A. Preparation of master alloys

The Ti and Cu powders have been mechanical alloyed by a high energy ball milling machine. Figure 1(a) shows the X-ray diffraction patterns of the Ti₃₀Cu₇₀ powder mixtures before mechanical alloying (MA), the powders after MA for 20 h, and the spark plasma sintered Ti₃₀Cu₇₀ alloys. The Ti₃₀Cu₇₀ powder mixtures before MA consist of α-Ti and Cu phases. After MA, the Ti₃₀Cu₇₀ powders are still α-Ti and Cu phases, but some Ti peaks of Ti (102), (110), (103), (112), and (201) have disappeared. After SPS processing at 800 °C for 30 min, all the α-Ti and Cu phases are disappeared and Cu₃Ti₂ (ICDD No. 18–0459) and Cu₄Ti₃ (ICDD No. 18–0460) phases become dominate. Figures 1(b)–1(d) shows the SEM micrographs of the mechanical alloyed Ti₃₀Cu₇₀ powders and fractured surface of the spark plasma sintered alloys with EDX spectrum analysis. The powders are agglomerates with mean particle size of 10 μm in diameter [Fig. 1(b)]. The fractured surface of the spark plasma sintered Ti₃₀Cu₇₀ alloy shows a homogenous microstructure and grain size of about 1 μm with a fracture mode of intergranular cracking [Fig. 1(c)]. The relative density of the Ti₃₀Cu₇₀ alloy is measured to be 99.9%. The EDX spectrum of the area scanning in the sintered alloys shows elemental compositions of 32:68 (at.%), which corresponds to the initial designed 30:70 (at.%) in the powders.

FIG. 1.

X-ray diffraction patterns (a) and SEM morphologies of the mechanical alloyed Ti₃₀Cu₇₀ powders (b) and fractured surface of the spark plasma sintered Ti₃₀Cu₇₀ master alloy (c) with EDX spectrum exhibiting the elemental compositions in the master alloys (d).

B. Preparation and characterization of nanoporous α-Ti

Figure 2(a) shows the schematic illustration of the solid state dealloying process. The slice of TiCu master alloy was immersed in magnesium powders, cold pressed, and finally heat treated in a pressureless sintering furnace. The temperatures are in the range of 500–600 °C for 10–60 min in argon atmosphere. Figure 2(b) shows the fractured cross section of the dealloyed TiCu alloys in the Mg powders after heat-treatment at 600 °C for 30 min. It can be seen that the TiCu slice is imbedded into the Mg matrix. A high magnification SEM micrograph on the central core of the TiCu slice is shown in Fig. 2(c). It shows a uniform and smooth fractured surface. The EDX spectrum in the red square area indicates that the elements are Ti, Cu, and Mg. It exhibits that the composition of Cu has reduced to 2.38 at.% but the content of Mg has increased to 57.46 at.% [Fig. 2(d)]. It means that the Cu atoms in the TiCu alloys have been selectively removed and dissolved into the Mg powder matrix during the solid state heat-treatment dealloying process.

FIG. 2.

Schematic illustration of the solid state dealloying process (a), SEM micrographs of fractured surface of the Ti₃₀Cu₇₀/Mg (b) and Ti₃₀Cu₇₀ alloy after heat-treatment in Mg powders (c) with EDX spectrum showing the elemental compositions in the TiCu sample (d).

The phase compositions of the TiCu alloys during the solid state dealloying process are revealed in the XRD patterns of Fig. 3(a). After heat-treatment in Mg powders at 600 °C for 15 min, α-Ti, Ti₂Cu (ICDD No. 15–0717), Cu₂Mg (ICDD No. 65–5831), and Mg phases can be seen in the XRD pattern of TiCu alloys. It means that the Cu atoms have been diffused into the Mg powders and formed new phase of Cu₂Mg. The Cu₃Ti₂ and Cu₄Ti₃ phases have transformed into Ti₂Cu phase due to the release of Cu atoms. Finally, the TiCu slice was heat treated for 30 min at 600 °C for a complete reaction. After etching in the dilute nitric acid, it shows pure hcp structured α-Ti phase. But there is a weak and broad Ti₂Cu peak at 2θ of 39.5°. It is because of the presence of residual Cu elements in the foams. Figures 3(b)–3(d) shows the SEM micrographs of the fractured surfaces of the heat-treatment and etched porous α-Ti with inserted EDX of the porous structure. The nano-pores are distributed in the cross section uniformly in the edge and center of the 300 μm wide foam. The open-cell nanoporous structure is uniformly spread throughout the entire cross section. As measured on the SEM image, the mean value of ligament size is around 650 nm. The EDX shows that it has just Ti element without Cu and Mg compositions in the Ti nanofoam. Therefore, relative pure nanoporous α-Ti foam was fabricated successfully.

