Interdiffusion studies of single crystal TiN/NbN superlattice thin films

Interdiffusion studies of single crystal TiN/NbN superlattice thin films C. Engstro¨m, J. Birch, and L. Hultmana) Thin Film Physics Division, Department of Physics, Linko¨ping University, S-581 83 Linko¨ping, Sweden

C. Lavoie, C. Cabral, and J. L. Jordan-Sweet IBM Thomas J. Watson Research Center, Yorktown Heights, New York 10598

J. R. A. Carlssonb) 1-139 Engineering Science Building, CSL, 1101 West Springfield Avenue, University of Illinois, Urbana, Illinois 61801

共Received 19 December 1997; accepted 16 April 1999兲 The interdiffusion of TiN and NbN in superlattice structures has been investigated by studying the evolution of superlattice satellite peaks in x-ray diffraction, as well as by cross-sectional transmission electron microscopy. Single crystal TiN/NbN superlattices with composition modulation periods of 4.4 and 12.3 nm were deposited, by reactive dual-cathode unbalanced magnetron sputtering in an Ar/N2 discharge, onto MgO共001兲 substrates held at a temperature of 700 °C. Isothermal annealings 共in the range of 750– 875 °C for 20 min兲 as well as a ramped annealing 共3 °C s⫺1 up to 1200 °C兲 were performed, and in situ x-ray diffraction spectra were continuously recorded using synchrotron light and a linear detector. The results pointed to a nonlinear diffusion in TiN–NbN couples. The structure maintained abrupt interfaces throughout annealing, while the position of the interfaces was continuously shifted into the TiN layers. A model is proposed where Ti diffuses into the NbN layer to form a NbTiN alloy, while the diffusion of Nb in the opposite direction is restricted. Within the temperature range from 750 to 875 °C, activation energies for metal interdiffusion were limited to 1.2 eV for the lower temperature end, and 2.5 eV for the higher temperature end. The expected lifetime against alloying has been determined using the random walk theory, and a TiN/NbN superlattice with a period of 4.4 nm is expected to sustain a layered structure for ⬃10 h at 750 °C and ⬃2 h at 850 °C. © 1999 American Vacuum Society. 关S0734-2101共99兲03805-1兴

I. INTRODUCTION Layered structures such as nitride superlattices are an emerging class of hard protective coatings. Increased hardness as compared to single-layered structures has been reported both for as-deposited single crystal superlattices1 and polycrystalline nanolayered thin films2 for the material systems TiN/NbN and TiN/VN. The hardness enhancement has been explained by dislocation hindering at the interfaces due to differences in shear moduli, and by coherency strain caused by lattice mismatch of the two materials.3 Use of superlattices or nanolayered thin films as coatings on cutting tools exposes the structure to elevated temperatures during the cutting process, which causes increased interdiffusion between the layers. A diffuse interface would reduce the hardness, which inversely depends on the interface width.3 Not much is known, however, concerning the thermal stability of nitride superlattices, or the diffusion behavior of metals in nitrides, and this apparent lack of data has motivated the presented work. The superlattice structure with its high density of interfaces is in a nonequilibrium state, and since the TiN–NbN system is miscible,4 the 共Ti,Nb兲N pseudobinary alloy is more stable from a thermodynamical point of view. The presence a兲

Electronic mail: [email protected] Present address: Thin Film S-583 30 Linko¨ping, Sweden.

b兲

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of interdiffusion, activated by thermal energy, results in a modification of the composition profile. While TiN/NbN multilayers are known to be stable up to 700 °C, 5 questions remain as to the degrading mechanism and the diffusion behavior at higher temperatures. The most sensitive technique for measuring diffusion today is the modulated film technique, i.e., the use of superlattice structures and studies of the decay of satellite peaks in x-ray diffraction. The theory of interdiffusion in artificial compositionally modulated materials is well developed, and in its simplest form, the equations are linearized by treating several parameters as composition independent.6 Although the majority of the work on interdiffusion to date has been performed on metallic systems, this diffusion theory is general and does not depend on the type of bonding within the structure. In this article, results are presented on intermixing in TiN/ NbN single crystal superlattices. Isothermal anneals in the temperature range from 750 to 875 °C, as well as a ramped anneal up to 1200 °C were performed. In order to consider lattice diffusion, not affected by grain-boundary diffusion, single crystal films were studied. The structure maintains abrupt interfaces throughout annealing, while the position of the interfaces is continuously shifted towards thinner TiN layers during annealing. This indicates different activation energies for Ti in NbN and Nb in TiN, respectively, and proposes that, as soon as Ti is freed from the TiN and crosses

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©1999 American Vacuum Society

