Christophe Dubourdieu

V Vallée

M Loup

O R Bonvalot

Hervé Joubert

O Roussel

Renault

C Durand

C Vallée

V Loup

M Bonvalot

O Joubert

H Roussel

O Renault

yttrium silicate thin films deposited by pulsed liquid-injection plasma-enhanced metal-organic chemical vapor deposition

I. INTRODUCTION

Materials with high dielectric constant have been introduced in integrated circuits as stated in the International Technology Roadmap for Semiconductors 1 for applications such as CMOS (complementary metal oxide semiconductor) gate dielectrics and MIM (metal-insulator-metal) capacitors. On the one hand, integration of MIM structure for the next generation of radio frequency chips is needed. Increasing the dielectric constant of the dielectric used in MIM structures allows an increase of the capacitance density. In addition to a high dielectric constant, the material must satisfy a large number of requirements different to CMOS gate oxide specifications such as a high degree of capacitance linearity, a high dielectric breakdown ͑E bd Ͼ 5M Vc m -1 ͒, and a low thermal budget ͑Ͻ450°C͒ specific to back-end processes. Efforts were first engaged to obtain MIM structures using Ta 2 O 5 . However, large leakage current has been observed [2][3][4] and MIM capacitors are tested rather with "medium " [such as HfO 2 , 5 Al 2 O 3 , 6 and Y 2 O 3 7 or with "very high " dielectrics (such as SrTiO 3 or BaTiO 3 ). On the other hand, the aggressive scaling of CMOS devices beyond the 0.1 m generation is driving the SiO 2 gate dielectrics to its physical limits, due to excessive direct tunneling leakage current. High dielectric constant (high-k) materials are attracting candidates for SiO 2 replacement. Indeed, they allow an increase of the physical thickness of the dielectric layer in the device wile preserving the capacitance value. 8 Most high-k studied materials are based on metal oxides ͑MO x ͒ and their silicates ͑M x Si y O z ͒. For instance, currently several dielectric materials such as Al 2 O 3 , ZrO 2 , and HfO 2 , as well as mixed oxides containing SiO 2 (silicates) or Al 2 O 3 (aluminates) are being studied extensively. Y 2 O 3 and rare earth oxides such as La 2 O 3 ,Pr 2 O 3 , and Gd 2 O 3 are also proposed as good candidates to replace SiO 2 thanks to their reasonably high permittivity ͑10Ͻ Ͻ 30͒, compatibility when in direct contact with Si, 9 good thermal stability, and relatively large conduca)

Electronic mail: cdurand@cea.fr 1 tion band offset (2.3 eV for Y 2 O 3 ). 10 Although previous studies have revealed that these oxide films absorb water [11][12][13] and present significant charge density responsible for deteriorated electrical characteristics, interesting results have been found on Y 2 O 3 and Y-based silicate. For instance, Chambers and Parsons pointed out that yttrium has a very high affinity for oxygen and showed that an annealing results in the thickness reduction of SiO 2 to form a silicate layer. 14,15 Copel et al. take advantage of the high stability of Y-O bondings to accurately control the silicate formation by developing a solid-state technique: the unwanted interfacial SiO 2 is consumed by yttrium oxide to form a silicate layer, which seems to be thermodynamically favorable. 16,17 In spite of a lower permittivity than metal oxides, silicates are expected to have beneficial properties with lower charge densities compared to metal oxides thanks to similarities with SiO 2 . To avoid interfacial reaction, the high-k layer can be capped to protect the film from water ambient absorption (OH species favor the interface formation), as proposed by Stemmer et al., 18 or a prenitrided sample can also be used to prevent Si diffusion from the substrate to form a silicate. 15 Concerning deposition methods, physical vapor deposition (PVD) processes were traditionally used to deposit metal oxides. Y 2 O 3 has been deposited by sputtering, [19][20][21][22] by ionized cluster beam deposition, [23][24][25] by ion beam sputtering, [26][27][28] by pulsed laser deposition, 29 by electron beam evaporation, [30][31][32][33] and by molecular beam epitaxy. 34,35 However, most of PVD techniques may not be suited in large scale CMOS processing on 200 or 300 mm Si wafers. Sputtering processes may not be suitable because of possible electrical damages induced in the extremely sensitive channel region of CMOS devices. Chemical vapor deposition (CVD) methods are usually preferred. However, a key point of CVD processes (named MOCVD when precursors are metal-organic compounds) deals with how to provide a stable and reproducible precursor supply to ensure a perfectly controlled growth of thin metal-oxide films. The use of a pulsed-liquid-injection source has been proposed to overcome this difficulty. [36][37][38][START_REF] Dubourdieu | Crystal Growth in Thin Solid Films: Control of Epitaxy[END_REF] The use of plasma assistance has been also reported for MOCVD process on HfO 2 , [START_REF] Choi | [END_REF]41 ZrO 2 , 42 Y 2 O 3 , 43 Hf, or Zr-based silicate 44 and for atomic layer deposition ͑ALD͒ process on ZrO 2 45,46 and SrTiO 3 . 47 Even if processes using plasma assistance are not well understood and optimized in terms of parameter adjustments, it looks very attractive. For instance, Kim et al. showed that the oxygen plasma use in an ALD process improved ZrO 2 film quality with a low carbon incorporation, a low leakage current, and negligible hysteresis compared to films deposited by conventional ALD. 46 Niu et al. achieved deposited yttrium oxide with oxygen plasma assistance and showed the presence of a three-layer structure Y 2 O 3 / silicate/ SiO 2 after annealing at 900°C under N 2 . 43 Advantages of plasma assistance are, first, a decrease of the deposition temperature for low thermal budget applications as MIM capacitors and, second, a possible use of plasma effects as predeposition or postdeposition treatments.

In this paper, we investigate the formation of yttrium oxide and yttrium silicate films by plasma-enhanced MOCVD ͑PE-MOCVD͒ using a pulsed-liquid-injection source. This type of source is hooked up to a plasma assisted MOCVD chamber. The role of plasma is to enable a lowering of the deposition temperature, which is particularly interesting for the deposition of dielectric films compatible with back-end thermal budget constraints ͑Ͻ450°C͒. We will indeed show that yttrium oxide deposition at below 380°C can be performed. The films chemical structure is analyzed by x-ray photoelectron spectroscopy ͑XPS͒ and by Fouriertransform infrared spectroscopy ͑FTIR͒ in the transmission mode. X-ray diffraction ͑XRD͒ is also used for the study of thick crystallized films. Results on thick ͑ϳ40 nm͒ and ultrathin ͑ϳ4n m ͒ films are presented. The interface nature and formation are analyzed and the influence of deposition parameters such as the reactive species incoming frequency is discussed. It is shown that the injection frequency plays an important role in the formation of the silicate and of the interfacial layer. The role of temperature is also investigated. Finally, electrical results are shown for as-deposited and annealed films.

II. EXPERIMENTAL PROCEDURE

A. The deposition process

Yttrium oxide films were grown by pulsed-injection plasma-assisted MOCVD. The experimental setup is shown in Fig. 1. The gas phase is generated using a pulsed-liquidinjection source. The liquid precursor is maintained at room temperature in a closed vessel pressurized under an inert gas (N 2 at 2 bars). The liquid reservoir is connected to a highspeed electrovalve (a car injector) driven by an electrical pulse, which is controlled by computer. Microamounts of liquid solution are injected into an evaporator held at 190°C and are "flash" evaporated. The precise amount of reactive species in the reactor is controlled accurately by the injection opening time and injection frequency. With this innovative delivery scheme, reasonable vapor pressures of thermally unstable precursors can be obtained and kept stable over the deposition run and thus complex oxides can be reproducibly grown. 37,[START_REF] Dubourdieu | Crystal Growth in Thin Solid Films: Control of Epitaxy[END_REF] For the time, this delivery scheme was combined, in this study, to a plasma-assisted CVD chamber. The plasma is driven by a 380 kHz power supply and pressure in the chamber can be adjusted between 27 and 670 Pa (0.2 and 5 Torr). A showerhead ensures a uniform distribution of gas stream over the heated wafer substrate (resistive heating).

