1. Introduction
Titanium alloys are widely used in various branches of mechanical engineering and for the manufacture of medical instruments and implants [1,2]. The advantages of titanium alloys include desirable specific mechanical properties, good fatigue life and ductility, and high corrosion resistance and biocompatibility [3]. At the same time, low tribological properties are the main disadvantage, which significantly limit the use of titanium alloys as structural materials for critical elements of machines and mechanisms [4,5,6]. Surface modification allows for overcoming this disadvantage while preserving of other useful properties [7,8,9,10,11,12,13,14,15]. Surface modification allows simultaneous improvement of several characteristics, such as surface hardness and fatigue properties [16]. One of the actively developed methods for this is the formation of the surface intermetallic layers based on the Ti-Al [17,18], Ti-Ni [19], Ti-Si [20], and Ti-Al-Cr [21] systems.
Intermetallic compounds of the Ti-Al system have good physical and mechanical properties [22], are resistant against oxidation up to the temperatures of 800–1000 °C [23,24,25], and possess high affinity with titanium, which provides high adhesive strength of the layers. In the Ti-Al system, intermetallic compounds with different stoichiometry can be formed, but usually there are three stable intermetallides: Ti3Al (with an ordered hcp-based structure), TiAl (with a tetragonal L10 lattice, c/a = 1.02), and TiAl3 (with tetragonal D022 lattice, c/a = 2.22) [2]. Also, alloying of titanium alloy with aluminum can lead to the formation of a substitutional solid solution for the hexagonal α-phase, and can change the lattice parameter of the alloy and, consequently, the c/a ratio. Alloying of Ti-alloy with aluminum up to the 27 at. % increases the c/a ratio to 1.61 (the initial value is 1.588, a = 2.950, c = 4.686) [26]. This is close to the optimal c/a ratio of 1.63, which provides the advanced tribological characteristics of titanium–aluminum alloys [27,28]. For the intermetallic compounds, γ-TiAl is a promising one due to the combination of high hardness, oxidation resistance, and satisfactory plasticity relative to the non-equiatomic phases [29]. Zhang et al. [30] showed that a γ-TiAl coating deposited on a TC4 titanium alloy had high wear resistance in dry friction against a steel ball.
The first approaches to obtaining the TiAl-based surface layers in titanium alloys were based on a vacuum annealing of titanium and aluminum coatings [31,32,33,34,35]. In this case, a reactive diffusion is the main mechanism for the formation of intermetallic layers, and it occurs both in the solid phase (up to the melting temperature of aluminum) [36,37,38,39,40] and in the liquid phase [41,42,43,44]. During the vacuum annealing, the TiAl3 phase is formed first because it has more negative Gibbs free energy relative to the TiAl and Ti3Al intermetallides, and the formation of TiAl2 and Ti2Al5 phases requires the TiAl one as an intermediate phase [45]. Xu et al. [31] have observed that the TiAl3 phase is the only one formed during the solid-state annealing of Ti-Al multilayer diffusion couples at temperatures of 520–650 °C. However, during the subsequent annealing of the composite Ti + TiAl3 layers at temperatures of 700 and 800 °C, the TiAl phase was formed at the expense of the reaction between Ti and TiAl3. Other studies [46,47,48] also confirmed the solid-state transformation of the TiAl3 phase into the TiAl one, when the aluminum-enriched intermetallide interacts with Ti. Even though higher temperature assists the desirable TiAl phase, the temperature for the formation of TiAl intermetallic coatings is recommended not to be higher than 500 °C to avoid the formation of Kirkendall voids [25]. Using a synchrotron radiation, Rogachev et al. [49,50] have shown the sequence of phase transformation during the heating of films with different periods of titanium and aluminum layers. The reaction of intermetallic formation begins at 600 K (327 °C), and at T = 760 K (487 °C) the first peaks of the TiAl phase appear. The formation of the Ti3Al intermetallic compound starts only at temperatures above 900 K (627 °C). In rather thick films with a large period of layer thickness, the TiAl3 phase was first formed, followed by the formation of the TiAl phase upon further heating. In thin Al and Ti films (about 8 nm per layer), the reaction proceeded instantly with the formation of the TiAl phase from α-Ti at a temperature of 700 K (427 °C). So, for lower (nanoscale) thickness of the titanium and aluminum layers, the TiAl intermetallic compound is formed without the intermediate phase due to the insufficient concentration of aluminum for the formation of TiAl3 intermetallide and lower energy of the phase transformation in nanoscale laminated structures compared to the bulk materials [51].
