Thomas Gheno

Brian Gleeson

On the Reaction Mechanism of MCrAlY Alloys with Oxide-Sulfate Deposits at 1100 • C

Keywords: MCrAlY coatings, Al 2 O 3 scale, Selective oxidation, Silicate deposits

The corrosion of γ-(Ni,Co) + β-(Ni,Co)Al NiCoCrAlY alloys exposed to CaO-rich, Na 2 SO 4 -containing deposits in air and CO 2 -H 2 O-O 2 at 1100 • C was studied with a focus on the mechanisms governing the accelerated attack of γ-rich compositions. Internal oxidation of Al resulted from the breakdown of an initially formed Al 2 O 3 scale, due to its reaction with the deposit. Compared to deposit-free conditions, this scale exhibited: (i) a delayed Al 2 O 3 structure transition from metastable θ to stable α; (ii) finer grains; (iii) a reduced adherence to the metal, to the extent that a gap developed at the metal/oxide interface; and (iv) significant permeability to S, as well as N during air exposures. The occurrence of sulfidation in S-free gases reflected the establishment of an elevated S activity under the molten Na 2 SO 4 , due to the locally reduced p O2 . The factors affecting the Al 2 O 3 -scale microstructure were examined, as well as the associated consequences on Al consumption and on the transition to internal oxidation. The roles of alloy and gas compositions in the identified mechanisms were also discussed.

Introduction

High-temperature structural materials are primarily designed for optimal mechanical properties, and often require coatings to withstand corrosion in the chemically harsh environments they are exposed to. In particular, MCrAlY coatings (M = Ni, Co, or both), based on the fcc-A1 solid solution γ-(Ni,Co) and the B2 compound β-(Ni,Co)Al, are used to protect Ni-base materials in aero, marine and land-based power generation turbines [START_REF] Reed | The Superalloys: Fundamentals and Applications[END_REF][START_REF] Bose | High Temperature Coatings[END_REF]. Environmental factors affecting their corrosion resistance include the presence of oxidants in gaseous form (e.g., O 2 , CO 2 , H 2 O, and SO 2 ) or as condensed phases, such as silicates or sulfates.

The oxidation behavior of MCrAlY coatings in deposit-free conditions, and specifically the mechanisms governing the establishment of a protective Al 2 O 3 scale, are relatively well documented (e.g., Refs. [START_REF] Clemens | TEM and SNMS studies of protective alumina scales on NiCrAlYalloys[END_REF][START_REF] Mennicke | Transient phase evolution during oxidation of a two-phase NiCoCrAlY bond coat[END_REF][START_REF] Levi | Alumina grown during deposition of thermal barrier coatings on NiCrAlY[END_REF][START_REF] Nijdam | On the microstructure of the initial oxide grown by controlled annealing and oxidation on a NiCoCrAlY bond coating[END_REF][START_REF] Sullivan | Vapor-phase-mediated phenomena associated with high temperature, high water content oxidation of MCrAlX bond coats[END_REF]). A considerable body of work has been devoted to the effects of H 2 O on the oxidation of Al 2 O 3 -forming alloys (see reviews in Refs. [START_REF] Birks | Introduction to the High Temperature Oxidation of Metals[END_REF][START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF][START_REF] Saunders | The oxidation behaviour of metals and alloys at high temperatures in atmospheres containing water vapour: a review[END_REF]). Although some aspects are system-dependent, it is generally found that H 2 O accelerates Al 2 O 3 growth and enhances scale cracking and spallation. Much less information is available regarding the role of CO 2 or SO 2 in these systems. Recent investigations on Fe-base [START_REF] Dryepondt | Effect of H 2 O and CO 2 on the oxidation behavior and durability at high temperature of ODS-FeCrAl[END_REF] and Ni-base [START_REF] Lance | The effect of cycle frequency, H 2 O and CO 2 on TBC lifetime with NiCoCrAlYHfSi bond coatings[END_REF] Al 2 O 3 -forming alloys in CO 2 -H 2 O found no significant effect of CO 2 . This is in contrast with Cr 2 O 3formers, which are subject to carburization in CO 2 , indicating that Al 2 O 3 is a much better barrier to carbon than Cr 2 O 3 is [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF].

Research on deposit-induced degradation of Al 2 O 3 -forming systems has mostly concerned hot corrosion, a form of accelerated corrosion by molten sulfates which can be particularly severe at 700-900 • C [START_REF] Goebel | Mechanisms for hot corrosion of nickel-base alloys[END_REF][START_REF] Stringer | Hot corrosion of high-temperature alloys[END_REF][START_REF] Rapp | Chemistry and electrochemistry of the hot corrosion of metals[END_REF][START_REF] Goward | Low-temperature hot corrosion in gas-turbines: a review of causes and coatings therefor[END_REF][START_REF] Birks | Chapter 8 in Introduction to the High Temperature Oxidation of Metals[END_REF][START_REF] Pettit | Hot corrosion of metals and alloys[END_REF]. Fly ash deposits relevant to gas turbines mainly contain some combination of SiO 2 , Al 2 O 3 , CaO, MgO, and Fe 2 O 3 [START_REF] Levi | Environmental degradation of thermal-barrier coatings by molten deposits[END_REF], with additions of Na, K, and S [START_REF] Bons | Highpressure turbine deposition in land-based gas turbines from various synfuels[END_REF]. The reaction of these constituents with bond coat materials has received little attention because the oxide melting points far exceed the operating temperature of any metal component. Yet liquid phases may form at relatively low temperatures upon reaction of SiO 2 with the alkali metal oxides (the lowest eutectic points in the Na 2 O-SiO 2 and K 2 O-SiO 2 systems are at 794 and 743 • C, respectively [START_REF] Wu | Optimization of the thermodynamic properties and phase diagrams of the Na 2 O-SiO 2 and K 2 O-SiO 2 systems[END_REF]), or of CaO with thermally grown oxides such as Cr 2 O 3 . The lowest eutectic point in the CaO-Cr 2 O 3 system is not known with precision (conflicting values of 1022 • C [START_REF] Levin | Phase diagrams for ceramists[END_REF] and 1061 • C [START_REF] Kaiser | The system CaO-CaCr 2 O 4 -CaAl 2 O 4 in air and under mildly reducing conditions[END_REF] are reported), but catastrophic corrosion of Ni-Cr alloys due to liquid Ca chromate formation has been observed down to 1000 • C [START_REF] Chiang | The effects of deposits of CaO, CaSO 4 , and MgO on the oxidation of several Cr 2 O 3 -forming and Al 2 O 3forming alloys[END_REF].

In recent papers, we studied the reactivity of cast γ-β NiCoCrAlY alloys with CaO [START_REF] Gheno | High temperature reaction of MCrAlY coating compositions with CaO deposits[END_REF] and various oxide-sulfate deposits in S-free gases at 1100 • C [START_REF] Gheno | Modes of deposit-induced accelerated attack of MCrAlY systems at 1100 • C[END_REF][START_REF] Gheno | Kinetics of Al 2 O 3 -scale growth by oxidation and dissolution in molten silicate[END_REF]. Alloys with a high Cr concentration and high γ fraction, which have optimal hot corrosion resistance, were found to be particularly susceptible to accelerated corrosion when exposed to CaO-rich, Na 2 SO 4 -containing deposits. In addition to liquid Ca chromate formation, two modes of degradation were identified, which involved Al 2 O 3 dissolution in molten Na silicate and solid-state Al 2 O 3 reaction with CaO to form Ca aluminate. Both led to enhanced Al consumption, destabilizing the external Al 2 O 3 and promoting internal oxidation. In the present paper, we study the reaction mechanism of NiCoCrAlY alloys exposed to CaO-rich, Na 2 SO 4 -containing deposits at 1100 • C, with particular attention to the onset of non-selective oxidation. The effects of the deposit constituents on Al 2 O 3 formation are assessed based on microstructural characterization of scales grown with and without deposit.

Materials and experiments

Cast NiCoCrAlY alloys of nominal compositions given in Table 1 were used in this study. These included a β-rich composition with relatively high Al and low Cr, and six γ-rich alloys with 30 Co (all alloy compositions in at. %) and variations of Al and Cr (12/16 Al, 27/30/33 Cr). All alloys had 0.1 Y. Ingots were made by argon-arc melting, followed by drop casting into 10-mm diameter rods which were then homogenized in vacuum for 6 h at 1200 • C plus another 48 h at 1150 • C. All processing was conducted at the Materials Preparation Center of the Ames Laboratory [28]. Approximately 1-mm-thick specimens were cut from the heattreated rods, ground using SiC paper to a P1200-grit finish, then degreased with detergent and ultrasonically cleaned in ethanol before exposure.

Microstructures were characterized after a further 50-h annealing at 1100 • C followed by air cooling. All alloys had a primary γ-β microstructure; volume fractions of β measured by image analysis using the ImageJ software [START_REF] Schneider | NIH image to ImageJ: 25 years of image analysis[END_REF] are included in Ta- (vol. %), using a horizontal tube furnace. Gas flow rates were set to 100 ml/min, with a linear velocity of 2 mm/s at reaction temperature, and a total pressure slightly over 1 atm. To obtain a wet gas, the dry air, CO 2 or CO 2 -O 2 was bubbled through demineralized water set at 62 • C, thereby producing a mixture with slightly more than the desired 0.2 atm H 2 O. The excess water vapor was subsequently condensed by cooling the wet gas in a distillation column at 60 • C. Prior to a given experiment, the furnace was stabilized at 1100 • C with the specimens positioned in a cold zone of the reaction tube, and the gas was allowed to flow for 2 h. To start the experiment, the specimens were magnetically pushed in the hot zone.