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Engstro¨m et al.: Interdiffusion studies of single crystal TiN/NbN

the interface, it moves quickly in the NbN to form a NbTiN alloy, while the opposite diffusion is limited. The interdiffusion limits the lifetime of the superlattice structure, which is critical for the use of these films as a hard protective coating. By using the random walk theory, the estimated lifetime for TiN/NbN superlattice structures has been determined. These calculations give an indication of useful operating temperatures in actual employment of nitride superlattices or nanolayered structures as wear protective coatings. II. EXPERIMENTAL DETAILS TiN/NbN superlattice thin films, with compositional modulation periods of ⌳⫽4.4 and 12.3 nm, were deposited by dual target reactive sputtering of Ti and Nb in an Ar/N2 discharge. Two 3-in.-circular planar magnetron cathodes with coupled magnetic fields7 were mounted on the deposition-chamber lid and tilted 25° with respect to the substrate surface normal. The targets were 99.9% pure Ti and Nb. Computer-controlled shutters, one for each magnetron, were used in addition to a rotating substrate table in order to get well-defined layers and interfaces. The gas inlets were positioned close to the bottom of the chamber. The substrates used were polished 0.5⫻10⫻10 mm3 MgO共001兲 single crystals. Before their insertion into the chamber, the substrates were cleaned in an ultrasonic cleaner with trichloroethylene, acetone, and ethanol, and subsequently blown dry with N2. The vacuum chamber was evacuated to a base pressure of 9⫻10⫺5 Pa (7⫻10⫺7 Torr), using a turbomolecular pump. The MgO substrates were heated by a resistive boron-nitride covered graphite heater, and their temperature was held at 900 °C for 1 h in order to degas and anneal the surface. The substrate temperature was then lowered to 700 °C and held constant during deposition. Prior to deposition, the targets were sputter cleaned for ⬃5 min in pure Ar at a pressure of 0.26 Pa 共2.0 mTorr兲. Thereafter, N2 was let into the chamber to a partial pressure of 6.2⫻10⫺2 Pa 共0.47 mTorr兲 while the Ar pressure was kept at 0.26 Pa 共2.0 mTorr兲, as controlled by a capacitance manometer. The discharges were established with constant-current power supplies at 0.5 and 0.7 A for the Ti and Nb target, respectively, giving a target voltage of 390 V for Ti and 360 V for Nb. Under these conditions, deposition rates of ⬃1.7 and ⬃2.0 Å/s for the TiN and NbN, respectively, were achieved. The total superlattice thickness was ⬃1 ␮ m. A TiN/NbN film consisting of MgO/150 nm TiN/50 nm NbN/20 nm TiN was also prepared under the same conditions, for elemental analysis in cross-sectional scanning transmission electron microscopy 共STEM兲. A Philips powder diffractometer equipped with a 1830 goniometer, line focus Cu source and focusing graphite monochromator on the secondary side was used to measure the x-ray diffraction 共XRD兲 from the TiN/NbN superlattices at low 共1 – 12°2 ␪ , i.e., reflectivity兲 and high 共36– 48°2 ␪ , around the 002 TiN/NbN average Bragg peak兲 angles. The experimental results from the powder diffractometer were compared to x-ray reflectivity simulations that used an optical formalism originally developed by Parrat,8 equivalent to JVST A - Vacuum, Surfaces, and Films

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the dynamical theory of XRD. A Philips materials research diffractometer 共MRD兲 x-ray diffractometer, equipped with a point focus Cu source, soller slits, and a flat graphite monochromator on the secondary side was used for texture measurements. Pole plots were obtained with ␾ and ␺ step lengths of 2.5°. The interdiffusion of TiN/NbN as performed by annealing superlattice films in a purified helium atmosphere, while the evolution of the superlattice satellite peaks was recorded continuously using synchrotron XRD. The annealing chamber was evacuated to 10⫺4 Pa (10⫺6 Torr) and then backfilled to 1 atm with purified He. The samples were either ramped at 35 °C s⫺1 and held at a constant temperature for 20 min 共isothermal anneal兲, or ramped at 3 °C s⫺1 to 1200 °C 共ramped anneal兲. The measurements were performed at beamline X-20C at the National Synchrotron Light Source in Brookhaven, NY. Synchrotron radiation was monochromatized using wide bandpass artificial multilayers resulting in an energy resolution of 1.5% at 6.9 keV 共1.797 Å兲. At these conditions, an area of approximately 2 ⫻2 mm2 on the sample received a typical photon flux of 3 ⫻1012 photons s⫺1. XRD patterns were collected using a position sensitive detector 共linear diode array兲 with millisecond time resolution. The linear detector simultaneously received diffracted x rays in a range of 10°2 ␪ . The diffraction spectrum was taken every 2 s during the 20 min isothermal anneal. Samples for transmission electron microscopy 共TEM兲 were prepared by ion milling 共ion etching兲 in a Bal-Tec RES 010 equipment. Samples for cross-sectional TEM 共XTEM兲 were glued together using Araldite glue cured at 180 °C, mechanically thinned to a thickness of ⬃60 ␮ m and subsequently polished using 0.25 ␮m diamond paste in the final polishing step. For both sides of the sample, the initial step of ion etching was carried out using two ion guns operated at 9.9 kV and 4.2 mA, with Ar as sputtering gas. The incident angle of the ion beams was kept at 6°–7° with respect to the sample surface plane and the sample was rotated continuously around its surface normal until a depression was formed in the sample. Subsequent etching was carried out using only one ion gun oriented orthogonal to the film surface, while the sample was continuously rocked ⫾50° about the projection of the beam axis. In the final etching stage (⬃5 min), the gun was operated at 3.0 kV and 1.5 mA in order to remove residual amorphous material which was created during milling at higher power settings. TEM analysis was performed in a Philips CM20 ultra twin 共UT兲 microscope. Compositional profiles across the interface of XTEM foils were obtained using a Vacuum Generators HB501 STEMenergy-dispersive x-ray microscope equipped with a field emission source and operated at 100 kV. The samples were scanned with an electron beam focused to a diameter of approximately 1 nm. Emitted x rays were collected at a take-off angle of 25° by an energy-dispersive detector through a thin polymer window. Corrections for atomic number were carried out following Ref. 9.

