Desorption of Al and Phase Transformation of Ti2AlN MAX Thin Film

Article pubs.acs.org/JPCC

Desorption of Al and Phase Transformation of Ti2AlN MAX Thin Film upon Annealing in Ultra-High-Vacuum Zheng Zhang,† Hongmei Jin,*,‡ Jianwei Chai,† Lu Shen,† Hwee Leng Seng,† Jisheng Pan,† Lai Mun Wong,† Michael B. Sullivan,‡ and Shi Jie Wang*,† †

Institute of Materials Research and Engineering, A*STAR (Agency for Science, Technology and Research), 3 Research Link, 117602 Singapore ‡ Institute of High Performance Computing, A*STAR (Agency for Science, Technology and Research), 1 Fusionopolis Way, Connexis, 138632 Singapore S Supporting Information *

ABSTRACT: Phase stability of single-crystalline Ti2AlN thin film in ultra-high vacuum has been studied in situ by X-ray photoelectron spectroscopy as a function of annealing temperature and ex situ by atomic force microscopy, secondary ion mass spectroscopy, energy-dispersive X-ray spectroscopy, X-ray diffraction, and nanoindentation. Ti2AlN is stable up to 600 °C. At 700 °C, Al is preferentially desorbed from the surface and becomes nearly undetected at 900 °C by XPS, where single-crystalline Ti2AlN with terrace morphology transforms into polycrystalline δ-TiN1−x and ξ-TiN0.75−y phases with voids on the surface and reduced film thickness. Mechanical properties including hardness and Young’s modulus are also observed to have deteriorated. Density functional theory calculation shows that Al atoms prefer to diffuse out from the Ti2AlN horizontally along the Al basal planes. The subsequent desorption of Al from surface due to its high vapor pressure results in the decreased Al composition, the void formation on the surface, and the decomposition of Ti2AlN. A kinetic model involving diffusion and desorption processes is proposed to describe the Al behavior and voids formation above 700 °C.

I. INTRODUCTION Ternary nitride, Ti2AlN, belongs to the interesting Mn+1AXn or MAX phase family, which concurrently possesses both ceramic and metallic properties (M: an early transition metal, A: an element in groups IIA and IVA, X: N or C, n = 1−3).1 As a representative of MAX phase, Ti2AlN exhibits a layered hexagonal structure (space group P63/mmc) that is constructed by vertically repeated two Ti2N layers intercalated by one atomic Al layer in between. The Ti2N layer can be seen as composed of horizontally repeated edge-sharing Ti6N octahedral and is identical to the NaCl-type structure of TiN.2 The strong covalent-ionic nature of Ti−N bonds gives Ti2AlN typical ceramic properties such as high melting point and hightemperature oxidation resistance, while the weak metallic nature of Ti−Al bonds renders Ti2AlN typical metallic properties such as good electrical and thermal conductivity, high ductility, and ease of machinability.1,3 As a result, Ti2AlN has great potentials to be applied in various fields, especially in high-temperature environments such as automobile/aircraft engine components, heating elements, heat exchangers, gas burner nozzles, wear- and corrosion-protective surface coatings, and so on. Fabrication of Ti2AlN can be traced back to 1997, when bulk polycrystalline Ti2AlN was fabricated through hot isostatic pressing of Ti and AlN powders at 1600 °C for 4 h under a pressure of 40 MPa.4 Single-crystalline Ti2AlN thin films deposited on MgO and Al2O3 substrates through ultra-high © 2014 American Chemical Society

vacuum (UHV) direct-current (DC) magnetron sputtering were only reported from year 2005 onward.3,5−9 Since their discovery, both polycrystalline and single-crystalline Ti2AlN phases have been subjected to extensive ex situ characterization for their bulk properties,2−8,10that is, structural, electrical, thermal, and mechanical properties. Little work has been carried out to analyze the Ti2AlN’s properties in situ (without oxidizing the films in air),3,6,7 especially for the thermal stability of its phase and composition, which are crucial for hightemperature applications previously mentioned. Although it is generally accepted that MAX phases do not melt congruently but instead decompose via desorption of A (mainly) and/or M elements, which have higher vapor pressures, the decomposition pathway and temperature can be significantly affected by various factors including sample thickness (bulk sample or thin films), environment (ambient air or vacuum), underlying substrates, and presence of impurities.6,10 For example, MAX phases are known to decompose at lower temperatures when they are in form of thin films instead of bulk materials due to shorter diffusion length, or when they are inside vacuum chambers instead of in ambient environment due to the absence of surface oxides as diffusion barrier, or when there is presence of impurities inside the bulk materials. Received: June 2, 2014 Revised: August 8, 2014 Published: August 14, 2014 20927