The kinetics of phase transformations in the TiAl system, the transformation rate, and the characteristic temperatures typical of classical vacuum annealing can be varied by the application of external fields [52,53]. In several papers [54,55,56], gradient intermetallic layers of Al2O3-Al3Ti-TiAl-TiAl3 were successfully obtained by a deposition of the magnetron Al coating and its subsequent oxidation in glow discharge plasma at 680 °C for 3 h. Romankov et al. [57] treated the Ti–Al coating with argon plasma pulses with different energies up to 185 J/cm2. Treatment with pulse energy higher than 135 J/cm2 caused a melting of the surface layer, and the liquid-phase interaction of Ti and Al provided the formation of the TiAl3 phase. No other intermetallic phases were found even when the irradiation parameters changed. To date, the detailed studies on the mechanism of formation of intermetallic layers during treatment in gas discharge plasma have not been completed. This scientific problem of the mechanisms of formation of intermetallic layers under a non-self-sustaining low-pressure arc discharge in plasma is not still solved.
Most of the mentioned above research papers were focused on the development of the methods of the surface modification of the titanium alloys, or the approaches to form desirable intermetallic phases in Ti-Al system. For the initial characterization of the microstructure, metallography, scanning electron microscopy, and X-ray diffraction methods have been successfully used [7,8,9,10,13,17,30,31,33,35,36,42,51]. However, detailed study of the microstructure in gradient materials requires the application of analytical methods with high resolution, such as transmission electron microscopy, which was successfully used previously [14,22,32,34,42,46,54,58]. The intermetallic layers fabricated by the non-self-sustaining low-pressure arc discharge in plasma were not studied in detail. The purpose of this article is to characterize the microstructure and phase composition of the coating and modified surface layers obtained in specimens of a Ti-6Al-4V alloy subjected to the complex ion-plasma treatment, including the deposition of the Al coating and subsequent processing in the low-pressure non-self-sustained arc discharge plasma, by the method of transmission electron microscopy.
2. Materials and Methods
Titanium alloy Ti-6Al-4V (86.5–91.2) Ti, (5.3–6.8) Al, (3.5–5.3) V, <0.3 Fe, <0.3 Zr, <0.15 Si, <0.05 N, <0.2 O (in mass. %) was used as a base material for processing (in as-received state after hot rolling in two-phase region, followed by an annealing at 760 °C). The alloy was supplied in the form of rolled bars with a diameter of 15 mm. The bars were cut into cylinders with a diameter of 15 mm and a thickness of 4 mm. The surfaces of the samples were polished using an abrasive paper to a roughness of Ra = 1.6 µm. Before treatment, an ion cleaning of the specimens’ surfaces with argon and aluminum plasmas was performed. The process of a complex ion-plasma treatment (CIPT) consisted of two stages. At the first stage, a coating of pure aluminum was deposited on the surface of the samples made of the Ti-6Al-4V alloy. The coating was obtained by a vacuum-arc plasma-assisted method in an argon atmosphere. The substrate temperature during the deposition did not exceed 400 °C. As a result of the first stage of the CIPT, an Al coating with a thickness of ~2.5 µm was deposited. At the second stage, the specimen with Al coating was treated in high-density plasma, which is generated by a plasma source with a filament cathode based on a non-self-sustained low-pressure arc discharge (“PINK” cathode, Institute of High Current Electronics, Siberian Branch of the Russian Academy of Sciences, Tomsk, Russia). The process temperature was (500 ± 15) °C, the duration—1 h, the pressure in the chamber—0.5–0.7 Pa. During the second stage, an interdiffusion of the components of the base material (titanium alloy) and the coating (Al) occurs and phase transformations are realized in the coating and the surface layers of the titanium alloy, which are stimulated both by the temperature (500 °C) and ion-plasma treatment. The CIPT was performed in argon atmosphere, and the specimens were placed on a turntable table in the center of the vacuum chamber and rotated relative to the plasma sources at a speed of 2 rpm. Both stages of the CIPT were performed in a single processing (vacuum) cycle.
