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Stress and oxidation in CuNi thin films
Thin Solid Films 355±356 (1999) 316±321 www.elsevier.com/locate/tsf
Stress and oxidation in CuNi thin ®lms W. BruÈckner*, S. Baunack Institute of Solid State and Materials Research Dresden, P.O. Box 270016, D-01171 Dresden, Germany
Abstract The stress evolution of a 400 nm thick Cu0.57Ni0.42Mn0.01 resistive ®lm on oxidized silicon substrate was investigated during a thermal cycle to 5508C using a laser-optical substrate-curvature technique. Cycle-stop-prepared samples were used to clarify the correlation between stress development and oxidation, especially by concentration-depth pro®ling using Auger electron spectroscopy. In the stress curve, the striking feature associated with oxidation is a tensile stress component of about 500 MPa between 300 and 3808C. This stress component is semiquantitatively explained by grain-boundary diffusion of Ni to the surface and the induced contraction of the CuNi(Mn) layer due to the material lost. With progressive oxidation, a CuO/NiO double layer forms, growing without an essential contribution of diffusion within CuNi(Mn) and, therefore, without oxidation-induced stress development. q 1999 Elsevier Science S.A. All rights reserved. Keywords: Stress; Oxidation; Alloys
1. Introduction Thin ®lms of CuNi with about 40 wt.% Ni and sometimes 1 wt.% Mn are used for resistive applications [1±3]. Like the commercial alloy Constantan (CuNi44Mn1), they have a resistivity (r ) of about 50 £ 1028 V m and a small temperature coef®cient of resistance (TCR) of , 50 £ 1026 K 21. After sputtering, the microstructure is stabilized by annealing. During this annealing under inert atmosphere (e.g. argon), (1) surface oxidation occurs due to reaction of CuNi(Mn) with the remaining oxygen in the Ar atmosphere and (2) tensile stress is developing due to atomic rearrangement in the ®lm interior. Whereas the as-sputtered ®lm is under tensile stress of about 300 MPa, the stress is about 1000 MPa after heat treatment [4,5]. The high stress is a problem regarding both mechanical stability (®lm cracking and delamination) and electrical stability (resistance degradation associated with plastic deformation) [6]. The present paper focuses on the role of surface oxidation in the stress development. The stress development in Cu0.57Ni0.42Mn0.01 thin ®lms with a thickness of 400 nm was studied during thermal cycling to maximum 5508C (Section 3.1). By cycle-stop preparation (preparation of samples cycled to lower maximum temperatures, Tmax), specimens were produced for characterization of the oxidation especially by Auger electron spectroscopy (AES, Section 3.2). As the non-conducting oxide layer reduces * Corresponding author. Tel.: 149-351-4659-532; fax: 149-351-4659537. E-mail address: [email protected] (W. BruÈckner)
the thickness of the conducting CuNi layer, additional results concerning the thickness of the oxidized layer were obtained by in-situ electrical-resistance measurements (Section 3.3). The development of a strong tensile-stress contribution between 300 and 3808C during the thermal ramp is semiquantitatively explained by grain-boundary diffusion of Ni to the surface and formation of an NiO layer (Section 3.4). 2. Experimental details The ®lms were deposited by dc-magnetron sputtering without external heating. The 6-in. sputter target had a nominal copper content of 59.0 wt.% (57.5 at.%) and an addition of 1.0 wt.% (1.1 at.%) Mn. The substrates were oxidized 3-in. silicon wafers with an oxide-layer thickness of 1.2 mm. The sputter conditions of the commercial sputtering system were: base pressure 1 £ 1024 Pa, working pressure 0.2 Pa Ar, target-substrate distance 55 mm. The sputtering power amounted to 3.2 kW. The deposition rate was about 80 nm/min. The substrate-temperature increase due to ion bombardment during the sputter process was , 30 K, as determined by temperature-indicating colored pencils. The ®lm thickness was measured by a Dektak stylus pro®ler on a wafer with resistance-measuring structures. Averaged over the wafer area, it was (381 ^ 7) nm. The chemical analysis of the ®lm material by titrimetry and atomic absorption analysis was (58:7 ^ 0:2) wt.% Cu, (40:5 ^ 0:2) wt.% Ni, and (0:8 ^ 0:1) wt.% Mn. Thus, the ®lm composition nearly equals the nominal target composi-
0040-6090/99/$ - see front matter q 1999 Elsevier Science S.A. All rights reserved. PII: S00 40-6090(99)0044 8-4
W. BruÈckner, S. Baunack / Thin Solid Films 355±356 (1999) 316±321
tion. As it has been determined recently by transmission electron microscopy (TEM), the grain morphology in the as-deposited state is columnar [5,7]. The lateral grain dimension is about 20 nm. The biaxial stress (s ) was determined as function of temperature by laser-optical substrate-curvature measurements using a Flexus FLX-2410 system. It is obtained by the well-known Stoney's equation [8] revised for biaxial stress Es ts2 1 1 ÿ 2 s r0 6 1 2 ns t r
1
where Es = 1 2 ns is the biaxial modulus of the substrate, ts and t are the thicknesses of the substrate and the ®lm, and r and r0 denote the curvature radii after and before ®lm deposition, respectively. The (100) oriented Si substrates have a thickness of 375 mm and a biaxial modulus of 180.5 GPa. The force per unit width s ´t is a quantity which is independent of ®lm thickness. It is suitable for description of stress in ®lms having changing thickness, e.g. due to oxidation. The heating and cooling rate during the thermal cycle was 4 K/min in Ar. The Ar gas had an oxygen contamination of 3 £ 1026 vol.%. The ¯ow rate of Ar was 0.15 m 3/h. Auger electron spectra (AES) were obtained in a Perkin± Elmer PHI660 Auger Microprobe operated at 10 kV and 50 nA. The sample normal was tilted 308 to the electron beam. For depth pro®ling, a scanned beam of Ar ions [4 keV, 1.4 A/m 2, in the surface region (corresponding to a normalized sputter dose in our measuring curves of , 0.05) for a better resolution, 0.3 A/m 2] impinging under 558 to the surface normal was used. Depth pro®les for Ni, Cu, Mn, O, Si, and C were recorded in the N £ N E mode. The magnitude of the Mn signal was of the order of the background noise (corresponding to about 1 wt.% Mn). Due to the overlap of Ni and Cu LMM Auger transitions, the weak L2,3MM transition was used for Ni. The spectra recorded during depth pro®ling were evaluated by principal component analysis [9] which made it possible to separate the Ni(L2,3MM) transition from spectral features belonging to Cu and, thus, to improve the detection limit for Ni. The concentrations were determined using the model of homogeneously-mixed elements with the composition of the as-deposited layer as standard. The electrical resistance (R) was determined as function of temperature with a four-terminal resistance-measuring structure arranged within the Flexus furnace. The measuring structure was contacted by tungsten springs. Using a constant-current source and a voltmeter with a high input resistance, the contact resistance does not affect the resistance measurement. Temperature rate and atmosphere agree with the data for stress measurements.