For the present study, tri(2,2,6,6-tetramethyl-3,5heptanedionate) yttrium [abbreviated as Y͑tmhd͒ 3 ] was used as a precursor. The metal-organic precursor, which is a solid powder, was dissolved in a solvent (cyclohexane). The gaseous reactive species in the evaporator were transferred via a heated line ͑195°C͒ to the reaction chamber using a neutral carrier gas [Ar at 30 SCCM (SCCM denotes cubic centimeter per minute at STP)]. The plasma gases were formed from Ar and O 2 . The reaction chamber was maintained at low pressure, typically 0.27 kPa ͑2 Torr͒. The films were deposited on p-type 8 in. Si͑100͒ substrates with 0.8 nm thermally grown SiO 2 . The standard deposition conditions are given in Table I. Postdeposition annealing were performed on selected samples to investigate their electrical properties.

B. Characterization techniques

XPS, FTIR, and XRD were used to identify the bulk as well as the interface chemistry and structure of the films. The XPS spectra were collected with an Al K␣ x-ray source ͑hv = 1486.6 eV͒ and an electron energy analyzer operating in a constant pass energy mode (50 and 20 eV for the survey and core level spectra, respectively). All studies were done at 45°as takeoff angle except for the angle-resolved studies. Spectral deconvolution was performed to extract the Si 2p, O 1s, C 1s, and Y 3d peak intensities. Chemical compositions were derived from the areas of the different XPS spectra weighted by corresponding sensitivity factors available from the database of the acquisition software. To check the XPS quantification, both XPS an RBS (Rutherford backscattering spectroscopy) were done on an Y 2 O 3 layer, which was elaborated on silicon by thermal MOCVD at 700°C. The [O]/[Y] ratio was found identical in RBS and XPS analyses using the Y 3d and O 1s sensitivity factors from XPS database. For the interfacial oxide thickness, we used a SiO 2 escape depth model from Si 2p spectra. Si 2p spectra were fitted with a Shirley background and five chemical states (Si 0 ,S i 1+ ,S i 2+ ,S i 3+ ,S i 4+ ), which are separated by approximately 1 eV each, with an increasing FWHM (the full width at half maximum). The fitting peak characteristics are presented in Table II. The peak from the substrate is located at 99.3 eV and is denoted Si 0 . The interfacial layer thickness of oxidized silicon t ox is usually calculated using Eq. (1), 48

t ox = ox sin͑͒lnͩ R R ϱ +1ͪ , ͑1͒
where is the photoelectron takeoff angle, R is the ratio A ox /A Si of the sample ͑A Si being the relative areas of the Si 0 component and A ox the sum of the relative area of the Si n+ components͒, R ϱ is the ratio A ox /A Si in the case of a thick thermal SiO 2 and bulk Si sample. The photoelectron inelastic mean-free path ox is difficult to estimate for our interface. Nevertheless, the ratio R evolution gives a good qualitative image of the interface thickness because ox , and R ϱ can be considered constant. Therefore, the ratio R will be used to describe the oxidized Si-based compounds' thickness evolution. High-k compounds are insulating, which leads to charging effects as evidenced by a shift to higher binding energies of XPS peak positions. This effect has been systematically corrected by positioning, when detectable, the Si 2p peak from the substrate at 99.3 eV. This correction leads to C 1s peak position from surface contamination at 285.7 eV. When no Si 2p contribution from the substrate could be detected by XPS, the charging correction was carried out relative to this surface C 1s peak position. The samples' surfaces were treated by an Ar + ion beam ͑1.5 keV @ 10 A during 30 s͒ only for thick layers before XPS analysis to remove the organic surface contamination. We initially determined Y 2 O 3 peak positions with a ceramic Y 2 O 3 sample ͑sintered Y 2 O 3 powder annealed at 1500°C during 24 h͒. The Y3 d ͑3/2͒-͑5/2͒ doublet is located at 156.7 and 158.7 eV in agreement with literature data on Y 2 O 3 reference compounds. [START_REF]Handbook of x-ray Photoelectron Spectroscopy[END_REF] Also, the solid precursor powder of Y͑tmhd͒ 3 was analyzed by XPS with a monochromatic source at 35°takeoff angle in order to locate the Y energy binding in the precursor molecule environment. A Study of the bonding environments in the grown material and at the interface was also carried out using FTIR. A Bruker IFS 55 Fourier-transform infrared spectrometer was used.

The resolution and number of scans were set at 4 cm -1 and 500, respectively. Capacitance-voltage ͑C-V͒ and currentvoltage ͑I-V͒ measurements were conducted on a HP 4284 A precision LCR meter and HP 4156 A meter, respectively. The structure was metal-oxide semiconductor formed by evaporating gold on the films through shadow masks. Four different surfaces ͑0.354ϫ 10 -2 , 0.694ϫ 10 -2 , 1.71ϫ 10 -2 , 2.78ϫ 10 -2 mm 2 ͒ were defined. The equiva- lent oxide thickness ͑EOT͒ was extracted from the accumulation region ͑no quantum mechanical correction was applied͒.

III. RESULTS AND DISCUSSION

A. Plasma enhanced MOCVD: Plasma effects

In order to work out the influence of the plasma and of the substrate temperature on the process, we performed three different experiments: (a) deposition using plasma assistance ͑70 W͒ but without heating the substrate, (b) deposition at 350°C without plasma, and finally (c) deposition using both thermal ͑350°C͒ and plasma assistance ͑70 W͒. The XPS analyses of the three samples are shown in Fig. 2. As expected, there is no deposited yttrium for the run performed without heating the substrate chuck (the actual substrate temperature is estimated to be lower than 150°C)[Fig. 2(a)].At 350°C, when no plasma is used, yttrium is not detected as well [Fig. 2(b)]. Clearly the plasma is needed to form an yttrium oxide film from the ␤-diketonate Y͑tmhd͒ 3 precursor at such a low temperature [Fig. 2(c)]. Indeed, the Y 3d, 3p, and 3s signatures are only observed in this case. A study of several key parameters (pressure, O 2 partial pressure and flux, plasma power) was carried out to understand deposition mechanisms and to optimize the process. Apparently, the plasma promotes the precursor decomposition but also induces two other opposite effects: (i) carbon deposition from the solvent and from the precursor organic chains dissociation and (ii) carbon etching. For several sets of deposition conditions, we determined threshold values beyond which a thick carbon layer is deposited instead of yttrium oxide. When the plasma density is high, it strongly decomposes carbon chains from the solvent and the precursor. If the O 2 partial pressure is too low, all carbon groups on the substrate surface cannot be etched and thus carbon species are incorporated during the growth. The use of oxygen in the plasma gas phase ensures minimum carbon contamination and oxygen stoichiometric adjustment of the deposited layer. However, a too high oxygen partial pressure in the plasma and a too high plasma power can also lead to an undue increase of deleterious SiO 2 at the interface. Therefore, optimization of the deposition conditions should take these different points into account (Table I). The plasma power was fixed to 70 W in order to limit the solvent decomposition and wafer oxidation. The working pressure was fixed at 0.27 kPa ͑2 Torr͒ in order to maintain a reasonable thermal exchange for the substrate heating, while, at the same time, to decrease energetic ion bombardment at the surface of the growing film, which may induce a high interfacial state density. Indeed, the plasma is localized near the showerhead and the sheaths are collisional. In this case, the plasma is similar to a "remote plasma." the reference material, the Y 3d 5/2-3/2 doublet position (respectively, 156.7 and 158.7 eV) and O 1s peak ͑529.5 eV͒ are consistent with literature data for Y 2 O 3 . The peak shoulder around 531.8 eV has also been observed by several authors and has been attributed to O-H bonding 14 or oxygen trapping. 29 Before sputtering the reference ceramic, this peak was higher relatively to the main one at 529.5 eV, indicating that surface bonds are involved. For the PE-MOCVD film grown at 350°C, the Y 3d ͑5/2͒-͑3/2͒ doublet is located at 158.5-160.5 eV [Fig. 3(a)]. This could correspond to Y-O-Si bonds, however, as we do not observe any peak in the Si 2p region, it is attributed to Y-O-C bonding and H contamination. The Y-O-H bond has already been reported to explain the shifted Y 3d doublet. [START_REF] Ulrich | [END_REF] However, in our case, FTIR analysis (Fig. 4) does not show any O-H signature ͑3000-3700 cm -1 ͒. The carbon bonding to oxygen is expected to shift the Y 3d peak toward higher binding energies as compared to Y-O binding in Y 2 O 3 by considering the electronegativity of carbon. The shift of Y-O-C bonding should also be larger than for Y-O-Si because C has a higher electronegativity compared to the Si one (respectively, 2.5 and 1.8 eV). For the organic precursor Y͑tmhd͒ 3 , the shift is indeed quite important as yttrium is bonded to three tmhd ligands as shown in the insert of the Y 3d region, Fig. 3(a The second peak at 533.0 eV comes from O-C bonds in the deposited layer and also from organic surface contamination. Indeed, after sputtering the surface, this peak decreases to the same level as the 530.6 eV one. The C 1s region [Fig. 3(c)] exhibits two distinct contributions located, respectively, at 285.7 and 291 eV. The first one comes only from organic surface contamination, because this peak disappears after sputtering the surface. On the contrary, the second peak corresponds to carbon in the film. As a conclusion, the thick films grown at 350°C using an injection frequency of 5 Hz contain Y-O bonds with no indication of silicate-type bonds but rather C and H environments. Incorporation of organic residue in the film is likely to occur during deposition at low temperature due to incomplete decomposition of the precursor. This is confirmed by the FTIR spectrum (Fig. 4) that shows a band at 840 cm -1 and two other broad bands centered on 1520 and 1400 cm -1 . These bands are assigned to C-O and C-H stretching modes.

