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Thermal stability of Ti3SiC2 thin films
Acta Materialia 55 (2007) 1479–1488 www.actamat-journals.com
Thermal stability of Ti3SiC2 thin films Jens Emmerlich a,*, Denis Music b, Per Eklund a, Ola Wilhelmsson c, Ulf Jansson c, Jochen M. Schneider b, Hans Ho¨gberg a, Lars Hultman a a
c
Linko¨ping University, Department of Physics, IFM, Thin Film Physics Division, SE-581 83 Linko¨ping, Sweden b Materials Chemistry, RWTH Aachen University, Kopernikusstraße 16, D-52074 Aachen, Germany Uppsala University, Department of Materials Chemistry, The Angstro ¨ m Laboratory, P.O. Box 538, SE-751 21 Uppsala, Sweden Received 26 May 2006; received in revised form 4 October 2006; accepted 5 October 2006 Available online 18 December 2006
Abstract The thermal stability of Ti3SiC2(0 0 0 1) thin films is studied by in situ X-ray diffraction analysis during vacuum furnace annealing in combination with X-ray photoelectron spectroscopy, transmission electron microscopy and scanning transmission electron microscopy with energy dispersive X-ray analysis. The films are found to be stable during annealing at temperatures up to 1000 C for 25 h. Annealing at 1100–1200 C results in the rapid decomposition of Ti3SiC2 by Si out-diffusion along the basal planes via domain boundaries to the free surface with subsequent evaporation. As a consequence, the material shrinks by the relaxation of the Ti3C2 slabs and, it is proposed, by an in-diffusion of O into the empty Si-mirror planes. The phase transformation process is followed by the detwinning of the as-relaxed Ti3C2 slabs into (1 1 1)-oriented TiC0.67 layers, which begin recrystallizing at 1300 C. Ab initio calculations are provided supporting the presented decomposition mechanisms. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ti3SiC2 thin films; Phase transformations; X-ray diffraction; Transmission electron microscopy; Ab initio electron theory
1. Introduction Ti3SiC2 is a Mn+1AXn phase (n = 1–3), where M = early transition metal, A = an A-group element (usually group 13–14), and X = C or N. The MAX phases exhibit a hexagonal structure (space group P63/mmc). In Ti3SiC2, Ti3C2 layers (isostructural to the binary TiC phase, where C is at octahedral sites) are intercalated with atomic layers of Si, which act as mirror planes [1]. The Si is weakly bonded to the Ti, whereas the covalent Ti–C bonds are much stronger [2,3]. MAX phases in general exhibit a high thermal stability, but little is known about the phase transformation pathways at temperatures above the equilibrium temperature range of the different MAX phases. In particular, there is a need to determine the potential phase decomposition in *
Corresponding author. Tel.: +46 13 281281. E-mail address: [email protected] (J. Emmerlich).
applications where the MAX phase material is in contact with another chemical medium. For example, there is a growing interest in thin films of Ti3SiC2 as electrical contacts [4]. For high-power applications, the thermal stability of the MAX phase is an issue as high thermal loads in the contact area to metals like Cu and Ni or to air can be expected. Furthermore, some of the phases are suitable for heating elements [5] and/or as protective coatings, in which case the material is subject to a wide range of interfaces and environments. Considering the Ti–Si–C phase diagrams at 1100 and 1250 C [6], Ti3SiC2 is in equilibrium with TiC or possible (Ti,Si)C [7]. Tie-lines can be observed to the Si-rich phases of SiC, TiSi2 and Ti5Si3. Ti3SiC2 bulk material was found to be thermally stable in Ar atmosphere to at least 1800 C [8,9]. Immersing Ti3SiC2 in molten cryolite at 960 C, however, transforms the phase topotactically into TiC with some retained Si (partially ordered cubic Ti(C0.67,Si0.06)) [10]. According to Barsoum et al., the reaction is triggered by
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.10.010
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the chemically induced out-diffusion of Si into the cryolite. As Zhou et al. reports [11], Ti3SiC2 decomposes in the presence of Cu at temperatures above 900 C. Depending on the temperature and the amount of Cu, the formed reaction products are Cu(Si) (solid solution of Si in Cu), TiC, Cu5Si and Cu15Si4. Fe and V impurities of only 1% in hot isostatic pressed Ti3SiC2 lead to a highly porous structure, which reduces the decomposition temperature of the MAX-phase material to below 1600 C [12]. Racault et al. [13] found that polycrystalline Ti3SiC2 bulk material decomposes at 1300 C in vacuum triggered by the chemical reaction between Si and the surrounding C crucible. An alumina crucible in an Ar atmosphere of 100 kPa increases the decomposition temperature to 1450 C [13]. The above-mentioned studies underline the importance of the diffusion of the relatively weakly bonded Si in the decomposition process of Ti3SiC2 and the chemical reactions between the MAX phase and the material in contact. That includes grain boundaries, a high density of dislocations and impurity phases surrounding the MAX-phase grains within polycrystalline bulk material typically produced by hot isostatic pressing [14] and self-propagating high-temperature synthesis [15]. Furthermore, the fact that Ti3SiC2 exhibits two polymorphs in which Si shifts from 2b to 2d (Wyckoff position) at approximately 1000 K [16] supports the notion of Si mobility. Phase-pure and singlecrystal material should therefore constitute suitable choices for studies of the mechanisms of Si out-diffusion and subsequent phase transformations in Ti3SiC2. Here, we report on the thermal stability of (0 0 0 1)oriented single-crystal Ti3SiC2 thin films grown epitaxially on Al2O3(0 0 0 1) wafers using DC magnetron sputtering from three elemental targets. The phase composition and microstructure of the films was studied in situ by X-ray diffraction (XRD) during annealing and ex situ by crosssectional transmission electron microscopy (XTEM) and X-ray photoelectron spectroscopy (XPS). Using direct observations supported by ab initio calculations, we derive a decomposition model for Ti3SiC2. The results show that Ti3SiC2 thin films exhibit surface decomposition in vacuum at 1000 C, but are stable over its bulk thickness to 1100 C, above which they transform into TiCx (x 0.67) upon accelerated Si out-diffusion. Results also show that the removal of Si basal planes in Ti3SiC2 promotes C rearrangement and O uptake likely located at the as-formed Ti3C2 interfaces. The phase transformation process becomes completed at 1300 C with the recrystallization of the TiC.
2. Experimental details For the experiment, 680-nm-thick Ti3SiC2(0 0 0 1) films were grown on a 20-nm-thick TiC(1 1 1) seed layer on Al2O3(0 0 0 1) substrates under ultra-high-vacuum conditions by DC magnetron sputtering from three singleelement targets of Ti, Si and C, respectively, in an Ar atmo-
sphere. Detailed information about the deposition and growth processes can be found elsewhere [17]. The XRD measurements were carried out in a Philips X’pert MPD XRD system (theta/theta). As-deposited films were annealed in vacuum at room temperature (RT) at a pressure of 3 · 107 mbar in situ. The diffractometer used a water-cooled furnace equipped with a Be entrance window. The studied temperature range was 800–1400 C. Heating from room temperature to 800 C was carried out in steps of 50 C in 20 s time intervals. At 800 C the annealing temperature, Ta was kept constant for 1 h and then increased in steps of 100 C with a 1 h holding period after each step, during which four cycles of h–2h scans were conducted between the angles 33 and 44. The reason for the selected scan range was the proximity of the following high-intensive film and substrate peaks: Ti3SiC20 0 0 8; TiC1 1 1; Ti5Si3Cx0 0 0 2; Al2O30 0 0 6. Prior to each h–2h scan, 2h and x optimization scans around 2h = 35.5 (approximate peak position of TiC1 1 1) were performed. The samples were placed on a resistively heated Ta filament providing a chemically inert environment. On the backside of the filament a calibrated Pt/Rh thermocouple was clamped to monitor the temperature. A second Ta filament surrounding the sample in a 2 cm radius with a similar thermocouple was used to prevent temperature gradients across the sample. Space constraints made it necessary to cut the samples into dimensions of 4 · 5 mm performed with a diamond wheel saw. Epitaxial 2.5-lm-thick Ti3SiC2(0 0 0 1) films were put mechanically face-to-face and annealed in an alumina crucible using a Netzsch STA 409C thermal analysis instrument. The as-deposited samples were heated up to 1400 C with a heating rate of 10 K/min in an Ar atmosphere at 1 bar. Calibrated Pt/Rh thermocouples at the top and bottom of the crucible were used for temperature measurement. Bonding energies of the elements were obtained through X-ray photoelectron spectroscopy measurements using a PHI Quantum instrument with monochromatic Al Ka radiation. The peak areas in the spectra were weighed with internal sensitivity factors in order to acquire elemental film compositions. Ar-ion sputter cleaning of the samples was done for 100 s, within which sputtering was interrupted in regular intervals to carry out spectral measurements to obtain depth profiles for every sample. Measurements by XPS showed that as-deposited Ti3SiC2 films contain 52% Ti, 13% Si and 35% C. A composition of 49% Ti, 36% C and 15% of Si for the as-deposited films was measured with elastic recoil detection analysis using high-energetic (35 MeV) Cl7+ ions with an incidence angle of 15. Considering the measurement error limits and the limited compositional stability range of Ti3SiC2, we conclude that the films used in the present experiments are effectively stoichiometric. Microstructural analysis on the samples was carried out using transmission electron microscopy using a FEI Tecnai G2 TF 20 UT 200 kV FEG microscope. All samples were
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thinned to 60 lm through grinding with diamond paper and subsequently ion milled to electron transparency at 5 keV with a Gatan precision ion polishing system (PIPS). The final polishing step was performed at 2.5 keV. Lowand high-resolution TEM images were obtained in microprobe mode, whereas scanning transmission electron microscopy (STEM) imaging was performed in nanoprobe ˚ using a high-angle mode with a beam spot size of 5 A annular dark-field (HAADF) detector with variable camera length. Compositional analysis in STEM mode was done by energy-dispersive X-ray (EDX) analysis in STEM mode with applied drift correction. 3. Theoretical methods Ab initio calculations were carried out using density functional theory [18], as implemented in the Vienna ab initio simulation program, wherein projector augmented wave potentials are employed [19]. The generalized-gradient approximation was applied in all calculations, both electronic and phonon related, with the so-called Blo¨chl corrections for the total energy [20]. The integration in the Brillouin zone is done on special k-points, which is determined after Monkhorst and Pack [21]. All configurations were studied on a mesh of 7 · 7 · 7 irreducible k-points. The convergence criterion for the total energy was set to 0.01 meV within a 500 eV cut-off and the convergence test with respect to the configuration size was performed on 1 · 1 · 2 and 2 · 2 · 1 supercells (0.58 meV/atom energy difference was obtained). All configurations were relaxed with respect to atomic positions (internal free parameters), lattice parameter a and hexagonal c/a ratio. After relaxing the atomic positions and volumes, the cohesive energy was calculated from the obtained minimum configurations. We used the difference in cohesive energy with respect to Ti3SiC2 to estimate the stability of the configuration probed. Ti3SiC2, Ti3C2/Ti3C2 (Si removed from Ti3SiC2) and Ti3C2/O/Ti3C2 (O at Si site in Ti3SiC2) were treated with the unit cell containing 12 atoms, while Ti3C2was modelled with 24 atoms (1 · 2 · 1 supercell). Ti3C2 designates cubic TiC0.67 containing 33% of C-vacancies and it was described with six layers of TiC(1 1 1) stacked along [0 0 0 1] hexagonal direction. For the calculations of TiC0.67, four C vacancy positions were randomly chosen and the energy difference between two configurations probed was 11 meV/atom.
