Compositional characterization of very thin SiO2/Si3N4/SiO2 stacked films by x-ray photoemission spectroscopy and time-of-flight-secondaryion-mass spectroscopy techniques S. Santucci,a) L. Lozzi, L. Ottaviano, M. Passacantando, and P. Picozzi INFM Unita’ L’Aquila, Dipartimento di Fisica, Universita` di L’Aquila, via Vetoio 67010 Coppito, L’Aquila, Italy
G. Moccia, R. Alfonsetti, A. Di Giacomo, and P. Fiorani Texas Instruments Italia, Nucleo Industriale Avezzano, 67051 Avezzano (AQ), Italy
~Received 1 October 1996; accepted 13 January 1997! The chemical composition of ultrathin oxide–nitride–oxide multilayer films grown onto p-type silicon substrates and subjected to different annealing processes and to various oxidation times of the nitride layer has been studied by means of x-ray photoelectron spectroscopy and time-of-flight– secondary-ion-mass spectroscopy. Our results show that the annealing process strongly influences the bottom SiO2/Si interface allowing the saturation of the dangling bonds present at this interface and decreasing the concentration of free hydrogen. By increasing the oxidation time, a better silicon dioxide layer is formed in the topmost layer of this structure. © 1997 American Vacuum Society. @S0734-2101~97!03503-3#
I. INTRODUCTION The gate–dielectric–Si~100! interface has been one of the main issues in the development of submicron ultralarge-scale integrated circuits technology. Due to their superior electrical properties, nitrided oxides or oxynitrides are currently the subject of intensive study for gate dielectrics or metal– oxide–semiconductor capacitors. In particular, multilayered structures composed of silicon oxide–nitride–oxide ~ONO! have been introduced, in order to increase the dielectric strength of the insulator film in the DRAM storage cell capacitors1 and of the interpolydielectrics in nonvolatile memory devices.2 A lot of studies have been published on the electric characterization of this multilayer structure as a function of the ONO thickness. It has been shown that it is possible to produce a good quality ONO structure with the whole thickness down to 8–10 nm.3–5 It has been also demonstrated that for these ultrathin films it is crucial to produce the bottom silicon dioxide layer in a very controlled manner, to avoid the formation at the interface of a broad silicon oxynitride.5 Only a few papers have been published on the elemental and chemical characterization of the ONO structures.6,7 The knowledge of the chemical composition and of the depth distribution of the elements in the multilayer structure is fundamental to understanding how to modify the sample preparation in order to improve the electrical characteristics. In this work the effects of the thermal annealing in hydrogen ambient and of the thickness of the top silicon oxide layer have been analyzed by studying the distribution and the chemical bonds of the elements forming the sample as a function of the sample depth. The elemental composition has been followed by means of the time-of-flight-secondary-ionmass spectroscopy ~TOF-SIMS!, which allows us to control a!
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the atomic concentration of all the elements, including hydrogen. The existing chemical bonds have been studied by using x-ray photoemission spectroscopy ~XPS!. The analysis of the XPS data has been performed using the common concentration determination and the Auger parameter method. II. EXPERIMENT The ONO structures were grown, in a fused quartz tube reactor, onto B-doped Si~100! wafers ~10 V cm!, submitted to a standard cleaning procedure. The preparation for the four samples ~A, A*, B, and B*! is the following ~see Table I!: ~a! loading in the controlled atmosphere of the reactor where the first oxidized silicon layer is formed; ~b! low pressure chemical vapor deposition ~LPCVD! of Si3N4 , ~c! oxidation of this last layer, and ~d!, only for A* and B* samples, annealing in H2 . All the loading and discharging operations were done in ultradry nitrogen atmosphere ~,1 ppm of humidity!. After the preparation, the surface roughness and the thickness were checked by using atomic force microscopy and transmission electron microscopy ~TEM!, respectively. Both techniques showed that the samples have the same characteristics. The mean-surface roughness was 0.2560.05 nm, and the layer thickness ~measured by TEM on the sample cross section! was bottom SiO2 layer 1.860.3 nm, Si3N4 5.060.3 nm, and top SiO2 layer 2.460.3 nm. These values give us a rough estimation of the whole thickness of the structure, confirmed also by ellipsometric data, but do not allow us to appreciate, due to the quite large relative error, the very small thickness variations of the two topmost layers, induced by the different oxidation conditions of the nitride layer, as it will be evidenced in the XPS and SIMS depth profiles. The XPS measurements were performed in a PHI 1257 ESCA apparatus working at a base pressure of 531028 Pa equipped with a Mg x-ray source ~h n 51253.6 eV!, a PHI
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TABLE I. Sample preparation conditions. O Native oxide growth during the loading procedure
O Oxidation of the topmost layers of Si3N4
N LPCVD of Si3N4 layer pressure51 Torr
Annealing in H2
Sample
Time ~min!
