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Surface & Coatings Technology
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Formation of Ti2AlN phase after post-heat treatment of Ti–Al–N films deposited bypulsed magnetron sputtering
Ying Yang a,b, Martin Keunecke b, Christian Stein b, Li-Jun Gao a, Jun Gong a, Xin Jiang c,Klaus Bewilogua b, Chao Sun a,⁎a State Key Laboratory of Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR Chinab Fraunhofer Institut für Schicht- und Oberflächentechnik, Bienroder Weg 54E, 38108 Braunschweig, Germanyc Institute of Materials Engineering, University of Siegen, Paul-Bonatz-Str. 9–11, Siegen 57076, Germany
Article history:Received 5 October 2010Accepted in revised form 9 November 2011Available online 20 November 2011
Keywords:Magnetron sputteringHeat treatmentMicrostructureTi2AlN phaseTi–Al–N films
Ti–Al–N films with different compositions were prepared by varying the nitrogen flow rate using pulsedmagnetron sputtering and subsequently annealed in air at 700 °C. The microstructure and microhardnessof the films were investigated before and after heat treatment. The XRD results indicate that no Ti2AlNMAX phase occurs in all the as-deposited films. Ti2AlN phases could be acquired only after annealing viasolid state reaction. The formation condition of Ti2AlN phase is strongly dependent on the film composition.It has been found that the nitrogen concentration in the films plays a key role for the phase formation, and thecritical value is in the range of 22.5–29.6 at.%. The fractured cross-section SEM images show that all the as-deposited films exhibit a fine columnar morphology. The annealed films with Ti2AlN phase componentschange the morphology from columnar to fine equiaxed polycrystalline. In addition, the microhardness ofthe films containing Ti2AlN phase is also explored, which is within the range of 18–24 GPa.
Mn+1AXn phases (abbreviated as MAX) are a family of inherentnanolaminated compounds, where M is an early transition metal, Ais an A-group element (mostly IIIA and IVA), X is nitrogen and/or car-bon, and n is equal to 1–3. On account of its unique structure, MAXphases exhibit good combinations of metallic and ceramic behavior,such as good thermal and electrical conductivity, plasticity, excellentthermal shock resistance and high temperature oxidation resistance[1–3]. Owing to these characteristic properties, MAX phases areexpected to be promising for thin film applications; hence, researchesof films with MAX phases have attracted much attention in recentyears [4–6]. In a set of researches of MAX phase materials, Ti2AlNMAX is far away from being well explored, including its microstruc-ture and mechanical properties.
With respect to the fabrication of Ti2AlN phase thin films, effortshave been particularly focused on the direct epitaxial growth modeon single-crystal MgO and Al2O3 substrates using reactive magnetronsputtering [7–9]. However, from the perspective of industrial applica-tion, key problems still remain to be solved, such as high synthesistemperature, rigorous requirements of substrates, limited depositionrate and film thickness. Solid state reaction has also been utilized tothe preparation of Ti2AlN phase thin films. Höglund et al. [10]
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reported Ti2AlN film fabrication by solid state reaction in AlN/Ti(0001) multilayer thin films. Recently, Wang et al.[11] and Garkaset al. [12] synthesized Ti2AlN phase by vacuum annealing Ti–Al–Nfilms at 800 °C for 1 h. However, little information has been publishedon the phase formation condition, and further investigation is stillneeded.
In the present study, we aim to explore the application feasibilityof Ti2AlN phase films on an industrial scale, and also investigate itsformation condition by means of solid state reaction. The influenceof composition on the formation of Ti2AlN phase was investigated,and an appropriate composition window was found. Furthermore;microhardness and abrasive wear rate of Ti2AlN phase have alsobeen studied to predict the potential application of the films.
2. Experimental details
2.1. Deposition parameters
Ti–Al–N films were fabricated by pulsed magnetron sputtering(CC800/9 sinox ML, CemeCon AG, Würselen, Germany) in an Ar/N2
atmosphere with N2 as reactive gas. Single crystalline silicon,100Cr6 ball bearing steel discs (Ø=32 mm, t=4 mm) and highspeed steel (HSS) discs (Ø=32 mm, t=3 mm) were used as sub-strates. Two Ti targets and two TiAl (50 at.%) mosaic targets (size:500×88 mm, purity: 99.99%) were employed. Prior to deposition,the substrates were polished and pre-cleaned with ethanol, and thechamber was evacuated to a pressure of approximately 8×10−4 Pa.
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In order to remove the contaminant layers and to ensure a good adhe-sion of the deposited films, the substrates were etched by argon ionsfor 15 min before deposition.