Fig. 1. X-ray diffractogram from an as-deposited Ti3SiC2(0 0 0 1) film grown on Al2O3(0 0 0 1) with a 20-nm-thick TiC seed layer.
The related Ti5Si3Cx(0 0 0 1) was observed as a minority phase. Fig. 2 shows the XRD peak intensity evolution over annealing time, ta, for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray), and Ti5Si3Cx0 0 0 2 (light gray) peaks from the Ti3SiC2(0 0 0 1) film. The sample was annealed at 800 C for ta = 1 h, after which the annealing temperature, Ta, was increased in 100 C steps with 1 h annealing at constant temperature after each step. It can be seen in Fig. 2 that the intensity of Ti3SiC20 0 0 8 decreases slightly with Ta increasing up to 1100 C. The pronounced decrease in the MAX-phase-peak intensity during 1 h annealing at Ta = 1200 C shows that decomposition of the Ti3SiC2 phase takes place at 1100 6 Ta 6 1200 C. The TiC1 1 1 peak exhibits a constant intensity between 800 and 1000 C. Above this temperature, however, the peak gradually gains intensity until Ta = 1300 C, where
4. Results and discussion 4.1. X-ray diffraction Fig. 1 is an X-ray diffractogram from an as-deposited Ti3SiC2(0 0 0 1) film. The exclusive presence of 0 0 0 ‘ indices confirm the epitaxial growth of the MAX phase as expected from the chosen growth conditions [17]. The presence of a TiC1 1 1 peak is due to the seed layer being deposited first.
Fig. 2. XRD peak intensity evolution over time for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray) and Ti5Si3Cx0 0 0 2 (light gray) peaks during annealing of a Ti3SiC2(0 0 0 1) film between 800 and 1300 C with temperature steps of 100 C and a 1 h holding period after each step.
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the intensity increase becomes definite. The intensity of the Si-rich Ti5Si3Cx phase remains low and constant throughout the annealing process. In order to provide insight into the dynamics of the decomposition process of Ti3SiC2, the temperature steps, DTa, were decreased to 20 C with the annealing temperature in each step kept constant for ta = 1 h. The XRDintensity evolution for this experiment is shown in Fig. 3. Between 1100 and 1160 C, the increase of the Ti3SiC20 0 0 8-peak intensity can be interpreted as increased crystalline quality due to annihilation of growth defects in the film. At Ta = 1180 C the Ti3SiC20 0 0 8 intensity decreases exponentially with an initial rate of DI/Dta(Ti3SiC2) of 19 counts/min, while TiC1 1 1 rises continuously at a rate of DI/Dta(TiC) of 0.2 counts/min. At the onset of Ti3SiC2 decomposition, Ta was kept constant for 20 h until the Ti3SiC20 0 0 8 peak was no longer detected. After that, Ta was ramped up to 1300 C, where the TiC1 1 1 intensity rises distinctively with a DI/Dta(TiC) of 7 counts/min (see Fig. 3). This indicates that TiC is growing as a decomposition product of Ti3SiC2. Again, no increase of the Ti5Si3Cx0 0 0 2 peak intensity was observed during the annealing experiment. A scan over 25 h was also performed at 1100 C (not shown). After a ta of 200 min, a steady decrease in the Ti3SiC20 0 0 8-peak intensity at a rate of DI/Dta of 0.8 counts/min was seen, indicating that the decomposition of Ti3SiC2 was activated. The rate of the change of the peak intensity in the XRD diffractograms is an indicator of the decomposition kinetics. The annealing experiments at Ta = 1100 C showed a very slow decomposition kinetics. After ta = 25 h more than 90% of the Ti3SiC2 film was still intact, as shown by TEM below. With a temperature increase of 80 C DI/Dta(Ti3SiC2) has drastically decreased, from 0.8 to 19 counts/min. This shows that the Ti3SiC2 decomposition is thermally activated.
Fig. 4. Time evolution of the X-ray FWHM for the TiC1 1 1 peak during annealing of a Ti3SiC2(0 0 0 1) film between 1100 and 1180 C with temperature steps of 20 C and 1 h holding periods after each step. After a 20 h holding period at 1180 C the temperature was increased to 1300 and 1400 C, respectively, each for 1 h.
The XRD full width at half maximum (FWHM) values for the TiC1 1 1 peak is shown in Fig. 4 vs. ta and Ta. There is a small decrease in the FWHM from 0.26 to 0.24 during 25 h at 1180 C. This is followed by a distinct decrease of FWHM to 0.2 within 1 h as Ta is raised to 1300 C. Since the FWHM value is an inverse measure of grain size, the observed drop in FWHM indicates that TiC layers recrystallize subsequent to the MAX-phase decomposition. 4.2. X-ray photoelectron spectroscopy The O1s peak at a binding energy of 532.3 eV and the Si2p peak at 103 eV indicate that the surface is mainly terminated by SiO2 [22,23]. After sputter cleaning the surface, O was no longer detected. The Si binding energy of the
Intensity [a. u.]
Si 2p
sputter cleaned for 30 s
T a = 1300 ˚C for 1 h
Fig. 3. XRD peak intensity evolution over time for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray) and Ti5Si3Cx0 0 0 2 (light gray) peaks during annealing of a Ti3SiC2(0 0 0 1) film between 1100 and 1180 C with temperature steps of 20 C and 1 h holding periods after each step. After 20 h holding period at 1180 C the temperature was increased to 1300 and 1400 C, respectively, each for 1 h.
108
106
104 102 100 Binding Energy [eV]
98
96
Fig. 5. X-ray photo-electron spectroscopy spectra showing the region of the Si 2p binding energy from the as-deposited Ti3SiC2(0 0 0 1) sample and a film annealed from 800 to 1300 C in steps of 100 C held at each temperature level for 1 h.
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as-deposited film of 99.2 eV, shown in Fig. 5, corresponds to the typical Si bonding environment in Ti3SiC2 [24]. Measurements on samples annealed at 1300 C and above showed no traces of Si in the film (see Fig. 5), but an O content [O] > 9%. XPS measurements on the sample annealed at 1100 C for 25 h showed a Si gradient into the film. The top-few nanometers did not contain Si, whereas the Si content increased noticeably after depth profiling to 50– 60 nm to the nominal 1/3 Si/Ti ratio. At the same time, an O diffusion profile was detected over the same depth with a significant 8% content at a depth of 30 nm. With the absence of Si, the surface is composed of TiO2 (530.1 eV) [25], with a large number of C–O bonds in the presence of Ti (531.1 eV) [26]. The peak attributed to TiO2 vanishes after 10 s sputter cleaning. 4.3. Transmission electron microscopy Fig. 6 is an XTEM micrograph from the sample annealed to 1100 C for 25 h. A 3–4-nm-thick TiO2 oxide layer evidenced by XPS can be seen on the surface of the sample. This is a thin protective oxide layer (too thin to attenuate X-rays in the above experiments) typically observed on Ti3SiC2 material [27,28]. Underneath the oxide layer, the MAX phase has decomposed and the extension of a TiC0.67 (nominal composition) layer can be seen to a depth of 60 nm. There are two cubic orientations found in this layer (see the Fourier transform inset in Fig. 6), which are twinned along ½1 2 1. In the XTEM image (Fig. 6), the twinned regions are separated by planar faults
Fig. 6. Cross-sectional TEM micrograph showing a Ti3SiC2(0 0 0 1) film annealed at 1100 C for 25 h with a Fourier transform from the TiC0.67 layer decomposed from Ti3SiC2. Twinned orientations along ½1 2 1 (horizontal line) can be seen distinguished by the prime for the twinned orientation.