The deposit compositions studied were SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 and SiO 2 -20Al 2 O 3 -33CaO-8MgO-6Fe 2 O 3 -5Na 2 SO 4 -1K 2 SO 4 (wt. %). These were prepared by grinding laboratory-grade reagents using a mortar and pestle, and were given no heat treatment prior to corrosion testing. The powders were mixed with ethanol, and the resulting slurry was applied to one face of a given specimen using a dropper. The procedure was adjusted to produce approximately uniform 30 ± 5 mg/cm 2 deposits after evaporation of the ethanol. The experiments were all isothermal, with durations ranging from 5 min to 250 h, as indicated in the text. The same amount of deposit, 30 mg/cm 2 , was applied once prior to each experiment, regardless of test duration.

After exposure, reaction products were examined by XRD, using a PANalytical Empyrean instrument with a Co radiation (K α1 = 1.789 Å), and by photostimulated luminescence spectroscopy (PSLS), which allows γ, θ, and α polymorphs of Al 2 O 3 to be distinguished [START_REF] Tolpygo | Microstructural study of the theta-alpha transformation in alumina scales formed on nickel-aluminides[END_REF], using a 633 nm HeNe laser in a Renishaw inVia Raman microscope. Surfaces were observed by optical microscopy with a Keyence VHX-600 and by SEM with a JEOL JSM-6510. In selected cases, thin foils were extracted from the specimen surface by focused ion beam (FIB) milling with a FEI Nova NanoLab 200, and characterized by transmission electron microscopy (TEM) on a Zeiss Libra 200FE and a JEOL JEM-2100F. In this case, deposits were washed off the surface prior to FIB milling. Polished cross-sections were prepared using conventional metallographic techniques. When a deposit was present, water-free cutting fluids and polishing suspensions were first used to retain any water-soluble products such as Na 2 SO 4 . Observations after short-term exposures showed that the sulfate had either reacted or evaporated, and so it was deemed appropriate to polish subsequent specimens in water-based media. The cross-sections were characterized by scanning electron microscopy (SEM) with a JEOL JSM-6510, and by EPMA with a JEOL JXA-8530F.

Results

Oxidation with no deposit

All alloys (Table 1) were oxidized 50 h in dry air, air-20H 2 O, CO 2 -20H 2 O, and CO 2 -20H 2 O-1.6O 2 with no deposit. External Al 2 O 3 scales formed in all cases; selected micrographs are shown in Fig. 1. Small regions of Cr, Co, and Ni oxides were occasionally found atop the scales, and the Al 2 O 3 contained up to 1 at. % Cr in its outer part. The β-rich Ni-19Co-15Cr-24Al-0.1Y alloy also formed Y aluminate precipitates within the scale that locally extended into the alloy (Fig. 1e). These were less numerous if not absent in the γ-rich alloys. No particular effect of N 2 , CO 2 or H 2 O such as nitridation, carburization or enhanced spallation was observed in these isothermal exposures.

Scales formed on the alloys with 12 at. % Al, which had the lowest β fractions, were relatively uniform. Average scale thicknesses were found to be consistently lower in the H 2 O-containing gases than in dry air, as shown in Fig. 2. As the β fraction, f β , increased, the effect of alloy microstructure became more prominent; specifically, the oxide located atop prior β domains was thinner than that formed on γ. For a simple mass balance reason, the depth of β dissolution decreased with increasing f β , and therefore the effect of phase constitution on scale thickness was mostly visible on the alloy with highest f β , as shown in Fig. 1e. Local thickness variations associated with alloy microstructure were less marked but also existed in the 16Al alloys. On average, these scales tended to be thinner than those formed on the 12Al alloys in dry air, and the difference between dry and H 2 O-containing gases was less marked. Two alloys, the β-rich Ni-19Co-15Cr-24Al-0.1Y and the γ-rich Ni-30Co-30Cr-12Al-0.1Y, were selected for shorter (5 and 20 min)

oxidation in CO 2 -20H 2 O-1.6O 2 .
As shown in Fig. 3, different scale thicknesses on γ and β outlined the β-rich alloy microstructure. The scale grown on γ was thinner (see interference colors) and smooth, while that grown on β was thicker and granular. Both were Al 2 O 3 , as verified by SEM-EDS and XRD. Analysis by PSLS further showed that after 5 min, the oxide grown on β contained θ and α-Al 2 O 3 in comparable proportions, while that formed on γ was fully transformed to α-Al 2 O 3 (Fig. 4). After 20 min, the transformation was complete or nearly complete for the oxide grown on both alloy phases. Cross-sectional observation confirmed that the scale was significantly thicker on β (Fig. 5).

The scale grown on the γ-rich alloy was mostly thin Al 2 O 3 , identified as α by PSLS. The underlying γ-β microstructure was not apparent. Some Cr-rich oxide was present atop the Al 2 O 3 with a surface coverage that varied with alloy grain orientation (Fig. 6a). This transient oxide was porous (Fig. 6b). Analysis by XRD showed that both Cr 2 O 3 and spinel phases were present along with Al 2 O 3 (Fig. 7). A TEM foil was prepared, which contained a zone of exclusive Al 2 O 3 and a zone with some transient oxide, as shown in Fig. 8. When present as a single oxide, the Al 2 O 3 had a typical two-layer microstructure: small equiaxed grains on the top, containing small amounts of Cr, Co, and Ni (1-6 at. %), and elongated grains at the bottom, with less than 1 at. % foreign species (Fig. 8b). When covered by transient oxide, the Al 2 O 3 consisted of a single layer with large grains (Fig. 8d). The outer layer was a porous Co-Cr spinel, while small grains containing Al, Cr, and Co were present between the two layers.

Early-stage reaction with mixed oxide-sulfate deposit

In order to investigate the mechanism of deposit-induced accelerated attack, short exposures (20 min and 1 h) were conducted using selected alloys and one quaternary deposit:

SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 in CO 2 -20H 2 O-1.6O 2 .
The present section focuses on the 20-min reaction of the β-rich Ni-19Co-15Cr-24Al-0.1Y and the γ-rich Ni-30Co-30Cr-12Al-0.1Y.

Surface observations after the deposit was washed showed that the corrosion product had poor adherence to the alloys, as a significant part had spalled. The remaining surface oxide was either Al-rich or Cr-rich, with local variations. The underlying γ-β microstructure could not be distinguished on either alloy. Analysis of the surface oxides by PSLS showed that the Al-rich oxide was mostly α-Al 2 O 3 , but a significant amount of θ-Al 2 O 3 remained, on both alloys. An example of a spectrum recorded on Ni-19Co-15Cr-24Al-0.1Y is included in Fig. 4.

Cross-sectional observations showed that the β-rich alloy mostly formed an ex- ternal Al 2 O 3 scale, with a few isolated zones where Al 2 O 3 intrusions less than 1-µm deep had grown into the alloy. By contrast, a transition to internal Al 2 O 3 growth, although localized in nature, was consistently observed on the γ-rich alloy. Varying degrees of degradation were observed along the specimen cross-section. Figure 9 shows the EPMA maps of a region in the earliest stage of the transition, where small but numerous Al Detailed characterization of the Ni-30Co-30Cr-12Al-0.1Y specimen was conducted by FIB and TEM. Searching for regions of interest by FIB milling confirmed that porosity was always present in the early stage of the external/internal Al 2 O 3 transition. Figure 11 shows the TEM analysis of the scale at a stage where intrusions have not quite yet formed. In fact, the scale was separated from the alloy by a gap, and only connected in places by bridging grains. The scale was remarkably different in its microstructure from that grown in the absence of a deposit (compare with Fig. 8). In particular, the oxide grains were much finer, and no columnar zone was present. Analysis by EDS (Fig. 11b) showed that the scale contained a Ca aluminosilicate layer, a mixed Al-Cr oxide layer, and Al 2 O 3 (from top to bottom). While most bridging grains across the alloy-scale gap were Al 2 O 3 , some were deduced to be Cr sulfide. Particles of the latter were also found within the scale and at the metal surface. Figure 12 shows the corrosion product in a more advanced stage of the transition, with a more variable and complex microstructure. Alumina intrusions developed inward and fine-grained Cr cluding within the Cr 2 O 3 . Analysis by EDS (Fig. 12b) showed that an Al 2 O 3 layer was still present locally at the metal surface, while the sides of a large pore were lined with Cr-rich oxide and some mixed Cr-Ni sulfide. On the scale surface was a silicate layer that contained Ca and Al.

Effect of atmosphere and evolution of the reaction morphology

Potential effects of the gas atmosphere were studied using a SiO 2 -20Al 4.3 ×10 -5 atm from available thermodynamic data [START_REF] Barin | Thermochemical Data of Pure Substances[END_REF]. No particular effect of the p O 2 or of the presence of CO 2 or H 2 O was found. The same reaction morphology was observed in all cases, and it was also equivalent to that obtained after 20-min reaction with the quaternary deposit (see previous section). Figure 13 Figure 14 shows the reaction products obtained after 1, 50, and 250 h exposures in dry air, with the same alloy and deposit. The presence of N 2 caused internal nitridation of Al, but this did not apparently make the deposit-induced degradation worse. The reaction morphology observed early was maintained over time. From 1 to 50 h, internal oxidation progressed, the nitrides coarsened and penetrated deeper, and a nitride-free zone developed below the scale, where Al depletion was most severe. Yet the new Al 2 O 3 layer formed at the base of the internal oxides provided effective repassivation, since the reaction product did not significantly thicken between 50 and 250 h. It is noted that nitridation was observed throughout the crosssections, even in zones where an apparently intact Al 2 O 3 scale was present and no intrusions had formed, which still existed after 250 h. No nitridation occurred when the same alloy was oxidized in air with no deposit.

