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Evolution of stress and microstructure in NiFe (20 wt.%) thin films during annealing
Thin Solid Films 385 Ž2001. 225᎐229
Evolution of stress and microstructure in NiFe Ž 20 wt.%. thin films during annealing U W. Bruckner , J. Thomas, C.M. Schneider ¨
Institute of Solid State and Materials Research Dresden, P.O. Box 270016, D-01171 Dresden, Germany Received 23 June 2000; received in revised form 3 December 2000; accepted 14 December 2000
Abstract The stress evolution on NiFe Ž20 wt.%. thin films with thickness of 170 nm was studied during thermal cycles up to maximal 530⬚C. The results are correlated to microstructural analyses carried out on samples cycled to various maximum temperatures by means of transmission electron microscopy, X-ray diffraction, and Auger electron spectroscopy. In addition, resistance-vs.-temperature measurements yielded more information about the microstructural evolution. The atomic rearrangement in grain boundaries is held responsible for a first irreversible stress contribution appearing between 150 and 300⬚C. A second, more distinct irreversible tensile stress contribution of approximately 400 MPa occurs between 300 and 400⬚C. It is explained by abnormal grain growth which can be observed in the same temperature range. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Stress; Resistivity; Structural properties; Alloys
1. Introduction Thin films made of NiFe with approximately 20 wt.% Fe ŽPermalloy. have superior soft magnetic properties, high anisotropic magnetoresistance, and low magnetostriction w1,2x. Therefore, they are, e.g. used in magnetic reading heads and in magnetic multilayers with giant magnetoresistance properties. Stress affects the reliability and influences the magnetic properties w3x. Furthermore, the knowledge of the stress evolution in single NiFe films during annealing is also interesting in comparison to the stress evolution in thin-film configurations, e.g. CurNiFe multilayers w4x or NiFerCurNiFe trilayers w5x. However, it should be taken into account that the mechanisms of strain reduction and relief in these very thin films are not necessarily the same as for thick films. The present article focuses on the stress evolution in U
Corresponding author. Tel.: q49-351-4659-532; fax: q49-3514659-745. .. E-mail address: [email protected] ŽW. Bruckner ¨
NiFe Ž20 wt.%. films with a thickness of 170 nm during thermal cycling. For comparison some films with 90 nm thickness were also investigated. In order to correlate stress and microstructure, microstructural analyses were made on samples after stress measurements up to various maximum temperatures. The microstructure was mainly studied by transmission electron microscopy ŽTEM. and additionally by X-ray diffraction ŽXRD. and Auger electron spectroscopy ŽAES.. Furthermore, resistance-vs.-temperature measurements were carried out in order to obtain additional information about the microstructural evolution. 2. Experimental details The NiFe films were deposited at room temperature ŽRT. by magnetron sputtering from a NiFe Ž20 wt.%. target onto rotating substrates. The following sputter conditions were applied: base pressure 1 = 10y6 mbar, sputter pressure 6 = 10y3 mbar Ar, target-substrate distance 93 mm, and sputter power 210 W. The substrates were oxidized 3-inch silicon wafers with Ž100.
0040-6090r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 1 . 0 0 7 5 4 - 4
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
226
orientation and 0.6 m oxide thickness. The film thickness was measured using a Dektak stylus profiler. It amounted to 171 nm for the primarily investigated films and 90 nm for the thinner ones. The stress in the as-deposited film was determined using a FLEXUS FLX-2410 measurement system. The stress-vs.-temperature behaviour was measured on 7mm-wide and 60-mm-long slabs of the coated silicon substrate. We employed a sensitive laser-optical stressmeasurement system described elsewhere w6x. The measurements were carried out under vacuum Ž- 1 = 10y5 mbar. in order to avoid oxidation of the films during the thermal treatment. The stress evolution was studied during thermal cycling up to a maximum temperature of approximately 520⬚C. For microstructural investigations, individual samples were cycled to maximum temperatures Tmax of approximately 200, 300, 360 and 420⬚C. The heating and cooling rate was 4 Krmin. In the cooling branch, a rate of 4 Krmin could be realized only down to 200⬚C due to the heat capacity of the apparatus. The film stress was determined from Stoney’s equation s
Es ts2 1 1 y r0 6 Ž 1 y s . t r
ž
/
Ž1.