FIG. 3.

X-ray diffraction patterns of the Ti₃₀Cu₇₀ alloy after heat-treatment and etching (a) and SEM micrographs of fractured surfaces of the nanoporous α-Ti foam after heat-treatment and etching with inserted EDX spectrum (b–d).

Figure 4 shows the TEM micrographs of the Ti₃₀Cu₇₀ alloys after solid state dealloying in Mg powders with corresponding selected area diffraction (SAD) patterns. Figure 4(a) exhibits the TEM morphology of the Ti/Mg nanocomposite. The SAD pattern in Fig. 4(c) confirms the presence of α-Ti and Mg phases. The Cu atoms have diffused into the Mg after the solid state heat-treatment dealloying, as shown in Fig. 4(b). The SAD pattern on the dark skeleton in Fig. 4(b) reveals that it is Mg₂Cu single crystal phase [Fig. 4(d)]. The bright matrix in Fig. 4(b) is Mg phase [Fig. 4(e)]. However, the Mg₂Cu phase has not been detected in the XRD result in Fig. 3(a), where just the Cu₂Mg phase was found. It means that the diffusion of Cu atoms into the Mg formed Mg₂Cu phase firstly. With increasing of Cu concentration, it was transformed into Cu₂Mg phase. Figures 4(e) and 4(f) shows the TEM micrograph and SAD of the nanoporous α-Ti foam. It shows that the particle size of the Ti is about 20–50 nm and the pore size is in the range of 5–20 nm. The SAD pattern indicates that it is polycrystalline α-Ti phase. The particle sizes of the raw Ti and Cu powders are 325 mesh (45 μm). After MA, the TiCu powders show particles size about 10 μm [Fig. 1(b)]. After SPS processing, the TiCu alloys show grain size around 1.0 μm [Fig. 1(c)]. However, the Ti shows 20–50 nm particles size after the heat-treatment solid state dealloying. It is owing to the dealloying of Cu element from the Cu₃Ti₂ and Cu₄Ti₃ solid solutions where the Ti and Cu atoms have nanometer compounding after MA and SPS. It is also due to the formation of Ti/Mg nanocomposite after the solid state dealloying [Fig. 4(a)].

FIG. 4.

TEM micrographs of the Ti₃₀Cu₇₀ alloys after heat-treatment in Mg powders with corresponding SAD patterns showing the Ti/Mg nanocomposite (a, c), Mg₂Cu phase (b, d) and Mg phase (b, e); the obtained nanoporous α-Ti foam (f) with SAD pattern (g).

The porous structure and specific surface area of the Ti foams were characterized by N₂ absorption–desorption isotherms. Figures 5(a) and 5(b) shows the nitrogen adsorption/desorption isothermal plot of the nanoporous Ti foam, BJH adsorption dV/dD pore volume and pore integral distribution in the Ti foams. The adsorption isotherm for the resulting foam has hysteresis loops along with a relative pressure (p/p ₀) from 0.0 to 0.99. These isotherms are type IV, typical of mesoporous character materials. The BET surface area was calculated based on the isotherm results. The Ti nanofoam exhibits a BET surface area of (34.4 ± 0.8) m²/g. The total porosity is calculated to be about 60%. The porous properties of the Ti foams were further confirmed by pore size distribution analysis determined by the Barrett-Joyner-Hallender (BJH) method. The BJH desorption result showed that the smallest pores in the Ti foams are 2 nm. The values of pore width are from 2 nm up to 125 nm around. However, most of the pore sizes are distributed in the range of 2–50 nm, as shown in Fig. 5(b) of the BJH desorption dV/dD pore volume with a function of pore diameter. It is displayed in Fig. 5(b) that there is a sharp peak at pore size of 3.20 nm for sample of the nanoporous Ti foams. It means that the Ti foams are in nanoporous structures within mesoporous range (2–50 nm). Therefore, it is a nanoporous foam with a relative uniform pore size distribution.

FIG. 5.

Nitrogen adsorption/desorption isothermal plot of the nanoporous α-Ti foam (a), BJH desorption of pore differential and integral distributions with a function of pore diameter (b).

