nanotubes, as compared to the bare substrate [10]. It was also
found that the phase composition and crystallinity degree of
the nanotubular structures affect the hydroxyapatite coating
formation [10, 24]. On the other hand, from the corrosion
point of view, the implanted materials are in direct contact
with extracellular body fluids which could promote corrosion
[25, 26]. Considering the formation of a native TiO2 film
which protects the metal from further oxidation, titaniumbased alloys have high stability and high in vitro corrosion
resistance [27–29]. Though, this protective oxide layer is thermodynamically stable; nevertheless, metal containing species
can still be released through a passive-dissolution mechanism
[30]. Therefore, the enhancement of corrosion resistance of
these alloys is very important. It was reported that the
nanoporous titania showed excellent bioactivity and corrosion
resistance than the bare substrate after immersion in Hank’s
solution for 7 days, which was probably due to the growth of
HAp layer on their surface [31]. Besides, Huang et al. [32]
189
recently found that the TiO2 nanotubes developed on 316 L
stainless-steel substrate provides improved corrosion resistance and hemocompatibility for stainless-steel implants.
Despite the success of modern implant therapy, implant in
biomedical applications is still restricted by a range of risk factors
such as implant biofunctionality, age, insufficient quantity, and
quality of the host bone and other systemic circumstances [33].
Hereon, the structural and morphological features, corrosion protection efficiency, as well as in vitro apatite formability of the
nanotubular configurations were reappraised as a complementary
study. For this purpose, the surface modification of biomedical
grade Ti64 alloy was performed by electrochemical anodization
technique using a DC power source. Thermal treatment was
performed after the anodization process to crystallize the nanotubes and improve the adhesion of the coating.
Materials and experimental methods
Preparation of the substrate for anodizing
Ti64 plates (15 mm × 15 mm × 2 mm, E Steel Sdn. Bhd,
Malaysia) were used as substrates. The specimens were initially polished using SiC emery paper (800–2400 grit) each
size for 2 min, followed by wet polishing in a diamond slurry.
The polished samples were then sonicated in acetone for
10 min at 40 °C. Finally, all the specimens were washed three
times with distilled water and dried at 100 °C for 1 h.
Fabrication of nanotubular layer
The prepared samples were anodized in ammonium fluoride
(NH4F, Sigma-Aldrich CO., 0.35 wt%) electrolyte dissolved
in a 90:10 ethylene glycol (EG, J.T. Baker CO.) and water
mixture at room temperature, using a DC power source
(Model E3641A, Agilent Technologies, Palo Alto, USA).
The substrate was connected to the positive terminal (anode),
and a graphite rod (D = 5 mm) was connected to the negative
terminal (cathode) of the power source. In all experiments, the
distance between the cathode and anode was fixed at roughly
20 mm. The electrochemical anodization was performed under a steady potential of 60 V for 1 h. Finally, the as-anodized
samples were annealed in an argon gas furnace at 500 and
700 °C for 1.5 h with a heating and cooling rate of 5 °C min−1.
Characterization methods
Phase composition and morphological variations
X-ray diffraction (XRD) measurements were performed using
a PANalytical Empyrean system (Netherlands) with Cu–Kα
radiation over a 2θ range from 20° to 80°. To examine the
XRD patterns, BPANalytical X’Pert HighScore^ software
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J Aust Ceram Soc (2019) 55:187–200
was used and the profiles were compared to standards compiled by the Joint Committee on Powder Diffraction and
Standards (JCPDS). The morphological variations of the coating were investigated using a high-resolution FEI Quanta 200
F field emission scanning electron microscope (FESEM).
Raman spectroscopy and XPS analysis
The X-ray photoelectron spectroscopy (XPS) of the specimens was performed using ULVAC-PHI Quantera II equipped
with monochromatic Al Kα X-ray source (hν = 1486.6 eV).
Raman spectroscopy was performed using Renishaw in Via
Raman microscope with a laser wavelength of 514 nm to
observe the structural information of the specimens.
each one was fully dissolved in 700 ml of water. A total of
40 ml of 1 M HCl solution was used to adjust the pH during
the SBF preparation. A 15-ml aliquot of this amount was added
followed by the addition of CaCl2·2H2O, and the second portion of the HCl solution was used in the remainder of the titration procedure. Following the addition of the (CH2OH)3CNH2,
the solution temperature increased from ambient to 37 °C.