dx.doi.org/10.1021/jp505428a | J. Phys. Chem. C 2014, 118, 20927−20939

The Journal of Physical Chemistry C

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samples by setting the Au 4f7/2, Ag 3d5/2, and Cu 2p3/2 peaks at binding energies (BEs) of 83.96 ± 0.02, 368.21 ± 0.02, and 932.62 ± 0.02 eV, respectively. As Ti2AlN film is conducting in nature, shift in BE due to charging after photoelectron emission is negligible, and no further BE correction is performed. The error of BE is estimated to be within ±0.2 eV. The sample was then annealed through resistive heating from 400 to 900 °C at a step size of 100 °C with the temperature being monitored by a pyrometer. At each temperature, the sample was held constant for a manually controlled period (typically 1 h) until XPS results showed no significant change in the composition before the temperature was raised further. After the end of the annealing experiment, the sample was cooled to room temperature and taken out of the UHV chamber for further characterization. The surface morphology of the films was examined by Digital Instrument Nanoscope IV AFM. The distribution of the elements inside the resulting thin film was analyzed by time-of-flight secondary ion mass spectroscopy (TOF-SIMS) using a TOF-SIMS IV instrument from IONTOF GmbH. A 3 keV Ar+ ion beam was used for sputtering, and a pulsed 25 keV Bi+ ion beam was used as analysis beam with detection of positive secondary ions. The crystalline structures of thin films were investigated using Bruker general area detector diffraction system X-ray diffraction (GADDS XRD) operated at a voltage of 40 kV and a current of 40 mA (Cu Kα X-ray, λ = 1.54 Å). Hardness and Young’s modulus of the films were determined from a MTS Nano Indenter XP using a continuous stiffness measurement (CSM) technique with a Berkovich indenter (three-faced pyramid diamond with a 20 nm radius of curvature at the apex). During the indentation test, the indenter was pressed into the film from the sample surface to 300 nm deep into the sample at a constant strain rate of 0.05 s−1. This method allows stiffness to be recorded continuously along the indentation depth. Modulus and hardness values can then be derived from stiffness data, and their profiles can be obtained after plotting the values as a function of displacement into surface. The highest hardness and modulus values obtained slightly below the surface are reported as properties of the thin film.16 The electronic structures of all materials are calculated using Vienna ab initio simulation package (VASP).17 The generalized gradient approximation (GGA-PBE)18 scheme is used for electron exchange and correlation, while the frozen-core projector-augmented wave (PAW) method is used to describe the interaction between ions and electrons.19 A plane-wave cutoff of 450 eV is employed, and Brillouin-zone integration is performed on a Monkhorst−Pack grid by using different grid densities. The structure is relaxed until the convergence of Hellmann−Feynman forces is 1500 °C and eventually into TiN0.75 at >1600 °C via desorption of Al only.12−14 Under a compressive pressure of 5 GPa, Kou et al. detected that their Ti2AlN disk (11 mm in diameter and 2 mm in height) decomposes into AlTi and TiN at 1500 °C through ex situ XRD study, and that the presence of Zr impurity would promote the decomposition of Ti2AlN.15 We have recently grown single-crystalline Ti2AlN thin film on MgO(111) substrate at 750 °C using DC magnetron sputtering from a Ti2Al alloy target in a mixed N2/Ar plasma.9 Compared with amorphous and polycrystalline phases, singlecrystalline phase has long-range order in crystal structure and is free of grains/domains. Thus, the single-crystalline thin film would help understand the intrinsic material properties of Ti2AlN free from the effects of dislocation, grain boundary, and grain size as well as the influence from oxidation in the ambient air. In view of the controversial reports about decomposition of Ti2AlN in the literature, we would like to shed more light on the thermal stability of Ti2AlN thin film through an in situ Xray photoelectron spectroscopy (XPS) study as a function of annealing temperature as well as ex situ characterization involving atomic force microscopy (AFM), secondary ion mass spectroscopy (SIMS), energy-dispersive X-ray spectroscopy (EDS), XRD, nanoindentation, and density functional theory (DFT) calculation. We also aim to bridge the gap between the very thin 50 nm film from Beckers et al. and the bulk sample from Low et al. and Kou et al. by investigating 400 nm single-crystalline Ti2AlN thin film’s stability in situ in a UHV environment. In addition, the decomposition pathway and kinetics will also be discussed.

II. EXPERIMENTAL METHODS 400 nm Ti2AlN (0002) thin films were grown on MgO(111) substrates at 750 °C after 1.5 h deposition in a growth chamber with a base pressure of 5.0 × 10−9 mbar through DC magnetron sputtering of a 3 in. Ti2Al alloy target (99.99% purity) at a power of 90 W in a mixture of Ar (3.2 × 10−3 mbar) and N2 (6.0 × 10−5 mbar) plasma.9 After deposition, the samples were transferred in situ (without exposing to air) to the analysis chamber of VG ESCALAB 220i-XL XPS, which has a base pressure of 1.0 × 10−9 mbar. A monochromatic Al Kα (1486.7 eV) X-ray with a diameter of 700 μm was employed, while the photoelectrons were collected at a normal takeoff angle with respect to surface plane. The electron analyzer is calibrated with polycrystalline gold, silver, and copper standard

III. RESULTS AND DISCUSSION A. Experimental Results. The high-resolution XPS spectra of Ti 2p, Al 2p, N 1s, the valence band (VB), Mg 1s, and O 1s from as-deposited Ti2AlN (which has been cooled to 25 °C) and after annealing in situ in XPS’s UHV chamber to 400, 500, 600, 700, 800, and 900 °C are shown in Figure 1. The spectra in each figure are scaled to have the same maximum height with the scaling factor indicated near the left axis of each spectrum for easy comparison. A larger scaling factor implies a weaker intensity in the original spectrum and vice versa. The peak shape and intensity of Ti 2p remained roughly unchanged from 25 to 900 °C. It is worth noting that while the peak BE of Ti 2p3/2 remained constant at 454.4 eV up to 600 °C, it shifted to 20928



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Figure 1. XPS high resolution spectra of (a) Ti 2p, (b) Al 2p, (c) N 1s, (d) valence band (VB), and (e) Mg 1s and O 1s from as-deposited Ti2AlN after cooling down to room temperature (25 °C) and after annealed in situ in ultra-high vacuum (UHV) chamber to 400, 500, 600, 700, 800, and 900 °C. The spectra in each figure are normalized to have the same maximum peak heights for each comparison. The resulting change in composition of Ti2AlN during in situ annealing is calculated based on the respective peak areas from panels a−c and is shown in panel f.