317
3. Results and discussion 3.1. Stress development The force per unit width versus temperature curves s ´t(T) measured during both the temperature cycle up to 5508C and the cycle-stop preparation are given in Fig. 1. The right axes give the stress for an assumed constant ®lm thickness (equal to the as-deposited value, t 381 nm). The heating branch starts with stress decrease. This thermal stress (s th) is due to the difference in thermal-expansion coef®cients of ®lm and substrate (a and a s, respectively) according to ÿ E a 2 as d sth 2 dT 12n
2
where E and n are Young's modulus and Poisson's ratio of the ®lm, respectively. The value of ds th/dT, as determined from the slopes of the heating curves in Fig. 1 between room temperature (RT)
Fig. 1. Development of the force per unit width in CuNi(Mn) thin ®lms during temperature cycles (4 K/min) up to (a) 5508C and (b) different maximal temperatures (Tmax, preparation of cycle-stop samples for microstructural characterization) in Ar atmosphere. The biaxial-stress axes refer to the thickness of the as-deposited ®lm. The large stress increase between 300 and 3808C is considered in the present article and ascribed to surface oxidation.
318
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and 1008C, is 2 3.0 MPa/K in accordance with previous results on CuNi(Mn) ®lms [10]. Deviations from thermal stress due to irreversible tensilestress contributions occur already above 1208C. These deviations up to 3008C have been analyzed in a recent paper [5]. They were explained by material densi®cation associated with grain-boundary relaxation. Above 3008C, a strong stress change occurs. An irreversible tensile stress contribution of about 500 MPa is developing. This leads to a transition from compressive to tensile stress. On thicker CuNi(Mn) ®lms (t $ 1000 nm) having tensile stress at 3008C, an analogous behavior was observed with an irreversible increase of the tensile stress [11]. Now, the question is whether the sharp transition from compressive to tensile stress is correlated with oxidation. In order to answer it, stress-temperature measurements of the CuNi(Mn) ®lms were carried out in an N2/H2(5 vol.%) gas mixture with and without air addition by a de®ned leak. The curves in Fig. 2 con®rm such a correlation. It was also shown by investigation of CuNi(Mn) ®lms using a recently constructed vacuum apparatus for stress measurements [12] that the sharp stress change between 300 and 3808C does not occur under vacuum [7]. Thus, the striking feature in the s (T) curve which is associated with oxidation is the development of high tensile stress. The reason for this effect will be discussed in Section 3.4. Above 4008C, the tensile stress is decreasing with about the slope of thermal stress, Eq. (2). Above 4508C, the stress decrease is increased in comparison to thermal stress. We assume that it is caused by stress-relaxation processes, e.g. grain-boundary diffusional ¯ow as occurring in CuNi ®lms [7,13] or by cracking of the growing oxide ®lm. The cooling branch shows similarities to the course of the stress-vs.temperature curves measured under vacuum [7,13]. Differences concerning lower stresses under Ar atmosphere result
Fig. 2. Development of the force per unit width in CuNi(Mn) ®lms during temperature ramps (4 K/min) in an N2/H2(5 vol.%) gas mixture without and with addition of air. The biaxial-stress axis refers to the thickness of the asdeposited ®lm. The experiment con®rms that the large stress increase above 3008C originates from oxidation.