where EsrŽ1 y s . is the biaxial modulus of the substrate Ž180.5 GPa., ts and t denote the substrate thickness Ž375 m. and film thickness, and r and r 0 are the substrate-curvature radii after and before deposition, respectively. The measurements of the electric resistance R were performed on samples with a four-terminal geometry in the furnace of the FLEXUS stress-measurement system as described elsewhere w7x. The atmosphere was defined by a flowing N2rH 2 Ž5 vol.%. gas mixture with a flow rate of 120 lrh. In order to resolve irreversible changes, the temperature-vs.-time profiles were chosen in analogy to the stress measurements with various maximum temperatures. XRD studies were carried out using a Philips-X’Pert diffractometer and CoK ␣ radiation. A Seifert XRD 3000 diffractometer equipped with a texture cradle was used for texture investigations. Concentration-depth profiles were recorded using a PHI 600 Auger microprobe with primary electrons of 10 keV Ž100 nA. incident at 30⬚ to the surface normal. For depth profiling, argon ions of 1.5 keV at 60⬚ to the surface normal were used. TEM investigations were done on both plan-view and cross-sectional specimens. The preparation was carried out by conventional techniques Žmechanical thinning and mechanical and ion dimpling.. For TEM observations, a Philips CM20FEG was employed.
Grain-size distributions were determined by means of the Digital Micrograph ŽGATAN. software. 3. Results The stress-vs.-temperature curves ŽFig. 1. show thermal stress due to the difference of the thermal expansion of film and substrate up to 150⬚C. Subsequently, irreversible tensile stress contributions occur which are most distinct between 300 and 370⬚C. Between 370 and 400⬚C, thermal stress is dominant again. Above 420⬚C, we assume that stress relaxation caused by plastic flow leads to the gradually decreasing compressive stress. The ŽT . curve shows a hysteresis above 250᎐300⬚C during a second thermal cycle thereby confirming the plastic flow in this temperature range. Fig. 2 shows the relative resistance change ⌬ RrR 0 s Ž R y R 0 .rR 0 during the temperature-time profiles, where R 0 is the RT value. The resistivity at RT in the as-deposited state amounts to 0 s 27 = 10y8 ⍀ m. A first irreversible resistivity decrease occurs for Tmax s 200 and 300⬚C. A further, strong decrease is observed between 330 and 380⬚C and is correlated to the ŽT . curves in Fig. 1. After the thermal cycle up to 530⬚C, the resistivity at RT reaches a value of 19 = 10y8 ⍀ m. This value corresponds closely to the bulk value of NiFe Ž20 wt.%., i.e. 0,bulk s 16 = 10y8 ⍀ m w8x. The thin-film resistivity approaching the bulk value points to defect annealing and grain coarsening during the heat treatment, as discussed below. For the 90-nm-thick NiFe Ž20 wt.%. films, very simi-
Fig. 1. Stress-vs.-temperature curves of 171-nm-thick NiFe films during heating and cooling at 4 Krmin in vacuum. The thinner lines demonstrate the irreversible stress contributions during cycles up to various maximum temperatures Tma x , as indicated. After the measurement, the samples were used for microstructural investigations.
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
Fig. 2. Relative resistance change vs. temperature of 171-nm-thick NiFe films during heating and cooling at 4 Krmin in N2rH 2 Ž5 vol.%. gas mixture. The thinner lines demonstrate the irreversible stress contributions during cycles up to various maximum temperatures Tma x , as indicated.
lar ŽT . and RŽT . curves were obtained. This indicates that the film thickness in the considered range does not affect the stress and microstructure evolution. The diffraction patterns of all samples show only reflections of an fcc solid solution without additional reflections. Superlattice reflections of FeNi 3 , which is the equilibrium phase below 517⬚C w9x, could not be detected. However, even if the film structure was chemically ordered, the small difference in the scattering amplitudes of Fe and Ni atoms renders the observation of ordering very difficult. The lattice parameters for the stress-free state, astress-free , are determined from the measured lattice parameters, ameas , using the Ž220. reflections and astressyfree s ameasr Ž 1 y 2 rE .