Based on the preparatory investigation of deposition parameters,an appropriate target power and argon flow rate were chosen andfixed in the present work, and the nitrogen flow rate was varied toprepare Ti–Al–N films with different compositions. The selected Titarget power and TiAl mosaic target power were 1.4 kW and 2.1 kWrespectively, and the argon flow rate was 300 sccm. The total pressureof mixed Ar/N2 was in the range of 0.27–0.29 Pa. During deposition,the substrates, powered with a constant d.c. voltage at a value of−60 V, were rotated during the process to maintain a homogenousfilm and the deposition temperature was approximately 250 °C. Dur-ing the ramp of Ti–Al–N film deposition, the applied target power wasgradually increased from 500 W up to a given value in an argon atmo-sphere in a few minutes, and then the gas flow changed from pureargon (400 sccm) to Ar/N2; therefore, a thin TixAly interlayer was pre-pared. The detailed deposition parameters are listed in Table 1.
2.2. Heat treatment condition
The specimens with HSS substrates were treated in air or in pro-tective H2 atmosphere after deposition in order to study the effectof heat treatment on the film properties. The samples of 1 #–4 #and 6 #–7 # were annealed in air at 700 °C for 30 min. Sample 5 #were annealed in air at 700 °C for 30 and 60 min respectively, andalso in protective H2 atmosphere (the pressure of H2 was 200 Pa) at680 °C for 90 min after deposition.
2.3. Characterization of the films
The thickness of the films was measured by a surface profilometer(Dektak-3). The phases of the films were identified by X-ray diffrac-tion (XRD,Philips X' pert ) using Cu Kα radiation with an incidentbeam angle of 2°. θ–2θ measurement with simplified texture analysiswas also conducted using the same facility. Scanning electron micros-copy (SEM, LEO 1530) was used to examine the morphology of thefilms. The chemical composition of the films was characterized byelectron-probe microanalysis (EPMA, CAMECA SX 100). Depth profil-ing of the elements in the films was determined by secondary ionmass spectroscopy (SIMS, Cameca ims 5f) with a Cs+ ion beam.
The hardness of the films was evaluated from the loading andunloading curves using a computer controlled microhardness tester(Fischerscope H100). Vickers indenter was applied in the present ex-periment. In order to ensure that the indentation depth does not ex-ceed 1/10th of the total film thickness, an indenter load of 40–60 mNwas employed depending on the thickness and the hardness of thefilms. The abrasive wear rate was carried out by a commercial ball-cratering device (KaloMAX NT) with a glycerin/alumina suspension(particle diameter approx. 1 μm) to grind a calotte. A 100Cr6 steelball (Ø=30 mm) was used as the counterpart. The abrasion timewas 3 min, and the rotation speed of the ball was 60 rpm. The abra-sive wear rate was derived from the number of turns of the ball andthe depth of the calotte.
Table 1Detailed deposition parameters of the films.
The deposition rates of the films are listed in Table 1. We can seethat the deposition rate decreases gradually with increasing the nitro-gen flow rate, which is closely related to the poisoning of the targetsby nitrogen gas. The composition of the as-deposited films was deter-mined by utilizing samples with silicon substrates and the results areillustrated in Table 2. The penetration depth was approximately 1 μm.It can be seen that with increasing the nitrogen flow rate, the titaniumconcentration decreases, while the nitrogen concentration increases,and the concentration of aluminum keeps almost constant. At agiven nitrogen flow rate of 25 sccm, the film composition is close tothat of Ti2AlN (see sample 5#). Other elements in the films are Arand O, which are less than 5% in total.
3.2. Morphology — fracture cross-section
The fracture cross-section morphology of the as-deposited sam-ples of 1 #–7 # was investigated and the results indicate that all theas-deposited films exhibit a columnar structure. Fig. 1 shows the mor-phology of sample 2 # as an example. It can be seen that the film has aclear and strong columnar structure.
In order to study the microstructure evolution of the films duringheat treatment, SEM observation of fractured specimens was con-ducted and their images were illustrated in Fig. 2. Fig. 2(a)–(h) dem-onstrates the morphology of the samples annealed in air at 700 °C for30 min. Fig. 2(e) shows the SEM image of sample 5 # annealed in airat 700 °C for 60 min, and Fig. 2(f) displays the microstructure of thesample 5 # annealed in protective H2 atmosphere at 680 °C for90 min. It can be seen that 1 #–3 # samples display a homogenousmorphology with equiaxed grains. It is noteworthy that an interestingphenomenon has been found in Fig. 2(d)–(f). The upper part of thefilms shows a fine equiaxed crystalline structure, while the lowerpart of the films indicates a strong columnar structure. Fig. 2(g) and
Fig. 1. Fracture cross-section SEM image of the as-deposited sample 2 #.
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(h) displays a columnar structure of sample 6 # and sample 7 #,respectively.