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(PF) (see arrowheads in Fig. 6). No Si islands or other deposits of Si were found in the transformation layer, which is consistent with the XRD and XPS results above. The bulk of the sample (lower part of Fig. 6) consists of Ti3SiC2. Fig. 7 shows an STEM image from the Ti3SiC2 film annealed at 1100 C for 25 h. The TiC0.67 transformation layer, Ti3SiC2 and TiC seed layer are well resolved in bright contrast. Also, the PFs are seen in the TiC0.67 layer indicated by the arrowheads in Fig. 7. In addition, plate-like voids (magnified in inset in Fig. 7) parallel to the (1 1 1) planes are present in this layer. An EDX line scan performed across the TiC0.67 layer into the MAX-phase region is also presented in Fig. 7 (corresponding to the S–F line in the STEM image). The variation in elemental concentration in the inserted graph confirms the phase identification made above. A vertical dashed line in the graph depicts the interface between the transformation layer and the unreacted MAX phase. A residual amount of Si can be found in the transformation layer as well as a high O concentration, which decreases as the Si concentration increases with depth into the film. The planar faults in the transformation layer discussed above are likely locations for residual Si, as seen in XPS, in the form of locally retained Si layers in the decomposed Ti3SiC2 film. For higher annealing temperatures, the XRD analysis showed an accelerated MAX-phase decomposition. The XTEM micrograph in Fig. 8a confirms that the Ti3SiC2 film has completely transformed into TiC0.67 at Ta = 1200 C. The as-transformed layer, however, contains yet larger voids compared with the transformed volume at Ta = 1100 C (confer inset in Fig. 7). Near the surface the
Fig. 7. Scanning TEM image in nanoprobe mode with a high-angle annular dark-field detector with a long camera length showing a Ti3SiC2(0 0 0 1) film annealed at 1100 C for 25 h with two voids magnified in the inset. Elemental concentrations from the energy dispersive X-ray line scan along SF are shown in the graph inserted at the lower right.
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Fig. 8. Cross-sectional TEM micrograph (a) showing a Ti3SiC2(0 0 0 1) film containing transformed TiC0.67 layers after annealing to 1200 C for 1 h, and the corresponding selected area diffraction pattern (b) including Ti5Si3Cx (S), the Al2O3(0 0 0 1) substrate (A) and twinned TiC0.67 (T, T 0 ) plus double diffraction spots (D).
voids have coarsened and make contact with the surface at places. Thus, the initially smooth Ti3SiC2(0 0 0 1) surface (see Figs. 6 and 7) has become uneven, with low-index TiC0.67 crystal facet surfaces. As a consequence, there are additional O-diffusion channels, which may in part explain the large O concentration observed in the annealed films. A typical selected area electron diffraction (SAED) pattern from the sample annealed to 1200 C is shown in Fig. 8b. The phases present are TiC0.67, (T), Ti5Si3Cx, (S) and Al2O3 (substrate, denoted A). The crystallographical relationships determined from the SAED pattern are TiC0.67(1 1 1)//Al2O3(0 0 0 2) and TiC0:67 ½1 0 1==Al2 O3 3 2Þ and TiC0:67 ½1 0 1 0, and TiC0:67 ð1 1 1Þ==Ti5 Si3 Cx ð21 ½1 0 1==Ti5 Si3 Cx ½ 211 3, respectively. Spots denoted T 0 in Fig. 8b are caused by transformed TiC0.67 layers, which are twinned at the PFs. This instigates twin Moire´ patterns at overlapping grains in twin relationships resulting in double diffraction spots (denoted D) in the SAED pattern in Fig. 8b. Fig. 9 shows an HREM image of the transformed layer in a Ti3SiC2 film annealed to 1300 C. The transformed TiC0.67 layer again contains Ti5Si3Cx inclusions, in this case with the alternative crystallographic relationship to TiC0.67 of TiC0.67(1 1 1)//Ti5Si3Cx(0 0 0 2) and TiC0:67 ½1 0 1== Ti5 Si3 Cx ½ 1 0 1 0 also observed in the XRD measurements (Fig. 1). We deduce that this is the predominant crystallographic orientational relationship, due to the larger volume of film material probed in the XRD analysis compared with the electron diffraction analysis above. Initial annealing experiments with subsequent XRD measurements of the Ti3SiC2(0 0 0 1) films mechanically face-to-face showed that the MAX phase was stable to at least 1400 C. This is in agreement with the literature data for high-temperature stability of bulk Ti3SiC2 [9]. The observed increase in decomposition temperature for the samples put mechanically face-to-face compared with the vacuum-annealed samples can be attributed to the absence of an interface to vacuum as well as the Si entrapment caused by the initial formation of an oxide at the perimeter of the contacting films. These results indicate that the thermal stability of the MAX phase can be increased by preventing Si removal from the surface.
Scrutiny of the films annealed to 1200–1300 C reveals numerous PFs in the transformed layer, which are often arranged in a layered structure as in Ti3SiC2, though without any long-range order. Typical PF features can be seen in Fig. 10a–c, which are HREM images from a sample annealed to Ta = 1300 C. Fig. 10a reveals a set of three periodically spaced PFs. Their separation distance is ˚ , which is 2.3 A ˚ smaller than the unit cell of the 14.5 A MAX phase presented in Fig. 11. This can be explained if Ti3C2 slabs relax during the removal of the Si basal plane. A magnification of region A (Fig. 10b) shows sets ˚ . Also, inspecof atomic layers with a separation of 1.7 A tion of region B in Fig. 10a reveals a larger plane separa-
Fig. 9. High-resolution image from a Ti3SiC2(0 0 0 1) film annealed to 1300 C for 1 h showing a Ti5Si3Cx(0 0 0 2) grain embedded in the TiC0.67(1 1 1) transformed layer.
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˚ and Fig. 11. Ti3SiC2 unit cell with cell parameters of a = 3.0749 A ˚ , including the distance dI = 3.86 A ˚ between Ti(I) layers (Ti c = 17.7574 A ˚ between Ti(I) and Ti(II) layers (Ti not adjacent to Si) and dII = 2.49 A bound to Si; dII is similar to distance between the C layers).
4.4. Decomposition model Fig. 12a–d is a schematic illustration of our proposed mechanism of Ti3SiC2 decomposition based on the above results from thermal vacuum annealing experiments.
Fig. 10. Cross-sectional high resolution TEM images from a Ti3SiC2(0 0 0 1) film after phase transformation into TiC0.67(1 1 1) during annealing at 1300 C showing (a) an overview with different planar faults of types A and B, presented at higher magnification in (b) and (c), respectively.
tion and stacking sequence corresponding to a layer of locally retained Ti3SiC2 phase (marked by an arrow). A HREM micrograph of that area is shown in Fig. 10c. The distance between the layers on each side of the mirror ˚ . This corplane, marked with an arrow in Fig. 10c, is 4 A ˚ responds to the distance dI = 3.9 A (Ti(I)–Si–Ti(I); see unit cell in Fig. 11) in Ti3SiC2. It can thus be inferred that these mirror planes are occupied by Si as a residual part of locally retained Ti3SiC2.
4.4.1. Stage I: Si out-diffusion and evaporation We find that the Ti3SiC2 decomposition is initiated by out-diffusion and entropy-driven evaporation of Si from the top (0 0 0 1) film surface toward vacuum during annealing at Ta P 1100, and possibly even at 1000 C (see Fig. 2) on a nanometer depth scale. For 25 h annealing at 1100 C, the MAX phase is decomposed to a depth of 60 nm. Our recent in situ photoemission studies [28] of Ti3SiC2(0 0 0 1) thin films annealed in ultra-high vacuum revealed, with high surface sensitivity, Si enrichment on the free surface at 1000 C. The Si evaporation at the free film surface takes place since the Si vapor pressure of 8 · 106 mbar at 1200 C [29] is of the same order of magnitude as the pressure range used during annealing. The Si out-diffusion can be explained by its bond strength. Theoretical calculations of Zhou and Sun [30] show that Si is partially metallically bonded to Ti(I) (Ti atoms adjacent to Si) and thus the weakest bonded element in the Ti3SiC2 phase compared with strongly covalently and to some extent ionically bonded Ti(II) (Ti bonded to C but not Si; see Fig. 11) and C. Also, the Ti(I) to Ti(I) distance with Si mirror planes in between is the largest distance found in the Ti3SiC2 phase and allows Si to diffuse easier than C, which is smaller, but placed at interstitial sites. In addition, Barsoum et al. [31] observed Si displacement about its atomic equilibrium position at 900 C preferentially along the a-axis, due to relatively weak Si–Si bonds, which alleviates the Si diffusion. The diffusing Si
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Fig. 12. Schematic illustration describing the different stages of phase transformations occurring during the decomposition of Ti3SiC2(0 0 0 1).
atoms are likely to emerge on the Ti3SiC2(0 0 0 1) surface via domain boundaries or other defects in the film as well as from surface steps. Our earlier work [17] using atomic force microscopy showed that the Ti3SiC2(0 0 0 1) surface exhibits a step-like topography with step heights of full and half unit cells from a helicoidal growth mode around threading screw dislocations. At the side of each step, the Si planes are exposed to the ambient. We infer that Ti3SiC2 material exposing surfaces of other crystallographic orientations are likely to be even more prone to release Si.
