Influence of substrate temperature on atomic layer

Thin Solid Films 340 (1999) 110±116

In¯uence of substrate temperature on atomic layer growth and properties of HfO2 thin ®lms J. Aarik a,*, A. Aidla a, A.-A. Kiisler a, T. Uustare a, V. Sammelselg b È likooli Street, EE2400 Tartu, Estonia Institute of Materials Science, University of Tartu, 18 U b Institute of Physics, University of Tartu, 142 Riia Street, EE2400 Tartu, Estonia

a

Received 1 April 1998; accepted 21 August 1998

Abstract Atomic layer growth of hafnium dioxide from HfCl4 and H2O has been studied at substrate temperatures ranging from 180±6008C. A quartz crystal microbalance was used for the real-time investigation of deposition kinetics and processes affecting the growth rate. It was shown that the layer-by-layer growth was self-limited at temperatures above 1808C. The data of ex situ measurements revealed that the structure, density and optical properties of the ®lms depended on the growth temperature. The absorption coef®cient of amorphous ®lms grown at 2258C was below 40 mm 21 in the spectral range of 260±850 nm. The refractive index of the ®lms grown at 2258C was 2.2 and 2.0 at 260 and 580 nm, respectively. The polycrystalline ®lms with monoclinic structure grown at 5008C had about 5% higher refractive index but more than an order of magnitude higher optical losses caused by light absorption and/or scattering. q 1999 Elsevier Science S.A. All rights reserved. Keywords: Atomic layer growth; Hafnium; Oxides

1. Introduction Hafnium dioxide (HfO2) is a material which has been extensively studied because of its hardness, high chemical stability and excellent dielectric properties. Due to its hardness and good transparency in the visible and ultraviolet ranges of the spectrum, HfO2 has been often used in optical coatings (e.g. [1,2]). However, it can be also applied as an insulator in electronic and optoelectronic devices. Growing the thin ®lms for these applications must be usually carried out with accurate thickness control while the layers themselves have to be very thin. Atomic layer deposition (ALD) is one of the techniques used for growing HfO2 thin ®lms [3,4]. The main advantage of ALD is a high-precision thickness control and uniform thickness of grown ®lms. In addition, it has been demonstrated that ALD enables one to deposit the high-density tetragonal polymorph of HfO2 [4] and the orthorhombic TiO2-II polymorph of titania [5] which have never been obtained by other thin ®lm deposition methods. HfO2 thin ®lms have been grown by ALD at substrate temperatures of 5008C [3] and 3258C [4] while HfCl4 and H2O have been used as the precursors. The ®lms obtained have been crystalline with monoclinic [3] and/or tetragonal [4] structure. * Corresponding author. Tel.: 1 372 7 375877; fax: 1 372 7 375540.

Unfortunately, as revealed by atomic force microscopy studies, the surfaces of the ®lms are rather rough while the surface roughness increases together with the ®lm thickness [3]. Because of light scattering on the surface roughness, the application of this kind of ®lm in optical coatings is limited. Therefore it is of interest to ®nd some ways to modify the deposition conditions so that ®lms with higher surface ¯atness can be grown. It has been shown that the surface morphology of ALDgrown oxide ®lms is sensitive to the substrate temperature during the ®lm growth. For instance, growth temperatures of about 1008C have produced amorphous titania ®lms with absorption and/or scattering losses below 10 mm 21 whereas absorption exceeding the latter value by 1.5±3 orders of magnitude has been measured for polycrystalline ®lms grown at 165±5008C [6]. Rather similar dependence was observed in the case of tantalum oxide ®lms grown by ALD at temperatures of 80±5008C [7]. However, no data about such a comparison have been published for HfO2. To our knowledge the dependence of growth rate, crystallinity and optical properties of the ®lms on deposition temperature has not been studied either. For this reason we performed experiments described below. The aim of the work was to investigate the atomic-layer deposition kinetics of HfO2 at different temperatures and study the structure and optical parameters of the ®lms obtained.

