High-temperature alkali corrosion of Kyocera SN282 silicon nitride

Corrosion Science 91 (2015) 68–74

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High-temperature alkali corrosion of Kyocera SN282 silicon nitride Madeleine Kant Jordache ⇑, Henry Du 1 Department of Chemical Engineering and Materials Science, Stevens Institute of Technology, 1 Castle Point Terrace, Hoboken, NJ 07030, USA

a r t i c l e

i n f o

Article history: Received 29 June 2014 Accepted 30 October 2014 Available online 6 November 2014 Keywords: A. Ceramic A. Glass A. Rare earth elements C. Alkaline corrosion C. Oxidation C. High temperature corrosion

a b s t r a c t The role of Cs in corrosion of SN282 Si3N4 with 5.35 wt.% Lu2O3 was examined at 1200 °C for up to 20 h in oxygen with 220 ppm CsNO3 to determine particulars of SN282 corrosion under Cs, and high-temperature Cs+ behavior in glass silicate hosting 137Cs. Cs accelerates SN282 oxidation. The pronounced adverse effect of Cs on Si3N4 corrosion resistance is predicted thermodynamically more strongly for SiO2 dissolution in Cs silicate than in other alkali silicates. While oxidation kinetics remained parabolic, Cs oxide layers were thick, uneven, heterogeneous and cracked. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction Silicon nitride (Si3N4) ceramics are proposed to replace super alloys for aerospace and industrial gas turbines, due to their being better in terms of both high temperature stability and corrosion resistance. Corrosion of Si3N4 structural components can contribute to markedly deteriorate ceramic mechanical properties, such as reduction of strength [1,2] and aggravation of creep [3]. Structural materials are frequently exposed to alkali corrosive condensed phases and gases. There have been a large number of studies on corrosion of Si3N4 by molten alkali salts [4–8]. To study alkali corrosion of Si3N4, other investigators introduced the alkali element either through ion implantation into the ceramic substrate [9] or as an alkali salt vapor mixed in the oxidation gas [10]. In the presence of alkali species, the amorphous silica layer grown on Si3N4, which has a role in slowing down further oxygen diffusion, would become less protective since alkali ions would disrupt the glass network connectivity by introducing non bridging oxygen ions [11]. These alkali generated non-bridging oxygen ions would render the glass network a more open structure to oxygen permeation, which would lead to an accelerated oxidation process dependent on the ceramic alkali corrosion resistance. Alkali corrosion of Si3N4 is a problem that needs to be addressed and diminished for high temperature structural applications of this ceramic.

Our investigation of oxidation of lately developed turbine – grade Kyocera SN 282 Si3N4 sintered with 5.35 wt.% Lu2O3 has been examined in an earlier paper [12] that reports on the specific ratelimiting mechanism in SN282 oxidation in dry oxygen. The present work investigates oxidation behavior of SN282 Si3N4 in an environment of O2 contaminated with Cs. Cesium alkali species were introduced via vapor phase. Investigation of surface chemical and phase distribution as well as characteristics of corrosion products for the particular corrosion condition was correlated with ceramic corrosion resistance. This work presents morphological development of oxide layers grown on SN282 silicon nitride in oxygen contaminated with Cs alkali. Investigation on oxidation behavior of SN282 when environment is contaminated with cesium not only offers basic information on particulars of SN282 corrosion under a large size alkali ion; this baseline research also helps to study the mixed alkali effect (MAE) – upon addition of a dissimilar alkali – in improving the alkali corrosion resistance of SN282 Si3N4, which makes the subject of another study. The amorphous cesium silicate grown on SN282 Si3N4 upon high-temperature oxidation in the present work is a system that is useful to study in order to determine information on glassy silicates as the base glass to incorporate radioactive 137Cs from high-level nuclear waste (HLW). 2. Materials and Experiments

⇑ Corresponding author at: 124 East 84th Str. Apt. 6B, New York, NY 10028, USA. Tel.: +1 (212) 988 2852; cell: +1 (646) 644 0156. E-mail addresses: [email protected] (M.K. Jordache), henry.du@ stevens.edu (H. Du). 1 Tel.: + 1 (201) 216 5262. http://dx.doi.org/10.1016/j.corsci.2014.10.050 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.

