Strengthening Adhesion of the Hydroxyapatite and Titanium Interface

Article Cite This: ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

www.acsanm.org

Strengthening Adhesion of the Hydroxyapatite and Titanium Interface by Substituting Silver and Zinc: A First Principles Investigation Jin P. Sun and Yan Song* School of Materials Science and Engineering, Harbin Institute of Technology at Weihai, 2 West Wenhua Road, Weihai 264209, China

ACS Appl. Nano Mater. Downloaded from pubs.acs.org by 185.250.43.230 on 09/05/18. For personal use only.

S Supporting Information *

ABSTRACT: A basic understanding of the influences of substituted silver and zinc on the adhesion between hydroxyapatite (HA) and α-Ti surfaces is obtained through first-principles electronic structure calculations. Since the (0001) plane presents the most thermodynamically stable surface of HA, the HA/Ti interfaces with Zn or Ag dopants were constructed between the HA(0001) with dopants and Ti(0001) surfaces. Applying the universal binding energy relation (UBER), the optimal interfacial distance for HA/Ti interface models are estimated and then the optimized interfacial models are determined after full geometrical relaxation, and the affections of Zn dopants on the interfacial distances for the HA/Ti interface are studied. The work of adhesion of interfaces with various stoichiometries was evaluated via the optimized HA/Ti interface models with dopants. It is strongly affected by substituted silver and zinc dopants and reaches the largest value of −3.52 J/m2 for Zn doping and of −3.48 J/m2 for Ag doping in the PO4-terminated interfaces, which is significantly larger than the one of the metal-atom-terminated stacked interfaces (1.31−2.55 J/m2) and of the stoichiometric HA(0001)/Ti(0001) interfaces of without Ca termination (−2.33 J/m2). Therefore, the bond strengths of HA/Ti interfaces are increased by silver and zinc dopants. Analysis of electronic structure reveals that charge transfer between HA and Ti slabs is increased by doping Ag or Zn in the HA. More charge transfers lead to stronger the Ti−O bonds and drive the HA/Ti interface system to be more metallic. KEYWORDS: HA/Ti interfaces, Ag and Zn dopants, work of adhesion, plane-average charge difference, first principles

1. INTRODUCTION Owning to their structural and chemical similarities to the inorganic composition of teeth and bone, including the same bioactivity and osteoconductivity, hydroxyapatite (2(Ca5(PO4)3(OH)), HA) is the important biological active materials coated on titanium and its alloys for medical devices, implants, and tissue engineering.1,2 In such applications, especially in total hip joint replacements and artificial teeth sockets, HA-coated Ti implants suffer from sustained shock and stress, with which coating exfoliation, interfacial cracks, and debris granules occur frequently due to huge differences in physical attributes and the low adhesion strength of the interface between the HA film and Ti substrate. 3−5 Furthermore, infection caused by either bacterial colonization at the wound site or foreign body response to the implant material is a main reason causing the failure of the HA-coated Ti implants in clinical applications.6−10 Clearly, the establishment of a sufficiently strong and durable mechanical link between the bone and implant is highly desirable, which will be the main focus of the current paper. One of the main structural characteristics of biological and synthetic HA is the ability to © XXXX American Chemical Society

incorporate a great variety of isomorphic substitutions, such as cations (Zn2+, Mg2+, Sr2+, Y3+, Ag+)7,9,11−15 and anions (SiO44−, CO32−),16−18 while keeping its hexagonal symmetry. These substitutions can be expected to affect the thermal stability, surface reactivity, crystallinity, and crystal shape of HA as well as biological, antibacterial, and bonding properties of HA coatings. For example, in biological hydroxyapatite, Mg, Sr, and Y are readily substituted for calcium, and the presence of Sr improves its bioactivity and Y enhances its density and crystallinity.7,18−20 With the goal of optimum bone−implant interaction, we presently explore the capabilities of bonemimetic HA coatings to mediate the mechanical coupling between the bone and implant. In doing so, the bonding strength between HA with dopants and Ti will be considered in depth. Many studies were devoted to the HA coating structures with or without dopants on metal substrate, including coating Received: June 29, 2018 Accepted: August 7, 2018 Published: August 7, 2018 A

DOI: 10.1021/acsanm.8b01103 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

Article

ACS Applied Nano Materials

reports were focused on the interfacial properties including the interfacial bonding strength, interface structure, and the phase composition just at the interface between HA coatings and the Ti substrate. Ma et al.37 fabricated zinc-substituted hydroxyapatite/silk fibroin composite coatings on titanium substrates at room temperature by electrophoretic deposition, and they reported that the obtained coatings maintained the phase of HA. In another article from the same group, they suggested that the zinc ions influence the crystal formation of hydroxyapatite and the synergy effect of zinc and alginate ions led to the preferred growth along ⟨002⟩ and ⟨211⟩ orientations.38 Nanocrystalline zinc-substituted HA films were deposited on the titanium-coated silicone by the electrophoretic deposition method and the TEM images of nanocrystals with different Zn ion concentrations show smaller crystalline sizes less than 100 nm.9 Zn dopants tend to inhibit the growth of the HA crystallites and thus form low crystallinity of the HA grains on the surface of substrate, which may have a negative effect on interfacial interaction between Zn-doped HA and the substrate.39 However, silver is a well-known bactericide that can inhibit a wide range of bacteria, even antibiotic-resistant bacteria.28,36,40,41 In spite of the advantages, cytotoxic effects on human cells at high concentrations have been shown in some investigations.42−44 In this view, there have been considerable efforts on the synthesis of Ag-substituted HA with small amounts of Ag.38,45,46 Roy et al.47 suggested that HA with 2 and 4 wt % Ag dopants has excellent antimicrobial properties while the adhesive bond strength between the substrate and coating is more than 15 MPa ensuring the mechanical integrity of the coatings. Some literature reported that the chemical bonds between the Ag-substituted HA coating and Ti substrate within the ion diffusion zone improved adhesion strength, while only cohesive failure was detected,48 but others noted that the bond strength on the Ti substrate is slightly lower with the Ag being incorporated into the HA films, which is possibly caused by the looser Ag-HA coating.49 Most recently, Li et al.50 prepared the Ag/Sr-substituted HA coatings on pure titanium via a one-step hydrothermal deposition method. The HA coatings with dopants characterized by HRTEM and selected area electron diffraction (SAED) were composed of nanoparticles whose main component was HA with good crystallinity. Since the biological response of doped HA and interfacial adhesion strength between HA coatings and Ti substrates are significantly changed by these substitutions, fundamental understanding of the defect chemistry and mechanical properties of substituted HA and the influence of substitutions on the chemistry and interface affinity from atomic and electronic levels is essential not only for understanding how nature works but also for designing synthetic biomaterials. Recent advances in electron microscopy have enabled atomic-resolution imaging of interface structures. However, understanding the observed images is not always straightforward because of abrupt structure changes and electronic discontinuity at interfaces.51 Ab initio calculations can provide complementary understanding of the interfacial structure and the electronic properties of the interface, especially when these properties are hardly accessed by experiment for the buried interface layer. One such example can be found in the work of Surmenev and Neyts et al.,34 where a comprehensive study of structure, energetic stability, and interface bonding nature of HA/rutile TiO2 was conducted. Their work also highlighted

methods, morphology, chemical compositions, crystallinity, and crystal orientations. There are a variety of methods to coat HA on Ti metal substrates, including sol−gel,21 solvothermal method,22,23 electrochemical deposition method,24−26 hydrothermal deposition method,14,26 electrophoretic deposition,27,28 and Rf sputter coating.28 The chemical composition and morphology of the coating varies greatly depending on the synthesis conditions and it is particularly true for the phase composition just at the interface. Relatively early, Ban et al.29 deposited hydroxyapatite crystals on a titanium electrode using the hydrothermal−electrochemical method and the crystallinity of the deposited hydroxyapatite increased continuously with the electrolyte temperature (at 200 °C the crystallinity of the deposited HA reached to about 90%). Carrado30 reported that HA films on titanium matrix prepared by plasma-spray method was hexagonal closed-packed (hcp) phase with slight amorphous background determined by high-energy synchrotron X-ray diffraction in energy dispersive. High-crystallinity HA crystal layers with a single phase were fabricated on titanium by the low-temperature flux coating method.31 Using the right angle radio frequency magnetron sputtering technique (RAMS), López et al.32 produced high-quality HA thin coatings on the Ti substrate. At the first stage of deposition, the CaP layer just at the interface between the coating and substrate was composed of crystalline nanodomains (∼2 nm) coexisting with highly disordered regions. The physicochemical characteristics of the HA coating revealed a higher sensitivity to the RAMS conditions and transformed these primary species into amorphous calcium phosphate building units. The growth of the induced particle and the subsequent transformation into HA crystalline grain domains could be possible by controlling and adjusting the RAMS conditions. Solla et al.33 reported the polycrystalline nature of HA coatings produced by pulsed laser deposition on the Ti substrate using HRTEM and focused ion beam. They revealed that the HA crystal shows a clear ⟨002⟩ oriented growth in columns from bottom to top and, in some cases, the preferred orientation appears to start at a certain distance from the substrate. Grubova et al.34 calculated the formation energies and work of adhesion of an amorphous HA surface onto amorphous as well as stoichiometric and nonstoichiometric crystalline TiO2. Carbonated25 and yttrium-substituted24 HA coatings with aligned crystal domains were produced on the seeded Ti substrates by electrochemical seeding followed by the hydrothermal crystal growth, and the crystal domains oriented with the crystallographic c-axis perpendicular to the substrate. Recently, Mg-substituted HA samples15 were hydrothermally synthesized and the characteristics of these nanoparticle by XRD, field emission scanning (FE-SEM), and high-resolution transmission electron microscopies (HRTEM) were that Mg2+ incorporation into HA nanorods gives rise to a tailored crystallinity degree, morphology, surface hydration, solubility, and degradation properties. Among all the antimicrobial metal ions, Zn is an essential trace element involved in osteoblast activity, collagen synthesis, and alkaline phosphatase activity, which plays an important role in the regulation of bone deposition and resorption.35 The incorporation of zinc ions into hydroxyapatite could add more biological function to the nanoparticles, such as antibacterial ability28,36 and osteoblast response.7 There are extensive investigations on the fabrication, morphology, and bioactivity properties of Zn-substituted HA coatings. However, only a few B