0040-6090/99/$ - see front matter q 1999 Elsevier Science S.A. All rights reserved. PII: S00 40-6090(98)0135 6-X

J. Aarik et al. / Thin Solid Films 340 (1999) 110±116

Fig. 1. Schematic of the ALD reactor.

2. Experimental HfO2 ®lms were grown in a ¯ow-type hot-wall ALD reactor (Fig. 1) using HfCl4 and H2O as the precursors. Solid HfCl4 was volatilized in an effusion cell and carried into the reaction zone with pure N2 gas. The pressure of HfCl4 was controlled by the effusion cell temperature (TC) which was varied from 130 to 1508C. The deviation of TC from any set-point value did not exceed 18C. The injection of HfCl4 into the reactor was switched on and off by changing the direction of the carrier gas ¯ow between the effusion cell and reaction zone. H2O was volatilized in a container kept at 208C. The ¯ow of H2O vapour into the

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reactor was controlled by a calibrated needle valve and solenoid valve. The precursors were led into the reaction zone alternately and the solid ®lm was formed as a result of subsequent surface reactions. In order to avoid gas phase reactions caused by intermixing the precursors, the reactor was purged with pure carrier gas after each precursor pulse. The deposition kinetics was studied using real-time quartz crystal microbalance (QCM) measurements. In these experiments, the thin ®lm was deposited onto the mass sensor and the mass change as a function of time was recorded. The time resolution of the measurements was 0.25 s. As the pulse and purge times were usually not shorter than 2 s in our experiments, the measurements enabled us to study the behaviour of the mass during a single ALD cycle which consisted of a HfCl4 pulse, purge time, H2O pulse and another purge time. To reduce the effect of mass sensor surface on the deposition kinetics, a HfO2 buffer layer with a thickness of 5±10 nm was grown on the sensor at 3008C before starting the real measurements. As the mass sensor was sensitive to the reactor temperature (TR) the latter was stabilized with the accuracy of 0.18C. The maximum temperature at which we could perform the QCM measurements was 4008C. The ®lms for ex situ measurements were grown on fused silica and single-crystal silicon substrates. Re¯ection high energy electron diffraction (RHEED) and X-ray diffraction (XRD) were used to determine the structure of the ®lms. The surface morphology was studied by scanning electron microscopy (SEM). Auger electron spectrometry (AES) and electron-probe microanalysis (EPMA) were applied to study the composition. The refractive indices and thicknesses of the ®lms grown on silicon substrates were obtained from ellipsometry measurements. The latter were carried out at the wavelength of 632.8 nm. The transmission spectra of the ®lms gown on silica substrates were measured by a spectrophotometer in the wavelength range 190±850 nm. The refractive index and ®lm thickness were calculated from the spectral oscillations of transmission curves using the method proposed by Swanepoel [8].

3. Results 3.1. Deposition kinetics

Fig. 2. Mass sensor signal as a function of deposition time in conditions when the mass sensor is exposed to alternating HfCl4 and H2O pulses (lower curve) and to subsequent HfCl4 pulses (upper curve).

Fig. 2 demonstrates the time dependence of the mass sensor signal recorded during ALD growth. In the conventional ALD mode (lower curve in Fig. 2) the HfCl4 pulse results in a mass increase. No signi®cant mass change can be observed during the purge time. The H2O pulse, in contrast, causes a mass decrease and relaxation of the signal at a new level whereas no changes occur after switching off the H2O pulse. This behaviour is typical for the chloride± water ALD process [9]. It shows that the surface reactions are completed during precursor pulses and no desorption of adsorbed species takes place during the purge times. Repla-

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Fig. 3. Dependence of mass increase per cycle on dose H2O (a) and on pressure of HfCl4 in the effusion cell (b).

