Fabrication of textured Ti₂AlC lamellar composites with improved mechanical properties

1. Introduction

Ti₂AlC belongs to the family of nanolayered ternary carbides or nitrides (MAX phases). These compounds have a general chemical formula of M_n+1AX_n, in which M is a transition metal, A is an A-group element (III_A and IV_A), and X is either C or N [[1], [2], [3]]. The unique nano laminated structure constituted by covalent-ionic M—X bonds and metallic M—A bonds endows them with a combination of metallic and ceramic properties [4]. MAX phases exhibit properties such as good electrical and thermal conductivity [5], high resistance to thermal shock [6,7], damage tolerance [8], corrosion and oxidation resistance [[9], [10], [11], [12], [13]], self-healing under high temperatures [[14], [15], [16], [17]], high resistance to radiation [[18], [19], [20], [21], [22], [23]] etc. Among the MAX phases, Ti₂AlC has been continually drawing attentions from researchers due to the advantages of relative light weight (4.11 g/cm³), low cost to achieve high purity, and superior oxidation resistance. The combination of the properties Ti₂AlC makes it very promising for corrosive, high temperature, nuclear and other applications.

However, Ti₂AlC has relatively lower compressive strength than conventional structure ceramics with the value only ranges from 393 MPa to 1263 MPa in quasi-static compression test [24]. The reasons for low compressive strength are related to its nano laminated structure, in which complex dislocations activities have been observed in previous works. Study of oriented polycrystalline Ti₃SiC₂ by Barsoum et al. [25] revealed that room temperature compressive deformation started with the generation and motion of dislocations in basal planes which could further be organized in arrays or walls. Dislocation structure of typical low-angle boundary associated with kink band was reported by Farber et al. [26] in room temperature deformed sample of Ti₃SiC₂. Guitton et al. conducted compression experiments of Ti₂AlN under confining gas pressure at room temperature and reported dislocation reactions in MAX phases for the first time [27]. Footprints of a high lattice friction were also observed in the work of Guitton et al. [27] and were detailed by Gouriet et al. [28] with first-principles calculations and a Peierls-Nabarro model. Moreover Guitton et al. [29] also reported out-of-basal-plane dislocations interaction and slip from basal plane to prismatic or pyramidal planes. Additional deform mechanisms including lamellar kinking, delamination [[30], [31], [32]] were also related to dislocation activities. The significant influence of grain microstructure and orientation on dislocation activity was noticed by investigators much latter. Bei et al. first reported the important contribution of grain shape and orientation relative to the compression axis [33] to compressive strength. Their results showed that grain microstructure played a vital role in the compressive deformation process. Then Jones et al. [34] reported that residual strains were compressive in grains with a low Schmid factor and tensile in grains with a high Schmid factor. In situ compression tests combined with neutron diffraction performed on Ti₂AlN by Guitton et al. [35] revealed that lamellar grains parallel to compression axis remain elastic while lamellar grains perpendicular to compression deform plastically, and a Baushinger effect was also proposed in that work. According to all the research works concerning compressive deformation behaviors above, deformation of individual grains depends on grain orientation and load direction. The grains of high Schmid factors especially 45° to the loading direction in compression test first reach the critical shear stress and deform [36]. Basing on the above results, we suppose a method by texturing grains to a preferred orientation to achieve high compressive strength of polycrystalline Ti₂AlC.

In addition to compressive strength, the fracture toughness of MAX phases could also be improved by texturing methods. One major breakthrough comes from texturing of MAX phases by strong magnetic field alignment followed by spark plasma sintering (SPS). This method has been proven to be very effective in boosting the fracture toughness of MAX phases including Nb₄AlC₃ [[37], [38], [39]], Ti₃SiC₂ [40] and Ti₃AlC₂ [41]. Encouraging results in fabricating textured MAX phases in recent years also focused on deforming MAX phases in SPS facilities [[42], [43], [44]]. However, a simple and more efficient fabrication method, i.e, hot pressing (HP) sintering, by which load induced grain texture in MAX phases was first realized [3], attracted less attention. Recently, we developed a new method to fabricate textured Ti₂AlC ceramic by utilizing the effect of loose packed nano flake (NF) powders aligning their basal planes perpendicular to load direction (//c-axis) when they are pressed and deformed. In light of the degree of texture, this method is comparable to spark plasma sintering. This fabrication route can reduce the requirements for texturing equipment and scale up engineering fabrication.

