Journal of the Korean Physical Society, Vol. 52, No. 4, April 2008, pp. 11031108
Atomic Layer Deposition of HfO2 Thin Films on Ultrathin SiO2 Formed by Remote Plasma Oxidation Seokhoon Kim, Sanghyun Woo, Hyungchul Kim, Inhoe Kim, Keunwoo Lee, Wooho Jeong, Taeyong Park and Hyeongtag Jeon Division of Materials Science and Engineering, Hanyang University, Seoul 133-791
(Received 17 August 2007) An ultrathin SiO2 layer was grown on a H-terminated Si substrate by using remote plasma oxidation. The subsequent HfO2 deposition on the ultrathin SiO2 was achieved by a remote plasma atomic layer deposition (RPALD). During the HfO2 lm deposition and rapid thermal annealing (RTA), the ultrathin SiO2 bu er layer e ectively suppressed the formation of Hf silicate layers in the interfacial region. The Hf silicate layer in the interfacial region grew with increasing RTA temperature. The positive xed oxide charges in the HfO2 lm were reduced with increasing RTA temperature. The thickness of the HfO2 lms with an ultrathin SiO2 bu er layer showed a lower e ective xed oxide charge density (Qf;ef f ) due to the shift of the at band voltage (VF B ) toward the positive direction, compared to those with H-terminated Si. The thin interfacial layer of the HfO2 lms with a thin SiO2 bu er layer resulted in a low equivalent oxide thickness (EOT) value. The leakage current densities of the HfO2 lms increased because of the crystallization of the HfO2 lm after RTA. PACS numbers: 77.55.+f, 78.40.Fy Keywords: ALD, High-k dielectrics, HfO2 , Interfacial layer, Bu er layer I. INTRODUCTION
The scaling down of conventional SiO2 lms in sub0.1-m metal oxide semiconductor eld e ect transistors (MOSFETs) is reaching its physical and electrical limits from the viewpoint of gate leakage current [1]. High dielectric constant (high-k) dielectrics, such as HfO2 , ZrO2 and Al2 O3 , have been investigated as possible replacements for SiO2 to suppress the leakage current. Particularly, HfO2 thin lms are currently of great interest due to a potential application as high-k materials in future CMOS devices. HfO2 exhibits a high dielectric constant (k 25), a large energy band gap (5.68 eV) and thermodynamic stability in contact with Si [2{4]. The atomic layer deposition (ALD) method has been studied for the deposition of high-k materials and has many practical advantages, such as nano-scale thickness control and the ability to grow uniform lms over a large substrate area [5, 6]. The ALD process, employing halogen precursors, provides good step coverage and a low impurity concentration in thin lms [7, 8]. However, it has several problems, such as the presence of halogen atom residues in the lms, the corrosion of gas delivery lines and the generation of particles. It has been suggested that metal-organic atomic layer deposi E-mail:
[email protected]; Tel: +82-2-2220-0387; Fax: +82-2-2292-3523
tion (MOALD) can overcome these problems. The use of metal-organic precursors has many advantages in depositing lms, but it also has some problems, such as low lm densities and high carbon contamination. To resolve these problems, the plasma-enhanced atomic layer deposition (PEALD) method has been used to deposit high-k thin lms. The PEALD method has been found to increase the reactivity of the precursors, minimize carbon contamination, widen the process window and densify the lms [9, 10]. However, high-energy species in plasma can induce damage, which degrades the interface quality and induces a high interface state density and a high xed charge [11]. A degraded interface reduces carrier mobility and lowers the operating speed of a device. To prevent the degradation of a lm's characteristics by plasma e ects, the remote plasma atomic layer deposition (RPALD) method has been studied to deposit high-k oxide and di usion barrier materials [12]. The RPALD method can minimize the damage caused by plasma generation. However, during RPALD or postdeposition annealing, an interfacial layer, composed of compound, such as SiOx and Hf silicate, is grown at the HfO2 /Si interface and this limits the decrease in the equivalent oxide thickness (EOT) [13]. Therefore, prior to the RPALD of HfO2 , the formation of a bu er layer on the Si substrate is required to prevent a reaction at the HfO2 /Si interface and to improve the quality of the interface [14,15]. In this study, a thin SiO2 layer was grown
-1103-
-1104-
Journal of the Korean Physical Society, Vol. 52, No. 4, April 2008
on a Si substrate by using in-situ remote plasma oxidation. This SiO2 layer is expected to act as a reaction barrier at the HfO2 /Si interface during HfO2 deposition. The physical and the electrical properties of a HfO2 lm deposited on a ultrathin SiO2 formed by remote plasma oxidation were investigated.
