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Microstructure and Nonbasal-Plane Growth

of Epitaxial MAX Phase Ti2AlN Thin Films

M. Beckers, N. Schell, R.M.S. Martins, A. Mcklich, W. Mller and L. Hultman1

1Thin Film Physics Division, Department of Physics, Chemistry and Biology (IFM),Linkping University, Linkping SE-581 83, Sweden

The Mn+1AXn phases [1] are a family ofnanolaminated compounds, where M is a transitionmetal, A an A-group element and X is nitrogenand/or carbon. Their hexagonal structure consistsof [MX6] octahedrons, interleaved by A-groupelement layers. In M2AX1 every third layer is

constituted by A-group atoms, every fourth forM3AX2, every fifth for M4AX3. For the Ti-Al-Nsystem Ti2AlN [2] and Ti4AlN3 [3] have beenreported. Here, the [MX]6 slabs are identical withthe structure of wear protective TiN. Since the M-X bonds are of strong covalent-ionic nature, theMn+1AXn phases exhibit typical ceramic propertieslike high melting points and thermal stabilities [4].On the other hand the M-A bonds are rather weak.Therefore, Mn+1AXn phases also exhibit metallic properties like good electrical-thermal conduct-ivity [5]. Under force effect they deform by kink bands which yields high ductility andmachinability. Due to this unique combination of properties, applications for Mn+1AXn phase aremanifold, one of them being wear protectivefunctional coatings for high temperature applica-tions. The first Mn+1AXn phase thin films, in theTi-S-C system, were deposited using magnetronsputtering [6]. Mn+1ANn phases are harder toobtain by sputter deposition, since the nitrogen hasto be introduced via the gas phase. Hence, onlyvery recently Joelsson et al. demonstrated thedeposition of Ti2AlN thin films [7]. They usedreactive sputtering from a compound 2Ti:Al target,

and deposited onto MgO(111) single-crystals atsubstrate temperatures of 830C.

The present work reports on successfulTi2AlN deposition by reactive cosputtering usingelemental Ti and Al targets, at considerably lowersubstrate temperatures of 690C studied by in-situx-ray diffraction. The temperature was measuredusing a thermocouple clamped onto the substrateholder. MgO(111) wafers were used as substrates.The base pressure at the deposition temperaturewas ~ 810-5 Pa. Prior to the Ti2AlN, a thin fcc(Ti1-xAlx)N seed layer was deposited at Ar and N2fluxes of 2.76 and 1.38 sccm at a working pressureof 0.35 Pa. The Ti and Al magnetrons were set to60 and 20 W, yielding a composition of

(Ti0.63Al0.37)N. For Ti2AlN growth the deposition pressure was increased to 0.8 Pa at Ar and N2fluxes of 7.94 and 0.48 sccm, respectively. The Tiand Al magnetron powers were changed to 80 and26 W.

The whole deposition chamber is mounted

into the goniometer of the ROBL beamline at theEuropean Synchrotron Radiation Facility (ESRF),and equipped with x-ray windows [8]. Monochro-matized x-rays of 12.915 keV were employed tostudy the Ti2AlN film growth in-situ in two differ-ent scattering geometries: (1) low angle specularreflectivity (XRR) either time-resolved at a fixedincidence angle, or scanned in coplanar scatteringgeometry for the determination of the filmthickness; (2) large angle x-ray diffraction (XRD)in Bragg-Brentano geometry in order to determinethe off-plane lattice parameter. The time resolved

XRR signal can be used to observe changes in thesurface morphology on an atomic scale, fromwhich the growth mode of the developing film canbe derived. The microstructure and texture of thefinal film was investigated using a combination oflab source pole figure measurements and cross-sectional transmission electron microscopy(XTEM). The final film composition was exa-mined by Rutherford backscattering spectroscopy(RBS), using a 1.7 MeV He+ beam.