4.4.2. Stage II: O uptake and SiO evaporation In the presence of an oxygen ambient (residual gas in the evacuated furnace in the present case), the Si-depleted highly defective structure formed under Stage I (see Fig. 12a) is liable to O in-diffusion. A large uptake of O (up to 10%) in our films was detected by XPS. During its way into the Ti3SiC2 film, O is likely to react with Si and form SiO gaseous species [13], thus accelerating the decomposition process as illustrated in Fig. 12b. O may eventually be incorporated in the material, e.g., by saturating
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the dangling bonds created by the Si-deficiency and thus creating O mirror planes between Ti3C2 slabs (Ti3C2/O/ Ti3C2).
Table 1 Calculated unit cell dimensions and cohesive energies for Ti3SiC2 and decomposition products ˚) a (A c/a DE (eV/atom)
4.4.3. Stage III: Ti3C2 relaxation, detwinning, and TiC0.67 formation by C-redistribution with void formation Fig. 12c shows the next stage in the decomposition process, which is activated at Ta P 1100 C. Here, the twinned Ti3C2 slabs (Ti3C2/Ti3C2) formed in Stages I and II relax towards each other. The transformation of the MAX-phase stacking is accompanied by a large reduction of volume as the Si-depleted hexagonal unit cell shrinks with or without the presence of O. A film entirely composed of a twinned Ti3C2 structure may give rise to superstructure peaks comparable to the 0 0 0 ‘ peaks in Ti3SiC2, but at slightly higher 2h angles considering the shrinkage of the unit cell. We find, however, no superstructure of the slabs in their repeated T and T 0 twin ordering. Also, it cannot be excluded based on these initial experiments that the Ti3C2T and T 0 slab relaxation is simultaneous with detwinning and recrystallization (see Stage IV below). This can be understood by considering the almost constant TiC1 1 1-peak intensity during 20 h annealing at 1180 C, as exemplified by Fig. 3 for Ta = 1180 C and ta = 20 h, which suggests – in such a case – that the coherent volume of the Ti3C2T/T 0 is too small to give rise to constructive interference during XRD measurements. The relaxation process includes also C redistribution from within the Ti3C2 slabs to the Si-depleted mirror planes. This process is in agreement with the known stability regime of TiCx (x = 0.97–0.47 [32,33]). Thus, homogenization in the C content to form TiC0.67 is what can be expected.
Ti3SiC2 Ti3iC2/Ti3iC2 recovery in z-direction Ti3iC2/Ti3iC2 recovery in x–y plane Ti3iC2/O/Ti3SiC2 TiC0.67
4.4.4. Stage IV: TiC0.67 layer growth and recrystallization From the XTEM study and the increase of TiC1 1 1peak intensity as well as the decrease in FWHM for TiC1 1 1 at Ta P 1300 C it can be concluded that the transformed layer undergoes both growth and recrystallization within the film as the final stage of the Ti3SiC2 decomposition process (see Fig. 12d). This is accompanied by an increased pore size observed during annealing between Ta = 1200 C and 1300 C (see Fig. 8a). Residual mirror planes containing Si as well as C and/or O may be present even up to Ta = 1300 C, as indicated by the large number of planar faults with different distances to their neighbor planes (see Fig. 10a–c). This may be the effect of local crystal perfection with no available diffusion paths for the Si. Note that the annealing temperatures used are 30% of the decomposition temperature of TiC [32]. Thus, only limited bulk diffusion in the TiC0.67 layers is expected. The recrystallization may instead be assisted by the enthalpy released in the above described reactions as well as by the relatively high planar-fault density and lattice point defects. The latter are, e.g., C vacancies and O interstitials, which eventually migrate to form the voids, but also any residual Si substitutional species. Clearly, the
3.075 3.078 3.087 3.062 3.066
5.775 4.934 4.845 4.967 4.894
0.000 0.352 0.377 0.564 0.392
chemical potential of Si in the various environments will determine the effective Ti3SiC2 stability. 4.5. DFT calculations Ab initio calculations were used to test our decomposition model of Ti3SiC2 in terms of the stability of the decomposition products, while kinetic and dynamic effects cannot be unraveled from these results. Table 1 contains the structural and stability data for Ti3SiC2, Ti3C2/Ti3C2, Ti3C2/O/ Ti3C2 and TiC0.67. Our ab initio data are consistent with the literature: the calculated equilibrium volume of Ti3SiC2 ˚ 3/atom, which is 0.9% larger than the experimenis 12.117 A ˚ 3/atom [1]. Furthermore, the calculated tal value of 12.003 A bulk modulus of Ti3SiC2 using the Birch–Murnaghan equation of states [34] is 209 GPa, which is 1.4% larger than the experimental value of 206 GPa [35]. In terms of the structure of the Ti3SiC2 decomposition products, it can generally be noted that the c/a ratio decreases upon decomposition, which is also consistent with our experimental data presented above. The unit cell volume of Ti3C2/Ti3C2 and Ti3C2/O/Ti3C2 is 15% smaller compared with Ti3SiC2, which is in agreement with the observations in Ref. [10]. The removal of Si layers from the Ti3SiC2 structure (Stage I in Fig. 12a) is associated with a decrease in the cohesive energy by 0.352 and 0.377 eV/atom, when the recovery is allowed in the z-direction and in the x/y plane, respectively. After the Si layer has been removed, two possible decomposition routes are probed. First, in the absence of oxygen, twins in Ti3C2 may be annealed out and C redistribution by short-range diffusion may take place, resulting in the formation of a TiC0.67 structure (Stage III). This is associated with the decrease in the cohesive energy by 0.392 eV/atom with respect to Ti3SiC2. Secondly, if oxygen is introduced into the remaining structure (Stage II), the cohesive energy decreases by 0.564 eV/atom with respect to Ti3SiC2. This stability increase suggests that oxidation is likely to occur in Ti3SiC2 upon removal of Si in an O-containing ambient. Therefore, the ab initio calculations support our decomposition model. 5. Conclusions We find that Ti3SiC2(0 0 0 1) thin films are effectively stable during vacuum furnace annealing up to 1100 C, above which they decompose into TiCx (x 0.67) following Si out-diffusion and evaporation. A sequence of phase
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transformation steps is determined, supported by DFT calculations, comprising relaxation and detwinning of Ti3C2 slabs to form TiC0.67, which undergoes recrystallization at 1200 6 Ta 6 1300 C. The Si out-diffusion may be accelerated by an O ambient through reactions to form gaseous SiO species. Thus, decomposition initiated at the surface on a nanometer scale is observed already at 1000 C. The transformed TiC0.67 layers contain voids and exhibit rough and facetted surfaces due to the involved phase transformations taking place under conditions of high lattice point defect content. The observed onset for the Ti3SiC2 decomposition initiated at the surface over a 10–60 nm depth at Ta = 1000–1100 C is in apparent contrast to the reported decomposition temperatures for bulk material of 1800– 2300 C [8,36]. We conclude that the discrepancy in the decomposition temperature is in part a result of the difference in detection sensitivity or chosen thickness criteria for a given transformation process between the studies. The role of the chemical environment is as important to the decomposition activation of Ti3SiC2 [10,13] as it is the activity of the A-element Si that is found to be the predominant factor in the decomposition of Ti3SiC2 and likely other MAX phases. Acknowledgements This work was supported by the Swedish Foundation of Strategic Research (SSF), Strategic Research Centre in Materials Science for Nanoscale Surface Engineering, the VINNOVA project on Industrialization of MAX-Phase Coatings, ABB Ltd. and Kanthal AB. Herbert Willmann is gratefully acknowledged for the measurements performed on the Netzsch thermal analysis instrument. The authors acknowledge Assoc. Prof. Jens Birch for valuable discussion as well as Dr. Ulrich Kreissig for performing ERDA measurements. JMS gratefully acknowledges support from Deutsche Forschungsgemeinschaft (DFG) (Schn 735/9-1). References [1] Jeitschko W, Nowotny H. Mh Chem 1967;98:329. [2] Barsoum MW. Prog Solid St Chem 2000;28:201.