Article


constituent and parameters are the same as in ref 57. Table S1 presents the atomic charges together with nonbonded force field parameters,58 and Table S2 gives the intra-atomic force field parameters,59,60 which were all extensively used in the systems containing HA.61 Biological and synthetic HA both occur in the hexagonal structure, having the P63/m space group with two formula units of Ca5(PO4)3OH.62 For amorphization of HA (aHA) MD simulations, it is more expedient to transform it into an orthorhombic cell (x, y, and z axes) with the z-axis parallel to the c-axis and double b-axis direction in order to maintain the periodicity of the lattice. The supercell of 3 × 2 × 6 was constructed with a vacuum layer of 30 Å to simulate the amorphization of the HA. The middle layer (3 × 2 × 2) of the supercell was kept unchanged to simulate the true internal HA crystal and other layers, those are upper and lower layers, were relaxed in MD simulations. A canonical (NVT) ensemble, a Nose−Hoover thermostat, and a short time step of 0.1 fs were adopted in our MD simulations using the LAMMPS program.63 Figure S6 illustrated the quenching process. First, we relaxed the system at 300 K for 50 ps and then heated it with a stepped heating mode to the target temperatures. In this mode, the system was heated at 100 K/20 ps speed, and then balanced 20 ps for every 100 K before reaching the target temperature. Since the melting point of the crystalline HA is 1923.16 K,64 four target temperatures, 1600, 2000, 2400, and 2800 K, were chosen. Balanced at the target temperature equilibrium for 50 ps, the system was then cooled to 300 K at a speed of 10 K/ps. Finally, the quenched structure of the system was obtained at an equilibration of 300 K for 50 ps. On the bais of the quenched structures, four interface models were designed to compare the interfacial interactions between the amorphization HA(0001) [aHA(0001)] and Ti(0001) surfaces by using density functional theory calculations. Both atomic and electronic structures in this manuscript were visualized with VESTA 3 software program.65 2.1. aHA(0001) Faces. In order to verify the BMH parameters in the simulated quenching method, the total energy and melting point of the bulk HA were first simulated under the above-mentioned steps for surface heating and cooling simulation. The value of total energy of HA bulk is −709.51 eV at 300 K, which is in agreement with that of the ab initio calculation (−625.78 eV)66 with the difference less than 12%. At 2200 K as shown in Figure S7, the system reaches its melting point, which is higher than that of measurements.64 For the generation of the amorphous HA structure from the crystalline HA(0001) surface, the temperatures of 1600, 2000, 2400, and 2800 K were chosen to investigate the affections of temperature on surface structures. Figure S8 presents the quenching surface structures. The HA(0001) surface at the room temperature (Figure S8a) was also given for comparison. In the melt and quench technique, the middle layer (3 × 2 × 2) of the supercell was fixed, while the two topmost and two bottom layers (3 × 2 × 4) of the supercell which are adjacent vacuum layer were relaxed. As shown in Figure S8b, quenching at temperature 1600 K (lower than the HA melting point), the atoms are disordered remarkably within the range of about 7Å thickness near the vacuum layer. The basic characteristic of this disorder is that the rotation of PO4 group and dramatic changes of the Ca ions coordinates at the surface. At 2000 K (Figure S8c), the obtained structure is similar to that of quenching at 1600 K. The main differences are that the thickness of disordered zone along the z axis is increased to 10Å and the atomic displacements increase. When the quenching temperature reaches at 2400 K higher than the melting temperature of the crystalline HA (Figure S8d), the distortions of the PO4 group increase remarkably and part of them near the surface move outward to the vacuum layer. The reorientation of hydroxyl groups aggravates and some of Ca atoms move toward the surface. At 2800 K, the reorientation and displacement of OH group escalate and this is especially true for atoms near the surface, and some of them move into the vacuum leaving vacancies, which is reported in some experimental studies.67 2.2. Interface Models. The Zn (or Ag) doped-HA(0001)/ Ti(0001) interface was simulated considering terminations of the constituted surfaces, dopants occupying sites, and boundary

the interfacial bonding mechanism of amorphous HA surface onto amorphous TiO2, since the interfacial HA coated by RF sputter method has an nanocrystal/amorphous structure.52 Motivated by similar objectives, in this research, we investigate the role of the Ag and Zn impurities in HA side on interfacial adhesion between HA and Ti slabs. We calculate [0001]oriented HA/Ti interfaces with dopants, focusing on the electronic mechanism of the interfacial strength, while Ag or Zn doping is modeled within the supercell approach. Four different stacks and two interface models for the HA/Ti interface with Zn dopants and three different stacks for Ag dopants were constructed to investigate the influences of stoichiometries of interfaces on the work of adhesion and the atomic arrangements in the interfacial zone. We also employ DFT calculations to investigate the interaction between the amorphous HA(0001) surface [aHA(0001)] and the Ti(0001) slab because the experimental data have shown that HA at the interface has the crystalline structures coexisting with amorphous structures. We are focusing on the electronic mechanism of the interfacial affinity of the HA/Ti interface modified by impurities of Zn or Ag and the aHA/Ti interface, determining the stable interface structures, and shedding light on the origin of the experimentally observed enhanced bioactivity, density and decreased crystallinity by Zn or Ag dopants as well as the bonding nature . The rest of the paper is organized as follows. We first describe the structural models and give the details of our calculations. In a next step, amorphization of HA(0001) surfaces was performed using a molecular dynamics simulated quenching approach (MD) and the interfacial separation is determined by the universal binding energy relation curves (UBER)53 using semicycle boundary interface model for HA/ Ti interface with Zn dopants. This optimized interfacial separation was used as an input for the ab initio modeling. On the basis of the quenched HA(0001) structures, the interfaces between the amorphous HA(0001)[aHA(0001)] and Ti(0001) surfaces were constructed and calculated using density functional theory calculations. Next, the optimal structures of the interfaces are given as well as their work of adhesion and interfacial energy. Finally, the electronic structures of Ag and Zn doped HA/Ti interfaces were investigated to clarify the intrinsic mechanisms of the impact of dopants on the adhesion of the interface.

2. METHODS OF CALCULATION The Vienna ab initio simulation package (VASP) was employed to perform all calculations.54,55 The details of the calculations are referred to our previous work. 56 The 3d 10 4p 2 and 4d 10 5s 1 configurations were regarded as the valence electrons of Zn and Ag, respectively. The convergence of the total energy of HA unit cell on the cutoff energy and surface energy of the HA(0001) and Ti(0001) surfaces on the cutoff and vacuum space were examined. Surface 1 energy is defined asEsurf = 2A [Eslab − NE bulk ], where Eslab and Ebulk refer to the total energies of the slab and the unit cell of bulk crystal, respectively, and N is the number of the unit cell contained in the slab model. From the convergence curves listed Figures S1−S5 in Supporting Information I, 15 Å vacuum layer and 450 eV planewave cutoff energy were used for studying the Ti(0001) and HA(0001) surfaces and interface properties using ab initio calculations. Amorphization of HA is performed using a molecular dynamics simulated quenching approach (MD). Hauptmann’s57 potential energy functions developed specifically for HA and BMH force field were used in our simulation. The details about the BMH force field C


Article


Figure 1. Top and side views of the relaxed HA/Ti with Zn dopant interfaces models: (a) 1Zn-HA(1)-Zn/Ti, (b) 1Zn1Ca-HA(1)-Zn/Ti, (c) 3PO4-HA(2)-Zn/Ti, and(d) 1Ca-HA(2)-Zn/Ti.

Figure 2. Top and side views of single-interface models of the HA/Ti interfaces with Ag dopants: (a) 3PO4-HA(2)-Ag/Ti, (b) 1Ca-HA(2)-Ag/Ti, and (c)1Ag-HA(1,2)-Ag/Ti after three Ti layers and one HA unit relaxation. periodicity. Detailed descriptions on the crystal structure of HA without dopants were given in our previous study66 and the model of HA crystal was shown in Figure S9a. On the basis of the relaxed structures of HA crystals with Zn or Ag dopant, the doped HA (0001) surface models are cleaved as shown in Figure S9b,c. On the basis of the optimized structures of different atom (group)-terminated HA(0001) surfaces with Zn or Ag dopants (as shown Figures S10 and S11 in Supporting Information III), the interface models with dopants are created showing in Figures 1 and 2 and the details are given in follows. Taking Zn-substituted interfaces as an example, considering the sites of substitutions, the cross sections of both Zn substitutes for Ca(1) and Ca(2), denoted as Zn(1) and Zn(2), respectively,

hereafter, are given in Figure S9b,c. The solid color lines in Figure S9 show how to cut the models. Taking Figure S9b, for instance, removal of particles above the green line produces a zinc atom termination surface (named 1Zn), removing particles from the 1(Zn +Ca) black line produces a surface model where two metal atoms, Zn and Ca, terminated, and so forth. As a matter of convenience, hereafter, they are called 1Zn, 1(Zn+Ca), OH, (PO4+OH), and 3PO4 for Zn substitutes for the Ca(1) atom, and named as 1Ca, 2Ca, (3PO4+OH), (2Ca+Zn), and 3PO4 for Zn located at the Ca(2) site. Every one corresponds to the crystal under one of the solid color lines in Figure S9. As shown in Figure S9b,c, five possible stacks exist between HA(0001) and Ti(0001) surfaces for each Zn substitution at the D


Article


Figure 3. Top and side views of aHA/Ti interface models: (a) 2Ca-aHA/Ti-OH, (b) PO4-aHA/Ti, (c) 2Ca-aHA/Ti-OH2, and (d) 2Ca-OHaHA/Ti after two Ti layers and one or half HA unit relaxation. interface between HA(0001) and Ti(0001) surfaces,56 and four interface models are constructed by integrating aHA on the base of Ti stacking along the [0001] direction with a 1.6 nm vacuum layer at the top of the interface (along the z axis). As shown in Figure S8, one or more topmost and bottom atomic layers close to the surface of the (3 × 2 × 6) HA(0001) systems after the “melt and quench” possess a noncrystal structure with random atom distributions. To identify the possible atomic structures of aHA(0001)/Ti(0001) interfaces, the different zones in the upper layer of the (3 × 2 × 6) HA(0001) systems were scanned. In this step, the structure quenched at 2800 K serves as a good example since it has been reported in some experiments.67 The structure of the interface between the aHA surfaces and Ti(0001) surface at all scanning zones was analyzed, and four different stacking configurations were selected for a subsequent DFT optimization. All interface models were subsequently relaxed in the DFT optimization. First, we simulated the aHA(0001)/Ti(0001) supercell with the size (1 × 1 × 2) unit orthorhombic cell of HA in directions a and b (a = 0.955 nm and b = 1.654 nm) and about 4.0 nm in the c direction with a 1.6 nm vacuum layer and 272 atoms. However, the VASP simulation is large and quite time-consuming. Thus, a (1 × 1 × 1) slab model of orthorhombic HA in all the a, b and c directions and about a 1.6 nm vacuum layer in the c direction and 184 atoms were used as shown in Figure 3. Four interface models, (a) 2Ca-aHA/Ti-OH, (b) PO4-aHA/Ti, (c) 2Ca-aHA/Ti-OH2, and (d) 2Ca-OH-a HA/Ti are calculated. Among them, the 2Ca-aHA/Ti-OH corresponds to the -O-O-Ca-Ca/Ti-Ti-··· atomic stacking at the interface, OH presents that there are hydroxyls moving into the vacuum layer, and the aHA shows the amorphous nature. PO4-aHA/ Ti is the same as the 3PO4-HA(2)-Zn/Ti interface except that there are no dopants in HA, and the HA is anoncrystal. The differences between 2Ca-aHA/Ti-OH2 and 2Ca-aHA/Ti-OH are the size of interface model and the orientation of OH, which is in the vacuum layer. 2Ca-OH-aHA/Ti presents the atomic stacking of -O-O-Ca-CaOH/Ti-Ti-; that is to say, there are one OH group and two Ca atoms at the interface.