Figure 1 shows time dependent XRR data for both the seed layer (a) and the Ti2AlN Mn+1AXn

phase layer (b) deposited onto MgO(111). Clearintensity oscillations are depicted for the seedlayer. The oscillation amplitude does not diminishduring deposition, hence a smooth layer-by-layer(step flow) growth onto MgO can be deduced, asexpected from reported epitaxial growth of pureTiN onto MgO(111) [9]. XRR characterizationduring a deposition interruption yielded a layerthickness of 121 , corresponding to a low growthrate of 0.3 /s, further promoting the layer-by-layer growth. Each oscillation in the timedependent XRR can be attributed to one lattice

constant of (Ti0.63Al0.37)N, hence a perfect cube-on-cube heteroepitaxial relation(Ti0.63Al0.37)N{100} // MgO {100}with only 0.74% lattice mismatch between

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permutation of the first three Miller indices yieldsfive additional Ti2AlN{1013} poles, positioned asindicated by the indices and the symbol ().Additionally, this figure must also be turnedaround by 120 and 240 due to the threefold

symmetry of the Ti2AlN(0002) peak. This makesup for another fourteen [2 times (6+1)] clusterpoints. For clarity, the indexing of these collectivepoints shown in Fig. 2(d) has been omitted. Eachin-plane turn of 120 is indicated by the symbols() and (). The picture obtained describes allthe features of the experimental pole figures.

The azimuthal orientational relationship canbe found by looking at the poles for= 90, whereobviously the MgO(110) and Ti2AlN(1210) polesoverlap. For the cubic MgO the correspondingdirection is [110] and for the Ti2AlN it is [1210].

Since there are three possible MgOdirections on the MgO(111) surface, the threefoldin-plane symmetry of the Ti2AlN can thus bedescribed by the orientational relationMgO{111} // Ti2AlN{1012}.

Figure 3(a) shows an XTEM micrographfrom Ti2AlN recorded along the MgO[112] zoneaxis. The layer is not single-crystalline, but exhi-bits equiaxed morphology with individual epitaxialgrains corresponding to Fig. 2. Their sizes are inthe dimension of the film thickness. Individualgrains show strong diffraction contrast, which is

due to the threefold in-plane symmetry describedabove. In accordance with time-resolved XRR thelayer displays high surface roughness, caused bythe tilted nonbasalplane growth.

100 nm

Ti( Al )N0.63 0.37

Ti AlN2

Ti AlN2

(0004)

(0006)

(0008)

(0002)

Ti AlN2

10 nm

(a)

(b) (c)

MgO

MgO

[112][111]

Fig. 3: XTEM micrographs of Ti2AlN grown on MgO(111)along the MgO[112] zone axis. (a) shows the film morpho-logy consisting of large crystal regions. (b) shows a highresolution micrograph at the interface with the typicalMAX phase 2-1-1 layered structure. It is made up by the(000l) planes as confirmed from the d-spacing calculationby Fast Fourier Transformation in (c).

Fig. 4: Atomic model illustration of the different(Ti0.63Al0.37)N and Ti2AlN crystallographic planes anddirections together with the geometric reconstructions forthe pseudomorphic fitting of Ti2AlN onto(Ti0.63Al0.37)N(111). (a) shows the (Ti0.63Al0.37)N(111)surface and (b) the three different Ti2AlNorientations of the Ti2AlN(0001) MAX phase basal-planealong the (Ti0.63Al0.37)N directions. (c) shows thesame threefold Ti2AlN orientations for theTi2AlN(101 2) planes. (d) shows cross-sectional views of

the Ti2AlN/(Ti0.63Al0.37)N interface for basal-plane andnonbasal-plane growth.

Figure 3(b) shows the interface between onegrains and the MgO substrate. The tilted Ti2AlNbasal-planes can be depicted from the layered 2-1-1 structure being typical for Mn+1AXn phases. Thelattice plane spacing calculations by Fast FourierTransformation confirm multiple Ti2AlN(000l)planes as shown in Fig. 3(c). The tilting angle ofthe basal planes is ~70, in agreement with the pole figure measurements. Hence, each of the

grains in Fig. 3(a) represents one possible in-planeorientation of the Ti2AlN(000l) plane.