Subsequently, the solution was titrated with 1 M HCl to
pH 7.4 at 37 °C. Moreover, the solution was diluted progressively with successive additions of deionized water during the
titration process to reach a final volume of 1 l. Lastly, the
500 °C annealed sample was immersed in SBF solution for
14 days. To observe the apatite layer and superficial changes,
the immersed samples were rinsed in ethanol and then analyzed
by FESEM/EDS.
Corrosion measurement
Corrosion analysis was carried out by means of a potentiostat/
galvanostat AutoLab PGSTAT30 from Ecochemie (the
Netherlands) outfitted with a frequency response analyzer
(FRA). The polarization experiments were performed using
a three-electrode arrangement where the bare undercoat, asanodized coating and annealed specimen were the working
electrodes, with a surface of 0.5 cm × 0.5 cm exposed into
the phosphate buffer solution (PBS, pH 7.2) for 30 min. A
platinum wire and saturated calomel electrode (SCE) were the
counter and reference electrodes, respectively. From the Tafel
plots, the corrosion parameters, i.e., the corrosion potential
(Ecorr; VSCE), the corrosion current density (Icorr; μA cm−2),
and polarization resistance (Rp; Ω cm−2), were obtained.
Besides, the corrosion protection efficiency (P.E.) was determined using the following equation [34]:
P:E:ð%Þ ¼
I 0corr −I ccorr
100
I 0corr
ð5Þ
where I 0Corr and I cCorr are the corrosion current of the bare Ti64
and the coated sample, respectively.
Apatite formability in SBF
To assess the apatite formability of the samples, two approaches
of the SBF solution were made. In the first approach [35], the
desired precursors were dissolved in distilled water and the pH
was adjusted to 7.3. Subsequently, the 500 °C annealed sample
was immersed in 10 ml SBF solution and held for 1 to 14 days
at a steady temperature of 37 °C and pH 7.3. In the second
approach, reagent grades of NaCl, Na2HPO4·2H2O, Na2SO4,
NaHCO3, CaCl2·2H2O, MgCl2·6H2O, (CH2OH)3CNH2, KCl,
and HCl were used in the preparation of the SBF media, in
accordance with the method explained by Tas [36]. In this regard, desired amounts of the above chemicals were dissolved in
deionized water. The precursors were added consecutively as
Results and discussion
Morphological variations
As the morphological features of the nanotubular structures
affect the biofunctionality of the metallic implants, the microstructural description is much needed [15]. Figure 2 demonstrates the FESEM micrographs of the substrate after one-step
anodization for 1 h in 0.35% NH4F electrolyte solution (90
EG:10 water) under a steady potential of 60 V. As described
above, during the initial steps of anodization, irregular pits were
formed by the localized dissolution of oxide film. This is
followed by the pits conversion to larger pores while most of
the areas were still covered by the oxide layer. In this regard,
Beranek et al. [37] proposed that the pore formation occurs
primarily at random locations, and self-ordering is merely a
product of the competition between the growing pores. On
the other hand, Raja et al. [38] suggested that the ordering of
pores may be a result of local surface perturbations.
From Fig. 2a, b, substantial alterations from nanoporous to
nanotubular structure were observed after 1 h of anodization.