Ca(1) site and the Ca(2) site. However, due to the variance of atomic densities at different surfaces, we reduce the considered structures to two sequences for both Zn atoms at the Ca(1) and Ca(2) sites, namely, 1Zn-HA(1)-Zn/Ti, 1(Zn+Ca)-HA(1)-Zn/Ti, 1Ca-HA(2)Zn/Ti, and 3PO4-HA(2)-Zn/Ti, denoted as nX-HA(m)-Y/ Ti (m, n = 1 or 2), where X is the terminated metal atom or PO4 group, Y is the dopant, n is the number of interfacial metal atoms (or groups), and m = 1 and 2 represents the site of dopants occupying on Ca(1) and Ca(2) site, respectively. One of them, the 1(Zn+Ca)-HA(1)-Zn/ Ti interface corresponds to the interfacial atomic stacks of -O-O-CaZn/Ti-Ti-··· standing for the oxygen-deficiency stacking. For the 1CaHA(2)-Zn/Ti interface model, one Ca atom terminates in the HA(0001) surface and a Zn dopant lies in the subinterface at the Ca(2) site. As for 1Zn-HA(1)-Zn/Ti, the dopant Zn replaces the Ca(1) atom and terminates in the HA(0001) surface. For the 1Zn-HA(1)-Zn/Ti case (Figure S12 of Supporting Information III), there are three nonequivalent sites on the Ti(0001) surface for the zinc atom at the interface in the HA side to face: (a) the top site of interface Ti atom (labeled as T), (b) the hexagonal center site surrounded by the Ti atoms in the topmost layer of Ti(0001) surface (labeled as C), and (c) the bridge site between Ti atoms in the topmost layer (labeled as B). The 3PO4-HA(2)-Zn/Ti interface is composed of the Ti(0001) surface and the PO4−PO4− PO4-terminated HA(0001) surface with Zn dopant in the subinterface at the Ca(2) site represented oxygen-rich stacking with tight coordinate. For Ag doped-HA/Ti interfaces (Figure 2), three translational states, 1Ag-HA(1,2)-Ag/Ti, 1Ca-HA(2)-Ag/Ti, and 3PO4-HA(2)Ag/Ti are calculated. Among them, the 1Ag-HA(1,2)-Ag/Ti corresponds to the interfacial atom stacking as O-O-Ag/Ti-Ti-··· 1Ag and Ag(1,2) present that there are two Ag atoms occupying two different Ca sites in the interface of the HA side and one of them is at the nearest neighbor interface. 1Ca-HA(2)-Ag/Ti is the same as 1CaHA(2)-Zn/Ti; that is, one Ca atom terminates the HA(0001) surface and a Ag impurity lies in the subinterface at the Ca(2) site. 3PO4HA(2)-Ag/Ti is the same as 3PO4-HA(2)-Zn/Ti except in different substitution atom. For the aHA(0001)/Ti(0001) interface, the original interface distance is set to 2.2 Å, according to our previous studies about the

3. RESULTS AND DISCUSSION 3.1. Optimization of Interface Structures. The surfaces with different terminations were obtained by cleaving the HA E


Article


Thus, a conclusion may be drawn that the fracture strength of the HA/Ti interface is enhanced by the Zn dopant at Ca(1) sites of the HA crystal. Second, the original interface models with the “optimal” interfacial separations have been relaxed, allowing three Ti atomic layers and one HA unit of the slabs relaxation to study the relaxed properties. The relaxed separations for 1Zn-HA(1)-Zn/Ti interface are listed in column of “full relaxed” of Table 1. For the d0 values of 1Zn-HA(1)-Zn/Ti shown in Table 1, the fully relaxed interfacial separations for bridge- and center-connected interfaces are around 0.18 nm, smaller than those obtained by the UBER method due to the rotation of the PO4 group in the HA slab and because the Ti atoms at the interface approach the HA slab upon the relaxation, while the UBER method just calculates the static energy of the system without account of the lattice mismatch between the two surfaces that make up the interface. For the top-connected interface, the interlayer distance is about 0.26 nm, much larger than those of bridge- and center-connected interfaces due to the higher repulsion between Ti and Zn atoms at the interfaces. Compared with the 1Ca-HA/Ti interface without dopant,56 the 1Zn-HA(1)-Zn/Ti interface with Zn dopant at the Ca(1) site has slightly shortened interlayer distances. However, because doped atoms at the subinterface and/or at the interface will be fully relaxed, it is valid to take the separation of 0.20 nm between the HA slab and Ti slab and the metal atoms at the interface of the HA side locating at the hexagonal center, which is formed by six Ti atoms from the two topmost Ti layers, as the original interslab distance and stacking configurations for other interface models with Ag or Zn dopants, respectively. As for the 1Ag-HA(1,2)-Ag/Ti case, the interfacial Ag is also occupied the hexagonal center of the interfacial Ti atoms with the original interface distance of 0.20 nm, since only the terminated metal atom type was different between 1Ag-HA(1,2)-Ag/Ti and 1Zn-HA(1)-Zn/Ti(1) interfaces. 3.1.2. Optimized Interface Structures. The interface structures of HA/Ti interfaces with Zn and Ag dopants were optimized, and the final structures are shown in Figures 1 and 2, respectively. Different numbers of chemical bonds are crossing the interface, which are formed between Ti atoms from Ti slab and O atoms from HA slab, and the characters of the bonds are significant affected by the dopants. There are several fascinating features in the relaxed interface structures with dopants. First, for nX-HA(1)-Zn/Ti cases, five Ti−O chemical bonds less than 0.24 nm in length crossing the interface were formed regardless the termination of either a single Zn metal atom or two metal atoms (one Zn and one Ca). The metal-against-metal stacking 1(Zn+Ca)-HA(1)-Zn/ Ti shows slightly different repulsive force between metal atoms from HA slab and Ti atoms of the Ti slabs compared to the single Zn-terminated 1Zn-HA(1)-Zn/Ti case and quite stronger bonding strength than that of the undoped HA/Ti interface with two Ca atom termination. Second, on the affection of dopants, the HA and Ti structures possess better coordination and therefore more strength bonding than undoped HA/Ti interface, especially for the 3PO4-HA(2)-Y/ Ti interfaces (Y = Zn or Ag). Third, Zn-terminated interfaces possess more strength bonding than a Ca-terminated interface as more Ti−O bonds crossing the interface were formed in the 1Zn-HA(1)-Zn/Ti than in the 1Ca-HA(2)-Zn/Ti interface models. Fourth, five Ti−O bonds with the length ranging from 0.2 to 0.24 nm were generated in Ca and Ag (at the Ca(2)

crystal with dopants along the colored lines in Figure S9. Optimized structures and surface energies were calculated and shown in Supporting Information III. Then, the appropriate separation between the two subslabs, the HA(0001) and Ti(0001) slabs, for interfacial supercells was obtained using the UBER method53 for 1Zn-HA(1)-Zn/Ti. On the basis of the results of UBER, the original interface models with the “optimal” interfacial separations have been fully relaxed, keeping atoms in the bottom three layers of Ti side and the topmost one HA unit layers fixed during the relaxation. 3.1.1. Interfacial Separation. Under the UBER method, fitted curves of the adhesion energy Ead versus the interfacial separation d for the three possible stacking structures of 1ZnHA(1)-Zn/Ti (in Figure S12) are obtained as shown in Figure S13.53 The detailed description about this method can be found in ref 56. As shown in Figure S13, the ideal adhesive energy of the bridge connection is slightly larger than that of the center connection, and the top connection owns the lowest ideal adhesive energy, which results from a larger interface spacing compared with that of the B and the C arrangements as shown in Figure S12. These results are as expected from the fact that the Zn and Ti atoms are all electropositive and repel each other. In order to make quantitative comparisons between the results and those of the relevant constituents, the ideal adhesive energies (E0), the equilibrium separation d0, the ideal peak interfacial stress (σmax), and the scaling constant s of all the interfacial configurations are listed in column of “UBER” of Table 1. Table 1. Interface Bonding Parameters for Three Translational Configurations of the 1Zn-HA(1)-Zn/Ti with Static and Fully Relaxations static (UBER) d0 (nm) E0 (J/m2) Wad (J/m2) σmax (GPa) γint (J/m2) s (nm) number of Ti−O

full relaxed

B

C

T

∼0.21 ∼0.88

∼0.20 ∼0.76

∼0.23 ∼0.47

∼11.69

∼7.81

6.74

0.06

0.07

0.05

B

C

T

0.18

0.19

0.26

−2.22

−2.20

−2.03

1.19

1.21

1.38

4

6

5

In the UBER column of Table 1 the d0 of the considered configurations is slightly larger than 0.2 nm and the optimal interface distance of the top stacking is larger than those of the bridge and center stacking arrangements. The translational configurations of the HA/Ti interface with Zn dopant weakly influence the ideal adhesion energy (E0) and its values for center- and top-connected interfaces are 0.76 and 0.47 J/m2, respectively. And the interfacial Zn atom prefers to locate at the bridge site of interfacial Ti atoms with an ideal adhesive energy of 0.88 J/m2, which is 0.11 and 0.41 J/m2 larger than it locates at the hexagonal center and the top of the interfacial Ti atom, respectively. Consequently, there is stronger bonding of the interface between the Zn doped HA and Ti in the bridgeconnected interfaces. However, we notice the bridgeconnected interface with an equilibrium separation d0 of 0.21 nm, which is 0.1 nm wider than the center-connected interface. The values of σmax are 11.69, 7.81 and 6.74 GPa for bridge-, center-, and top-connected interfaces, respectively, which is much larger than that of the HA/Ti interface without dopants. F


Article

a

Ti−O distance (×10−1 nm) number of Ti−O bonds

PO4 in HA slab

Δ Δd12 Δd23 Δd34 Δd12 Δd23

1Ag-HA(1,2)-Ag/Ti

[Variations of the interlayer distance Δdij as the percentage of the distance in the crystal Δdij = (dij − d0ij)/d0ij × 100%. 3PO4-, 1Ca-, 1Zn-, 1(Zn+Ca)- and 1Ag-mean 3PO4-terminated, 1-Ca-terminated, 1Zn-terminated, 1Zn1Ca-terminated, and 1Ag-terminated of the HA side in the interfacial supercell. The number in parentheses represents the corresponding quantitative value of the same metal-atomor group-terminated HA/Ti interfaces without dopants in the vertical columns of 3PO4-HA(2)-Zn/Ti and 1Ca-HA(2)-Zn/Ti, and that in the 1Zn1Ca-HA(1)-Zn/Ti vertical column is the corresponding value of the two Ca-atom-terminated HA/Ti interface without dopants.