MAX phases (Ti₂AlC in particular) are also characterized by low hardness and flexural strength, which limit their structural and wear applications [7]. One effective method to reinforce MAX phases is to add high hardness and modulus Al₂O₃ to increase their hardness and strength [45,46]. In addition, other benefits include remarkable chemical stability in MAX phases and compatible thermal expansion coefficient can also be obtained. The results showed that addition of Al₂O₃ could improve mechanical properties and wear resistance of MAX phases [[47], [48], [49]]. Moreover the reinforcing effect is more pronounce by decreasing the ceramic particle sizes [50]. Previous studies reported that the highest and lowest values of the compressive strength correspond to Ti₂AlC samples with the smallest and largest grain sizes, respectively. Therefore, it is worth trying to refine the grain size to nano level. In this work, coarse grain (CG) Ti₂AlC powders were ball milled to 200 nm reactively in oxygen containing atmosphere to in-situ synthesize nano scale dispersive Al₂O₃ phase, and partially transforming into Ti₃AlC₂ in the sintering process leading to textured grain (TG) composites.

2. Experimental

2.1. Fabrication process

A schematic representation of the fabrication process for oriented Ti₂AlC nano layer composites is demonstrated in Fig. 1. The Ti₂AlC coarse grain bulks at the beginning of the fabrication process were synthesized by HP sintering. To fabricate Ti₂AlC coarse grain bulks, intermetallic TiAl powders (99.5 wt%, -300 mesh) and TiC powders (99 wt%, 2-4 μm) were used as raw materials. A molar ratio of TiAl: 0.95TiC was adopted. These raw materials were fully mixed and filled into a graphite die, which was loaded to a vacuum HP sintering furnace. High purity Ti₂AlC coarse grain bulks were obtained by sintering the mixture at 1400 °C for 2 h at 30 MPa in an argon atmosphere. The as-sintered Ti₂AlC bulk products were machined to remove surface contaminants. The bulk was then crushed and sieved to collect coarse grain powders (Fig. 2(a)) with average diameter of 40 μm. The coarse grain powders were used to fabricate Ti₂AlC nano flakes. In this work, we utilized a self-made gas reactive ball milling machine to bring about Ti₂AlC lamellar kinking and delamination in grains. Ti₂AlC coarse grain powders disintegrated to nano flakes as the damages accumulated. Reactive gas with volume ratio of O₂:Ar = 1:4 was used to facilitate disintegration of Ti₂AlC grains and introduce oxygen to Ti₂AlC nano flakes, which were converted into in-situ synthesized nano dispersive Al₂O₃ phase. The milling process was conducted at 500 rpm with Si₃N₄ balls and jar for 48 h in dispersant of ethanol. Then the slurry with Ti₂AlC nano flakes of 200 nm and concentration of 30 wt% was dried in an oven to obtain powders (Fig. 2(b)). The textured Ti₂AlC composites were fabricated by HP sintering process by densifying and rearranging the nano flakes under vacuum condition for 2 h at 1200 °C and 30 MPa.

View original graphic| Download| PPT slide

Fig. 1.   Schematic illustration of fabrication process for highly textured Ti₂AlC composite ceramic.

View original graphic| Download| PPT slide

Fig. 2.   SEM images of the Ti₂AlC coarse grain powder (a), and nano flake powder (b).