probably from the roughness and brittleness of the surface oxide layer. In the case of a double layer (metallic CuNi and an assumed homogeneous oxide layer), the average ®lm stress (s a), as measured by the substrate-curvature technique, is
sa
tCuNi t s 1 ox sox ttot CuNi ttot
3
where tCuNi, tox, s CuNi, and s ox are the thicknesses of and the stresses within the corresponding layers, respectively, and ttot is the total ®lm thickness. The total ®lm thickness for an oxidized ®lm is higher than the initial value t for an unoxidized ®lm according to ttot 1 2 at 1 bat
4
Eq. (4) describes that in a portion (a) of the ®lm thickness, the layer thickness increases by the so-called Pilling± Bedworth ratio (b ) of the oxidation theory [14]. As the oxide layer grows with increasing temperature, the average force per unit width in the ®lm is given in Figs. 1 and 2. 3.2. AES concentration-depth pro®les The ®lm oxidation was studied by AES on the cycle-stopprepared samples. Selected AES depth pro®les are given in Fig. 3. Because the ®lms exhibit a distinct roughening due to both sputter effects and oxidation, a depth scale was not established. The concentrations are plotted vs. the relative sputter dose necessary for reaching the oxidized silicon substrate. The dose increases with Tmax mainly due to the lower sputter rate in oxides. No ®lm/substrate reaction or interdiffusion was detected. The pro®le for T max 3008C shows only small changes compared to the as-deposited sample. The surface is mainly covered by a very thin skin of nickel oxide. However, there are some small surface cracks. In these regions, as observed by secondary electron imaging, the uppermost layer is rich in copper and carbon. For T max 4508C, the nickel concentration in the CuNi(Mn) is distinctly reduced, and a laterally-homogeneous Ni-oxide layer forms. At T max 5008C, islands of a copper-rich oxide form within the nickel oxide, and at 5508C, the ®lm is covered by a rough double layer of copper oxide on nickel oxide. Such a double oxide was also observed on oxidized bulk CuNi [15±17]. The Ni concentration in the CuNi(Mn) layer decreases with the formation of the oxide layer. For T max 3008C, no concentration shift (DcNi) could be detected. For the samples with T max 4508C and 5508C, the shift amounts to DcNi 25 at.% and 2 7 at.%, respectively. Especially for T max 4508C, this corresponds to the fact that the Ni content is reduced as an NiO surface layer is formed. 3.3. Electrical resistance The resistivity of the as-deposited ®lm was measured to be 52 ^ 1 £ 1028 V m. The relative change in electrical resistance (DR/R0) with V m temperature is shown in Fig. 4,
W. BruÈckner, S. Baunack / Thin Solid Films 355±356 (1999) 316±321
319
Fig. 3. AES concentration-depth pro®les for CuNi(Mn) ®lms (a) in the as-deposited state and (b)±(d) after temperature cycles up to the indicated maximum temperatures (Tmax) in Ar atmosphere. The diagrams show that up to T max 3008C no distinct surface oxidation occurs, for T max 4508C (1) an NiO layer is formed on the top of the ®lm and (2) the Ni concentration within the ®lm is reduced, and for T max 5008C a Cu±O/Ni±O double layer is formed.
where R0 is the RT magnitude. Up to 1208C, an effective TCR of 253 £ 1026 K 21 was found. This value corresponds well to the bulk value (230 £ 1026 K 21) for CuNi with 41.62 at.% Ni [18]. Above 1208C, the resistance decreases more steeply. A sharp change to a resistance increase was observed in Ar at 3808C. This is above the temperature of the steep stress change. We attribute it to oxidation in the stage II (see Section 3.4). For comparison, the DR/R0-vs.-T curve measured in a N2/H2 (5 vol.%) gas mixture is also given in Fig. 4. In such an atmosphere, the steep resistance increase between 300 and 3808C does not occur. This characterizes the described feature in the curve under Ar atmosphere as oxidation-induced. Below 3808C, the relative
resistance change in N2/H2 atmosphere is slightly smaller than that in Ar atmosphere, showing slight oxidation at this temperatures under Ar atmosphere. As already previously observed [4], only a thin Mnoxide-rich surface layer is forming during a temperature ramp to 5508C in N2/H2 atmosphere. This fact enables us to consider the measured curve in Fig. 4 for N2/H2 atmosphere as practically without CuNi(Mn) oxidation at least to 5008C. 3.4. Model for the oxidation-induced stress change The oxides NiO and CuO are immiscible [19]. Therefore,
320
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area of the entire ®lm. The thickness of the metallic CuNi(Mn) shrinks, but no stress is developed in the ®lm. Now the stress produced in stage I shall be evaluated. It is assumed that a nickel layer of width wd diffuses out of the perpendicular grain boundaries to the surface. With the outdiffusion of Ni, the lateral shrinkage (1lat) in both the x and y direction of the ®lm plane is
1lat wd =L
5
where L is the lateral grain dimension of assumed square columns. This shrinkage results in a biaxial tensile-stress contribution Ds 1
Fig. 4. Relative change in electrical resistance in CuNi(Mn) ®lms during temperature cycles (4 K/min) in Ar atmosphere and, for comparison, in an N2/H2(5 vol.%) gas mixture. Up to 3808C under Ar atmosphere, only a small resistance contribution occurs due to oxidation, above 3808C, the resistance increases distinctly by oxidation.