[3] Ahuja R, Eriksson O, Wills JM, Johansson B. Appl Phys Lett 2000;76:2226. [4] Tuller, HL, Spears, MA, Mlcak, R. United States Patent No. US 2002/0 068 488 A1, 2002. [5] Sundberg, M, Lindgren, KJ, El-Raghy, T, Malmqvist, G. International Patent No. WO 2004/069 745 A1, 2004. [6] Wakelkamp WJJ, van Loo FJJ, Metselaar R. J Eur Ceram Soc 1991;8:135. [7] Witthaut M, Weiß R, Zimmermann E, von Richthofen A, Neuschu¨tz D. Z Metallkd 1998;89:9. [8] Radhakrishnan R, Williams JJ, Akinc M. J All Comp 1999;285:85. [9] Barsoum MW, El-Raghy T, Radovic M. Interceram 2000;49: 226. [10] Barsoum MW, El-Raghy T, Farber L, Amer M, Christini R, Adams A. J Electrochem Soc 1999;146:3919. [11] Zhou Y, Gu W. Z Metallkd 2004;95:50. [12] Tzenov N, Barsoum MW, El-Raghy T. J Eur Ceram Soc 2000;20:801. [13] Racault C, Langlais F, Naslain R. J Mater Sci 1994;29:3384. [14] Barsoum MW, El-Raghy T. J Am Ceram Soc 1996;79:1953. [15] Riley DP, Kisi EH, Hansen TC, Hewat AW. J Am Ceram Soc 2002;85:2417. [16] Sun Z, Zhou J, Music D, Ahuja R, Schneider JM. Scripta Mater 2006;54:105. [17] Emmerlich J, Palmquist J-P, Cziga´ny Zs, Ho¨gberg H, Szasva´ri Sz, Molina-Aldareguia JM, et al. J Appl Phys 2004;69:4817. [18] Hohenberg P, Kohn W. Phys Rev 1964;136:B864. [19] Kresse G, Joubert D. Phys Rev B 1999;59:1758. [20] Blo¨chl PE. Phys Rev B 1994;50:17 953. [21] Monkhorst HJ, Pack JD. Phys Rev B 1976;13:5188. [22] Sprenger D, Bach H, Meisel W, Gutlich P. J Non-cryst Solids 1990;126:111. [23] Laczka M, Beier W, Stoch L. Glastechn Berichte 1989;62:320. [24] Stoltz SE, Starnberg HI, Barsoum MW. J Phys Chem Solids 2003;64:2321. [25] Nefedov VI, Gati D, Dzhurinskii BF, Sergushin NP, Salyn YV. Zh Neorg Khimii 1975;20:2307. [26] Rajumon MK, Hegde MS, Rao CNR. Catal Lett 1988;1:351. [27] Kisi EH, Crossley JAA, Myhra S, Barsoum MW. J Phys Chem Solids 1998;59:1437. [28] Eklund P, Virojanadara C, Emmerlich J, Ho¨gberg H, Hultman L, Johansson LI. Phys Rev B 2006;74:045417. [29] Honig RE, Kramer DA. RCA Rev 1969;30:285. [30] Zhou Y, Sun Z. J Phys: Condens Matter 1969;12:L457. [31] Barsoum MW, El-Raghy T, Rawn CJ, Porter WD, Wang H, Payzant EA, et al. J Phys Chem Solids 1999;60:429. [32] Pierson HO. Handbook of refractory carbides and nitrides. Westwood, NJ: Noyes Publications; 1996. [33] Zueva LV, Gusev AI. Phys Solid State 1999;41:1032. [34] Birch F. J Geophys Res 1978;83:1257. [35] Onodera A, Hirano H, Yuasa T, Gao NF, Miyamoto Y. Appl Phys Lett 1999;74:3782. [36] Du Y, Schuster JC, Seifert H, Aldinger F. J Am Ceram Soc 2000;83:197.
Thermal stability of Ti3SiC2 thin films Jens Emmerlich a,*, Denis Music b, Per Eklund a, Ola Wilhelmsson c, Ulf Jansson c, Jochen M. Schneider b, Hans Ho¨gberg a, Lars Hultman a a
c
Linko¨ping University, Department of Physics, IFM, Thin Film Physics Division, SE-581 83 Linko¨ping, Sweden b Materials Chemistry, RWTH Aachen University, Kopernikusstraße 16, D-52074 Aachen, Germany Uppsala University, Department of Materials Chemistry, The Angstro ¨ m Laboratory, P.O. Box 538, SE-751 21 Uppsala, Sweden Received 26 May 2006; received in revised form 4 October 2006; accepted 5 October 2006 Available online 18 December 2006
Abstract The thermal stability of Ti3SiC2(0 0 0 1) thin films is studied by in situ X-ray diffraction analysis during vacuum furnace annealing in combination with X-ray photoelectron spectroscopy, transmission electron microscopy and scanning transmission electron microscopy with energy dispersive X-ray analysis. The films are found to be stable during annealing at temperatures up to 1000 C for 25 h. Annealing at 1100–1200 C results in the rapid decomposition of Ti3SiC2 by Si out-diffusion along the basal planes via domain boundaries to the free surface with subsequent evaporation. As a consequence, the material shrinks by the relaxation of the Ti3C2 slabs and, it is proposed, by an in-diffusion of O into the empty Si-mirror planes. The phase transformation process is followed by the detwinning of the as-relaxed Ti3C2 slabs into (1 1 1)-oriented TiC0.67 layers, which begin recrystallizing at 1300 C. Ab initio calculations are provided supporting the presented decomposition mechanisms. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ti3SiC2 thin films; Phase transformations; X-ray diffraction; Transmission electron microscopy; Ab initio electron theory
1. Introduction Ti3SiC2 is a Mn+1AXn phase (n = 1–3), where M = early transition metal, A = an A-group element (usually group 13–14), and X = C or N. The MAX phases exhibit a hexagonal structure (space group P63/mmc). In Ti3SiC2, Ti3C2 layers (isostructural to the binary TiC phase, where C is at octahedral sites) are intercalated with atomic layers of Si, which act as mirror planes [1]. The Si is weakly bonded to the Ti, whereas the covalent Ti–C bonds are much stronger [2,3]. MAX phases in general exhibit a high thermal stability, but little is known about the phase transformation pathways at temperatures above the equilibrium temperature range of the different MAX phases. In particular, there is a need to determine the potential phase decomposition in *
Corresponding author. Tel.: +46 13 281281. E-mail address: [email protected] (J. Emmerlich).
applications where the MAX phase material is in contact with another chemical medium. For example, there is a growing interest in thin films of Ti3SiC2 as electrical contacts [4]. For high-power applications, the thermal stability of the MAX phase is an issue as high thermal loads in the contact area to metals like Cu and Ni or to air can be expected. Furthermore, some of the phases are suitable for heating elements [5] and/or as protective coatings, in which case the material is subject to a wide range of interfaces and environments. Considering the Ti–Si–C phase diagrams at 1100 and 1250 C [6], Ti3SiC2 is in equilibrium with TiC or possible (Ti,Si)C [7]. Tie-lines can be observed to the Si-rich phases of SiC, TiSi2 and Ti5Si3. Ti3SiC2 bulk material was found to be thermally stable in Ar atmosphere to at least 1800 C [8,9]. Immersing Ti3SiC2 in molten cryolite at 960 C, however, transforms the phase topotactically into TiC with some retained Si (partially ordered cubic Ti(C0.67,Si0.06)) [10]. According to Barsoum et al., the reaction is triggered by
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.10.010
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the chemically induced out-diffusion of Si into the cryolite. As Zhou et al. reports [11], Ti3SiC2 decomposes in the presence of Cu at temperatures above 900 C. Depending on the temperature and the amount of Cu, the formed reaction products are Cu(Si) (solid solution of Si in Cu), TiC, Cu5Si and Cu15Si4. Fe and V impurities of only 1% in hot isostatic pressed Ti3SiC2 lead to a highly porous structure, which reduces the decomposition temperature of the MAX-phase material to below 1600 C [12]. Racault et al. [13] found that polycrystalline Ti3SiC2 bulk material decomposes at 1300 C in vacuum triggered by the chemical reaction between Si and the surrounding C crucible. An alumina crucible in an Ar atmosphere of 100 kPa increases the decomposition temperature to 1450 C [13]. The above-mentioned studies underline the importance of the diffusion of the relatively weakly bonded Si in the decomposition process of Ti3SiC2 and the chemical reactions between the MAX phase and the material in contact. That includes grain boundaries, a high density of dislocations and impurity phases surrounding the MAX-phase grains within polycrystalline bulk material typically produced by hot isostatic pressing [14] and self-propagating high-temperature synthesis [15]. Furthermore, the fact that Ti3SiC2 exhibits two polymorphs in which Si shifts from 2b to 2d (Wyckoff position) at approximately 1000 K [16] supports the notion of Si mobility. Phase-pure and singlecrystal material should therefore constitute suitable choices for studies of the mechanisms of Si out-diffusion and subsequent phase transformations in Ti3SiC2. Here, we report on the thermal stability of (0 0 0 1)oriented single-crystal Ti3SiC2 thin films grown epitaxially on Al2O3(0 0 0 1) wafers using DC magnetron sputtering from three elemental targets. The phase composition and microstructure of the films was studied in situ by X-ray diffraction (XRD) during annealing and ex situ by crosssectional transmission electron microscopy (XTEM) and X-ray photoelectron spectroscopy (XPS). Using direct observations supported by ab initio calculations, we derive a decomposition model for Ti3SiC2. The results show that Ti3SiC2 thin films exhibit surface decomposition in vacuum at 1000 C, but are stable over its bulk thickness to 1100 C, above which they transform into TiCx (x 0.67) upon accelerated Si out-diffusion. Results also show that the removal of Si basal planes in Ti3SiC2 promotes C rearrangement and O uptake likely located at the as-formed Ti3C2 interfaces. The phase transformation process becomes completed at 1300 C with the recrystallization of the TiC.