2.2. Characterization

In order to determine phase contents and degrees of texture, X-ray diffraction (XRD) patterns of all samples were recorded on a diffractometer (Rigaku Smartlab Japan) with Cu Kα radiation at 40 kV and 40 mA. The scanning speed was set at 2°/min and 0.02° per step for all samples in XRD. Quantitative analysis of XRD patterns to retrieve phase contents information was based on Rietveld refinement method. A software of MAUD [51,52] was used in XRD quantitative analysis. Grain orientation and phase distribution maps were constructed by electron backscattered diffraction (EBSD, TESCAN MAIA3, Czech) technique for information of phase morphology and orientation relationships. Before EBSD analysis, surface stress of samples was eliminated by vibratory polishing and etching of broad argon ion beam milling (LEICA, EM RES102, Germany). Scanning electron microscope (SEM, TESCAN MAIA3, Czech) equipped with an energy dispersive spectrometer (EDS, OXFORD X-Max^N, UK) was also used for microstructures characterization. A laser diffraction analyzer (Mastersizer2000, Malvern, UK) was used to measure powder particle size. The Archimedes method was used for density measurement. The analysis of oxygen concentration was done by O/N/H analyzer (LECO TCH600, USA). Image analysis of grain size and phase content was performed with Image-Pro Plus 6.0 software. A universal testing machine (INSTRON 5582, USA) was employed for room temperature mechanical property testing. Cylindrical samples with dimensions of 6 mm in diameter and 9 mm in height were used in compression test. Three-point bending method with single edge notched beam (SENB) samples (2 mm × 4 mm × 20 mm) was used for fracture toughness measurement. Fracture toughness notches in samples were around 2 mm deep and 0.1 mm wide with a 16 mm testing span. Three-point bending test samples (2 mm × 2 mm × 20 mm) with the same span were also used for flexural strength measurement. A hardness indenter (LECO AMH 43, USA) was used to measure microhardness of all samples.

3. Results and discussion

3.1. Effects of sintering temperature on microstructures

EBSD characterization of Ti₂AlC coarse grain bulks fabricated in this work shows that it contains around 3.6 vol% of TiAl_x as impurity phase. The average grain size, based on more than 100 individual grains on the band contrast map of EBSD characterization, is 16.0 ± 11.1 μm in length and 6.5 ± 3.6 μm in thickness. Fig. 2 shows the SEM images of Ti₂AlC coarse grain powders which is crushed and sieved from bulks and Ti₂AlC nano flakes obtained from reactive milling process. Apparently the Ti₂AlC coarse grain powders are dense and polycrystalline with multiple facets. In some coarse grains we could observe cleavage steps. These cleavage steps stem from the nanolaminated structure and the growth of (0001) basal planes which give Ti₂AlC coarse grain flake-like substructure. Those flake-like substructures would disintegrate in subsequent reactive ball milling process and produce large quantities of Ti₂AlC nano flakes.

In order to study the densifying and texturing process of the nano flakes in the HP sintering and to evaluate the effect of sintering temperature, we investigated bulks sintered at 1000, 1100 and 1200 °C by SEM and XRD characterization. Fig. 3 reveals the microstructure of bulk samples sintered at 1000 and 1100 °C. It is clear that the Ti₂AlC phase retained its nano powder structure at 1000 °C with intact spindle-shaped Ti₂AlC phase. The relative density of samples sintered at 1000 °C reached 76.8% of the sample sintered at 1200 °C. With sintering temperature increased to 1100 °C, diffusion in Ti₂AlC nano powder becomes active and powder particles are fused together. The relative density of samples sintered at 1100 °C improved to 89.6% of the sample sintered at 1200 °C.

View original graphic| Download| PPT slide

Fig. 3.   SEM micrograph of fractured TSS surface from samples sintered at 1000 °C (a), and 1100 °C (b).

Fig. 4 shows SEM images of polished and fractured surface of TSS sample sintered at 1200 °C. SEM image of polished surface shows the sample sintered at 1200 °C is fully dense with micro-pores that exist in 1000 and 1100 °C sintered samples completely removed. In Fig. 4(a) the dark dispersive particles were identified as Al₂O₃ grains by EDS method, while light contrast with spindle-shape still can be observed is Ti₂AlC. In the fractured surface of Fig. 4(b), those plate-like MAX phase grains are delaminated in trans-granular mode with angular-grained Al₂O₃ particles exist between the lamellas. We could note that the basal planes of most MAX phase grains are perpendicular to the c-axis manifesting its highly textured nature.

View original graphic| Download| PPT slide

Fig. 4.   SEM micrographs of samples sintered at 1200 °C. (a) Polished TSS plane of textured composite ceramic. The dark contrast indicates Al₂O₃ phase. (b) SEM image of fractured surface of TSS plane of textured composite ceramic.