only NiO or a double layer CuO/NiO is observed. The formation enthalpy for NiO (240 kJ/mol) is higher than that for CuO (165 kJ/mol) [20]. This is the energetic reason, that NiO forms preferentially up to 4508C. If during oxidation the metal atoms penetrate preferentially the oxide layer, then, in principle, the oxidation continues on the top layer without stress generation. If, vice versa, oxygen atoms penetrate preferentially the oxide layer, then the oxidation occurs on the oxide/metal interface and a compressive stress is produced. The preferential penetration of metal atoms through the growing oxide is typical for metals [21]. The behavior is often more complicated by grain-boundary diffusion and grain-boundary oxidation, connected with microcrack formation and other effects [22,23]. In the present paper, the following model is considered for the stress evolution during oxidation. In a ®rst stage (300±4008C, stage I), nickel diffuses to the surface. This mass transport occurs by grain-boundary diffusion. Using the melting temperature T m 1550 K, the temperature range of 300±4008C corresponds to homologous temperatures of T=T m 0:37±0:43, i.e. the typical range for processes associated with grain-boundary diffusion. As the grain boundaries are perpendicular to the ®lm plane in the columnar grain morphology, the mass transport out of grain boundaries to the surface results in an elastic strain in the grains and, thus, in a biaxial tensile-stress contribution. The situation is, therefore, similar to that of grain-boundary diffusional creep (Coble creep) or hillocking in metallization ®lms [24,25]. When at higher temperatures (.4008C) volume diffusion starts becoming effective (stage II), the transport of metal atoms from the ®lm interior to the surface occurs over the
lat E= 1
2 n
6
Since the material of the out-diffused volume is oxidized and the shrinkage in both the x and y direction is 1lat, a portion of 21lat related to the total number of metallic atoms is oxidized. The grown oxide thickness (tox) is given by tox 21lat bt
7
Concerning the change in resistance due to oxidation, the NiO can be considered as insulator. Then, the fractional resistance change is equal to the fractional thickness change of the metal, i.e. DRox 21lat R
8
Now, the lateral shrinkage 1lat, the width wd of the outdiffusing layer of the grain boundaries, the thickness tox of the forming NiO layer, and the resulting resistance change DRox/R shall be evaluated for the case that the developing stress is the observed stress change of 500 MPa between 300 and 3808C. The following material data are used: E= 1 2 n 214 GPa [10], b 1:6 for NiO [14], and L 20 nm. Using Eqs. (5)±(8), one obtains 1lat 0:0023, wd 0:047 nm, tox 3:2 nm, and DRox =R 0:47%, respectively. The value of wd 0:047 nm means that about 20% of an atomic layer (atomic diameter is 0.25 nm) is diffused out of the grain boundaries of the columnar grains. In the evaluation, (i) the stress contribution of the growing NiO layer which is small for thin oxide layers according to Eq. (3) and (ii) the depth inhomogeneity of the ®lm stress due to different diffusion distances from different depths and the attachment of the ®lm onto the substrate, are not taken into account. However, the resistance changes evaluated from the stress changes correspond fairly well to the situation up to 3808C. Therefore, we believe, that our grainboundary-diffusion model is con®rmed for temperatures below 3808C by the given estimation. Above 3808C, the situation is more complicated because of various reasons. ² Lattice diffusion starts being more and more effective. This is con®rmed by the observed changes of the nickel concentration in the CuNi(Mn) sublayer (see Section 3.2)
W. BruÈckner, S. Baunack / Thin Solid Films 355±356 (1999) 316±321
as well as by recently reported lattice-parameter changes correlated to the composition shift [2,5]. Therefore, the oxide layer can grow by diffusion of the metal atoms from the uppermost CuNi(Mn) layers into the oxide without distinct stress generation. ² Stress relaxation by plastic deformation occurs. Such stress relaxation (diffusional ¯ow, mechanical twinning) was observed in CuNi ®lms above 3008C [7,13]. ² Grain coarsening and oxidation lead to microstructural and topographical changes which affect the stress evolution. Although we could not observe grain coarsening in TEM studies on in-plane sections of the cycle-stopprepared sample with T max 4508C, recrystallization may already start in this sample at the ®lm/substrate interface as found during thermal cycling of CuNi(Mn) ®lms in vacuum [7]. ² Stress contributions by (also inhomogeneous) growing oxide layer(s) in¯uence the averaged ®lm stress (force per unit length) according to Eq. (3), but are complicated. Because of these numerous in¯uences to the stress development, the stress-vs.-temperature curves above 3808C can not be analyzed in detail. 4. Conclusions During heat treatment under argon, oxidation of CuNi(Mn) resistive ®lms starts already at 3008C by reaction with the remaining oxygen under formation of a thin NiO layer. Above 4008C, the growth of the NiO layer increases. Above 5008C, a double layer of copper oxide on nickel oxide forms. Between 300 and 4008C, the oxidation in the initial stage gives rise to tensile stress of about 500 MPa. This is the striking feature in the stress-temperature curve associated with oxidation. The tensile-stress contribution is explained semiquantitatively by a model in which nickel diffuses in the columnar structure to the surface by grain-boundary diffusion and forms NiO. The out-diffusion of Ni leads to ®lm contraction and, because of the attachment of the ®lm onto the substrate, to the measured tensile-stress component. One can learn, from the investigations, that the formation of an only a few nanometers thick oxide layer on thin metallic ®lms may possibly lead to high stresses.
321
Acknowledgements The authors are indebted to Th. Knuth at microtech Teltow for ®lm preparation, V. Michel for chemical analysis, J. Thomas for TEM studies, V. Weihnacht for stimulating discussions, and R. Vogel for technical assistance.