The samples were cut perpendicular to the modified layer, mechanically ground using abrasive papers, and polished using a SiO-based suspension. To reveal the microstructure, the samples were subjected to chemical etching in a solution: 16 mL of hydrofluoric acid, 16 mL of nitric acid, 68 mL of glycerol.
Microstructural studies were carried out by scanning and transmission electron microscopy using Tescan Mira (SEM, Tescan, Brno, Czech Republic) and Jeol JEM-2100 (TEM, JEOL, Tokyo, Japan) electron microscopes equipped with an Energy-Dispersive Spectroscopy (EDS) unit. To study the structure by TEM, foils were cut from the samples perpendicular to the modified layer, mechanically ground, and thinned with a Jeol EM 09100IS ion slicer (JEOL, Japan). Indexing of selected area electron diffraction patterns was performed as detailed elsewhere [59].
3. Results and Discussion
Figure 1 shows characteristic SEM images for the original titanium alloy Ti-6Al-4V without any surface treatment. In the as-received state, large primary α-Ti grains of 2–10 μm in diameter are seen (the mean grain size is 6.9 μm). Coarse grains of β-Ti lie primarily near the grain boundaries of the α-Ti, and recrystallized β-Ti particles are seen inside of the primary α-Ti grains. Bright areas (β-Ti) on the SEM image (Figure 1b), obtained by detection of the backscattered electrons, clearly show the distribution and morphology of the β-phase in the as-received microstructure. Therefore, two-phase microstructure is observed, which is typical of annealed (α + β) titanium alloys [60].
Figure 2 shows the common cross-sectional view and some magnified fragments of the microstructure of the specimen after the deposition of the aluminum coating (first stage of the CIPT).
The aluminum coating has a complex structure and consists of two characteristic regions. Next to the interface between the titanium alloy and the coating, a nanostructural region is observed (Figure 2a,b). Its thickness is about 300 nm. The selected area electron diffraction (SAED) pattern, corresponding to this area, is represented mainly by rings with the interplanar distances belonging to the pure aluminum. These rings consist of many point reflections, which are not diffused either in the radial nor in the azimuthal direction. Both these facts indicate that this region consists of many crystallites (grains) of pure aluminum with high-angle misorientations between them. Dark-field images, obtained in reflections (1) and (2) (Figure 2c), show that grain size can reach several tens of nanometers in size, but the mean size is 28 ± 9 nm. The formation of the nanocrystalline structure (with the size of the elements lower than 100 nm) is due to the presence of many defects on the surface as a result of surface activation during the ion cleaning [61]. These defects served as crystallization centers, which contribute to the formation of many nuclei of the grains. Simultaneously, ion cleaning stimulates diffusion of the main components of the titanium alloy into the nanocrystalline coating. Despite the fact that the treatment temperature is high enough for recrystallization to begin in pure aluminum [62], this nanocrystalline structure is stable. This question needs further study in detail, but we can assume that grain boundary segregations can suppress grain growth near the “titanium alloy-coating” interface [63]. It could be Ti and V from the base material, or aluminum oxides from the cathode surface that are sprayed in the beginning of the coating deposition. H. Garbacz et al. [54] reported that a similar nanocrystalline zone formed under the magnetron sputtering of Al coating on the titanium alloy. They reported that the impurities in the chamber can be adsorbed and segregated on the surface of the growing grains, inhibiting their coarsening.
The region of the coating next to the outer surface of the specimen is characterized by an ultrafine-grained structure (typical grain size range is 100 nm–1 μm). The grain size in this region is about 300 nm, but some grains can reach the 1 μm size (Figure 2a, the upper part). These grains produce bright point reflections in SAED patterns (Figure 2c). In the middle part of the coating, thin interlayers of the nanocrystalline aluminum are seen between the ultrafine grains. When one obtains a dark-field image in combined reflection of ultrafine-grained and nanocrystalline aluminum (reflections 1 and 2 in Figure 2c), the appearance of the nanocrystalline grains between the ultrafine-grained ones are clearly seen (Figure 2d). however, moving to the outer surface of the specimen, they are less frequent. Therefore, during the deposition of the outer part of the coating, a recrystallization is activated. This is possible because the coating temperature (400 °C) is enough for recrystallization in pure aluminum [62]. On the other hand, the concentrations of titanium and vanadium is low (or they are absent) in this part of the coating to suppress grain boundary migration [63].