cement of chlorine adsorbed during the HfCl4 pulse with signi®cantly lighter hydroxyl groups and/or oxygen is an explanation for the decrease of ®lm mass during the H2O pulse. The upper trace in Fig. 2 shows the behaviour of ®lm mass when the H2O pulses are not supplied. As one can see, the mass sensor signal signi®cantly changes during the ®rst HfCl4 pulse only. Consequently a saturated adsorbate layer is obtained during a single chloride pulse. As there is no noticeable desorption between subsequent pulses, the following exposure of the sensor to HfCl4 does not cause any more mass changes. Both the saturation of adsorption during a precursor pulse and the absence of desorption during purge times are main requirements for ALD growth. Evidence of perfect ALD-type growth was also the fact that the growth rate determined as an increase of the mass sensor signal per ALD cycle was independent of purge times when the latter were varied from 2 to 10 s at HfCl4 and H2O pulse times of 2 s. At purge times below 2 s, however, the growth rate increased with decreasing purge times. This indicates that, in the latter conditions, the gas phase reactions caused by overlapping of precursor pulses most probably also contribute to the ®lm growth. Thus after switching

off the ¯uxes the precursors can still reside in the reactor up to 2 s. The dependences of the growth rate on the precursor doses are shown in Fig. 3a,b. As one can see, the growth rate saturates with increasing precursor dose. However, the saturation behaviour signi®cantly depends on the reactor temperature. At 1808C the saturation is not complete and occurs at signi®cantly higher growth rates and precursor doses than at 3008C. This demonstrates that the surface concentration of adsorption sites for precursor molecules depends on the temperature. A possible reason for such a dependence is the in¯uence of temperature on the abundance of surface hydroxyl groups. On the one hand, it is known that metal chlorides are generally more reactive towards surface OH-groups than towards oxygen bridges (see e.g. [9]). On the other hand, chlorine can be released from adsorbing HfCl4 molecules when exchange reactions with surface hydroxyl groups take place. As a result of exchange reactions the dimensions of an adsorbed molecule may be smaller than those of the impinging molecule. The higher is the number of hydroxyl groups reacting with an impinging chloride molecule, the smaller is the average size of the adsorbed molecule and the higher is the maximum possible surface concentration of adsorbed molecules in a surface monolayer. As shown earlier [10], the mass changes corresponding to chloride and H2O pulses enable one to obtain additional information about the most probable mechanism of surface reactions. Provided that after each H2O pulse the surface is terminated by OH-bonds and no desorption of surface intermediate species takes place, one can describe the surface reactions as ÿ ÿ n±OHs 1 HfCl4 g ! ±O±n HfCl42n s 1 nHCl g 1 and

ÿ ±O±n HfCl42n s 1 4 2 nH2 O g ÿ ! ±O±n Hf OH42n s 1 4 2 nHCl g

2

during HfCl4 and H2O pulses, respectively. Formally n, which shows the average number of OH-groups reacting with an adsorbing chloride molecule, can range from 0 to 4. However, in the conditions when stable ALD growth is achieved (i.e. the mass increase per completed ALD cycle does not depend on the number of cycles applied) and no OH-groups stay in the ®lms, the number of OH-groups consumed in the ®rst reaction (Eq. (1)) must be equal to the number of those formed in the second reaction stage (Eq. (2)). This requirement is ful®lled when n 2. Assuming that n 2 and using the molar masses of substances, one can ®nd from Eqs. (1) and (2) that Dm0 =Dm1 0:85, where Dm0 is the mass change during a completed ALD cycle and Dm1 is the mass change during a HfCl4 pulse. The values of Dm0/Dm1 measured by QCM for reactor temperatures 180, 225, 300 and 4008C were equal to 0.86, 0.80, 0.78 and 0.74, respectively. As one can see, the experi-