2.1. Materials The silicon nitride ceramic used in this investigation was Kyocera SN 282 b-Si3N4 with 5.35 wt.% Lu2O3 sintering aid by Kyocera

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2.2. Corrosion conditions Cesium corrosion of SN282 was investigated at 1200 °C for a duration of up to 20 h in 101.325 kPa dry oxygen flowing at 500 SCCM mixed with 220 ppm cesium nitrate vapor. The continuous supply of alkali vapor was ensured by evaporation of the alkali salt throughout the entire duration of corrosion experiments. The weight loss of the 4-alumina crucibles loaded with CsNO3 powder was measured after each run to make sure that alkali concentration in O2 in the oxidation environment was the same from one run to another. The laboratory set-up for corrosion experiments was a horizontal mullite tube furnace with two independently controlled temperature regions: a high temperature region (CM Furnace, Bloomfield, NJ) – where SN 282 samples were corroded – and a low temperature region (Marshall Furnace Co. Inc., Scotts Valley, CA) – where cesium nitrate salt (of 99.99% purity, Aldrich, ACS Reagent) was evaporated. Upon evaporation, the alkali salt vapors mixed into the flowing oxygen and were carried by the flowing gas to the high temperature region containing SN282 test samples. The evaporation temperature was 655 °C and 4 alumina crucibles were used with alkali salt. An alkali concentration of 220 ppm CsNO3 in O2 was calculated under the assumption that the alkali vapors behave as ideal gases and using the weight loss rate of CsNO3 and the flow rate of the oxygen gas.

Oxidation of Kyocera SN 282 Si3N4 takes place about three times faster in 220 ppm CsNO3–O2 gas mixture than in oxygen ambient at 1200 °C [12]. Notably thicker oxides grew on SN 282 when CsNO3 was mixed in O2 environment than in solely O2 ambient at this temperature. Signs of SN282 oxidative degradation under alkali contamination are apparent through oxide cracking, uneven oxide thickness and multiple phase separation. 3.1. Oxidation rate Shown in Fig. 1 is the increase of oxide thickness function of time for oxidation of SN 282 at 1200 °C in O2 with 220 ppm CsNO3. The trends of oxide growth of SN 282 in 220 ppm CsNO3–O2 mixture follow a parabolic relation with time indicating a diffusion controlled oxidation mechanism. Cesium contamination in O2 environment accelerates SN 282 oxidation compared to oxidation in pure oxygen at the same temperature and duration. At 1200 °C, the SN282 oxide growth rate is about three times faster than in pure oxygen, as shown in Table 1 by the ratio of rate constants in the two different ambients, oxygen with 220 ppm CsNO3, and pure oxygen: KpSN282 in CsNO3–O2/KpSN282 in O2 = 2.99–3. 3.2. Oxide layer morphology Fig. 2(a)–(c) presents morphology of SN 282 oxidized in a 220 ppm CsNO3–O2 gas mixture at 1200 °C for 20 h. For 1 0.9

O2 w/ 220 ppm CsNO3 O2

0.8 0.7 0.6

x 2 (μm 2)

Industrial Ceramics Corporation. SN 282 was processed through gas–pressure liquid phase sintering followed by heat treatment [13]. The vendor indicated the next impurity species present in the bulk ceramic: 0.058 wt.% Yb2O3, 0.045 wt.% Al2O3, 0.053 wt.% Na2O, 0.016 wt.% CaO, and 0.002 wt.% MgO. Ceramic samples were polished to a root mean square roughness of about 10 nm through tribochemical polishing in 3 wt.% chromic acid solution, followed by ultrasonically cleaning in deionized water, acetone, and methanol to remove surface contaminants prior to corrosion exposure. Samples of silicon single crystal were cut from a silicon wafer that was (1 1 1) surface oriented, and were also oxidized simultaneously with Kyocera as witness samples for oxidation/corrosion. The phase characteristics of as-received (and of corroded) samples were examined using a Siemens D 5000 X-ray Diffractometer with Cu Ka radiation source.

0.5 0.4 0.3

2.3. Microstructural investigation

0.2 3

The as-corroded Si3N4 coupons (6  3  2 mm ) were analyzed by XRD for phase identification. The corroded layers microstructure was examined at Scanning Electron Microscope (SEM), lateral chemical elemental mapping was pursued through Energy Dispersive Spectrometry (EDS) analyses on the corroded surface, and thickness of the corroded layer was measured through SEM cross-section imaging, with at least 9–10 measurements per sample. A Leo 982 field-emission-gun SEM, Thornwood, NY equipped with EDS was used for morphological and compositional analyses as well as cross-section oxide thickness measurement. 3. Results and discussion High temperature oxidation of SN282 in 220 ppm CsNO3–O2 gas mixture has shown a decrease in ceramic oxidation resistance considering both the found oxidation rate value and the oxide morphology standpoints compared to ceramic oxidation resistance in solely oxygen environment. Oxidation rate, morphology evolution and phase changes indicate acceleration of oxidation, morphological degradation and multiple phase formation in the oxide for SN 282 oxidized under a cesium contaminated ambient when compared to ceramic oxidation in only dry oxygen, free of alkali.