3.3. Composition — SIMS results
With the aim of elaboration of the SEM results in Fig. 2(d)–(f),sample 5 # was selected to perform SIMS experiments, and the re-sults are shown in Fig. 3. Fig. 3(a) displays the composition depthprofile of the as-deposited film. Fig. 3(b) and (c) indicates the
Fig. 2. The fracture cross-section SEM micrograph of the annealed films. Notes: (e) sample 5sphere at 680 °C for 90 min; other samples were annealed in air at 700 °C for 30 min.
composition distribution along the depth after annealing in air at700 °C for 60 min and in protective H2 atmosphere at 680 °C for90 min respectively. It can be seen from Fig. 3(a) that there is aTi3Al interlayer near the film/substrate interface. With respect to theTi–Al–N film, the composition is not homogenous along the filmdepth at the first stage of deposition. The content of Ti increasesand the content of N decreases gradually with increasing the filmthickness, and the content of Al maintains almost constant. With afurther increase in deposition time, the film composition keeps
# annealed in air at 700 °C for 60 min; (f) sample 5 # annealed in protective H2 atmo-
Fig. 3. SIMS results of sample 5 # at different states. (a) The as-deposited sample(substrate: 100Cr6). (b) The sample annealed in air at 700 °C for 60 min (substrate:HSS). (c) The sample annealed in protective H2 atmosphere at 680 °C for 90 min(substrate: HSS).
Fig. 4. The phases of the annealed films deposited at various nitrogen flow rate (sub-strate: HSS). (a) The annealed samples 1 #–4 # and (b) the annealed samples 5 #–7 #.
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invariable as shown in the second stage, reaching the ratio of Ti:Al:Nnearly equals to 2:1:1, which is also consistent with the EPMA result.Other elements are Ar and C, which are less than 5 at.% in total. Sim-ilar results can also be found in Fig. 3(b) and (c). On account of thelimited amount of the nitrogen flow rate and the relatively largesize of the chamber (size: 850×850×1000 mm3), the process takesmore time until the film composition becomes constant. It is assumedthat the composition depth distribution in the other films is also nothomogeneous as a result of the same reason. It can be seen from
Fig. 3(b) that oxygen is only present at the surface of the films, andthere is little amount of oxygen in the inner films.
By comparing the three figures, it can be seen that no significantcomposition change in the upper part of the film has been identifiedafter heat treatment. Diffusion only occurs between the interlayerand the film near the interface, leading to the changes of compositionwithin the interlayer, and the diffusion becomes more abundantwhilst the duration of heat treatment is increased.
3.4. Structure of the films — X-ray diffraction
The XRD results of the as-deposited films are also presented inTable 2. It can be seen that the crystal structure of the as-depositedfilms is highly dependent on the film composition. In addition toTi3AlN inverse perovskite phase, the films are composed of α-Ti metal-lic phase, hex.-Ti3Al intermetallic phase and fcc-(Ti,Al)N ceramic phase.Similar results have been found in Ref [12]. No hex.-Ti2AlN phase wasdetected in all the as-deposited films, probably due to the relativelylow synthesis temperature. The inherent complex structure of MAXphases is important in determining the growth temperature [13].
The phase structure of the annealed films is described in Fig. 4. Asfor sample 1 #, the peak at 13° can only be assigned to the hex.-Ti2AlNphase (13° is the characteristic peak of Ti2AlN phase). Besides Ti2AlNphase, α-Ti, Ti3Al intermetallic phase and possibly a few amount ofTi3AlN are still present in the films. Compared to sample 1 #, Ti2AlNphase is more pronounced in the annealed sample 2 #, and Ti3AlNphase has almost vanished. The XRD patterns of the annealed filmsof sample 3 # and sample 4 # are similar, which are composed mainly
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of Ti2AlN phase within the range of the X-ray penetration depth. Withrespect to sample 5 # annealed at atmosphere for 30 min, Ti2AlNphase with broad peaks of fcc-TiN phase (solid solution of (Ti,Al)N)and Ti3AlN phase have been detected.
With the aim to explore the impact of annealing time on the phasestructure of the films, sample 5 # was annealed at atmosphere for an-other 30 min (in total 60 min). The result shows that the Ti2AlN peaksare more pronounced than that after the 30 min treatment. Broadpeaks of fcc-TiN phase and Ti3AlN phase still exist. The XRD resultsof annealed sample 6 # and sample 7 # are similar with those of theas-deposited samples, and fcc-TiN phase and Ti3AlN phase fit withthe patterns.
It can be seen that Ti2AlN phase was identified in the annealedsamples of 1 #–5 #. No Ti2AlN phase was observed in the annealedspecimens of 6 # and 7 #. The possible interpretation is that with in-creasing the nitrogen flow rate, the nitrogen concentration in thefilms increases, and the film is prone to form (Ti,Al)N phase. (Ti,Al)N is a relatively thermal stable phase and has strong Ti–N bonding.Hence, it is difficult for Al element layer to be interleaved betweenTi–N slabs and therefore the formation of Ti2AlN phase is suppressed.