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III. DIFFUSION THEORY A superlattice structure, which is a periodic composition modulation, can be described by a Fourier series10 ⫹⬁

c共 x 兲⫽

兺 A m sin m⫽⫺⬁

冉

m

冊

2␲ x , ⌳

共1兲

where A m is the amplitude of the mth order Fourier component. A 1 is the amplitude of the basic sine wave 共also called the fundamental component兲, which then is modulated into a square wave by adding higher order coefficients. The amplitude of the Fourier coefficients are related to the intensity of 2 11 , and thus, an inthe superlattice satellite peaks: I m ⬀A m creased diffusion, which leads to a decrease in the modulation amplitude, can be viewed as a decrease in the intensity of the satellite peaks. The decrease of the satellite intensity is used to calculate the diffusivity in the so called continuum approach of the linearized diffusion equation6

冉冊

d 2␲m ˜ 共 ln I m 兲 ⫽⫺2D dt ⌳

2

,

共2兲

˜ is the effective interdiffusion coefficient. This is a where D standard method for studying interdiffusion in superlattices, and is valid except when the repeated layer thickness ⌳ is of the same order as the interatomic spacing.6 From the temperature dependence of the diffusivity, an activation energy E a can be extracted by an Arrhenius expression

冉冊

˜ ⫽D ˜ 0 exp ⫺ D

Ea E ˜ ⫽ln D ˜ 0⫺ a , ⇔ln D kT kT

共3兲

where k is the Boltzmann constant. Thus, a plot of the loga˜ as a function of 1/T yields the activation energy rithm of D from the slope of the graph. For a more detailed description of the linearized diffusion equation, the reader is referred to, e.g., Refs. 6 and 12. IV. RESULTS The superlattices exhibited flat and shiny macroscopic surfaces, and were single crystal as seen by XRD and TEM. The as-deposited superlattices exhibited a large number of x-ray satellites, indicating abrupt interfaces and a strong compositional modulation. This is demonstrated, using the powder diffractometer, for the as-deposited superlattice with a periodicity ⌳ of 12.3 nm in Figs. 1共a兲 and 1共b兲, showing experimental and simulated x-ray reflectivity scans, and a high-angle x-ray diffractogram recorded around the TiN/ NbN 002 Bragg peak, respectively. A TiN/NbN 兵110其 XRD pole plot of ␺ ⫽0° – 90° for the very same film is shown in Fig. 1共c兲, as recorded by the MRD diffractometer. The film shows four sharp peaks at the same ␺ angle of ⬃54°, separated from each other by 90° in ␾. The corresponding 兵110其 peaks from the MgO were also detected since the divergence of the detector was 0.3°2 ␪ . Thus, the superlattice is epitaxial with the same orientation as the MgO substrate. J. Vac. Sci. Technol. A, Vol. 17, No. 5, Sep/Oct 1999

FIG. 1. 共a兲 Reflectivity and 共b兲 high-angle ␪ –2␪ x-ray diffractograms, as well as a simulation of the reflectivity data from the as-deposited TiN/NbN superlattice with a periodicity of 12.3 nm. ␪ B denotes the average so called ‘‘Bragg peak’’ (m⫽0). The satellites are labeled with their m values. 共c兲 XRD pole plot of the TiN/NbN 兵110其 planes from the same film.