component between 1 and 3 eV (labeled as “B”) was ∼1.4 times higher than the component between 0 and 1 eV (denoted as “A”) for the as-deposited film at 25 °C (Figure 1d). However, these two components had similar height at 800 °C, while the intensity of the component A far exceeded that of component B at 900 °C. As the VB is a collection of the valence electrons that participate in bond formation and reflects the total-density-of-states of all the atoms on the surface, a change in VB implies a change in the surface chemical environment, which is coherent with the shift of Ti 2p and N 1s at 900 °C, and the change in VB will be discussed in more detail in Section B (Figure 7). Mg 1s signals were not detected, while the O 1s signals were weak throughout the annealing, suggesting that the thin film after annealing is still covering the underlying MgO(111) substrate and no considerable oxides are formed during annealing (Figure 1e). The relative change in the composition of Ti2AlN as a function of annealing temperature can be calculated from the spectra shown in Figure 1a−c after a Shirley-type background subtraction with consideration of the corresponding Scofield

a higher BE of 454.7 eV at 800 and 900 °C (Figure 1a). Upon heating to 400 °C, an additional component around 75.2 eV appeared in Al 2p spectra together with the main component at 72.3 eV (Figure 1b). The main component at 72.3 eV is the characteristic of Ti2AlN phase,9,20,21 while the small component at 75.2 eV can be attributed to either AlN or Al2O3. Because there was insignificant presence of oxygen throughout annealing (Figure 1e), the formation of Al2O3 is unlikely, and we assign the component at 75.2 eV to AlN. Although the peak shape and intensity of Al 2p remained alike until 600 °C, its peak intensity decreased significantly at 800 °C evidenced by a much bigger scaling factor of 1.9 versus 1.1 at 600 °C. At 900 °C, Al 2p signals were only weakly detected, indicating a pronounced loss of Al from surface at this temperature. The peak shape of N 1s stayed constant at 397.4 eV until 600 °C but shifted to a lower BE of 397.3 eV at 900 °C. Because the BE of Al 2p remained unchanged, the opposite direction in the shift of Ti 2p and N 1s BE at 900 °C indicates a change in the chemical environment of Ti and N and an additional charge transfer from Ti to N at 900 °C. In the VB spectra, the 20929



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Figure 2. 5 μm × 5 μm AFM images of (a) the as-deposited Ti2AlN (0002) thin film and (b) after annealing in situ to 900 °C in UHV. (c,d) Higher magnification of the surfaces (2 μm × 2 μm) from the respective areas marked by black squares in panels a and b, respectively. The line profiles of the green lines marked in panels c and d are shown in panels e and f, respectively.

the decease in Al% after annealing in UHV. The first one is that Al segregates to the surface and forms Al-rich 3D islands. When such Al islands grow higher than the core-level photoelectron attenuation depth (3λ) of Al 2p, Al atoms below 3λ from surface effectively do not contribute to the Al XPS signals. This process also exposes more Ti and N within the analysis area. As a result, there is a decrease in the Al% when compared with an initially flat surface. The second possibility is that Al desorbs from surface, which will cause a loss of materials from the

photoionization cross sections, the transmission function of the spectrometer, and the energy compensation factor.22 It can be seen clearly from the resulting plot in Figure 1f that the composition of Ti, Al, and N remained stable at around 38.1, 26.8, and 35.2%, respectively, up to 600 °C. Al % slightly increased to 28.6% at 700 °C before it decreased to 20.9% at 800 °C. After annealing 1 h at 900 °C, Al % reduced to only ∼4.3%, while Ti and N % increased to 44.7 and 51.0%, respectively. There are three plausible scenarios pertaining to 20930



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Figure 3. TOF-SIMS depth profiles of (a) the as-deposited Ti2AlN (0002) thin film and (b) after annealing in situ to 900 °C in UHV.

TOF-SIMS intensity profiles of Ti, Al, N, Mg, and O in the as-deposited and after annealing Ti2AlN are shown in Figure 3 as a function of sputtering time. Three regions can be identified in the as-deposited Ti2AlN thin film on MgO(111) substrate (Figure 3a). In the first region from surface until ∼300 s sputtering, intensities of Ti, Al, and N stayed consistent, and this region is attributed to the single-crystalline Ti2AlN layer. In the second region from 300 to 450 s, intensities of Ti and N started to decrease, while intensity of Al dipped initially but climbed up before it decreased again. Intensities of Mg and O started to appear and rise sharply. This second region is the interfacial layer between single-crystalline Ti2AlN and MgO, and spinel phase, that is, MgAl2O4, is likely formed here. The third region from 450 s onward, where intensities of Mg and O stayed stable while intensities of Ti and N were nearly not detected, corresponds to the MgO(111) substrate. After annealing in vacuum at 900 °C, three regions are still visible in the depth profile (Figure 3b). In the first region from surface until ∼200 s sputtering, intensities of Ti and N stayed consistent and were comparable to those in the as-deposited film. However, the intensity of Al was nearly negligible, which indicates that Al is depleted from this first region. In the second region from 200 to 450 s, intensities of Ti and N started to decrease, while intensities of Mg and O started to increase, similar to their behaviors in as-deposited film. The intensity of Al increased from nearly zero to two climaxes with the first one slightly higher than second one before it declined. It is worth noting that the second climax is close in intensity to the only Al climax in the as-deposited film, and thus it is likely produced during deposition. The newly formed higher climax of Al after annealing can be attributed to inward diffusion of Al. The third region, where intensities of Mg and O were stable, corresponds to the MgO(111) substrate. Al and Ti inward diffusion coupled to Mg outward diffusion were also observed by Beckers et al. during annealing of 50 nm Ti2AlN grown on MgO(111) at 750 °C.6 Our TOF-SIMS results clearly demonstrate an inward diffusion of Al after annealing in vacuum at 900 °C, while the inward diffusion of Ti and outward diffusion of Mg are less pronounced. However, if inward diffusion of Al is the only reason accounting for depletion of Al signals from surface and formation of holes on the surface, then overall Al composition inside the film should stay roughly the same. To verify this, we examine the film’s