It is clear that highly ordered TiO2 nanotubes were formed
after 1 h of anodization, where the nanotubular arrays are
uniformly distributed over the anodized surface. According
to the higher magnification FESEM micrograph (Fig. 2b),
the average inner diameter of the nanotubes is 120 nm. Due
to the high electric field across the electrode, electric fieldassisted dissolution prevails over chemical dissolution during
the initial stages of anodization. When the anodization proceeds and the oxide layer thickens, the chemical dissolution
predominates over the field-assisted dissolution. Accordingly,
the size and density of the pores are notably increased. After
this, the enlargement and proliferation of the pores take place
through an internal motion at the oxide/metal interface, which
results in the formation of hollow cylindrical oxide particles
that develop into the nanotubular structure. From Fig. 2c, the
J Aust Ceram Soc (2019) 55:187–200
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Fig. 2 FESEM micrographs of
the substrate after one-step anodization for 1 h in 0.35% NH4F
electrolyte solution (90 EG:10
water) under a steady potential of
60 V. a, b Top view at different
magnifications. c Bottom view. d
Cross-sectional view
bottom of the TiO2 nanotubes illustrates a series of evenly
spaced Bbumps^ that signify the pore tips of each individual
nanotube [39, 40]. Since the barrier layer at the bottom of the
nanotubes is scalloped, this layer can be divided into the upper
and lower parts, which are known as the pure barrier layer and
interface barrier layer, respectively [17]. The upper part is
considered to be pure oxide, while the lower part consists of
a mixture of oxide and bare substrate. The oxide layer at the
bottom of the pores becomes thinner with time due to the
chemical dissolution process. When the thickness declines,
the electric field-assisted dissolution could occur again in this
region and, consequently, the pores could penetrate into the
substrate and the nanotubes grow increasingly longer. If the
anodization is conducted in fluoride concentrated solution, the
rate of attack could be faster; thus, the thinning process is also
faster. As the thickness decreases, the electric field-assisted
dissolution could reoccur at this region. Due to this process,
the pores penetrate into the Ti substrate and the nanotubes
grow deeper. Nevertheless, as the voltage is constantly applied, the anodization process could reoccur at the bottom of
the pores, which could lead to the formation of nanotubes with
closed bottoms. Therefore, the expansion of the nanotubular
arrays is influenced by the fluoride content in the bath and the
electric field dissolution degree. As the anodization voltage is
very high, the rate of electric field dissolution at the barrier
layer in the nanotubes could be higher; hence, longer nanotubes could be produced [41]. From the FESEM crosssectional micrograph in Fig. 2d, the average length of the
nanotubes is 843 nm.
To investigate the effect of subsequent heat treatment on
the microstructural evolution, thermal annealing was
performed at low heating and cooling rate of
5 °C min −1 at 500 and 700 °C for 1.5 h in an argon gas
furnace. Figure 3 displays the top view FESEM micrographs of the annealed samples after annealing at 500
and 700 °C.
As can be seen in Fig. 3a, highly ordered TiO2 nanotubular
arrays were formed during annealing at 500 °C for 1.5 h. From
Fig. 3b, there are no main alterations in the morphological
characteristics after annealing at this temperature and thus
the mean inside diameter of the nanotubes is 120 nm.
According to literature, the nanotubes might collapse
due to high temperature annealing and longer annealing
time [42]. As illustrated in Fig. 3c, the nanotubular arrays
were completely demolished after annealing at 700 °C
for 1.5 h. As a consequence of this unfavorable annealing, the TiO 2 nanotubes were transformed into a coarse
structure (Fig. 3d) which is consistent with previous reports [43].
Phase analysis
Figure 4 displays the XRD patterns of the bare substrate, the
as-anodized specimen and the 500 °C annealed sample, as
well as the tetragonal lattice constants and unit cell volume
of TiO2 nanotubes. As can be seen in Fig. 4a, the XRD
profile of the bare substrate illustrates only the diffraction
peaks of Ti (JCPDS#005-0682) located at 2θ = 35.1°, 38.4°,
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Fig. 3 FESEM micrographs of
the annealed samples at different
temperatures. a, b 500 °C. c, d
700 °C
40.2°, 53.1°, 63.1°, 70.6°, and 76.4°, which are attributed to
the (1 0 0), (0 0 2), (1 0 1), (1 0 2), (1 1 0), (1 1 2), and (2 0 1)
planes, respectively. As shown in Fig. 4b, an amorphous
coating was formed after 1 h of anodization. However, the
as-prepared coating was not composed of entirely amorphous nanotube structure. There are some characteristic
peaks attributed to the presence of TiO2 with anatase crystalline phase (JCPDS01-071-1166, crystal system: tetragonal; space group: I41/amd; space group number: 141) including (1 0 1) at 2θ = 25.22°, (0 0 4) at 2θ = 37.67°, and (1
0 5) at 2θ = 53.34°. A highly crystalline anatase phase was
present after the thermal treatment at 500 °C due to the
thermally induced crystallization. Consequently, new diffraction peaks including (2 0 0) plane at 2θ = 48.1°, (2 0
1) plane at 2θ = 55.2°, (2 0 4) plane at 2θ = 62.8°, (1 1 6)
plane at 2θ = 68.8°, and (2 1 5) plane at 2θ = 75.3° attributed
to the anatase phase became obvious in the XRD
diffractogram (Fig. 4c).