The position of the reflectivity peaks were used to determine ⌳ by linear regression of m 2 versus sin2 ␪, using the modified Bragg’s law m␭⫽2⌳ sin ␪ ⇔ sin2 ␪ ⫽

冉冊 ␭ 2⌳

冑

1⫹

¯␩ 2 ⫺1 sin2 ␪

共4兲

2

m 2 共 ¯␩ 2 ⫺1 兲 ,

共5兲

where ␪ is the diffraction angle, ␭ is the wavelength of the x rays, m is the satellite index, and ¯␩ is the average refractory index of the materials.13 Furthermore, simulations of reflectivity scans yielded the individual layer thicknesses. In the example shown in Fig. 1共a兲, the layer thicknesses for TiN and NbN were 6.75⫾0.05 and 5.55⫾0.05 nm, respectively,

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FIG. 3. Example of the raw data collected by the linear detector during a 1200 s 共20 min兲 isothermal anneal at 875 °C, showing the intensity of highangle satellite peaks on the negative flank of the TiN/NbN 002 Bragg peak from the 4.4 nm sample. 共Note that the wavelength of the beam is 1.8 Å.兲 FIG. 2. XTEM image, with a higher magnification inset from a 1-␮m-thick superlattice film with the periodicity ⌳⫽12.3 nm. A well-defined layered structure is viewed throughout the film. The film exhibited a columnar structure along the growth direction. The higher-magnification inset demonstrates the well-defined interfaces revealed by the elemental contrast of the TiN 共light兲 and NbN 共dark兲 layers.

with an interface roughness of ⬃0.1 nm. No shift of the satellite peak positions was observed in any of the samples during isothermal anneals, nor after the ramp anneal, demonstrating that the periodicity of the superlattices remained constant. A XTEM image of the layered structure of the 12.3-nmperiod film is shown in Fig. 2. The film exhibited an apparent columnar microstructure along the growth direction, with relatively flat and large terraces separated by regions where the superlattice layers curved towards the substrate. The film, however, maintained a continuous crystal lattice despite the local layer curvature, such that no grain boundaries were present. At the film top surface, cusps were present in between the emerging columns, evidently showing that the columnar structure had its origin in an undulated surface during growth. The higher magnification inset in Fig. 2 demonstrates the well-defined interfaces between the TiN and NbN layers, revealed by the elemental contrast. As analyzed by Rutherford backscattering 共RBS兲, as well as depth profiling Auger, no contaminants were detected down to ⬍1 at. % . Isothermal anneals of 4.4 nm superlattice samples were carried out at several temperatures ranging from 750 to 875 °C. The intensity of high-angle satellite peaks on the negative flank of the TiN/NbN 002 Bragg peak, was recorded during the anneals. An example of the raw data collected by the linear detector is shown in Fig. 3, which demonstrates the evolution of three order satellites during an isothermal anneal at 875 °C for 20 min. The decrease of the first satellite peak (m⫽⫺1) as a function of time, for all isothermal anneals, is shown in Fig. 4. In accordance with the linear diffusion theory, as derived from Eq. 共2兲, a plot of ln(I/I0) versus isothermal annealing time is linear with a slope proportional to the diffusion coefficient. An effective JVST A - Vacuum, Surfaces, and Films

˜ was extracted from the linear part (t⬎5 min) diffusivity D of each curve. ˜ values obtained from Fig. 4 are plotted in an The D Arrhenius plot in Fig. 5. A linear fit to these data points yields the activation energy for the interdiffusion. However, no linear fit can satisfyingly be made to all the data points. Even so, for the higher temperature range, an activation energy extracted from Eq. 共3兲, with a linear fit in Fig. 5 yields a value of 2.5⫾0.1 eV. Correspondingly, a value of 1.2 ⫾0.1 eV was obtained for the activation energy from the lower temperature range. Obviously, this derivation suffers from fitting only three data points, but nevertheless places limits for the activation energies of processes responsible for metal interdiffusion inside of this temperature regime. The ˜ 0 in the two temperature regimes was also obprefactor D ˜ 0 ⬃1 tained from the fittings to the data in Fig. 5: D ⫺17 2 ⫺1 ⫺11 2 ˜ 0 ⬃1⫻10 m s for T⬍830 °C, and D m s⫺1 for ⫻10 T⬎830 °C. Figure 6 shows the behavior of the first (m⫽⫺1) and second (m⫽⫺2) order satellites around the TiN/NbN 002 Bragg peak, as a function of annealing time for three differ-

FIG. 4. Decay of the first high-angle XRD satellite reflection (m⫽⫺1) during isothermal anneals of the ⌳⫽4.4 nm film, at temperatures ranging from 750 to 875 °C.