The dealloying method, present here, in the solid state powders is effective and powerful. The mechanism for the heat-treatment solid state dealloying is a solid state chemical reaction based on mixing enthalpy relationship of elements. The entropy between Mg and Cu is negative (−3 kJ/mol) while for the case of Ti and Mg is positive (16 kJ/mol). ^{Reference Takeuchi and Inoue29} The selective removal of components takes place by the attractive (negative entropy) or repulsive force (positive entropy) among the constituent elements. It is independent of the nobilities of constituent elements. During the heat-treatment, the Cu elements have been diffused into the Mg powder compacts and selectively dissolved. Nanocomposite of Ti/Mg is formed because of the attractive force for Cu–Mg but the repulsive force for Ti–Mg. The solid state chemical reaction process proceeds as that the solid solutions of Cu₃Ti₂ and Cu₄Ti₃ are reacted with Mg during heat-treatment and formed to be Ti₂Cu and Mg₂Cu phases, finally to be Cu₂Mg phase and Ti/Mg nanocomposite. In the end, the Ti/Mg nanocomposite was rapidly etched in dilute HNO_3. Ti has higher value (Ti²⁺, −1.628 V) of standard electrode potential than that of Mg (Mg²⁺, −2.375 V). Therefore, the Ti has better corrosion resistance compared with Mg and Cu₂Mg in the dilute HNO₃. At last, the Mg and Cu₂Mg phases have been selectively removed by the etching in the acid and the pure nanoporous α-Ti foams have been fabricated. It is noticed that the nanoporous structures were formed actually by the etching of the Ti/Mg nanocomposite in nitric acid. But during this step, no dynamic rearranging of the atoms occurs as in the dealloying, so it is merely an etching/dissolution process that chemically removes one of the two phases from the nanocomposite. During the solid state dealloying in Mg powders, there are dynamic rearranging of the atoms, in addition some constituent elements in the master alloy have been selectively removed and dissolved into the Mg powders, which is similar to an aqueous dealloying process. Therefore, the Mg powder method can be treated as a dealloying process to a certain extent.

C. Mechanisms of the solid state dealloying

The surface diffusion is the mechanism for dealloying in aqueous solution by the conventional electrochemical dealloying. ^{Reference Jin, Wang, Parida, Wang, Seo and Weissmuller2–Reference Erlebacher, Aziz, Karma, Dimitrov and Sieradzk6} For example, Erlebacher et al. ^{Reference Erlebacher, Aziz, Karma, Dimitrov and Sieradzk6,Reference Erlebacher and Seshadri30} revealed that the formation of nanoporous Au structures is caused by diffusion of Au atoms at the Au–Ag solid/electrolyte interface. The mechanism for the formation of the Ti nanoporous structures in Mg melt is surface and volume diffusion. ^{Reference Tsuda, Wada and Kato20} The pore-forming mechanism in the solid state dealloying using powders here need investigations.

To study the pore-forming mechanism, disk-shaped Ti₃₀Cu₇₀ samples with diameter of 10 mm and thickness of 1.2 mm were used. One disk-shaped Ti₃₀Cu₇₀ sample was dealloyed by heat-treatment at 600 °C for holding time of 10 min, and then chemical dealloyed in dilute nitric acid. It can show an uncompleted dealloying state. There are three parts in the sample: dense Ti₃₀Cu₇₀ (inner core), porous TiCu (Ti₂Cu, Ti₃Cu) intermediate part, and nanoporous α-Ti part (out surface). The total thickness of the α-Ti and TiCu layers can be used to study the depth of Mg atoms diffusion into the Ti₃₀Cu₇₀ alloy. Therefore, the average thickness of the porous α-Ti (l _α-Ti) and porous TiCu (l _TiCu) can be calculated with the following equation, ^{Reference Meguro, O and Kajihara31}

(1) $${l_{\rm{x}}} = {{{A_{\rm{x}}}} \over {{W_{\rm{x}}}}}\quad ,$$

where, l _x is the thickness of Mg diffusion into the TiCu alloy, A _x means the total area of the compound x layer on the cross section, W _x refers to the total length parallel to the compound x interface. Figure 6 shows the plots of logarithm of layer thickness [lg(l _x)] with a function of logarithm of holding time t [lg(t/t ₀)] at various heat-treatment temperatures of 510, 540, 570, and 600 °C. The total thickness, the thickness of porous TiCu and α-Ti are noted with triangle, circle, and square shape, respectively. As can be seen, the thickness lg(l _x) increase linearly as a function of lg(t/t ₀) at various temperatures. It indicates that all the logarithm of the Mg diffusion thickness increase linearly with the logarithm of holding time. The relationship between the thickness l _x and the holding time t obeys the following equation,