The above analysis has shown that the deposition of the Al coating on the Ti-6Al-4V alloy provides the formation of the layered coating with the gradient microstructure—mainly nanocrystalline (<100 nm) near the interface between the coating and the Ti-based matrix, and ultrafine-grained (<1 μm) in the outer part of the coating.
The surface of the treated specimen has a wavy relief due to the large amount of the droplet fraction during the coating by Al (Figure 2a). The interface between the titanium alloy and Al coating is not straight either. The cleaning of the specimens’ surfaces before the coating provides its etching. As the grains of titanium alloy are differently oriented relative to the treated surface, they are etched to different depths and form the steps in the intergranular boundaries (Figure 2a). The α-stabilized region of ≈5 µm thick is observed in the near-surface layer in the titanium-based part of the specimen. This region can be distinguished from the base material by the dissolution of the small recrystallized β-Ti particles inside of α-Ti grains and rather large β-Ti grains near the grain boundaries (lower part of the Figure 2a). Aluminum diffuses into titanium alloy at the initial stage of the deposition, when the surface is activated by the ion cleaning. During the process, the temperature of the specimen can reach 400 °C, which is sufficient for diffusion of aluminum along grain boundaries and in defects of the crystal lattice. The ion cleaning accelerates the diffusion process [61].
Figure 3 shows TEM images of the Ti-6Al-4V alloy matrix at the depth where aluminum diffusion does not affect the microstructure of the alloy. Typical (α + β) grain structure is observed (Figure 3a,b,e,f), which consists of the globular grains bounded by thin equilibrium boundaries, randomly distributed dislocations in α-Ti grains (Figure 3a). The selected area diffraction patterns and dark-field analysis confirm the presence of both α (Figure 3a–d) and β phases (Figure 3e,f).
After the CIPT, the coating preserved its layered morphology, but its phase composition and microstructure changed drastically. The common cross-sectional view of the microstructure in the specimen is shown in Figure 4. The magnified sections of the Figure 4a are presented in Figure 4b–e. They demonstrate the typical microstructure in the outer and inner parts of the coating (Figure 4b,c respectively), of the modified intermetallic layers in base material near the interface “coating/substrate” (Figure 4d), and α-Ti layer in the base material near the intermetallic modified layer (Figure 4e).
The former aluminum coating has been saved after the treatment, and it consists of two characteristic zones (the upper part of the image in Figure 4). The first zone is located deep in the coating, near the “coating/substrate” interface, and it has a thickness of 200 nm (Figure 4a,c). Apparently, this zone is inherited from the nanocrystalline part of the aluminum coating and represents an extremely fine microstructure (Figure 4c). The SAED pattern corresponding to this region gives diffuse rings that have strong radial diffusions (Figure 5a,b). The separate reflections are not identified on each ring, and the azimuthal diffusion of the reflections is present as well. The ratios of the interplanar distances (dhkl) in Figure 5a show that the phase possesses a FCC crystal lattice, and the dhkl-values are close to those for aluminum (Figure 5b). However, dhkl-values in Figure 5c are close to those for α-Ti (Table 1). Given the nature of the radial diffusions of the electron diffraction patterns [58], it can be assumed that there can be a solid-solution hardening of the nanocrystalline aluminum grains with different (and gradient in single crystallite) concentrations of the solutes (Ti,V) and α-Ti phase in grain-boundary regions. The complex spotty contrast on the TEM images, captured from this thin layer, and the random distribution of the reflections in the rings, all testify to the nanocrystalline microstructure. Wide diffused boundaries between crystallites and continuous (low-angle) misorientations are typical of such structures, possessing high distortions or high density of the defects both in the crystal lattice and the boundaries [64,65,66]. The areas of similar morphology are also seen in the upper part of the coating, and lie as the interlayers between rather coarse grains of an Al3Ti intermetallic phase. These interlayers show the ring diffraction patterns, but the grains of the Al3Ti phase between them show point reflections (Figure 5c).
The second characteristic zone is the main one in the former aluminum coating, is located closer to the outer specimen surface, and is characterized by a grain structure without clearly defined grain boundaries in bright-field images (Figure 4b). The interpretation of the SAED patterns obtained from this region show the presence of the TiAl3 phase only (Figure 6a,b). Dark-field images obtained in reflections of the TiAl3 phase are shown in Figure 6c,d. The relatively large grains are seen there.