2.05 2.41 0.36 −11.8% −6.0% −7.1% 11.6% 1.8% 2.01−2.24 4

1Ca-HA(2)-Ag/Ti 3PO4-HA(2)-Ag/Ti

2.05 2.23 0.18 −17.6% −4.3% −7.0% 15.5% 3.1% 2.05−2.12 7 2.10 (2.40) 2.20 (3.69) 0.10 (1.29) −14.4% (−11.8%) −6.7% (−5.1%) −6.5% (−6.1%) 5.3% (0.01%) 2.4% (9.1%) 2.09−2.23 (2.3) 4(1)

1(Zn+Ca)-HA(1)-Zn/Ti 1Zn-HA(1)-Zn/Ti

2.05 1.81 0.24 −12.5% −0.3% −6.4% 5.3% −2.9% 1.93−2.4 4 2.20 (2.31) 2.39 (2.38) 0.19 (0.07) −13.3% (−10.7%) −4.3% (−2.9%) −7.0% (−6.0%) 11.2% (19.8%) −1.4% (8.7%) 2.05−2.08 (2.17−2.35) 3(4)

1Ca-HA(2)-Zn/Ti 3PO4-HA(2)-Zn/Ti interlayer system

2.02 (2.4) 1.64 (2.48) 0.28 (0.08) −10.1% (−9.7%) −6.8% (−6.2%) −8.3% (−7.6%) 10.6% (10.1%) 2.6% (6.1%) 2.10−2.33 (2.0−2.07) 10(4)

HA/Ti−Ag HA/Ti−Zn

Table 2. Variation of Interfacial Distance of HA/Ti Interfaces with Dopants after Relaxationsa G

original d1 (10−1 nm) relaxed d0 (10−1 nm) 10−1 nm Ti slab

site) coterminated interface, the 1Ca-HA(2)-Ag/Ti interface, while there are only two Ti−O bonds in the single Ag-atomterminated interface (the 1Ag-HA(2)-Ag/Ti interface). Fifth, dramatic changes of coordinates of Ti atoms occurred upon the relaxation, whereas the positions of atoms in the HA slab side are rare to be distorted no matter what terminations of HA(0001) surfaces are. The structural details including the changes of interlayer spacing Δdij of the considered HA/Ti interfaces with dopants are presented in Table 2. The definition of atomic layers and the calculation method of the interlayer separation changes are described in ref 56. Since 3PO4-HA(2)-Zn/Ti and 1CaHA(2)-Zn/Ti are the most promising interface structures among the Zn doped HA/Ti interface, while 3PO4-HA(2)-Ag/ Ti and 1Ag-HA(1,2)-Ag/Ti are the most likely interface structures in Ag doped interfaces, their geometric details are analyzed here (Figures 1 and 2). In 3PO4-HA(2)-Zn/Ti, there are ten Ti−O bonds, the largest number of chemical bonds of the considered systems, the interface with bond length shorter than 0.233 nm, showing a strong chemical bonding of the Zn doped HA/Ti interface. The strong Ti−O bonds formed at the interface have significant impacts on the geometrical configurations of Ti and HA layers. For instance, the interlayer spacing between the closest interfacial PO4 layer and the second closest interfacial PO4 layer in the c direction is significantly increased by 10.6% comparing to that of bulk HA and slightly larger than that of the 3PO4-HA/Ti interface without dopants, as shown in the brackets in Table 2. The deviation of the interlamellar distances of the center PO4 layer is 2.6%, which is much smaller than that of the interfacial PO4 layer. The c dimension of the HA lamella is enlarged by 3.3%. In contrast, the interlayer distances of the Ti(0001) slab are contracted and the deviations, Δd12, Δd23, and Δd34, are −10.1%, −6.8%, and −8.3%, respectively. For the 1Ca-HA(2)Zn/Ti interface, three Ti−O bonds with lengths around 0.20 nm, closing to the bond length of Ti−O bonds in rutile titanium dioxide crystal (0.198 nm, I4̅1/amd),68 are crossing the interface, showing the quite strong chemical bonding at the interface. By comparing the 1Zn-HA(1)-Zn/Ti interface with the 1(Zn+Ca)-HA(1)-Zn/Ti interface, we find that the shortening Ti−O bond delivers the larger interfacial adhesive energy. For the 1Ca-HA(2)-Zn/Ti interface, the interlamellar separation is increased by 11.2% for the interfacial PO4 layer while the interlamellar spacing is compressed by 1.4% for subinterfacial PO4 layer. The separation deviations of Ti slab from interface to bulk are −13.3%, −4.3%, and −7.0%, respectively. For 3PO4-HA(2)-Ag/Ti interface, there are seven pairs of Ti−O atoms with the distance less than 0.22 nm. It has the strongest interface combination as listed in Table 2 for HA/Ti interface with Ag dopant, which is consistent with the largest adhesive energy showing in Table 3. Upon relaxation, deviations for the respective (0001) interlamellar separations of PO4 layer from interface to center are 15.5% and 3.1%, respectively. The dimension of the HA lamella in the [0001] direction is increased about 6.2%. For the Ti(0001) side, the interlayer separations are significantly decreased and the deviations from the interface to center are −17.6%, −4.3%, and −7.0, respectively. As for the 1Ag-HA(1,2)-Ag/Ti case, two Ti−O bonds with shorter lengths than that of other interfaces (Table 2) were generated across the interface. The shorter length of the Ti−O bond corresponds to the larger interface adhesion energy comparing to the case of the 1Ca-

2.09 2.27 0.18 −12.0% −6.3% −6.8% 13.4% 4.1% 2.02−2.03 2



Article


Table 3. Work of Adhesion and Interfacial Energy for Four Translational States of Zn Doped HA/Ti Interfaces, Three Translational States of Ag Doped HA/Ti Interfaces, and Four Interface Systems of aHA Stacking on Crystal Ti Slab system

Etotal (eV)

EZn‑HA(0001) (eV)

ETi(0001) (eV)

Wad (J/m2)

γint (J/m2)

1Zn-HA(1)-Zn/Ti 3PO4-HA(2)-Zn/Ti 1Ca-HA(2)-Zn/Ti 1(Ca+Zn)-HA(1)-Zn/Ti 1Ag-HA(1,2)-Ag/Ti 1Ca-HA(2)-Ag/Ti 3PO4-HA(2)-Ag/Ti 2Ca-aHA/Ti-OH2 2Ca-OH-aHA/Ti 2Ca-aHA/Ti-OH PO4-aHA/Ti

−1005.530 −1004.334 −1004.532 −1004.320 −1006.328 −1007.002 −1010.033 −1291.258 −1275.033 −1867.345 −1890.265

−604.466 −599.184 −608.808 −605.073 −608.538 −610.487 −603.797 −570.300 −572.219 −1165.580 −1165.689

−390.257 −388.002 −383.313 −390.906 −389.465 −389.389 −388.936 −699.128 −684.269 −681.592 −681.883

−2.22 −3.52 −2.55 −1.71 −1.67 −1.43 −3.48 −1.97 −1.88 −2.046 −4.33

1.19 0.21 0.91 2.03 1.61 1.99 0.21

bonding than the 1Ca-HA(2)-Ag/Ti. Comparing the work of adhesion of the 3PO4-HA(2)-Ag/Ti interface with that of the PO4-HA/Ti interface without dopants,56 we can draw the conclusion that Ag dopants occupying the Ca(2) site with the Ag/(Ag + Ca) = 5% enhance the bonding strength of HA/Ti interface. From Table 3, we can concluded that the Zn or Ag dopant enhances the bonding strength of the crystal HA/Ti interface, especially when it occupies the Ca(2) site. Previously, similar results are reported that the chemical bonds between the Agsubstituted HA coating and Ti substrate within the ion diffusion zone improved adhesion strength, while only cohesive failure was detected.48 But others noted that the bond strength on the Ti substrate is slightly lower with the Ag being incorporated into the HA films49 and Zn dopant tends to inhibit the growth of the HA crystallites, which may have a negative effect on the interfacial interaction between Zn-doped HA and the substrate.39 It is obviously that there is a little bit inconsistency between the results of our ab inito calculation with some experimental data. The difference may arise from the looser Ag-HA or Zn-HA coatings prepared experimentally and from the lower crystallinity of the HA grains on the surface of substrate. Table 3 also presents the calculated work of adhesion for four interface systems of aHA slab stacking on the crystal Ti slab. The more negative the value obtained Wad is, the greater the bonding strength at the interface is. It appears from Table 3 that stacking configurations have a strong influence on the interfacial adhesion values and the PO4-aHA/Ti interface has the most negative value of Wad −4.33 J/m2. Comparison of results obtained for crystalline HA(0001) without dopants56 and aHA(0001) reveals that a strong interaction is obtained only for the latter. Previously, rather similar values of the work of adhesion, in the ranges of −0.018 and −2.429 J /m2, were obtained for the aHA/aTiO2 interface system using MD and DFT by Grubova et al.34 3.3. Interface Energy. The calculated interface energies γint are calculated via the methodology in ref 56 and listed in Table 3 for HA/Ti interfaces with Zn and Ag dopants. All considered interfaces own positive γint and are in the order of 1Ca1Zn-HA(1)-Zn/Ti > 1Zn-HA(1)-Zn/Ti > 1Ca-HA(2)Zn/Ti > 3PO4-HA(2)-Zn/Ti for the HA/Ti semicycle interface with Zn dopants, while the order is 1Ca-HA(2)Ag/Ti > 1Ag-HA(1.2)-Ag/Ti > 3PO4-HA(2)-Ag/Ti for the HA/Ti interfaces with Ag dopant. Like interfaces composed of different compounds, the interfaces between HA and Ti are consist of quite different crystalline structures and chemical

terminated interface with an Ag dopant. Upon relaxation, the interlayer separations are increased about 13.4% and 4.1% for the interface and subinterface PO4 layers, respectively. The interlayer separation deviations of the Ti(0001) slab from the interface to central are −12.0%, −6.3%, and −6.8%, respectively. In general, the Ti−O bonds across the HA/Ti interfaces form and cause an expansion along the c axis, rotations of PO4 groups, and off-interface movements of metal atoms (Ca and Zn or Ag dopant) in the HA(0001) slab, but a shrinkage of the interlayer distances in the Ti slab. The relaxed interface structures of aHA/Ti are presented in Figure 3. Similar to the HA/Ti interface, the aHA/Ti interfaces also forms different numbers of Ti−O bonds between the Ti and the O from the PO4 groups in aHA at the interface. After optimization, the top two Ti layers show a zigzag structure. Moreover, a strong distortion in the atomic positions occurred in the aHA structure due to its amorphous structure, especially for Ca and O atoms at the interface and subinterface. 3.2. Work of Adhesion. On the basis of the supercell calculations, works of adhesion Wad for interfaces with Zn and Ag dopants were listed in Table 3, together with total energies of the considered interfaces, HA(0001) slabs with dopants, and Ti(0001) slabs. The work of adhesion Wad is calculated using the same method described in ref 56. From Table 3, for Zn doped interfaces, it is obvious that the combination of 3PO4HA/Ti is stronger than that of other interfaces with metal atom terminations. The 1Zn-HA(1)-Zn/Ti with one Zn atom termination shows a relatively weak combination, but much stronger than the corresponding undoped interface and partly doped 1Ca1Zn-HA(1)-Zn/Ti interface. Comparing with our previous theoretical study on the interfaces between stoichiometric HA(0001) slab and Ti(0001) slab,56 the results in Table 3 indicate that the Zn dopants enhance the bonding strength of the HA/Ti interface, especially when Zn dopants occupy the Ca(2) site of HA crystals. For instance, the most negative adhesive work of zinc doped interfaces (Wad) is −3.52 J/m2 while it is −2.33 J/m2 for the undoped interfaces, and less negative Wad is −1.71 and −0.8 J/m2 for HA/Ti interfaces with Zn occupying the Ca(2) site and HA/Ti without dopant, respectively. As shown in Table 3, the calculated work of adhesion of the PO4-terminated HA/Ti interface with Ag/(Ag + Ca) = 5% (Ag dopants occupy at the Ca(2) site) are more negative than 1CaHA(2)-Ag/Ti and 1Ag-HA(1,2)-Ag/Ti interfaces and therefore gives stronger adhesion. The one Ag-atom-terminated interface, 1Ag-HA(1,2)-Ag/Ti, reveals a slightly stronger H