II. EXPERIMENTS
HfO2 lms were grown on Si (100) p-type substrates. The Si (100) substrates were cleaned with a piranha solution (H2 O2 : H2 SO4 = 1 : 4) for 10 min, washed in deionized (DI) water and dipped in a dilute HF solution for 2 min (HF : H2 O = 1 : 100). An approximately 0.5-nm-thick SiO2 layer was grown on the Si substrate by using remote plasma oxidation for 10 s at 0.1 Torr. The process temperature and the plasma power were 250 C and 100 W, respectively. The HfO2 thin lms were deposited at 250 C on the 0.5-nm-thick SiO2 layer using a tetrakis(diethylamino)hafnium (TDEAH) precursor with the RPALD process. A TDEAH precursor was introduced into the reactant chamber by using a bubbler with Ar carrier gas. The basic cycle of the HfO2 deposition consisted of supplying the Hf precursor and then exposing a remote oxygen plasma. An Ar purging process was performed for the complete separation of the precursor injection and the oxygen plasma process. Rapid thermal annealing (RTA) after HfO2 deposition was performed at 600 and 800 C for 30 s in a N2 ambient. The chemical state and bonds of the deposited HfO2 thin lm were investigated by using X-ray photoelectron spectroscopy (XPS). High-resolution transmission electron microscopy (HRTEM) was used to determine the thickness of the HfO2 lm and interfacial layer. Metal-oxide-semiconductor (MOS) capacitors with 100 nm-thick Pt electrodes deposited using an e-beam evaporator were fabricated to investigate the electrical properties of the HfO2 lms. Then, post-metallization annealing (PMA) was carried out in a 97 % N2 /3 % H2 mixed atmosphere at 450 C for 30 min. The electrical properties such as the equivalent oxide thickness (EOT), the at band voltage (VF B ) and the e ective xed oxide charge density (Qf;ef f ), were investigated by using capacitance-voltage (C -V ) measurements with a Keithley 590 C -V analyzer. The leakage current density was measured using an HP4155A semiconductor parameter analyzer.
III. RESULTS AND DISCUSSION
Figure 1 shows the Si 2p XPS spectra of an approximately 0.5-nm-thick SiO2 lm formed on a Si substrate
Fig. 1. Si 2p XPS spectrum of (a) 0.5-nm-thick SiO2 lm formed on a Si substrate by using the remote plasma oxidation and Si 2p XPS spectra of (b) - (d) 5 nm-thick HfO2 lms grown on 0.5 nm-thick SiO2 for various RTA temperatures. The RTA was performed at 600 and 800 C for 30 s in a N2 ambient.