Discussion

Effects of alloy and gas compositions in the absence of a deposit

The influence of alloy microstructure on the oxide formed in the absence of a deposit was mostly visible on the β-rich alloys. The features observed here are consistent with descriptions of γ-β systems oxidation by previous investigators, e.g., Refs. [START_REF] Clemens | TEM and SNMS studies of protective alumina scales on NiCrAlYalloys[END_REF][START_REF] Mennicke | Transient phase evolution during oxidation of a two-phase NiCoCrAlY bond coat[END_REF][START_REF] Levi | Alumina grown during deposition of thermal barrier coatings on NiCrAlY[END_REF][START_REF] Nijdam | On the microstructure of the initial oxide grown by controlled annealing and oxidation on a NiCoCrAlY bond coating[END_REF][START_REF] Sullivan | Vapor-phase-mediated phenomena associated with high temperature, high water content oxidation of MCrAlX bond coats[END_REF]. After 5 min, the oxide grown on β still retained some transient θ-Al 2 O 3 , while that formed on γ was fully α-Al 2 O 3 (Fig. 4). The faster θ → α transformation is usually attributed to the higher Cr content of the γ phase: the transient Cr 2 O 3 it forms is thought to act as a template and to facilitate the nucleation of the isostructural α-Al 2 O 3 [START_REF] Brumm | The oxidation behavior of NiAl. 1. Phasetransformations in the alumina scale during oxidation of NiAl and NiAl-Cr alloys[END_REF][START_REF] Kitajima | Rapid formation of α-Al 2 O 3 scale on an Fe-Al alloy by pure-metal coatings at 900 °C[END_REF]. The parabolic constant of θ is two orders of magnitude larger than that of α at 1100 • C [START_REF] Brumm | The oxidation behavior of NiAl. 1. Phasetransformations in the alumina scale during oxidation of NiAl and NiAl-Cr alloys[END_REF][START_REF] Rybicki | Effect of the θ-α-Al 2 O 3 transformation on the oxidation behavior of β-NiAl+Zr[END_REF], which is why the scale grown on β was thicker than that formed on γ after 20 min (Figs. 3 and5). Yet, following the analysis of Brumm and Grabke [START_REF] Brumm | The oxidation behavior of NiAl. 1. Phasetransformations in the alumina scale during oxidation of NiAl and NiAl-Cr alloys[END_REF], a delayed θ → α transformation also favors grain growth, which implies that the α-Al 2 O 3 eventually present on β should have a coarser microstructure than that grown on γ. Since α-Al 2 O 3 growth at this temperature is primarily by grain boundary diffusion, a larger-grained scale is expected to grow slower. Indeed, after 50 h, the oxide formed on β was significantly thinner than that formed on γ (Fig. 1e).

Although H 2 O is known to enhance Al 2 O 3 scale spallation [START_REF] Birks | Introduction to the High Temperature Oxidation of Metals[END_REF][START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF][START_REF] Saunders | The oxidation behaviour of metals and alloys at high temperatures in atmospheres containing water vapour: a review[END_REF], this did not affect the present isothermal exposures, as the 0.1 at. % Y provided good metal/oxide adherence even in wet atmospheres. Similarly, while H 2 O may trigger internal oxidation of marginal Al 2 O 3 -formers [START_REF] Zhao | Assessment of the detrimental effects of steam on Al 2 O 3 -scale establishment[END_REF], this was not the case here due to the relatively high Al and Cr concentrations. Chromia volatilization in wet atmospheres via Cr hydroxide formation is significant at 1100 • C [START_REF] Opila | Predicting oxide stability in high-temperature water vapor[END_REF]. While this is not the case of Al 2 O 3 (Al(OH) 3 formation that is negligible below about 1300 • C [START_REF] Opila | Predicting oxide stability in high-temperature water vapor[END_REF]), NiAl 2 O 4 was shown to sublime at significant rates in the presence of H 2 O at 1100 • C [START_REF] Sullivan | Vapor-phase-mediated phenomena associated with high temperature, high water content oxidation of MCrAlX bond coats[END_REF]. In the present work, transient Cr 2 O 3 and Cr-containing spinels formed on the γ phase, and their volatilization likely accounts for the reduction in scale thicknesses observed in H 2 O-containing gases compared to dry air on the γ-rich alloys (Fig. 2). In the case of CO 2 -20H 2 O, the reduced p O 2 might also have contributed to produce thinner scales, since less transient oxides are typically observed at low p O 2 [START_REF] Nijdam | Promoting exclusive α-Al 2 O 3 growth upon high-temperature oxidation of NiCrAl alloys: experiment versus model predictions[END_REF]. Finally, no effect of CO 2 was observed, which is consistent with observations made by other investigators after longer term testing in CO 2 -H 2 O [START_REF] Dryepondt | Effect of H 2 O and CO 2 on the oxidation behavior and durability at high temperature of ODS-FeCrAl[END_REF][START_REF] Lance | The effect of cycle frequency, H 2 O and CO 2 on TBC lifetime with NiCoCrAlYHfSi bond coatings[END_REF].

Reaction mechanism of deposit-induced degradation

Exposure of the γ-rich alloys Ni-30Co-27Cr-12Al-0.1Y and Ni-30Co-30Cr-12Al-0.1Y to SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 and SiO 2 -20Al 2 O 3 -33CaO-8MgO-6Fe 2 O 3 -5Na 2 SO 4 -1K 2 SO 4 deposits caused a transition to the internal oxidation of Al. The reaction products were equivalent regardless of gas composition. The same features were also observed after reaction of various MCrAlY compositions at 1100 • C with similar deposits, i.e., containing CaO and Na 2 SO 4 [START_REF] Gheno | Modes of deposit-induced accelerated attack of MCrAlY systems at 1100 • C[END_REF]. We conclude that for a large part, the deposit-induced degradation of MCrAlY systems in S-free atmospheres at this temperature proceeds according to a common mechanism.

As mentioned in Section 3, the degradation was localized. The morphological continuity between intact Al 2 O 3 scale, shallow intrusions and internal oxidation (Figs. 10 and13) suggests that these were evolutionary steps of the same process. This process is shown schematically in Fig. 15a, and can be summarized as follows. In examining the cause of the transition from external to internal Al oxidation, we note that the reaction morphology adopted after Al 2 O 3 failure is, precisely, related to the "mode of failure," and ultimately to the conditions underlying Al 2 O 3 stability as an external scale, which is dictated by kinetics. Due to the high thermodynamic stability of Al 2 O 3 , the Al activity required for Al-Al 2 O 3 equilibrium is vanishingly small at practical p O 2 values. Any Al-containing alloy is thus expected to form Al 2 O 3 ; yet it may fail to form it as an external scale if the outward Al flux is not sufficient to counter the inward O flux, in which case Al oxidizes internally instead [START_REF] Wagner | Reaktionstypen bei der Oxydation von Legierungen[END_REF][START_REF] Rapp | Kinetics, microstructures and mechanism of internal oxidation -its effect and prevention in high temperature alloy oxidation[END_REF]. In accordance with the so-called gettering effect [START_REF] Wagner | Passivity and inhibition during the oxidation of metals at elevated temperatures[END_REF], if a Cr 2 O 3 scale is present, the low p O 2 at the metal surface generates a limited O flux, and the Al concentration required for its external oxidation is also found to be very low. When an Al 2 O 3 scale is formed, the Al flux still needs to match that required for the oxide to grow. Assuming parabolic oxidation kinetics and a constant Al diffusion coefficient in the alloy (D Al ), the minimum bulk Al fraction to sustain Al 2 O 3 growth is given by [START_REF] Wagner | Theoretical analysis of the diffusion processes determining the oxidation rate of alloys[END_REF]:

N * Al = F 1 2 k c D Al ( 1
)
where F is defined by F (u) = √ πu(1 -erf u) exp u 2 and k c is the parabolic constant for metal recession, defined by d 2 = 2k c t with d the depth of the metal/oxide interface relative to the original alloy surface. For N Al < N * Al , a Cr-rich alloy would form external Cr 2 O 3 and internal Al 2 O 3 , as shown in Fig. 15a. This is the case of certain wrought austenitic Ni-base alloys with 2-5 at. % Al at temperatures above

∼ 850 • C [43-45]. For N Al > N *
Al , an Al 2 O 3 scale is kinetically stable but will eventually fail after the alloy Al reservoir is fully exhausted. When the depletion profiles generated on opposite sides of a specimen reach its center, the Al concentration at the metal/oxide interface starts to decrease, and the interfacial p O 2 rises until a Cr 2 O 3 sublayer forms, i.e., Cr reduces the Al 2 O 3 . In this case, no internal oxidation of Al is possible, as its concentration profile is essentially flat, and Al 2 O 3 is not thermodynamically stable even below the surface. This "intrinsic chemical failure" [START_REF] Evans | Mechanisms of breakaway oxidation and application to a chromia-forming steel[END_REF], observed for long or high-temperature exposures of thin specimens [START_REF] Quadakkers | Oxidation induced lifetime limits of thin walled, iron based, alumina forming, oxide dispersion strengthened alloy components[END_REF][START_REF] Wright | High-temperature oxidation life characteristics of ODS-Fe 3 Al[END_REF][START_REF] Stott | Breakdown of protective scales during the oxidation of thin foils of Fe-20Cr-5Al alloys at high temperatures[END_REF][START_REF] Al-Badairy | A comparison of breakaway oxidation in wedge-shaped and parallel sided coupons of FeCrAl alloys[END_REF][START_REF] Wright | Influence of aluminum depletion effects on the calculation of the oxidation lifetimes of FeCrAl alloys[END_REF][START_REF] Tatlock | Air oxidation of commercial FeCrAlRE alloy foils between 800 and 950[END_REF][START_REF] Strehl | Two calculations concerning the critical aluminium content in Fe20Cr5Al alloys[END_REF][START_REF] Young | Slow transition from protective to breakaway oxidation of Haynes 214 foil at high temperature[END_REF][START_REF] Chyrkin | Effect of thermal cycling on protective properties of alumina scale grown on thin Haynes 214 foil[END_REF], is represented schematically in Fig. 15b.