References [1] I. Nishino, Y. Ichinose, Y. Sorimachi, I. Tsubata, Int. J. Hybrid Microelectronics 8 (1985) 18. [2] W. BruÈckner, S. Baunack, D. Elefant, G. Reiss, J. Appl. Phys. 79 (1996) 8516. [3] H. Kiyokawa, S. Yonezawa, K. Horita, T. Unishi, M. Takashima, Denki Kagaku 63 (1995) 648. [4] W. BruÈckner, S. Baunack, G. Reiss, G. Leitner, Th. Knuth, Thin Solid Films 258 (1995) 252. [5] W. BruÈckner, W. Pitschke, S. Baunack, J. Thomas, J. Mater. Res. 14 (1999) 1286. [6] W. BruÈckner, J. Edelmann, H. Vinzelberg, G. Reiss, Th. Knuth, Thin Solid Films 280 (1996) 227. [7] W. BruÈckner, V. Weihnacht, W. Pitschke, J. Thomas, S. Baunack, J. Mater. Res. (1999) submitted for publication . [8] G.G. Stoney, Proc. R. Soc. (London) A 82 (1909) 172. [9] S.W. Gaarenstrom, Appl. Surf. Sci. 7 (1981) 7. [10] W. BruÈckner, Phys. Stat. Sol. (a) 148 (1995) K89. [11] W. BruÈckner, F. Macionczyk, G. Reiss, Mater. Res. Soc. Symp. Proc. 436 (1997) 47. [12] W. BruÈckner, H. Grieûmann, Rev. Sci. Instrum. 69 (1998) 3662. [13] W. BruÈckner, V. Weihnacht, J. Appl. Phys. 85 (1999) 3602. [14] M. SchuÈtze, Die Korrosionsschutzwirkung oxidischer Deckschichten unter thermisch-chemisch-mechanischer WerkstoffbeanspruchungGebruÈder Borntraeger, Berlin, 1991, p. 19. [15] K. Tsuji, K. Hirokawa, Surf. Interface Anal. 17 (1991) 819. [16] R. Souchet, M. Lenglet, P. Miche, S. Weber, S. Scherrer, Analusis Lett. 21 (1993) 173. [17] P. Druska, H.-H. Strehblow, Surf. Interface Anal. 23 (1995) 440. [18] B. Svensson, Annalen der Physik, 5 Folge 25 (1936) 263. [19] N.G. Schmahl, F. MuÈller, Z. Anorg. Allg. Chemie 332 (1964) 217. [20] H. Landolt, R. BoÈrnstein, Zahlenwerte und Funktionen. Vol. 2, Springer, Berlin, 1961, p. 179. [21] M.V. Speight, J.E. Harris, Acta Metall. 26 (1978) 1043. [22] F.N. Rhines, J.S. Wolf, Metallurgical Trans. 1 (1970) 1701. [23] W. Jaenicke, S. Leistikow, A. Stadler, J. Electrochem. Soc. 111 (1964) 1031. [24] M.S. Jackson, Che-Yu Li, Acta Metall. 30 (1982) 1993. [25] M.F. Doerner, W.D. Nix, Critical Reviews in Solid State and Materials Sciences 14 (1988) 225.
Stress and oxidation in CuNi thin ®lms W. BruÈckner*, S. Baunack Institute of Solid State and Materials Research Dresden, P.O. Box 270016, D-01171 Dresden, Germany
Abstract The stress evolution of a 400 nm thick Cu0.57Ni0.42Mn0.01 resistive ®lm on oxidized silicon substrate was investigated during a thermal cycle to 5508C using a laser-optical substrate-curvature technique. Cycle-stop-prepared samples were used to clarify the correlation between stress development and oxidation, especially by concentration-depth pro®ling using Auger electron spectroscopy. In the stress curve, the striking feature associated with oxidation is a tensile stress component of about 500 MPa between 300 and 3808C. This stress component is semiquantitatively explained by grain-boundary diffusion of Ni to the surface and the induced contraction of the CuNi(Mn) layer due to the material lost. With progressive oxidation, a CuO/NiO double layer forms, growing without an essential contribution of diffusion within CuNi(Mn) and, therefore, without oxidation-induced stress development. q 1999 Elsevier Science S.A. All rights reserved. Keywords: Stress; Oxidation; Alloys
1. Introduction Thin ®lms of CuNi with about 40 wt.% Ni and sometimes 1 wt.% Mn are used for resistive applications [1±3]. Like the commercial alloy Constantan (CuNi44Mn1), they have a resistivity (r ) of about 50 £ 1028 V m and a small temperature coef®cient of resistance (TCR) of , 50 £ 1026 K 21. After sputtering, the microstructure is stabilized by annealing. During this annealing under inert atmosphere (e.g. argon), (1) surface oxidation occurs due to reaction of CuNi(Mn) with the remaining oxygen in the Ar atmosphere and (2) tensile stress is developing due to atomic rearrangement in the ®lm interior. Whereas the as-sputtered ®lm is under tensile stress of about 300 MPa, the stress is about 1000 MPa after heat treatment [4,5]. The high stress is a problem regarding both mechanical stability (®lm cracking and delamination) and electrical stability (resistance degradation associated with plastic deformation) [6]. The present paper focuses on the role of surface oxidation in the stress development. The stress development in Cu0.57Ni0.42Mn0.01 thin ®lms with a thickness of 400 nm was studied during thermal cycling to maximum 5508C (Section 3.1). By cycle-stop preparation (preparation of samples cycled to lower maximum temperatures, Tmax), specimens were produced for characterization of the oxidation especially by Auger electron spectroscopy (AES, Section 3.2). As the non-conducting oxide layer reduces * Corresponding author. Tel.: 149-351-4659-532; fax: 149-351-4659537. E-mail address: [email protected] (W. BruÈckner)
the thickness of the conducting CuNi layer, additional results concerning the thickness of the oxidized layer were obtained by in-situ electrical-resistance measurements (Section 3.3). The development of a strong tensile-stress contribution between 300 and 3808C during the thermal ramp is semiquantitatively explained by grain-boundary diffusion of Ni to the surface and formation of an NiO layer (Section 3.4). 2. Experimental details The ®lms were deposited by dc-magnetron sputtering without external heating. The 6-in. sputter target had a nominal copper content of 59.0 wt.% (57.5 at.%) and an addition of 1.0 wt.% (1.1 at.%) Mn. The substrates were oxidized 3-in. silicon wafers with an oxide-layer thickness of 1.2 mm. The sputter conditions of the commercial sputtering system were: base pressure 1 £ 1024 Pa, working pressure 0.2 Pa Ar, target-substrate distance 55 mm. The sputtering power amounted to 3.2 kW. The deposition rate was about 80 nm/min. The substrate-temperature increase due to ion bombardment during the sputter process was , 30 K, as determined by temperature-indicating colored pencils. The ®lm thickness was measured by a Dektak stylus pro®ler on a wafer with resistance-measuring structures. Averaged over the wafer area, it was (381 ^ 7) nm. The chemical analysis of the ®lm material by titrimetry and atomic absorption analysis was (58:7 ^ 0:2) wt.% Cu, (40:5 ^ 0:2) wt.% Ni, and (0:8 ^ 0:1) wt.% Mn. Thus, the ®lm composition nearly equals the nominal target composi-