The tetragonal lattice parameters, i.e., a, b, and c, are normally specified using Eq. (6) as follows:
1
h2 þ k 2 l 2
¼
þ 2
a2
c
d2
ð6Þ
where h, k, and l are the Miller indices of the reflection planes.
In addition to the lattice parameters, the unit cell volume of (V)
was determined using Eq. (7):
V ¼ a2 c
ð7Þ
Besides, the in-plane and out-of-plane strains were
examined by comparing the experimental results of
the lattice parameters (a and c) with the unstrained
lattice constants, i.e., a0 and c0, using Eqs. (8) and
(9) [44]:
εa ¼
ða−a0 Þ
ðIn‐plane strainÞ
a0
ð8Þ
εc ¼
ðc−c0 Þ
c0
ð9Þ
ðOut‐of ‐plane strainÞ
where a0 and c0 for the standard TiO2 (JCPDS01-071-1166)
are 3.7842 and 9.5146 Å, respectively.
The standard crystal structure constants for anatase (#01071-1166), a and c, are 3.7842 and 9.5146 Å, respectively. As
shown in Fig. 4d, e, the lattice constants of the 500 °C
annealed sample along the a- and c-axes have changed to
3.7796 and 9.4706 Å, respectively. As a result of these
differences in the lattice parameters, the unit cell volume
shrunk from 136.25 to 135.29 Å3 after annealing at 500 °C
(Fig. 4f). The decrease in the unit cell volume is due to the
thermally induced lattice relaxation. Based on the obtained data, the in-plane and out-of-plane strains of the specimens are − 0.12 and − 0.46%, respectively. The low
amounts of in-plane and out-of-plane strains are most likely due to the small lattice mismatch after the thermal treatment at 500 °C. It should be noted that the negative
amount of strain shows that the in-plane strain in anatase
was compressive.
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Fig. 4 XRD patterns of the a bare substrate, b the as-anodized specimen and c the 500 °C annealed sample as well as the d, e tetragonal lattice constants
and f unit cell volume
XPS analysis
XPS was utilized to investigate the chemical species in the
coatings. The XPS spectra and high resolution of the C 1s,
Ti 2p, and O 1s regions of the 1 h anodized sample before and
after the thermal treatment at 500 °C are shown in Figs. 5 and
6, respectively. The C 1s regions in Figs. 5b and 6b are
deconvoluted into four peaks: the peak with the lowest binding energy (284.41 eV), represents the graphitic carbon phase,
followed by the three types of carbon–oxygen bonds of the
carboxyl group with the highest binding energy of 289.24 eV
[45]. The binding energies of Ti 2p3/2 and Ti 2p1/2 components
of the 1 h anodized sample are located at 458.8 and 464.4 eV,
showing the presence of the Ti4+ oxidation state [46–48].
Similar results are observed for the 500 °C annealed sample
(Figs. 5c and 6c). The deconvoluted oxygen XPS spectra are
shown in Figs. 5d and 6d. In addition to the primary peaks for
the C–O and O–C=O bonds (already detected in the C 1s
region), the O 1s region shows the presence of a single peak
at 529.77 eV which corresponds to the lattice oxygen of the
Ti–O bonds. The identification of Al 2s and Al 2p regions in
the XPS spectra confirms the migration of the alloy species to
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J Aust Ceram Soc (2019) 55:187–200
Fig. 5 a XPS spectra and high resolution of b C 1s, c Ti 2p, d O 1s, and e F 1S regions of the 1 h anodized sample
the coating during anodization and subsequent annealing. In
accordance with XPS data, the Ti/Al/O atomic ratio is
18.07:1.76:56.70 for the 1 h anodized sample and
21.27:4.62:66.33 for the 500 °C annealed specimen.