B. Properties of thick yttrium oxide films

The previous film was annealed under Ar or O 2 flow at 700°C or under H 2 flow at 450°C for 30 min. The /2 x-ray diffraction diagrams (Fig. 5) indicate crystallization of the film at 700°C, whereas it is still amorphous (or nanocrystallized) at 450°C. XPS spectra are shown in Fig. 6 shifted to 529.5 eV as in Y 2 O 3 crystal. Moreover, the 700°C annealings reduce the C impurities in the layer, since the intensity of broadbands around 1520 and 1400 cm -1 and band at 840 cm -1 are considerably attenuated on the FTIR spectra (Fig. 4). This effect is also visible with XPS analysis on the O 1s region [Fig. 6(b)] where the contribution attributed to O-C bonds (at 533 eV) is reduced compared to O -Y peak near 530 eV. We note on FTIR spectra a significant increase of the peak at 1080 cm -1 corresponding to transversal optical mode of Si-O bond in silicon dioxide for 700°C O 2 annealing. To conclude, a high temperature annealing leads to the crystallization of the as-deposited films and to a decrease in carbon content.

C. Analysis of thin films: Effects of injection frequency on the interface layer and silicate formation 1. XPS analysis

By varying the injection frequency (from1upto10Hz), we intended to vary the reactive species supply rate at the substrate surface, which might influence the interface layer growth. The standard conditions were used for all other parameters (Table I). For each experiment, the injected volume was the same (1000 injections), which corresponds to a layer thickness thin enough ͑Ͻ10 nm͒ to detect the silicon substrate by XPS, except for the 10 Hz film. The thickness was found to increase with increasing injection frequency. XPS was used to study both the film and the interface (no Ar sputtering was used prior to measurements in order to prevent film or interface degradation). Figure 7 presents the Si 2p, Y 3d, and O 1s XPS regions of layers grown at increasing injection frequency. The Si 2p spectra [Fig. 7(a)] are normalized to the Si 0 peak area (the substrate peak Si 0 is at 99.3 eV). If not normalized, the Si 0 peak decreases with increasing injection frequency, indicating an increase of the film thickness. A bare SiO 2 ͑0.8 nm͒ / Si substrate is also shown for comparison. The corresponding Si-O peak is located at 103.3 eV, which is the expected binding energy for SiO 2 . Clearly from Fig. 7(a), the Si-O peak for the different films is shifted towards lower energies than the binding energy of SiO 2 , which can be assigned to Y-O-Si type of bonds. The expected range of Si 2p in Y-O-Si binding is indeed in the 101.5-103 eV range, depending on the exact silicate composition. [START_REF] Moulder | Handbook of x-ray Photoelectron Spectrocopy[END_REF] Also the FWHM of the Si-O Si 2p peak, which is of 1.6 eV for SiO 2 in the bare substrate, is strongly, increased for the films (to 2.2 eV), which indicates a wide range of Si-O bonding in the films. Finally, the normalization to the silicon substrate Si 2p peak area clearly emphasizes the increasing content of silicate-based compounds in the films with decreasing injection frequency. This point will be specified later (Sec. III C 2). In the Y 3d region [Fig. 7(b)], the Y 3d 5/2 and Y 3d 3/2 peaks are located at, respectively, 158.7 and 160.7 eV for the 1 Hz film. These positions are coherent with Y-O-Si bonds, which have been evidenced also on the Si 2p region. However, Y-O-C bindings cannot be precluded at this point as we showed that the corresponding Y 3d doublet appears for similar energies [Fig. 3(a)]. The Y 3d splitting is progressively less pronounced with increasing injection frequency (whereas its lo-cation is only slightly changed), indicating a much wider Y-O bonding environment, probably arising from both Y-O-Si and Y-O-C type of bonds. On the other hand, the 1 Hz sample consists of a well-defined bonding (presumably a silicate compound with no secondary yttrium carbonate phase as shown hereafter in the following section). No evidence for Y-Si bonds from silicides was found on the Y 3d or Si 2p regions (expected at ϳ156.5 and 98.7 eV, respectively). On the O 1s region [Fig. 7(c)], the peak shape also changes progressively with the injection frequency. For high frequencies (5or10Hz), two separated peaks appear. One is at 533 eV ͑FHWM= 2.0 eV͒, which corresponds to the binding energy in SiO 2 and, as we showed it previously [Fig. . Si 2p spectra are normalized on Si 0 peak area (peak from substrate) in order to estimate the oxidized Si-based compound thickness, since the well-known method for measuring the thickness of oxide film by XPS is based on the area ratio between the oxidized silicon A Si-O and substrate silicon A Si . The spectrum of initial substrate ͓Si/ SiO 2 ͑0.8 nm͔͒ is also done for comparison. contribution around 531.5 eV is consistent with a silicate compound 14 and confirms the observation made on the Y 3d and Si 2p regions.

Angle-resolved XPS analysis

In order to get some insight on the nature and chemical structure of both the interface (in contact with the silicon substrate) and film, we performed angle-resolved XPS (AR-XPS) on the 1 and 5 Hz films. Three different takeoff angles were studied (90°and two grazing incidence anges: 30°and 15°). Note that, when the takeoff angle decreases, the probed depth decreases as well, so that surface contributions are enhanced. Figure 8 shows the Si 2p, Y 3d, and O 1s XPS regions for both samples. The Si 2p spectra intensity is normalized on the Si-O peak area. Thus, the substrate Si 0 peak clearly decreases with decreasing takeoff angle as expected (for the 5 Hz sample, the substrate is even not probed at 15°). A quantitative analysis of the Si-O Si 2p peak was carried out by considering four Si oxidized states (Si 1+ ,S i 2+ ,S i 3+ , and Si 4+ , see Table II). The result of the fitting procedure is given in the inserts on the Si 2p regions [Figs. 8(a) and 8(b)] by a relative intensities bar graph of each Si n+ (n=1-4) contribution. The Si 3+ contribution is shifted by 2.8-3.0 eV compared to Si 0 . This value is larger than the common value reported for the Si 3+ suboxide ͑2.6 eV͒ in a conventional SiO 2 / Si interface. [START_REF] Green | [END_REF] The fitting using a imposed 2.6 eV shift position value for the Si 3+ component was clearly not adapted. The presence of such a component of oxidized Si, abnormally shifted relative to Si 3+ in Si/ SiO 2 , is attributed mainly to silicate-type bondings Si-O-Y (this does not mean that we exclude the presence of substoichiometric SiO x , but the corresponding contribution is overlapped by the silicate one due to our limited energy resolution). Other reasons support the Si 3+ assignment to mainly silicate bonds. First, the Y 3d peak position at 158.7 eV is consistent with silicate Y-O-Si bonds. Second, the Si 3+ intensity from substoichiometric SiO x compared to the other oxidized contributions (i.e., Si 1+ ,S i 2+ , and Si 4+ ) is limited, since it represents only 11.3% of the Si-O signal for the initial SiO 2 ͑8Å ͒ / Si substrate. Even when the initial substrate is exposed under O 2 plasma during 1500 s (under the standard plasma chamber parameters presented in Table I) without injection, an increase of contribution from substiochiometric SiO x (i.e., Si 1+ ,Si 2+ , and Si 3+ ) is not observed and the Si 3+ peak represents 8.2% of the Si-O signal. On the contrary, when the yttrium precursor is injected, the Si 3+ is strongly enhanced, since the Si 3+ signal represents 46.5% and 39.8% of the Si-O peak on the 5 and 1 Hz sample, respectively. Such an increase is likely to arise mainly from a silicate contribution. Electron energy-loss spectroscopy measurements have been done and confirmed the presence of silicate (not shown here). It is worth noting that a similar fitting procedure was employed to study Al 2 O 3 / Si interface in Ref. 53. As the deposited film is amorphous, silicate-type bindings is most likely in multiple bonding environment. Therefore, a part of the Si 2+ contribution is certainly assigned to silicate-type bondings. As the O 2 plasma under the standard conditions is not energetic, the incompletely oxidized silicate presence is a reasonable hypothesis.