0040-6090/99/$ - see front matter q 1999 Elsevier Science S.A. All rights reserved. PII: S00 40-6090(99)0044 8-4
W. BruÈckner, S. Baunack / Thin Solid Films 355±356 (1999) 316±321
tion. As it has been determined recently by transmission electron microscopy (TEM), the grain morphology in the as-deposited state is columnar [5,7]. The lateral grain dimension is about 20 nm. The biaxial stress (s ) was determined as function of temperature by laser-optical substrate-curvature measurements using a Flexus FLX-2410 system. It is obtained by the well-known Stoney's equation [8] revised for biaxial stress Es ts2 1 1 ÿ 2 s r0 6 1 2 ns t r
1
where Es = 1 2 ns is the biaxial modulus of the substrate, ts and t are the thicknesses of the substrate and the ®lm, and r and r0 denote the curvature radii after and before ®lm deposition, respectively. The (100) oriented Si substrates have a thickness of 375 mm and a biaxial modulus of 180.5 GPa. The force per unit width s ´t is a quantity which is independent of ®lm thickness. It is suitable for description of stress in ®lms having changing thickness, e.g. due to oxidation. The heating and cooling rate during the thermal cycle was 4 K/min in Ar. The Ar gas had an oxygen contamination of 3 £ 1026 vol.%. The ¯ow rate of Ar was 0.15 m 3/h. Auger electron spectra (AES) were obtained in a Perkin± Elmer PHI660 Auger Microprobe operated at 10 kV and 50 nA. The sample normal was tilted 308 to the electron beam. For depth pro®ling, a scanned beam of Ar ions [4 keV, 1.4 A/m 2, in the surface region (corresponding to a normalized sputter dose in our measuring curves of , 0.05) for a better resolution, 0.3 A/m 2] impinging under 558 to the surface normal was used. Depth pro®les for Ni, Cu, Mn, O, Si, and C were recorded in the N £ N E mode. The magnitude of the Mn signal was of the order of the background noise (corresponding to about 1 wt.% Mn). Due to the overlap of Ni and Cu LMM Auger transitions, the weak L2,3MM transition was used for Ni. The spectra recorded during depth pro®ling were evaluated by principal component analysis [9] which made it possible to separate the Ni(L2,3MM) transition from spectral features belonging to Cu and, thus, to improve the detection limit for Ni. The concentrations were determined using the model of homogeneously-mixed elements with the composition of the as-deposited layer as standard. The electrical resistance (R) was determined as function of temperature with a four-terminal resistance-measuring structure arranged within the Flexus furnace. The measuring structure was contacted by tungsten springs. Using a constant-current source and a voltmeter with a high input resistance, the contact resistance does not affect the resistance measurement. Temperature rate and atmosphere agree with the data for stress measurements.
317
3. Results and discussion 3.1. Stress development The force per unit width versus temperature curves s ´t(T) measured during both the temperature cycle up to 5508C and the cycle-stop preparation are given in Fig. 1. The right axes give the stress for an assumed constant ®lm thickness (equal to the as-deposited value, t 381 nm). The heating branch starts with stress decrease. This thermal stress (s th) is due to the difference in thermal-expansion coef®cients of ®lm and substrate (a and a s, respectively) according to ÿ E a 2 as d sth 2 dT 12n
2
where E and n are Young's modulus and Poisson's ratio of the ®lm, respectively. The value of ds th/dT, as determined from the slopes of the heating curves in Fig. 1 between room temperature (RT)
Fig. 1. Development of the force per unit width in CuNi(Mn) thin ®lms during temperature cycles (4 K/min) up to (a) 5508C and (b) different maximal temperatures (Tmax, preparation of cycle-stop samples for microstructural characterization) in Ar atmosphere. The biaxial-stress axes refer to the thickness of the as-deposited ®lm. The large stress increase between 300 and 3808C is considered in the present article and ascribed to surface oxidation.
318
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and 1008C, is 2 3.0 MPa/K in accordance with previous results on CuNi(Mn) ®lms [10]. Deviations from thermal stress due to irreversible tensilestress contributions occur already above 1208C. These deviations up to 3008C have been analyzed in a recent paper [5]. They were explained by material densi®cation associated with grain-boundary relaxation. Above 3008C, a strong stress change occurs. An irreversible tensile stress contribution of about 500 MPa is developing. This leads to a transition from compressive to tensile stress. On thicker CuNi(Mn) ®lms (t $ 1000 nm) having tensile stress at 3008C, an analogous behavior was observed with an irreversible increase of the tensile stress [11]. Now, the question is whether the sharp transition from compressive to tensile stress is correlated with oxidation. In order to answer it, stress-temperature measurements of the CuNi(Mn) ®lms were carried out in an N2/H2(5 vol.%) gas mixture with and without air addition by a de®ned leak. The curves in Fig. 2 con®rm such a correlation. It was also shown by investigation of CuNi(Mn) ®lms using a recently constructed vacuum apparatus for stress measurements [12] that the sharp stress change between 300 and 3808C does not occur under vacuum [7]. Thus, the striking feature in the s (T) curve which is associated with oxidation is the development of high tensile stress. The reason for this effect will be discussed in Section 3.4. Above 4008C, the tensile stress is decreasing with about the slope of thermal stress, Eq. (2). Above 4508C, the stress decrease is increased in comparison to thermal stress. We assume that it is caused by stress-relaxation processes, e.g. grain-boundary diffusional ¯ow as occurring in CuNi ®lms [7,13] or by cracking of the growing oxide ®lm. The cooling branch shows similarities to the course of the stress-vs.temperature curves measured under vacuum [7,13]. Differences concerning lower stresses under Ar atmosphere result
Fig. 2. Development of the force per unit width in CuNi(Mn) ®lms during temperature ramps (4 K/min) in an N2/H2(5 vol.%) gas mixture without and with addition of air. The biaxial-stress axis refers to the thickness of the asdeposited ®lm. The experiment con®rms that the large stress increase above 3008C originates from oxidation.