As shown in Fig. 5e, the F 1s signal was characterized by a
single band at 685.2 eV. It was found that the F 1s signal may
have two contributing bands. The lower binding energy
(684.0 eV) is attributed to the surface Ti–F species (Fsurf),
while the main band (690.0 eV, Fsub) is assigned to the lattice
F from the O–Ti–F moieties in the TiO2–xFx solid solutions,
due to the nucleophilic substitution by the F− ions. The predominance of the Fsub component is typical of the rutile phase,
while the anatase or multiphase TiO2 show the predominance
of the Fsurf species [49]. Hereon, the detection of F 1s region
shows the presence of a high-level organic/halide content in
the tubes, where the loss by decomposition during thermal
annealing could destroy the nanotubes. However, from the
FESEM images of the annealed samples, the TiO 2
nanotubular arrays remained stable up to 500 °C. This stability
can be attributed to the surface-fluorinated Ti–F species,
where its removal is not accompanied by dramatic structural
changes. On the contrary, due to the partial or complete transformation from the anatase to rutile phase, the Fsub component
(the presence of lattice Ti–F bonds) is predominant after
annealing at 700 °C [50, 51]. Therefore, significant changes
are expected due to the removal of lattice F during
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195
Fig. 6 a XPS spectra and high resolution of b C 1s, c Ti 2p, and d O 1s regions of the 1 h anodized sample after thermal treatment at 500 °C
annealing at 700 °C for 1.5 h. This argument is in accord a nc e w i t h t h e F E S E M ob s e r v a t i on s , w h er e t h e
nanotubular arrays were completely destroyed after annealing at 700 °C for 1.5 h.
Raman spectral analysis
Generally, useful information can be extracted from the
Raman spectra, for instance, the crystallinity degree, the structure, and composition of the product from the molecular vibration modes. According to the literature [52], TiO2 has six
or four transition spectrum for the anatase and rutile phases,
respectively. The Raman active modes of anatase are A1g
(519 cm−1), B1g (399 and 519 cm cm−1), and Eg (144, 197,
and 639 cm cm −1 ). In addition, A 1g (612 cm −1 ), B 1g
(143 cm−1), B2g (826 cm−1), and Eg (447 cm−1) are the
Raman active modes of the rutile phase. As can be seen in
Fig. 7 of the 1 h anodized sample after thermal treatment at
500 °C, the typical peaks are attributed to the anatase phase.
Moreover, the TiO 2 nanotube layer show sharp Raman
peaks, indicating a high degree of crystallinity [53]. This
result is in good agreement with the XRD data, where a
highly crystalline anatase phase was formed after annealing
at 500 °C.
Assessment of corrosion behavior
It is known that the TiO2 nanotubular arrays with suitable
dimensions can enhance the corrosion resistance and biological activity [54, 55]. Hereon, the potentiodynamic polarization tests were performed on the bare substrate, where the asanodized sample and the 500 °C heat-treated coating were
immersed in PBS solution as shown in Fig. 8. The resultant
corrosion potential (Ecorr), corrosion current density (Icorr),
and polarization resistance (Rp) are summarized in Table 1,
where Ecorr represents the substrate’s tendency to corrode and
Icorr represents the corrosion rate and protection efficiency
(P.E.) values. According to the obtained data, the Ecorr of
the bare substrate is − 0.143 ± 0.005 V. This value reached
− 0.135 ± 0.005 V for the as-anodized sample and − 0.089
± 0.005 V for the heat-treated specimen. On the other hand,
the I corr of bare substrate is 4.334 × 10 −6 ± 0.005 ×
10 −6 μA cm −2 , while this value declined to 5.498 ×
10−7 ± 0.005 × 10−6 μA cm−2 for the heat-treated sample.