Article


Figure 4. Total and partial density of states for Zn doped HA/Ti interfaces: (a) 1Ca-HA(2)-Zn/Ti and (b) 3PO4-HA(2)-Zn/Ti interfaces. The insets in part a show Zn, Ti, and PO4 groups and an OH unit and in part b show the interfacial PO4 group and Zn.

dopants, there are some new features occurring in Figure 4a. First, the bonding states are distributed throughout the energy rang from Fermi energy to −8.6 eV and the interfacial Ti d electrons show many new peaks at around −8.1 and −8.6 eV and in the energy range from −7.1 to −3.5 eV, which interact with interfacial PO4 group p electrons at the same energy levels of the 1Ca-HA(2)-Zn slab. These interaction states contribute to a pd hybridization between interfacial O and Ti atoms, definitely verifying the formation of valence bonds across the Zn doped interface. They are the major bonding peaks of the Ti d and O p electrons of the Ca-HA(2)-Zn/Ti system. Second, the density of states for Ca 3s, 3p electrons around Fermi energy is very small, indicating that the Ti-Ca interaction is weak. Finally, the O p electrons from the interfacial OH unit distributed around −8.6 and −8.1 eV interact with interfacial Ti 3d electrons that appear as peaks at the same energies. A conclusion can be drawn that the Ti−O interactions are the main valence bonding at the interface as there is higher coordination among Ti and O atoms than in the Ca or Zn dopants. The DOSs for interfacial PO4 groups show strong interactions among interfacial PO4 group p states and Zn 3d in the valence band at about −7.1 eV in Figure 4a. Comparing the DOSs of same atoms (or groups) in interfacial layers with that in interior bulk (Figure 4a), one can find that the aggregation and delocalization of valence charges of the interfacial atoms are more distinct than that of atoms in

bonds. As a result, interface energies can be considered relatively high. That implies the interfacial alloying is not likely to happen and that the HA/Ti interface has the small possibility of forming a new phase. However, for PO4terminated HA/Ti interfaces with dopants, the γint values are relatively lower than that of undoped PO4-terminated HA/Ti interfaces;56 that is to say, Zn and Ag dopants decrease the HA/Ti interface energy and therefore enhance the possibility of interfacial alloying or forming new phase. One reason for this result can be attributed to the reduced lattice mismatch between the (0001) interfaces by Zn and Ag dopants. 3.4. Electronic Structures and Bonding Interactions. The electronic structures of Zn and Ag doped interfaces were calculated and some of them as examples were shown in Figures 4 and 5. Parts a and b of Figure 4 show the density of states (DOSs) for the systems of 1Ca-HA(2)-Zn/Ti and 3PO4-HA(2)-Zn/Ti, most strong interfacial interactions configurations for Zn doped HA/Ti, while electronic structures of the 3PO4-HA(2)-Ag/Ti and 1Ag-HA(1.2)-Ag/Ti configurations for HA/Ti with Ag dopants are elaborated (Figure 5a,b). Figure 4a is the density of states for the Ca-HA(2)-Zn/Ti system, where the partial density of states (PDOS) of PO4 groups in the energy range −3.5 to +2.0 eV and the Ti atoms in the energy range −8.0 to −2.0 eV are highlighted in the inset. Although the metallic characteristics was not alter by the I


Article


Figure 5. Total and partial DOSs of (a) 1Ag-HA(1,2)-Ag/Ti and (b) 3PO4-HA(2)-Ag/Ti interfaces.

We further evaluated the charge density difference using equation Δρ(r) = ρHA/Ti(r) − ρHA(0001)(r) − ρTi(0001)(r) and the meaning of parameters in the formula is referred to in ref 56. The calculated charge rearrangement and charge different distributions (CDD) projected on (101̅0) of 1Ca-HA(2)-Zn/ Ti interface is shown in Figure S14a,b, which reveal nonuniform interaction behavior at the interface between HA and Ti slabs. Charge transfer can be observed from Ti (light blue) to the HA (yellow) near the interface confined to two atomic layers in Figure S14a. The charge transfer also alters the charge distribution around the interfacial Ca and Ti atoms, which are circled in a charge distribution ball (yellow). As shown in Figure S14b, the O-sites (red) are electron accumulation and Ti-sites (blue) are electron depletion. Thus, the Ti atom was an electron donor as the charge different distributions around it were negative and the O atom formed an ionic bond with Ti1 atom at the interface. Positive CDD was found in the region around interfacial Ca and Ti2 atom, which means that interfacial Ca atom formed a metallic bonding with Ti atom. The total and partial DOSs of the 3PO4-HA(2)-Zn/Ti system are illustrated in Figure 4b. The insets highlight the partial DOSs of interfacial PO4 group in the energy range from −2.0 to +2.0 eV and that of Zn atom in the energy region of −3.0 to +2.0 eV. The partial DOSs of the PO4 group, Ti atoms, and Zn atom show the difference between the 1Ca-HA(2)-Zn/ Ti and 3PO4-HA(2)-Zn/Ti systems. The magnitude of the interface PO4 p states and Ti d states were increased between

interior bulk. For instance, the integral valence electron of the nine Ti atoms at the interface layer is much smaller than that of the nine Ti atoms at the subinterface layer, and the PO4 group shows the contrast consequence. Thus, there are more valence electrons transferred from Ti atoms at the interface to other atoms than those at the subinterface layer. For the PO4 group in the central interface layer, the valence charge accumulations are quite obvious by comparing the integral valence electron of the interfacial PO4 group with that in the interior bulk of the HA side. That shows the interfacial bonding is mainly contributed by the interactions among atoms which are just at the interface from HA and Ti sides. The partial DOSs show that there are electron accumulations of both PO4 and OH groups at the interface in the energy range from −8.9 to −3.0 eV. Conversely, the partial DOS of the interface Ti atom shows that there are quite a few depletion states in the energy range from −2.0 to +2.0 eV. The interfacial Ti 4s and 3d have made great contribution by charge transfers across the HA/Ti interface and their DOSs are delocalization. Therefore, the valence band is mainly composed of s, p, and d states of the HA(0001) side with an admixture of Ti 3d states for HA/Ti interfaces, the conduction bands are dominated by Ti 3d states, and Ti−O bonds formed on the 1Ca-HA(2)-Zn/Ti interface are probably ionic. The interaction between p orbitals from PO4 group and the Zn 3d orbital is obviously stronger than that between PO4 group and Ca atom and the ionic characteristics of HA(0001) slab are strengthened. J


Article

ACS Applied Nano Materials −2.0 eV and the Fermi energy. The partial DOSs of Zn p d states appear at −2.6 and −2.0 eV. These changes indicate that the PO4 unit strongly interacts with the Ti atom owning to the increased overlaps between the two groups in the energy range from −2.0 eV to the Fermi energy. The partial DOSs of Ca p and Ti d electrons appear at the Fermi energy, showing the metallic characteristics of the interface. It is worthwhile to note that the partial DOS of the Ti-interface 4s states was depleted compared with that of the bulk Ti atoms in the energy range from −5.5 to −2.8 eV, while a large amount of PO4 p states occurs in the same energy region. That is to say, the Ti 4s electrons were transferred to O atom causing the interfacial Ti to form strong ionic bonds with O atoms. This may be why the 3PO4-HA(2)-Zn/Ti system has a more negative adhesion energy than the 1Ca-HA(2)-Zn/Ti interface as listed in Table 3. Charges are transferred from interfacial Ti atom to interfacial O atom, and thus negative PO4-terminated HA with Zn dopants at the Ca(2) site on Ti(0001) surface are generated, as shown in Figure S14c,d. In contrast to the case of 1Ca-HA(2)-Zn/Ti interface, the number of O atoms at the interface is increased and the amount of charge transfer is much larger. The most important feature with PO4 group at the interface is that the region of charge redistribution is tiny, which results in the more concentrated interaction among interface atoms (group). Charge transfer between the Zn dopants and Ti at the interface is much larger than that between interfacial Ca and Ti in the 1Ca-HA(2)-Zn/Ti interface. And the metallic characteristic of the 3PO4-HA(2)Zn/Ti interface is slightly strengthened. Therefore, the bonding interaction across the interface is strong, mainly arising from O, Ti, and Zn atoms near the interface. In Figure 5a, we provide the DOSs of 1Ag-HA(1,2)-Ag/Ti and the partial DOSs of the interfacial PO4 group, interfacial Ti atoms, PO4 units in interior bulk layer, and interfacial Ag atom are stressed in the inset. Comparing with the DOSs of the 1CaHA(2)-Zn/Ti interface, the DOSs of the 1Ag-HA(1,2)-Ag/Ti interface and the PDOS of the interfacial PO4 group, Ca atoms, and Ti atom are changed. There is a new sharp peak of interfacial Ti 3d states at about −7.1 eV and it overlaps with the peak of interfacial PO4 p states at the same energy level. Although this overlap will contribute to the interaction between Ti and PO4, the main part of Ti d orbitals distribute around the Fermi energy, while the PDOS of PO4 p electrons at the Fermi level is low enough to be ignored, as shown in the insets in Figure 5a, thus the weak PO4 interaction with the Ti atom. The partial DOSs of Ag dopants show a sharp peak at the Fermi energy and overlaps with the Ti 3d state at the same energy level, indicating Ag−Ti metallic bonding across the interface. It is worthwhile to mention that the partial DOS of the interfacial Ag dopant is more localized than that of Ag in the interior bulk layer. Furthermore, the interfacial Ag 4d has a sharp peak at about −4.4 eV and it overlaps with the PO4 p state at the same energy level. Thus, there is a stronger interaction between interfacial Ag and PO4 units than the interaction between interiors Ag and PO4 group. The threedimensional charge rearrangement and CDD on the (101̅0) surface of this system are shown in Figure S15a,b. In contrast to the case of the 1Zn-HA(1)-Zn/Ti interface, the amount of charge transfer is much smaller. Similarly, the charges accumulate in the vicinity of interfacial Ag and O atoms and deplete around the interfacial Ti atom. Clearly, the ionic Ti−O bond across the interface can be observed in Figure S15b.