For the 5 Hz sample, no shift of the Si-O peak is observed when the incidence angle changes from 90°to 30°. No significant change is also observed in the relative intensity of the Si 3+ and Si 4+ contributions between 90°and 30°[ insert of Fig. 8(a)]. Hence, the 5 Hz film is composed of an yttrium silicate, which is probably present at the interface (or on top of a very thin SiO 2 layer that cannot be resolved by changing the takeoff angle), Since the Si 2p peak disappears at 30°, no silicon atom is present at the top surface. Thus, the silicate film does not extend up to the top surface and the top layer consists of Y x O y C z film. Indeed, the Y 3d angleresolved peak exhibits a slight shift towards higher energies (+0.5 eV from 158.5 to 159.0 eV) with decreasing takeoff angle that corresponds to an Y-carbonate signature enhancement compared to the Y-silicate signal, since at 15°no silicate is involved in the measured response (accordingly to the Si 2p region). These observations are consistent with the presence of two kinds of bindings distributed differently in the sample: Y-O-C (top surface) and Y-O-Si (interface).As we already mentioned it for the thick sample analysis (Fig. 3), the binding energies of Y-O-Si and Y-O-C are quite close. The peak located at 532.8 eV, which does not shift with decreasing angle, indicates that the C contamination is uniformly distributed in the film.

For the 1 Hz sample, the Si-O peak from the Si 2p region is significantly shifted ͑-0.5 eV͒ towards lower binding energies for decreasing takeoff angle, which clearly indicates a distribution of different silicon-oxide based compounds within the depth of the film. The quantitative analysis [insert of Fig. 8(d)] shows a strong increase of silicate signals (Si 2+ and Si 3+ ) and a clear decrease of the SiO 2 contribution ͑Si 4+ ͒ at grazing angles. The sample appears composed of a SiO 2 interface with the silicon substrate and of a silicate layer on top of it. The increasing contribution at 531.5 eV [Fig. 8(f)] with decreasing takeoff angle is consistent with this picture. Since no shift is observed on the Y 3d peak [Fig. 8(e)], there is no Y x O y C z layer.

These characterizations lead us to propose two different microstructures for the 1 and 5 Hz films, which are depicted in Fig. 9. At low injection frequency ͑1H z ͒, the film is composed of an yttrium silicate layer with an interfacial SiO 2 layer in contact with silicon ͑Si/ SiO 2 ͑ϳ4n m ͒ /Si x O y Y z ͒.At higher frequencies, the interfacial layer is rather an yttrium silicate (silicon rich) with an yttrium oxide layer on top of it with carbon impurities ͑Si/ Si x O y Y x ͑ϳ2n m ͒ /Y x O y C z ͒. This picture was confirmed by transmission electron microscopy and electron energy-loss spectroscopy investigations. However, if the 1 Hz sample is grown thicker ͑ϳ40 nm͒, a top Y x O y C z film is grown too.

Discussion

The injection frequency (i.e., the flux of reactive species) has a strong effect on the plasma density and composition, which may influence the growth mechanisms. We propose the schematic reactions (shown below) during the deposition. These reactions are inspired by those proposed by Niu et al., 43 Si

+ O* + R → SiO 2 +R, ͑2͒ Y-O-RЈ +O* → Y-O-C + RЉ, ͑3͒ Y-O-RЈ +Si+O* → Si-O-Y + RЉ. ͑4͒
The reactions are written schematically, and are not balanced. Reaction (2) is the silicon oxidation by the energetic oxygen species (denoted O*) from the oxygen plasma that is enhanced by volatile organic compounds, denoted R. Indeed, oxidation rate is strongly limited ͑0.03 nm/ min͒ when experiments are carried out without injection under standard process conditions (Table I), whereas it is strongly increased when solvent (cyclohexane) is injected into the oxygen plasma (0.12 nm/ min at 5 Hz). 54 Reaction (3) is the formation of yttrium oxide carbonate by decomposition and oxidation reaction of the organic Y͑tmhd͒ 3 precursor molecule thanks to the oxygen plasma. The precursor molecule is represented by Y-O-RЈ, where RЈ is the organic ligand [reactions (3) and (4)]. Only a schematic reaction is proposed here since decomposition paths of precursor molecules such as Y͑tmhd͒ 3 is far too complicated to be described and is not known in detail. RЉ denotes compounds resulting from the precursor decomposition, i.e., the organic ligand RЈ itself and oxidation by-products of the organic ligand. Indeed, the analysis of the light emitted by plasma using optical emission spectroscopy (not shown here) indicates clearly the CO, CO 2 , and H 2 O signatures in the plasma phase. Reaction (4) describes the silicate formation by decomposition of the precursor and reaction with Si, which is enhanced by the O 2 plasma. Silicate films can contain a slight carbon contamination. Reaction (2) and (4) occur at the substrate interface, whereas reaction (3) is the main deposition mechanism far from the interface. For a high reactant flow, i.e., at 5 Hz injection frequency, a high reactive species quantity is present on the substrate surface and reactions that consume precursor molecules are favored, i.e., reactions (3) and (4). The interface mainly results in silicate compounds, since re- action (4) is enhanced compared to reaction (2). Since the deposition rate increasing frequency as found previously by XPS (because the precursor quantity on the surface is higher), the interface reaction (4) is rapidly stopped because Si atoms diffusion is limited due to the rapid formation of the YOC layer and reaction (3) that leads to the formation of yttrium oxide carbonates becomes the only deposition mechanism. On the contrary, at 1 Hz, i.e., with a low reactant flux, the growth rate is lower (because the precusor quantity on the surface is lower) and reaction (4) will occur for a longer time than during the 5 Hz process because here the YOC layer growth is slower and the Si diffusion is less limited. In this case, reactions that consume silicon near the interface are more likely to occur, i.e., reaction (2), which leads to the formation of a thick SiO 2 layer, and reaction (4), which gives a silicate layer (on top of SiO 2 ). Beyond a critical thickness, the deposited layer at 1 Hz will be yttrium oxide (with carbon contamination) without silicon (as found for 40 nm thick layers deposited at 1 Hz), since reaction (4) is limited by the presence of Si atoms. Beyond a critical distance from the substrate, no silicon is available and only reaction (3) occurs. From this study, it appears that a high reactive species flux should be preferred to limit the formation of the interfacial SiO 2 layer.