surface and subsequent formation of voids/holes on the surface. The third likelihood is that Al diffuses inward to react with MgO at the interface and forms MgAl2O4 spinel phase, as recently observed by Beckers et al.6 In this way, Al diffusing more than 3λ away from the surface is unable to be detected by XPS. In the meantime, surface morphology is likely to change as a result of a change in crystalline phase from Ti2AlN into MgAl2O4. All three of these mechanisms will modify the surface but in different ways; the surface morphology of the Ti2AlN after annealing was thus examined under AFM. The as-deposited 400 nm Ti2AlN thin films appeared to be silver in color, which is consistent with Joelsson et al.’s observation.8 When examined under AFM, the morphology shows continuous terraces flow and a hexagonal surface plateau, which suggests a layer-by-layer growth mode where the growth rate along basal plane far exceeds that across basal plane (or along the c axis, Figure 2a,c). The step height is measured to be 1.4 nm (Figure 2e), close to Ti2AlN’s lattice constant of 1.3614 nm on the c axis.1 After annealing, the silver-colored Ti2AlN film changed to a golden color, which is typically observed for titanium nitrides. Accordingly, the microscopic AFM images showed that the step-flow surface evolved into irregular flat-top 2D islands with occasional triangle shape at the edges and random voids/holes (regions with black color) clearly seen in between the islands (Figure 2b,d). The depth of the voids is measured to be in the range of 10−20 nm (Figure 2f). There is no formation of visible 3D islands on the surface. Hence, we exclude the first possibility of Al-rich 3D island formation previously discussed. The formation of holes and voids implies the loss of materials from the surface, which can occur through desorption, inward diffusion, or both. For example, holes have been observed previously on the surfaces of various materials during annealing in vacuum due to desorption of elements with relative higher vapor pressures, that is, Al in Ti2AlN,11−13 Ti4AlN311,13 and Ti2AlC,23 Si in Ti3SiC2,24 Ge in Si0.8Ge0.2,25 Sb and Te in Ge2Sb2Te5,26 and so on. Voids and cavities can also be created underneath the surface alumina oxide scale when NiAl and FeAl alloy are oxidized in air with the inward diffusion of Ni and Fe.27 To differentiate materials loss through desorption from inward diffusion, we further characterized distribution of Al along the depth and the composition of Al inside the resulting thin film by TOF-SIMS and EDS, respectively. 20931



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Figure 4. EDS point scans of the as-deposited Ti2AlN (0002) thin film (black curve) and after annealing in situ to 900 °C in UHV (red curve).

Figure 5. XRD patterns of the as-deposited Ti2AlN (0002) thin film (bottom black curve) and after annealing in situ to 900 °C in UHV (top brown curve).

annealing (1.0 × 10−8 mbar). Al desorption and porous surfaces have been observed following vacuum annealing and decomposition of cylindrical bars of Ti2AlN, Ti4AlN3, and Ti2AlC.11−13,23 In addition, there is a slight reduction in film thickness from 414 ± 87 nm of the as-deposited film to 359 ± 40 nm of the film after annealing at 900 °C, as measured by surface profiler. Therefore, the decrease in Al% as probed by XPS, a reduced film thickness, and the formation of holes on the surface in our work should be similarly ascribed to Al desorption at 900 °C. Some minor Ti desorption might also take place at 900 °C as Ti’s vapor pressure is calculated to be 3.8 × 10−7 mbar at 900 °C,28 which is slightly higher than the base pressure of 1.0 × 10−8 mbar during annealing. To probe the chemical structure of the remaining Al-deficient Ti−N film, we perform a XRD scanning on the thin film after annealing. The incident X-ray angle (Ω) is carefully set to avoid the diffraction peaks from MgO(111) substrate as their extremely high intensities would permanently saturate the

composition using energy dispersive X-ray spectroscopy (EDS) installed inside a JEOL thermal field-emission scanning electron microscope (JSM-7600F) with X-MAX50 silicon drift detector operating at an accelerating voltage of 8 kV. This accelerating voltage is carefully selected to probe as thick of a layer as possible within the Ti2AlN thin film while avoiding the signals from MgO substrate. In the EDS spectra (Figure 4), although the intensity of N Kα peak remained roughly the same, there was a dramatic decrease in the intensities of Al Kα and Al Kβ together with a small decrease in the intensities of Ti Kα and Ti Kβ. The Al and Ti atomic compositions (at %) in the resulting thin film after annealing decreased from 17.5 to 1.0% and from 51.3 to 47.9%, respectively, while N composition increased from 31.2 to 51.2%. The bulk composition after annealing probed by EDS clearly indicates a reduction in Al concentration. Al is known to have a high vapor pressure, which can be calculated to be 4.5 × 10−3 mbar at 900 °C28 and far exceeds the base pressure in the UHV chamber during 20932



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Figure 6. Hardness and Young’s modulus of the as-deposited Ti2AlN (0002) thin film and after annealing in situ at 900 °C in UHV as measured by nanoindentation. The literature values for single-crystalline TiN (001) and TiN(111)30 as well as polycrystalline Ti2AlN4 are also included with * mark for reference.

curve) and 900 °C (the purple curve) pronouncedly. Because the intensity of VB is closely associated with the total density of states (TDOS) of the surface, DFT calculation is performed to understand the changes behind XPS and XRD results with special attention being paid to the region between 0 and −3 eV. The partial density-of-state (PDOS) and total DOS (TDOS) of perfect Ti2AlN were calculated and are shown in Figure 7b. It can be seen from TDOS and PDOS that the component A around EF from 0 to −1 eV in VB (fitted with a blue peak) is mainly contributed by Ti 3d with minor contribution from Al 3p and N 2p states. The component B from −1 to −3 eV in VB (fitted with red and green peaks) is mainly contributed by Ti 3d and Al 3p with minor contribution from Ti 3p, Al 3s, and N 2p. These observations agree with the previous reports.31−33 To have a better understanding of the VB at 800 °C, we calculated TDOS of various Al-deficient Ti2AlN cells, which were built by deliberately removing Al atoms one by one from a supercell of 2 × 2 × 1 Ti2AlN. The main development after removing Al atoms is a gradual decrease in the TDOS between −1 and −3 eV. This is understandable because the TDOS from −1 to −3 eV is contributed mainly by Ti 3d and Al 3p states (Figure 7b). A decrease in number of Al remaining in the structure will lead to a decrease in total Al 3p states and thus a reduction in TDOS between −1 and −3 eV regions. The TDOS after removing four out of a total of eight Al atoms from the 2 × 2 × 1 Ti2AlN best matches VB at 800 °C and is plotted in Figure 7c. It can be seen that the TDOS of component B has decreased significantly as a result of a decrease in PDOS of Al 3p and Ti 3d states. The contribution from Ti 3d state to component A stays nearly unchanged. As a result, the TDOS of component B relative to component A decreases when compared with that in the perfect Ti2AlN (Figure 7b). At 900 °C, the component B in VB seemed to disappear while the component A around EF seemed to develop from a Gaussian shape into an asymmetric Lorentzian shape (Figure 7d). We have calculated the TDOS of TiN, TiN0.9, and TiN0.75, which are shown in Figure S1 of the Supporting Information. The TDOS between 0 and −3 eV is mainly contributed from Ti 3d states, while the N vacancy in TiN (which results in formation of TiN0.9 or TiN0.75) seems to have trivial effect on