by using remote plasma oxidation and those of 5-nmthick HfO2 lms grown on 0.5-nm-thick SiO2 for various RTA temperatures. RTA was performed at 600 and 800 C for 30 s in a N2 ambient after HfO2 lm deposition. The XPS was measured at a take-o angle of 90 . All peaks were calibrated with the C 1s peak position of 285 eV. The Si 2p spectra in Figure 1 consist of two peaks, a peak at 99 eV, corresponding to the Si substrate and a peak at a higher binding energy of 102 to 103.3 eV, related to the formation of the interfacial layer. The binding energies between 102 and 103 eV correspond to the formation of silicate [14]. SiO2 was shown at a binding energy of 103.3 eV [14]. Figure 1(a) shows the formation of 0.5-nm-thick SiO2 layer on the Si substrate after remote plasma oxidation. After HfO2 deposition, intensity of the bulk Si peak at 99 eV in Figure 1(a) is sharply decreased compared to that of the bulk Si peak in Figure 1(b). This was due to the increased lm thickness after HfO2 growth. The Si 2p peak, at a binding energy of 103.3 eV in Figure 1(a), shifted to a lower binding energy after HfO2 deposition, as shown in Figure 1(b). This showed that the silicate layer was formed in the interfacial region during HfO2 deposition by using RPALD. The intensity of the Hf silicate peak increased while that of the bulk Si peak at 99 eV in Figures 1(b)-(d) slightly decreased as the RTA temperature was increased. This revealed that the thickness of the Hf silicate layers increased with increasing annealing temperature. As the RTA temperature was increased from 600 to 800 C, the binding energy related to the formation of silicate shifted higher from 102.4 to 102.6 eV. The residual oxygen gas di used more rapidly into the interfacial layer of the HfO2 lm in the RTA process at higher temperature. This results in the growth of the Hf silicate layer and additional oxidation in the
Atomic Layer Deposition of HfO2 Thin Films on { Seokhoon Kim et
Fig. 2. Hf 4f XPS spectra of HfO2 lms grown on 0.5 nm-thick SiO2 for various RTA temperatures. The RTA was performed at 600 and 800 C for 30 s in a N2 ambient.
Fig. 3. Cross-sectional TEM images of HfO2 lms grown on 0.5-nm-thick SiO2 : (a) as-grown lm and lms annealed at (b) 600 and (c) 800 C. RTA was performed for 30 s in a N2 ambient.
interfacial region. Figure 2 shows the Hf 4f XPS spectra of the HfO2 lms grown on 0.5-nm-thick SiO2 after RTA at 600 and 800 C for 30 s in a N2 ambient. As shown in Figure 2(a), the Hf 4f7=2 peak is located at a binding energy of 16.5 eV for the as-grown HfO2 lms, with a 0.5-nmthick SiO2 layer. Previously we reported that the Hf 4f7=2 peak position for an as-grown HfO2 lm on Hterminated Si by using RPALD was 16.6 eV [11]. The very thin SiO2 bu er layer may cause a shift of the Hf 4f7=2 peak to a binding energy lower than that for the HfO2 lm without a bu er layer because it retarded the growth of Hf silicate in the interfacial region during HfO2 deposition. The Hf 4f7=2 peak position for Hf silicate is 1 eV higher than it is for HfO2 , for which the peak was located at 16.5 to 17 eV [4]. As the RTA temperature was increased, the position of the Hf 4f7=2 peak shifted to a higher binding energy because of the growth of Hf silicate in the interfacial region. After annealing, the Hf 4f7=2 peaks of the HfO2 lms showed binding energies of 16.7 eV for 600 C and 17.4 eV for 800 C. Figure 3 shows cross-sectional TEM images of HfO2 lms grown on 0.5-nm-thick SiO2 before and after annealing. The RTA was performed for 30 s in a N2 ambi-
al.
-1105-
Fig. 4. High-frequency C -V curves of as-grown HfO2 lms for various numbers of RPALD process cycles. HfO2 lms were deposited by 10, 20, 30 and 40 process cycles. RTA was performed at 800 C for 30 s in a N2 ambient.