To conclude, both pole figure and XTEMresults show that the orientational relationship between the MgO and the Ti2AlN isMgO{111} // Ti2AlN{1012}. This isin contrast to the basal-plane growth of Ti2AlN forhigh-temperature deposition conditions, withMgO{111} // Ti2AlN{0001} [7].This basal-plane growth can be expected due to thelow lattice mismatch of only 0.77 % between theTi2AlN(0001) and MgO(111) surfaces, as

indicated in Figs. 4(a-d). During nucleation theTi2AlN(0001) may arrange itself in three possibleorientations along the MgO directions.During coalescence, adjacent grains adjoin

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M. Beckers et al.; Microstructure and Nonbasal-Plane Growth ...12

(Fig. 4(b)), resulting in the reported single-crystalfilm morphology [7].

In our case the MgO(111) is parallel to theTi2AlN(1012) plane, which is projected inFig. 4(c). As indicated by the polygon on the

surface nitrogen atoms constituting this plane, alow lattice mismatch of 0.77% along oneMgO direction is possible. However, single-crystalline epitaxial growth is prevented. First, aninitial tilted nucleation with threefold in-planeorientation does not allow coalescence atincreasing film thickness. This is in agreementwith the polycrystalline equiaxed morphologyobserved in XTEM, and the pole figure results.Second, in the other MgO directions theadaptation can only be achieved by taking alsoover next surface nitrogen atoms into account, as

shown by the polygon in Fig. 4(a). Still, thereconstruction shows a quite large misfit of 7.88 to 7.43 , corresponding to 6%. Besides, thecross-sectional view in Fig. 4(d) demonstrates thatthe interface cannot be configured only by Ti-Nbondings, as it would be the case for basal-planegrown Ti2AlN(0001)//MgO(111). In fact, a ternaryTi-Al-N atomic reconfiguration of more thannearest-neighbor-atoms is necessary.

This complicated strained interfacial bondingcan be attributed to the following: Whendepositing on the (Ti0.63Al0.37)N seed layer, its

(111) surface tends to be N-terminated [11].Comparing with Fig. 4(d), this implies that theinitial atomic layer of Ti2AlN(0001) has to consist purely of Ti atoms. However, the employed co-sputtering process provides an incoming flux ofboth Ti and Al atoms, so the initial elementarycomposition would need to be reconfigured. Thisreconfiguration will be hindered, since the(Ti0.63Al0.37)N(111) surface is one with a relativelyhigh adatom bonding energy and thus low mobility[11]. Moreover, the comparably low substratetemperatures of only 690C will aggravate this

deficiency of adatom mobility. Hence, the textureevolution is kinetically prevented from achievingmaximum thermodynamic stability, e.g. basal-plane orientation. A deposition on MgO substrateswithout a (Ti0.63Al0.37)N seed layer at 830C [7]allows for the correct partitioning of the elementsat the substrate-film interface. Decreasing thesubstrate temperature as in the present case causesdifferent interfacial adaptation due to kineticalrestrictions as described above.

In summary, Ti2AlN single-phase thin filmshave been successfully deposited onto MgO(111)

substrates at growth temperatures of 690C. Thegrowth process has been characterized by in-situXRR and shows kinetic roughening of thedeveloping films. The orientational relationship

between MgO substrate and Ti2AlN as derived by pole figure measurements isMgO{111} // Ti2AlN{1012}. Thissuggests a tilted, nonbasal-plane growth of theTi2AlN which leads to a non single-crystal

morphology and rough, facetted surfaces as proven by XTEM. The nonbasal-plane growth can beexplained in terms of insufficient adatom mobilityduring deposition leading to a kinetically restrictedgrowth mode and to the observed tilted basal-planegrowth.

Acknowledgments

We gratefully acknowledge technicalassistance by U. Strauch during experiments andfinancial support from the Deutsche Forschungs-gemeinschaft under contract SCHE 682.

The presented results have been recentlypublished in a more detailed version as M. Beckerset al., J. Appl. Phys. 99 (2006) 34902.

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