The DOSs of 3PO4-HA(2)-Ag/Ti are depicted in Figure 5 together with the partial DOSs of interfacial atoms and units as the insets. The partial DOSs of the PO4 group, Ti atoms, and Ag atom show differences between the systems of 1AgHA(1,2)-Ag/Ti and 3PO4-HA(2)-Ag/Ti. A stronger bond at the 3PO4-HA(2)-Ag/Ti interface is found owing to p−d hybridization of the PO4-interface p states and Ti-interface d states around the Fermi energy level as displayed in the insets. Ca s, p states and Ag s, d states appear across the Fermi energy and make their contributions to the metallic bonding across the interface. In the 3PO4-HA(2)-Ag/Ti interface, the partial DOS of interfacial Ag at the Ca(2) site is quite weakly localized compared with that of interfacial Ag dopants at the Ca(1) site in the 1Ag-HA(1,2)-Ag/Ti interface. It is worthwhile to note that the amplitude of interfacial Ca p sharp peaks in the energy range −6.4 to +5.6 eV is much larger than that of Ca in the bulk layer and overlaps with PO4 p states in the same energy range. The charge rearrangement is illustrated in Figure S15c, and CDD on the (101̅0) surface of this system is shown in Figure S15d. Charge transfer can be observed from Ti (light blue color) to the O (yellow color) and charge accumulation is obviously around the Ag dopant near the interface. In comparison with the charge transfer of the Ag-terminated system (1Ag-HA(1,2)-Ag/Ti), the amount of electron transfer of the PO4 termination interface is much larger and the number of Ti−O ionic bonding is increased. In summary, for the Zn and Ag doped HA(000)/Ti(0001) systems, the 3PO4-terminated interfaces are more stable and preferable than the systems with metal atom terminations. The Ti−O ionic bonds are the dominant interactions at the interface. There are also metallic interactions across the interface, mainly arising from interfacial metal atom, such as Ti, Ag, and Zn atoms in metal-atom-terminated HA/Ti interfaces. Adhesion between HA(0001) and Ti(0001) slabs are enhanced by Zn and Ag dopants. In Figure 6, the difference plane-average charge (PAC) was provided for the considered systems. The PAC presents the spatial distribution of charge and its rearrangement along the specified direction. Since we are chiefly interested in the charge distribution across the interface and in its vicinity, the direction perpendicular to the interface plane was selected. To get an insight into the bonding mechanism and charge distribution in the vicinity of the interface, the PACs of the whole interface and its individual slabs are calculated. Difference of planeaverage charge is obtained by the PAC of the whole structure subtracting the sum of PACs of the constituents. Figure 6a shows the difference of PAC of the 1Ca-HA(2)Zn/Ti system. The first remarkable feature of the changes of electronic density across the interface is its giant inhomogeneity in the HA slab as compared to that in the titanium slab. The second characteristic is a relatively small charge rearrangement in the region around the interface (in Figure 6a). For the 1Ca-HA(2)-Zn/Ti system, the difference of PAC at the interface region was the averaged charge transfer arising mainly among interfacial O, Ca, and Ti atoms. However, it is a completely contrary charge transfer behavior for O and Ca atoms on the HA side just at the interface, i.e., O accepting electron and Ca donating electron, which results in the small rearrangement of charge at the interface. The charge rearrangement happens mostly in a 0.5 nm region around the interface. In Figure 6b, we provide the difference of PAC of the 3PO4HA(2)-Zn/Ti interface. It is clear that charges are accumulated K


Article


interface, the PAC difference is smaller than for 3PO4-HA(2)Zn/Ti and 1Ca-HA(2)-Zn/Ti interfaces; however, the charge rearrangement also happens in a region of 0.5 nm thickness around the interface. Figure 6d is the calculated PAC difference of the 3PO4HA(2)-Ag/Ti system. The spatial charge distribution is quite similar to 3PO4-HA(2)-Zn/Ti except for the strong inhomogeneous charge distribution in the Ti side. While the charge rearrangement across the interface also resembles the pattern of the 3PO4-HA(2)-Zn/Ti with lower amplitude. The main feature is charge accumulation in the interfacial O layer and charge depletion of interfacial Ti layer. To summarize, the difference PACs of the four considered structures uncover a varying degree of charge rearrangement, depending on stacking and dopants. The change of the charge distribution primarily occurs in a 0.3−0.5 nm thick area across the interface. The interfacial binding strength represented by adhesive work is related to the difference of PACs. In Figure S16, this correlation is quite obvious when the work of adhesion is plotted as the integral of charge transfer. The latter is obtained by integrating the absolute values of the difference PACs along the c axis and could be regarded as a rough measurement of the total amount of charge transfer. The value of work of adhesion for 1Ag-HA(1,2)-Ag/Ti system is −1.71 J/m2 as its charge transfer is quite small. The 3PO4HA(2)-Zn/Ti and 3PO4-HA(2)-Ag/Ti structures possess greater charge transfers, and thus works of adhesion for them are of −3.52 and −3.48 J/m2, respectively. The 1Ca-HA(2)Zn/Ti interface shows a slightly larger charge transfer than the 1Ag-HA(1,2)-Ag/Ti interface, and its work of adhesion is −2.55 J/m2. From Figure S16 one can note that the two interfaces of 3PO4-HA(2)-Zn/Ti and 3PO4-HA(2)-Ag/Ti show similar works of adhesion, but a large amount of charge transfer occurs in the former, while it is relatively small in the latter. To clarify the insight of this phenomenon, the Bader charge analysis was performed to quantitatively analyze the charge transfer between atoms at the interface zone. The Bader atomic charges of elements in interfaces 3PO4-HA(2)-Zn/Ti and 3PO4-HA(2)-Ag/Ti are listed in Table 4. In the 3PO4-HA(2)Table 4. Bader Atomic Charges of Elements in the First Layer of Interfaces 3PO4-HA(2)-Zn/Ti and 3PO4-HA(2)Ag/Ti in Units of e

Figure 6. Difference of plane-average charge (PAC) of 1Ca-HA(2)Zn/Ti interface (a), 3PO4-HA(2)-Zn/Ti interface (b), 1Ag-HA(1,2)Ag/Ti interface (c), and 3PO4-HA(2)-Ag/Ti interface (d).

interface

P

O

Ca

3PO4-HA(2)-Zn/Ti

−2.7 to −2.5 −3.65 to −3.59

1.20 to 1.35 1.45− 1.48

−1.47 to −1.43 −1.6

3PO4-HA(2)-Ag/Ti

Zn/ Ag −0.51 −0.39

Ti −0.45 to −0.26 −0.66 to −0.25

Zn/Ti interface, 9 Ti atoms in the Ti side face up to 5 O atoms, 3 P atoms, 2 Ca atoms, and 1 Zn atom in the HA side, making 10 Ti−O bonds crossing the interface. Our calculation results show that Ti, P, Ca, and Zn are electron donators, the nine Ti atoms lose charges in the range of 0.26 to 0.45 e, the P atom loses about 2.6 e, the Ca atom loses about 1.5 e, and the Zn atom loses 0.5 e, while five O atoms gain charges about 1.35 e, 1.24 e, 1.20 e, 1.30 e, and 1.25 e, respectively. In the 3PO4-HA(2)-Ag/Ti interface, six Ti atoms bonded with 4 O atoms forming 7 Ti−O bonds crossing the interface. And, Ti, P, Ca, and Ag lose electrons, while O atoms gain electrons. Six Ti atoms lose about 0.66 e, 0.34 e, 0.29 e, 0.32e, 0.31 e, and

at the interface O atom layer on the HA side. In contrast, charge depletion was found at the subinterface atoms layer on the HA slab side, including P, Ca, and Zn atoms. Similarly, the charge distributions in the Ti slab are quite homogeneous compared with that of the HA side. The glaring feature of the charge distribution in this system is the quite significant charge rearrangement on the 0.4 nm thick region in the vicinity of the interface. As shown in Figure 6c, the difference of PAC in the 1AgHA(1,2)-Ag/Ti interface is quite different from that of the 1Ca-HA(2)-Zn/Ti. Owing to the weaker bonding at the L


Article


the HA/Ti interface is enhanced by those two kinds of dopants.

0.25 e, while the O atoms gain about 1.45 e, 1.45 e, 1.48 e, and 1.49e, respectively. The P atoms lose charges about 3.65 e, 3.64 e, and 3.58 e, respectively, which are quite larger than that of P atoms at the 3PO4-HA(2)-Zn/Ti interface. That is to say, interfacial O atoms at the 3PO4-HA(2)-Ag/Ti interface get more charges from P than those at the 3PO4-HA(2)-Zn/Ti interface. Thus, interfacial O atoms at the 3PO4-HA(2)-Ag/Ti interface gain less charge from the Ti slab side. This may explain why the large difference of charge transfer between the 3PO4-HA(2)-Zn/Ti and 3PO4-HA(2)-Ag/Ti interfaces has no effect on the final value of the work of adhesion.

■

ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsanm.8b01103. Convergence checks regarding the surface energy, the cutoff energy, and the vacuum spacing of the Ti(0001) and HA (0001), parameters of the force field for hydroxyapatite, heating and quenching steps, and energy−temperature curves for simulated quenching molecular dynamics simulations, models of HA(0001) surfaces with Zn or Ag dopants and the atomic structures of different atom terminations, 1Ca, 1Zn, PO4, and 1(Ca+Zn) of HA(0001) surfaces with Zn dopants after relaxation (Figure S10), 1Ag-HA (1.2)-Ag, the 1Ca-HA (2)-Ag, and the PO4-HA (2)-Ag surface models (Figure S11) and their surface energies of the all considered surface (Table S3), three possible stacking structures of 1Zn-HA(1) -Zn/Ti (Figure S12) and their fitted curves of the adhesion energy Ead vs the interfacial separations (Figure S13), three dimensional charge arrangement and the projection of the CDD on the (101̅0) plane for HA/Ti interfaces with Zinc dopants (Figure S14) and those of HA/Ti interface with Ag dopants (Figure S15), work of adhesion as a function of integral charge transfer (Figure S16) (PDF)

4. CONCLUSION First-principles calculations are performed to investigate the influence of Zn and Ag dopants on the interfacial adhesive properties of HA(0001)/Ti(0001) systems. The work of adhesion and interface energies of various interface structures with different stacking configurations, terminations, and different doped sites and amorphous structures are evaluated. Both Zn and Ag dopants strengthen the adhesion properties of the HA/Ti interfaces, and the four most stable structures have been identified as 3PO4-HA(2)-Zn/Ti, 1Ca-HA(2)-Zn/Ti, 3PO4-HA(2)-Ag/Ti, and 1Ag-HA/Ti(1.2)-Ag. On the basis of the quenched HA(0001) structures, the amorphous HA/Ti interfaces were also constructed and studied using firstprinciples calculations. Influence of dopants on the bonding of the valence bonds across HA/Ti interface is investigated. Zn and Ag contained interfaces with PO4 terminations showing larger numbers of Ti−O bonds crossing the interface than the pure HA/Ti systems with the same termination. A maximum of ten Ti−O bonds with lengths from 0.21 to 0.24 nm occur in the 3PO4-HA(2)-Zn/Ti systems and seven Ti−O bonds with lengths about 0.20 nm formed in the 3PO4-HA(2)-Ag/Ti structure. Both Zn and Ag dopants strengthen the HA/Ti interface energetically. The works of adhesion for 3PO4HA(2)-Zn/Ti and 3PO4-HA(2)-Ag/Ti interfaces are −3.52 and −3.48 J/m2, respectively, which are more negative than that of the pure HA/Ti interface with the same termination. The 1(Ca+Zn)-HA(1)-Zn/Ti interface and one-metal-terminated stacking with Ag dopants, the 1Ca-HA(2)-Ag/Ti and 1Ag-HA(1,2)-Ag/Ti interfaces, have the weak adhesion but still stronger than adhesion in the corresponding the pure HA/ Ti interfaces. The stronger adhesion properties induced by Zn and Ag dopants will enhance shock and stress resistances of the HA-coated Ti implants. Zn and Ag dopants also have the function of antibacterial and promoting osteoblast response, thus doping of Zn or Ag will improve the clinical application performances of HA/Ti implants. The work of adhesion and stacking configurations for the interfaces between amorphous HA and Ti (0001) are also studied. The PO4-aHA/Ti interface stacked as O−O−|Ti−Ti−··· has the most negative value of Wad −4.33 J/m2. Electronic structure analysis of the HA/Ti interfaces has shown that the atomic coordination and dopants at the interface result in many characteristics in DOSs, which range from metallization to p−d hybridization of O p states and Ti d states around the Fermi energy level for all considered interfaces. We found that the key factors to determine the work of adhesion are the changes of charge distribution around the interface and the chemical bonding characteristics (ionic or covalent); i.e., the interface that has great charge rearrangement and covalent bonding across the interface shows the strong adhesion. The amount of charge transfer across the interface is increased by Zn or Ag dopants and the affinity of