1200°C

0.1 0

0

5

10

15

20

time (hr) Fig. 1. Oxide layer thickness as a function of time for SN 282 after oxidation at 1200 °C in 220 ppm CsNO3–O2 gas mixture and oxide layer thickness as a function of time for SN282 oxidized in clean O2 (lower plot) previously reported in another publication [12].

Table 1 Parabolic rate constant (lm2/hr) for oxidation of SN 282 in 220 ppm CsNO3–O2 mixture and respectively in O2 at 1200 °C. Ambient

Parabolic rate constant (lm2/hr)

220 ppm CsNO3–O2 mixture O2

(4.18 ± 0.49)  102a (1.4 ± 0.015)  102b

a Errors in parabolic rate constant are generated by the variation of SN282 oxide thickness that is larger in SN282 oxidized in O2 mixed with CsNO3 than in SN282 oxidized in O2. b Parabolic rate constant for oxidation of SN282 in pure O2 at 1200 °C was reported in a previous paper [12].

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silica

Cs silicate

Lu2Si2O7

Lu2Si2O7

(e)

silica

relative intensity (arbitrary units)

70

(a) β -Si3N4 Lu2Si2O7 Lu2SiO5

20

25

30

35

40

0.95 μm

50

55

60

65

70

Lu2Si2O7

0.52 μm

relative intensity (arbitrary units)

Lu2Si2O7

45

two theta

Cs silicate

(b)

20

Fig. 2. SEM micrographs of SN 282 oxidized at 1200 °C for 20 h in 220 ppm CsNO3– O2 gas mixture (a) and (b) plan-view and (c) cross-section and in O2 gas (d) and (e) plan-view and (f) cross-section [12].

25

β-Si3N4 Lu2Si2O7 Lu2SiO5 cristobalite

30

35

40

45

50

55

60

65

70

two theta

comparison, Fig. 2(d)–(f) shows morphology of SN 282 oxidized in dry O2 in the same conditions of temperature and duration of oxidation as reported in a previous paper [12]. SEM investigation of surface morphology of SN 282 oxidized in 220 ppm CsNO3–O2 revealed marked changes compared to SN 282 oxidized in dry O2 ambient. In the two different environments, alkali-contaminated oxygen and pure oxygen, SN 282 oxidized layers are cesium silicate and silica, respectively, in which discrete Lu2Si2O7 grains (identified based on XRD and EDS) are embedded. Fig. 2(b) and (e) indicate an increased size and a lower density of Lu2Si2O7 particles in oxide when SN 282 was oxidized in O2 with CsNO3 compared to SN 282 oxidized in dry oxygen. SEM cross-section images of SN 282 oxides grown in the two different environments, alkali contaminated and free of alkali, in Fig. 2(c) and (f) indicate a notably thicker oxide, of large thickness variation, on SN 282 oxidized in O2–CsNO3 mixture compared to a thin, smooth, oxide of even thickness grown on SN282 in O2. Through cross-section examination, the SN 282 oxides grown in O2 with Cs show no subsurface damage, neither cracking, nor subsurface oxidation. SN 282 oxide grown under cesium contamination presents signs of oxide degradation through cracking, visible on SN282 oxide surface in Fig. 2(b). Lu2Si2O7 grains are visible to a lesser extent, through oxide crosssection, raised above the oxide surface on SN282 oxidized under cesium contamination, compared to their more elevated appearance above the oxide surface on SN 282 oxidized in clean oxygen, as illustrated in Fig. 2(c) and (f). 3.3. Phase assessment Fig. 3(a) presents X-ray diffraction pattern of as-received Kyocera SN282 as being a b-Si3N4 ceramic with largely crystallized phases in grain boundaries as a major Lu2Si2O7 phase and reduced amounts of Lu2SiO5. Fig. 3(b) shows X-ray diffraction spectrum of SN 282 oxidized at 1200 °C for 20 h in dry O2 as previously reported [12]. Fig. 4 illustrates the X-ray pattern of SN 282 oxidized

relative intensity (arbitrary units)

Fig. 3. XRD pattern of Kyocera SN 282 Si3N4 (a) as-received (b) oxidized at 1200 °C for 20 h in dry O2.

20

25

30

35

40

45

50

55

60

65

70

two theta Fig. 4. XRD pattern of SN 282 oxidized at 1200 °C for 20 h in 220 ppm CsNO3–O2 gas mixture.