D. Effect of substrate temperature

Five different substrate temperatures (in the range 200-380°C) were investigated for films grown at 1 Hz in the standard conditions given in Table I. A same precursor liquid volume was injected (1000 injections). The films were investigated by XPS (no Ar sputtering was performed prior to measurements since the films were very thin ͑Ͻ10 nm͒). Figure 10 shows the XPS spectra of the Si 2p, Y 3d, and O 1s lines (Si 2p spectra are normalized on Si 0 ). The XPS quantification from C 1s spectra (not shown here) shows that the C content increases with decreasing substrate temperature, as it is expected due to incomplete precursor decomposition. First, it appears that deposition can occur at temperature as low as 200°C. However, the sample grown at 200°C is significantly different with high carbon content with limited yttrium content (mainly in the form of Y-O-C and limited silicate). In the range 250-380°C, the Y 3d spectra are unchanged. The binding environment of Y is mainly Y-O-Si, as shown previously for the film grown at 350°C ͑158.8 eV͒. The O 1s and Si 2p regions indicate that the yttrium silicate thickness increases as compared to the SiO 2 one, when the deposition temperature increases (O 1s and Si 2p contributions around, respectively, ϳ531.5 and ϳ102.4 eV increase with increasing temperature). The quantitative analysis of the Si 2p region was done, as previously described (Table II), and the peak area of each Si n+ component (n =1-4), normalized to the Si 0 one, is plotted as a function of temperature in Fig. 11(a). The graph shows that the Si 4+ ratio stays roughly constant, which indicates that the SiO 2 thickness is not significantly modified in the studied temperature range. On the contrary, the Si 2+ and especially the Si 3+ components increase with increasing temperature. These two contributions arise from a wide range of silicon bonding environments available in Y-O-Si. Such an increase is related to the increase of the yttrium thickness with increasing temperature. In Fig. 11(b), the area ratio A Si-O /A Si component (A Si-O is the sum of all area silicon oxidized components: Si 1+ ,S i 2+ ,S i 3+ , and Si 4+ ) is represented in an Arrhenius-type plot, which shows the extent of the yttrium silicate thickness versus the inverse of the deposition temperature. The activation energy ͑E a ͒ for the silicate formation during deposition is equal to 12.5 kJ/ mol. This value is rather low, which indicates that thermal activation is not significant in this temperature range, as could be expected (precursor's decomposition requires higher temperatures).

E. Electrical characterizations

A typical C-V curve measured with an electrode surface of 0.694ϫ 10 -2 mm 2 is shown in the insert of Fig. 12 for a film grown at 5 Hz ͑ϳ40 nm͒. The EOT was extracted from the accumulation capacitance and is plotted in Fig. 12 for three film thicknesses (estimated by x-ray reflectometry). The intercept of the linear fit is equal to 1.7 nm ͑±1.3 nm͒. The effective permittivity of the films extracted from the slope of the linear fit is of 7.5͑±1.2͒. This depressed value compared to Y 2 O 3 is consistent with the presence of carbonate and silicate. Annealing of the thick films does not change significantly the EOT. For example, the EOT values extracted from the 1 kHz C-V measurements are 18.8, 17.2, 18.2 and, 16.9 nm for, respectively, as-deposited, 450°C H 2 annealed, Ar and O 2 700°C annealed samples. Although the 700°C thermal post-treatment significantly reduces the organic contamination and leads to the crystallization of the Y 2 O 3 phase, the EOT value stays constant probably due to the increase of the interfacial layer. The current density measured as a function of applied voltage is plotted in Fig. 13 for a 40 nm thick film (grown at 5 Hz) and annealed samples. Electrical breakdown for the as-deposited film occurs at ϳ34 V (which corresponds to an electrical breakdown field E bd of the order of 8 MV/ cm) and the film exhibits a low leakage current density (below 10 -7 A/cm 2 for a bias volt-age ranging from -20 to 0 V). The dielectric breakdown field values are in good agreement with those expected from the relationship proposed by McPherson al. that links E bd and the dielectric constant ͓E bd ϳ͑͒ -1/2 ͔. 55 No significant change is observed on the J(V) curve for the 450°C H 2 annealed film. The 700°C annealed films present a lower E bd and a higher leakage current density, which are attributed to the film crystallization. To conclude, the as-deposited and 450°C-annealed films exhibit promising electrical properties, especially for MIM devices, which requires a low thermal budget. 7

IV. CONCLUSION

Yttrium oxide and yttrium thin films were grown by pulsed-liquid-injection PE-MOCVD from Y͑tmhd͒ 3 precursor. The plasma allows a significant reduction of the thermal budget as compared to conventional thermal MOCVD and yttrium silicate films can be grown at temperatures as low as 250°C. Depending on the incoming reactive species flow rate (which is tuned by the injection frequency), a SiO 2 or an yttrium silicate interface layer is favored. Beyond a critical thickness, yttrium oxide is formed but the low deposition temperature implies carbon contamination, which can, however, be reduced upon annealing at 700°C. The SiO 2 interfacial layer can be limited by increasing the injection frequency (i.e., at 5 Hz). Finally, the as-deposited films exhibit a high electrical breakdown field ͑E bd =8 MV/cm͒ and a low leakage current density ͑I Ͻ 10 -7 A/cm 2 ͒, which is a very promising result for MIM applications.

The International Technology Roadmap for Semiconductors, Semiconductors Industry Association, see also http://public.itrs.net/ for the most recent updates.

FIG. 1 .

 1 FIG. 1. Schematic of the PE-MOCVD experimental setup used to deposit yttrium oxide with the pulsed liquid-injection precursor source connected to the plasma chamber.

Figure 3

 3 Figure3presents the Y 3d, O 1s, and C 1s XPS spectra for a thick ͑ϳ40 nm͒ film deposited in the standard condition (TableI) using an injection frequency of 5 Hz. The corresponding Y 3d and O 1s spectra for an yttrium oxide reference ceramic are also shown, after surface sputtering. In

FIG. 2 .

 2 FIG. 2. XPS survey scans for (a) process with only plasma enhanced at 70 W (as PECVD), (b) process with only substrate heating at 350°C (as MOCVD),a n d(c) process with both plasma and substrate heating at 70 W and 350°C, respectively ͑PE-MOCVD͒. This graph shows the synergy between the plasma and the substrate heating to deposit Y 2 O 3 layers using the PE-MOCVD technique.

  ) (165.0 eV for the main component, the smaller one probably correspond to additional yttrium bindings due to air exposure). If a constant splitting of 2.0 eV is assumed for the Y 3d doublet, the fitting of the Y 3d spectrum requires at least two components for each peak, which indicates a multiple yttrium bonding environment (C and H presumably). Consistent with this, the O 1s peak at 530.6 eV can be attributed to Y-O-C bonds with H contamination [Fig. 3(b)].

  FIG.4. The graph shows FTIR spectra for the thick sample and three annealing post-treatments. The as-deposited film is denoted by curve with squares (), the H 2 annealed at 450°C for 30 min by the curves with circles ͑b͒, the O 2 annealed at 700°C for 30 min by the curves with up triangles ͑᭡͒, and finally the Ar annealed at 700°C for 30 min by curves with down triangles (᭢).

FIG. 5 .

 5 FIG. 5. XRD diagrams for (1) as-deposited sample, (2) H 2 annealed at 450°C, (3) Ar annealed at 700°C, and (4) O 2 annealed at 700°C.

  FIG. 7. XPS spectra [(a) Si 2p, (b) Y 3d, and (c) O 1s regions] for samples with increasing injection frequency [1H z (b),2H z(•),5H z͑᭡͒,a n d 10 Hz (᭢)]. Si 2p spectra are normalized on Si 0 peak area (peak from substrate) in order to estimate the oxidized Si-based compound thickness, since the well-known method for measuring the thickness of oxide film by XPS is based on the area ratio between the oxidized silicon A Si-O and substrate silicon A Si . The spectrum of initial substrate ͓Si/ SiO 2 ͑0.8 nm͔͒ is also done for comparison.

FIG. 8 .

 8 FIG. 8. Angle-resolved XPS analyses (Si 2p, Y 3d, and O 1s regions) at three takeoff angles [90°(),30°͑b͒ and 15°͑᭡͒] of the 5 Hz sample [(a), (b),and(c)] and the 1 Hz sample [(d), (e), and (f)]. Note that the intensity of Si 2p spectra is normalized on the Si-O peak. The relative intensity of each Si n+ contribution for the different takeoff angle analyses extracted from the Si 2p fit model is plot in the insert of Si 2p regions [(a) and (d) corresponding to the 5 and 1 Hz samples, respectively].