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˜ ) obtained from Fig. 4 plotFIG. 5. Temperature-dependent diffusivities (D ted in an Arrhenius plot.

ent temperatures: T⫽830, 850, and 875 °C. The second order satellite decays very rapidly at initial stages of the annealing, and for the higher temperatures 共T⫽840, 875 °C兲, its intensity decreases to a minimum and then increases again. This indicates a nonlinear diffusion behavior as will be discussed in Sec. V. A study of the composition profile evolution during a temperature ramp anneal (3 °C s⫺1) to 1200 °C was performed on the superlattice film with a period of 12.3 nm. Figure 7 displays a comparison between the reflectivity scans of the as-deposited and the ramp-annealed samples, including a simulated diffractograms. The as-deposited superlattice was asymmetric, with a thicker TiN layer, and the decrease in intensity of every even-numbered satellite peak in the diffractogram of the annealed sample is consistent with a formation of a more symmetric superlattice. Simulation results show a change in the layer ratio during annealing from 67:55

FIG. 6. Decay of the first (m⫽⫺1) and second (m⫽⫺2) order high-angle XRD satellite reflections during isothermal anneals of the ⌳⫽4.4 nm film, at 共a兲 T⫽830, 共b兲 850, and 共c兲 875 °C. J. Vac. Sci. Technol. A, Vol. 17, No. 5, Sep/Oct 1999

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FIG. 7. Reflectivity scans from 共a兲 an as-deposited as well as 共b兲 a rampannealed 共3 °C s⫺1 to 1200 °C兲 TiN/NbN superlattice sample with ⌳ ⫽12.3 nm, also including x-ray simulations.

共TiN/NbN兲 to 60:63 共TiN/NbTiN兲 with well-defined interfaces and a constant roughness of ⬃0.1 nm, for both films. The shift of the interface and the unaltered superlattice periodicity are corroborated by XTEM images of the asdeposited and annealed samples, as shown in Fig. 8. The elemental contrast in Fig. 8 reveals the well-kept layered structure and the more symmetric appearance of the annealed sample at this stage of annealing. Titanium and Nb concentrations across the interface of a diffusion-couple sample were studied by STEM. The diffusion couple consisted of MgO/150 nm TiN/50 nm NbN/20 nm TiN. Figure 9 shows the diffusion profile across the TiN–NbN interface, after the film had been annealed for 1 h at 850 °C. Linescans, as well as individually probed positions across the interface, were recorded. The measurements reveal a Ti concentration of ⬃6 at. % throughout the NbN layer, while no Nb was detected in the TiN layers. The detected Ti concentration was not a result of redeposition during sample preparation 共ion etching at low angles兲, since no Ti was detected in the MgO. The interface broadening observed in Fig. 9 is due to instrumental factors.14

FIG. 8. XTEMs of 共a兲 as-deposited and 共b兲 ramp-annealed 共3 °C s⫺1 to 1200 °C兲 single crystal TiN/NbN superlattice with ⌳⫽12.3 nm.

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TABLE I. Reported data of activation energies for metal diffusion in nitrides. Diffusion species/material N in TiN Ti in AlN Nb in AlN Pu in PuN Pu in Pu共C,N兲 Metal interdiffusion in TiN/NbN Metal interdiffusion in TiN/NbN

FIG. 9. STEM measurements of the 共a兲 Ti and 共b兲 Nb concentrations across the TiN–NbN interface in a MgO/150 nm TiN/50 nm NbN/20 nm TiN diffusion couple annealed for 1 h at 850 °C. Linescans, as well as individually probed positions 共circles and triangles, respectively兲 across the interface are shown.

V. DISCUSSION The presence of local deviations from planar growth of the as-deposited TiN/NbN films observed in XTEM images 共see Fig. 2兲 can be explained by coherency-strain relaxation close to the film/substrate interface.15 Also, a tendency for three-dimensional growth of the NbN layers on TiN is possible.16 Steady-state microstructure evolution was characterized by surface cusps with trailing columnar boundaries from each subsequent nitride layer replicating the former surface, while maintaining a continuous crystal lattice across the apparent boundary.16 Surface cusps developed from initial film surface asperities as an effect of self-shadowing of the deposition flux in combination with a limited adatom mobility close to the epitaxial temperature.15 For this chamber geometry, it is noted that the two magnetrons each gave a metal vapor flux directed 25° with respect to the substrate surface normal. The films nevertheless exhibited the most defined layered structure for a nitride, as seen from XRD. This is ascribed to the employment of substrate table rotation that gives relatively even layer thickness distribution over the film area analyzed. During isothermal anneals, the first superlattice satellite decayed exponentially except for the initial annealing time where some deviation was observed 共see Fig. 4兲. This initial nonlinearity is attributed to microstructural changes since no corresponding nonlinearity was observed in comparable scans of low-angle satellites.23 The exponential decay of the first satellite is expected since the first component in the Fourier series description of the periodic superlattice structure is mainly sensitive to the average decay of the modulation amplitude. However, according to the linear diffusion theory, an initial square wave composition profile will diffuse to become a sinusoidal modulation wave, and the higher-order satellites would also decrease exponentially, with a slope m 2 times that of the first satellite 关see Eq. 共2兲兴. The behavior of the second satellite during annealing 共see Fig. 6兲 is thus not described by the linear theory, and the diffusion for the TiN/NbN system is to be viewed as nonlinear. The deviation of the higher Fourier components from the JVST A - Vacuum, Surfaces, and Films

