Fracture toughness and structural evolution in the TiAlN system upon annealing

Hard coatings used to protect engineering components from external loads and harsh environments should ideally be strong and tough. Here we study the fracture toughness, K IC, of Ti1−xAlxN upon annealing by employing micro-fracture experiments on freestanding films. We found that K IC increases by about 11% when annealing the samples at 900 °C, because the decomposition of the supersaturated matrix leads to the formation of nanometer-sized domains, precipitation of hexagonal-structured B4 AlN (with their significantly larger specific volume), formation of stacking faults, and nano-twins. In contrast, for TiN, where no decomposition processes and formation of nanometer-sized domains can be initiated by an annealing treatment, the fracture toughness K IC remains roughly constant when annealed above the film deposition temperature. As the increase in K IC found for Ti1−xAlxN upon annealing is within statistical errors, we carried out complementary cube corner nanoindentation experiments, which clearly show reduced (or even impeded) crack formation for annealed Ti1−xAlxN as compared with their as-deposited counterpart. The ability of Ti1−xAlxN to maintain and even increase the fracture toughness up to high temperatures in combination with the concomitant age hardening effects and excellent oxidation resistance contributes to the success of this type of coatings.

We carried out cantilever deflection (and cube corner nanoindentation experiments) to study the evolution of the fracture toughness of up to 1000 °C ex-situ vacuum annealed Ti 1−x Al x N free-standing films and correlated them with the film structural evolution and the mechanical properties, hardness (H) and Young's modulus (E), obtained from independent experiments. The mechanical properties were corroborated with HRTEM investigations to give atomic scale insights into the thermally decomposed Ti 1−x Al x N structure. TiN coatings are used as a benchmark, as no decomposition processes are active that would lead to the formation of new nm-sized domains.

Results
Structural evolution. Energy dispersive X-ray spectroscopy (EDXS) analysis rendered a chemical composition of Ti 0.40 Al 0.60 N. Due to the specific sputter condition of the Ti 0.5 Al 0.5 compound target, the coatings prepared are slightly richer in Al than the target for the deposition parameters used 15 . The oxygen content within the coatings is below 1 at.%, as obtained by elastic recoil detection analysis of coatings prepared under comparable conditions 15 . Figure 1a shows the X-ray diffraction patterns of our Ti 0.40 Al 0.60 N films grown onto Al 2 O 3 (1102) substrates after vacuum annealing at different annealing temperatures, T a , for 10 min. Up to 750 °C, Ti 0.40 Al 0.60 N maintains its single phase face-centered cubic (rock-salt-type, B1) structure. The slight peak shift to higher 2θ angles and decrease in peak broadening indicate recovery of built-in structural point and line defects, which results in a lattice parameter decrease in the films. The peak shift to higher 2θ angles also suggests B1 AlN formation (its lattice parameter is smaller as compared to Ti 0.40 Al 0.60 N 16 , hence the diffraction peaks occur at higher 2θ angles). Between 850 and 1000 °C, an asymmetric peak broadening is observed, which indicates isostructural formation of cubic AlN-and TiN-rich domains. Especially, the right shoulder in vicinity of the cubic (200) peak -indicative for cubic AlN formation -is clearly visible and becomes more pronounced with increasing temperature. Hexagonal (wurtzite-type, B4 structured) AlN first emerges at 850 °C and its phase fraction increases with increasing temperature. The shift of the XRD reflections from the major cubic structured Ti 1−x Al x N matrix phase to lower 2θ angles is a result of decreasing Al content (hence, the XRD peaks shift towards the lower 2θ position of TiN). On the other hand, compressive stresses, e.g., induced by the B1 to B4 phase transformation of AlN 17 under volume expansion of ~26% 16 or by thermal stresses, contribute to the peak shift to lower 2θ angles (for the thermal expansion coefficients, α, holds as α B1-(Ti,Al)N > α Al2O3 > α B4-AlN , see refs [18][19][20].