Temp. ~°C!
Gas
Time ~min!
Temp. ~°C!
SiH2Cl2
NH3
Time ~min!
Temp. ~°C!
O2
H2
Time ~min!
Temp. ~°C!
Gas flow ~1/min! H2
A B A* B*
15 15 15 15
650 650 650 650
Nitrogen Nitrogen Nitrogen Nitrogen
9 9 9 9
780 780 780 780
15 15 15 15
45 45 45 45
17 26 17 26
900 900 900 900
9 9 9 9
4.5 4.5 4.5 4.5
no no 30 30
no no 450 450
no no 10 10
Gas flow ~1/min!
10-360 hemispherical analyzer, and an Ar1 ion gun. The depth profiles were performed by sputtering the samples using a quite low ion energy ~1.5 keV! in order to avoid any preferential sputtering, which can artificially change the composition of the analyzed surface.8 To perform the XPS analyses, we followed the features of the following elements: Si(2p), O(1s), N(1s), C(1s) core peaks and Si KL 2,3L 2,3 Auger peak. The high-energy Si KL 2,3L 2,3 Auger transition was excited by using the bremsstrahlung radiation of the source. TOF-SIMS depth profiles were collected with a PHIEVANS TFS2000 TOF-SIMS instrument. A 1 keV Cs1 primary ion beam was rastered over an area of 2503250 mm2. The secondary ions were extracted at 3000 V from the 60 mm diam central area of the crater using a 15 keV gallium ion gun and were mass analyzed with the time-of-flight spectrometer at a resolution of 9000. In this way, a high depth profile resolution ~less than 2 nm! was obtained and a very low sputtering rate, which minimizes the secondary effects of the sputtering, like diffusion of nitrogen,9 was acquired. III. RESULTS AND DISCUSSION In the Auger parameter method the energetic distance a8 between the core level ~2p for silicon! and a sharp Auger peak ~KL 23L 23 for silicon! is measured a 8 5BE ~Si2 p)1KE ~Si KL 23L 23!. This parameter is very sensitive to the chemical bond of the excited atom, for example, it goes from 1716.0 eV for pure silicon to 1713.9 eV for Si3N4 , to 1711.5 eV for SiO2 . In the case of these silicon compounds, this strong variation is mainly due to a quite large shift of the Auger peak, which shifts about 10 eV when going from Si to SiO2 , while the Si 2p variation is smaller ~about 4 eV!. The change of a8 is strictly related to the screening effects following the emission of an electron. In fact, Da852 R ex ,10 where R ex is the extra-atomic relaxation after the emission of a core electron. Therefore, any variation of a8 is a fingerprint of the change of the core-hole screening. An important experimental characteristic of this method is that any charging problem does not influence the analysis, because it influences both peaks but does not change their difference. This is actually very important for applying this technique to the study of silicon compounds, oxides, and oxynitrides. Following the J. Vac. Sci. Technol. A, Vol. 15, No. 3, May/Jun 1997
Gas flow ~1/min!