E a 关eV/at.兴 2.1–2.2 8.0 3.7 4.3 3.4 1.2⫾0.1 2.5⫾0.1

T 关 °C兴

Ref.

1000–1500 28, 29, 30, 31 1280–1400 32 1400–1600 32 19 33 750–830 present work 830–875 present work

linear theory is attributed to their role in modifying the composition profile. Similar profile evolution have been reported in earlier studies of interdiffusion in, e.g., the Ag–Au binary system12 and in modulated Cu–Ni foils.12 The decrease of every even-numbered satellite peak experienced in the ramp-annealed sample 共see Fig. 7兲 indicates a shift of the interfaces to a more symmetric superlattice, which causes extinction of even-numbered satellite peaks due to destructive interference. An interface shift is also evident in the isothermal-annealed samples 共see Fig. 6兲, and the evolution of the second order satellite peak during anneals at 850 and 875 °C reveals a shift of the interface past the symmetric structure as the peak intensity passes a minimum. These observations point to a nonbalanced diffusion leading to a Kirkendall shift of the interface17 such that more diffusion takes place into the NbN layers than into the TiN layers. However, no void formation was observed in the TiN layers by XTEM imaging of the ramp-annealed sample. This suggests that the diffusion is balanced by variations in the vacancy concentration on the metal sublattice in the nitrides, or that Ti diffuses from near the interface while the created vacancies are filled by Nb such that the NbN:Ti layer moves into the TiN layer. The predominant diffusion of Ti into the NbN layers is also confirmed by STEM measurements of TiN/NbN diffusion couples 共see Fig. 9兲, in which no Nb is traced in the TiN layer, while Ti is observed with an even concentration throughout the NbN layer. The above results infer that Ti is the dominating metal atom in the diffusion process, and thus, the activation energies E a discussed below, primarily refer to Ti diffusion in NbN. The two activation energies obtained in the temperature range T⫽750– 830 °C indicates the existence of different diffusion mechanisms with E a ⫽1.2⫾0.1 eV for the lower temperature end and E a ⫽2.5⫾0.1 eV for the higher temperature end. Literature data on activation energies for metals in nitrides, as well as for nitrogen, are presented in Table I. The data are scarce, and concerning the metal diffusion, there are no reports for the TiN/NbN materials system. For some other nitrides, activation energies for metal diffusion range from 3.7 to 8.0 eV, while the activation energy for nitrogen diffusion 共in TiN兲 is lower and ranges from 2.1 to 2.2 eV. Nitrogen diffuses significantly faster than the metal atoms,18 which diffuse via a vacancy mechanism on the metal sublattice.19 A relatively high-temperature stability for TiN has been reported by Mader, Fischmeister, and Bergmann,20 who detected no change in defect structure dur-

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˜ t⫽⌳/2. l⫽ 冑6D

FIG. 10. Phenomenological model for diffusion in TiN/NbN couples. The model assumes that Ti atoms and vacancies in TiN are smaller than Nb atoms or vacancies in NbN by virtue of the difference in lattice parameter 共a 0,TiN⫽4.24 Å and a 0,NbN⫽4.39 Å兲. An order of activation energies for in NbN in TiN ⬍E Nb . diffusion is assigned: E Ti a a