2. Experimental details For the experiment, 680-nm-thick Ti3SiC2(0 0 0 1) films were grown on a 20-nm-thick TiC(1 1 1) seed layer on Al2O3(0 0 0 1) substrates under ultra-high-vacuum conditions by DC magnetron sputtering from three singleelement targets of Ti, Si and C, respectively, in an Ar atmo-
sphere. Detailed information about the deposition and growth processes can be found elsewhere [17]. The XRD measurements were carried out in a Philips X’pert MPD XRD system (theta/theta). As-deposited films were annealed in vacuum at room temperature (RT) at a pressure of 3 · 107 mbar in situ. The diffractometer used a water-cooled furnace equipped with a Be entrance window. The studied temperature range was 800–1400 C. Heating from room temperature to 800 C was carried out in steps of 50 C in 20 s time intervals. At 800 C the annealing temperature, Ta was kept constant for 1 h and then increased in steps of 100 C with a 1 h holding period after each step, during which four cycles of h–2h scans were conducted between the angles 33 and 44. The reason for the selected scan range was the proximity of the following high-intensive film and substrate peaks: Ti3SiC20 0 0 8; TiC1 1 1; Ti5Si3Cx0 0 0 2; Al2O30 0 0 6. Prior to each h–2h scan, 2h and x optimization scans around 2h = 35.5 (approximate peak position of TiC1 1 1) were performed. The samples were placed on a resistively heated Ta filament providing a chemically inert environment. On the backside of the filament a calibrated Pt/Rh thermocouple was clamped to monitor the temperature. A second Ta filament surrounding the sample in a 2 cm radius with a similar thermocouple was used to prevent temperature gradients across the sample. Space constraints made it necessary to cut the samples into dimensions of 4 · 5 mm performed with a diamond wheel saw. Epitaxial 2.5-lm-thick Ti3SiC2(0 0 0 1) films were put mechanically face-to-face and annealed in an alumina crucible using a Netzsch STA 409C thermal analysis instrument. The as-deposited samples were heated up to 1400 C with a heating rate of 10 K/min in an Ar atmosphere at 1 bar. Calibrated Pt/Rh thermocouples at the top and bottom of the crucible were used for temperature measurement. Bonding energies of the elements were obtained through X-ray photoelectron spectroscopy measurements using a PHI Quantum instrument with monochromatic Al Ka radiation. The peak areas in the spectra were weighed with internal sensitivity factors in order to acquire elemental film compositions. Ar-ion sputter cleaning of the samples was done for 100 s, within which sputtering was interrupted in regular intervals to carry out spectral measurements to obtain depth profiles for every sample. Measurements by XPS showed that as-deposited Ti3SiC2 films contain 52% Ti, 13% Si and 35% C. A composition of 49% Ti, 36% C and 15% of Si for the as-deposited films was measured with elastic recoil detection analysis using high-energetic (35 MeV) Cl7+ ions with an incidence angle of 15. Considering the measurement error limits and the limited compositional stability range of Ti3SiC2, we conclude that the films used in the present experiments are effectively stoichiometric. Microstructural analysis on the samples was carried out using transmission electron microscopy using a FEI Tecnai G2 TF 20 UT 200 kV FEG microscope. All samples were
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thinned to 60 lm through grinding with diamond paper and subsequently ion milled to electron transparency at 5 keV with a Gatan precision ion polishing system (PIPS). The final polishing step was performed at 2.5 keV. Lowand high-resolution TEM images were obtained in microprobe mode, whereas scanning transmission electron microscopy (STEM) imaging was performed in nanoprobe ˚ using a high-angle mode with a beam spot size of 5 A annular dark-field (HAADF) detector with variable camera length. Compositional analysis in STEM mode was done by energy-dispersive X-ray (EDX) analysis in STEM mode with applied drift correction. 3. Theoretical methods Ab initio calculations were carried out using density functional theory [18], as implemented in the Vienna ab initio simulation program, wherein projector augmented wave potentials are employed [19]. The generalized-gradient approximation was applied in all calculations, both electronic and phonon related, with the so-called Blo¨chl corrections for the total energy [20]. The integration in the Brillouin zone is done on special k-points, which is determined after Monkhorst and Pack [21]. All configurations were studied on a mesh of 7 · 7 · 7 irreducible k-points. The convergence criterion for the total energy was set to 0.01 meV within a 500 eV cut-off and the convergence test with respect to the configuration size was performed on 1 · 1 · 2 and 2 · 2 · 1 supercells (0.58 meV/atom energy difference was obtained). All configurations were relaxed with respect to atomic positions (internal free parameters), lattice parameter a and hexagonal c/a ratio. After relaxing the atomic positions and volumes, the cohesive energy was calculated from the obtained minimum configurations. We used the difference in cohesive energy with respect to Ti3SiC2 to estimate the stability of the configuration probed. Ti3SiC2, Ti3C2/Ti3C2 (Si removed from Ti3SiC2) and Ti3C2/O/Ti3C2 (O at Si site in Ti3SiC2) were treated with the unit cell containing 12 atoms, while Ti3C2was modelled with 24 atoms (1 · 2 · 1 supercell). Ti3C2 designates cubic TiC0.67 containing 33% of C-vacancies and it was described with six layers of TiC(1 1 1) stacked along [0 0 0 1] hexagonal direction. For the calculations of TiC0.67, four C vacancy positions were randomly chosen and the energy difference between two configurations probed was 11 meV/atom.
Fig. 1. X-ray diffractogram from an as-deposited Ti3SiC2(0 0 0 1) film grown on Al2O3(0 0 0 1) with a 20-nm-thick TiC seed layer.
The related Ti5Si3Cx(0 0 0 1) was observed as a minority phase. Fig. 2 shows the XRD peak intensity evolution over annealing time, ta, for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray), and Ti5Si3Cx0 0 0 2 (light gray) peaks from the Ti3SiC2(0 0 0 1) film. The sample was annealed at 800 C for ta = 1 h, after which the annealing temperature, Ta, was increased in 100 C steps with 1 h annealing at constant temperature after each step. It can be seen in Fig. 2 that the intensity of Ti3SiC20 0 0 8 decreases slightly with Ta increasing up to 1100 C. The pronounced decrease in the MAX-phase-peak intensity during 1 h annealing at Ta = 1200 C shows that decomposition of the Ti3SiC2 phase takes place at 1100 6 Ta 6 1200 C. The TiC1 1 1 peak exhibits a constant intensity between 800 and 1000 C. Above this temperature, however, the peak gradually gains intensity until Ta = 1300 C, where
4. Results and discussion 4.1. X-ray diffraction Fig. 1 is an X-ray diffractogram from an as-deposited Ti3SiC2(0 0 0 1) film. The exclusive presence of 0 0 0 ‘ indices confirm the epitaxial growth of the MAX phase as expected from the chosen growth conditions [17]. The presence of a TiC1 1 1 peak is due to the seed layer being deposited first.
Fig. 2. XRD peak intensity evolution over time for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray) and Ti5Si3Cx0 0 0 2 (light gray) peaks during annealing of a Ti3SiC2(0 0 0 1) film between 800 and 1300 C with temperature steps of 100 C and a 1 h holding period after each step.
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the intensity increase becomes definite. The intensity of the Si-rich Ti5Si3Cx phase remains low and constant throughout the annealing process. In order to provide insight into the dynamics of the decomposition process of Ti3SiC2, the temperature steps, DTa, were decreased to 20 C with the annealing temperature in each step kept constant for ta = 1 h. The XRDintensity evolution for this experiment is shown in Fig. 3. Between 1100 and 1160 C, the increase of the Ti3SiC20 0 0 8-peak intensity can be interpreted as increased crystalline quality due to annihilation of growth defects in the film. At Ta = 1180 C the Ti3SiC20 0 0 8 intensity decreases exponentially with an initial rate of DI/Dta(Ti3SiC2) of 19 counts/min, while TiC1 1 1 rises continuously at a rate of DI/Dta(TiC) of 0.2 counts/min. At the onset of Ti3SiC2 decomposition, Ta was kept constant for 20 h until the Ti3SiC20 0 0 8 peak was no longer detected. After that, Ta was ramped up to 1300 C, where the TiC1 1 1 intensity rises distinctively with a DI/Dta(TiC) of 7 counts/min (see Fig. 3). This indicates that TiC is growing as a decomposition product of Ti3SiC2. Again, no increase of the Ti5Si3Cx0 0 0 2 peak intensity was observed during the annealing experiment. A scan over 25 h was also performed at 1100 C (not shown). After a ta of 200 min, a steady decrease in the Ti3SiC20 0 0 8-peak intensity at a rate of DI/Dta of 0.8 counts/min was seen, indicating that the decomposition of Ti3SiC2 was activated. The rate of the change of the peak intensity in the XRD diffractograms is an indicator of the decomposition kinetics. The annealing experiments at Ta = 1100 C showed a very slow decomposition kinetics. After ta = 25 h more than 90% of the Ti3SiC2 film was still intact, as shown by TEM below. With a temperature increase of 80 C DI/Dta(Ti3SiC2) has drastically decreased, from 0.8 to 19 counts/min. This shows that the Ti3SiC2 decomposition is thermally activated.
Fig. 4. Time evolution of the X-ray FWHM for the TiC1 1 1 peak during annealing of a Ti3SiC2(0 0 0 1) film between 1100 and 1180 C with temperature steps of 20 C and 1 h holding periods after each step. After a 20 h holding period at 1180 C the temperature was increased to 1300 and 1400 C, respectively, each for 1 h.
The XRD full width at half maximum (FWHM) values for the TiC1 1 1 peak is shown in Fig. 4 vs. ta and Ta. There is a small decrease in the FWHM from 0.26 to 0.24 during 25 h at 1180 C. This is followed by a distinct decrease of FWHM to 0.2 within 1 h as Ta is raised to 1300 C. Since the FWHM value is an inverse measure of grain size, the observed drop in FWHM indicates that TiC layers recrystallize subsequent to the MAX-phase decomposition. 4.2. X-ray photoelectron spectroscopy The O1s peak at a binding energy of 532.3 eV and the Si2p peak at 103 eV indicate that the surface is mainly terminated by SiO2 [22,23]. After sputter cleaning the surface, O was no longer detected. The Si binding energy of the
Intensity [a. u.]
Si 2p
sputter cleaned for 30 s
T a = 1300 ˚C for 1 h
Fig. 3. XRD peak intensity evolution over time for Ti3SiC20 0 0 8 (black), TiC1 1 1 (gray) and Ti5Si3Cx0 0 0 2 (light gray) peaks during annealing of a Ti3SiC2(0 0 0 1) film between 1100 and 1180 C with temperature steps of 20 C and 1 h holding periods after each step. After 20 h holding period at 1180 C the temperature was increased to 1300 and 1400 C, respectively, each for 1 h.
108
106
104 102 100 Binding Energy [eV]
98
96
Fig. 5. X-ray photo-electron spectroscopy spectra showing the region of the Si 2p binding energy from the as-deposited Ti3SiC2(0 0 0 1) sample and a film annealed from 800 to 1300 C in steps of 100 C held at each temperature level for 1 h.