(2) $${l_{\rm{x}}} = k{\left( {{t \over {{t_0}}}} \right)^n}$$

where, t ₀ is the unit time, 1 s, that is adopted to make the argument t/t ₀ of the powder function dimensionless, k _x is the rate constant, n is the reaction exponent and it is the slope in the logarithmic formula. This formula is usually used for analyzing reaction diffusion interfaces. ^{Reference Meguro, O and Kajihara31} The values of n and k _x are determined by the least-square method using the data in Fig. 6 and summarized in Table I. The exponent n is affected by the rate-controlling process. When the layer growth is controlled by volume diffusion, n is equal to 0.5. The index n is equal to 0.25 when only grain boundary diffusion occurred. ^{Reference Corcoran, King, Lanerolle and Kim32} The value of n is equal to 1 when growth of the phase is controlled by interface reaction. Figure 7 shows the relationship between reaction exponent n versus heat-treatment temperature of the disk-shaped sample for α-Ti phase, TiCu phase and total layer. The n value for growth of the total porous phase in Fig. 7(a) is determined by the α-Ti layer and TiCu layer. In Fig. 7(b), the value of n for the porous TiCu phase decreases with increasing of the dealloying temperature. The value of n is equal to 0.56 at dealloying temperatures of 510 and 540 °C, but around 0.5 at 570 and 600 °C. Thus, the growth of the porous TiCu phase is controlled by interface reaction at 510 and 540 °C, and by volume diffusion at relative higher temperatures of 570 and 600 °C. In Fig. 7(c), the n of the porous α-Ti phase is greater than 0.5, close to 0.6. So, the growth of the α-Ti layer is mainly controlled by the interface reaction. Consequently, the pore-forming process in the solid state dealloying using powders is controlled by both of interface reaction and volume diffusion. The dealloying within metallic powders is a water-free process. During the heat-treatment, diffusion is activated for mass transport. The constituent of Cu in the TiCu alloy was dissolved into the Mg powder compacts thanks to the mechanism of interface reaction and volume diffusion.

FIG. 6.

Plots of thicknesses of compound layers lg(l _x) with a function of lg(t/t ₀) at various heat-treatment temperatures of 510 °C (a), 540 °C (b), 570 °C (c) and 600 °C (d).

TABLE I.

Values of rate constant k and exponent n of disk-shaped samples at 510, 540, 570, and 600 °C.

FIG. 7.

The relationship between exponent n versus heat-treatment temperature of the disk-shaped sample for total layer (a), porous TiCu (b) and porous α-Ti phase (c).

To better understand the mechanism for the formation of the porous layers, the activation energy for the growth of the porous layers was analyzed. Figure 8 shows the relationship between lnk of the total porous phase, the porous TiCu phase and porous α-Ti phase versus reciprocal heat-treatment temperature in Kelvin rating for the growth of each layers (1/T). As can be seen from Fig. 8, the plotted points are well located on straight lines. Thus the activation energy can be obtained by fitting an Arrhenius-type equation,

(3) $$k = {k_0}\exp \left( { - {E \over {RT}}} \right)\quad ,$$

where k ₀ is the pre-exponential factor, R is the gas constant, and E is the activation energy. The activation energy of the total porous phase, the porous TiCu and the porous α-Ti are calculated to be 117.80, 127.01, and 104.99 kJ/mol, respectively. These values are higher than those in the Mg melts (35–45 kJ/mol). ^{Reference Tsuda, Wada and Kato20} It indicates that the dealloying in solid state powders need much higher activation energies. Taguchi et al. ^{Reference Taguchi, Iijima and Hirono33} reported that the activation energies are 79–125 kJ/mol for the growth of reaction layers in the reaction diffusion of a solid TiCu couple. Our values are very close to these in the solid TiCu couple. It means that the formation of α-Ti during the heat-treatment dealloying in Mg powders can be attributed to a solid-state interdiffusion phenomenon based on mixing enthalpy relationship among constituent elements. Finally, the nanoporous structured α-Ti are formed by subsequent etching in an acidic solution.

FIG. 8.

Plot of rate constant k versus reciprocal of heat-treatment temperature T in Kelvin rating.