The alloys used here formed external Al 2 O 3 in the absence of a deposit (N Al > N * Al ), but developed a reaction morphology typical of alloys with sub-critical Al fraction (N Al < N * Al ) with an oxide-sulfate deposit present. The reaction of Al 2 O 3 with CaO to form Ca aluminates or its dissolution in molten Na 2 O-SiO 2 each leads to an increased rate of Al consumption, potentially exceeding the Al flux available from the alloy and triggering internal oxidation. This amounts to a modification of the mass balance underlying Eq. ( 1), where k c is replaced by a larger constant accounting for Al 2 O 3 consumption [START_REF] Gheno | High temperature reaction of MCrAlY coating compositions with CaO deposits[END_REF][START_REF] Gheno | Kinetics of Al 2 O 3 -scale growth by oxidation and dissolution in molten silicate[END_REF]. Our study of various alloy and deposit compositions [START_REF] Gheno | Modes of deposit-induced accelerated attack of MCrAlY systems at 1100 • C[END_REF] showed that deposits containing CaO, SiO 2 , and Na 2 O (as Na 2 SO 4 ) were susceptible to induce both modes of attack, solid-state reaction and liquid dissolution, and that Al 2 O 3 failure in γ-rich MCrAlY systems resulted from a combination of the two. While the assertion is justified, the present results indicate that Al 2 O 3 scales formed under the oxide-sulfate deposit were markedly different from those grown in the absence of a deposit, and suggest that additional factors played a role in their failure.

Deposit effects on Al 2 O 3 properties

The presence of the oxide-sulfate deposit affected the Al 2 O 3 scale in several ways. The possible causes and consequences of each aspect are discussed in the following.

Al 2 O 3 phase transition
The θ → α transition was delayed with the deposit present, compared to deposit-free exposures (see Fig. 4 for Ni-19Co-15Cr-24Al-0.1Y, equivalent results were obtained for Ni-30Co-30Cr-12Al-0.1Y in locations where an Al 2 O 3 scale was still present). The crystalline phases initially present in the deposit (α-quartz, α-Al 2 O 3 , B1-CaO, B1-MgO, α-Fe 2 O 3 ) or formed by reaction with the scale (Ca and Mg aluminates) could in principle stabilize θ or α-Al 2 O 3 via epitaxial relationships. Alternatively, the development of a reduced p O 2 beneath the deposit, in particular in the presence of molten phases (Na 2 SO 4 or Na 2 O-SiO 2 ), may also affect the transition. According to Nijdam et al. [START_REF] Nijdam | On the microstructure of the initial oxide grown by controlled annealing and oxidation on a NiCoCrAlY bond coating[END_REF], a reduced p O 2 can either stabilize or suppress θ-Al 2 O 3 on a NiCoCrAlY alloy at 1100 • C, depending on the p O 2 level and the heating procedure. Alvarado-Orozco et al. [START_REF] Alvarado-Orozco | Kinetic study of the competitive growth between θ-Al 2 O 3 and α-Al 2 O 3 during the early stages of oxidation of β-(Ni,Pt)Al bond coat systems: Effects of low oxygen partial pressure and temperature[END_REF] reported that a p O 2 of 10 -5 atm stabilized θ-Al 2 O 3 on β-NiAl at 1000-1200 • C, compared to atmospheric pressure.

The mechanism of the θ → α transition is complex and system-dependent, and why a deposit would stabilize θ-Al 2 O 3 remains unclear at present. However, a consequence is readily identified: as the parabolic constant of θ is two orders of magnitude larger than that of α at 1100 • C [START_REF] Brumm | The oxidation behavior of NiAl. 1. Phasetransformations in the alumina scale during oxidation of NiAl and NiAl-Cr alloys[END_REF][START_REF] Rybicki | Effect of the θ-α-Al 2 O 3 transformation on the oxidation behavior of β-NiAl+Zr[END_REF], delaying the transition would maintain a higher rate of Al consumption, enhance Al depletion in the alloy and thereby favor internal oxidation. As discussed in Section 4.1, a delay is also expected to eventually produce a coarse-grained, slow-growing α-Al 2 O 3 .

Al 2 O 3 microstructure
The microstructure of the Al 2 O 3 grown beneath the deposit was significantly different from that of the scale grown in deposit-free conditions (compare Figs. 8 and11). In particular, the Al 2 O 3 grain size was much smaller with the deposit present, which is at variance with what could be expected from the delayed θ → α transition. This suggests that another aspect prevailed in determining the scale microstructure; specifically, the deposit either favored a faster nucleation or hindered grain growth or sintering, independently of the oxide crystal structure. Furthermore, the deposit caused the development of an interfacial gap, which was not observed in deposit-free conditions, and which indeed was much more extensive than the porosity usually observed in Al 2 O 3 scales [START_REF] Pint | On the formation of interfacial and internal voids in α-Al 2 O 3 scales[END_REF]. Both the smaller grain size and the interfacial gap can be interpreted as the result of a poisoning of interfacial reactions. As argued below, this is thought to be related to the presence of S in the system.

At a temperature where Al 2 O 3 grows mainly by grain boundary diffusion, a smaller grain size is expected to yield a faster growth and enhanced Al consumption, adding to the delayed θ → α transition and further favoring internal oxidation. Exactly how much each contributed to Al 2 O 3 failure is difficult to determine. Scale thicknesses measured after 20-min exposure with and without deposit, if converted to parabolic constants, both yield k p values about 10 times larger than those obtained after 50-h exposure. This reflects the rapid evolution of Al 2 O 3 microstructure in the first hours of oxidation, as discussed above, and as shown by thermogravimetric analysis of similar NiCoCrAlY compositions [START_REF] Gheno | High temperature reaction of MCrAlY coating compositions with CaO deposits[END_REF]. Such an evolution precludes a quantitative analysis of the effect of short-term deposit-induced microstructural changes on scaling kinetics, all the more so since these kinetics were also affected by the scale reaction with the deposit, independently of its microstructure. The presence of the deposit rendered the Al 2 O 3 scale permeable to secondary oxidants: S which was present as deposited sulfate, and N which was present as N 2 in air exposures. Chromium sulfide formation requires both an elevated S activity and a low p O 2 (high p O 2 stabilize sulfates rather than sulfides). Thus, the presence of such sulfides within the scale and at the metal/oxide interface (Fig. 11) indicates that the scale integrity was preserved to some extent, and that S permeated via small defects such as grain boundaries or sub-micron channels, rather than large cracks. Indeed, the scale appeared compact in the TEM. A high p S 2 could then be obtained by local equilibrium between incoming S-bearing molecules and the surrounding oxide, as discussed by Birks et al. [START_REF] Birks | Corrosion mechanisms of metals and alloys in multicomponent oxidative environments[END_REF][START_REF] Pope | The penetration by sulfur of NiO scales growing on nickel[END_REF] in the case of Ni exposed to CO-CO 2 -SO 2 . Similar cases of attack by a secondary oxidant have been reported in various systems, e.g., carburization below Cr 2 O 3 grown in CO-CO 2 [START_REF] Giggins | Corrosion of metals and alloys in mixed gas environments at elevated temperatures[END_REF][START_REF] Zheng | High-temperature corrosion of Cr 2 O 3 -forming alloys in CO-CO 2 -N 2 atmospheres[END_REF][START_REF] Gheno | Carburisation of ferritic Fe-Cr alloys by low carbon activity gases[END_REF]; yet the secondary oxidant was always present in the gas environment. On the other hand, Goebel and Pettit [START_REF] Goebel | Na 2 SO 4 -induced accelerated oxidation (hot corrosion) of nickel[END_REF] showed that under a molten Na 2 SO 4 deposit, O consumption due to oxidation could produce an elevated p S 2 . Recent work on Ni reaction at 700 • C [START_REF] Gheno | On the hot corrosion of nickel at 700 • C[END_REF] examined how this applied in S-containing and S-free atmospheres, with and without Na 2 SO 4 . Sulfidation was observed below Na 2 SO 4 , even in air. By analogy, we conclude that in the present case, a high p S 2 was generated at the oxide surface due to oxidation underneath the molten Na 2 SO 4 , and further amplified as S-bearing molecules traversed the scale, equilibrating with the surrounding p O 2 .