FIG. 9 .

 9 FIG. 9. Diagram illustrating the difference between the 5 and 1 Hz structure consistent with AR-XPS.

FIG. 10 .

 10 FIG. 10. XPS spectra [(a) Si 2p, (b) Y 3d, and (c) O 1s regions] for samples with increasing deposition temperature [380°C (), 350°C ͑b͒,3 0 0 °C ͑᭡͒, 250°C (᭢), and 200°C (छ)]. As Fig. 7, Si 2p spectra are normalized on Si 0 peak area (peak from substrate) in order to estimate the oxidized Si-based compound thickness evolution. The spectrum of initial substrate ͓Si/ SiO 2 ͑0.8 nm͔͒ is also done for comparison.

  2 P. C. Joshi and N. W. Cole, J. Appl. Phys. 86,8 7 1(1999).

  3 S. Blonkowski, M. Regache, and A. Halimaoui, J. Appl. Phys. 90, 1501 (2001).

4T.

  FIG. 11. Graph (a). Area ratios determined from the Si 2p fitting model vs the deposition temperature. The graph gives the area ratio Si-O / Si and also area ratio of each oxidation state. Graph (b). Arrhenius diagram of Si-O / Si area ratio vs 1000/T. The activation energy is extracted from the slope.

FIG. 13 .

 13 FIG. 13. Leakage current density vs voltage for thick Y 2 O 3 film ͑ϳ35 nm͒: as-deposited film (), 450°C annealed films under H 2 ͑b͒, and 700°C annealed films under Ar ͑᭡͒ and O 2 (᭢).

TABLE I .

 I The standard process conditions of the plasma chamber and of the pulsed-liquid precursor source, respectively.

	Standard plasma	Pressure	Gas	Input plasma	Wafer
	chamber parameters			power	temperature
		0.27 kPa	O 2	70 W	350°C
		͑2 Torr͒ 250 SCCM
	Standard injection	Dilution	Injection	Frequency Opening time
	parameters		number
		0.02 mol/ l	1-5000	1 -10 Hz	2 ms

TABLE II .

 II Binding energy, FWHM (full width at half maximum) and peak shape (Lorentzian-Gaussian percentage) for each Si oxidation state to fit the Si 2p XPS region.

				Lorentzian (L) Gaussian (G)
	Oxidation	Binding		percentage (%)
	state	energy (eV)	FWHM	(L:100 and G:0)
	Si 0	99.3	0.9	25
	Si 1+	100.3	1.0	0
	Si 2+	101.1	1.2	0
	Si 3+	102.1Ϫ102.3 a	1.3Ϫ1.5 a	0
	Si 4+	103.3	1.4	0

a Variable values, which depend on fitting results.

ACKNOWLEDGMENTS

The authors would like to thank warmly N. Rochat for helpful expertise on IR-analysis (FTIR), D. Muyard, G. Rolland, and M. Burdin for XRD analyses. F. Ducroquet is also thanked for helpful discussion on electrical characterizations. L. Vallier is also gratefully thanked for XPS expertise and deposition equipment support.


William Rapin

Gilles Dromart

Juergen Schieber

Benton C Clark

Linda C Kah

David M Rubin

Sanjeev Gupta

Amelie Roberts

Gwénaël Caravaca

Not always wet: an aridification sequence in the orbital clay-sulfate transition of aeolis mons

des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

NOT ALWAYS WET: AN ARIDIFICATION SEQUENCE IN THE ORBITAL CLAY-SULFATE

TRANSITION OF AEOLIS MONS. W. Rapin 1 , G. Dromart 2 , J. Schieber 3 , B.C. Clark 4 , L. Kah 5 , D. Rubin 6 , S. Gupta 7 , A. Roberts 7 , G. Caravaca 1 , L. Edgar 8 , R. Y. Sheppard 9 , E. Dehouck 2 , S. Le Mouélic 10 , A. Bryk 11 , W. E. Dietrich 11 , P. Gasda 12 , O. Gasnault 1 , N. Lanza 12 . 1 IRAP/UPS/CNRS, Toulouse, France (william.rapin@irap.omp.eu), 2 Univ. Lyon, LGL-TPE, 3 DGS, Indiana Univ., 4 Space Science Institute, CO, 5 University of Tennessee, 6 UC Santa Cruz, 7 Imperial College London, 8 USGS, 9 PSI 10 LPG, Nantes, 11 UC Berkeley, 12 LANL.

Introduction:

On Mars, prominent thick-layered sulfate-bearing deposits are observed at a number of late Noachian to late Hesperian locations (~3.5 Ga) [1]. Their apparent absence in older strata has led to the hypothesis that they may represent ancient evaporites related to the global aridification [2]. In Gale crater, the Curiosity rover is exploring the sulfate-bearing unit, a regional package hundreds of meters thick of yet mostly undefined origin [3]. Curiosity's investigations can test whether the formation processes of this unit are similar to a form of primary evaporation, such as early diagenetic salt crystallization, or later stage precipitations with different conditions.

Entering the transition: After exploring the mudstone-rich and clay-bearing strata of the Murray and Carolyn Shoemaker fm. for about seven years, the Curiosity rover has explored the orbital clay-sulfate transition starting in 2021 (Fig. 1). Bedrock in the basal section is marked by diverse diagenetic overprints where sedimentary structures are less visible. Further up, the rover imaged butte-forming outcrops and revealed a >100-m-thick interval with a transition into large-scale trough cross-bedded structures (Fig 2a) [4].

Geochemistry and distribution of sulfates: The base of the clay-sulfate transition succession is marked by a sharp increase in the occurrence of bedrock with nodular texture associated with sulfate enrichments (Fig. 1). The rover analyzed the composition of three bedrock lithologic components: (i) a smooth host bedrock, which was drilled at multiple locations and revealed the disappearance of clay minerals in X-ray diffraction (XRD) [START_REF] Rampe | pdf 54th Lunar and Planetary Science Conference 2023[END_REF] -the evolved gas analyses (EGA) also suggest minor Mg-sulfates and an isotopic change in sulfur compared to the clay-bearing unit below [START_REF] Franz | [END_REF]; (ii) a bedrock with mm-scale nodular texture which was observed in the upper section of the clay-sulfate transition, and became abundant within the marker band valley at and around the Canaima drill locality (Fig. 2d) -revealing significant Mg-sulfate content according to both XRD and EGA [START_REF] Rampe | pdf 54th Lunar and Planetary Science Conference 2023[END_REF][START_REF] Franz | [END_REF]; (iii) a cm-scale nodular bedrock, which was not drilled due to its uneven surface, but shows diverse sulfate-enriched compositions as measured by ChemCam (Fig. 1). Casulfates are present mostly in smooth host bedrock in variable amounts across the stratigraphy, but also sporadically as submillimeter heterogeneity within nodules. Overall, the bulk bedrock (combining smooth and nodular) in the transition shows an increase in sulfate content relative to prior strata, nodular textures being a key component of that change (Fig. 1).

Unique nodular lithologies: Unusual cm-scale and mm-scale textures have been observed within the claysulfate transition. At the base a regular pattern of centimetric polygonal ridges (Fig. 2) is clearly crosscut by later stage fractures. These sulfate-enriched polygonal ridges may represent the first evidence of a paleosol on Mars formed by sustained wet-dry cycles at the surface [START_REF] Goehring | [END_REF]. Near the top of the section, in the marker band valley, bedrock has the highest abundance of sulfates and density of nodules. At and near the Canaima drill sample hydrated Mg-sulfates are a major component of the bulk bedrock with mm-scale nodular texture, ChemCam hydrogen signal indicates ~21 wt.% associated water (Fig. 1). This enhanced hydration is also observed in the cm-scale nodular bedrock leading to the marker band surface (Fig. 2e).