3.2. Phase evolution and texture

XRD patterns were presented in Fig. 5(a) and (b). The JCPDS files we used for XRD phase identification and calculation were the following: #29-0095 for Ti₂AlC, #52-0875 for Ti₃AlC₂ and #10-0173 for Al₂O₃. The results show that the raw Ti₂AlC coarse grain powders we fabricated are of high purity. Impurity phase peaks which belong to TiC and Ti₃AlC₂ that commonly exist in commercial Ti₂AlC were not discernible. For this study high purity Ti₂AlC as root material is particularly important to discriminate phase changes in following fabrication process. The peak zone in the 2θ of 37.5° to 41° of all results in Fig. 5 was enlarged to illustrate subsequent phase peak changes more clearly. In the enlarged zone of Fig. 5(a), the first obvious peak change comes from the diffraction peaks of Ti₂AlC nano flakes, which broadened due to grain size effect. Ti₂AlC coarse grain powder diffraction pattern shows (103) plane as the main peak and (006) plane as secondary while in the same zone of Ti₂AlC nano flakes only one gentle single-peak was observed. In the enlarged zone, diffraction peak changes for TSS and TTS samples are complicated. First, both (103) and (006) peak of Ti₂AlC reemerged in TTS sample sintered at 1000 °C, suggesting that recrystallization initiated at 1000 °C. Secondly, results of TSS samples sintered at 1000 °C only showed (103) peak while (006) peak almost disappeared, which means that some degrees of texturing exists. But the texturing is not strong enough to cover (103) plane and other {hk0} planes in TTS sample considering the intensity of (103) plane is usually five times stronger than (006) plane. In Fig. 5(b), it seems that there is little difference between samples sintered at 1100 °C and 1000 °C. However, for samples sintered at 1200 °C, strong Ti₃AlC₂ (referred as 312 in Fig. 5 enlarged zone) peak emerged in both TSS and TTS samples indicating some of the Ti₂AlC flakes transformed into Ti₃AlC₂ during the sintering process. Moreover, peak intensity exhibited intense preference with Ti₂AlC (103) and Ti₃AlC₂ (104) almost disappeared in TTS sample. And the Ti₂AlC (006) and Ti₃AlC₂ (008) disappeared in TSS sample. The explanation for the above phenomenon is that both TSS and TTS samples are highly textured which is confirmed by the peaks of other planes in the Fig. 5(b) with {00l} peaks almost disappeared in TSS sample and the peak intensity of {hk0} decreased dramatically in TTS sample.

View original graphic| Download| PPT slide

Fig. 5.   (a) XRD patterns of the Ti₂AlC coarse grain powders (CG), ball milling produced nano flakes (NF), textured top surface (1000 °C-TTS) and textured side surface (1000 °C-TSS) of a sample sintered at 1000 °C. (b) textured top surface (1100 °C-TTS) and textured side surface (1100 °C -TSS) of a sample sintered at 1100 °C, textured top surface (1200 °C-TTS) and textured side surface (1200 °C-TSS) of a sample sintered at 1200 °C.

In order to assess the degree of texture, the Lotgering orientation factor (f_l) was calculated by the equation of f_l=(P-P₀)/(1-P₀). In the calculation of f_(00l), P and P₀ were the ratio of ΣI(00l)/ΣI(hkl) in textured samples and standard JCPDS sample respectively. For calculation of f_(hk0), P and P0 were replaced by the ratio of ΣI(hk0)/ΣI(hkl) correspondingly. ΣI(00l), ΣI(hk0) and ΣI(hkl) are the sums of the peak intensities of the (0 0 l), (h k 0) and (h k l) planes from XRD patterns. The Lotgering orientation factor of Ti₂AlC for {00l} planes in TTS and Ti₃AlC₂ {00l} planes in TTS is 0.74 and 0.49, respectively, suggesting that Ti₂AlC grains are more concentrated along c-axis than Ti₃AlC₂ grains. Al₂O₃ was not observed in the XRD patterns in both NF samples and sintered samples, possibly due to its main peaks overlapping with those of Ti₂AlC and Ti₃AlC₂.