■

AUTHOR INFORMATION

Corresponding Author

*Y. Song. E-mail: [email protected]. Telephone: +86-6315687772. ORCID

Yan Song: 0000-0002-9081-6518 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are thankful to Gang Chen from Harbin Institute of Technology at Weihai for his generous assistance with MD investigations. This work was supported by The Key research and development plan of Shandong, China, Grant No. 2018GGX103020, The National Key Research and Development Program of China, Grant No. 2016YFB0701301, the Natural Science Foundation of Shandong, China, Grant No. ZR2014EMM013, and the Fundamental Research Funds for the Central Universities, Grant No. HIT.NSRIF.2017013. Simulations were performed using HPC resources in CAS Shenyang Supercomputing Center.

■

REFERENCES

(1) Dorozhkin, S. V. Self-setting Calcium Orthophosphate Formulations. J. Funct. Biomater. 2013, 4, 209−311. (2) Dorozhkin, S. V. Calcium Orthophosphates Occurrence, Properties, Biomineralization, Pathological Calcification and Biomimetic Applications. Biomatter. 2011, 1, 121−164. (3) Zhao, Z.; Du, L.; Tao, Y.; Li, Q.; Luo, L. Enhancing the Adhesion Strength of Micro Electroforming Layer by Ultrasonic Agitation Method and the Application. Ultrason. Sonochem. 2016, 33, 10−17. M


Article


Hydroxyapatite Prepared by Sol-Gel Method. Sci. World J. 2014, 2014, 165351. (22) Zhang, Y. G.; Zhu, Y. J.; Chen, F.; Sun, T. W. Biocompatible, Ultralight, Strong Hydroxyapatite Networks Based on Hydroxyapatite Microtubes with Excellent Permeability and Ultralow Thermal Conductivity. ACS Appl. Mater. Interfaces 2017, 9, 7918−7928. (23) Yang, R. L.; Zhu, Y. J.; Chen, F. F.; Dong, L. Y.; Xiong, Z. C. Luminescent, Fire-Resistant, and Water-Proof Ultralong Hydroxyapatite Nanowire-Based Paper for Multimode Anticounterfeiting Applications. ACS Appl. Mater. Interfaces 2017, 9, 25455−25464. (24) Wei, X.; Fu, C.; Savino, K.; Yates, M. Z. Fully Dense YttriumSubstituted Hydroxyapatite Coatings with Aligned Crystal Domains. Cryst. Growth Des. 2012, 12, 217−223. (25) Wei, X.; Fu, C.; Savino, K.; Yates, M. Z. Carbonated Hydroxyapatite Coatings with Aligned Crystal Domains. Cryst. Growth Des. 2012, 12, 3474−3480. (26) Shepherd, D.; Best, S. M. Production of Zinc Substituted Hydroxyapatite using Various Precipitation Routes. Biomed. Mater. 2013, 8, 025003. (27) Shalom, H.; Feldman, Y.; Rosentsveig, R.; Pinkas, I.; KaplanAshiri, I.; Moshkovich, A.; Perfilyev, V.; Rapoport, L.; Tenne, R. Electrophoretic Deposition of Hydroxyapatite Film Containing ReDoped MoS(2) Nanoparticles. Int. J. Mol. Sci. 2018, 19, 657. (28) Bakhsheshi-Rad, H. R.; Hamzah, E.; Ismail, A. F.; KasiriAsgarani, M.; Daroonparvar, M.; Parham, S.; Iqbal, N.; Medraj, M. Novel Bi-Layered Nanostructured SiO2 /Ag-FHAp Coating on Biodegradable Magnesium Alloy for Biomedical Applications. Ceram. Int. 2016, 42, 11941−11950. (29) Ban, S.; Maruno, S. Hydrothermal−Electrochemical Deposition of Hydroxyapatite. J. Biomed. Mater. Res. 1998, 42, 387−395. (30) Carrado, A. Structural, Microstructural, and Residual Stress Investigations of Plasma-Sprayed Hydroxyapatite on Ti-6Al-4 V. ACS Appl. Mater. Interfaces 2010, 2, 561−565. (31) Teshima, K.; Wagata, H.; Sakurai, K.; Enomoto, H.; Mori, S.; Yubuta, K.; Shishido, T.; Oishi, S. High-Quality Ultralong Hydroxyapatite Nanowhiskers Grown Directly on Titanium Surfaces by Novel Low-Temperature Flux Coating Method. Cryst. Growth Des. 2012, 12, 4890−4896. (32) Lopez, E. O.; Mello, A.; Sendao, H.; Costa, L. T.; Rossi, A. L.; Ospina, R. O.; Borghi, F. F.; Silva Filho, J. G.; Rossi, A. M. Growth of Crystalline Hydroxyapatite Thin films at Room Temperature by Tuning the Energy of the RF-Magnetron Sputtering plasma. ACS Appl. Mater. Interfaces 2013, 5, 9435−9445. (33) Solla, E. L.; Rodríguez-González, B.; Aguiar, H.; RodríguezValencia, C.; Serra, J.; González, P. Revealing the Nanostructure of Calcium Phosphate Coatings using HRTEM/FIB Techniques. Mater. Charact. 2016, 122, 148−153. (34) Grubova, I. Y.; Surmeneva, M. A.; Huygh, S.; Surmenev, R. A.; Neyts, E. C. Density Functional Theory Study of Interface Interactions in Hydroxyapatite/Rutile Composites for Biomedical Applications. J. Phys. Chem. C 2017, 121, 15687−15695. (35) Brandão-Neto, J.; Stefan, V.; Mendonça, B. B.; Bloise, W.; Castro, A. V. B. The Essential Role of Zinc in Growth. Nutr. Res. (N. Y., NY, U. S.) 1995, 15, 335−358. (36) Jin, G.; Qin, H.; Cao, H.; Qiao, Y.; Zhao, Y.; Peng, X.; Zhang, X.; Liu, X.; Chu, P. K. Zn/Ag Micro-Galvanic Couples Formed on Titanium and Osseointegration Effects in the Presence of S. aureus. Biomaterials 2015, 65, 22−31. (37) Zhong, Z.; Ma, J. Fabrication, Characterization, and in Vtro Study of Zinc Substituted Hydroxyapatite/Silk Fibroin Composite Coatings on Titanium for Biomedical Applications. J. Biomater. Appl. 2017, 32, 399−409. (38) Ma, J.; Qin, J. Graphene-like Zinc Substituted Hydroxyapatite. Cryst. Growth Des. 2015, 15, 1273−1279. (39) Yuan, Q.; Wu, J.; Qin, C.; Xu, A.; Zhang, Z.; Lin, Y.; Chen, Z.; Lin, S.; Yuan, Z.; Ren, X.; Zhang, P. One-Pot Synthesis and Characterization of Zn-Doped Hydroxyapatite Nanocomposites. Mater. Chem. Phys. 2017, 199, 122−130.

(4) Huang, Y.; Hao, M.; Nian, X.; Qiao, H.; Zhang, X.; Zhang, X.; Song, G.; Guo, J.; Pang, X.; Zhang, H. Strontium and Copper CoSubstituted Hydroxyapatite-Based Coatings with Improved Antibacterial Activity and Cytocompatibility Fabricated by Electrodeposition. Ceram. Int. 2016, 42, 11876−11888. (5) Han, C. D.; Shin, K. Y.; Lee, H. H.; Park, K. K.; Yang, I. H.; Lee, W. S. The Results of Long-Term Follow-up of Total Hip Arthroplasty using Hydroxyapatite-Coated Cups. Hip & pelvis 2015, 27, 209−215. (6) Zhang, X.; Chaimayo, W.; Yang, C.; Yao, J.; Miller, B. L.; Yates, M. Z. Silver-Hydroxyapatite Composite Coatings with Enhanced Antimicrobial Activities through Heat Treatment. Surf. Coat. Technol. 2017, 325, 39−45. (7) Tao, Z. S.; Zhou, W. S.; He, X. W.; Liu, W.; Bai, B. L.; Zhou, Q.; Huang, Z. L.; Tu, K. K.; Li, H.; Sun, T.; Lv, Y. X.; Cui, W.; Yang, L. A Comparative Study of Zinc, Magnesium, Strontium-Incorporated Hydroxyapatite-Coated Titanium Implants for Osseointegration of Osteopenic Rats. Mater. Sci. Eng., C 2016, 62, 226−232. (8) Xie, C. M.; Lu, X.; Wang, K. F.; Meng, F. Z.; Jiang, O.; Zhang, H. P.; Zhi, W.; Fang, L. M. Silver Nanoparticles and Growth Factors Incorporated Hydroxyapatite Coatings on Metallic Implant Surfaces for Enhancement of Osteoinductivity and Antibacterial Properties. ACS Appl. Mater. Interfaces 2014, 6, 8580−8589. (9) Peñaflor Galindo, T. G.; Kataoka, T.; Fujii, S.; Okuda, M.; Tagaya, M. Preparation of Nanocrystalline Zinc-Substituted Hydroxyapatite Films and Their Biological Properties. Colloid and Interface Science Communications 2016, 10−11, 15−19. (10) Mirzaee, M.; Vaezi, M.; Palizdar, Y. Synthesis and Characterization of Silver Doped Hydroxyapatite Nanocomposite Coatings and Evaluation of Their Antibacterial and Corrosion Resistance Properties in Simulated Body Fluid. Mater. Sci. Eng., C 2016, 69, 675−684. (11) Moreira, M. P.; de Almeida Soares, G. D.; Dentzer, J.; Anselme, K.; de Sena, L. A.; Kuznetsov, A.; dos Santos, E. A. Synthesis of Magnesium- and Manganese-Doped Hydroxyapatite Structures Assisted by the Simultaneous Incorporation of Strontium. Mater. Sci. Eng., C 2016, 61, 736−743. (12) Wang, M.; Wang, Q.; Lu, X.; Wang, K.; Ren, F. Computer Simulation of Ions Doped Hydroxyapatite: A Brief Review. J. Wuhan Univ. Technol., Mater. Sci. Ed. 2017, 32, 978−987. (13) Š upová, M. Substituted Hydroxyapatites for Biomedical Applications: A review. Ceram. Int. 2015, 41, 9203−9231. (14) Hidalgo-Robatto, B. M.; López-Á lvarez, M.; Azevedo, A. S.; Dorado, J.; Serra, J.; Azevedo, N. F.; González, P. Pulsed Laser Deposition of Copper and Zinc Doped Hydroxyapatite Coatings for Biomedical Applications. Surf. Coat. Technol. 2018, 333, 168−177. (15) Andres, N. C.; D’Elia, N. L.; Ruso, J. M.; Campelo, A. E.; Massheimer, V. L.; Messina, P. V. Manipulation of Mg2+-Ca2+ Switch on the Development of Bone Mimetic Hydroxyapatite. ACS Appl. Mater. Interfaces 2017, 9, 15698−15710. (16) Ratnayake, J. T. B.; Mucalo, M.; Dias, G. J. Substituted Hydroxyapatites for Bone Regeneration: A Review of Current Trends. J. Biomed. Mater. Res., Part B 2017, 105, 1285−1299. (17) Nakamura, M.; Hiratai, R.; Hentunen, T.; Salonen, J.; Yamashita, K. Hydroxyapatite with High Carbonate Substitutions Promotes Osteoclast Resorption through Osteocyte-like Cells. ACS Biomater. Sci. Eng. 2016, 2, 259−267. (18) Huang, T.; Xiao, Y.; Wang, S.; Huang, Y.; Liu, X.; Wu, F.; Gu, Z. Nanostructured Si, Mg, CO32− Substituted Hydroxyapatite Coatings Deposited by Liquid Precursor Plasma Spraying: Synthesis and Characterization. J. Therm. Spray Technol. 2011, 20, 829−836. (19) Wei, X.; Fu, C.; Savino, K.; Yates, M. Z. Fully Dense YttriumSubstituted Hydroxyapatite Coatings with Aligned Crystal Domains. Cryst. Growth Des. 2012, 12, 217−223. (20) Lin, K.; Liu, P.; Wei, L.; Zou, Z.; Zhang, W.; Qian, Y.; Shen, Y.; Chang, J. Strontium Substituted Hydroxyapatite Porous Microspheres: Surfactant-Free Hydrothermal Synthesis, Enhanced Biological Response and Sustained Drug Release. Chem. Eng. J. 2013, 222, 49−59. (21) Iconaru, S. L.; Chapon, P.; Le Coustumer, P.; Predoi, D. Antimicrobial Activity of Thin Solid Films of Silver Doped N