in same conditions of temperature and duration, however, in an O2 ambient contaminated with alkali as 220 ppm CsNO3. X-ray pattern indicates increased crystallization in the oxidized SN 282 when ceramic oxidation took place in O2 containing Cs compared to SN282 oxidized in pure O2. SN282 oxidized in 220 ppm CsNO3– O2 shows X-ray spectral peaks of Lu2Si2O7 and Lu2SiO5, the silicates formed in SN282 oxide from Lu sintering additive. These Lu silicates are present in the oxide layers of both SN282 oxidized in O2 with CsNO3 and in SN 282 oxidized in pure O2, and in either of these oxides, Lu2Si2O7 quantitatively dominates over Lu2SiO5. However, the ratio Lu2Si2O7/Lu2SiO5 appears larger in SN282 oxidized in CsNO3–O2 than in SN282 oxidized in dry oxygen and in SN 282 as-received. The height of the Lu2SiO5 peak of maximum intensity at 2h = 31.126° is considerably diminished relative to the height of the Lu2Si2O7 peak of maximum intensity at

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2h = 27.946° in SN 282 oxidized in oxygen with Cs, compared to their relative peak heights in SN282 oxidized in dry oxygen. Overall, the X-ray spectrum of SN282 oxidized in CsNO3–O2 shows tall, strong spectral lines of Lu2Si2O7 that are dominant over reduced, low-intensity Lu2SiO5 spectral lines. In the X-ray spectrum of SN 282 oxidized in a CsNO3–O2 mixture, Lu2Si2O7 lines appear to be the strongest after the spectral lines of SN282 ceramic substrate. In SN282 oxidized in pure O2, the peaks of Lu2Si2O7, although taller than those of Lu2SiO5, are still of comparable intensity with the later ones, the peaks of both phases being much lower than the substrate peaks. The comparative X-ray patterns of SN282 oxidized in CsNO3–O2 in Fig. 4, SN 282 oxidized in dry O2 in Fig. 3 (b) and asreceived SN282 in Fig. 3(a), indicate that the strong Lu2Si2O7 lines in XRD pattern of SN 282 oxidized in CsNO3–O2 are due to a very large amount of Lu2Si2O7 formed in SN282 oxide when the ceramic was oxidized in an environment of O2 contaminated with Cs. In other words, the amount of Lu2Si2O7 in the oxide grown under cesium contamination appears to far exceed the amount of Lu2Si2O7 in SN282 oxide grown in O2, which at its turn surpasses the quantity of Lu2Si2O7 crystallized in as-received ceramic. Regarding crystalline silica phases, cristobalite was found in both SN282 oxides grown on SN282 oxidized in the two different environments: in CsNO3–O2 mixture, and in dry oxygen. However, only SN282 oxidized in cesium contaminated ambient shows crystallization of tridymite. 3.4. Lateral elemental distribution Fig. 5 illustrates SEM micrograph (a) and corresponding EDS spectra (b) and (c) of regions of oxide for SN 282 oxidized at 1200 °C for 20 h in a 220 ppm CsNO3–O2 gas mixture. For the dark oxide area between bright particles presented in SEM picture in Fig. 5(a), the EDS spectrum shown in Fig. 5(b) indicates oxygen and silicon as major peaks, corresponding to silica in the oxide layer, and minor peaks for Lu and Cs that show formation of Lu silicate and Cs silicate in SN 282 oxide. For lightly colored particles visible on the SN282 oxide surface in the SEM picture from Fig. 5(a), the EDS from Fig. 5(c) indicates strong Lu, O and Si spectral lines, which are the peaks expected for the discrete Lu silicate phases nucleated and grown in the oxidized layer. Pairing information furnished through XRD and EDS, it can be assessed that the lightly colored particles in the SN 282 oxidized layer are predominantly Lu2Si2O7 crystalline phases and a reduced amount of crystalline Lu2SiO5 phases. 3.5. Effect of cesium in oxidation of SN282 The effect of cesium ions in the oxidation environment on oxidation resistance of SN282 was investigated from the oxidation rate, morphological, structural and chemical distribution points of view. A parabolic rate law was found for oxidation of Kyocera SN 282 in O2 mixed with 220 ppm CsNO3, which is a similar oxidation type rate law like for oxidation of SN282 in clean oxygen, previously found [12]. Direct comparison of the two parabolic rate constants found for oxidation of SN282 in O2 with Cs and respectively in clean oxygen only, shows an about three times faster oxidation rate of SN282 in O2 contaminated with Cs than in pure oxygen ambient, indicating that Cs contamination accelerates SN282 high temperature oxidation. This kinetic result, as well as oxide morphology changes, like cracking of oxide surface and an increase in the number of crystalline phases that formed in SN282 oxide grown in O2 mixed with Cs salt, compared to the SN282 oxide, dense, stable, of largely amorphous silica, with reduced crystallinity grown in clean O2, show that the oxidation resistance SN 282 exhibited in clean O2 is lost when the ceramic is exposed to oxidation in O2 contaminated with