In accordance with the polarization curves, the anodized
specimen exhibited a stable surface film since the immersed
coating for 30 min had a very low Icorr value. As compared
to the bare substrate, the as-anodized and the 500 °C heattreated specimens showed a 71 and 87% increase in the
P.E., respectively. This suggests that the thermal annealing
196
J Aust Ceram Soc (2019) 55:187–200
Fig. 7 Raman spectrum of TiO2
nanotubes after annealing at
500 °C for 1.5 h
of the TiO2 nanotubular arrays at 500 °C could improve the
in vitro corrosion resistance of the Ti64 alloy in corrosive
environments. This effect can be attributed to the thermally
induced crystallization of the product which resulted in the
formation of a nanotubular arrays with higher stability than
the amorphous structure. In this regard, Saji et al. [56] examined the corrosion behavior of the Ti–35Nb–5Ta–7Zr
alloy with amorphous nanoporous and nanotubular oxide
in Ringer’s solution at 37 °C. They reported that the nanotube layer could form an instantaneous and efficient passivation. In addition, Yu et al. [54] studied the corrosion behavior of titanium with nanotube layers in naturally aerated
Hank’s solution using open circuit potentials (OCP), electrochemical impedance spectroscopy (EIS), and
Fig. 8 Polarization curves of the
substrate, the as-anodized sample,
and the 500 °C heat-treated coating soaked in PBS solution for
30 min
potentiodynamic polarization tests. The electrochemical results showed that the TiO2 nanotube layers on titanium exhibited enhanced corrosion resistance in simulated biofluid
compared to the smooth Ti. More recently, Li et al. [55]
fabricated TiO2–Mg nanotubes with different sizes on Ti–
Mg via the anodization of magnetron-sputtered Ti coatings
on AZ91D magnesium alloy by varying the anodization
voltage and time. They found that the corrosion resistance
of TiO2–Mg NTs was efficiently improved compared to the
AZ91D. Therefore, it can be concluded that, although native TiO2 passivation layer is responsible for good corrosion resistance of the Ti alloys, the improved corrosion
resistance is due to the formation of TiO2 nanotubular arrays on the Ti64 substrate.
J Aust Ceram Soc (2019) 55:187–200
Table 1 Corrosion potential
(Ecorr), corrosion current density
(Icorr), polarization resistance
(Rp), and effectiveness of
corrosion protection (P.E.) values
197
Sample
Ecorr (V)
Icorr (μA cm−2)
Rp (Ohm)
Density (g cm−3)
P.E. (%)
Substrate
Anodized
500 °C
− 0.143 ± 0.005
− 0. 135 ± 0.005
− 0.089 ± 0.005
4.334 × 10−6 ± 0.005 × 10−6
1.272 × 10−6 ± 0.005 × 10−6
5.498 × 10−7 ± 0.005 × 10−6
7.674 × 103
1.799 × 104
4.978 × 104
4.51
1.40
1.40
–
71
87
In vitro bioactivity in SBF
After being implanted into the human body, the implant surface directly contacts and interacts with the tissues and cells.
Hence, the in vitro bioactivity of the metallic implants should
be investigated before utilized in orthopedic applications because this feature plays a vital role in influencing the biological responses [57]. Based on literature survey, no stable apatite layer was observed on amorphous TiO2 films [17]. In the
present case, the rapid development of bone-like apatite layer
is very likely seeing that the as-anodized sample is not
completely amorphous and some characteristic peaks corresponding to the anatase TiO2 are present [11]. However, there
are many challenges associated with the in vitro bioactivity of
the anodized samples after subsequent thermal treatment. In
this regard, Li et al. [58] prepared long TiO2 nanotubes in
fluorinated DMSO solution and annealed at 500 °C. They
immersed the TiO2 nanotubes in SBF solution and observed
a thick apatite layer over the whole surface of the nanotubular
arrays after 14 days, signifying their outstanding in vitro bioactivity. This effect was mainly attributed to the high specific
surface area and the anatase phase. In another report, Ma et al.
[59] examined the precalcification of nanotube array surfaces
to hasten the apatite formation. Nanotube array samples were
annealed at 500 °C for 2 h for crystallization. After rinsing in
DI water, the precalcified and non-precalcified nanotubes
were immersed in a supersaturated calcium phosphate solution (SCP) for Ca–P deposition at 37.5 °C. They reported that
the surface of the precalcified nanotube was completely covered by a flake-like layer after 7 days of immersion in SCP.