ing an anneal at 900 °C, as seen in TEM. During annealing at 1300 °C for 1 h, however, void growth at TiN grain boundary triple points took place,20 which should correspond to bulk diffusion, generally activated at approximately half TiN of the melting temperature in units of degrees Kelvin (T m ⫽2950 °C). From the above comparison, we infer that the higher E a value (⬃2.5 eV) observed for the temperatures 830– 875 °C corresponds to bulk-like diffusion of Ti in NbN:Ti. The low E a (⬃1.2 eV) for T⬍830 °C, however, may correspond to diffusion mediated by point defect arrangements in the nitrides. As a matter of fact, Elstner, point defects ⫽1.23 eV for Kupter, and Richter,21 obtained E TiN, a lattice relaxation from compressive stress state in sputterdeposited films deposited at 200 °C and annealed at 450 °C. A model is proposed to describe the diffusion in the TiN/ NbN system. We assume different activation energies for Ti in NbN in NbN and for Nb in TiN—namely E Ti a Nb in TiN ⬍E a —and that Ti atoms and vacancies in TiN are smaller than Nb atoms or vacancies in NbN by virtue of the difference in lattice parameter (a 0,TiN⫽4.24 Å⬍a 0,NbN ⫽4.39 Å) The model is pictured in Fig. 10: As soon as Ti is freed from the TiN and crosses the interface, it moves rapidly in the NbN to form a NbTiN alloy, while the opposite diffusion is limited. The superlattice will thus keep a welldefined layered structure throughout the interdiffusion process while the interfaces will shift towards thinner TiN layers. When all the TiN is consumed by alloying the NbN layers, the superlattice structure will cease to exist, and the film will have transformed into a homogeneous Nbx Tiy N alloy. The diffusion data from Fig. 4 were used for an estimation of the lifetime t of the superlattice structure with respect to alloying as a function of temperature. The lifetime was calculated using the random walk theory for a face-centeredcubic 共fcc兲 lattice22 where complete interdiffusion was ex˜ denotes pected for a diffusion length l of ⌳/2, and where D the temperature-dependent diffusivity J. Vac. Sci. Technol. A, Vol. 17, No. 5, Sep/Oct 1999

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共6兲

The calculations yield an estimated lifetime for the 4.4 nm superlattice structure of ⬃10 h at 750 °C and ⬃2 h at 850 °C. Initial results23 show comparable diffusivities for the two superlattice periods of 4.4 and 12.3 nm at ⬃900 °C, and thus, larger period superlattices, or multilayers with dense grain boundaries, should be expected to exhibit longer effective lifetimes against alloying, according to Eq. 共6兲. Corresponding calculations yield an ⬃18 h lifetime for the 12.3 nm TiN/NbN superlattice, at 850 °C. The results from the calculations of the superlattice lifetime may serve as input for selecting cutting data in applications with superlattices as wear-protective coatings. As a matter of fact, the temperature distribution on a cutting edge of a cemented carbide cutting insert has been estimated by Dearnley.24 Depending on cutting data 共cutting speed, feed rate, depth of cut, workpiece material, etc.兲 different temperature distribution will be obtained. Generally, the maximum temperature is reached at the rake face ⬃0.2– 0.4 ␮ m from the edge, and may reach 1000– 1200 °C, while the temperature at the flank face is normally ⬃200– 400 °C lower. The flank wear is characterized by abrasive wear, such that a high hardness and abrasive wear resistance of the coating is desired. Coatings of TiN/NbN are reported to exhibit significant abrasive wear resistance at room temperature,25,26 and recent results by Selinder et al.27 indicate that TiN/NbN multilayer-coated WC/Co perform better than TiN reference coatings in milling machine tests, with workpiece materials including both general steel and stainless steel. The present results of the thermal stability of TiN/NbN multilayers at ⬃850 °C thus substantiates the findings in Ref. 27.

VI. CONCLUSIONS The interdiffusion of TiN/NbN superlattice thin films was studied by XRD and electron microscopy. The results show that the diffusion is nonlinear. The structure sustains abrupt and well-defined interfaces throughout annealing, however, the interfaces are continuously shifted towards thinner TiN layers. This indicates different activation energies for Ti in NbN and for Nb in TiN, respectively. A model is suggested where Ti, as soon as it is freed from TiN, moves rapidly in the NbN to form a NbTiN alloy, while the opposite diffusion of Nb is limited due to a higher activation energy. Activation energies for the metal interdiffusion, dominated by Ti diffusion in NbN, show a low value of 1.2⫾0.1 eV for temperatures ⬍830 °C, which may correspond to point-defect assisted diffusion. For higher temperatures and up to 875 °C, the activation energy is 2.6⫾0.1 eV, corresponding to bulklike diffusion. As determined by the random walk theory for a fcc lattice, a TiN/NbN superlattice with a period of 4.4 nm has an expected lifetime before alloying of ⬃10 h at 750 °C and ⬃2 h at 850 °C, while a TiN/NbN superlattice with a period of 12.3 nm would last for ⬃18 h at 850 °C.