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of a surface layer which is partially terminated by OHgroups after the H2O pulse could explain the experimental Dm0/Dm1 ratios observed at 225±4008C. One can consider that both oxygen bridging between adjacent Hf atoms and OH-groups may appear on the surface during H2O treatment. Assuming that the adsorption of HfCl4 is described by Eq. (1), the reaction between the chlorinated surface and H2O can be expressed as ÿ ±O±n HfCl42n s 1 2H2 O g ! ±O±n HfO22n OHn s ÿ 1 4 2 nHCl g 5

Fig. 4. Mass increase per cycle and thickness increase per cycle as functions of reactor temperature. HfCl4 pressure is 50 Pa in the effusion cell and H2O dose is 2.1 mPa m 3.

mental value for 1808C is close to that found from Eqs. (1) and (2) for n 2. The data measured for higher reactor temperatures, however, are signi®cantly lower than this value. A possible reason for reduced experimental values of Dm0/Dm1 is that the number of OH-groups reacting with an adsorbing HfCl4 molecule is lower than 2. At n , 2, however, Eqs. (1) and (2) predict an increase of the surface concentration of OH-groups during each ALD cycle. Thus the growth rate should also increase with the number of cycles applied. As one can see in Fig. 2, this kind of dependence has not been recorded. Thus another mechanism causing changes in Dm0/Dm1 should be considered. One can assume that dehydroxylation of the HfO2 surface, similar to that of SiO2 [9] and TiO2 [11], occurs at elevated temperatures. In this process adjacent OHgroups react with each other forming oxygen bridges or unsaturated bonds on the surface (see e.g. [12]). If the reaction between HfCl4 and oxygen bridges is similar to that proposed by Ritala et al. [12] for the TiCl4/H2O ALD process, it could be written as ÿ ÿ Hf ±O±2 Hf s 1 HfCl4 g ! Hf Cl±O±Hf OHfCl3 s 3 Proceeding from the same analogy [12,13], we can express the overall surface reaction which occurs during H2O pulse and following purge time: ÿ ÿ Hf Cl±O±Hf OHfCl3 s 1 2H2 O g ÿ 4 ! Hf ±O±2 Hf ±O±2 Hf s 1 4HCl g The Dm0/Dm1 ratio calculated from Eqs. (3) and (4) equals 0.66 and is signi®cantly lower than the experimental values presented above. Consequently the reaction mechanism described by Eqs. (3) and (4) is unlikely to be dominant at 180±4008C. Comparison of calculated Dm0/Dm1 values with experimental data indicates, however, that formation

where 0 # n # 2 and, for the reasons discussed above, the number of OH-groups recovered during the H2O pulse (Eq. (5)) is equal to those consumed during HfCl4 adsorption (Eq. (1)). One can see that at n 2 Eq. (5) is identical with Eq. (2) while at n 0 it gives the same mass change per adsorbed HfCl4 molecule as Eq. (4). In a general case the Dm0/Dm1 ratio calculated from Eqs. (1) and (5) depends on n and can be expressed as Dm0 =Dm1 210:6=320:6 2 36:5n

6

An agreement with experimental Dm0/Dm1 values measured at 180, 225, 300 and 4008C can be obtained at n values of 2.0, 1.6, 1.4 and 1.0, respectively. Correspondingly, the chlorine to hafnium ratio in the adsorbate layer formed during HfCl4 pulse (see Eq. (1)), should be equal to 2.2, 2.4, 2.5 and 3.0 at 180, 225, 300 and 4008C, respectively. Both the decrease of the number of OH-groups participating in surface reactions and the increase of Cl/Hf ratio with the increasing deposition temperature are well consistent with the observed decrease of deposition rate. Indeed, the deposition temperature increase from 180 to 4008C causes a reduction of growth rate by a factor of 1.8 (Fig. 4). At the same time the model based on Eqs. (1) and (5) and on experimental values predicts that the number of OHgroups participating in the surface reactions decreases by a factor of 2 and the Cl/Hf ratio increases by a factor of 1.5. Thus the reaction mechanism proposed rather well explains the relationship between growth rate and Dm0/Dm1 ratio. In this way it supports the idea that changing roles of surface OH-groups and/or steric effects are responsible for the dependence of deposition rate on the substrate temperature. Comparing the growth rates measured by QCM and those obtained from optical thickness measurements (Fig. 4) one can see that the thickness change per cycle diminishes faster with increasing growth temperature than the mass increase per cycle. This is evidence that the atomic density of ®lms increases with growth temperature. Rough estimation made on the basis of data presented in Fig. 4 shows that the ®lm density increases by a factor of 1.2 with the increase of deposition temperature from 180 to 4008C. Such behaviour is probably due to the changes in the ®lm structure.