Ž2.
where E is Young’s modulus and is Poisson’s ratio of thin film. The isotropic bulk value was used for NiFe Ž20 wt.%., i.e. Es 210 GPa w10x, which nearly equals the value for Ni w11x, and s 0.31, as measured for bulk Ni w11x. The stress-free lattice parameter amounts to 0.3542 nm for the as-deposited sample. It does not change for the annealed samples within the experimental error of the technique Ž ⌬ as "0.0005 nm.. The value obtained corresponds closely to the lattice parameter of bulk NiFe Ž20 wt.%., as 0.3541 nm w12x. The texture was investigated on the as-deposited sample. The microstructure has a predominant Ž111. fibre texture parallel to the film normal Žfull width at half maximums 16⬚.. Deviations concern a preference
227
of 1004 planes which are inclined by approximately 20⬚ to the surface normal. From the AES concentration-depth profiles, the atomic-concentration ratio of Fe to Ni was found to be in accordance with the target value. For the samples with Tmax G 360⬚C, surface oxidation showed up despite the vacuum annealing. The concentrations in the surface layer are approximately 40᎐50 at.% Ni, 20᎐30 at.% Fe and 20᎐30 at.% O. Electron diffraction in connection with TEM investigations Žsee below. shows additional oxide reflections for the near-surface regions of the samples with Tmax G 360⬚C. The additional reflections correspond to a spinel lattice with a lattice parameter of approximately 0.84 nm and result from Fe᎐Ni oxides, as reported w13x. Such diffraction spots were also observed for the oxidation of Permalloy w14x. For understanding the ŽT . and RŽT . curves, the information from TEM studies as to the grain morphology is important. Plan-view images of all samples and a cross-sectional image of the specimen with Tmax s 360⬚C were recorded. From the plan-view images near the surface, average grain dimensions of 11 nm were obtained in the near-surface region for the asdeposited sample and the samples with Tmax s 200 and 300⬚C. TEM images of the as-deposited sample and the sample with Tmax s 300⬚C are shown in Fig. 3. It can be seen that the grain boundaries in the annealed sample are straightened and more often hexagonal cross-section grains form. For Tmax s 360⬚C, large grains Žaverage size ds 490 nm. were found besides the small grains ŽFig. 4.. The average size of the small grains was determined as 16 nm, i.e. these grains have slightly grown in comparison to Tmax F 300⬚C. The cross-sectional image showed columnar morphology of the small grains. The appearance of the large grains indicates abnormal grain growth, as recently also reported for CuNi thin films w15x. After the 530⬚C cycle, the abnormal grain growth is complete. The large grains are heavily twinned and of an averaged grain diameter of ds 490 nm. It is not yet understood why the matrix is stagnant and only some grains grow abnormally. 4. Discussion In the following, the volume reduction and the corresponding stress development due to grain growth will be estimated. The morphology of the Žoften twinned. grains is on the whole columnar perpendicular to the film plane and round in the film plane in both the as-deposited and annealed state. Therefore, hexagonal columns are used for the following estimation. Then, the grain-boundary area per unit volume, agb , is given by agb f
2 d
Ž3.
228
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
¨gb ,exc s agb ⌬Vexc f
2⌬Vexc d
Ž4.
Assuming the case that the excess volume in the grain boundaries is translated into lateral shrinkage in the film plane, then the volume shrinkage results in an in-plane strain of 1 1 1 grain growth f ¨gb ,exc f ⌬Vexc y 2 d0 d1
ž
/
Ž5.
where d 0 and d1 are the average grain diameters before and after grain growth, respectively. The corresponding biaxial tensile stress contribution is grain growth s
Fig. 3. Plan-view TEM images of the as-deposited sample Ža. and of the sample after thermal cycling up to 300⬚C Žb.. The comparison of both images indicates to the evolution of hexagonal cross-section grains during the thermal treatment.
where d is the average grain diameter. If ⌬Vexc is the excess volume per unit area of the grain boundaries, then the total excess volume per unit volume, ¨gb,exc , is
Fig. 4. Plan-view TEM image of the sample after thermal cycling up to 360⬚C. Abnormal grain growth occurs consuming the small columnar grains, which do not undergo significant morphological changes.
E 1 y grain growth
Ž6.