Article

ACS Applied Nano Materials (40) Gao, A.; Hang, R.; Chu, P. K. Recent Advances in AntiInfection Surfaces Fabricated on Biomedical Implants by PlasmaBased Technology. Surf. Coat. Technol. 2017, 312, 2−6. (41) Qiao, Y.; Zhang, W.; Tian, P.; Meng, F.; Zhu, H.; Jiang, X.; Liu, X.; Chu, P. K. Stimulation of Bone Growth Following zinc Incorporation into Biomaterials. Biomaterials 2014, 35, 6882−6897. (42) Vukomanović, M.; Repnik, U.; Zavašnik-Bergant, T.; Kostanjšek, R.; Š kapin, S. D.; Suvorov, D. Is Nano-Silver Safe within Bioactive Hydroxyapatite Composites? ACS Biomater. Sci. Eng. 2015, 1, 935−946. (43) Xie, C. M.; Lu, X.; Wang, K. F.; Meng, F. Z.; Jiang, O.; Zhang, H. P.; Zhi, W.; Fang, L. M. Silver Nanoparticles and Growth Factors Incorporated Hydroxyapatite Coatings on Metallic Implant Surfaces for Enhancement of Osteoinductivity and Antibacterial Properties. ACS Appl. Mater. Interfaces 2014, 6, 8580−8589. (44) Zhao, Y.; Cao, H.; Qin, H.; Cheng, T.; Qian, S.; Cheng, M.; Peng, X.; Wang, J.; Zhang, Y.; Jin, G.; Zhang, X.; Liu, X.; Chu, P. K. Balancing the Osteogenic and Antibacterial Properties of Titanium by Codoping of Mg and Ag: An in Vitro and in Vivo Study. ACS Appl. Mater. Interfaces 2015, 7, 17826−17836. (45) Simchi, A.; Tamjid, E.; Pishbin, F.; Boccaccini, A. R. Recent Progress in Inorganic and Composite Coatings with Bactericidal Capability for Orthopaedic Applications. Nanomedicine 2011, 7 (1), 22−39. (46) Erakovic, S.; Jankovic, A.; Veljovic, D.; Palcevskis, E.; Mitric, M.; Stevanovic, T.; Janackovic, D.; Miskovic-Stankovic, V. Corrosion Stability and Bioactivity in Simulated Body Fluid of Silver/ Hydroxyapatite and Silver/Hydroxyapatite/Lignin Coatings on Titanium Obtained by Electrophoretic Deposition. J. Phys. Chem. B 2013, 117, 1633−1643. (47) Roy, M.; Fielding, G. A.; Beyenal, H.; Bandyopadhyay, A.; Bose, S. Mechanical, in Vitro Antimicrobial, and Biological Properties of Plasma-Sprayed Silver-Doped Hydroxyapatite Coating. ACS Appl. Mater. Interfaces 2012, 4, 1341−1349. (48) Gokcekaya, O.; Webster, T. J.; Ueda, K.; Narushima, T.; Ergun, C. In vitro Performance of Ag-Incorporated Hydroxyapatite and its Adhesive Porous Coatings Deposited by Electrostatic Spraying. Mater. Sci. Eng., C 2017, 77, 556−564. (49) Yan, Y.; Zhang, X.; Huang, Y.; Ding, Q.; Pang, X. Antibacterial and Bioactivity of Silver Substituted Hydroxyapatite/TiO2 Nanotube Composite Coatings on Titanium. Appl. Surf. Sci. 2014, 314, 348− 357. (50) Geng, Z.; Wang, R.; Zhuo, X.; Li, Z.; Huang, Y.; Ma, L.; Cui, Z.; Zhu, S.; Liang, Y.; Liu, Y.; Bao, H.; Li, X.; Huo, Q.; Liu, Z.; Yang, X. Incorporation of Silver and Strontium in Hydroxyapatite Coating on Titanium Surface for Enhanced Antibacterial and Biological Properties. Mater. Sci. Eng., C 2017, 71, 852−861. (51) Jin, H.; Zhang, Z.; Nie, Y.; Zeng, Y.; Shen, L.; Sullivan, M. B.; Wang, S. J. Interfacial Structure of Ti2AlN Thin Films on MgO(111). J. Phys. Chem. C 2013, 117, 16515−16522. (52) Surmeneva, M. A.; Surmenev, R. A.; Nikonova, Y. A.; Selezneva, I. I.; Ivanova, A. A.; Putlyaev, V. I.; Prymak, O.; Epple, M. Fabrication, Ultra-Structure Characterization and in Vitro Studies of RF Magnetron Sputter Deposited Nano-Hydroxyapatite Thin Films for Biomedical Applications. Appl. Surf. Sci. 2014, 317, 172− 180. (53) Banerjea, A.; Smith, J. R. Origins of the Universal BindingEnergy Relation. Phys. Rev. B: Condens. Matter Mater. Phys. 1988, 37 (12), 6632−6645. (54) Kresse, G.; Furthmuller, J. Efficiency of Ab-Initio Total Energy Calculations for Metals and Semiconductors using a Alanewave Basis Set. Comput. Mater. Sci. 1996, 6, 15−50. (55) Kresse, G.; Hafner, J. Ab Initio Molecular Dynamics for Liquid Metals. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 47, 558−561. (56) Sun, J. P.; Dai, J.; Song, Y.; Wang, Y.; Yang, R. Affinity of the Interface between Hydroxyapatite (0001) and Titanium (0001) Surfaces: a First-Principles Investigation. ACS Appl. Mater. Interfaces 2014, 6, 20738−20751.

(57) Hauptmann, S.; Dufner, H.; Brickmann, J.; Kast, S. M.; Berry, R. S. Potential Energy Function for Apatites. Phys. Chem. Chem. Phys. 2003, 5, 635−639. (58) Wu, H.; Xu, D.; Yang, M.; Zhang, X. Surface Structure of Hydroxyapatite from Simulated Annealing Molecular Dynamics Simulations. Langmuir 2016, 32, 4643−4652. (59) Nair, A. K.; Gautieri, A.; Chang, S. W.; Buehler, M. J. Molecular Mechanics of Mineralized Collagen Fibrils in Bone. Nat. Commun. 2013, 4, 1724. (60) Bhowmik, R.; Katti, K. S.; Katti, D. Molecular Dynamics Simulation of Hydroxyapatite−Polyacrylic acid Interfaces. Polymer 2007, 48, 664−674. (61) Shen, J. W.; Wu, T.; Wang, Q.; Pan, H. H. Molecular Simulation of Protein Adsorption and Desorption on Hydroxyapatite Surfaces. Biomaterials 2008, 29 (5), 513−532. (62) Calderín, L.; Stott, M. J.; Rubio, A. Electronic and Crystallographic Structure of Apatites. Phys. Rev. B: Condens. Matter Mater. Phys. 2003, 67, 134106. (63) Plimpton, S. Fast Parallel Algorithms for Short-Range Molecular Dynamics. J. Comput. Phys. 1995, 117, 1−19. (64) Lu, Y. P.; Li, M. S.; Li, S. T.; Wang, Z. G.; Zhu, R. F. PlasmaSprayed Hydroxyapatite+Titania Composite Bond Coat for Hydroxyapatite Coating on Titanium Substrate. Biomaterials 2004, 25, 4393− 4403. (65) Momma, K.; Izumi, F. VESTA 3 for Three-Dimensional Visualization of Srystal, Volumetric and Morphology data. J. Appl. Crystallogr. 2011, 44, 1272−1276. (66) Sun, J. P.; Song, Y.; Wen, G. W.; Wang, Y.; Yang, R. Softening of Hydroxyapatite by Vacancies: a First Principles Investigation. Mater. Sci. Eng., C 2013, 33, 1109−1115. (67) Ben Osman, M.; Diallo-Garcia, S.; Herledan, V.; Brouri, D.; Yoshioka, T.; Kubo, J.; Millot, Y.; Costentin, G. Discrimination of Surface and Bulk Structure of Crystalline Hydroxyapatite Nanoparticles by NMR. J. Phys. Chem. C 2015, 119, 23008−23020. (68) Popov, M. N.; Spitaler, J.; Muhlbacher, M.; Walter, C.; Keckes, J.; Mitterer, C.; Draxl, C. TiO2(100)/Al2O3(0001) Interface A Firstprinciples study supported by experiment. Phys. Rev. B: Condens. Matter Mater. Phys. 2012, 86, 205309.

O


Strengthening Adhesion of the Hydroxyapatite and Titanium Interface

Recommend Documents