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as-deposited film of 99.2 eV, shown in Fig. 5, corresponds to the typical Si bonding environment in Ti3SiC2 [24]. Measurements on samples annealed at 1300 C and above showed no traces of Si in the film (see Fig. 5), but an O content [O] > 9%. XPS measurements on the sample annealed at 1100 C for 25 h showed a Si gradient into the film. The top-few nanometers did not contain Si, whereas the Si content increased noticeably after depth profiling to 50– 60 nm to the nominal 1/3 Si/Ti ratio. At the same time, an O diffusion profile was detected over the same depth with a significant 8% content at a depth of 30 nm. With the absence of Si, the surface is composed of TiO2 (530.1 eV) [25], with a large number of C–O bonds in the presence of Ti (531.1 eV) [26]. The peak attributed to TiO2 vanishes after 10 s sputter cleaning. 4.3. Transmission electron microscopy Fig. 6 is an XTEM micrograph from the sample annealed to 1100 C for 25 h. A 3–4-nm-thick TiO2 oxide layer evidenced by XPS can be seen on the surface of the sample. This is a thin protective oxide layer (too thin to attenuate X-rays in the above experiments) typically observed on Ti3SiC2 material [27,28]. Underneath the oxide layer, the MAX phase has decomposed and the extension of a TiC0.67 (nominal composition) layer can be seen to a depth of 60 nm. There are two cubic orientations found in this layer (see the Fourier transform inset in Fig. 6), which are twinned along ½1 2 1. In the XTEM image (Fig. 6), the twinned regions are separated by planar faults
Fig. 6. Cross-sectional TEM micrograph showing a Ti3SiC2(0 0 0 1) film annealed at 1100 C for 25 h with a Fourier transform from the TiC0.67 layer decomposed from Ti3SiC2. Twinned orientations along ½1 2 1 (horizontal line) can be seen distinguished by the prime for the twinned orientation.
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(PF) (see arrowheads in Fig. 6). No Si islands or other deposits of Si were found in the transformation layer, which is consistent with the XRD and XPS results above. The bulk of the sample (lower part of Fig. 6) consists of Ti3SiC2. Fig. 7 shows an STEM image from the Ti3SiC2 film annealed at 1100 C for 25 h. The TiC0.67 transformation layer, Ti3SiC2 and TiC seed layer are well resolved in bright contrast. Also, the PFs are seen in the TiC0.67 layer indicated by the arrowheads in Fig. 7. In addition, plate-like voids (magnified in inset in Fig. 7) parallel to the (1 1 1) planes are present in this layer. An EDX line scan performed across the TiC0.67 layer into the MAX-phase region is also presented in Fig. 7 (corresponding to the S–F line in the STEM image). The variation in elemental concentration in the inserted graph confirms the phase identification made above. A vertical dashed line in the graph depicts the interface between the transformation layer and the unreacted MAX phase. A residual amount of Si can be found in the transformation layer as well as a high O concentration, which decreases as the Si concentration increases with depth into the film. The planar faults in the transformation layer discussed above are likely locations for residual Si, as seen in XPS, in the form of locally retained Si layers in the decomposed Ti3SiC2 film. For higher annealing temperatures, the XRD analysis showed an accelerated MAX-phase decomposition. The XTEM micrograph in Fig. 8a confirms that the Ti3SiC2 film has completely transformed into TiC0.67 at Ta = 1200 C. The as-transformed layer, however, contains yet larger voids compared with the transformed volume at Ta = 1100 C (confer inset in Fig. 7). Near the surface the
Fig. 7. Scanning TEM image in nanoprobe mode with a high-angle annular dark-field detector with a long camera length showing a Ti3SiC2(0 0 0 1) film annealed at 1100 C for 25 h with two voids magnified in the inset. Elemental concentrations from the energy dispersive X-ray line scan along SF are shown in the graph inserted at the lower right.
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Fig. 8. Cross-sectional TEM micrograph (a) showing a Ti3SiC2(0 0 0 1) film containing transformed TiC0.67 layers after annealing to 1200 C for 1 h, and the corresponding selected area diffraction pattern (b) including Ti5Si3Cx (S), the Al2O3(0 0 0 1) substrate (A) and twinned TiC0.67 (T, T 0 ) plus double diffraction spots (D).
voids have coarsened and make contact with the surface at places. Thus, the initially smooth Ti3SiC2(0 0 0 1) surface (see Figs. 6 and 7) has become uneven, with low-index TiC0.67 crystal facet surfaces. As a consequence, there are additional O-diffusion channels, which may in part explain the large O concentration observed in the annealed films. A typical selected area electron diffraction (SAED) pattern from the sample annealed to 1200 C is shown in Fig. 8b. The phases present are TiC0.67, (T), Ti5Si3Cx, (S) and Al2O3 (substrate, denoted A). The crystallographical relationships determined from the SAED pattern are TiC0.67(1 1 1)//Al2O3(0 0 0 2) and TiC0:67 ½1 0 1==Al2 O3 3 2Þ and TiC0:67 ½1 0 1 0, and TiC0:67 ð1 1 1Þ==Ti5 Si3 Cx ð21 ½1 0 1==Ti5 Si3 Cx ½ 211 3, respectively. Spots denoted T 0 in Fig. 8b are caused by transformed TiC0.67 layers, which are twinned at the PFs. This instigates twin Moire´ patterns at overlapping grains in twin relationships resulting in double diffraction spots (denoted D) in the SAED pattern in Fig. 8b. Fig. 9 shows an HREM image of the transformed layer in a Ti3SiC2 film annealed to 1300 C. The transformed TiC0.67 layer again contains Ti5Si3Cx inclusions, in this case with the alternative crystallographic relationship to TiC0.67 of TiC0.67(1 1 1)//Ti5Si3Cx(0 0 0 2) and TiC0:67 ½1 0 1== Ti5 Si3 Cx ½ 1 0 1 0 also observed in the XRD measurements (Fig. 1). We deduce that this is the predominant crystallographic orientational relationship, due to the larger volume of film material probed in the XRD analysis compared with the electron diffraction analysis above. Initial annealing experiments with subsequent XRD measurements of the Ti3SiC2(0 0 0 1) films mechanically face-to-face showed that the MAX phase was stable to at least 1400 C. This is in agreement with the literature data for high-temperature stability of bulk Ti3SiC2 [9]. The observed increase in decomposition temperature for the samples put mechanically face-to-face compared with the vacuum-annealed samples can be attributed to the absence of an interface to vacuum as well as the Si entrapment caused by the initial formation of an oxide at the perimeter of the contacting films. These results indicate that the thermal stability of the MAX phase can be increased by preventing Si removal from the surface.
Scrutiny of the films annealed to 1200–1300 C reveals numerous PFs in the transformed layer, which are often arranged in a layered structure as in Ti3SiC2, though without any long-range order. Typical PF features can be seen in Fig. 10a–c, which are HREM images from a sample annealed to Ta = 1300 C. Fig. 10a reveals a set of three periodically spaced PFs. Their separation distance is ˚ , which is 2.3 A ˚ smaller than the unit cell of the 14.5 A MAX phase presented in Fig. 11. This can be explained if Ti3C2 slabs relax during the removal of the Si basal plane. A magnification of region A (Fig. 10b) shows sets ˚ . Also, inspecof atomic layers with a separation of 1.7 A tion of region B in Fig. 10a reveals a larger plane separa-
Fig. 9. High-resolution image from a Ti3SiC2(0 0 0 1) film annealed to 1300 C for 1 h showing a Ti5Si3Cx(0 0 0 2) grain embedded in the TiC0.67(1 1 1) transformed layer.
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˚ and Fig. 11. Ti3SiC2 unit cell with cell parameters of a = 3.0749 A ˚ , including the distance dI = 3.86 A ˚ between Ti(I) layers (Ti c = 17.7574 A ˚ between Ti(I) and Ti(II) layers (Ti not adjacent to Si) and dII = 2.49 A bound to Si; dII is similar to distance between the C layers).
4.4. Decomposition model Fig. 12a–d is a schematic illustration of our proposed mechanism of Ti3SiC2 decomposition based on the above results from thermal vacuum annealing experiments.
Fig. 10. Cross-sectional high resolution TEM images from a Ti3SiC2(0 0 0 1) film after phase transformation into TiC0.67(1 1 1) during annealing at 1300 C showing (a) an overview with different planar faults of types A and B, presented at higher magnification in (b) and (c), respectively.
tion and stacking sequence corresponding to a layer of locally retained Ti3SiC2 phase (marked by an arrow). A HREM micrograph of that area is shown in Fig. 10c. The distance between the layers on each side of the mirror ˚ . This corplane, marked with an arrow in Fig. 10c, is 4 A ˚ responds to the distance dI = 3.9 A (Ti(I)–Si–Ti(I); see unit cell in Fig. 11) in Ti3SiC2. It can thus be inferred that these mirror planes are occupied by Si as a residual part of locally retained Ti3SiC2.