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temperature was diminished to 2258C. Also chlorine contamination of the ®lms deposited at 2258C was observed by AES and EPMA. The chlorine concentration reached 4.3 at.% in these ®lms. This result indicates that adsorbed chlorine has not been completely exchanged at 2258C although saturation of adsorption has been achieved during the H2O pulse. Polycrystalline ®lms deposited at 3008C contained about 1 at.% chlorine while no chlorine residues could be recorded in the ®lms grown at 400±6008C. Chlorine contamination is probably one reason for formation of amorphous ®lms at low substrate temperatures. 3.3. Optical properties Fig. 5. X-ray diffraction pattern of a HfO2 ®lm deposited on silicon substrate at 5008C. The ®lm thickness is 200 nm. Arrows indicate the positions of peaks which can be attributed to the orthorhombic phase [14,15]. All the other peaks belong to monoclinic HfO2.

3.2. Structure and morphology The ®lms grown at substrate temperatures below 3008C were amorphous. At 3008C amorphous ®lms grew in the initial stage of deposition whereas the ®lms thicker than 25±30 nm contained crystallites with monoclinic structure. The ®lms grown at 4008C and higher substrate temperatures were polycrystalline at the thicknesses of 8±10 nm already. According to RHEED data the crystallinity of ®lms improved with increasing the growth temperature up to 5008C. Further raising the growth temperature to 6008C did not yield signi®cant changes in RHEED patterns. All the crystalline ®lms were mainly of monoclinic structure with some preferred orientation of crystallites in the ®lms grown at 500±6008C. However, three re¯ection peaks not belonging to the monoclinic phase, were observed in the XRD rocking curves of the ®lms grown at 5008C (Fig. 5). These re¯ections corresponding to the 2u values of 30.28, 59.06 and 75.888 and interplanar distances of 0.2949, 0.1563 and 0.1252 nm, respectively, can be attributed to a highpressure orthorhombic polymorph of HfO2 [14,15]. It should be noted that the ®rst re¯ection mentioned above is very close to the one earlier observed by Ritala et al. [3] and Kukli et al. [4] at 2u 30:38 and ascribed to the tetragonal phase of HfO2 [4]. SEM micrographs (Fig. 6) show that the surface morphology of the ®lms depends on the deposition temperature. No roughness can be seen on the surface of an amorphous ®lm grown at 2258C (Fig. 6a). In contrast, crystallites with dimensions close to 100 nm can be observed on the surface of polycrystalline ®lms deposited at 5008C (Fig. 6c). This kind of behaviour is in agreement with some earlier results which show a very similar effect in the case of titanium oxide [6] and tantalum oxide [7] ®lms. According to AES and EPMA data the composition of the ®lms was independent of deposition temperature when the latter was varied from 400 to 6008C. However, the O/Hf ratio decreased by the factor of 1.06 when the growth