According to literature, ⌬Vexc amounts to approximately 0.02= 10y9 m3rm 2 for an equilibrium Al-⌺3 grain boundary w16x. By taking ⌬Vexc s 0.03= 10y9 m3rm 2 for our non-equilibrium grain boundaries and d 0 s 11 nm and d1 s 490 nm, grain growth is estimated by Eqs. Ž5. and Ž6. to be approximately 800 MPa which is somewhat higher than the irreversible stress changes found between 300 and 400⬚C in Fig. 1 Ž400 MPa.. As to the accuracy of the above estimation, one has to take into account that: Ži. ⌬Vexc is only an estimated value; Žii. a portion of the excess volume in the grain boundaries may be translated into a reduction of the film thickness without stress generation; and Žiii. stress relaxation due to plastic deformation may also occur. Due to the low film stress in the considered temperature range, remarkable plastic deformation is not expected. Thus, we assume that densification by grain growth leads to both lateral and thickness shrinkage and that the lateral shrinkage is the main reason for the irreversible stress change between 300 and 400⬚C. Unfortunately, if, e.g. 50% of the densification goes back to thickness reduction, the thickness reduction amounts to approximately 1 nm, which is below the detection limit. Up to 250⬚C, an irreversible tensile stress of 100 MPa is developed. Since only grain-boundary diffusion can be effective at this temperature, we speculate, although we have no clear evidence, that the 100 MPa results from a reduction by approximately 20% of the excess volume in the grain boundaries, estimated according to the above model. Such a value may be expected for densification by atomic rearrangement in non-equilibrium grain boundaries. This may be grainboundary relaxation Žatomic rearrangement without grain-boundary diffusion. or the evolution of straight 120⬚ grain boundaries by grain-boundary diffusion or migration, as indicated in Fig. 3a,b. The resistance decrease in the considered temperature range corre-
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
sponds qualitatively to this mechanism, too, which was already reported for CuNi w17x. In conclusion, densification by atomic rearrangement in grain-boundaries ŽRTy 250⬚C. and abnormal grain growth Žespecially between approx. 300 and 400⬚C. may be the reason for the generation of tensile stress in NiFe Ž20 wt.%. films. Acknowledgements The authors are indebted to L. van Loyen for film preparation, W. Pitschke and M. Hecker for XRD measurements, S. Baunack for AES studies and R. Vogel and B. Arnold for technical assistance. References w1x C. Nishimura, Y. Nagai, K. Yanagisawa, T. Toshima, IEEE Trans. Magn. 23 Ž1987. 2728. w2x R.M. Valletta, G. Guthmiller, G. Gorman, J. Vac. Sci. Technol. A 9 Ž1991. 2093. w3x D.-H. Han, J.-G. Zhu, J.H. Judy, J.M. Siversen, J. Appl. Phys. 81 Ž1997. 4519. w4x L. van Loyen, D. Elefant, D. Tietjen, M. Hecker, J. Thomas, C.M. Schneider, J. Appl. Phys. 87 Ž2000. 4852.
229
w5x W. Bruckner, S. Baunack, M. Hecker, J.-I. Monch, L. van ¨ ¨ Loyen, C.M. Schneider, Appl. Phys. Lett. 77 Ž2000. 358. w6x V. Weihnacht, W. Bruckner, C.M. Schneider, Rev. Sci. Instrum. ¨ 71 Ž2000. 4479. w7x W. Bruckner, G. Sobe, H. Grießmann, S. Baunack, G. Reiss, ¨ Thin Solid Films 261 Ž1995. 90. w8x K. Schroder, CRC Handbook of Electrical Resistivities of Bi¨ nary Metallic Alloys, CRC, Boca Raton, FL, 1983, p. 247. w9x Th.B. Massalski, Binary Alloy Phase Diagrams 2, ASM International, Metals Park, OH, 1990, p. 1735. w10x Gmelins Handbuch der anorganischen Chemie 59 part D2, Verlag Chemie, Weinheim, 1959, p. 391. w11x F. Kohlrausch, Praktische Physik 3, B.G. Teubner, Stuttgart, 1996, p. 300. w12x R.M. Bozorth, Ferromagnetism, D. van Nostrand, Princeton, NJ, 1968, p. 104. w13x Gmelins Handbuch der anorganischen Chemie 59 part D2, Verlag Chemie, Weinheim, 1959, p. 116. w14x M. Kitada, J. Mater. Sci. 26 Ž1991. 4150. w15x W. Bruckner, V. Weihnacht, W. Pitschke, J. Thomas, S. Bau¨ nack, J. Mater. Res. 15 Ž2000. 1062. w16x A.P. Sutton, R.W. Ballufi, Interfaces in Crystalline Materials, Clarendon Press, Oxford, 1995, p. 353. w17x W. Bruckner, W. Pitschke, S. Baunack, J. Thomas, J. Mater. ¨ Res. 14 Ž1999. 1286.