Discussion and interpretations: The terrains orbitally defined as part of the clay-sulfate transition and leading upward to the sulfate-bearing unit show multiple signs consistent with marked changes towards drier paleoenvironments both in the sedimentary and geochemical record. Within the large-scale crossbedded strata of most likely eolian origin, sulfates occur in the form of diverse lithologies mostly related to nodular bedrock. The abundance of nodules and related sulfates is highly variable within the stratigraphy.

Within the transition, features within specific intervals have already been identified as indicators of locally and intermittently wet environments, such as polygonal ridges in the basal section or lenses with wave ripple laminae in the upper section, and within the marker band itself [8]. The surrounding dry eolian strata adjacent to these features are also marked by the enhanced occurrence of nodular lithologies (Fig. 1). This association and the general stratigraphic control on the distribution of sulfate-enriched nodules is consistent with their formation within the sediments due to evaporation where a sulfate-rich water table reached or came close to the surface in a variably dry climate.

Figure 1 :

 1 Figure 1:The clay-sulfate transition (red interval) as recorded by the ChemCam instrument on bedrock with smooth and nodular textures. Left: General stratigraphy (credit: MSL sed-strat team) against sulfur signal[9], with the marker band (red arrow) and identified features marking intermittently wet paleoenvironments (shaded blue). Dark red curve represents bedrock average (smooth host and nodular). Right: Sulfur and hydrogen as function of MgO and CaO content with hydrated Mg/Ca-sulfates references. MgO and CaO content in the transition are preliminary compositions derived from MOC[10] and renormalized to accommodate estimated SO3 and H2O content based on S and H signal[9,11].

Figure 2 :

 2 Figure 2: Mirador butte with large-scale crossbedding and diverse bounding surfaces (dashed red) consistent with eolian deposition and example of contact with interdune lens deposits (dashed yellow) (a). Remarkable nodular bedrock (b-e) also highlights the diversity of textures observed with distinct sulfate-enriched composition. Regular pattern of polygonal ridges (b); concretions (c); drill sample at Canaima (arrow) with mmnodular texture (d); bedrock with a saturated density of hydrated Mg-sulfate-rich cm-scale nodules below the marker band (e).

  Planetary Science Conference 2023 (LPI Contrib. No. 2806)


J D Béguin

D Adrian

J A Petit

J P Rivière

C Vahlas

S Vaillant

Improvement of salt corrosion resistance of titanium alloys by PVD and CVD coatings

The corrosion resistance of two kind of coating systems was investigated to protect a Ti-6242 titanium under hot salt stress corrosion conditions: PVD silicon coatings and CVD alumina coating. All these coatings were tested under hot Salt Stress Corrosion (SSC) with a constant static mechanical stress and cyclic temperature conditions, and have been brought near to the real industrial work conditions. Although all coated specimens were damaged in salted area, both SiC and Al 2 O 3 coatings showed no SSC rupture during 100 h, while Si x N y lifetime under SSC conditions was increased by a factor 2 compared to noncoated one. As well as these interesting quantitative results, coated specimens exhibit a good adherence on titanium substrate. 1

INTRODUCTION

The use of the Ti-6242 titanium alloy was largely developed in compressor section in aero engines, because of its good specific mechanical resistance, high toughness and structural stability. However, under severe conditions, the presence of salt combined with strong mechanical stress at high temperature, can lead to an untimely failure: It is the well-known hot salt stress corrosion phenomena. Consequently, it is necessary to improve the corrosion resistance by a protective coating, with respect to initial mechanical performance of the substrate. In this study, SiC, Si x N y , Al 2 O 3 coatings were investigated to protect a titanium alloy substrate under SSC conditions. At first unloaded coated samples were corroded with salt deposit without mechanical loading at 540°C during 100 h to characterize salted damaged areas and corrosion products by Energy Dispersive Spectroscope (EDS) and X-Ray Diffraction (XRD). Afterwards salted samples were evaluated using a creep test bench. High temperature cycling was performed in a mobile furnace under tensile loading. The salt deposit was formed by dipping calibrated droplets of saline solution onto the coated sample surface.

MATERIALS AND EXPERIMENTAL PROCEDURE

Experiments are performed on flat samples (10mm diameter) for salt exposure without mechanical loading, and cylindrical samples (4.5mm diameter, 45 mm length), for salt stress corrosion test bench. The tested material is Ti-6242 with duplex microstructure (α coarse grain, in a β transformed matrix). Flat surface samples were ground using a 600 grit abrasive SiC paper before surface treatment. The salt deposit is obtained by depositing a calibrated droplet of saline solution on sample surface. Saline solution was prepared with natural sea salt. The salt crystallisation was achieved in an oven for 5 minutes at 110°C. This procedure allows to obtain repeatable deposit. Flat samples have been placed in furnace with air circulation. Cylindrical samples were evaluated using a creep test bench.Thermal cyclic conditions 520°C/24h and free air cooling 150°C/1h, were applied four times (100 h total duration), under 450 MPa tensile stress.

Silicon coatings: SiC and Si x N y

These coatings were deposited by a PVD technique with an assisted ion beam bombardment (120 keV energy Ar + ions ) under 5.10 -3 Pa, for the first 40 nm, to create a chemical gradient composition by an ionic mixing between the coatings and the substrate, in order to improve their adherence (Rivière 2005). Theses dense coatings were very dense and of amorphous structure; their thicknesses are respectively 0.4 µm for Si x N y and 0.33 µm for SiC.

Alumina coating

This coating was processed by metal organic chemical vapor deposition (MOCVD) at atmospheric pressure (Valhas 2006). The corresponding process involves two reactors operating in series. The first one ensures high, reliable, and reproducible flux of vapors of aluminium acetylacetate (Al(acac) 3 ) precursor through sublimation in fluidized bed. The fluidized bed sublimator is connected at its exit to a vertical, resistively heated CVD reactor. Samples were maintained welded on a regulation thermocouple, which is positioned close to the sample along the central axis of the reactor. Preliminary experiments were performed in sublimation conditions 393K, with nitrogen as fluidizing gas. More details on the setup are provided. Sublimation conditions ensure feeding of the CVD reactor with 0.19 sccm [Al(acac) 3 ]. An additional gas line was used to feed the CVD reactor with 100 sccm of oxygen bubbling through a H 2 O bath maintained at 100°C. Deposition temperature was fixed at 580°C in order to maintain the microstructural integrity of the Ti6242 alloy. 300 nm thick amorphous alumina films were deposited in 3 h runs. Films are smooth and do not contain any heteroatoms, namely carbon.

CORROSION RESISTANCE OF COATED SPECIMENS

Salt attack without tensile stress

Salt coated specimens (10.5 mg/cm 2 ) were heated at 540°C for in a tubular furnace during 100h. In order to confirm the degradation mechanism, salted samples were observed means optical microscopy after distilled water immersion to eliminate residual salt. The observation of corroded area for SiC and Si x N y coatings showed that pitting corrosion is the principal mode of degradation, while an accelerated oxidation occured at and near the salt deposit for the SiC coating. Corrosion degradation of Al 2 O 3 coating was not significant (Fig. 1).

Salt attack under tensile stress (SSC conditions)

Pre-oxidation of SiC and Si x N y coated specimens were conducted in air at 600°C under dry air for 100h, to increase oxide barrier thickness, in order to improve salt corrosion resistance. Salted samples were evaluated using a creep test bench under 450 MPa tensile stress at 520°C during 100 h. Salt deposit was formed by dipping calibrated droplets of saline solution onto samples (120 µg or 2,1 mg/cm 2 approximatively). Initially, non-coated Ti-6242 specimen was evaluated in the same condition for reference (Fig. 2).

Under these conditions, lifetime of non-coated specimen was about 39 h. All coated specimens were damaged in salted area (Fig. 3), but the results showed no SSC failure for both SiC and Al 2 O 3 coatings during 100 h and an increase of lifetime for Si x N y coating (96 h).

After SEM examination of Al 2 O 3 cracking, it was clear that the main damage origin was thermomechanical and not SSC phenomena (Fig. 4). It was due to a local imperfection like machining ridges on Ti-6242 sample surface, and/or low thickness of Al 2 O 3 coating, which increases local stress under these severe conditions. S.E.M observations of corroded areas, for both silicon coatings exhibited a non-uniform degradation with pits and cracks and corrosion products. X-ray maps (Fig. 5) showed that Na, O and Si were the major chemical elements.