3.5. Microhardness and abrasive wear rate
This section addresses the microhardness and abrasive wear rateof the films before and after heat treatment. It can be seen fromFig. 5(a) that the microhardness of the as-deposited films mainly in-creases with the nitrogen flow rate, and there is no major change inthe microhardness after annealing. The microhardness in the films
Fig. 5. Film properties as a function of nitrogen flow rate (substrate: HSS). (a) Micro-hardness and (b) abrasive wear rate.
with Ti2AlN phase inclusion is within the range of 18–24 GPa (seethe annealed 1 #–5 # samples). The measured value is not the micro-hardness of the pure Ti2AlN phase film because of the influences ofthe inclusions of other phases and the inhomogenous microstructurealong the depth. Nevertheless, this result is comparable with otherstudies [6,14] and 4 to 5 times higher than that of bulk materials(4.3 GPa [15]). The higher hardness in the films compared to that ofbulk materials can be attributed to the Hall–Petch effect [16] due tothe much smaller grain size in the films (normally order of magnitudeμm in the bulk materials).
The abrasive wear rate mainly decreases with increasing the nitro-gen flow rate both in the as-deposited and annealed films and the wearrate of the annealed films is higher than that in the as-deposited filmsas shown in Fig. 5(b). In order to facilitate comparison, the abrasivewear rate of (Ti,Al)N film prepared under the same experimental con-dition is approximately 3 x 10−15 m3 m−1 N−1, which is nearly onethird of the films with Ti2AlN phase components. However, morework needs to be carried out to better understand and exploit thewear properties of the films with Ti2AlN phase inclusion.
4. Discussion
The process window and composition window which is appropri-ate for the formation of Ti2AlN phase have been deduced based on theXRD results mentioned above. As the nitrogen flow rate increasesfrom 10 sccm to 25 sccm, the atomic percentage of N increases from8 to 23%. Ti2AlN phase can be formed in this composition region bymeans of solid state reaction, which possibly occurs among Ti3AlN,TiN, Ti3Al, etc. Similar phenomena have been reported by Garkas etal. [12]. With a further increase of the nitrogen flow rate, no Ti2AlNphase could be detected in the annealed films. It can be seen thatthe nitrogen concentration in the films plays a key role in theTi2AlN phase formation, and no Ti2AlN phase can be identified whenthe nitrogen concentration exceeds a critical value, which is between22.5 and 29.6 at.%. Unfortunately, the detailed formation mechanismis still unclear, and further investigation would be needed. In spiteof this, it is demonstrated that the fabrication of Ti2AlN phase film isnot only limited to laboratory-scale studies, but also can be per-formed using an industrial-size coater, which is a big step to the ap-plication of MAX phase film at an industrial scale. Furthermore,Ti2AlN phase can be acquired by simply annealing Ti–Al–N films at at-mosphere. This indicates that Ti2AlN phase can be self-formed duringperformance in a given condition and be expected to exhibit promis-ing corrosion and oxidation resistance [11,17].
In combination with the XRD, SEM and SIMS results, Ti2AlN phasewas identified in the annealed samples 1 #–5 #; meanwhile, the mi-crostructure changed from columnar to fine equiaxed polycrystallinedue to annealing. With respect to the annealed samples 6 # and 7 #,the morphology remains the same as it is in the as-deposited films,and no Ti2AlN phase was determined. Therefore, it is speculatedthat Ti2AlN phase formation is accompanied by the microstructureevolution from columnar to fine equiaxed crystals in the process ofsolid state reaction.
5. Conclusion
Ti–Al–N films were synthesized by pulsed magnetron sputteringand Ti2AlN phase was formed by means of annealing. The microstruc-ture and microhardness of the films before and after heat treatmentwere investigated. The main results are summarized as follows:
1. There is no Ti2AlN phase observed in all the as-deposited films, andit can be formed by means of annealing in air at 700 °C for 30/60 min or in protective H2 atmosphere at 680 °C for 90 min.Ti2AlN phase formation is strongly related to the film composition.An appropriate process window and composition window have
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been found, and the critical nitrogen content is in the range of22.5–29.6 at.%. Furthermore, it is demonstrated that Ti2AlN phasefilm can be acquired using industrial coating equipment, which isinspiring for the industrial-scale application of Ti2AlN phase film.
2. The fractured cross-section SEM images indicate that all the as-deposited film exhibit a columnar morphology, and it changes tofine equiaxed crystals following with the phase formation ofTi2AlN phase.
3. The microhardness of the films composed of Ti2AlN phase displaysa good value of 18–24 GPa.
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