Alumina is generally believed to be a good barrier to secondary oxidants, more effective than Cr 2 O 3 in particular [START_REF] Young | High Temperature Oxidation and Corrosion of Metals[END_REF], and reports of sulfidation below apparently compact Al 2 O 3 scales are not frequent. Cases of Al 2 O 3 permeation by S were reported on β-NiAl [START_REF] Brumm | The oxidation of NiAl. 3. Internal and intergranular oxidation[END_REF][START_REF] Schramm | Sulfidation behaviour of nickel aluminides[END_REF][START_REF] Grabke | Accelerated oxidation, internal oxidation, intergranular oxidation, and pesting of intermetallic compounds[END_REF][START_REF] Gheno | Reaction morphologies developed by nickel aluminides in type II hot corrosion conditions: The effect of chromium[END_REF], where it led to the fast growth of porous Al 2 O 3 nodules. Similar nodules were observed in cases where the S was present in the metal before oxidation [START_REF] Goebel | Influence of sulfides on oxidation behavior of nickel-base alloys[END_REF][START_REF] Smialek | Non-protective alumina growth in sulfur-doped NiAl(Zr)[END_REF]; the widespread porosity was then attributed to SO 2 evolution. The present situation is different, as the scale remained thin and apparently compact above the gap separating it from the metal. Interfacial segregation of S was found in many systems [START_REF] Funkenbusch | Reactive element-sulfur interaction and oxide scale adherence[END_REF][START_REF] Smeggil | A relationship between indigenous impurity elements and protective oxide scale adherence characteristics[END_REF][START_REF] Smialek | Adherent Al 2 O 3 scales formed on undoped NiCrAl alloys[END_REF][START_REF] Hou | Oxide scale adhesion and impurity segregation at the scale/metal interface[END_REF][START_REF] Smialek | Effects of hydrogen annealing, sulfur segregation and diffusion on the cyclic oxidation resistance of superalloys -a review[END_REF][START_REF] Smialek | Maintaining adhesion of protective Al 2 O 3 scales[END_REF][START_REF] Hou | The effect of H 2 -anneal on the adhesion of Al 2 O 3 scales on a Fe 3 Al-based alloy[END_REF][START_REF] Smialek | Optimizing scale adhesion on single crystal superalloys[END_REF][START_REF] Hou | Sulfur segregation at Al 2 O 3 /γ-Ni + γ'-Ni 3 Al interfaces: Effects of Pt, Cr and Hf additions[END_REF] to reduce the oxide adherence. This was in part attributed to a reduction in the surface energy of interfacial pores [START_REF] Grabke | Segregation of sulfur during growth of oxide scales[END_REF][START_REF] Grabke | Segregation beneath oxide scales[END_REF][START_REF] Hou | Interfacial segregation, pore formation, and scale adhesion on NiAl alloys[END_REF] (these pores occur naturally due to vacancy injection, but remain much smaller in the absence of S). First-principles calculations also show that S weakens an intact Ni/Al 2 O 3 bond [START_REF] Hong | Sulfur at nickel-alumina interfaces: Molecular orbital theory[END_REF][START_REF] Zhang | Influence of sulfur on the adhesion of the nickel/alumina interface[END_REF][START_REF] Jiang | Pt effects in γ-Ni(Al)/α-Al 2 O 3 adhesion[END_REF].

In accordance with these findings, we conclude that S ingress through nanosized defects of an initially compact scale, and its interfacial segregation, led to the observed gap; the latter might have been further stabilized by the establishment of an internal S 2 -SO 2 mixture, which, like CO-CO 2 and H 2 -H 2 O mixtures in Fe oxide scales [START_REF] Fujii | The mechanism of the high-temperature oxidation of iron-chromium alloys in water vapor[END_REF][START_REF] Rahmel | Einfluss von Wasserdampf und Kohlendioxyd auf die Oxydation von Eisen in Sauerstoff bei Hohen Temperaturen[END_REF][START_REF] Fujii | Carburization of Fe-Cr alloys during oxidation in dry carbon dioxide[END_REF][START_REF] Ehlers | Enhanced oxidation of the 9%cr steel P91 in water vapour containing environments[END_REF], would substantially increase the total pressure while a reduced p O 2 is maintained. A practical consequence of the peculiar microstructure observed in Fig. 11, and further indication that it was indeed related to S, is that all scales formed under Na 2 SO 4 -containing deposits showed significant spallation.

Furthermore, the fact that the deposit also made Al 2 O 3 scales permeable to N, even in locations where they were apparently intact and the transition to internal oxidation had not started, leads to the conclusion that the scale microstructural changes were accompanied by a loss of protective properties. Specifically, despite the scale compact appearance in the TEM, the density of nano-scale defects capable of transmitting secondary oxidants such as N 2 or SO 2 must have been greatly increased. The grain boundary density increased simply due to the smaller grain size, but this would hardly account for the increased permeability. Instead, the grain boundary properties must have themselves been affected.

Finally, in terms of scale stability, a common effect of the reaction with secondary oxidants is that it reduces the availability of the oxide-forming element, which may lead to non-selective oxidation [START_REF] Giggins | Corrosion of metals and alloys in mixed gas environments at elevated temperatures[END_REF]. Here sulfides were either Cr-rich (Fig. 11) or mixed Cr-Ni (Fig. 12), which did not affect the Al supply. Internal nitrides were Al-rich, but no significant effect on the extent of internal oxidation was observed.

Effects of alloy and gas compositions on deposit-induced attack

Both the delayed θ → α transition and the smaller grain size of the Al 2 O 3 grown with the deposit present are expected to accelerate Al consumption and favor internal oxidation. This would affect γ-rich alloys more than β-rich alloys, because the former have a smaller Al reservoir to start with. Furthermore, the θ → α transition is more rapid on γ than on β in the absence of a deposit (Fig. 4), which makes γ-rich alloys all the more susceptible to the delaying action of the deposit.

Alloys with larger Cr concentrations tend to sustain more degradation by CaOrich deposits, because of the rapid reaction of transient Cr-rich oxide with CaO to form a liquid Ca chromate [START_REF] Gheno | High temperature reaction of MCrAlY coating compositions with CaO deposits[END_REF]. While the specifics of the reaction mechanism are not fully understood at present, larger Cr concentrations were also found to locally accelerate the initiation of Al 2 O 3 dissolution in molten Na 2 O-SiO 2 , promoting Al 2 O 3 failure [START_REF] Gheno | Modes of deposit-induced accelerated attack of MCrAlY systems at 1100 • C[END_REF][START_REF] Gheno | Kinetics of Al 2 O 3 -scale growth by oxidation and dissolution in molten silicate[END_REF]. Therefore, the fact that the γ-rich alloys used here had high Cr levels and formed relatively large amounts of Cr-rich transient oxide, compared to the Cr-lean β-rich alloy (see Section 3.1), presumably contributed to the nonselective oxidation exhibited under the deposit. Accordingly, the reduced transient Cr oxidation observed in deposit-free conditions in wet atmospheres, compared to dry air, and due to H 2 O-mediated volatilization (see Section 4.1), could have mitigated the deleterious effect of the large Cr levels. This was in fact not the case, as results observed with a deposit present in wet atmospheres and dry air were equivalent. A possible reason for this is that the deposit may have affected the gas flow accessible to the specimen surfaces, hampering volatilization; in addition, the reaction of Cr-rich oxide with CaO or Na 2 O-SiO 2 in the very early stage of the exposure may have been so fast that it was insensitive to volatilization.

Finally, the fact that after Al 2 O 3 breakdown, an external Cr 2 O 3 layer formed and remained stable, favoring repassivation (Section 4.2), may seem to be in contradiction with the tendency of Cr 2 O 3 to react very rapidly with CaO in the transient stage [START_REF] Gheno | High temperature reaction of MCrAlY coating compositions with CaO deposits[END_REF]. Molten Ca chromate was observed to form under mixed deposits despite the fact that adding Al 2 O 3 or SiO 2 reduced the CaO activity, compared to a pure CaO deposit [START_REF] Gheno | Modes of deposit-induced accelerated attack of MCrAlY systems at 1100 • C[END_REF]. The apparent difference in Cr 2 O 3 reactivity between the transient stage and steady state is likely due to a CaO depletion in the regions of the deposit closest to the specimen surface. Experiments where some fresh deposit is periodically re-applied would be useful to study the stability of the external Cr 2 O 3 layer in more realistic conditions.

Conclusions

The deposit-induced accelerated degradation of γ-rich NiCoCrAlY alloys involved, in its early stage, the breakdown of an initially formed Al 2 O 3 scale, the development of Al 2 O 3 intrusions and internal precipitates, and the establishment of an external Cr 2 O 3 layer. This was followed by a repassivation stage, where a new Al 2 O 3 layer formed at the internal oxidation front. This morphological evolution is in part similar to that observed on more dilute NiCrAl-base alloys; its occurrence on compositions which form Al 2 O 3 scales in the absence of a deposit reflects the "mode of failure" associated with the deposit-induced attack: the Al 2 O 3 scale fails because of its reaction with the deposit, but once the alloy is isolated from the deposit by a Cr 2 O 3 layer, it returns to forming a continuous Al 2 O 3 layer. The protecting role of the external Cr 2 O 3 may, however, be challenged if fresh CaO-rich deposit is re-applied, because of the formation of liquid Ca chromate.

Detailed analysis of the reaction products formed during the early stage of the reaction, before the transition from external to internal Al 2 O 3 growth, showed that in addition to reacting in the solid state (to form Ca aluminates) or dissolving in a liquid phase (molten Na 2 O-SiO 2 ), the Al 2 O 3 scale microstructure was affected by the deposit. Compared to deposit-free conditions, the θ → α transformation was delayed, the scale grain size was reduced, both factors further increasing Al consumption and favoring internal oxidation. In addition, the scale adherence was reduced, to an extent where a nearly continuous gap developed at the metal/oxide interface. Sulfur ingress through nano-sized defects of the initially compact scale, and its interfacial segregation, were concluded to be the cause of both the interfacial gap and the reduced grain size. The observation that scales grown under the deposit were permeable to N also led to the conclusion that the modified microstructure involved a loss of protective properties, specifically, that the grain boundary ability to block molecules of secondary oxidants was reduced.

