Oxidation of a Fe–18Cr–8Ni austenitic stainless steel at 850 °C in O2: Microstructure evolution during breakaway oxidation

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Oxidation of a Fe–18Cr–8Ni austenitic stainless steel at 850 ◦ C in O2 : Microstructure evolution during breakaway oxidation Audrey Col, Valérie Parry, Céline Pascal ∗ Univ. Grenoble Alpes, CNRS, SIMaP, F-38000 Grenoble, France

a r t i c l e

i n f o

Article history: Received 7 July 2016 Received in revised form 8 October 2016 Accepted 16 October 2016 Available online xxx Keywords: A. Stainless steel B. Raman spectroscopy B. SEM B. TEM C. Internal oxidation C. High temperature corrosion

a b s t r a c t Oxidation tests of AISI 304L in breakaway conditions (850 ◦ C in O2 ) were performed up to 312 h. The evolution of the oxidation affected zone microstructure was investigated using a combination of composional/elemental (TEM, Raman spectroscopy) and structural (EBSD) mapping techniques as well as thermodynamic calculations. The formation of a dense and continuous Cr2 O3 healing layer at the border of the internal oxidation zone happens along the grain boundaries and is linked to their more efficient Cr supply. The propagation of the oxidation front is related to a local conversion of Cr2 O3 in less protective FeCr2 O4 . © 2016 Elsevier Ltd. All rights reserved.

1. Introduction Austenitic stainless steels, such as AISI 304L, are protected from rapid oxidation by the selective consumption of chromium which leads to the formation of a protective chromia scale, Cr2 O3 , characterised by a slow growth rate. Numerous investigations related to chromia forming alloys report the formation of a two-layer oxide scale composed of a dense and adherent inner sub-layer of Cr2 O3 and an outer sub-layer of spinel type oxide MnCr2 O4 [1–4]. Chromia is a barrier against further oxidation owing to its very low diffusion coefficients for oxygen and metals, and then offer a high oxidation resistance [5]. The formation of protective Cr-rich oxide scale, Cr2 O3 , depends on the chromium concentration at the alloy/oxide interface and also on the chromium diffusion coefficient in the metallic substrate from the core towards the interface. Analytical solutions of the Wagner’s equations have been used to model the Cr depletion profiles [6,7]. According to Evans et al. experiments [8], about 16 wt.% in Cr is required at the alloy/scale interface to form a healing layer in the temperature range from 750 to 900 ◦ C for a Fe–20Cr–25Ni steel. However, the steel grain size [2,9–14], the silicon content in the alloy [12,15–17] and the surrounding atmosphere [18] influence the Cr critical value for a given temperature. Decreasing the

∗ Corresponding author. E-mail address: [email protected] (C. Pascal).

alloy grain size allows a quicker formation of a continuous chromia layer due to a faster effective diffusion of Cr in the alloy and to a smaller lateral growth distance required to form a continuous layer. Presence of silicon in the nominal composition of the steel induces the formation of silica at the alloy/oxide interface that acts as an additional diffusion barrier. In water vapour [19–22] or carbon dioxide [23,24] containing atmospheres, the growth rate of chromia scale is increased causing a greater depletion of Cr in the subsurface zone. In both cases, the subsequent loss of chromium tends to convert the protective Cr-rich oxide, initially formed, into a poorly protective Fe-rich fast growing oxide [25,26]. Moreover, part of Cr can be either trapped in M23 C6 or M7 C3 carbides in carburizing atmospheres [27,28], or lost through Cr2 O3 volatilization in gaseous CrO2 (OH)2 (g) in water vapour containing atmospheres [29,30]. According to Huntz et al. [3], for austenitic stainless steels, 18 wt.% of Cr is not sufficient to sustain the formation of chromia when temperature (or time) increases because lattice diffusion coefficients in face-centered cubic crystal structure (fcc-austenite) are smaller than the one in body-centered cubic crystal structure of ferritic stainless steel (bcc-ferrite). At high temperature (>850 ◦ C) in high oxygen partial pressure (O2 or air), even if a thin chromia scale forms in the early stage of oxidation, the chromium concentration and diffusivity cannot feed the further growth of Cr2 O3 . Then, the nodular growth of iron oxides (spinel-like phase and hematite) takes place which induces a sudden increase of the oxidation rate. The phenomenon is usually called “breakaway oxidation”.