To complete these analyses, corrosion products were examined using an X-ray diffractometer under a low incidence angle (3° and 5°) with a pipe collimator to focus the incident beam (800 µm) on corroded area. In spite of a low critallinity of corrosion products, the main detectable phases were sodium silicates and silicon oxides, with remaining NaCl phases (Fig. 6).

In the case of the Al 2 O 3 coated sample, S.E.M observations exhibited no visible pits with a small amount of corrosion products (Fig. 7). Combining XRD and EDS analyse, corrosion products contained sodium aluminate (NaAlO 2 ) and spinel phase (MgAl 2 O 3 type). The main difference compared to silicon coatings was the chemical reaction with the Mg element present in sea salt. This element seems to play an important role in the corrosion process. In this case again, no rich Cl products were detected by XRD and EDS analyse.

× × × × Na 2 SiO 3 × × × × * • • • • Na 2 O • • • • ♦ ♦ ♦ ♦Ti ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ × × × × ◊ ◊ ◊ ◊ NaCl ◊ ◊ ◊ ◊ SiC Coating * * * * * Na 4 SiO 4 × × × × Na 2 SiO 3 × × × × * • • • • ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦Ti ♦ ♦ ♦ ♦ × × × × ◊ ◊ ◊ ◊ NaCl ◊ ◊ ◊ ◊ Si x N y Coating DISCUSSION

Silicon coatings

Similar corrosion phenomena were observed for both silicon coatings, when the samples were coated with NaCl after exposure to dry air at high temperature. This leads to a sodium silicate formation (Pickrell 1994). The possible reactions to produce sodium silicate are the following: Discontinous sodium silicates layer was contributed to the formation of some cracks and breackdown of SiO 2 (Guangming 2005) in the corroded area, because of volume change (Fig. 8). Cracks could serve as short path to lead to a serious hot corrosion of SiC or Si x N y coating (Si 3 N 4 for simplification) and the following reactions can occur at 520°C: Some of the Cl 2 then reacts with silicon coatings to form silicon cloride:

♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ ♦ ◊ ◊ ◊ ◊ Al 2 O 3 Coating ◊ ◊ ◊ ◊ NaCl ♦ ♦ ♦ ♦Ti ♦ ♦ ♦ ♦ × × × × NaAlO 2 × × × × × × × × × × × × × × × × × × × × × × × × ♦ ♦ ♦ ♦ × × × × × × × × × × × × ◊ ◊ ◊ ◊ ◊ ◊ ◊ ◊ ∆ ∆ ∆ ∆G<o ∆ ∆ ∆ ∆G<o ∆ ∆ ∆ ∆G<o ∆ ∆ ∆ ∆G<o <Na 2 SiO 3 > 2 <Na> + 1/2(O

Figure 1 :

 1 Figure 1: Surface morphology of water washed surface of SiC (a), SixNy (b) and Al 2 O 3 (c) after salt attack at 540°C/100h.

Figure 2 :Figure3:Figure 4 :

 24 Figure 2: Calibrated salt deposit on creep specimen before test (a), non-coated Ti6242 specimen after test (b).

Figure 6 :

 6 Figure 6: X-ray diffraction spectra under low incident angle (3°) on correded silicon coatings.

Figure 5 :

 5 Figure 5: X-ray map on corroded silicon coatings.

Figure 7 :Figure 8 :

 78 Figure 7: X-ray map and XRD analysis (low incident angle 5°) after salt attack of Al 2 O 3 . coating

2 ) <Na 4 SiO 4 > <Na 4 SiO 4 > <Na 2 O> + <SiO 2 > <Na 2 SiO 3 > + <Na 2 O> <Na 2 O> 2<Na 2 O> + <SiO 2 >

  SiC and Si x N y and CVD-Al 2 O 3 was evaluated under severe conditions. Lifetime of all coated specimens under SSC conditions were significantly improved by a factor 2.5 compared to uncoated specimen. Silicon coatings were more damaged than Al 2 O 3 coating. The formation of sodium silicate seems to be the key factor in the overall corrosion reactions. The Al 2 O 3 /Ti6242 system showed the best salt attack resistance. It only appeared as small clusters without pits presence. Surface defects of Ti6242 substrate seem to be the principal cause, by initiating corrosion phenomena after a mechanical damage of Al 2 O 3 layer, under tensile sollicitations. Surface and coatings technology; (article in press). Valhas, C., Caussat, B., Senocq, F., Gladfelter, W.L., Aloui, L.,Moersch, T., (2006) An original device for the delivery of vapors to a deposition apparatus based on the sublimation in a fluidized bed. Chem.Vap. Deposition; submitted. Yao, Z., Maeek, M.,(1995) NaCl induced hot corrosion of a titanium aluminide alloy. MaterialScience and Engineering; A192/193, 994-1000. Zheng , C.,Wu, N.Q., Mao, S.X., (2003) Effect of Al 2 O 3 averlay on hot-corrosion behaviour of yttriastabilized zirconia coating in molten sulphate-vanadate salt. Thin solidfilm, 443, 46-52.

	5.2 Alumina coating
	Chemical analysis combinated with thermodynamic calculation showed that sea salt can
	react with Al 2 O 3 coating (Zheng 2003) to form sodium aluminate (Yao 1995) and spinel
	phase which may occur at 520°C according to following reactions:
		2 <NaCl> + <Al 2 O 3 > + 1/2(O 2 )	2<NaAlO 2 > + (Cl 2 ) ∆ ∆ ∆ ∆G<o	(10)
		<Mg> + 1/2(O 2 )	2<MgO>		∆ ∆ ∆ ∆G<o	(11)
		<MgO> + <Al 2 O 3 >	<MgAl 2 O 4 >		∆ ∆ ∆ ∆G<o	(12)
	6	CONCLUSION
	Corrosion resistance of PVD-7 BIBLIOGRAPHIES
			(1)
			(2)
			(3)
			(4)
	Keywords:
	6242 alloy; corrosion protection; PVD silicon coatings; Interfacial ion dynamic mixing; MOCVD alumina
	coating; organo-metallic precursor; Salt deposit; thermochemical degradation; Hot salt stress
	corrosion; oxychloridation; chloridation, oxididation.
		2 <NaCl> + 2<SiC> + 3 (H 2 O) + (O 2 )	<SiO 2 > + <Na 2 SiO 3 > + (Cl 2 ) +3(H 2 ) + 2C
					∆ ∆ ∆ ∆G<o	(5)
		4 <NaCl> + 2<SiC> + 2 (H 2 O) + 2(O 2 )	<SiO 2 > + <Na 4 SiO 4 > + 2(Cl 2 ) +4(H 2 ) +2 C
					∆ ∆ ∆ ∆G<o	(6)
	2 <NaCl> + <Si 3 N 4 > + 2 (H 2 O) + 5/2(O 2 )	2<SiO 2 > + <Na 2 SiO 3 > + (Cl 2 ) +2(H 2 ) +2(N 2 )
					∆ ∆ ∆ ∆G<o	(7)
		<SiC> + 2 (Cl 2 )	(Si Cl 4 ) + <C>	∆ ∆ ∆ ∆G<o	Volatil products (8)
		<Si 3 N 4 >+ 6 (Cl 2 )	3 (Si Cl 4 ) + 2 (N 2 )	∆ ∆ ∆ ∆G<o	(9)

Guangming, L.,

Meishnan, L., Yanchun, Z., (2005)

Influence of the hot corrosion of Ti 3 SiC in the mixture of Na 2 SO 4 -NaCl melts. Corrosion science;(artcle in press). Pickrell, G.R., Sun, T., Brown, J.J.,(1995) High alcali corrosion of SiC and Si 3 Ni 4 . Fuel processing technology; 44, 213-236. Rivière, J.P., Pichon, L., Drouet, M.,

Galdikas, A., Poquillon, D., (2005)

Silicon based oxidation resistant coating on Ti6242 alloy corrosion by dynamic ion mixing.