probably from the roughness and brittleness of the surface oxide layer. In the case of a double layer (metallic CuNi and an assumed homogeneous oxide layer), the average ®lm stress (s a), as measured by the substrate-curvature technique, is
sa
tCuNi t s 1 ox sox ttot CuNi ttot
3
where tCuNi, tox, s CuNi, and s ox are the thicknesses of and the stresses within the corresponding layers, respectively, and ttot is the total ®lm thickness. The total ®lm thickness for an oxidized ®lm is higher than the initial value t for an unoxidized ®lm according to ttot 1 2 at 1 bat
4
Eq. (4) describes that in a portion (a) of the ®lm thickness, the layer thickness increases by the so-called Pilling± Bedworth ratio (b ) of the oxidation theory [14]. As the oxide layer grows with increasing temperature, the average force per unit width in the ®lm is given in Figs. 1 and 2. 3.2. AES concentration-depth pro®les The ®lm oxidation was studied by AES on the cycle-stopprepared samples. Selected AES depth pro®les are given in Fig. 3. Because the ®lms exhibit a distinct roughening due to both sputter effects and oxidation, a depth scale was not established. The concentrations are plotted vs. the relative sputter dose necessary for reaching the oxidized silicon substrate. The dose increases with Tmax mainly due to the lower sputter rate in oxides. No ®lm/substrate reaction or interdiffusion was detected. The pro®le for T max 3008C shows only small changes compared to the as-deposited sample. The surface is mainly covered by a very thin skin of nickel oxide. However, there are some small surface cracks. In these regions, as observed by secondary electron imaging, the uppermost layer is rich in copper and carbon. For T max 4508C, the nickel concentration in the CuNi(Mn) is distinctly reduced, and a laterally-homogeneous Ni-oxide layer forms. At T max 5008C, islands of a copper-rich oxide form within the nickel oxide, and at 5508C, the ®lm is covered by a rough double layer of copper oxide on nickel oxide. Such a double oxide was also observed on oxidized bulk CuNi [15±17]. The Ni concentration in the CuNi(Mn) layer decreases with the formation of the oxide layer. For T max 3008C, no concentration shift (DcNi) could be detected. For the samples with T max 4508C and 5508C, the shift amounts to DcNi 25 at.% and 2 7 at.%, respectively. Especially for T max 4508C, this corresponds to the fact that the Ni content is reduced as an NiO surface layer is formed. 3.3. Electrical resistance The resistivity of the as-deposited ®lm was measured to be 52 ^ 1 £ 1028 V m. The relative change in electrical resistance (DR/R0) with V m temperature is shown in Fig. 4,
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319
Fig. 3. AES concentration-depth pro®les for CuNi(Mn) ®lms (a) in the as-deposited state and (b)±(d) after temperature cycles up to the indicated maximum temperatures (Tmax) in Ar atmosphere. The diagrams show that up to T max 3008C no distinct surface oxidation occurs, for T max 4508C (1) an NiO layer is formed on the top of the ®lm and (2) the Ni concentration within the ®lm is reduced, and for T max 5008C a Cu±O/Ni±O double layer is formed.
where R0 is the RT magnitude. Up to 1208C, an effective TCR of 253 £ 1026 K 21 was found. This value corresponds well to the bulk value (230 £ 1026 K 21) for CuNi with 41.62 at.% Ni [18]. Above 1208C, the resistance decreases more steeply. A sharp change to a resistance increase was observed in Ar at 3808C. This is above the temperature of the steep stress change. We attribute it to oxidation in the stage II (see Section 3.4). For comparison, the DR/R0-vs.-T curve measured in a N2/H2 (5 vol.%) gas mixture is also given in Fig. 4. In such an atmosphere, the steep resistance increase between 300 and 3808C does not occur. This characterizes the described feature in the curve under Ar atmosphere as oxidation-induced. Below 3808C, the relative
resistance change in N2/H2 atmosphere is slightly smaller than that in Ar atmosphere, showing slight oxidation at this temperatures under Ar atmosphere. As already previously observed [4], only a thin Mnoxide-rich surface layer is forming during a temperature ramp to 5508C in N2/H2 atmosphere. This fact enables us to consider the measured curve in Fig. 4 for N2/H2 atmosphere as practically without CuNi(Mn) oxidation at least to 5008C. 3.4. Model for the oxidation-induced stress change The oxides NiO and CuO are immiscible [19]. Therefore,
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area of the entire ®lm. The thickness of the metallic CuNi(Mn) shrinks, but no stress is developed in the ®lm. Now the stress produced in stage I shall be evaluated. It is assumed that a nickel layer of width wd diffuses out of the perpendicular grain boundaries to the surface. With the outdiffusion of Ni, the lateral shrinkage (1lat) in both the x and y direction of the ®lm plane is
1lat wd =L
5
where L is the lateral grain dimension of assumed square columns. This shrinkage results in a biaxial tensile-stress contribution Ds 1
Fig. 4. Relative change in electrical resistance in CuNi(Mn) ®lms during temperature cycles (4 K/min) in Ar atmosphere and, for comparison, in an N2/H2(5 vol.%) gas mixture. Up to 3808C under Ar atmosphere, only a small resistance contribution occurs due to oxidation, above 3808C, the resistance increases distinctly by oxidation.