pixels in the 2D detector. As a result, MgO(111) and (222) peaks were not present in the XRD patterns. Compared with a single series of (002) peaks in diffraction pattern of asdeposited single-crystalline Ti2AlN, there are many more peaks in the film after annealing (Figure 5). While a small peak at 23.2° is likely to be attributed to Ti2N(101), the other diffraction peaks at 36.6, 42.5, 61.7, 73.7, and 77.5° can be assigned to four possible structures, namely, TiN0.9, TiN0.76, (TiN)0.96, and TiN. The presence of Ti2N (101) peak at 23.2° suggests that the first step of decomposition of Ti2AlN at vacuum occurs through the following pathway: 800 ° C

Ti 2AlN(s) ⎯⎯⎯⎯⎯→ ε‐Ti 2N(s) + Al(g)↑

(1)

According to the phase diagram of Ti−N, ε-Ti2N decomposes into δ-TiN1−x and ξ-TiN0.75−y between 1080 and 1290 °C and further into α-Ti(N) and δ-TiN1−x between 1290 and 2300 °C before it melts into liquid from 2300 °C and above.29 Therefore, it is likely that ε-Ti2N is further decomposed in vacuum when the temperature increases through the reaction 900 ° C

ε‐Ti 2N(s) ⎯⎯⎯⎯⎯→ δ‐TiN1 − x(s) + ξ‐TiN0.75 − y(s)

(2)

The thin film after annealing is likely to be a mixture of δTiN1−x (e.g., TiN0.9) and ξ-TiN0.75−y (values of x and y denoting N-deficient phases). Because both of them share the same cubic structure with very close lattice constants, it is difficult to differentiate them apart from XRD. Nevertheless, with a change in thin film’s nature from single-crystalline Ti2AlN MAX phase to a mixture of polycrystalline Ti nitrides, the thin film’s hardness and Young’s modulus are measured to be lower than both the single-crystalline phases of Ti2AlN9 and TiN30 but still higher than that of polycrystalline Ti2AlN4 (Figure 6). B. Density Functional Theory Calculation Result. The change in crystalline structure is manifested not only in the XRD pattern but also in the XPS’s VB spectra. By stacking all VB spectra in Figure 1d without scaling, it can be seen that the VB stayed roughly unchanged from as-deposited to 700 °C (Figure 7a). However, the VB in the region between 0 and −3 eV decreased in intensity and changed in shape at 800 (the pink 20933



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Figure 7. (a) Stacked valence band (VB) spectra of Ti2AlN during annealing display discernible changes between 0 and −3 eV at 800 and 900 °C. Such changes are correlated with DFT calculations of partial density-of-state (DOS) of Ti, Al, and N and total DOS inside (b) a perfect 1 × 1 × 1 Ti2AlN (0002) unit, (c) four Al atoms removed from 2 × 2 × 1 Ti2AlN (0002) structure, and (d) one N atom removed from a 1 × 1 × 1 TiN (001) unit, which are representing as-deposited Ti2AlN (0002) thin film, after annealing at 800 and 900 °C, respectively.

TDOS in this region. Therefore, the TDOS and PDOS of TiN0.75 are plotted as a representative in Figure 7d. It can be seen that the TDOS decreased almost linearly from 0 to −3 eV and is close to the VB captured by XPS as a reflection of change in crystalline structure from Ti2AlN to a mixture of TiN1−x and TiN0.76−y. Bader charge analysis of the charge transfer between Ti and N in Ti2AlN, TiN, TiN0.9, and TiN0.75 is listed in Table 1. Each Ti loses 1.230 electrons, while each N gains −1.750 electrons in Ti2AlN. However, Ti can lose 1.200 to 1.697 electrons, while each N gains −1.766 electrons in TiN0.75. By transforming from Ti2AlN into TiN0.75 or a mixture of TiN0.75

with TiN0.9, Ti transfers more electrons to N. This agrees well with the opposite movement of Ti 2p’s BE toward a higher value and N 1s’s BE toward a lower value at 900 °C as shown in Figure 1. By now, it is clear that whole temperature range up to 900 °C can be broadly divided into two regimes according to the stability of Ti2AlN MAX phase. When the annealing temperature is not >600 °C (herein defined as Regime I), Ti2AlN MAX phase is stable and no Al desorption is detected. At temperatures above 600 °C (defined as Regime II), Al desorption from surface together with interfacial reaction of 20934



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Table 1. Charge Transfer (e) of Ti and N in Ti2AlN, TiN, TiN0.9, and TiN0.75 calculated by Bader Charge Analysis Based on the Supercell Structures Described in Figure S1 of the Supporting Informationa materials

charge on each Ti atom

charge on each N atom

Ti2AlN TiN TiN0.9 TiN0.75

1.230 1.670 1.116−1.674 1.200−1.697

−1.750 −1.670 −1.610 ∼ −1.740 −1.766

a Range of values indicates different amounts of charge transfer for Ti and N atoms at different positions in TiN0.9 and TiN0.75.