The structural evolution of single phase cubic structured TiN, Fig. 1b, is dominated by recovery of built-in structural point and line defects and results in smaller lattice parameters. Accordingly, the peaks are shifted to larger 2θ angles and become sharper with increasing temperature. Both, Ti 0.40 Al 0.60 N and TiN crystallized in a TEM/HRTEM study. TEM studies were performed on the sample annealed at 900 °C using cross-section samples. A low-magnified image presents an overview of the coating morphology (Fig. 2a), where columnar grains are clearly visible. At this annealing temperature, AlN based hexagonal phases emerge. An atomic resolution TEM image of one portion of grain interfaces are shown in Fig. 2b, the corresponding fast Fourier transforms (FFTs) are seen on the right-hand side. Analysis indicates that a cubic structured Ti 1−x Al x N grain is oriented along the [001] direction while the adjacent hexagonal AlN grain is close to [2110] direction, with an orientation relationship of Ti 1−x Al x N (220)//AlN (0001). This implies that hexagonal AlN (0001) grows on Ti 1−x Al x N (220) planes along this direction. The corresponding FFTs clearly signify the plane relationship between these two phases. This has also been proved by tilting the grains to another orientation. Figure 2c shows one hexagonal AlN grain, grown in between two cubic Ti 1−x Al x N grains, viewed along the [1120] direction while Ti 1−x Al x N is off [001] zone axis, as illustrated in the corresponding FFTs (inserted). Here, only a series of planes appear. The orientation relation is Ti 1−x Al x N (220)//AlN (1100) for this case. It is further noted that the planes in hexagonal AlN are severely distorted or inclined which means that internal stress is strongly involved during the phase transformation. There are numerous defects present in the hexagonal AlN regions, for instance stacking faults and nano-twins marked exemplarily with white arrows in Fig. 2c. In some cases, the AlN phase seems to form in the Ti 1−x Al x N matrix, i.e. Fig. 2b, since the FFT from AlN contains Ti 1− x Al x N spots. However, hexagonal AlN frequently forms at the grain boundary as demonstrated in Fig. 2c, in which the hexagonal AlN and Ti 1−x Al x N phases are separated and formed in between two Ti 1−x Al x N grains. Consequently, the AlN phase transformation (from cubic to hexagonal) can take place in the matrix and also at the grain boundaries, in agreement with earlier studies 21 .
Nanoindentation. The mechanical properties as a function of annealing temperature are presented in Fig. 3 and are in line with previous studies reported in literature 9 . The indentation hardness (H), Fig. 3a, increases for Ti 0.40 Al 0.60 N (red curves) by ~9% from 34 ± 1 GPa in the as-deposited state to 37 ± 2 GPa at 900 °C, before it decreases again down to 28 ± 2 GPa at 1000 °C. The Young's modulus (E), Fig. 3b, shows a similar trend. Contrarily, the hardness of TiN (blue curves) steadily decreases with increasing T a , from 32 ± 1 GPa at room temperature to 27 ± 1 GPa at 850 °C, (Fig. 3a), while the Young's modulus marginally decreases (Fig. 3b). The chosen deposition conditions used in the present study resulted in coatings with excellent mechanical properties in the as-deposited state. In general, age hardening effects are more pronounced for softer coatings, e.g., a relative increase of ~25% was observed for Ti 1−x Al x N with an as-deposited hardness of 'only' ~26 GPa 21 .
The elastic strain to failure [22][23][24][25][26] , (H/E), which is often used to qualitatively rate materials for their failure resistance, suggests superior properties of Ti 0. 40  (Please note that the actual cantilever dimensions, lever arms, and pre-notch depths differ from sample to sample. Hence, Fig. 4a, does not allow direct ranking of the samples with respect to their stiffness and fracture toughness). Figure 4b shows a typical free-standing cantilever. The substrate material had been removed by focused ion beam milling to avoid the influence of residual stresses and substrate interference. Scanning electron micrographs of the post-mortem fracture cross-sections, Fig. 4c,d, do not show discernible changes of the film morphology upon annealing. However, the structure of TiN (Fig. 4d) appears more columnar-grained in comparison with Ti 0.40 Al 0.60 N (Fig. 4c). The K IC values, as calculated from the maximum load at failure, the actual pre-notch depth, and cantilever dimensions using a linear elastic fracture mechanics approach 27 , are presented in Fig. 5. The data suggest an increase in K IC from 2.7 ± 0.3 MPa•√m in the as-deposited state to 3.0 ± 0.01 MPa•√m at 900 °C followed by a decreases to 2.8 ± 0.4 MPa•√m at 1000 °C (red curve). The relative increase of ~11% in fracture toughness of Ti 0.40 Al 0.60 N is similar to the relative increase in hardness of ~9%. Please note, however, that strictly speaking the increase in fracture toughness is within statistical error. Interestingly, the pronounced decrease in hardness at 1000 °C due to wurtzite AlN formation is not observed for K IC , which -in agreement with the H/E criterion-only slightly decreases. Lower K IC values of ~1.9 MPa•√m are found for as-deposited and annealed TiN (blue curve).