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Cs. As result of acceleration of SN282 oxidation in Cs contaminated ambient, the oxide layer appears in cross-section considerably larger and of variable thickness that did not resemble anymore the thin, smooth, and uniform oxide grown on SN282 in clean oxygen. Cesium silicate formed upon incorporation of cesium ions into the silica layer appeared to disrupt the stability of the growing oxide layer, of silica with enclosed discrete lutetium disilicate phases, as was established during oxidation of SN282 in clean oxygen [12]. At the level of Cs concentration in O2 ambient used in present study, cesium silicate formed on SN282 upon oxidation clearly contributed to decrease the number of Lu2Si2O7 phases apparent in the oxidized layer, while markedly increasing their sizes. Cesium silicate formed upon oxidation on SN282 appeared to be very reactive, contributing to dissolve Lu2Si2O7 particles of small diameter, while accelerating the growth of larger ones, as seen in Fig. 2(b) as compared to Fig. 2(e). XRD and EDS indicate increased amounts of Lu2Si2O7 with comparatively more diminished Lu2SiO5 phases in SN282 oxide grown in O2 containing cesium than in SN282 oxidized in clean O2. 3.5.1. Thermodynamic aspects for oxidation of Si3N4 in O2 with Cs The detrimental effect of cesium on ceramic oxidation resistance, the alteration in morphology and increased crystallinity in the SN 282 oxide layer can be interpreted according to the following discussion. Cs corrosion takes place during oxidation of Si3N4 in an oxygen environment contaminated with Cs, due to dissolution of the silica layer upon Cs incorporation. Possible oxidation and dissolution reactions include:

Si3 N4 þ 3O2 ! 3SiO2 þ 2N2

ð1Þ

xSiO2 þ Cs2 O ! Cs2 O  xSiO2

ð2Þ

Cs is incorporated into amorphous silica as the stable Cs2O oxide supplied steadily during oxidation of Si3N4 upon continuous evaporation of CsNO3 salt, which is subsequently mixed in flowing oxygen. Thermodynamic analysis by the use of HSC Program (HSC CHEMISTRY Ver. 4.0 Copyright (C) Outokumpu Research Oy, Pori, Finland, A. Roine) indicated that CsNO3 vapor is not thermodynamically predicted to decompose in the oxidation environment from room temperature (RT) to 1200 °C, since DG0T , the value of the standard Gibbs free energy change for the chemical reaction at temperature T is positive, where T has values between RT and 1200 °C. The HSC program shows decomposition reactions of CsNO3–O2 mixture to provide the stable Cs2O are to happen upon reaching Si3N4. Upon decomposition of cesium nitrate salt, Cs2O is incorporated into the growing silica layer on SN 282. Cs2O dissolves the SiO2 layer on Si3N4, forming a cesium silicate. Possible oxidation/dissolution reactions with formation of stable cesium silicates from Cs2O– SiO2 system in SN 282 oxide include:

2Si3 N4 þ 25=2O2 þ 6CsNO3 ¼ 3Cs2 Si2 O5 þ 14NO2

DG01200 C ¼ 2; 551:597 kJ=mol 4Si3 N4 þ 9=2O2 þ 6CsNO3 ¼ 3Cs2 Si4 O9 þ 11N2 DG01200 C ¼ 7; 742:194 kJ=mol

ð3Þ ð4Þ

The pertinent thermochemical data for the systems presented in Eqs. (3) and (4) are given in the next references to the standard data: (a) Si3N4, CsNO3, Cs2Si2O5 and Cs2Si4O9 data in Knacke 91, Ref. [14], (b) O2 (g) and NO2 (g) thermodynamic data in Frenkel 94, Ref. [15], used by HSC to calculate thermochemical data, and (c) N2 (g) data in JANAF 85, Ref. [16]. Both reactions (3) and (4) for cesium silicates formation are predicted thermodynamically, with a large negative value of the standard Gibbs free energy change DG01200 C for the reaction. However, the strongest predicted thermodynamically among the two reactions is the last reaction

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(b)

Intensity (a.u.)

Si

O

Lu

(a)

Cs

0

1

2

3

4

5

Energy (keV)

Si

(c)

Lu

Intensity (a.u.)

5 μm O

N

0

1

2

3

4

5

Energy (keV) Fig. 5. SN 282 oxidized at 1200 °C for 20 h in 220 ppm CsNO3–O2 gas mixture (a) SEM micrograph, (b) EDS of region of dark contrast of the oxide between lightly colored particles and (c) EDS of particles of light contrast of oxide in (a).