Sulfide formation in the alloys exposed to the deposit showed that Na 2 SO 4 allowed sulfidizing conditions to establish, even in S-free atmospheres. Sulfidation and nitridation were consequences of the deposit-induced degradation, but neither was found in itself to significantly affect the corrosion resistance. Similarly, no interaction between the other gas constituents (CO 2 , H 2 O) and the deposits was observed.

Figure 1 :

 1 Figure 1: Oxide scales formed during exposure of (a), (b) Ni-30Co-27Cr-12Al-0.1Y, (c), (d) Ni-30Co-30Cr-16Al-0.1Y, and (e) Ni-19Co-15Cr-24Al-0.1Y in (a, c) dry air and (b, d, e) CO 2 -20H 2 O-1.6O 2 , with no deposit, 50 h at 1100 • C.

Figure 3 :

 3 Figure 3: Surface micrographs of oxide grown on Ni-19Co-15Cr-24Al-0.1Y after 20 min in CO 2 -20H 2 O-1.6O 2 at 1100 • C, with no deposit. In top-left (light micrograph), the dark and neon phases are the oxides grown on β and γ, respectively. Top-right is an SEM-BSE image; at bottom (SEM-SE), the granular and smooth phases are the oxides grown on β and γ, respectively.

Figure 4 :

 4 Figure 4: PSLS spectra recorded on the surface of Ni-19Co-15Cr-24Al-0.1Y reacted in CO 2 -20H 2 O-1.6O 2 at 1100 • C, with and without a SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 deposit. Note that θ-Al 2 O 3 luminescence is 10-12 times weaker than that of α-Al 2 O 3[START_REF] Tolpygo | Microstructural study of the theta-alpha transformation in alumina scales formed on nickel-aluminides[END_REF]. The proportions of θ and α varied on the specimen surfaces; the spectra included here are representative of ∼ 10 spectra recorded for each phase on each specimen. With the deposit, the underlying γ-β microstructure could not be identified from the surface: the spectra were recorded at random locations.

Figure 5 :

 5 Figure 5: Cross-sectional view of Al 2 O 3 scale grown on Ni-19Co-15Cr-24Al-0.1Y after 20 min in CO 2 -20H 2 O-1.6O 2 at 1100 • C, with no deposit. The scale is thicker above prior β regions (white arrows), where a thin band of γ has formed between the scale and the underlying β due to Al consumption.

Figure 6 :

 6 Figure 6: Surface micrographs of surface oxide grown on Ni-30Co-30Cr-12Al-0.1Y after 20 min in CO 2 -20H 2 O-1.6O 2 at 1100 • C, with no deposit. In the light image (left), Al 2 O 3 appears neon, while Cr-rich oxide appears black. In the SEM image (right), Al 2 O 3 appears dark-gray, while Cr-rich oxide appears light-gray.

Figure 7 :Figure 8 :

 78 Figure 7: XRD patterns of oxide grown on Ni-30Co-30Cr-12Al-0.1Y after 20 min in CO 2 -20H 2 O-1.6O 2 at 1100 • C, with no deposit. Co radiation, grazing angle (Ω = 0.5 • ).

  2 O 3 intrusions developed from the scale inward. An external Cr-rich oxide layer that included small amounts of Ni and Co, denoted Cr 2 O 3 , formed above the Al 2 O 3 . Aluminosilicates containing Ca and Na were present above the Cr 2 O 3 . Remarkably, S-rich zones were detected near the Al 2 O 3 layer. A micrograph representative of the morphologies observed along the cross-section is shown in Fig. 10, where regions of apparently intact Al 2 O 3 are seen side-by-side with Al 2 O 3 intrusions and regions of internal Al 2 O 3 /external Cr 2 O 3 . Some degree of porosity was consistently observed at the base of the scale where the intrusions initiated.

  2 O 3 locally replaced Al 2 O 3 as the external layer. Pores were constantly present throughout the corrosion product, in-

Figure 9 :Figure 10 :

 910 Figure 9: SEM view and WDS maps of corrosion product formed in the early stage of accelerated degradation on Ni-30Co-30Cr-12Al-0.1Y after 20-min exposure to SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 deposit in CO 2 -20H 2 O-1.6O 2 at 1100 • C.

Figure 11 :

 11 Figure 11: TEM analysis of oxide scale formed on Ni-30Co-30Cr-12Al-0.1Y after 20-min exposure to SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 deposit in CO 2 -20H 2 O-1.6O 2 at 1100 • C, in region where accelerated corrosion has not started. (a) STEM high angle annular dark field (HAADF) overview; (b) STEM-HAADF view of oxide scale and alloy-scale interfacial gap, and (c) corresponding EDS maps.

  presents the EPMA maps recorded after 1 h in CO 2 -20H 2 O-1.6O 2 . While the alloy was still partially covered by an external Al 2 O 3 scale, in other regions, Al 2 O 3 intrusions penetrated about 5 µm into the metal. These partially joined at the reaction front to form a subsurface Al 2 O 3 layer. Sulfur enrichments were found at locations corresponding to the initially formed Al 2 O 3 and to the intrusions. A Cr 2 O 3 layer covered by a mixed silicate was present above the internal oxidation zone.

Figure 12 :

 12 Figure 12: TEM analysis of corrosion product formed on Ni-30Co-30Cr-12Al-0.1Y after 20-min exposure to SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 deposit in CO 2 -20H 2 O-1.6O 2 at 1100 • C, in region of accelerated corrosion. The scale contains Al 2 O 3 and Cr 2 O 3 layers surmounted by mixed silicates, and Al 2 O 3 fingers grow internally. (a) STEM-HAADF images; (b) EDS maps.

Figure 13 :Figure 14 :

 1314 Figure 13: SEM view and WDS maps of corrosion product formed on Ni-30Co-27Cr-12Al-0.1Y after 1 h exposure to SiO 2 -20Al 2 O 3 -33CaO-8MgO-6Fe 2 O 3 -5Na 2 SO 4 -1K 2 SO 4 deposit in CO 2 -20H 2 O-1.6O 2 at 1100 • C. Repassivation starts as an Al 2 O 3 sublayer forms at the base of the internal oxidation zone.

Figure 15 :

 15 Figure 15: Schematic representation of transition from Al 2 O 3 to Cr 2 O 3 scaling in situations of (a) deposit-induced attack of Al 2 O 3 -former (present work) or oxidation of marginal Al 2 O 3former, and (b) "intrinsic chemical failure" due to complete exhaustion of Al in Al 2 O 3former.

Al 2 O

 2 3 permeability to secondary oxidants

Table 1 :

 1 Nominal composition of NiCoCrAlY alloys used in the present study (at. %). Volume fractions of β, f β , were measured by image analysis after annealing at 1100 • C. Isothermal corrosion experiments were conducted at 1100 • C in flowing dry air (< 1 ppm CO 2 and CO, < 2 ppm H 2 O), air-20H 2 O, CO 2 -20H 2 O and CO 2 -20H 2 O-1.6O

	#	Elemental concentration (at. %)	f β at 1100 • C
		Ni	Co	Cr	Al	Y	(vol. %)
	UT2	42	19	15	24	0.1	57
	UT3	31	30	27	12	0.1	13
	UT4	28	30	30	12	0.1	16
	UT5	25	30	33	12	0.1	18
	UT6	27	30	27	16	0.1	34
	UT7	24	30	30	16	0.1	40
	UT8	21	30	33	16	0.1	42

ble 1. The alloys precipitated a small amount of yttrium-rich intermetallics ( 0.1 vol. %), and some also contained σ-CoCr (1, 2, and 22 vol. % in Ni-30Co-33Cr-12Al-0.1Y, Ni-30Co-30Cr-16Al-0.1Y, and Ni-30Co-33Cr-16Al-0.1Y, respectively).

  [START_REF] Bose | High Temperature Coatings[END_REF]

  Figure 2: Average Al 2 O 3 scale thickness measured on γ-rich alloys after exposure to various gases, 50 h at 1100 • C.

	dry air
	air-H 2 O
	CO 2 -H 2 O
	CO 2 -H 2 O-O 2
	Ni-30Co-27Cr-12Al-0.1Y Ni-30Co-30Cr-12Al-0.1Y

  2 O 3 -33CaO-8MgO-6Fe 2 O 3 -5Na 2 SO 4 -1K 2 SO 4 deposit. This composition can be classified as class C fly ash, and has similar effects as the SiO 2 -25Al 2 O 3 -40CaO-10Na 2 SO 4 mixture [26]. The γ-rich alloy Ni-30Co-27Cr-12Al-0.1Y was exposed to SiO 2 -20Al 2 O 3 -33CaO-8MgO-6Fe 2 O 3 -5Na 2 SO 4 -1K 2 SO 4 for 1 h in dry air, CO 2 -20H 2 O-1.6O 2 , and CO 2 -20H 2 O. The equilibrium p O 2 in the latter gas mixture is calculated to be

Acknowledgements

This work was supported by the Department of Energy through the University Turbine Systems Research (UTSR) Program run by the National Energy Technology Laboratory, award number DE-FE0007271, Seth Lawson, Project Manager. The authors thank Maryam Zahiri Azar and Arthur Heuer (Case Western Reserve University), as well as Susheng Tan (University of Pittsburgh), for the FIB and TEM work.