ent. As shown in Figures 3 (b) and (c), the HfO2 lms were crystallized from an amorphous structure to a polycrystalline structure after RTA at 600 and 800 C for 30 s in a N2 ambient because the crystallization temperature of HfO2 was less than 500 C. The thickness of the interfacial layer for HfO2 lms grown on 0.5-nm-thick SiO2 was 1.3 nm, as determined from the TEM images shown in Figure 3(a) and it was not increased even after RTA at 600 C, as shown in Figure 3(b). After RTA at 800 C, the HfO2 lm had a 2.0 nm-thick interfacial layer, as shown in Figure 3(c). This means that the di usion of oxygen into Si and Si into the HfO2 lm occurred during RTA at 800 C. The growth of the interfacial layer results in the decreased thickness of the HfO2 layer. The thinner HfO2 layer at 800 C may be less crystallized compared to that at 600 C because the degree of crystallinity of HfO2 lm has a tendency to be in inverse proportion to the lm thickness. Previously, we also observed that the thickness of the interfacial layer for the HfO2 lm grown on H-terminated Si by using RPALD increased from 2.0 to 2.2 nm after annealing at 800 C [16]. The thicknesses of the interfacial layer for as-grown and annealed HfO2 lms with an ultrathin SiO2 bu er layer were less than those of the lms grown on H-terminated Si. The ultrathin SiO2 layer formed on a Si substrate by using remote plasma oxidation prior to HfO2 deposition can e ectively retard the growth of the interfacial layer. The SiO2 bu er layer can act as a reaction barrier for the growth of the Hf silicate layer. Figure 4 shows the high-frequency C -V curves of the as-grown HfO2 lms as a function of the number of RPALD process cycles. HfO2 lms were grown by using 10, 20, 30 and 40 process cycles, respectively. The C -V curves were measured at a frequency of 100 kHz. From the high-frequency C -V data of Figure 4, the maximum capacitance value in the accumulation region decreased
-1106-
Journal of the Korean Physical Society, Vol. 52, No. 4, April 2008
Fig. 5. Interface trap level con guration in a MOS structure on a p-type substrate at VF B .
with increasing number of RPALD process cycles due to the increased lm thickness. For evaluating EOT values, a quantum-mechanical correction was applied. The EOT values of HfO2 lms at for 10, 20, 30 and 40 cycles were 0.78, 1.01, 1.26 and 1.50, respectively. The accumulated capacitance didn't vary in direct proportion to the increased lm thickness due to a variation in the dielectric constant of the lm and the formation of an interfacial layer. The dielectric constants of HfO2 lms for 10, 20, 30 and 40 cycles were 8.9, 9.7, 9.1 and 10.4, respectively. If the mobile charge and oxide trapped charge are assumed to be negligible, a compensation e ect between the interface state charge (Qit ) and the xed oxide charge (Qf ) is considered to determine the actual magnitude of the charge defect. Figure 5 shows the interface trap level con guration in a MOS structure on a p-type substrate at VF B . Most of the interface trap levels are empty. Assuming that the interface traps in the upper and the lower halves of the band gap (Eg ) are acceptorand donor-type traps, respectively, the acceptor-type interface traps above the Fermi energy (EF ) are neutral when empty and the donor-type interface traps below EF are positive when empty [17]. The donor-type interface traps above EF appear as a positive charge and the xed oxide charge of HfO2 lm is usually positive. There is no compensation e ect because both Qit and Qf are positive and cause a shift of VF B in the negative direction. Therefore, the e ective xed oxide charge density (Qf;ef f ) value is the sum of Qf and Qit and can be calculated from the VF B shift of the C -V curve by using Eqs. (1) and (2). Qf Qit Qf;ef f VF B = ms t t = ms t (1) "ox ox "ox ox "ox ox ( VF B ) ( VF B ) Qf;ef f = ms "ox = ms (2) tox Cmax ms is the di erence in work function between the Pt electrode and the Si substrate and Cmax is the maximum accumulation capacitance. To determine the value of ms in our MOS samples, we set the e ective work
Fig. 6. High-frequency C -V curves of 5-nm-thick HfO2 lms on 0.5-nm-thick SiO2 : as-grown and after RTA at 800 C for 30 s in a N ambient. 2