D. Nanomechanical properties

The ligament coarsening of the nanoporous α-Ti foams with a function of the heat-treatment temperatures was investigated and illustrated in Fig. 9(a). It can be clearly seen that the ligament sizes grow up with increasing of the heat-treatment temperature at a constant holding time of 30 min. It coarsely obeys a linear increment phenomena. The mean values of ligament size are 492, 523, 602, and 650 nm at temperatures of 510 °C, 540 °C, 570 °C, and 600 °C, respectively. Figure 9(b) shows the nanoindentation loading/unloading curves of the nanoporous Ti foams. The curves in blue color are for the 600 °C heat treated foams with ligament size of 650 nm. The curves in black color are for the 540 °C processed foams (523 nm). The overlapping loading sections of each data set demonstrate the excellent reproducibility of the measurements. The contact depth which represents plastic deformation intensity, increased with ligament size. Based on the nanoindentation results, the mean values of nanohardness and elastic modulus are calculated. Figure 9(c) demonstrates the relationship between hardness, elastic modulus and ligament size of the nanoporous α-Ti foams. As shown, the hardness and elastic modulus are both reduced with the ligament size increasing. The 492 nm ligament sized sample prepared at 510 °C shows hardness of 259 MPa and elastic modulus of 4.15 GPa. However, the 650 nm sample (600 °C processed) displays hardness of 178 MPa and elastic modulus of 3.04 GPa. The hardness and elastic modulus both follow linear decay fit as a function of ligament size increasing. These mechanical values are higher than those obtained by Mg melt method. ^{Reference Wada, Setyawan, Yubuta and Kato21} Hakamada and Mabuchi reported that the hardness is equal to yield strength in the nanoporous foams. ^{Reference Hakamada and Mabuchi34} In the current study, we found that the hardness and elastic modulus of nanoporous α-Ti decrease as the ligament size increases following linear decay. It means that the yield strengths of the nanoporous α-Ti foams follow the same trend.

FIG. 9.

The ligament size with a function of heat-treatment temperature (a). Nanoindentation loading/unloading curves (b) and the relationship between hardness, Elastic modulus with ligament size of the nanoporous α-Ti foams (c).

This study provides a new dealloying route by heat-treatment in solid state powders. The heat-treatment temperatures (500–600 °C) are lower than the melting point of Mg (650 °C) and much lower than the Mg melt method (670–870 °C). ^{Reference Wada, Yubuta, Inoue and Kato16,Reference Tsuda, Wada and Kato20,Reference Wada, Setyawan, Yubuta and Kato21} The holding time in the solid state dealloying (10–60 min) is longer than the method of Mg melts (0.1–20 min). ^{Reference Wada, Yubuta, Inoue and Kato16,Reference Tsuda, Wada and Kato20,Reference Wada, Setyawan, Yubuta and Kato21} Recently, Wada et al. reported the preparation of nanoporous FeCr alloy via a solid-state interfacial dealloying reaction. ^{Reference Wada, Yubuta and Kato35} They used Mg plate (1 × 12 × 12 mm) as solid dealloying agent, where the plate was fixed tightly with precursor of (FeCr)Ni using stainless-steel bolts and heat treated at 440–480 °C for 6–72 h. ^{Reference Wada, Yubuta and Kato35} In our study, Mg powders are used as solid dealloying agent, where the precursor was immersed in the powders and heated treated at 500–600 °C for 10–60 min. Both of them using Mg plate and powders belong to the solid state dealloying method. But the powder method is relatively simple and easy to implement since it’s no need for fixing with the precursor tightly. Its dealloying time is much shorter (10–60 min) than that of the Mg plate method (6–72 h). The method of solid state dealloying using powders need just a normal pressureless sintering furnace. The mean particle sizes of the Mg powders are not fine about 45 μm without hazard to flame. This method allows the shape of master alloy to be largely maintained during dealloying. Therefore, it is a simple and convenient method. It can be used alternatively besides the Mg melt and plate method. It could be used for the fabrication of nanoporous metals those have lower standard electrode potentials and are not suitable with electro chemical dealloying method, for example, nanoporous Ti, Mo, Nd, W, V, Ta, Zr, Hf, and so on. These metals all have very high melting point and are very suitable for the SPS processing and solid state dealloying. People can choose other elemental powders to instead of the Mg powders. It depends on the mixing entropy of elements in the alloy compounds and the entropy of the alloy constitution with the dealloying element. From an application perspective, Ti is valuable because it has the highest specific strength of any elemental metals and highly corrosion resistant, as well as is considered to be biomedical inert. The nanopores may be able to be loaded with stored bioactive compounds such as proteins or antibacterial drugs for slow release such that nanoporous networks could act as reservoirs for these bioactive compounds and prevent infections and inflammations that lead to septic loosening of implants during healing. ^{Reference Panagiotopoulos, Moreira Jorge, Rebai, Georgarakis, Botta and Yavari17,Reference Necula, Apachitei, Fratila-Apachitei, van Langelaan and Duszczyk36} In addition, fabrication of nanoporous Ti foams would have implications in many fields. The Ti nanoporous metals can serve as precursors to other nanoporous ceramic materials, such as, titanium oxides, nitrides, and carbides. The nanoporous Ti foams may find very broad functional applications in the near future, especially in biomedical field as a new structured implant with drug release properties.