Michel Daher Mansour

Romain Parret

Laurence Masson

Nanoscale investigation of Si nanoribbon growth on Ag(110)

come   L'archive ouverte pluridisciplinaire

Electronic mail: masson@cinam.univ-mrs.fr

We present a nanoscale investigation by means of scanning tunneling microscopy of Si nanostructure growth on the anisotropic silver (110) surface, in the submonolayer range.

Four types of Si nanostructures are studied statistically as a function of the substrate temperature in the range 300 K to 500 K: isolated single-and double nanoribbons, which differ only by a factor of two in their width, and their respective bi-dimensional counterparts in the self-assembly regime. Our observations highlight different growth regimes controlled by kinetics. Below 320 K, the Si adatoms diffuse along the easy [11 0] direction, forming essentially isolated single nanoribbons randomly distributed on the silver terraces. At higher temperatures, transverse diffusion along the [001] direction is activated and a competition between the growth of self-assembled single nanoribbons and isolated double nanoribbons is observed. Above 440 K, a transition from one-to two-dimensional double nanoribbon growth is evidenced. At 490 K, the Si deposition results in the formation of massively self-assembled double nanoribbons. Based on Arrhenius analyses, activation barriers are found to be (125 + 15) meV and (210 + 20) meV for the formation of isolated Si nanoribbons and self-assembled Si double nanoribbons, respectively. Our results allow for a better understanding of the kinetic limiting processes which determine the submonolayer morphology and illustrate the role played by the missing row reconstruction of the Ag(110) surface in the formation of extended Si nanoribbon arrays.

I. INTRODUCTION

Nanostructuring of surfaces using their intrinsic structural properties has received great interest over the last past two decades, with the aim at creating atomically controlled large scale templates. This step of nanopatterning is a promising route to produce by templatedirected growth high densities of true atomic dimension nanostructures with controlled geometries, regular sizes and spacings and specific physical properties related to their structural properties. The method applies for inorganic nanostructures and more recently for molecular arrangements. [1][2][3][4][5][6][7][8] Moreover, the properties of interest of these highly dense and uniform surface-supported nanostructures such as electronic, optical or magnetic properties can be investigated using either local probes (scanning tunneling spectroscopy and spinpolarized scanning tunneling microscopy) or integrating techniques (angle-resolved photoelectron spectroscopy, surface differential reflectance spectroscopy and x-ray magnetic circular dichroism).

Growth proceeds in many cases away from thermodynamic equilibrium. The degree to which the growth occurs far from equilibrium conditions determines if the morphology of the grown structures is rather determined by thermodynamic quantities or by the growth kinetics.

The kinetically controlled growth has proven to be, in the last twenty years, an alternative strategy to produce a wide variety of new zero-dimensional (0D) 1,4 , one-dimensional (1D) 2 and two-dimensional (2D) 9 nanostructures on substrates by playing with the growth conditions and the interplay between the adsorbate and the substrate. The morphology of these metastable nanostructures is the result of the competition between different thermally activated atomic scale processes characterized by their activation energy. Scanning tunneling microscopy (STM) has proven to be a technique of choice to obtain valuable information on the atomic structure of these nanostructures and the activated atomic scale processes taking place on surfaces. 10,11 Among these nanostructures, 2D Si layers grown on silver substrates have recently attracted considerable interest for both their fundamental interest and expected promising properties analog to the electronic properties of graphene. 12,13 It has been mainly reported since 2005 that submonolayer Si deposition on the anisotropic Ag(110) surface leads to the spontaneous formation of flat lying Si nanoribbons (NRs), perfectly aligned with the troughs of the bare (110) surface (see [14][15][16] for the earliest articles). Si NRs with a width of 0.8 nm (denoted hereafter as single nanoribbons (SNRs)) are randomly distributed on silver terraces upon Si deposition at room temperature (RT) while larger Si NRs, 1.6 nm in width (denoted hereafter as double nanoribbons (DNRs)), self-assemble laterally to form regular DNR arrays with a pitch of 2 nm when the sample is heated at ∼ 470 K. Surface diffraction experiments have shown that the SNRs are locally organized in a 3×2 superstructure when grown at ∼ 430 K while a highly-ordered 5×2 superstructure composed of DNRs is formed upon Si deposition at higher temperature (∼ 470 K). 17,18 Despite the expected simplicity of the Si/Ag interface, regarding the negligible miscibility of Si and Ag below ∼ 1070 K, it took about 10 years to solve the atomic structure of these Si NRs. The first step was to evidence using STM and grazing incidence x-ray diffraction (GIXD) that the Ag(110) surface surprisingly undergoes a missing row reconstruction underneath the Si NRs during the growth of both SNRs and DNRs. 18 An original surface-supported Si arrangement was later proposed by Cerdá et al. 19 , based on density functional theory (DFT), rapidly confirmed by Prévot et al. 20 using a combined GIXD-STM-DFT study: the Si NRs correspond to pentamer chains lying in the missing row troughs of the reconstructed Ag(110) surface. Apart from these studies, the Si DNR array, covering uniformly the whole Ag(110) surface, has been successfully used as a template for the growth of identical 1D nanostructures made of Co 21 , Mn 22 , 9,10phenanthrenequinone 23 or C 60 24 and the study of their magnetic properties in the case of Co 25 .

The use of this Si array as a template takes advantage of its high structural quality (low defect density at the atomic scale and single domain orientation of the Si grating) and the electronic decoupling of adsorbates from the conductive substrate, which is crucial to study the magnetic properties of supported nanostructures.

While the atomic structure of the Si NRs has been solved, the growth mechanisms of their arrangement in 3×2 and 5×2 superstructures have received little attention. However, a detailed knowledge of the atomic scale processes involved is required for understanding and controlling growth morphologies in the kinetic growth regime. The control of the formation of the Si superstructures would allow for the nanofabrication of highly-ordered templates with different geometries by varying the width of the Si NRs and the pitch of the grating.

Moreover, questions remain regarding the early stages of growth, especially concerning the role and the atomic structure of the Si nanodots observed by STM, considered as the precursor for the formation of SNRs. 19,26 Here, we present a nanoscale investigation of the Si NR growth on Ag(110), for submonolayer Si coverage and substrate temperature ranging from 300 K to 500 K. The behavior of four types of Si nanostructures have been studied: isolated SNRs and isolated DNRs (denoted hereafter as 1D-SNRs and 1D-DNRs, respectively) and their respective 2D counterparts in the self-assembly regime (denoted hereafter as 2D-SNRs associated with the 3×2 superstructure and 2D-DNRs associated with the 5×2 superstructure, respectively). We focus our study on the statistical analysis of the Si NR density measured in STM images to gain a better understanding of the kinetic limiting processes in the different growth modes observed. Our results strongly suggest that the transition from Si nano-dots to the NR structure is the limiting process for the growth of both 1D-SNRs and 1D-DNRs while the growth of 2D-SNRs and 2D-DNRs is driven by the formation of close-packed silver missing rows.

II. EXPERIMENT

All the experiments were carried out in an ultrahigh vacuum (UHV) set-up. The (110) surface of a silver single-crystal was prepared by repeated cycles of Ar + sputtering and STM measurements were carried out at the CINaM in Marseille, using Scienta Omicron apparatus. All STM images were obtained at RT in the constant current mode with sample voltages from -1.2 V to -0.7 V and tunneling current of ∼300 pA. The X-Y piezoelectric scanner of the STM was carefully calibrated using the extended 5×2 Si DNR superstructure. The STM data have been processed using the Gwyddion software. 27 After Si deposition at substrate temperatures ranging from 300 K to 500 K, the surface was frozen down to RT, where no motion or modification of the Si NRs is observed. The Si NR counting was performed from STM images with typical size (300×300) nm 2 , insuring that SNRs and DNRs can be distinguished and the spacing between Si NRs can be measured.

III. Results

We have investigated the Si growth on Ag(110) for substrate temperatures ranging from 300 K to 500 K and in the submonolayer regime. For these growth conditions, the deposition of Si on Ag(110) leads to the formation of a Si ultrathin layer composed of Si NRs.

Different arrangements of the Si NRs, depending on the growth temperature and for the lower coverage studied, can be viewed in the STM images of Fig. 1. At RT, 1D-SNRs are essentially formed and are randomly distributed on the silver terraces [Fig. 1(a)], as already reported. 14,18 Si nano-dots, which are considered to be the precursor for the formation of SNRs, are also visible. 15,19,26 At 350 K, as illustrated in Fig. 1 chains, by a lateral repetition of the motif of SNRs. 19 We can conclude that no other Si nanostructures are formed in the temperature range of interest. In (a), the arrows indicate the formation of small Si nano-dots.

To get a deeper insight in the growth mechanisms of the Si NRs, we have plotted in Considering that the formation of the four Si species obeys to Arrhenius laws, we have interpreted our STM images by analyzing the density of anisotropic Si islands as a function of the substrate temperature, as frequently reported in the literature for both isotropic 2D islands and anisotropic 1D-2D islands. 10,28 In Figs. This behavior has also been observed by GIXD as illustrated in Fig. 4. In Fig. 4(a), the evolution of the Si NR TLD, expressed in SNRs/nm (i.e. each DNR is counted twice) is reported for θ Si ∼ 0.22 ML and 0.54 ML, as a function of the substrate temperature during Si deposition. Figure 4(b) displays the behavior, as a function of the annealing substrate temperature, of the intensity of three intense GIXD diffraction peaks associated with the ×5 periodicity of the 5×2 superstructure at completion (see 18 for more details concerning the GIXD experiments). This abrupt break at 490 K corresponds to a drastic change of the morphology of the Si/Ag interface, described in 29

IV. DISCUSSION

In our experiments, we clearly observe two transitions from 1D to 2D growth of anisotropic islands: (i) in the middle temperature range, with the formation of 1D-SNRs and 2D-SNRs self-assembled in a 3×2 superstructure above 370 K and (ii) in the higher temperature range, with the formation of 1D-DNRs and 2D-DNRs self-assembled in a 5×2 superstructure above 440 K. Similar transition has been reported in 28 in the growth of anisotropic Cu islands on Pd(110). This transition was attributed to the anisotropy of diffusion along and across the easy [11 0] direction. In our system, the transverse migration is activated above 320 K but it doesn't lead to the formation of 3×2 islands, which is observed above 370 K. The same behavior occurs at higher temperature with the formation of 1D-DNRs above 370 K and 5×2 islands above 440 K. This clearly demonstrates that the formation of the Si species on Ag(110) is driven by the activation barriers for their formation rather than by the diffusion of Si adatoms. From our results, we can easily rank from the lowest to the highest the activation energy for the formation of 1D-SNRs, 2D-SNRs, 1D-DNRs and 2D-DNRs.