Evolution of stress and microstructure in NiFe Ž 20 wt.%. thin films during annealing U W. Bruckner , J. Thomas, C.M. Schneider ¨
Institute of Solid State and Materials Research Dresden, P.O. Box 270016, D-01171 Dresden, Germany Received 23 June 2000; received in revised form 3 December 2000; accepted 14 December 2000
Abstract The stress evolution on NiFe Ž20 wt.%. thin films with thickness of 170 nm was studied during thermal cycles up to maximal 530⬚C. The results are correlated to microstructural analyses carried out on samples cycled to various maximum temperatures by means of transmission electron microscopy, X-ray diffraction, and Auger electron spectroscopy. In addition, resistance-vs.-temperature measurements yielded more information about the microstructural evolution. The atomic rearrangement in grain boundaries is held responsible for a first irreversible stress contribution appearing between 150 and 300⬚C. A second, more distinct irreversible tensile stress contribution of approximately 400 MPa occurs between 300 and 400⬚C. It is explained by abnormal grain growth which can be observed in the same temperature range. 䊚 2001 Elsevier Science B.V. All rights reserved. Keywords: Stress; Resistivity; Structural properties; Alloys
1. Introduction Thin films made of NiFe with approximately 20 wt.% Fe ŽPermalloy. have superior soft magnetic properties, high anisotropic magnetoresistance, and low magnetostriction w1,2x. Therefore, they are, e.g. used in magnetic reading heads and in magnetic multilayers with giant magnetoresistance properties. Stress affects the reliability and influences the magnetic properties w3x. Furthermore, the knowledge of the stress evolution in single NiFe films during annealing is also interesting in comparison to the stress evolution in thin-film configurations, e.g. CurNiFe multilayers w4x or NiFerCurNiFe trilayers w5x. However, it should be taken into account that the mechanisms of strain reduction and relief in these very thin films are not necessarily the same as for thick films. The present article focuses on the stress evolution in U
Corresponding author. Tel.: q49-351-4659-532; fax: q49-3514659-745. .. E-mail address: [email protected] ŽW. Bruckner ¨
NiFe Ž20 wt.%. films with a thickness of 170 nm during thermal cycling. For comparison some films with 90 nm thickness were also investigated. In order to correlate stress and microstructure, microstructural analyses were made on samples after stress measurements up to various maximum temperatures. The microstructure was mainly studied by transmission electron microscopy ŽTEM. and additionally by X-ray diffraction ŽXRD. and Auger electron spectroscopy ŽAES.. Furthermore, resistance-vs.-temperature measurements were carried out in order to obtain additional information about the microstructural evolution. 2. Experimental details The NiFe films were deposited at room temperature ŽRT. by magnetron sputtering from a NiFe Ž20 wt.%. target onto rotating substrates. The following sputter conditions were applied: base pressure 1 = 10y6 mbar, sputter pressure 6 = 10y3 mbar Ar, target-substrate distance 93 mm, and sputter power 210 W. The substrates were oxidized 3-inch silicon wafers with Ž100.
0040-6090r01r$ - see front matter 䊚 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 4 0 - 6 0 9 0 Ž 0 1 . 0 0 7 5 4 - 4
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
226
orientation and 0.6 m oxide thickness. The film thickness was measured using a Dektak stylus profiler. It amounted to 171 nm for the primarily investigated films and 90 nm for the thinner ones. The stress in the as-deposited film was determined using a FLEXUS FLX-2410 measurement system. The stress-vs.-temperature behaviour was measured on 7mm-wide and 60-mm-long slabs of the coated silicon substrate. We employed a sensitive laser-optical stressmeasurement system described elsewhere w6x. The measurements were carried out under vacuum Ž- 1 = 10y5 mbar. in order to avoid oxidation of the films during the thermal treatment. The stress evolution was studied during thermal cycling up to a maximum temperature of approximately 520⬚C. For microstructural investigations, individual samples were cycled to maximum temperatures Tmax of approximately 200, 300, 360 and 420⬚C. The heating and cooling rate was 4 Krmin. In the cooling branch, a rate of 4 Krmin could be realized only down to 200⬚C due to the heat capacity of the apparatus. The film stress was determined from Stoney’s equation s
Es ts2 1 1 y r0 6 Ž 1 y s . t r
ž
/
Ž1.