Fig. 7 shows the transmission spectra of ®lms grown at 225 and 5008C. The thicknesses of the ®lms are 300 and 370 nm, respectively. It can be seen that the lower growth temperature yields ®lm of signi®cantly higher transparency than the higher temperature at which a polycrystalline structure is formed. Moreover, the absorption edge of the polycrystalline ®lm grown at 5008C is at a higher wavelength than that of the amorphous ®lm grown at 2258C. Thus an optical band gap shrinkage, similar to that earlier observed for tantalum oxide [7] and titanium oxide [6,16] thin ®lms, takes place with transition from amorphous phase to the crystalline one. It can be seen in Fig. 7 that there is a speci®c feature at the wavelengths of 190±200 nm in the transmission spectrum of the ®lm grown at 5008C. The tail which appears in the range of strong absorption cannot be explained by interference [8]. Thus a modi®cation of the absorption edge shape takes place with changing deposition temperature. It should be noted that the tail appears at a growth temperature of 3008C, already, although at the latter temperature the effect is not so strong as that shown in Fig. 7. Probably the changes in the shape of absorption edge are related to the structure of ®lms. In order to understand the physical reasons for appearance of the absorption tail, however, additional studies are needed because a number of effects (co-existing phases of different structure and/or stoichiometry, size effects, in¯uence of grain surfaces and interfaces, intrinsic stresses etc.) can cause this phenomenon. The refractive index values calculated from the transmission spectrum of the ®lm grown at 2258C (Fig. 7), are equal to 2.2 and 2.0 at the wavelengths 260 and 580 nm, respectively, while the absorption coef®cient is below 40 mm 21 in the spectral range of 260±850 nm. In contrast, the absorption coef®cient calculated for the ®lm grown at 5008C is 300 mm 21 at 580 nm and as high as 800 mm 21 at 260 nm. Unfortunately, the transmission spectrum of the ®lm grown at 5008C (Fig. 7) is obviously affected by light scattering. For this reason the absorption coef®cient and refractive index values obtained from this curve are not reliable. In order to reduce the effect of light scattering the refractive index was measured by the ellipsometer for very thin (60±80 nm) ®lms grown on silicon substrates. The values


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Fig. 7. Transmission spectra of HfO2 ®lms grown at 225 and 5008C using 1400 and 3000 ALD cycles, respectively. The thicknesses of the ®lms are 300 and 370 nm.

the ®lms increases with growth temperature. This is not surprising and is most probably caused by the increase of atomic density discussed in Section 3.1. The refractive index value measured in this work for the ®lms grown at 5008C is close to that published for HfO2 thin ®lms deposited in the reactive ion-plating process [2]. 4. Conclusions

Fig. 6. SEM images of ®lms grown on silicon substrates at 225 (a), 300 (b) and 5008C (c). The thicknesses of the ®lms are 300, 370 and 370 nm, respectively.

obtained were 1.95, 2.00 and 2.10 for the ®lms grown at 225, 300 and 5008C. It should be noted, however, that the structure of the ®lms grown at 3008C was obviously not uniform in the growth direction. Indeed, as mentioned above, the structure of these ®lms depended on the thickness whereas the transition from the formation of amorphous phase to the growth of polycrystalline structure appeared at thicknesses of 25±30 nm. In the case of ®lms deposited at 225 and 5008C no dependence of the ®lm structure on the thickness was revealed. Therefore the refractive index values determined for these ®lms characterize amorphous and polycrystalline HfO2, respectively. The refractive index values presented above indicate that the optical density of

The results of our experiments show that HfO2 thin ®lms of high optical quality can be obtained at substrate temperatures below 3008C. The ®lms grown at these temperatures are amorphous and have rather ¯at surfaces. In addition, the optical band gap of amorphous ®lms is wider than that of polycrystalline ®lms grown at higher temperatures. However, deviations from the self-limited growth mode appear at deposition temperatures of 1808C and lower. Moreover, the ®lm density and refractive index decrease with decreasing growth temperature. Thus the appropriate choice of substrate temperature in ALD-processing of HfO2 thin ®lms is of importance. Also, the temperature deviations along the substrate should be minimized because the growth rate depends on the substrate temperature. Acknowledgements The authors are thankful to Professor Lauri NiinistoÈ for the access to XRD facilities at Helsinki University of Technology. We also wish to thank Kaupo Kukli and Hugo MaÈndar for performing XRD measurements. The work was supported by Estonian Science Foundation.

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