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ACKNOWLEDGMENTS The Materials Research Consortium for Thin Film Growth, funded by the Swedish Foundation for Strategic Research 共SSF兲 is acknowledged for financial support. Acknowledged are also Ulf Wahlsto¨m at IMC, Linko¨ping, Sweden, for technical support, Dr. Philip Yashar at Northwestern University, Evanston, USA, for x-ray reflectivity simulations, Dr. Bjo¨rgvin Hjo¨rvarsson at the Royal Institute of Technology, Stockholm, Sweden, for RBS measurements, ˚ bom at Linko¨ping University, Sweden, for Auger Lisa A measurements, and Professor Jan-Eric Sundgren, Thin Film Physics Division, Linko¨ping University, Sweden, for discussions. 1

P. B. Mirkarimi, M. Shinn, L. Hultman, and S. A. Barnett, J. Vac. Sci. Technol. A 10, 75 共1992兲. 2 X. Chu, M. S. Wong, W. D. Sproul, S. L. Rhode, and S. A. Barnett, J. Vac. Sci. Technol. A 10, 1604 共1992兲. 3 X. Chu and S. A. Barnett, J. Appl. Phys. 77, 4403 共1995兲. 4 H. Holleck, Binare und Ternare Carbid und Nitridsysteme der Ubersgangmetalle 共gebruder Borntreager, Berlin-Stuttgart, 1984兲. 5 S. Lopez, M.-S. Wong, and W. D. Sproul, J. Vac. Sci. Technol. A 13, 1644 共1995兲. 6 A. L. Greer and F. Spaepen, in Synthetic Modulated Structures, edited by L. L. Chang and B. C. Geissen 共Academic, New York, 1985兲. 7 B. Window and N. Savvides, J. Vac. Sci. Technol. A 4, 2 共1986兲. 8 L. G. Parratt, Phys. Rev. 95, 359 共1954兲. 9 G. Cliff and G. W. Lorimer, J. Microsc. 103, 203 共1975兲. 10 J. DuMond and J. P. Youtz, J. Appl. Phys. 11, 357 共1940兲. 11 A. Guinier, X-ray Diffraction 共Freeman, San Francisco, 1963兲. 12 T. Tsakalakos, Ph.D. thesis, Northwestern University, Evanston, Illinois, 1977兲.

JVST A - Vacuum, Surfaces, and Films

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B. K. Agrawall, X-ray Spectroscopy—An Introduction 共Springer, Berlin, 1979兲. 14 I. Petrov, E. Mojab, R. C. Powell, J. E. Greene, L. Hultman, and J.-E. Sundgren, Appl. Phys. Lett. 60, 2491 共1992兲. 15 L. Hultman, L. R. Wallenberg, M. Shinn, and S. A. Barnett, J. Vac. Sci. Technol. A 10, 1618 共1992兲. 16 L. Hultman, M. Shinn, P. B. Mirkarimi, and S. A. Barnett, J. Cryst. Growth 135, 309 共1994兲. 17 M. Ohring, The Materials Science of Thin Films 共Academic, San Diego, 1992兲. 18 H. Matzke, Defect Diffus. Forum 83, 111 共1992兲. 19 H. Matzke, J. Phys. 共Paris兲, Colloq. 4, 24 共1979兲. 20 W. Mader, H. D. Fischmeister, and E. Bergmann, Thin Solid Films 182, 141 共1989兲. 21 F. Elstner, H. Kupfer, and F. Richter, Phys. Status Solidi A 147, 373 共1995兲. 22 P. G. Shewman, Diffusion in Solids 共McGraw-Hill, New York, 1963兲. 23 C. Engstro¨m, J. Birch, L. Hultman, C. Lavoie, C. Cabral, and J. L. Jordan-Sweet 共unpublished兲. 24 P. A. Dearnley, Surf. Eng. 1, 43 共1985兲. 25 H. Ljuncrantz, C. Engstro¨m, L. Hultman, M. Olsson, X. Chu, M. S. Wong, and W. D. Sproul, J. Vac. Sci. Technol. A 16, 3104 共1998兲. 26 M. Nordin, M. Larsson, and S. Hogmark, Surf. Coat. Technol. 106, 234 共1998兲. 27 ˚. O ¨ stlund, and S. T. Selinder, M. E. Sjo¨strand, M. Nordin, M. Larsson, A Hogmark, Surf. Coat. Technol. 105, 51 共1998兲. 28 A. J. Perry, J. Vac. Sci. Technol. A 6, 2140 共1988兲. 29 Atomic Rev. Spec. Issue, edited by K. L. Komarek 共Int. Atomic Energy Agency, Vienna, 1993兲, Vol. 9. 30 R. J. Wasilewski, J. Inst. Met. 83, 94 共1954兲. 31 V. D. Repkin, G. V. Kurtukov, A. A. Kornilov, and V. V. Bespalov, Metalloterm, Protsessy Khim. Metall. 320 共1971兲. 32 P. Lamparter, S. Steeb, and A. Gokelberger, High Temp.-High Press. 3, 727 共1971兲. 33 M. H. Bradbury and H. Matzke, Nucl. Sci. Eng. 84, 291 共1993兲. 13