To qualitatively proof that K IC increases upon annealing, we carried out independent cube corner nanoindenation experiments on coated Al 2 O 3 (1102) substrates. Scanning electron microscopy images of the indents show aggravated (or even impeded) crack formation for annealed Ti 1−x Al x N samples as compared to the as-deposited counterpart, see Fig. 6. Please note that in the cube corner experiment, residual stresses (e.g., massive compressive residual stresses forming due to the cubic to wurtzite AlN phase transformation under volumes expansion) and the underlying substrate can influence the formation of cracks.

Discussion
The structural evolution of supersaturated cubic Ti 1−x Al x N upon annealing has been experimentally proven in the literature by atom probe tomography 11,28 , small angle X-ray scattering 29 , transmission electron microscopy 30 , and described by phase field simulations 30 : During the early stage, very few nanometer-sized B1 AlN-and TiN-rich domains form in a coherent manner (that is, the crystallographic orientation of the domains correspond to that of the Ti 1−x Al x N parent grain). With progressive annealing time, the domains gain in size and the compositional variations become more pronounced, so that the modulation amplitudes (Ti-and Al-rich) become larger. If the annealing is continued for too long or performed at higher temperatures, coherency strains are relieved by misfit dislocations. Eventually, cubic structured AlN-rich domains transform into the softer but thermodynamically stable (first (semi) coherent then incoherent) hexagonal AlN. The cubic to hexagonal AlN phase transformation is associated with a large volume expansion of ~26% 16 .
Thermally-induced hardening effects in the TiAlN system have been attributed to coherency strains 9 . Coherency strains hinder the movement of dislocations 31 , as it is more difficult for dislocations to passage through a strained than a homogenous lattice. In addition, the coherent domains differ in their elastic properties due to the strong compositional dependent elastic anisotropy of Ti 1−x Al x N 32 , which also hinders the dislocation motion and contributes to the hardness enhancement 32 .
The structural evolution observed in the present study is in line with the literature reports mentioned above. Additionally, we have evidenced severely distorted or inclined lattice planes and numerous defects (including stacking faults) in the hexagonal AlN phase by HRTEM investigations (Fig. 2). This could explain why the measured hardness at 900 °C is relatively high despite the presence of the "soft" hexagonal AlN phase, which is usually reported to deteriorate the hardness.
We have been able to show that besides age hardening effects, the fracture toughness increases upon annealing. Both properties show a similar relative increase of around 10% as compared to the as-deposited state and peak at the same temperature of 900 °C. This suggests that similar microstructural characteristics are responsible for the enhancement of the mechanical properties. We could demonstrate in an earlier study 3 that a coherent nanostructure composed of alternating materials has the potential to enhance the fracture toughness for a certain  bilayer period of a few nanometers. In the superlattice films, also coherency strains 33,34 and variations in the elastic properties are present. It should be mentioned, however, that in contrast to the hardness, the fracture toughness is not primarily governed by the hindrance of dislocation motion: the load-displacement data collected during the cantilever deflection experiments (Fig. 4a) suggest a linear elastic behavior until failure with no indications of plastic deformation.
In agreement with literature reports 21 , we found that cubic AlN forms preferentially at high diffusivity paths such as grain boundaries. If grain boundaries represent the weakest link where cracks preferentially propagate 35 , grain boundary reinforcement 36 has the potential to effectively hinder the crack propagation.