According to the EDS analysis (Figure 7, Table 2), the upper region with the coarse grains has a chemical composition close to the stoichiometric phase TiAl3. This directly confirms TEM data about the composition of this part of the coating. So, during the CIPT, the transformation of the ultrafine-grained pure aluminum into TiAl3 intermetallic phase occurred, and all aluminum reacted with titanium: 3Al + Ti → Al3Ti. This result is in line with the data of other authors about the solid-state reaction of titanium and aluminum [45,67]. Simultaneously with the phase transformation, grain growth is activated, causing the formation of the rather coarse structure in this intermetallic region.
The intriguing result is that no phase transformation or recrystallization occur in the nanocrystalline region near the “titanium alloy-coating” interface or in former nanocrystalline interlayers between large Al3Ti grains. These interlayers and the main nanocrystalline region consist of 97–98% aluminum and only about 2–3% titanium and vanadium (Figure 7, Table 2). Such concentrations of titanium are not sufficient for the formation of a TiAl-based intermetallic phase, but a mixture of the Al + TiAl3 is possible. TEM analysis shows the preservation of the Al-based and α-Ti phases. The concentration of 2–3% titanium exceeds the solubility of Ti in Al [2,67]. So, the main part of the titanium can be concentrated in or near the boundaries of the nanocrystallites saturated with lower concentration of Ti.
Grain boundaries in nanocrystalline structures can provide ultra-fast diffusion of the solutes [68], working as diffusion channels for fast transport of Ti and V to the upper part of the coating, and for Al atoms to the lower one, under the CIPT. We suppose that the small size of the nanocrystallites or grain boundary regions (where concentrations of Ti can be high) is insufficient for the nucleation of the intermetallic phase, but we have not found any available information in open sources about the critical size of the nuclei in the TiAl system under transformation. As the solid-phase reaction of titanium and aluminum did not occur in this layer (as in the main volume of the coating), one can also expect the formation of thin films or segregations at the boundaries of nanocrystalline grains, which prevent diffusion and hinder grain growth. In particular, the formation of an aluminum oxide film may occur due to the fact that, at the initial moment of deposition, an oxide film is always present on the aluminum cathode. Grain boundary segregations of α-Ti and solid-solution hardening can also suppress grain growth [63]. Therefore, the nanoscale elements of the microstructure obtained during the aluminum coating deposition in the first step of the treatment are preserved during further CIPT processing, and work as a dense net of diffusional channels for the transport of the Ti to the aluminum part of the specimen, and of Al to the Ti-based part of the specimen.
During the CIPT, phase transformations occur in the Ti-based part of the specimen, and a layered microstructure forms there. The common view of the Ti-based part of the specimen is seen on the lower part of the image on Figure 4a,d,e. Figure 8 shows the TEM images of the structure with more details. The whole affected surface can be divided into several characteristic zones.
The first zone is located next to the “coating/substrate” interface, and has an average thickness of 1.5 µm. The structure of this layer is represented by fine grains with predominantly equiaxed morphology (Figure 8a,b). The SAED patterns are presented by the numerous point reflections, most of them belonging to the TiAl3 phase. However, there are several point reflections with interplanar spacing which is characteristic of the TiAl phase as well. The reflections are not rounded into the diffraction rings, but are distributed rather homogeneously along them. Thus, the misoriented grain structure forms in this region during the CIPT. All reflections are diffused, both in radial and azimuthal directions (Figure 8c). Therefore, a low-angle misorientation and slightly variable chemical composition are peculiar for the grains of the TiAl3 phase. The dark-field image obtained in the indicated area on the Figure 8c reflections (including all possible phases) shows that the grain structure is not uniform throughout the layer; there are ultrafine grains (hundreds of nm) and nano-sized fragments. The average grain size of the TiAl3 phase in this layer (in the modified surface layer of Ti alloy) is much smaller than that of the same phase in the coating (Figure 6c,d and Figure 8d). This is apparently due to the difference in the recrystallization temperatures for aluminum and titanium alloy [2,63].