work of Wagner et al.,10 there were a lot of works showing the effects of the chemical bonds on a8. In particular, it has been demonstrated that in silicon oxide compounds ~SiOx ! there is a linear relationship between a8 ~calculated for the silicon features! and the oxygen concentration x.11,12 A similar behavior has been reported for the silicon Auger parameter in silicon oxynitrides ~SiOx Ny ! as a function of the oxygen and nitrogen concentrations.13 These results allow determining the relative concentration measuring a8 only. In Fig. 1 the application of the Auger parameter method on one of our samples, A*, is reported. The dashed lines represent the position of Si 2p ~on the left! and of Si KL 23L 23 ~on the right! for pure silicon ~shorter distance! and for Si3N4 ~larger distance!. In the lowest curves, recorded after 225 s of sputtering, the presence of two components in both the spectra is clearly evident: the predominant signal due to Si3N4 ~Si 2p at about 101.8 eV binding energy and Si KL 23L 23 at about 1612.1 eV kinetic energy! and the small structures due to pure silicon ~99.3 and 1616.7 eV, respectively!. This indicates that, after the reported sputtering time, we have a well stoichiometric silicon nitrides layer. The small clean silicon signals come from the substrate because the thickness of the very thin structure is lower than the electron escape depth. Increasing the sputtering time, we observe the growth of the Si signal and the shift and the lowering of the Si 2p and in the Si KL 23L 23 features due to the silicon compounds. Because of the small signal-to-noise ratio, after 450 s of sputtering no distinct peak due to any silicon compound is visible in the Auger spectra, but only a long tail in the low kinetic energy side of the pure silicon feature is present. Therefore, in these cases, in order to find the position of the silicon compound Auger peak, a fit procedure has been applied. After about 675 s, all the signals attributable to silicon compounds miss and the only visible features are due to pure silicon. For all the samples, the behavior of the Auger parameter, together with the intensity of the characteristic peaks of all the elements ~Si 2p, O 1s, and N 1s!, has been followed as a function of the sputtering time and has been reported in Fig. 2 for samples A and A* and in Fig. 3 for samples B and B*. For each sample a8 ~upper part! and the atomic concentration ~lower part! have been reported. The atomic concen-
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FIG. 1. Silicon 2p and KL 23L 23 Auger spectra as a function of the sputtering time for the A* sample. The dashed lines represent the distance between these two features for clean silicon ~upper line! and for Si3N4 .
tration has been determined by calculating the intensity under the characteristic peaks and then weighting them through the atomic sensitivity factors.14 For the Auger parameters there are several sputtering times where more than one point have been determined. This means that it was possible to separate the peaks due to the different silicon compounds. Starting from the top of the sample and looking at the a8 parameter, the presence of oxidized silicon is evident. Sample A does not show a SiO2 stoichiometric layer, because a8 has a higher value with respect to the silicon dioxide one, and it corresponds to a SiOx51.75 layer. After the annealing ~sample A*!, a8 is very close to the pure silicon one, showing that the upper layer is almost completely oxidized. Instead, for the B and B* samples, a8 has always the silicon dioxide value. This is surely the effect of the longer oxidation time of the topmost layers used for these two samples ~see Table I, 9th column!. After this oxidized layer, all the samples show a quite wide sputtering time zone in which a good Si3N4 layer has been formed. This zone goes up to about 450 s for A and B samples, while it stops to about 300–320 s of sputtering for A* and B* samples. The main differences between the annealed ~A* and B*! and nonannealed ~A and B! samples are observed at the interface of the Si3N4 layer with the bottom oxide. In the nonannealed samples, a8 shows a slight reduction from the Si3N4 value, suggesting the formation of a silicon oxynitride layer. This presence can be explained as due to the retarded JVST A - Vacuum, Surfaces, and Films
growth of the Si3N4 layer in the presence of nonstoichiometric silicon oxide, which produces the growth of a oxynitride layer.15 Instead, in the same depth position, for the A* and B* annealed samples, a8 grows over the Si3N4 line. An a8 value higher than the Si3N4 one cannot be caused by the formation of any silicon oxynitride compounds, because for these compounds a8 ranges between 1711.5 eV ~SiO2! and 1713.9 eV ~Si3N4!.13 An higher value could be attributed to a strong understoichiometric silicon oxide ~in our case, SiOx50.8! but, as recently reported,16 its morphology, with this oxygen concentration, cannot be represented by using the random bonding model anymore, but as islands of pure silicon embedded in a silicon oxide matrix. In this case, the insulating behavior of this layer would be almost completely destroyed and the electric reliability of the whole structure considerably lowered. Therefore, considering that the electrical behavior of annealed samples improves with the heat treatment,17 we are more induced to think that the reported behavior of a8 for the annealed samples A* and B* can be assigned to the presence of bonds between silicon and hydrogen. In fact, it has been shown that at the SiO2/Si interface a lot of Si bonds my be saturated by hydrogen atoms.18,19 In this case, the screening effects of the valence electrons on the Si 2 p core holes, created either by the photoemission process or by the Auger decay, should increase, and, as a consequence, a8 will grow up, too. After this inter-
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FIG. 2. XPS depth profile of A and A* samples ~lower curves!. The dashed, continuous, and dotted lines represent, for each sample, the relative concentrations of Si, O, and N, respectively. The upper curves show the modified Auger parameter of the same samples reported as a function of the sputtering time.