function of Pt to be 5.6 V and we set the doping concentration of the p-type Si substrate to be 5 1015 /cm2 . The ms was calculated to be 0.65 V. The value of VF B for the HfO2 lm were 0.64, 0.63, 0.62 and 0.60 V for the HfO2 lms with 10, 20, 30 and 40 process cycles, respectively. The calculated Qf;ef f of the HfO2 lms at 10, 20, 30 and 40 process cycles were 1.75 1011 , 3.2 1011 , 4.4 1011 and 7.0 1011 cm 2 , respectively. As the number of RPALD cycle increased, the at band voltage (VF B ) was shifted in the negative direction, which indicated that HfO2 had a positive xed oxide charge and the positive charges generated in the lms were increased by the growth of HfO2 layer. Figure 6 shows high-frequency C -V curves for approximately 5-nm-thick HfO2 lms grown on 0.5-nm-thick SiO2 before and after RTA at 800 C for 30 s in a N2 ambient. The hystereses of the HfO2 lms were both below 50 mV. After annealing at 800 C, the Cmax of the annealed lms was decreased because of the decreased series capacitance caused by the growth of the interfacial layer. This resulted in a decrease of the EOT value. The VF B for the annealed lms was shifted in the positive direction compared to that for the as-grown lms, which indicated that positive defect charges in the HfO2 lm were reduced by RTA. Figure 7 shows the VF B and the Qf;ef f values for the HfO2 lms grown on H-terminated Si and on 0.5-nmthick SiO2 , respectively. RTA was performed at 600 and 800 C for 30 s in a N2 ambient. The Qf;ef f was determined from the VF B shift with increasing annealing temperature. As the annealing temperature was increased, the VF B for all lms shifted in the positive direction. This meant that the positive defect charges were reduced with increasing RTA temperature. As the RTA temperature was increased up to 800 C, the VF B shifted from 0.51 to 0.55 for HfO2 lms grown on H-terminated Si and
Atomic Layer Deposition of HfO2 Thin Films on { Seokhoon Kim et
-1107-
al.
lms with a 0.5 nm-thick SiO2 bu er layer being thinner compared to the lms with H-terminated Si. The leakage current densities increased from 2.5 10 7 to 4.4 10 7 A/cm2 for H-terminated Si and from 3.3 10 7 to 4.2 10 7 A/cm2 for 0.5 nm-thick SiO2 as a result of the RTA process. The increased leakage current densities of the HfO2 lms with increasing RTA temperature are considered to be due to the crystallization of the amorphous HfO2 lm.
IV. CONCLUSION
Fig. 7. VF B and Qf;ef f values for the HfO2 lms grown on H-terminated Si and on 0.5 nm-thick SiO2 : as-grown and annealed at 600 and 800 C in a N2 ambient.
Fig. 8. EOT and J -V values for HfO2 lms grown on H-terminated Si and on 0.5-nm-thick SiO2 : as-grown and annealed at 600 and 800 C for 30 s in a N2 ambient.
shifted from 0.59 to 0.61 for those lms on 0.5-nm-thick SiO2 . Also, the Qf;ef f decreased from 1.6 1012 to 9.69 1011 cm 2 for H-terminated Si and from 7.67 1011 to 4.83 1011 cm 2 for 0.5-nm-thick SiO2 . The HfO2 lms with a thin SiO2 bu er layer showed a greater shift of the VF B toward the positive direction and a lower Qf;ef f than lms with H-terminated Si. Figure 8 shows EOT and J -V values for HfO2 lms grown on H-terminated Si and on 0.5-nm-thick SiO2 . RTA was performed at 600 and 800 C for 30 s in a N2 ambient. The leakage current density was measured at a gate bias voltage of VF B 1. As shown in Figure 8, the EOT values of the HfO2 lms for H-terminated Si and for 0.5-nm-thick SiO2 were 2.09 and 1.83 for the as-grown lm, 2.15 and 1.85 for the lm annealed at 600 C and 2.34 and 1.92 for the lm annealed at 800 C. The EOT values of HfO2 lms with a 0.5-nm-thick SiO2 bu er layer were lower than those of lms with Hterminated Si. This is due to the interfacial layer of the
In summary, an approximately 0.5-nm-thick SiO2 layer was formed on a Si substrate by using a remote plasma oxidation process. The HfO2 lms were deposited on Hterminated and pre-oxidized Si substrates, respectively, by using the RPALD method. A silicate layer was formed in the interfacial region during HfO2 deposition by using RPALD. The HfO2 lm with a thin SiO2 bu er layer e ectively retarded the formation of an Hf silicate layer during HfO2 lm deposition and post-deposition annealing. The Hf silicate layer in the interfacial region thickened as the RTA temperature was increased. The increase in the thickness of HfO2 generated positively xed oxide charges in the lms, which were reduced by RTA. The HfO2 lms with a thin SiO2 bu er layer showed a shift of the VF B in the positive direction, a lower Qf;ef f and a lower EOT than those with H-terminated Si. The leakage current densities of the HfO2 lms increased due to the crystallization of the HfO2 lm after RTA.