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In various experimental conditions, breakaway oxidation of austenitic stainless steels gives rise to the inward growth of a complex oxide composed of a (Ni,Fe)-rich phase and of a spinel type oxides: in carburizing atmospheres [23,24,26,27,31], in water vapour containing atmospheres [25,30,32–37] and in air at higher temperature [4]. In all cases, depletion of chromium in the subsurface alloy zone is the key factor leading to breakaway oxidation. Depending on the experimental conditions, Ni-rich phase trapped in FeCr2 O4 spinel oxide may be metallic or oxidised: (Ni,Fe,Cr)3 O4 spinel oxide [25,26,36], FeNi3 [31], Ni–Fe solid solution [4,23,33,35,37–39]. Some authors also report formation of a Cr-rich oxide layer at the interface with the substrate which can be Cr2 O3 [23,30,40] or FeCr2 O4 [35,37]. The objective of this work is to study the evolution of the AISI 304L microstructure after breakaway oxidation in oxygen at 850 ◦ C, to characterise the observed microconstituents and to describe the breakaway process in such experimental conditions. In particular, Raman spectral mapping allows to evidence an alteration of the composition from Cr2 O3 to FeCr2 O4 linked to a loss of protectiveness of the Cr-rich oxide scale formed at the interface with the substrate.

sections. Observations were performed using a Field Emission Gun (FEG) SEM Zeiss Ultra 55 equipped with a X-Ray (SSD Bruker) detector and an Electron Back Scattered Diffraction (EBSD, TSL-EDAX) detector. Identification of oxide composition was carried out using Raman spectral mapping. Due to the spatial resolution limits of the SEM and Raman spectroscopy techniques, elemental chemical analysis of the finest microconstituents was performed with transmission electron microscopy (TEM). The TEM thin lamella was prepared using focused ion beam thinning (SEM-FIB Zeiss Nvision 40) and observed using a TEM-FEG Jeol 2100F equipped with EDS analysis. The chemical composition in the underlying substrate was analysed using electron probe microanalysis (EPMA, Cameca SX50). 3. Results 3.1. Evolution of microstructure and composition of the outer oxides scale SEM surface views in SE mode of the AISI 304L specimens oxidized at 850 ◦ C in O2 for 48–312 h are displayed in Fig. 1. For the shorter oxidation times (48 h, 72 h and 96 h), the oxide surface is smooth with submicronic polygonal grains. Local thickenings occur after 96 h of oxidation. For oxidation time longer than 110 h, oxide morphology is strongly modified with the appearance of coarse angular crystals. Formation of necks between grains happens leading to a porous sintered-like microstructure of the base oxide. After 312 h at 850 ◦ C, the oxide scale is rough with formation of nodules and vertical growth of oxide grains. For oxidation time longer than 110 h, the oxide scale is thick and porous, sign of a rapid oxide growth. The thermal stresses, originated from the mismatch of the thermal expansion coefficients between the oxide and the alloy, associated with a low cohesive energy of the porous oxide result in oxide spallation during cooling. Surface Raman spectra of the samples oxidized for 72 h and 312 h at 850 ◦ C in O2 are displayed in Fig. 2. For each samples, ten spectra were randomly acquired. Raman peaks were identified using McCarthy and Boehme [41] and Hosterman spectra as references [42]. According to Fig. 2a, for 90% of spectra, the outer oxide scale after 72 h at 850 ◦ C is composed of corundum-structure solid solution Fe2-x Crx O3 with varying composition up to Fe-rich one. Small amounts of chromia Cr2 O3 and manganese chromite MnCr2 O4 are also detected (10% of spectra in Fig. 2b). Since a partial spallation of the oxide scale occurs for the longest oxidation time, two types of spectra were recorded for the sample oxidized for 312 h. The outer oxide scale, related to an unspalled area in Fig. 2c, is mainly composed of hematite Fe2 O3 with a small amount of Fe-rich corundum-type solid solution Fe2-x Crx O3 . According to Fig. 2d, spalled areas contain Fe–Cr spinel-type oxide which indicates that spallation happens within the oxide scale rather than at the alloy/oxide interface. These results show that the oxide scale is duplex with formation of hematite on the top surface and of Fe–Cr spinel-type oxide as inner scale, as typically reported in literature [1,3,4,43,44]. SEM outer surface study associated to Raman spectroscopy results allow to describe the chemical and morphological evolution of the outer surface of the oxide scale during oxidation of AISI 304L at 850 ◦ C in O2 . Up to 96 h, the oxide scale is made of regular platlet grains of corundum-structure solid solution Fe2-x Crx O3 . The iron content