Another important mechanism for increased fracture toughness is phase transformation toughening, which is omnipresent in partially stabilized zirconia bulk ceramics 12 , for example. For Ti 1−x Al x N coatings, the spinodally formed cubic structured AlN-rich domains represent the phase with the ability of a martensitic-like phase transformation from the metastable cubic structure to the stable wurtzite-type (w) variant. The associated volume expansion of ~26% 16 slows down or closes advancing cracks, leading to a significant K IC increase. Therefore, the evolution of K IC with T a of our Ti 0.40 Al 0.60 N coatings is not proportional to that of H with T a , especially at temperatures above 850 °C. The hardness significantly decreases for an increase of T a from 950 to 1000 °C, as also the w-AlN formation significantly increases (please compare Figs. 1 and 3a), but at the same time, the fracture toughness K IC only slightly decreases. The K IC value of 2.8 ± 0.4 MPa•√m after annealing at 1000 °C, is still above that of the as deposited state (with K IC = 2.7 ± 0.3 MPa•√m), whereas the hardness with H = 28 ± 2 GPa after annealing at 1000 °C is significantly below the as deposited value of 34 ± 1 GPa. Hence, effective other mechanisms are present in this type of material, especially when decomposition of the supersaturated matrix phase occurs and w-AlN based phases are able to form.
Note that in the chosen free-standing cantilever setup macro-stresses are relieved and thus do not contribute to the observed toughness enhancement. However, due to the extensive difference in the molar volume between cubic and wurtzite AlN, the thermally-induced formation of hexagonal AlN results in pronounced compressive stresses 17,37 in the application where the coatings are firmly attached to a substrate/engineering component. Compressive stresses result in apparent toughening of Ti 1−x Al x N, as the coating can withstand higher tensile stresses before cracks are initiated (the compressive stresses have to be overcome first before crack formation). The effect of compressive stresses on the fracture toughness is supposed to be much more pronounced than its influence on the hardness. This is why, in real application, the K IC increase upon annealing is expected to be significantly larger than the K IC enhancement found from free-standing micro-cantilever bending tests. This is reflected in the aggravated crack formation observed in the cube corner experiments, see Fig. 6.
As the 'inherent' fracture toughness enhancing effects are strongly connected with the spinodal decomposition, we anticipate that alloying [38][39][40][41] and other concepts to modify the spinodal decomposition characteristics (formation of coherent cubic AlN domains at lower temperatures but delayed formation of the thermodynamically stable phase wurtzite AlN, different shape and size of cubic AlN domains) are applicable to optimize the self-toughening behavior. In general, alloying has the potential to enhance the inherent toughness by modifying the electronic structure and bonding characteristics 42,43 .
The peak in hardness and fracture toughness at 900 °C corresponds to spinodally decomposed TiAlN with fractions of hexagonal AlN as indicated by XRD (Fig. 1) and TEM (Fig. 2). The severely distorted hexagonal AlN with multiple stacking faults suggests that also nano-twinning might become a relevant mechanism. The presence of twins impedes dislocation motion and induces strengthening, but multiple twinning systems can also enhance ductility by acting as a carrier of plasticity 14 .
Based on our results we propose that the additional functionality of Ti 1−x Al x N, i.e. the self-toughening ability at temperatures typical for many various applications, contributes to the outstanding performance of Ti 1−x Al x N coatings in e.g., dry or high speed cutting.

Methods
Sample preparation. Ti 0.40 Al 0.60 N films were deposited in a lab-scale magnetron sputter system (a modified Leybold Heraeus Z400) equipped with a 3 inch powder-metallurgical processed Ti 0.50 Al 0.50 compound target. Polished single crystalline Al 2 O 3 (1102) platelets (10 × 10 × 0.53 mm 3 ) were chosen as substrate materials due to their high thermal stability, inertness and to avoid interdiffusion between film and substrate materials upon annealing up to 1000 °C. Before the deposition, the substrates (ultrasonically pre-cleaned in aceton and ethanol) were heated within the deposition chamber to 500 °C, thermally cleaned for 20 min and sputter cleaned with Ar ions for 10 min. The deposition was performed at the same temperature in a mixed N 2 /Ar atmosphere with a gas flow ratio of 4 sccm/6 sccm and a constant total pressure of 0.35 Pa by setting the target current to 1 A (DC) while applying a DC bias voltage of −50 V to the substrates. The films were grown to a thickness of about 1.8 µm with an average deposition rate of about 75 nm/min. The base pressure was below 5·10 -6 mbar. TiN coatings of about 1.2 µm were synthesized by powering a 3 inch Ti cathode with 500 W within an N 2 /Ar gas mixture (flow ratio of 3 sccm/7 sccm, constant total pressure of 0.4 Pa) and applying a bias voltage of −60 V to the substrates. The deposition rate was about 13 nm/min.