Al with MgO(111) substrate is taking place. In addition, it is also observed by XPS that the intensity of Al 2p decreased initially but stabilized over time (after 1 h) at each temperature in Regime II. If only aluminum desorption takes place in Regime II, the Al% ratio should keep decreasing instead of stabilizing at each temperature after the initial drop. Therefore, besides Al desorption, it would be reasonable to expect a concurrent Al diffusion from bulk to surface in Regime II due to a concentration gradient (which serves as the driving force) between the bulk and the surface layers after the initial loss of Al from the surface. We have carried out DFT calculation to compare the diffusion path of Al from bulk to surface along the c axis (vertically) and along the Al basal plane (e.g, horizontal xy plane). It is found out that the activation energy of Al diffusion along the horizontal xy plane is 0.70 eV, which is much lower than that of Al diffusion along the c axis (∼10.77 eV). After Al diffuses horizontally to the step edge, it will be exposed to free surface immediately due to the flat terrace morphology and can then be desorbed to the vacuum easily at high temperatures. C. Kinetics of Ti2AlN Decomposition and Al Desorption−Diffusion above 600 °C. To quantitatively describe the kinetics of Ti2AlN decomposition, Al desorption, and diffusion in Regime II, we propose a model to fit the XPS data obtained during the steady states (where the Ti2AlN thin film composition is stable) in Regime II. λ is the inelastic mean free path of photoelectron in a sample and is typically 600 °C, both the surface layer and the bulk layer possess the same overall Al concentration (denoted as CAlo, in units of number of Al atoms per cm3, as shown in Figure 8a). In Regime II, when temperature is >600 °C, Al desorbs to the vapor phase (Al(v)) with a rate constant of ksdes, as shown in Figure 8b. As desorption of Al takes place, diffusion of Al from the bulk layer to the surface region denoted as Al(d) would arise due to Al concentration gradient between these two layers. This process occurs with a rate constant given by kBDif, as shown in Figure 8b. Thus, at a given temperature, the measured Al concentration (CAl) at a given time in the surface layer (i.e., 3λ) will depend on the relative rate of desorption and diffusion. Initially, the desorption rate is high while the diffusion rate is low due to small concentration gradient. As diffusion rate catches up with desorption rate due to higher concentration gradient, decrease in Al% slows down and Al% eventually stabilizes to a constant value after 1 h, where a steady state is achieved between the desorption and diffusion rate.

Figure 8. Schematic diagram of the model depicting Al desorption and diffusion process from Ti2AlN during annealing in (a) Regime-I (≤600 °C) and (b) Regime-II (>600 °C).

For a given element A, its concentration (CA, number of atoms of the A element per cm3) is proportional to its XPS intensity as measured within the XPS analysis depth (3λ) and is given by CA = IA /DA

(3)

where IA and DA are peak area and peak-related constant for A element’s XPS peak, respectively. Likewise, the values of Al concentration (CAl) at various temperatures can be related to the Al 2p’s peak area in Figure 1b using eq 3. The peak areas (I) for Ti 2p, Al 2p, and N 1s as a function of temperature are shown in Figure 9a. It can be seen that the intensities of Al 2p, Ti 2p, and N 1s stayed nearly unchanged in Regime I (from RT to 600 °C); therefore, the initial Al intensity (IAlo) is determined experimentally as the averaged IAl values in Regime I. In Regime II (>600 °C), the IAl values at steady state decreased as annealing temperature increased. This would suggest a decrease in CAl as temperature increases. In contrast, intensities of Ti 2p and N 1s consistently increased throughout Regime II. Because Ti2AlN(0001) is a layered structure composed of alternate stacking Al layer and Ti2N layer, the removal of Al layers through desorption would expose more Ti2N layers within the probing depth of XPS and therefore is expected to contribute to an increase in the intensities of Ti 2p and N 1s peaks. The steady-state Al concentration (CAl) measured at a given temperature can be normalized into a coefficient S, which is defined as a ratio of the measured concentration in the steadystate regime (CAl) to the original Al concentration (CAlo) and is written as 20935



Article

Figure 9. (a) Change of peak intensities of Al 2p, Ti 2p, and N 1s during annealing in UHV. (b) Plot of S as a function of temperature for Al 2p by taking the average peak intensity from 25 to 600 °C as IAlo. Plots of (c) ΔS and (d) S/(1 − S) as a function of temperature for Al 2p. The fitting results are represented by dashed lines.

S=

CAl I /D I = Al A = Al CAlo IAlo/DA IAlo

rDec =

(4)

IAl 700°C − IAl800°C IAlo

The variation of the steady state coefficients (S) with

=

temperature calculated from Al 2p’s peak intensity is shown in =

Figure 9b. Our aim is to describe the kinetics of these steady states in Regime II, which involves decomposition, desorption,

=

and diffusion processes through fitting the stabilized values of S

t ⎛ IAl700°C ×⎜ I − ⎝ Alo

IAl800°C ⎞ IAlo

⎟

⎠

t IAlo t IAlo t

× (S700°C − S800°C) × ΔS800°C

(5)

where ΔS800°C is a reduction in S when temperature increased from 700 to 800 °C. Following Low et al.’s work,11−14 we considered the decomposition rate (rDec) as

as a function of temperature. We will first consider the rate of decomposition. In Regime II, the decrease in Al 2p’s peak intensity is due to

⎛ E ⎞ rDec = A × exp⎜ − Dec ⎟ ⎝ kBT ⎠

decomposition of Ti2AlN and subsequently desorption of Al into the vacuum environment. Hence, the decomposition rate at each temperature can be defined as the decrease in Al 2p

(6)

where A is a constant, EDec is the decomposition activation energy, kB is the Boltzmann’s constant, and T is the absolute temperature. Combining eqs 5 and 6, ΔS can be related to EDec through the following equation

intensity divided by the annealing time (t). For example, the decomposition rate (rDec) at 800 °C can be expressed as 20936