where EsrŽ1 y s . is the biaxial modulus of the substrate Ž180.5 GPa., ts and t denote the substrate thickness Ž375 m. and film thickness, and r and r 0 are the substrate-curvature radii after and before deposition, respectively. The measurements of the electric resistance R were performed on samples with a four-terminal geometry in the furnace of the FLEXUS stress-measurement system as described elsewhere w7x. The atmosphere was defined by a flowing N2rH 2 Ž5 vol.%. gas mixture with a flow rate of 120 lrh. In order to resolve irreversible changes, the temperature-vs.-time profiles were chosen in analogy to the stress measurements with various maximum temperatures. XRD studies were carried out using a Philips-X’Pert diffractometer and CoK ␣ radiation. A Seifert XRD 3000 diffractometer equipped with a texture cradle was used for texture investigations. Concentration-depth profiles were recorded using a PHI 600 Auger microprobe with primary electrons of 10 keV Ž100 nA. incident at 30⬚ to the surface normal. For depth profiling, argon ions of 1.5 keV at 60⬚ to the surface normal were used. TEM investigations were done on both plan-view and cross-sectional specimens. The preparation was carried out by conventional techniques Žmechanical thinning and mechanical and ion dimpling.. For TEM observations, a Philips CM20FEG was employed.
Grain-size distributions were determined by means of the Digital Micrograph ŽGATAN. software. 3. Results The stress-vs.-temperature curves ŽFig. 1. show thermal stress due to the difference of the thermal expansion of film and substrate up to 150⬚C. Subsequently, irreversible tensile stress contributions occur which are most distinct between 300 and 370⬚C. Between 370 and 400⬚C, thermal stress is dominant again. Above 420⬚C, we assume that stress relaxation caused by plastic flow leads to the gradually decreasing compressive stress. The ŽT . curve shows a hysteresis above 250᎐300⬚C during a second thermal cycle thereby confirming the plastic flow in this temperature range. Fig. 2 shows the relative resistance change ⌬ RrR 0 s Ž R y R 0 .rR 0 during the temperature-time profiles, where R 0 is the RT value. The resistivity at RT in the as-deposited state amounts to 0 s 27 = 10y8 ⍀ m. A first irreversible resistivity decrease occurs for Tmax s 200 and 300⬚C. A further, strong decrease is observed between 330 and 380⬚C and is correlated to the ŽT . curves in Fig. 1. After the thermal cycle up to 530⬚C, the resistivity at RT reaches a value of 19 = 10y8 ⍀ m. This value corresponds closely to the bulk value of NiFe Ž20 wt.%., i.e. 0,bulk s 16 = 10y8 ⍀ m w8x. The thin-film resistivity approaching the bulk value points to defect annealing and grain coarsening during the heat treatment, as discussed below. For the 90-nm-thick NiFe Ž20 wt.%. films, very simi-
Fig. 1. Stress-vs.-temperature curves of 171-nm-thick NiFe films during heating and cooling at 4 Krmin in vacuum. The thinner lines demonstrate the irreversible stress contributions during cycles up to various maximum temperatures Tma x , as indicated. After the measurement, the samples were used for microstructural investigations.
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
Fig. 2. Relative resistance change vs. temperature of 171-nm-thick NiFe films during heating and cooling at 4 Krmin in N2rH 2 Ž5 vol.%. gas mixture. The thinner lines demonstrate the irreversible stress contributions during cycles up to various maximum temperatures Tma x , as indicated.
lar ŽT . and RŽT . curves were obtained. This indicates that the film thickness in the considered range does not affect the stress and microstructure evolution. The diffraction patterns of all samples show only reflections of an fcc solid solution without additional reflections. Superlattice reflections of FeNi 3 , which is the equilibrium phase below 517⬚C w9x, could not be detected. However, even if the film structure was chemically ordered, the small difference in the scattering amplitudes of Fe and Ni atoms renders the observation of ordering very difficult. The lattice parameters for the stress-free state, astress-free , are determined from the measured lattice parameters, ameas , using the Ž220. reflections and astressyfree s ameasr Ž 1 y 2 rE .