(4) that shows formation of Cs2Si4O9. Its negative value of the standard Gibbs free energy change for the reaction DG01200 C ¼ 7; 742:194 kJ=mol is about 3 times larger than the negative value of the standard Gibbs free energy change for the reaction of formation of Cs2Si2O5. In the present study, at the level of CsNO3 concentration in O2, the Cs silicate formed on SN282 upon oxidation in O2 with CsNO3 appears to be an amorphous phase; if any amount of Cs2Si4O9 formed in the oxidized layer, it should be probably under the X-ray Diffractometer detectable limit. It is to note that the negative value of the standard Gibbs free energy change DG0T ; where T value is between RT and 1200 °C for the reaction of Cs2O fluxing SiO2 appears to be significantly larger than the negative value of the standard Gibbs free energy change DG0T ; where T has values between RT and 1200 °C for the reaction of other alkali oxides, of smaller alkali cations fluxing SiO2, indicating a much stronger thermodynamic prediction for the fluxing reaction of SiO2 by the oxide of the large Cs+ ion than the fluxing reaction of SiO2 by alkali oxides of smaller alkali cation. The overall strong fluxing effect of Cs2O on SiO2 was indicated experimentally by SN 282 oxide morphology evolution when ceramic oxidation was pursued in O2 + 220 ppm CsNO3. The magnitude of the fluxing effect was shown not only through a significantly

increased oxide thickness of SN282 oxidized in CsNO3–O2, Fig. 2(c), than for SN282 oxidized in only clean oxygen, Fig. 2(f), but also, by degradation of oxide morphology, as illustrated in Fig. 2(b), through cracking due to cristobalite b to a transformation during cooling, indicating that a large amount of cristobalite was formed in the abundant, thick cesium silicate formed on SN282 in cesium contaminated ambient. As a result of the strong fluxing effect of Cs2O on SiO2, the transport of gaseous oxygen through the low viscosity cesium silicate was accelerated, with adverse effect on ceramic oxidation resistance, which was decreased due to the diminished protection against oxidation that the low viscosity cesium silicate offered to the ceramic substrate, compared to that of the amorphous silica grown on SN282 in clean oxygen as it appears in Fig. 2(a), (b) and (d), (e). The thick Cs corrosion layers grew unevenly, showing surface cracks and numerous crystalline phases. Cs silicate formation in the oxidized layer on SN 282 leads to a decrease in viscosity of the silica network associated with weakening of the glass network as a result of the known effect of breaking bridging oxygen ions (BO) by the alkali ions (Cs+) introduced in silica with formation of non-bridging oxygen ions (NBO). Bonds between NBO and the Cs+ modifier (Si–O–Cs) are ionic bonds that

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are weaker than the strong covalent bonds between BO and the network former Si4+ (Si–O–Si); these comparatively weak bonds between NBO and the modifier cation are not part of the tetrahedral network and their formation leads to decreasing glass network viscosity [17]. Cs+ is a network modifier and NBO introduced by Cs+ ions decrease alkali silicate viscosity. In the present study, physical evidence that results from decreased Cs silicate viscosity is the very thick SN282 oxide layer grown in CsNO3–O2, as compared to the thin amorphous SN282 silicon oxide grown in clean O2, since it is known that O2 permeates faster in a less viscous glass, i.e. in this case, increased O2 permeation probably took place through a low viscosity cesium silicate than through amorphous silicon oxide. In the amorphous silica network formed on SN282 upon oxidation in O2 only cristobalite formation was noticed, since is usually the first crystalline SiO2 phase formed due to a lower energy path because of an increased similarity of cristobalite structure with that of amorphous silica than the structure of tridymite with glass silica [18]. The defective, reactive Cs silicate network formed on SN282 created conditions for phase transformation to tridymite, the less protective form of SiO2 that is the stable crystalline SiO2 at 1200 °C; the presence of broken bonds in cesium silicate allowed displacement of the structural units, enhancing crystallization of tridymite nuclei and their growth. For alkali silicates, the modified random network (MRN) structural model with the modifiers clustered to form channels into the glass network [19] was advanced gradually, in agreement with experimental results, to replace the continuous random network model (CRN) microstructure in which alkali were randomly distributed in the silica glass. Alkali diffusivity in alkali silicates follow an Arrhenius type relation with absolute temperature with the same activation energy like the ionic conductivity of alkali silicate glasses, so that the activation energy Ea for alkali diffusion is modeled by the Anderson– Stuart relation:

Ea ¼ Eb þ Es

ð5Þ

where Eb is the electrostatic binding energy of alkali ion and Es is the strain energy to enlarge the glass network for the alkali ion to pass through [20]. According to the MRN oxide glass model the alkali are microsegregated in silica, which creates pathways, or alkali channels, in the microstructure along which alkali ions migrate [19]. Diffusion coefficients and activation energies for ionic transport determined experimentally [19] agree with the model of alkali channels in silicates and with alkali migration along these microsegregated pathways. A Raman spectroscopy study on the structure of rubidium and cesium silicate glasses [21] has calculated the relative concentration of Qi species function of composition, where Qi represents the Si environments in alkali silicate glasses and i is the number of BO per silicon; Q4: Si(OSi)4, Q3: Si(OSi)3 O R+ (R–Li, Na, Cs), Q2: Si(OSi)2(O)2(); (0 6 i 6 4). They found that increasing the alkali cation’s radius from Na < K < Rb < Cs will result in shifting to the left the equilibrium of reaction:

2Q 3 $ Q 4 þ Q 2

ð6Þ

That is, in Cs silicates, compared to silicates of lower alkali radius, Cs introduction will strongly shift the equilibrium of the above reaction to the left, toward producing more Q3 units (and respectively diminishing Q4 units); Q3 will position themselves alongside alkali conduction channels, participating so directly toward formation and developing of the Cs conduction channels, resulting in a decrease of the activation energy for Cs diffusion, which is governed by the least favorable jump of the Cs+ ion, which takes place between Cs

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conduction channels. Such a decreased activation energy would favor increased diffusion/incorporation of Cs+ into silica glass with the effect of increasing oxygen permeation through the more open structure of the Cs silicate formed. Cs introduction into silica shifts the equilibrium of the above reaction to the left, increasing the production of Q3 (with respectively diminishing Q4), which perhaps is a factor contributing to the presently noticed increased thickness of SN282 oxidized layer under Cs contamination. Multiple phases formed in SN 282 oxidized in O2 + 220 ppm CsNO3; the crystalline SiO2 phases, cristobalite and tridymite, were both found through XRD in the SN 282 oxidized layer, as shown in Fig. 4. Cesium silicate is probably present as an amorphous phase; if beginning of crystallization of Cs2Si4O9 took place, the amount of this crystalline silicate is under the X-ray Diffractometer detectable limit. Cesium incorporation in silica network led to a defective and reactive network through which diffusion of Lu3+ ions was increased, contributing to faster growth of Lu2Si2O7 grains in the oxidized layer; Lu2Si2O7 particles that were larger than critical nucleus size grew larger than the size of Lu2Si2O7 particles in the SN 282 oxidized layer in dry O2, Fig. 2(a), (d) and (b), (e). In the reactive cesium silicate, of low viscosity, through which O2 diffusion was accelerated, Lu2SiO5 was faster oxidized to Lu2Si2O7, so that the ratio Lu2Si2O7/Lu2SiO5 was larger than it was in silica on SN 282 oxidized in dry O2. An increased reactivity of cesium silicate as compared to other alkali silicates of lower alkali cation radius can be interpreted according to the next discussion. Weakening the silica network structure, i.e. the Si–O bonds with incorporation of the network-modifying alkalis leads to physical and chemical changes that result in an increased chemical reactivity in general [22]; a stronger fluxing effect of Cs2O than of other alkali oxides, of lower radius cation, on SiO2 leads to an accelerated gaseous transport in cesium silicate than in other alkali silicates, of smaller size modifier that translates into largely increased and of uneven thickness oxidized layers produced, and on account of the increased mesh size cesium silicate formed, larger than of other alkali silicates of smaller cation size, the presence of broken bonds eased more the displacement of structural units with an effect on increasing nucleation and growth of multiple crystalline phases. The kinetics result of parabolic reaction characteristics for oxidation of SN282 in CsNO3–O2 suggests that the oxide layer on SN 282 preserved a degree of protectiveness of the silica layer against the oxidation process. Despite the increased rate of ceramic oxidation under cesium contaminant, and the thicker, heterogeneous, with more abundant crystallization of phases SN 282 oxide grown in CsNO3–O2 than SN282 oxide formed in O2, oxidation of SN 282 in O2 containing 220 ppm CsNO3 at 1200 °C remains parabolic, i.e. it did not turn into a linear oxidation law that would basically mean the respective oxide totally lost its protective character.

4. Conclusions Oxidation of SN282 in O2 contaminated with 220 ppm CsNO3, despite following a parabolic oxidation law (indicating diffusion controlled oxidation mechanism), like a rate type law followed by the ceramic for oxidation in dry O2, shows significant differences from the ceramic oxidation in a pure oxygen environment. While cesium accelerates the oxidation rate of SN282 due to the fluxing effect of Cs2O on SiO2, a well-known fluxing effect of alkali oxides on silica, it appears that the thermodynamic prediction that goes for the fluxing of silica reaction by Cs2O is stronger than that of the fluxing effect on silica for other alkali oxides of smaller alkali cations. The fluxing reaction of Cs2O on SiO2 generates a low viscosity cesium silicate that allows for an increase in oxygen