Energy dispersive X-ray spectroscopy (EDXS) measurements of the films were performed with an EDAX Sapphire EDS detector inside a Philips XL-30 scanning electron microscope. Thin film standards characterized by elastic recoil detection analyses were used to calibrate the EDX measurements.
The films on Al 2 O 3 were annealed in a vacuum furnace (Centorr LF22-2000, base pressure <3·10 −3 Pa) at different maximum temperatures (T a ) between 750 and 1000 °C using a heating rate of 20 °C min −1 and passive cooling. At T a , the temperature was kept constant for 10 min.
Structural investigations of coated Al 2 O 3 substrates were performed by X-ray diffraction in symmetric Bragg-Brentano geometry using a PANalytical X'Pert Pro MPD diffractometer (Cu-K α radiation).
ScIEntIfIc REPORTS | 7: 16476 | DOI:10.1038/s41598-017-16751-1 Cross-sectional TEM specimens were prepared using a standard TEM sample preparation approach including cutting, gluing, grinding and dimpling. Finally, Ar ion milling was carried out. A JEOL 2100 F field emission microscope (200 kV) equipped with an image-side C S -corrector with a resolution of 1.2 Å at 200 kV was used. The aberration coefficients were set to be sufficiently small, i.e. C S ~ 10.0 μm. The HRTEM images were taken under a slight over-focus. The HRTEM images were carefully analysed using Digital Micrograph software.
Micromechanical testing. The mechanical properties, hardness and indentation modulus, were measured using a UMIS nanoindenter equipped with a Berkovich tip. At least 30 indents per sample, with increasing loads from 3 to 45 mN were performed. The recorded data were evaluated using the Oliver and Pharr method 44 . To minimize substrate interference, only indents with indentation depths below 10% of the coating thickness were taken into account. The cube corner experiments were carried with the UMIS nanoindenter using a peak indentation load of 150 mN. The high load needed to create cracks resulted in indentation depths of about 1.3 µm in the cube corner experiment.
The fracture toughness was determined from micromechanical cantilever bending tests of free-standing film material. As-deposited and annealed coated Al 2 O 3 samples were broken and their cross-sections carefully polished. The substrate material was removed by Focused Ion Beam (FIB) milling perpendicular to the film growth direction using a FEI Quanta 200 3D DBFIB work station. Then the sample holder was tilted 90° and cantilevers were milled perpendicular to the film surface. The cantilever dimensions of ∼t × t × 6t μm 3 , with t denoting the film thickness, were chosen based on guidelines reported in Brinckmann et al. 45 For the final milling step, the ion beam current was reduced to 500 pA, the initial notch was milled with 50 pA. To circumvent the problem of a finite root radii on the fracture toughness measurements, bridged notches according to Matoy et al. 27 were used (the notch length was chosen to be ∼0.75t).
The micromechanical experiments were performed inside a scanning electron microscope (FEI Quanta 200 FEGSEM) using a PicoIndenter (Hysitron PI85) equipped with a spherical diamond tip with a nominal tip radius of 1 μm. The micro-cantilever beams were loaded displacement-controlled with 5 nm/s with the loading axis perpendicular to the film surface. Per annealing temperature at least 3 tests were conducted. The fracture toughness, K IC , was determined using linear elastic fracture mechanics according to the formula given in ref. 27 : . In the equation, F max denotes the maximum load applied, L the lever arm (distance between the notch and the position of loading), B the width of the cantilever, W the thickness of the cantilever, and a the initial crack length (measured from the post mortem fracture cross-sections).