The Journal of Physical Chemistry C ⎛ E A×t 1⎞ ΔS = × exp⎜ − Dec × ⎟ IAlo T⎠ ⎝ kB

Article

where the activation energies EBDif and Esdes and the preexponential factors ABDif and Asdes are related to the rate constant for Al diffusion (kBDif) and Al desorption (ksdes) by the usual Arrhenius expressions, k = A exp (−E/kBT), where kB is the Boltzmann constant. The experimental plot of S/(1 − S) against 1/T is shown in Figure 9d, and a good exponential fit is obtained. Esdes − EBDif can be calculated to be 1.14 eV similarly using formula of kB/−t1 and the “t1” value from fitting. At the end of Section B, our DFT calculation shows that Al prefers to diffuse along the Al horizontal plane instead of crossing the Al and Ti2N planes. The lowest diffusion activation energy we obtain is 0.70 eV when diffusing along the Al basal plane. If we take this value as s diffusion activation energy (kBDif ̀ ), Edes is calculated to be 1.84 eV. Our recent paper shows that the surface of Ti2AlN tends to be Ti1- or Al-terminated due to the smaller surface energies in their configuration (Table 2).34 Hence, after an Al atom diffuses

(7)

An exponential function is used to fit the experimental plot of ΔS against 1/T, and a good fit is shown in Figure 9c. After obtaining the value of “t1” (the variable defined in the fitting equation) from fitting, the decomposition activation energy (EDec) can be calculated to be 1.29 eV using the formula kB/t1. This EDec value is higher than the EDec of Ti2AlC (0.89 eV or 85.7 kJ mol−1)14 and is reasonable considering that Ti2AlN is reported to be more stable than Ti2AlC.15,33 Next, we will consider the desorption and diffusion process. For simplicity, we assume that the rate of Al desorption at surface is first-order with respect to the measured Al concentration (CAl). Hence the rate of Al concentration change due to desorption (rdes) can be given as follows rdes =

dCAl s = kdes CAl dt

Table 2. Surface Energies (J/m2) of Ti2AlN Terminated by Ti1, Al, Ti2, and N Show That Surface of Ti2AlN is Most Stable When Terminated by Ti1 Layer or Al Layer Due to Their Smaller Surface Energy Values34a

(8)

where ksdes is the rate constant for desorption from the XPS analysis regime into vacuum, as shown schematically in Figure 8b. As prevously explained, the measured Al concentration (CAl) at a given temperature depends on both the loss of Al through desorption and the gain of Al through diffusion from bulk region. We assume that the rate for this diffusion process depends not only on the Al concentration in the bulk (i.e., CAlo) but also on the availability of empty sites created through the loss of Al (i.e., (CAlo − CAl)) within the detected layer. Thus, the rate of Al concentration change within these diffusing layers (rDif) can be written as follows rDif =

dCAl B = kDif CAlo(CAlo − CAl) dt

surface energy

Ti2 termination

N termination

2.20

2.21

5.07

2.46

The layers below the Ti1 and Ti2 layers are N layer and Al layer, respectively.

laterally to the step edge, it will require breaking lateral Al−Al bond or vertical Al−Ti bond before it can desorb into the vacuum. A value of 1.84 eV for desorption activation energy (Esdes) is considered to be reasonable given that it falls in between the bond dissociation energies of Al−Al and Al−Ti (1.38 and 2.73 eV, respectively).28 D. Discussion. In this work, 400 nm single-crystalline Ti2AlN thin film was grown at 700 °C in a mixed Ar/N2 plasma at a pressure of 3.2 × 10−3 mbar. The same Ti2AlN film, however, started decomposing at 700 °C in the UHV chamber of XPS. It is because the vapor pressure of Al at 700 °C (4.1 × 10−6 mbar) is lower than the pressure during growth (3.2 × 10−3 mbar) but far exceeds the base pressure of UHV chamber (1.0 × 10−9 mbar). The 400 nm Ti2AlN thin film subsequently decomposed completely at 900 °C. Our decomposition temperatures (700−900 °C) are much lower than the temperatures (1500−1800 °C) reported by Low et al. for bulk Ti2AlN11−14 but are close to values (750−800 °C) reported by Becker et al. for Ti2AlN thin films.6 This is expected as the starting temperature for surface-initiated phase decomposition during annealing is known to be the lowest for thin films in vacuum, followed by bulk material in vacuum and the highest for bulk materials in air.6,10−12 Decomposition of cylindrical bars of Ti2AlN after annealing 1 h in vacuum at 1550 °C actually only happened at the outermost 140 μm layer after examining the cross-section of the sample under SEM.12 Hence, factors including short diffusion distance from thin film to the surface (400 nm), surface sensitive detection technique (XPS), UHV environment, and the absence of oxide scale as a desorption barrier together contribute to the lower decomposition temperature of Ti2AlN thin film in our work. The very small film thickness can also help explain the different decomposition pathway Becker et al. observed for the same 50 nm Ti2AlN(0001) thin film but grown on MgO(111)

(9)

(10)

In the steady-state regime at each temperature where the Al concentration stabilizes with time dCAl =0 dt B s kDif CAlo(CAlo − CAl) − kdes CAl = 0 s ⎛ kdes CAl ⎞ C C B ⎜ ⎟ kDif 1 − − × Al = 0, as S = Al ⎟ ⎜ CAlo ⎠ CAlo CAlo CAlo ⎝

hence B kDif (1 − S) =

Al termination

a

where kBDif is the rate constant for diffusion from bulk into the XPS analysis regime, as shown schematically in Figure 8b. Thereafter, by taking into consideration the surface desorption and bulk diffusion process, the rate of Al concentration change at a given temperature is thus expressed by dCAl B s = kDif CAlo(CAlo − CAl) − kdes CAl dt

Ti1 termination

s kdes ×S CAlo

B B ⎛ Es − E B CAlo × kDif CAlo × ADif S 1⎞ Dif = = exp⎜ des × ⎟ s s 1−S kdes Ades kB T⎠ ⎝