only NiO or a double layer CuO/NiO is observed. The formation enthalpy for NiO (240 kJ/mol) is higher than that for CuO (165 kJ/mol) [20]. This is the energetic reason, that NiO forms preferentially up to 4508C. If during oxidation the metal atoms penetrate preferentially the oxide layer, then, in principle, the oxidation continues on the top layer without stress generation. If, vice versa, oxygen atoms penetrate preferentially the oxide layer, then the oxidation occurs on the oxide/metal interface and a compressive stress is produced. The preferential penetration of metal atoms through the growing oxide is typical for metals [21]. The behavior is often more complicated by grain-boundary diffusion and grain-boundary oxidation, connected with microcrack formation and other effects [22,23]. In the present paper, the following model is considered for the stress evolution during oxidation. In a ®rst stage (300±4008C, stage I), nickel diffuses to the surface. This mass transport occurs by grain-boundary diffusion. Using the melting temperature T m 1550 K, the temperature range of 300±4008C corresponds to homologous temperatures of T=T m 0:37±0:43, i.e. the typical range for processes associated with grain-boundary diffusion. As the grain boundaries are perpendicular to the ®lm plane in the columnar grain morphology, the mass transport out of grain boundaries to the surface results in an elastic strain in the grains and, thus, in a biaxial tensile-stress contribution. The situation is, therefore, similar to that of grain-boundary diffusional creep (Coble creep) or hillocking in metallization ®lms [24,25]. When at higher temperatures (.4008C) volume diffusion starts becoming effective (stage II), the transport of metal atoms from the ®lm interior to the surface occurs over the
lat E= 1
2 n
6
Since the material of the out-diffused volume is oxidized and the shrinkage in both the x and y direction is 1lat, a portion of 21lat related to the total number of metallic atoms is oxidized. The grown oxide thickness (tox) is given by tox 21lat bt
7
Concerning the change in resistance due to oxidation, the NiO can be considered as insulator. Then, the fractional resistance change is equal to the fractional thickness change of the metal, i.e. DRox 21lat R
8
Now, the lateral shrinkage 1lat, the width wd of the outdiffusing layer of the grain boundaries, the thickness tox of the forming NiO layer, and the resulting resistance change DRox/R shall be evaluated for the case that the developing stress is the observed stress change of 500 MPa between 300 and 3808C. The following material data are used: E= 1 2 n 214 GPa [10], b 1:6 for NiO [14], and L 20 nm. Using Eqs. (5)±(8), one obtains 1lat 0:0023, wd 0:047 nm, tox 3:2 nm, and DRox =R 0:47%, respectively. The value of wd 0:047 nm means that about 20% of an atomic layer (atomic diameter is 0.25 nm) is diffused out of the grain boundaries of the columnar grains. In the evaluation, (i) the stress contribution of the growing NiO layer which is small for thin oxide layers according to Eq. (3) and (ii) the depth inhomogeneity of the ®lm stress due to different diffusion distances from different depths and the attachment of the ®lm onto the substrate, are not taken into account. However, the resistance changes evaluated from the stress changes correspond fairly well to the situation up to 3808C. Therefore, we believe, that our grainboundary-diffusion model is con®rmed for temperatures below 3808C by the given estimation. Above 3808C, the situation is more complicated because of various reasons. ² Lattice diffusion starts being more and more effective. This is con®rmed by the observed changes of the nickel concentration in the CuNi(Mn) sublayer (see Section 3.2)
W. BruÈckner, S. Baunack / Thin Solid Films 355±356 (1999) 316±321
as well as by recently reported lattice-parameter changes correlated to the composition shift [2,5]. Therefore, the oxide layer can grow by diffusion of the metal atoms from the uppermost CuNi(Mn) layers into the oxide without distinct stress generation. ² Stress relaxation by plastic deformation occurs. Such stress relaxation (diffusional ¯ow, mechanical twinning) was observed in CuNi ®lms above 3008C [7,13]. ² Grain coarsening and oxidation lead to microstructural and topographical changes which affect the stress evolution. Although we could not observe grain coarsening in TEM studies on in-plane sections of the cycle-stopprepared sample with T max 4508C, recrystallization may already start in this sample at the ®lm/substrate interface as found during thermal cycling of CuNi(Mn) ®lms in vacuum [7]. ² Stress contributions by (also inhomogeneous) growing oxide layer(s) in¯uence the averaged ®lm stress (force per unit length) according to Eq. (3), but are complicated. Because of these numerous in¯uences to the stress development, the stress-vs.-temperature curves above 3808C can not be analyzed in detail. 4. Conclusions During heat treatment under argon, oxidation of CuNi(Mn) resistive ®lms starts already at 3008C by reaction with the remaining oxygen under formation of a thin NiO layer. Above 4008C, the growth of the NiO layer increases. Above 5008C, a double layer of copper oxide on nickel oxide forms. Between 300 and 4008C, the oxidation in the initial stage gives rise to tensile stress of about 500 MPa. This is the striking feature in the stress-temperature curve associated with oxidation. The tensile-stress contribution is explained semiquantitatively by a model in which nickel diffuses in the columnar structure to the surface by grain-boundary diffusion and forms NiO. The out-diffusion of Ni leads to ®lm contraction and, because of the attachment of the ®lm onto the substrate, to the measured tensile-stress component. One can learn, from the investigations, that the formation of an only a few nanometers thick oxide layer on thin metallic ®lms may possibly lead to high stresses.
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Acknowledgements The authors are indebted to Th. Knuth at microtech Teltow for ®lm preparation, V. Michel for chemical analysis, J. Thomas for TEM studies, V. Weihnacht for stimulating discussions, and R. Vogel for technical assistance.
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