Two models have been proposed for the structure of the Si nano-dot considered as the precursor of Si SNRs: one has a quasi-hexagonal geometry and the other corresponds to two Si pentagons sitting side by side, both structures are located in a Ag di-vacancy. 19,26 In both articles, a transition from unstable cluster precursors to SNRs is mentioned, when more Ag atoms are removed to form a trough. Let us remind that a DNR corresponds to not-linked twin Si SNRs running along a silver missing row. STM observations, not shown here, performed after Si growth at 490 K, when a highly perfect 5×2 superstructure is formed, reveal that when a DNR has nucleated, the twin Si SNRs grow simultaneously as the termination of one DNR corresponds to the terminations of both twin SNRs. This strongly suggests that double Si nano-dots can be considered as the precursor structure for the formation of DNRs. We can thus assess that the transition from Si nano-dot to NR structure is the limiting process for the growth of both 1D-SNRs and 1D-DNRs, with a higher activation barrier for the formation of double Si nano-dots. Single Si nano-dots, precursor structure of the Si SNRs, are frequently observed by STM, for instance after growth at temperatures slightly higher than RT. To our knowledge, observation of double Si nano-dots has never been reported. Their observation by STM is rather challenging as they statistically constitute very rare events for Si growth at low temperature and would require a fast cooling of the sample after growth at high temperature to quench any atomic scale processes occurring on the surface.

One can wonder why a transition from 1D to 2D NR growth doesn't occur once the transverse diffusion of Si adatoms is activated. One argument can be given if we consider the 3×2 and 5×2 superstructures, which have a silver missing row TLD of 1/3 and 2/5, respectively and remark that the denser 2×2 superstructure made of pentamer chains with a silver missing row TLD of 1/2 is not formed. This indicates that the formation of closepacked silver missing rows is a limited process for the formation of self-assembled Si pentamer chains. Between 370 K and 440 K, we observe a competition between the formation of local 3×2 arrangements composed of ∼ 3 SNRs and 1D-DNRs. This occurs when the number of close-packed silver missing rows of the local 3×2 arrangement induces an activation barrier higher than the one for the formation of 1D-DNRs. We can note that this transition occurs when the number of Ag missing rows of the 3×2 structure is higher than the one of 1D-DNRs, with similar spacing. This could explain why no extended 3×2

superstructures can be formed on Ag(110).

V. CONCLUSION

We have reported a nanoscale study of the Si NR growth which allows for a better description and understanding of the different growth regimes in the submonolayer range. Our results underline the role played by the silver substrate in the growth mechanisms occurring upon Si deposition. This highlights the strong interaction between Si and Ag in the formation of the surface-supported Si NRs, already reported for their structural 20,30 , electronic 31 and optical properties 32,33 . We show that, due to the competition between the growth of 2D-SNRs and 1D-DNRs, no extended self-assembled SNR array can be formed, in contrary to the case of self-assembled DNRs. The former array has thus a limited reliability for its use as a 1D template, in contrary to the latter one. In the temperature range where only one type of Si arrangement is observed, the activation barrier for the formation of 1D-SNRs and 2D-DNRs are extracted from the measurement of Si NR densities as a function of the substrate temperature. A detailed analysis of these densities using mean-field nucleation theory and kinetic Monte Carlo simulations would provide a profound understanding of the atomic scale processes involved and information about the precursor structure for the formation of Si NRs, which is still in debate.

1

  Nanoscale investigation of Si nanoribbon growth on Ag(110)Michel Daher Mansour, Romain Parret and Laurence Masson a) Aix Marseille Univ, CNRS, CINaM, Campus de Luminy, Case 913, 13288 Marseille Cedex 9, France a)

  annealing at 770 K. The cleanness of the Ag(110) surface was systematically checked by STM before Si deposition. Si was evaporated from a direct current heated piece of silicon wafer kept at 1520 K onto the clean Ag(110) surface, after preoutgassing at 1070 K for several hours. The clean Ag(110) sample was exposed to a Si flux at a constant deposition rate of 5×10 -4 monolayer (ML) per second. 1 ML corresponds to the Ag(110) surface atom density, which also equals the density of the perfect 3×2 superstructure completely covering the Ag(110) surface. We focused our experiments on two low Si coverages (θ Si ∼ 0.22 and 0.54 ML) to avoid coalescence of Si NR islands. The reproducibility of the deposition of this home-made Si evaporator is ∼ + 10 %. The substrate was heated at the rear side by radiation emitted by a commercial Scienta Omicron heater made of a W filament. The substrate temperature (T) during Si deposition was measured using a chromel-alumel thermocouple, spot welded very close to the Ag sample and calibrated prior to the set of experiments using a heating power-substrate temperature relationship established with the thermocouple spot welded instead of the Ag sample. The accuracy of the temperature measurement is + 10 K.

  For each temperature and each Si coverage, at least 2 STM images have been used to perform a statistical analysis of the Si NRs. The typical number of Si NRs in one STM image, counted along the direction perpendicular to the Si NRs, is ∼ 70 SNRs at low temperature and ∼ 35 DNRs at high temperature, for θ Si ∼ 0.22 ML. For each STM image, several measurements of the percentage of the different kinds of Si NRs formed and of the transverse linear density of Si NRs along the direction perpendicular to the NRs, have been performed using line profiles, except for the higher growth temperatures (T > 420 K), where the length of the Si NRs was longer than the size of STM images (300 nm) and thus only one measurement was possible. The error bar of the Si NR density is + 10%, corresponding to the mean statistical variation observed, for a given temperature, between the different measurements made. The lattice parameters of Ag(110) is 0.409 nm in the [001] direction, perpendicular to the Si NR axis and 0.289 nm in the [11 0] direction, parallel to the Si NR axis. GIXD experiments were carried out on the ID3 beamline of the European Synchrotron Radiation Facility (ESRF) storage ring with 17 keV x-rays and the incidence angle kept fixed at 0.22°. The basis is expressed in the cubic basis of Ag, = (0,0,1), = (1/2, -1/ 2,0) and = (1/2,1/2,0), where a = 0.409 nm. The corresponding (h,k,l) indices are used for indexing a reflection in reciprocal space.

  Fig. 1(c), 2D-SNRs are locally organized in a 3×2 superstructure while 1D-DNRs start to

FIG. 1 .

 1 FIG. 1. (Color Online) STM images ((70×70) nm 2 ) of Si deposited on Ag(110) at four different substrate temperatures and nearly the same Si coverage θ Si ∼ 0.22 ML, showing the

Fig. 2 (

 2 Fig. 2(b), as a function of the substrate temperature during Si deposition, the percentage of the

FIG. 2 .

 2 FIG. 2. (Color Online) (a) Transverse linear density (TLD) of Si NRs, perpendicular to the

  FIG. 3. (Color Online) (a) Arrhenius plot of the transverse linear density (TLD) of Si SNR islands (blue filled dots) and Si DNR islands (red open dots), in nm -1 , and the surface density of Si NRs (black dots), in nm -2 , for θ Si ∼ 0.22 ML as a function of 1/T. (b) Arrhenius plot of

  FIG. 4. (Color Online) (a) : transverse linear density (TLD) of Si NRs, expressed in SNRs/nm, as a function of the substrate temperature during Si deposition for θ Si ∼ 0.22 ML and 0.54

Finally, the presence

  in STM images of SNRs evolving in DNRs (or vice versa) during the growth of Si NRs, as observed in Fig.1(a), 1(b) or 1(c), has to be discussed. These features could result from the coalescence of growing SNRs and DNRs along the[11 0] direction, parallel to the NRs. Another explanation could be that a DNR evolves in a SNR due to the activation barrier for the formation of 2 close-packed silver missing rows. Lastly, let us consider that a SNR could evolve in a DNR. One mechanism could be that a single Si nanodot nucleates in the close vicinity of a SNR and leads to the growth of a twin SNR, forming thus a DNR. This mechanism is not consistent with the formation of double Si nano-dots, precursor for the formation of DNRs. Our STM observations made after Si growth at 490 K related to the termination of DNRs and the formation of 2D-SNRs in a large temperature range are in favor of the formation of double Si nano-dots. Nevertheless, we cannot exclude the mechanism of nucleation of single Si nano-dots in the close vicinity of SNRs to explain the formation of DNRs.

ACKNOWLEDGMENTS

The authors gratefully acknowledge Geoffroy Prévot and Romain Bernard from the INSP (Paris) for fruitful discussions and collaboration during the GIXD run performed on the ID3 beamline at ESRF (Grenoble). Support from the ID3 beamline staff is acknowledged.