The second zone is not wide and is located under the main layer of TiAl3 intermetallic compound (Figure 9). It consists of nonequiaxed grains of 250–300 nm in size. The SAED analysis shows that these grains belong to the TiAl intermetallic phase (Figure 9, areas 1 and 2). The thin layer of the Ti3Al phase was also found (Figure 9, areas 3 and 4). However, there are only single grains between the main intermetallic surface layer (TiAl3 + TiAl) and matrix titanium alloys, where phase transformation has not been activated. H. Garbacz and colleagues [54] treated the titanium alloy under glow discharge conditions after the magnetron sputtering of the Al coating on the surface of the specimen. They reported a similar sequence of the phases, which is shown in Figure 9. In their research, however, the pre-coating was completely diffused.
The reactive diffusion in solid state is responsible for the formation of the intermetallic layers described above (both in the coating and in base material) [36,37,38,39,40]. Vacuum annealing without the application of the ion-plasma treatment assists the TiAl3 phase formation because of the lower Gibbs free energy of this phase relative to the TiAl and Ti3Al phases [45]. The temperature of the CIPT (500 °C) also assists TiAl3 phase formation [31]. The thin layers of the TiAl and Ti3Al phases are observed only near the base material, where the concentration of the Ti is high. Therefore, the solid-state transformation of the TiAl3 phase into TiAl and titanium-rich Ti3Al phases [46,47,48] is possible during the CIPT process.
TEM EDS analysis confirms the presence and distribution of the phases identified during TEM diffraction study (Table 1). The increased aluminum content is found at a depth of up to 2.5 μm from the TiAl intertmetallic layer (Table 1). Diffusion of Al into the depth of the base alloy assists the dissolution of the grains of β-Ti phase (Figure 4a,e). The microstructure of the titanium alloy at depth, where aluminum does not influence the composition of the alloy, is similar to that on Figure 3.
According to TEM analysis, the following sequence of the phases forms as a result of the CIPT treatment:
[TiAl3 → TiAl3 + nc-(Al(Ti) + α-Ti)] → [TiAl3 → TiAl3 + TiAl → TiAl → Ti3Al → α-Ti alloy → (α + β)-Ti alloy].
The square brackets combine the phases in the coating and in the substrate, respectively. The important results of the above analysis are:
-. The CIPT processing allows formation of the intermetallic coatings and surface-modified layers in the titanium alloy. The main phase is the TiAl3 intermetallide, which correlates with previous research [31,45,46,47,48]. In Ti alloy, a desirable phase TiAl and titanium-enriched Ti3Al phase are also present. Therefore, the variation of the treatment regime is needed to reduce the content of the TiAl3 phase and to increase the fraction of the TiAl one.
-. Despite the fact that TiAl3 is the main intermetallic phase both in the coating and in the base material, it possesses different morphologies in either. The sequential variation of the microstructure (grain size and distribution) was observed when one moves from the surface of the specimen to its depth (IM—intermetallics, Al—Al-based solid solution, Ti—Ti-based solid solution):
Fine grains (IM) → fine grains (IM) + nanocrystallites (nc-(Al(Ti) + α-Ti)) → nanocrystallites (nc-(Al(Ti) + α-Ti)) → interface “coating/substrate” → ultrafine grains (IM) → fine grains (Ti)
During the phase transformation under the CIPT, the temperature-assisted (500 °C) grain growth of the intermetallic phase TiAl3 occurs in Al-based and Ti-based parts with principally different melting and recrystallization temperatures [2]. The grain size in the Al-based coating is obviously higher than that of the Ti-based one.
-. Phase transformation and recrystallization are not realized in nanocrystalline regions of the aluminum coating near the “titanium alloy/coating” interface. Elemental EDS analysis and TEM diffraction analysis were completed, and confirmed the preservation of the nanoscale-sized fragments and high concentrations of Al in this region. A high fraction of the grain boundaries in nanocrystalline regions can favor the diffusion of the elements in both directions and, therefore, stimulates phase transformation in coarser grains situated under and over the nanocrystalline layer. These diffusion flows can suppress grain boundary migration and support the stability of the nanocrystalline structure.
4. Conclusions
The complex processing of the two-phase (α + β) titanium alloy Ti-6Al-4V was performed, which included the deposition of the Al coating and subsequent surface treatment in the low-pressure non-self-sustained arc discharge plasma. The microstructure and phase composition of the coating and modified layers of base material were studied using TEM methods. The main results of the paper can be summarized as follows:
The deposition of Al on the Ti-6Al-4V alloy is accompanied by the formation of a layered aluminum coating with a gradient microstructure—nanocrystalline near the “coating/substrate” interface and fine-grained in the outer part of the coating. The α-stabilized region of ≈5 µm thickness is formed in the surface layer of the base titanium alloy due to the diffusion of the aluminum during the deposition of the coating.