2. Material and methods The commercial 1 mm-thick foil of AISI 304L (EN 1.4307) used in this work was supplied by Goodfellow. The chemical composition obtained by fluorescence spectroscopy analysis and optical emission spectrometry equipped with a gas analyser is reported in Table 1. Specimens were cut to 10 × 20 × 1 mm3 dimensions, then ground up to 1200 grit with SiC abrasive papers. After measurement of dimensions, the coupons were cleaned in acetone and ethanol. Oxidation experiments were performed in a tubular furnace at 850 ◦ C (heating rate 13 ◦ C/min) at atmospheric pressure in O2 gas flow (flow speed of 21 cm/min at room temperature) for 48–312 h. The samples were placed in several ceramic boats to take them out after a given time. Short-time oxidized samples (48 h, 72 h and 96 h) were quenched under a continuous flow of O2 in the cold extremity of the furnace. Long-time oxidized samples (110 h, 120 h, 148 h and 312 h) were naturally cooled down in the furnace under O2 flow in order to limit the spallation of their thicker oxide scales due to the mismatch of thermal expansion coefficients between the oxide scale and the metallic substrate. After oxidation, sample cross-sections were prepared for microstructural investigations. Samples were coated using epoxy resin to prevent oxide scale damages during metallographic preparation (cutting, grinding and polishing). The sample cross-sections were ground up to 1200 grit SiC paper and polished up to 1 ␮m diamond paste. A first set of two techniques was used to characterize of the oxide phases that are present on the outer surface of the oxide scale. Raman spectroscopy was performed on the oxidized sample surface in order to determine the chemical composition of the oxide phase (spatial resolution ∼1 ␮m2 with depth resolution ∼1 ␮m). For each sample, ten Raman spectra were randomly acquired with a Renishaw RM1000 spectrometer using a 514.532 nm Ar-laser. The morphology of the outer oxide scale was investigated using scanning electron microscopy (SEM, LEO S440). A second set of techniques was carried out to study the complex chemistry and morphology of the oxidation affected zone on crossTable 1 Chemical composition of austenitic stainless steel AISI 304L (in wt.%). wt.%

Fe

Ni

Cr

Mn

Si

Co

Cu

Mo

C

S

AISI 304L

Bal.

8.143

17.462

1.724

0.319

0.203

0.355

0.262

0.021

0.002

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Fig. 1. SEM-SE micrographs of the outer surface of oxide scale formed on AISI 304L oxidized at 850 ◦ C in O2 flow for (a) 48 h, (b) 72 h, (c) 96 h, (d) 110 h, (e) 120 h, (f) 148 h and (g) 312 h.

Fig. 2. Raman spectra obtained at the outer surface of the oxide scale grown on AISI 304L substrate after oxidation at 850 ◦ C in O2 flow for (a,b) 72 h and (c,d) 312 h.

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Fig. 3. SEM-BSE cross-section micrographs after oxidation of AISI 304L at 850 ◦ C in O2 flow for (a) 48 h, (b) 72 h, (c) 96 h, (d) 110 h, (e) 148 h, (f) 312 h and (g) 312 h (rectangle indicates location of image f).

in the solid solution oxide increases with oxidation time leading to a full conversion of the surface scale in porous hematite above 110 h at 850 ◦ C which is characteristic of breakaway oxidation. Large facetted crystals may be related to manganese chromite, MnCr2 O4 , which is known to nucleate from pre-existing corundum-type oxide [45] owing to the fast diffusion of Mn in chromia [46]. 3.2. Evolution of microstructure and composition of the bulk oxidation affected zone Fig. 3 shows SEM cross-section micrographs in BSE mode of AISI 304L samples oxidized at 850 ◦ C in O2 for 48–312 h. Up to 96 h of oxidation, the oxidation affected zone (refered as OAZ) is composed of a regular oxide scale with a uniform thickness, less than 5 ␮m, as well as intergranular oxides. According to EDX maps (not presented), oxidation patterns for the shortest oxidation times (up to 48 h) correspond to a well-described typical microstructure. The oxide scale is duplex with an outer (Mn,Cr)3 O4 spinel oxide and inner chromium-iron solid solution oxide (Cr,Fe)2 O3 [1–4,29,45,47–50]. A discontinous silica layer forms at the alloy/oxide interface with intrusions along the substrate grain boundaries as commonly observed for austenitic steels due to the slow lattice diffusion in fcc-matrix as compared to the grain boundaries diffusion of species [10,47,48,51]. For oxidation longer than 96 h, local thickening occurs which indicates that oxide nodules with homogeneous microstructure start to grow. After 110 h of oxidation, the OAZ reveals complex chemistry and microstructure. Since spallation of the oxide scale occurs not only during the cooling step but also during metallographic preparation, only the internal part of the OAZ