Ž2.
where E is Young’s modulus and is Poisson’s ratio of thin film. The isotropic bulk value was used for NiFe Ž20 wt.%., i.e. Es 210 GPa w10x, which nearly equals the value for Ni w11x, and s 0.31, as measured for bulk Ni w11x. The stress-free lattice parameter amounts to 0.3542 nm for the as-deposited sample. It does not change for the annealed samples within the experimental error of the technique Ž ⌬ as "0.0005 nm.. The value obtained corresponds closely to the lattice parameter of bulk NiFe Ž20 wt.%., as 0.3541 nm w12x. The texture was investigated on the as-deposited sample. The microstructure has a predominant Ž111. fibre texture parallel to the film normal Žfull width at half maximums 16⬚.. Deviations concern a preference
227
of 1004 planes which are inclined by approximately 20⬚ to the surface normal. From the AES concentration-depth profiles, the atomic-concentration ratio of Fe to Ni was found to be in accordance with the target value. For the samples with Tmax G 360⬚C, surface oxidation showed up despite the vacuum annealing. The concentrations in the surface layer are approximately 40᎐50 at.% Ni, 20᎐30 at.% Fe and 20᎐30 at.% O. Electron diffraction in connection with TEM investigations Žsee below. shows additional oxide reflections for the near-surface regions of the samples with Tmax G 360⬚C. The additional reflections correspond to a spinel lattice with a lattice parameter of approximately 0.84 nm and result from Fe᎐Ni oxides, as reported w13x. Such diffraction spots were also observed for the oxidation of Permalloy w14x. For understanding the ŽT . and RŽT . curves, the information from TEM studies as to the grain morphology is important. Plan-view images of all samples and a cross-sectional image of the specimen with Tmax s 360⬚C were recorded. From the plan-view images near the surface, average grain dimensions of 11 nm were obtained in the near-surface region for the asdeposited sample and the samples with Tmax s 200 and 300⬚C. TEM images of the as-deposited sample and the sample with Tmax s 300⬚C are shown in Fig. 3. It can be seen that the grain boundaries in the annealed sample are straightened and more often hexagonal cross-section grains form. For Tmax s 360⬚C, large grains Žaverage size ds 490 nm. were found besides the small grains ŽFig. 4.. The average size of the small grains was determined as 16 nm, i.e. these grains have slightly grown in comparison to Tmax F 300⬚C. The cross-sectional image showed columnar morphology of the small grains. The appearance of the large grains indicates abnormal grain growth, as recently also reported for CuNi thin films w15x. After the 530⬚C cycle, the abnormal grain growth is complete. The large grains are heavily twinned and of an averaged grain diameter of ds 490 nm. It is not yet understood why the matrix is stagnant and only some grains grow abnormally. 4. Discussion In the following, the volume reduction and the corresponding stress development due to grain growth will be estimated. The morphology of the Žoften twinned. grains is on the whole columnar perpendicular to the film plane and round in the film plane in both the as-deposited and annealed state. Therefore, hexagonal columns are used for the following estimation. Then, the grain-boundary area per unit volume, agb , is given by agb f
2 d
Ž3.
228
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
¨gb ,exc s agb ⌬Vexc f
2⌬Vexc d
Ž4.
Assuming the case that the excess volume in the grain boundaries is translated into lateral shrinkage in the film plane, then the volume shrinkage results in an in-plane strain of 1 1 1 grain growth f ¨gb ,exc f ⌬Vexc y 2 d0 d1
ž
/
Ž5.
where d 0 and d1 are the average grain diameters before and after grain growth, respectively. The corresponding biaxial tensile stress contribution is grain growth s
Fig. 3. Plan-view TEM images of the as-deposited sample Ža. and of the sample after thermal cycling up to 300⬚C Žb.. The comparison of both images indicates to the evolution of hexagonal cross-section grains during the thermal treatment.
where d is the average grain diameter. If ⌬Vexc is the excess volume per unit area of the grain boundaries, then the total excess volume per unit volume, ¨gb,exc , is
Fig. 4. Plan-view TEM image of the sample after thermal cycling up to 360⬚C. Abnormal grain growth occurs consuming the small columnar grains, which do not undergo significant morphological changes.
E 1 y grain growth
Ž6.