ACKNOWLEDGMENTS
This work was supported by the Korea Research Foundation Grant funded by the Korean Government (MOEHRD) (KRF-2004-005-D00166).
REFERENCES
[1] H. D. Kim and Y. Roh, J. Korean Phys. Soc. 49, S755 (2006). [2] G. D. Wilk, R. M. Wallace and J. M. Anthony, J. Appl. Phys. 87, 484 (2000). [3] M. H. Cho, Y. S. Roh, C. N. Whang and K. Jeong, Appl. Phys. Lett. 81, 472 (2002). [4] G. D. Wilk, R. M. Wallace and J. M. Anthony, J. Appl. Phys. 89, 5243 (2001). [5] M. Ritala, K. Kukli, A. Rahtu, P. I. Raisanen, M. Leskela, T. Sajavaara and J. Keinonen, Science 288, 319 (2000). [6] H. Kim and P. C. McIntyre, J. Korean Phys. Soc. 48, 5 (2006).
-1108-
Journal of the Korean Physical Society, Vol. 52, No. 4, April 2008
[7] M.-H. Cho, Y. S. Roh, C. N. Whang, K. Jeong, S. W. Nahm, D.-H. Ko, J. H. Lee, N. I. Lee and K. Fujihara, Appl. Phys. Lett. 81, 472 (2002). [8] J. Koo, J. Lee, S. Kim, Y. D. Kim, H. Jeon, D. S. Kim and Y. Kim, J. Korean Phys. Soc. 47, 501 (2005). [9] Y. Kim, J. Koo, J. Han, S. Choi, H. Jeon and C. G. Park, J. Appl. Phys. 92, 5443 (2002). [10] W. Lee, Y. Choi, K. Hong, N. H. Kim, Y. Park and J. Park, J. Korean Phys. Soc. 46, L756 (2005). [11] J. Kim, S. Kim, H. Kang, J. Choi, H. Jeon, M. Cho, K. Jung, S. Back, K. Yoo and C. Bae, J. Appl. Phys. 98, 094054 (2005). [12] J. Y. Kim, D. Y. Kim, H. O. Park and H. Jeon, J. Elec-
trochem. Soc. 152, G29 (2005). [13] M. Park, J. Koo, J. Kim, H. Jeon, C. Bae and Cristiano Krug, Appl. Phys. Lett. 86, 252110 (2005). [14] S. Kim, J. Y.Kim, J. Kim, J. Choi, H. Kang, H. Jeon and C. Bae, Electrochem. Solid-State Lett. 9, G40 (2006). [15] J. Choi, S. Kim, J. Kim, H. Kang, H. Jeon and C. Bae, J. Vac. Sci. Technol. A 24, 678 (2006). [16] S. Kim, J. Kim, J. Choi, H. Kang, H. Jeon, W. Cho, K. S. An, T. M. Chung, Y. Kim and C. Bae, Electrochem. Solid-State Lett. 9, G200 (2006). [17] C. Bae, C. Krug, G. Lucovsky, A. Chakraborty and U. Mishra, J. Appl. Phys. 96, 2674 (2004).