However, no such precipitation was observed on the nonprecalcified sample exposed to the same conditions during
SCP immersion. Hereon, due to differences in the electrolyte
composition and product specifications, the in vitro bioactivity behavior of the annealed TiO2 nanotubes was reevaluated.
Figure 9 shows the surface morphologies and EDS results of
the 500 °C annealed sample after immersion in SBF for 1 to
14 days.
After 1 to 7 days of immersion in SBF, a trace of apatite
deposition could be observed onto the TiO2 nanotubes, where
most of the top ends of the nanotubular arrays are not covered
with the bone-like apatite layer (Fig. 9a, b). As the immersion
time was increased to 14 days, the amount of apatite deposition increased sharply and thereby a thick layer of apatite with
Ca/P ratio of 2.88 was formed on the TiO2 nanotube surface
(Fig. 9c–e). Consistent with previous studies [60], the EDS
pattern shows that the deposited layer is a carbonated apatite
containing calcium, phosphorus, oxygen, and carbon elements
as the main constituents. To validate the in vitro bioactivity,
the apatite-inducing ability of the annealed sample was also
tested in different SBF solutions for 14 days according to
the method described by Tas [36] as shown in Fig. 9f–j. It
was found that very thick apatite deposition (Ca/P ratio of
2.59) composed of large agglomerates was formed on the
TiO2 nanotubes. It should be noted that there were no significant differences in the apatite formation on the
nanotubular arrays in the Tas-SBF solution compared to
the Kokubo method.
Conclusion
To conclude, the electrochemical anodization of Ti–6Al–4V
alloy was performed in ammonium fluoride electrolyte dissolved in a 90:10 ethylene glycol and water mixture at room
temperature under a steady potential of 60 V for 1 h. As the
anodization time was extended to 1 h, significant changes
from nanoporous to nanotubular structure were observed.
From the FESEM observations, highly ordered TiO2 nanotubes were formed during annealing at 500 °C for 1.5 h.
Further increase in the annealing temperature to 700 °C resulted in the destruction of the tubular configuration; thereby, a
coarse grain structure was formed. The F 1s region in the XPS
spectrum of the anodized sample confirmed the presence of
high organic/halide content in the tubes (surface-fluorinated
Ti–F species), where the removal was not accompanied by
significant changes in the morphological features. The
Raman spectra showed the formation of the anatase phase of
TiO2 after annealing at 500 °C for 1.5 h. The results of in
vitro corrosion tests indicated that the 500 °C annealed
sample had a significant improvement of P.E. and showed
the highest percentage of protection efficiency. The improved corrosion resistance was attributed to the formation of TiO 2 nanotubular arrays on the Ti64 substrate.
After 1 to 7 days of immersion in SBF, a trace of apatite
deposition was detected on the TiO2 nanotubes, where the
quantity of apatite deposition increased notably after
14 days of immersion and thereby a thick layer of apatite
was formed.
198
J Aust Ceram Soc (2019) 55:187–200
Fig. 9 Surface morphologies and EDS results of the 500 °C annealed sample after exposure to Kokubo-SBF for a 1 day, b 7 days, and c–e 14 days and to
Tas-SBF for f 1 day, g 7 days, and h–j 14 days
Acknowledgements The authors would like to acknowledge the
University of Malaya for providing the necessary facilities and resources
for this research. The authors are also grateful to Research Affairs of
Islamic Azad University, Najafabad Branch for supporting this research.
Funding information This research was fully funded by the University of
Malaya with the high impact research grant numbers of RP032C-15AET
and PG081-2014B.
Highlights
– Corrosion behavior and bioactivity of TiO2 nanotubes on Ti64 were
investigated.
– Themodified surface showed an outstanding capability in enhancing
the bioactivity.
– Corrosion protection efficiency increased remarkably after annealing
at 500 °C.
– Ti 2p3/2 and Ti 2p1/2 components confirmed the existence of Ti4+ state.
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