According to literature, ⌬Vexc amounts to approximately 0.02= 10y9 m3rm 2 for an equilibrium Al-⌺3 grain boundary w16x. By taking ⌬Vexc s 0.03= 10y9 m3rm 2 for our non-equilibrium grain boundaries and d 0 s 11 nm and d1 s 490 nm, grain growth is estimated by Eqs. Ž5. and Ž6. to be approximately 800 MPa which is somewhat higher than the irreversible stress changes found between 300 and 400⬚C in Fig. 1 Ž400 MPa.. As to the accuracy of the above estimation, one has to take into account that: Ži. ⌬Vexc is only an estimated value; Žii. a portion of the excess volume in the grain boundaries may be translated into a reduction of the film thickness without stress generation; and Žiii. stress relaxation due to plastic deformation may also occur. Due to the low film stress in the considered temperature range, remarkable plastic deformation is not expected. Thus, we assume that densification by grain growth leads to both lateral and thickness shrinkage and that the lateral shrinkage is the main reason for the irreversible stress change between 300 and 400⬚C. Unfortunately, if, e.g. 50% of the densification goes back to thickness reduction, the thickness reduction amounts to approximately 1 nm, which is below the detection limit. Up to 250⬚C, an irreversible tensile stress of 100 MPa is developed. Since only grain-boundary diffusion can be effective at this temperature, we speculate, although we have no clear evidence, that the 100 MPa results from a reduction by approximately 20% of the excess volume in the grain boundaries, estimated according to the above model. Such a value may be expected for densification by atomic rearrangement in non-equilibrium grain boundaries. This may be grainboundary relaxation Žatomic rearrangement without grain-boundary diffusion. or the evolution of straight 120⬚ grain boundaries by grain-boundary diffusion or migration, as indicated in Fig. 3a,b. The resistance decrease in the considered temperature range corre-
W. Bruckner et al. r Thin Solid Films 385 (2001) 225᎐229 ¨
sponds qualitatively to this mechanism, too, which was already reported for CuNi w17x. In conclusion, densification by atomic rearrangement in grain-boundaries ŽRTy 250⬚C. and abnormal grain growth Žespecially between approx. 300 and 400⬚C. may be the reason for the generation of tensile stress in NiFe Ž20 wt.%. films. Acknowledgements The authors are indebted to L. van Loyen for film preparation, W. Pitschke and M. Hecker for XRD measurements, S. Baunack for AES studies and R. Vogel and B. Arnold for technical assistance. References w1x C. Nishimura, Y. Nagai, K. Yanagisawa, T. Toshima, IEEE Trans. Magn. 23 Ž1987. 2728. w2x R.M. Valletta, G. Guthmiller, G. Gorman, J. Vac. Sci. Technol. A 9 Ž1991. 2093. w3x D.-H. Han, J.-G. Zhu, J.H. Judy, J.M. Siversen, J. Appl. Phys. 81 Ž1997. 4519. w4x L. van Loyen, D. Elefant, D. Tietjen, M. Hecker, J. Thomas, C.M. Schneider, J. Appl. Phys. 87 Ž2000. 4852.
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w5x W. Bruckner, S. Baunack, M. Hecker, J.-I. Monch, L. van ¨ ¨ Loyen, C.M. Schneider, Appl. Phys. Lett. 77 Ž2000. 358. w6x V. Weihnacht, W. Bruckner, C.M. Schneider, Rev. Sci. Instrum. ¨ 71 Ž2000. 4479. w7x W. Bruckner, G. Sobe, H. Grießmann, S. Baunack, G. Reiss, ¨ Thin Solid Films 261 Ž1995. 90. w8x K. Schroder, CRC Handbook of Electrical Resistivities of Bi¨ nary Metallic Alloys, CRC, Boca Raton, FL, 1983, p. 247. w9x Th.B. Massalski, Binary Alloy Phase Diagrams 2, ASM International, Metals Park, OH, 1990, p. 1735. w10x Gmelins Handbuch der anorganischen Chemie 59 part D2, Verlag Chemie, Weinheim, 1959, p. 391. w11x F. Kohlrausch, Praktische Physik 3, B.G. Teubner, Stuttgart, 1996, p. 300. w12x R.M. Bozorth, Ferromagnetism, D. van Nostrand, Princeton, NJ, 1968, p. 104. w13x Gmelins Handbuch der anorganischen Chemie 59 part D2, Verlag Chemie, Weinheim, 1959, p. 116. w14x M. Kitada, J. Mater. Sci. 26 Ž1991. 4150. w15x W. Bruckner, V. Weihnacht, W. Pitschke, J. Thomas, S. Bau¨ nack, J. Mater. Res. 15 Ž2000. 1062. w16x A.P. Sutton, R.W. Ballufi, Interfaces in Crystalline Materials, Clarendon Press, Oxford, 1995, p. 353. w17x W. Bruckner, W. Pitschke, S. Baunack, J. Thomas, J. Mater. ¨ Res. 14 Ž1999. 1286.