After the CIPT, the coating and the surface of the base titanium alloy have a layered morphology, and each of the layers possesses different grain structure and composition. Moving from the surface of the former Al coating to the depth of the specimen, the following evolution of the microstructure and phase composition is observed.
Phase composition:
TiAl3 → TiAl3 + nc-(Al(Ti) + α-Ti) → nc-(Al(Ti) + α-Ti) → TiAl3 → TiAl3 + TiAl → TiAl → Ti3Al → α-Ti alloy → (α + β)-Ti alloy
Microstructure:
Fine grains (IM) → fine grains (IM) + nanocrystallites (nc-(Al(Ti) + α-Ti)) → nanocrystallites (nc-(Al(Ti) + α-Ti)) → interface “coating/substrate” → ultrafine grains (IM) → fine grains (Ti),
IM—intermetallides, Al—Al-based solid solution, Ti—Ti-based solid solution.
The nanocrystalline aluminum layer, which is formed during the deposition of the aluminum coating near the “titanium alloy/coating” interface, does not undergo recrystallization under the CIPT. The layer can favor the diffusion of the elements in both directions and, therefore, stimulates phase transformation in coarser grains situated under and over the nanocrystalline layer.
Conceptualization, K.R. and E.A.; methodology, K.R. and E.A.; validation, A.N. (Aleksey Nikolaev) and A.N. (Almaz Nazarov); formal analysis, A.N. (Aleksey Nikolaev); investigation, A.N. (Aleksey Nikolaev), E.Z., A.N. (Almaz Nazarov) and V.M.; resources, K.R. and E.A.; data curation, A.N. (Aleksey Nikolaev); writing—original draft preparation, A.N. (Aleksey Nikolaev); writing—review and editing, K.R. and E.A.; visualization, A.N. (Aleksey Nikolaev); supervision, K.R. and E.A.; project administration, K.R. and E.A.; funding acquisition, K.R. and E.A. All authors have read and agreed to the published version of the manuscript.
Not applicable.
Not applicable.
Data available on request.
The equipment of the “Nanotech” center of the Institute of Strength Physics and Materials Science SB RAS was utilized. The authors thank M.Yu. Panchenko for help with TEM study.
The authors declare no conflict of interest.
Footnotes
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Enlarge this image.Figure 1. SEM images of the microstructure in as-received titanium alloy Ti-6Al-4V: (a) detection of secondary electrons; (b) detection of backscattered electrons. Brighter areas correspond to the β-Ti.
Enlarge this image.Figure 2. Cross-sectional view of the titanium alloy specimen coated with aluminum (a–e). Images (a,b)—bright field TEM images, (b) is a magnified image of the square region marked on (a). Image (c) is a selected area electron diffraction pattern corresponding to the image (b) (area of analysis is marked by circle). Dark-field images (d,e) are obtained in reflections 1 and 2 (marked on (c)).
Enlarge this image.Figure 3. Bright-field (a,e) and dark-field (c,d) TEM images of the microstructure in the titanium part of the specimen coated with aluminum. Images (b,f) are selected area electron diffraction patterns corresponding to the images (a,e), respectively (areas of analysis are marked by red and green circles). Dark-field images (c,d) are obtained in reflections 1 and 2 (marked in (b)).
Enlarge this image.Figure 4. TEM bright-field images of the cross-sectional view of the specimen after the CIPT (a–e). Images (b–e) are enlarged images of the areas marked in image (a).
Enlarge this image.Figure 5. TEM image of the microstructure of the coating after the CIPT (a) and selected area electron diffraction patterns from the indicated (a) areas (b,c). White contours on (a) show the interlayer of the nanocrystalline structure between coarse grains of the TiAl3 phase. NC—nanocrystalline region near the interface “coating/substrate”.
Enlarge this image.Figure 6. TEM bright-field (a) and dark-field (c,d) images and corresponding selected area electron diffraction pattern (b) with interpretation obtained for the coating (after the CIPT processing). The interpretation is performed for the two nearest grains of the TiAl3 phase in images (c,d). Image (c) was obtained in (112) reflection of TiAl3 phase, orange cell on (b); d—in (310) TiAl3, blue cell on (b).