FIG. 3. XPS depth profile of B and B* samples ~lower curves!. The dashed, continuous, and dotted lines represent, for each sample, the relative concentrations of Si, O, and N, respectively. The upper curves show the modified Auger parameter of the same samples reported as a function of the sputtering time.
face zone, only a8 for pure silicon is detectable for all the samples. Looking at the atomic concentration figures, a confirmation of the overall behavior of the silicon compounds in the ONO structures can be obtained. But, in this case, it is quite hard to understand how the different elements are bound. In fact, because of the charging problems after the photoemission process, it is not easy to understand what the binding energies are of the different peaks. Moreover, as we have already underlined, the shift of the Si 2 p peaks is not so big to allow distinguishing between silicon oxides and oxynitrides, while the peak positions of nitrogen and oxygen 1s are even less sensitive to the chemical bonds. In Figs. 4 and 5 are reported the TOF-SIMS results of our investigated samples. An accurate comparison of the results for samples A and A* can be summarized with the following observations: ~i! for the annealed sample, we note an evident
lowering of hydrogen in the whole profile and its depression at the SiO2/Si interface, with respect to the nonannealed one; and ~ii! the increase and the narrowing of the first SiO2 feature related to the top oxide layer, and the narrowing of the feature related to the bottom oxide, of sample A* with respect to A. This behavior is confirmed, even if in a less evident way, from the in-depth evolution of the oxygen content. For samples B and B*, the results can be summarized in the following way: ~i! a larger amount of hydrogen in the nonannealed sample due to the longer oxidation process in the O21H2 flow and its very considerable lowering in the annealed sample B*, going approximately towards the A* value; and ~ii! the two SiO2 maxima are closer than in samples A and A*, still imputable to the longer oxidation time, which provides a deeper penetration of the oxygen. For all the samples, the intensity of the SiN signal would suggest the presence of nitrogen in the whole structure, in
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FIG. 4. TOF-SIMS depth profile of A and A* samples. The different symbols refer to: ~s! oxygen, ~1! SiO2 , ~h! nitrogen, ~m! silicon, ~!! SiN, and ~d! hydrogen.
FIG. 5. TOF-SIMS depth profile of B and B* samples. The different symbols refer to: ~s! oxygen, ~1! SiO2 , ~h! nitrogen, ~m! silicon, ~!! SiN, and ~d! hydrogen.
particular, at the SiO2/Si interface and in depth of the silicon substrate. But, this is due to the different sputtering yield of the molecules. If we scale down the SiN signal taking into account the XPS depth profiles, its intensity is comparable with the SiO2 one, confirming its appreciable presence only between the SiO2 maxima while its presence can be excluded in the interfacial region of the SiO2/Si layer. The SiN tendency is confirmed by the shape of the N signals.
The present TOF-SIMS results compared with those by XPS allow us to conclude that the annealing process lowers the whole hydrogen content of the sample but favors the adjusting of the unsaturated dangling bonds into the SiOx Ny bottom layer. In the annealed A* and B* samples, hydrogen is present as bonded hydrogen, particularly at Si3N4/SiO2/Si interfaces because it increases the silicon Auger parameter. On the contrary, the more conspicuous number of atomic hydrogen in the not annealed samples can be considered as
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trapped hydrogen because its presence does not influence the silicon Auger parameter. IV. CONCLUSIONS An accurate XPS analysis using the ‘‘Auger parameter method’’ has been performed on four ultrathin ONO structures prepared in different conditions. This analysis has allowed us to characterize both the stoichiometry and the presence of various silicon compounds as a function of depth in the ONO structure, showing the effect of the different growth processes on the single layers and on the whole structure. In particular, we can point out that studying the variations of the silicon Auger parameter with the support of the TOFSIMS measurements, which allow the very important quantification of the hydrogen in the depth of the samples, we are able to put in evidence the most important difference between the annealed and not annealed samples, i.e., the presence of bonded hydrogen at the interfaces of the bottom oxide layer. The obtained results validate the annealing procedure in hydrogen as a correct technique able to reduce the presence of hydrogen related defects near the oxide interfaces, which strongly influences the reliability of the whole dielectric layer. 1
G. Q. Lo, S. Ito, Dim-Lee Kwong, V. K. Mathews, and P. C. Fazan, IEEE Electron Device Lett. 13, 372 ~1992!.
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