can be observed in Fig. 3d–g. The sample oxidized for 110 h displays at once both oxidation patterns, i.e. a regular oxide scale, homogeneous nodules and deep oxidation zones. This sample, reproduced three times, always gives rise to the same microstructure. In agreement with the chemical and morphological study of the outer surface oxides, significative evolution of the oxide scale occurs after oxidation for ∼96–110 h. In order to highlight the specific mechanisms of AISI 304L breakaway oxidation in O2 at 850 ◦ C, attention will be paid to the study of the inner oxidation affected zones obtained after oxidation for 110 h (restrained breakaway oxidation) and 312 h (massive breakaway oxidation). Figs. 4–6 present SEM-BSE cross-section micrographs of the oxidized sample and EDS maps of the OAZ observed after 110 h at 850 ◦ C in O2 . Three types of microconstituent are observed: regular oxide scale (Fig. 4), fully-oxydized nodules with homogeneous microstructure (Fig. 5) and deep zones of internal oxidation (Fig. 6). As shown in Fig. 6, the metallic matrix can be seen among oxide in the inner OAZ suggesting that the oxidation is internal. This zone is refered as internal oxidation zone (IOZ). According to literature [1–4,29,45,47–50] and to the EDS maps in Fig. 4, the regular oxide scale exhibits successive sublayers from alloy/oxide interface to the outer oxide surface: SiO2 /Cr2 O3 /(Cr,Fe)2 O3 /MnCr2 O4 . Electron probe microanalysis results (not presented here) indicate that the underneath substrate composition is modified up to 25 ␮m deep with a Cr depletion associated to Ni and Fe enrichments (13.0 at.% Ni near the metal oxide interface instead of 7.4 at.% in the bulk and 77.7 at.% Fe near the interface instead of 71.1 at.% in the bulk). Fully-oxidized nodules, shown in Fig. 5, reveal an homogeneous microstructure and are composed of an inner Cr-rich part and an

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Fig. 4. SEM-BSE cross-section micrographs and EDX element mapping after oxidation of AISI 304L at 850 ◦ C for 110 h: uniform-thick scale.

Fig. 5. SEM-BSE cross-section micrographs and EDX element mapping after oxidation of AISI 304L at 850 ◦ C for 110 h: fully-oxidized nodule.

outer Fe-rich part. The inner part of the nodules is slightly smaller than the corresponding top part and lateral cracks are observed within the nodule at the level of the former alloy/oxide interface. Nodules seem to be located between intergranular silica intrusions. Deep pits of IOZ are occasionally located below the fullyoxidized nodules (Fig. 6). Grey levels variations in the BSE micrograph and EDX mapping indicate that the pits are composed of Cr-rich oxides and cavities embedded in a Ni-rich metallic matrix depleted in Fe and Cr. Intergranular oxidation occurs with the formation of discontinuous Cr-rich oxide layers (denoted by an arrow in Fig. 6) with a penetration depth of the order of the oxidation propagation front. BSE-SEM, EDX and Raman spectral maps of the sample crosssection after 312 h at 850 ◦ C in O2 are displayed in Fig. 7. The