Enlarge this image.Figure 7. TEM image of the transverse sections of the sample combined with the marks showing the points of the EDS analysis (Table 2) (a), and corresponding EDS maps for Ti (b), Al (c) and V (d). Specimen was subjected to the CIPT. Image was obtained in scanning TEM regime.
Enlarge this image.Figure 8. TEM bright-field (a,b) and dark-field (d) images of the microstructure of the modified layer in titanium alloy Ti-6Al-4V. The selected area electron diffraction patterns (c) obtained for the area marked on (b). Rings on the image (c) show the radial distribution of the possible positions for the closest reflections of the Al3Ti phase. Dark-field image (d) was obtained in reflection (1) marked on (c).
Enlarge this image.Figure 9. Cross-sectional view (TEM image) of the microstructure in the modified layer in the titanium-based part of the specimen combined with the selected area electron diffraction patterns for the areas marked by the circles.
The experimentally obtained interplanar distances for the diffraction rings on
| Experimental Values | Calculated Values | |||||
|---|---|---|---|---|---|---|
| α-Ti | Al | Al3Ti | ||||
| dhkl, Å | hkl | d, Å | hkl | d, Å | hkl | d, Å |
| 2.556 | 100 | 2.558 | 111 | 2.338 | 002 | 4.291 |
| 101 | 3.515 | |||||
| 002 | 2.341 | 110 | 2.724 | |||
| 2.23 | 101 | 2.244 | 002 | 2.025 | 112 | 2.300 |
| 103 | 2.297 | |||||
| 004 | 2.146 | |||||
| 1.73 | 102 | 1.729 | - | - | 200 | 1.927 |
| 202 | 1.758 | |||||
| 211 | 1.690 | |||||
| 114 | 1.686 | |||||
| 1.48 | 110 | 1.475 | 022 | 1.432 | 105 | 1.568 |
| 213 | 1.476 | |||||
| 204 | 1.434 | |||||
| 006 | 1.431 | |||||
| 1.34 | 103 | 1.326 | - | - | 220 | 1.363 |
| 1.26 | 200 | 1.275 | 113 | 1.221 | 222 | 1.299 |
| 112 | 1.244 | 301 | 1.270 | |||
| 116 | 1.267 | |||||
Elemental composition of the areas (in atomic %) marked on the
| Spectrum | Al | Ti | V | Predicted Phase |
|---|---|---|---|---|
| Spectrum 1 | 77.59 | 21.47 | 0.94 | TiAl3 |
| Spectrum 2 | 76.44 | 22.89 | 0.67 | |
| Spectrum 3 | 97.85 | 2.15 | 0 | Al(Ti,V) + TiAl3 |
| Spectrum 4 | 98.16 | 1.65 | 0.19 | |
| Spectrum 5 | 97.24 | 2.76 | 0 | |
| Spectrum 6 | 80.69 | 18.93 | 0.38 | TiAl3 |
| Spectrum 7 | 78.10 | 21.06 | 0.84 | |
| Spectrum 8 | 77.78 | 21.04 | 1.18 | |
| Spectrum 9 | 79.29 | 20.09 | 0.62 | |
| Spectrum 10 | 70.98 | 28.23 | 0.79 | |
| Spectrum 11 | 77.62 | 21.17 | 1.21 | |
| Spectrum 12 | 52.70 | 46.41 | 0.89 | TiAl |
| Spectrum 13 | 57.68 | 42.18 | 0.14 | |
| Spectrum 14 | 53.95 | 46.05 | 0 | |
| Spectrum 15 | 12.73 | 85.95 | 1.32 | Ti(Al,V) |
| Spectrum 16 | 12.76 | 84.81 | 2.43 | |
| Spectrum 17 | 14.51 | 83.95 | 1.54 | |
| Spectrum 18 | 12.70 | 85.91 | 1.39 |
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1 Department of Mechanical Engineering Technology, Institute of Aviation Technologies and Materials, Ufa University of Science and Technology, Zaki Validi st. 32, 450076 Ufa, Russia;
2 Institute of Strength Physics and Materials Science, pr. Akademicheskii 2/4, 634055 Tomsk, Russia;