3.3. Crystal orientation and phase distribution

Further investigations of phase composition and grain orientations were carried out by employing EBSD techniques. The constructed orientation and phase distribution map of grains on both TTS and TSS planes are shown in Fig. 6 in which grain colors differ according to their crystal orientations. By examining grain orientation maps (Fig. 6(a) and (c)) with the phase distribution maps (Fig. 6(b) and (d)), the red grains in grain orientation maps are determined Ti₂AlC and Ti₃AlC₂ phases whose [0001] crystal orientation is parallel to c-axis. This means both Ti₂AlC and Ti₃AlC₂ grains are textured with {0001} planes normal to c-axis. Interestingly, some grains show a mixture of Ti₂AlC and Ti₃AlC₂. Apparently Ti₃AlC₂ has been transformed from Ti₂AlC nano flakes due to that fact that Ti₂AlC flakes became unstable with Al atoms diffuse away and are oxidized to form Al₂O₃ phase [53]. Then a portion of Ti₂AlC flakes transformed into Ti₃AlC₂ in the sintering process as the amount of Al vacancy increases. Evidence for this kind of order-to-disorder phase transformation was uncovered by scanning transmission electron microscopy investigation recently [23].

View original graphic| Download| PPT slide

Fig. 6.   Grain orientation and phase distribution map. Inverse pole figure (IPF) map in z-axis of TTS plane (a) and IPF map in y-axis of TSS plane (c) indicating grain orientations. Corresponding phase distribution map of TTS plane (b) and TSS plane (d).

EBSD characterizations also confirm the presence of α-Al₂O₃ phase with the red grain in Fig. 6(b) and (d). Nevertheless it is still a challenge to find the orientation relationships of such small grains on EBSD map so we give the TTS sample’s corresponding pole figures of each phase in Fig. 7. It is evident that Ti₃AlC₂ {0001} and Ti₂AlC {0001} are parallel (the white circles in Fig. 7). The above two phases can intergrowth with each other with the relationship of (0001) Ti₃AlC₂ // (0001) Ti₂AlC [54]. Al₂O₃ phase looks also textured with {0001} plane normal to c-axis but two independent directions were observed. Regular crystallographic orientations reported in previous studies [55] of (110) α-Al₂O₃ // (001) Ti₂AlC and (110) α-Al₂O₃ // (1 1- 0) Ti₂AlC in early oxidation stages were not observed in this work. The orientation relationships of (0001) Ti₃AlC₂ and (0001) Ti₂AlC parallel to (0001) α-Al₂O₃ were reported in MAX phase films [56]. And the orientation relationship between Al₂O₃ oxide scale and Ti₃AlC₂ substrate was (0001) Ti₃AlC₂ // (0001) α-Al₂O₃ [57]. So the orientation relationship of (0001) Ti₂AlC // (0001) Ti₃AlC₂ // (0001) α-Al₂O₃ is highly likely in this work. SEM image of Fig. 4(a) shows the microstructure of TSS sample sintered at 1200 °C, in which the dark dispersive phase was identified as Al₂O₃ grain by EDS method. In order to quantify the diameter and content of Al₂O₃ phase we performed image analysis on SEM images. The result is that the average diameter of Al₂O₃ grains is around 210 nm. That is the primary reason that XRD characterization failed to detect Al₂O₃ phase. The peak intensity of XRD diffraction reduced dramatically with nano grain effect. Moreover main peaks of Al₂O₃ (104) and (113) planes were overlap respectively with Ti₃AlC₂ (102) and Ti₂AlC (104) peaks. Image analysis based on EBSD band contrast map of TSS plane shows that the mean length of MAX phase gains is 1.09 ± 0.64 μm and the mean thickness is 0.37 ± 0.16 μm. According to SEM image analysis, the volume fraction of Al₂O₃ phase is 9.6 vol.% (9.3 wt%), which was supported by chemical analysis result with 4.2 wt% O detected. The Ti₂AlC and Ti₃AlC₂ phase fraction was estimated by EBSD phase calculation as around 60 wt% and 30 wt% respectively, agreeing well with those obtained by XRD quantitative analysis.