(11) 20937



Article

Figure 10. Schematic drawings of the cross-section structures of (a) as-deposited Ti2AlN thin film and after annealing at (b) 700 and (c) 900 °C.

raised to 900 °C, Al desorption speeds up due to higher thermal energy, while the diffusion rate cannot catch up due to limited Al supply in the 400 nm thin film. As a result, Al completely desorbs from the thin film after 1 h (Figure 10c). Subsequently, the remaining Ti2N layers transform into δTiN1−x and ξ-TiN0.75−y following the Ti−N phase diagram. They stack upon each other by filling the empty layer left by the Al atoms, leading to a decrease in film thickness. Desorption of Ti is also likely as its vapor pressure has increased above the base pressure in UHV environment.

and Al2O3(0001) during vacuum annealing, besides the influence from different substrates. As described in the Introduction, Becker et al. observed 50 nm Ti2AlN(0001) thin film grown on Al2O3(0001) substrate decomposed into Ti2N accompanied by desorption of Al at 800 °C.6 However, he did not observe Al desorption when the same 50 nm Ti2AlN thin film but grown on MgO(111) substrate decomposed at 750 °C into Mg2TiO4 and MgAl2O4 spinel phases through Ti and Al inward diffusion and Mg outward diffusion.6 In the decomposition of our Ti2AlN(0001) thin film grown on MgO(111), Al inward diffusion and possible Mg−Al−O spinel phase formation at the interface with MgO substrate are similarly observed from TOF-SIMS depth profiles. In addition, Al desorption is also observed from voids at the surface shown in the AFM image, a reduced film thickness, and a decrease in Al% detected by both XPS and EDS. The desorption of A elements in MAX phase is frequently observed during decomposition of the corresponding MAX due to its higher vapor pressure. We believe that it is because the Ti2AlN film is so thin (50 nm) in Becker et al.’s work that Al in the film has largely reacted with MgO to form spinel phase before Al can desorb from Ti2AlN during decomposition. As a result, Becker did not observe desorption of Al from 50 nm Ti2AlN(0001) grown on MgO(111) substrate. Our Ti2AlN film is eight times thicker than that of Becker et al. and hence sufficient for the loss of Al from Ti2AlN through both desorption from surface and diffusion into the MgO substrate. Our observation agrees largely with Low et al.’s decomposition path from Ti2AlN into TiN0.5 (>1500 °C) and eventually into TiN0.75 (>1600 °C) via desorption of Al,12−14 although the corresponding decomposition temperatures are much lower in our case. On the basis of the results, the desorption and diffusion of Al from Ti2AlN can be schematically described in Figure 10 (The interfacial reaction between Ti2AlN and MgO is not included because the reaction does not universally occur but occurs only on MgO substrate.). In Regime I, when temperature is not >600 °C, the phase and composition of Ti2AlN are stable (Figure 10a). In Regime II, when temperature is >600 °C, Al sitting on the step edge breaks bonds with Al horizontally and with Ti vertically and then desorbs into the vacuum as its vapor pressure at 700 °C exceeds the base pressure in UHV chamber (Figure 10b). Desorption of Al leads to an initial decrease in Al % and creates a concentration gradient, which drives Al in the bulk to diffuse preferably along horizontal Al plane rather than vertically across Ti2N layer to the step edge due to its weak bonding with the two adjacent Ti layers. A steady state is established when the diffusion rate catches up with the desorption rate and the Al composition stops declining and becomes stable. Small holes/voids appear as a result of Al desorption from the surface. When the temperature is further

IV. CONCLUSIONS In summary, the thermal stability of 400 nm single-crystalline Ti2AlN thin film in UHV has been studied by XPS, AFM, TOFSIMS, EDS, XRD, nanoindentation, and DFT calculations. In Regime I, when temperature is not >600 °C, phase and composition of Ti2AlN is stable. In Regime II, when temperature is >600 °C, Al desorption occurs due to its high vapor pressure, which leads to a decrease in Al surface concentration. As a result, decomposition of hexagonal Ti2AlN into cubic Ti2N is initiated from surface and is propagated inward. At 900 °C, Al is nearly undetected, while Ti2N further transforms into δ-TiN1−x and ξ-TiN0.75−y phases accompanied by holes formation on the surface and a reduced film thickness due to loss of Al. By modeling the intensity of Al 2p spectra from XPS data, we obtained values of 1.29 and 1.84 eV for Ti2AlN decomposition activation energy (EDec) and Al s desorption activation energy (Edes ), respectively, which substantiates the Al desorption and diffusion mechanism. After decomposition, the remaining polycrystalline δ-TiN1−x and ξ-TiN0.75−y film still possess considerable hardness and Young’s modulus (with values ∼50% of those of singlecrystalline Ti2AlN and higher than the hardness of polycrystalline Ti2AlN). Besides desorption from surface, Al at the Ti2AlN/MgO(111) interface region also diffuses inward and reacts with MgO to form Mg−Al−O spinel phase. The decomposition of single-crystalline Ti2AlN but in ambient air due to oxidation would be an interesting subject for further study and comparison.

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ASSOCIATED CONTENT

S Supporting Information *

Partial density of states and total density of states calculated by density function theory from (a) a perfect 1 × 1 × 1 TiN structure, (b) TiN0.9 represented by removing three N atoms from 2 × 2 × 2 TiN, and (c) TiN0.75 represented by removing one N atom from 1 × 1 × 1 TiN supercell structure. This material is available free of charge via the Internet at http:// pubs.acs.org. 20938



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AUTHOR INFORMATION

Corresponding Authors

*H.M.Jin: Tel: +65-64191332. Fax: +65-64632536. E-mail: [email protected]. *S.J.Wang: Tel: +65-68748184. Fax: +65-67744657. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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Desorption of Al and Phase Transformation of Ti2AlN MAX Thin Film

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