4.4.1. Stage I: Si out-diffusion and evaporation We find that the Ti3SiC2 decomposition is initiated by out-diffusion and entropy-driven evaporation of Si from the top (0 0 0 1) film surface toward vacuum during annealing at Ta P 1100, and possibly even at 1000 C (see Fig. 2) on a nanometer depth scale. For 25 h annealing at 1100 C, the MAX phase is decomposed to a depth of 60 nm. Our recent in situ photoemission studies [28] of Ti3SiC2(0 0 0 1) thin films annealed in ultra-high vacuum revealed, with high surface sensitivity, Si enrichment on the free surface at 1000 C. The Si evaporation at the free film surface takes place since the Si vapor pressure of 8 · 106 mbar at 1200 C [29] is of the same order of magnitude as the pressure range used during annealing. The Si out-diffusion can be explained by its bond strength. Theoretical calculations of Zhou and Sun [30] show that Si is partially metallically bonded to Ti(I) (Ti atoms adjacent to Si) and thus the weakest bonded element in the Ti3SiC2 phase compared with strongly covalently and to some extent ionically bonded Ti(II) (Ti bonded to C but not Si; see Fig. 11) and C. Also, the Ti(I) to Ti(I) distance with Si mirror planes in between is the largest distance found in the Ti3SiC2 phase and allows Si to diffuse easier than C, which is smaller, but placed at interstitial sites. In addition, Barsoum et al. [31] observed Si displacement about its atomic equilibrium position at 900 C preferentially along the a-axis, due to relatively weak Si–Si bonds, which alleviates the Si diffusion. The diffusing Si
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Fig. 12. Schematic illustration describing the different stages of phase transformations occurring during the decomposition of Ti3SiC2(0 0 0 1).
atoms are likely to emerge on the Ti3SiC2(0 0 0 1) surface via domain boundaries or other defects in the film as well as from surface steps. Our earlier work [17] using atomic force microscopy showed that the Ti3SiC2(0 0 0 1) surface exhibits a step-like topography with step heights of full and half unit cells from a helicoidal growth mode around threading screw dislocations. At the side of each step, the Si planes are exposed to the ambient. We infer that Ti3SiC2 material exposing surfaces of other crystallographic orientations are likely to be even more prone to release Si.
4.4.2. Stage II: O uptake and SiO evaporation In the presence of an oxygen ambient (residual gas in the evacuated furnace in the present case), the Si-depleted highly defective structure formed under Stage I (see Fig. 12a) is liable to O in-diffusion. A large uptake of O (up to 10%) in our films was detected by XPS. During its way into the Ti3SiC2 film, O is likely to react with Si and form SiO gaseous species [13], thus accelerating the decomposition process as illustrated in Fig. 12b. O may eventually be incorporated in the material, e.g., by saturating
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the dangling bonds created by the Si-deficiency and thus creating O mirror planes between Ti3C2 slabs (Ti3C2/O/ Ti3C2).
Table 1 Calculated unit cell dimensions and cohesive energies for Ti3SiC2 and decomposition products ˚) a (A c/a DE (eV/atom)
4.4.3. Stage III: Ti3C2 relaxation, detwinning, and TiC0.67 formation by C-redistribution with void formation Fig. 12c shows the next stage in the decomposition process, which is activated at Ta P 1100 C. Here, the twinned Ti3C2 slabs (Ti3C2/Ti3C2) formed in Stages I and II relax towards each other. The transformation of the MAX-phase stacking is accompanied by a large reduction of volume as the Si-depleted hexagonal unit cell shrinks with or without the presence of O. A film entirely composed of a twinned Ti3C2 structure may give rise to superstructure peaks comparable to the 0 0 0 ‘ peaks in Ti3SiC2, but at slightly higher 2h angles considering the shrinkage of the unit cell. We find, however, no superstructure of the slabs in their repeated T and T 0 twin ordering. Also, it cannot be excluded based on these initial experiments that the Ti3C2T and T 0 slab relaxation is simultaneous with detwinning and recrystallization (see Stage IV below). This can be understood by considering the almost constant TiC1 1 1-peak intensity during 20 h annealing at 1180 C, as exemplified by Fig. 3 for Ta = 1180 C and ta = 20 h, which suggests – in such a case – that the coherent volume of the Ti3C2T/T 0 is too small to give rise to constructive interference during XRD measurements. The relaxation process includes also C redistribution from within the Ti3C2 slabs to the Si-depleted mirror planes. This process is in agreement with the known stability regime of TiCx (x = 0.97–0.47 [32,33]). Thus, homogenization in the C content to form TiC0.67 is what can be expected.
Ti3SiC2 Ti3iC2/Ti3iC2 recovery in z-direction Ti3iC2/Ti3iC2 recovery in x–y plane Ti3iC2/O/Ti3SiC2 TiC0.67
4.4.4. Stage IV: TiC0.67 layer growth and recrystallization From the XTEM study and the increase of TiC1 1 1peak intensity as well as the decrease in FWHM for TiC1 1 1 at Ta P 1300 C it can be concluded that the transformed layer undergoes both growth and recrystallization within the film as the final stage of the Ti3SiC2 decomposition process (see Fig. 12d). This is accompanied by an increased pore size observed during annealing between Ta = 1200 C and 1300 C (see Fig. 8a). Residual mirror planes containing Si as well as C and/or O may be present even up to Ta = 1300 C, as indicated by the large number of planar faults with different distances to their neighbor planes (see Fig. 10a–c). This may be the effect of local crystal perfection with no available diffusion paths for the Si. Note that the annealing temperatures used are 30% of the decomposition temperature of TiC [32]. Thus, only limited bulk diffusion in the TiC0.67 layers is expected. The recrystallization may instead be assisted by the enthalpy released in the above described reactions as well as by the relatively high planar-fault density and lattice point defects. The latter are, e.g., C vacancies and O interstitials, which eventually migrate to form the voids, but also any residual Si substitutional species. Clearly, the
3.075 3.078 3.087 3.062 3.066
5.775 4.934 4.845 4.967 4.894
0.000 0.352 0.377 0.564 0.392
chemical potential of Si in the various environments will determine the effective Ti3SiC2 stability. 4.5. DFT calculations Ab initio calculations were used to test our decomposition model of Ti3SiC2 in terms of the stability of the decomposition products, while kinetic and dynamic effects cannot be unraveled from these results. Table 1 contains the structural and stability data for Ti3SiC2, Ti3C2/Ti3C2, Ti3C2/O/ Ti3C2 and TiC0.67. Our ab initio data are consistent with the literature: the calculated equilibrium volume of Ti3SiC2 ˚ 3/atom, which is 0.9% larger than the experimenis 12.117 A ˚ 3/atom [1]. Furthermore, the calculated tal value of 12.003 A bulk modulus of Ti3SiC2 using the Birch–Murnaghan equation of states [34] is 209 GPa, which is 1.4% larger than the experimental value of 206 GPa [35]. In terms of the structure of the Ti3SiC2 decomposition products, it can generally be noted that the c/a ratio decreases upon decomposition, which is also consistent with our experimental data presented above. The unit cell volume of Ti3C2/Ti3C2 and Ti3C2/O/Ti3C2 is 15% smaller compared with Ti3SiC2, which is in agreement with the observations in Ref. [10]. The removal of Si layers from the Ti3SiC2 structure (Stage I in Fig. 12a) is associated with a decrease in the cohesive energy by 0.352 and 0.377 eV/atom, when the recovery is allowed in the z-direction and in the x/y plane, respectively. After the Si layer has been removed, two possible decomposition routes are probed. First, in the absence of oxygen, twins in Ti3C2 may be annealed out and C redistribution by short-range diffusion may take place, resulting in the formation of a TiC0.67 structure (Stage III). This is associated with the decrease in the cohesive energy by 0.392 eV/atom with respect to Ti3SiC2. Secondly, if oxygen is introduced into the remaining structure (Stage II), the cohesive energy decreases by 0.564 eV/atom with respect to Ti3SiC2. This stability increase suggests that oxidation is likely to occur in Ti3SiC2 upon removal of Si in an O-containing ambient. Therefore, the ab initio calculations support our decomposition model. 5. Conclusions We find that Ti3SiC2(0 0 0 1) thin films are effectively stable during vacuum furnace annealing up to 1100 C, above which they decompose into TiCx (x 0.67) following Si out-diffusion and evaporation. A sequence of phase
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transformation steps is determined, supported by DFT calculations, comprising relaxation and detwinning of Ti3C2 slabs to form TiC0.67, which undergoes recrystallization at 1200 6 Ta 6 1300 C. The Si out-diffusion may be accelerated by an O ambient through reactions to form gaseous SiO species. Thus, decomposition initiated at the surface on a nanometer scale is observed already at 1000 C. The transformed TiC0.67 layers contain voids and exhibit rough and facetted surfaces due to the involved phase transformations taking place under conditions of high lattice point defect content. The observed onset for the Ti3SiC2 decomposition initiated at the surface over a 10–60 nm depth at Ta = 1000–1100 C is in apparent contrast to the reported decomposition temperatures for bulk material of 1800– 2300 C [8,36]. We conclude that the discrepancy in the decomposition temperature is in part a result of the difference in detection sensitivity or chosen thickness criteria for a given transformation process between the studies. The role of the chemical environment is as important to the decomposition activation of Ti3SiC2 [10,13] as it is the activity of the A-element Si that is found to be the predominant factor in the decomposition of Ti3SiC2 and likely other MAX phases. Acknowledgements This work was supported by the Swedish Foundation of Strategic Research (SSF), Strategic Research Centre in Materials Science for Nanoscale Surface Engineering, the VINNOVA project on Industrialization of MAX-Phase Coatings, ABB Ltd. and Kanthal AB. Herbert Willmann is gratefully acknowledged for the measurements performed on the Netzsch thermal analysis instrument. The authors acknowledge Assoc. Prof. Jens Birch for valuable discussion as well as Dr. Ulrich Kreissig for performing ERDA measurements. JMS gratefully acknowledges support from Deutsche Forschungsgemeinschaft (DFG) (Schn 735/9-1). References [1] Jeitschko W, Nowotny H. Mh Chem 1967;98:329. [2] Barsoum MW. Prog Solid St Chem 2000;28:201.
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