View original graphic| Download| PPT slide

Fig. 7.   Pole figures corresponding to Fig. 6(a), of Ti₂AlC, the Ti₃AlC₂ and Al₂O₃. Parallel crystallographic planes of the three phases are indicated by white circles.

3.4. Mechanical properties and strengthening mechanism

Polycrystalline Ti₂AlC with such textured microstructure is expected to possess some unusual mechanical properties. It is interesting to see if the predicted property improvement is realized over conventional coarse grain Ti₂AlC. Typical mechanical properties are listed in Table 1. Results for anisotropic TG samples are classified by the test load direction being parallel and perpendicular to c-axis. Because a small portion Al element turned into Al₂O₃ phase, and Ti₂AlC becomes unstable and transformed to Ti₃AlC₂ phase (4.24 g/cm³), the density of TG composite (4.15 g/cm³) is outweighed pure CG samples (4.10 g/cm³). The average fracture toughness is 8.5 MPa m^1/2 (//c-axis) and 8.2 MPa m^1/2 (⊥c-axis), about 30% higher than untextured CG samples (6.2 MPa m^1/2). Improvement of toughness partially comes from in-situ formed Al₂O₃ particles which could deflect and blunt cracks [41]. Plate grains pulling-out and bridging cracks could also increase fracture toughness. However the fracture toughness values parallel and perpendicular to c-axis are so close. This may result from the morphology of sintered Ti₂AlC grains being spindle like rather than thin flakes. Therefore effects of crack deflecting and bridging are not so anisotropic. Compressive strength of textured composites exceeds 2 GPa, more than 3 times the value of pure CG samples, which is the highest compressive strength for Ti₂AlC and Ti₂AlC composites. Flexural strength is almost 2 times of those of the untextured CG samples. Compared to untextured CG samples, hardness of TG composite samples also increased with the value parallel to c-axis (7.9 GPa) being slightly higher than that vertical to c-axis (6.8 GPa). SEM images in Fig. 8(a) and (b) are Vickers hardness indentations (10 N) of TTS and TSS surface. Apparently the indentation on TTS plane is square while indentation on TSS plane is rhombic. The diagonal length of indent along c-axis exceeds diagonal length perpendicular to the c-axis. Moreover delamination and local fracture were only observed along the direction parallel to c-axis, demonstrating its anisotropic mechanical response of textured microstructure.

View original graphic| Download| PPT slide

Fig. 8.   SEM image of Vickers indentation morphologies of textured composites: (a) TTS and (b) TSS sample.

4. Conclusion

Textured composite containing Ti₂AlC, Ti₃AlC₂ and Al₂O₃ has been successfully fabricated with {00l} planes of Ti₂AlC and Ti₃AlC₂ being highly oriented and vertical to c-axis. Such microstructure is designed according to a number of strengthening mechanisms, including blocking the deformation of Ti₂AlC, Hall-Petch relationship of grain size, and particulate reinforcement. The Lotgering orientation factor of Ti₂AlC and Ti₃AlC₂ {00l} planes in TTS reaches 0.74 and 0.49, respectively, suggesting that the basal planes normal to Ti₂AlC grains are more concentrated than Ti₃AlC₂ grains around c-axis. The orientation relationships of (0001) Ti₂AlC // (0001) Ti₃AlC₂ // (0001) α-Al₂O₃ are possible. The as-fabricated material showed toughness of 8.5 MPa m^1/2 (//c-axis), compressive strength of 2037 MPa (//c-axis), flexural strength of 735 MPa (//c-axis) and high hardness of 7.9 GPa (//c-axis). The synthesis temperature is reduced to 1200 °C due to high surface energy of nano flakes serving as driving force to reduce sintering temperature, which is substantially lower than normal HP synthesis temperature of 1400 °C. This work provides a new method to fabricate highly textured Ti₂AlC bulks with improved mechanical properties by HP sintering, offering a route for fabricating highly textured MAX phases different from the mainstream SPS method.

Acknowledgements

This work was financially supported by the National Key R&D Program of China (Nos. 2017YFB0306201 and 2016YFB0701303).