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transport. Cesium induces an advanced devitrification of the amorphous SN282 silicon oxide layer that appears to be favored by the increased mesh size of Cs silicate formed over those of other smaller size alkali silicates. SN 282 oxide grown in O2 mixed with a Cs salt vapor showed significant morphological instability in the form of surface cracking and uneven thickness in oxide cross-section. The very reactive cesium silicate formed on SN282 changed the distribution pattern of Lu2Si2O7 phases in the SN 282 oxidized layer, as was established under the dry oxygen ambient [12], by dissolving the small Lu2Si2O7 grains, while accelerating the growth of the relatively large Lu2Si2O7 grains. Acknowledgement This work was funded by the National Science Foundation (NSF) under Grant DMR-0102340. References [1] D.S. Fox, M.D. Cuy, T.E. Strangman, Sea salt corrosion and strength of an yttriacontaining silicon nitride, J. Am. Ceram. Soc. 80 (1997) 2798–2804. [2] J.J. Swab, G.L. Leatherman, Effects of hot corrosion on the room temperature strength of structural ceramics. , 1989. [3] R.E. Tressler, Environmental effects on long term reliability of SiC and Si3N4 ceramics, in: R.E. Tressler, M. McNallen (Eds.), Corrosion and Corrosive Degradation of Ceramics, The American Ceramic Society, Westerville, OH, 1989, pp. 99–124. [4] N.S. Jacobson, Corrosion of silicon-based ceramics in combustion environments, J. Am. Ceram. Soc. 76 (1993) 3–28. [5] W.L. Fielder, Oxidation and Hot Corrosion of Hot-pressed Si3N4 at 1000 °C, NASA Technical Memorandum, 1985. pp. 86977. [6] N.S. Jacobson, Sodium sulfate: deposition and dissolution of silica, Oxid. Met. 31 (1989) 91–103.

[7] N.S. Jacobson, D.S. Fox, Molten-salt corrosion of silicon nitride: II, sodium sulfate, J. Am. Ceram. Soc. 71 (1988) 139–148. [8] Y. Shinata, M. Hara, T. Nakagawa, C. Shimizu, Hot corrosion of reactionsintered Si3N4 in Molten Na2SO4, in: Y. Saito, B. Onay, T. Maruyama (Eds.), High Temperature Corrosion of Advanced Materials and Protective Coatings, Elsevier Science Publishers, B.V., 1992. [9] Z. Zheng, R.E. Tressler, K.E. Spear, A comparison of the oxidation of sodiumimplanted CVD Si3N4 with the oxidation of sodium-implanted SiC crystals, Corros. Sci. 33 (1992) 569–580. [10] G.R. Pickrell, T. Sun, J.J. Brown Jr., High temperature alkali corrosion of SiC and Si3N4, Fuel Process. Technol. 44 (1995) 213–236. [11] W.D. Kingery, H.K. Bowen, D.R. Uhlmann, Introduction to Ceramics, second ed., John Wiley & Sons, Singapore, New York, Chichester, Brisbane, Toronto, 1991. [12] M.K. Jordache, H. Du, Oxidation behavior of silicon nitride sintered with Lu2O3 additive, J. Mater. Sci. 41 (2006) 7040–7044. [13] H.-T. Lin et al., Evaluation of mechanical reliability of silicon nitride vanes after field tests in an industrial gas turbine, in: Proceedings of ASME Turbo Expo 4A, 2002, pp. 147–154. [14] O. Knacke, O. Kubaschewski, K. Hesselman, Thermochemical Properties of Inorganic Substances, second ed., Springer-Verlag, Berlin, 1991. [15] M. Frenkel, G.J. Kabo, K.N. Marsh, G.N. Roganov, R.C. Wilhoit, Thermodynamics of organic compounds in the gas state, 1994. [16] M.W. Chase et al. (Eds.), JANAF Thermochemical Tables, J. Phys. Chem. Ref., third ed., vol. 14, Hemisphere Publishing Corp., 1985. [17] J.F. Stebbins, Z. Xu, NMR evidence for excess non-bridging oxygen in an aluminosilicate glass, Nature 390 (1997). [18] J.R. Blachere, F.S. Pettit, High Temperature Corrosion of Ceramics, Publisher Noyes Data Corporation Mill Road, Park Ridge, New Jersey, 1989. 07656. [19] G.N. Greaves, W. Smith, E. Giulotto, E. Pantos, Local structure, microstructure and glass properties, J. Non-Cryst. Solids 222 (1997) 13–24. [20] O.L. Anderson, D.A. Stuart, Calculation of activation energy of ionic conductivity in silica glasses by classical methods, J. Am. Ceram. Soc. 37 (1954) 573–580. [21] V.N. Bykov, A.A. Osipov, V.N. Anfilogov, Structural study of rubidium and caesium silicate glasses by Raman spectroscopy, Eur. J. Glass Sci. Technol. Part B 41 (2000) 10–11. 2. [22] M.A. Lamkin, F.L. Riley, R.J. Fordham, Oxygen mobility in silicon dioxide and silicate glasses: a review, J. Eur. Ceram. Soc. 10 (1992) 347–367.