inner OAZ pattern, characterized by intermixing of cavities, oxide grains and Ni-rich phases, has common features with the pits of IOZ observed after 110 h (Fig. 6). However, it differs from IOZ pits after 110 h because of the formation of a dense and continuous Cr-rich healing layer (about 5 ␮m thick) which underlines the OAZ following a geometrical path. In order to investigate the fine microstructure and chemistry of the OAZ core, a thin TEM lamella extracted using FIB milling was studied. Fig. 8 presents a TEM micrograph and the associated EDX element mapping for O, Si, Ni, Cr, Mn and Fe. The oxide phase presents a large variation of composition with complementary decrease or increase content of metallic elements Cr and Fe: if one is low, the second is high. Submicronic in size Si containing inclusions, SiO2 and possibly fayalite Fe2 SiO4 , are also detected. Metallic phases, rich in Ni and poor in Fe and Cr, embedded inside the oxide network and cavities are also observed. According to Raman spectral map in Fig. 7, the composition of the continuous Cr-rich layer depends on its location with respect to the oxidation front. Corundum-type oxide, chromia Cr2 O3 , is detected where the oxidation front is stopped, while spinel-type oxide, iron chromite FeCr2 O4 , is detected where the oxidation front is extended beyond this dense Cr-rich oxide layer. From the EBSD crystallographic orientation map of the substrate at the vicinity of alloy/oxide interface, displayed in Fig. 9, the location of the substrate grain boundaries is clearly evidenced in relation with the “geometrical pattern” of the continuous Cr-rich oxide layer. Therefore, the growth of the dense Cr-rich oxide layer takes place along the grain boundaries of the substrate. According to electron probe microanalysis results (not presented here), the Cr-depleted zone below the dense Cr2 O3 layer, is thin (4–6 ␮m deep). 4. Discussion The resistance of a chromia-forming steel toward the breakdown of the protective chromia scale is determined by three main factors: (i) the concentration of chromium in the steel, (ii) the lattice diffusion of Cr in the steel and (iii) the grain size of the metallic substrate. In the present experimental conditions, the two first conditions are not favourable: (i) the steel Cr content (17.5 wt.%)

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Fig. 6. SEM-BSE cross-section micrographs and EDX element mapping after oxidation of AISI 304L at 850 ◦ C for 110 h: internal oxidation zone.

Fig. 7. SEM-BSE cross-section micrographs, SEM-EDX element mapping and Raman spectral mapping after oxidation of AISI 304L at 850 ◦ C for 312 h.

is insufficient to sustain the formation and growth of a fully protective chromia scale and (ii) the lattice diffusion is known to be lower for fcc-austenitic steel compared to bcc-ferritic steel. (iii) Although the substrate grain size ranges from 10 to 20 ␮m (not presented), the subsurface region of the alloy is finer grained as a result of recrystallization consecutive to surface cold working during sample preparation (grinding) prior to oxidation tests. The Cr supply to the metal/oxide interface is expected to be enhanced through Cr diffusion via dislocation pathways and numerous grain boundaries in the early stage of oxidation [31]. Nevertheless, at such high temperature, the effect is time-limited since grain growth occurs gradually and results in an increase in the average grain size (e.g. about 40–80 ␮m after 110 h at 850 ◦ C in the present study) [52]. The oxidation proceeds through an outward Cr cationic diffusion, the underlying steel supplies Cr to the protective oxide scale more rapidly at the steel grain boundaries than at the steel grain

centres due to slow lattice diffusion compared to grain boundary diffusion. At the centre of the underneath alloy grains, the chromium activity becomes insufficient and then chromia layer locally loses its protective feature as described for example in Refs. [26,34,53]. Local chemical failure results in fast outward iron diffusion and inward oxygen diffusion leading to the formation of two-layer fully-oxidized nodules (observed in Fig. 5). The present study focuses on the mechanisms responsible for the propagation of the inward oxidation front and the resulting microstructure rather than the formation of the fully-oxidized nodules already described in the literature. According to Raman spectral mapping in Fig. 7, the propagation of the oxidation front depends on the composition of the healing layer. Cr2 O3 is detected where the oxidation front is stopped while FeCr2 O4 is detected where the oxidation front is extended beyond this layer. In order to discuss the conversion from protective Cr2 O3 to non-protective FeCr2 O4 , thermodynamic calculations were per-

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Fig. 8. Bright field TEM image and TEM-EDX element mapping for O, Si, Ni, Cr, Mn, Fe elements in the oxidation affected zone (OAZ) after 312 h at 850 ◦ C.

Fig. 9. SEM-SE cross-section micrographs and EBSD orientation mapping after oxidation of AISI 304L at 850 ◦ C for 312 h.

formed using Thermo-Calc software with TCFE7 database [54,55]. Results are displayed in Fig. 10. According to Fig. 10a, once the Cr2 O3 is formed, the oxygen partial pressure at the metal/oxide interface decreases to 10−26 bar since oxygen partial pressure is then controlled by the decomposition pressure of Cr2 O3 which is very low. So, the formation of chromia decreases the content of oxygen dissolved in the metal. Elements such as Si and Mn, forming more stable oxides, can be oxidized selectively in the vicinity of Cr2 O3 oxide as SiO2 and MnO respectively. Indeed, the expected MnO sublayer beneath chromia scale is not observed for several reasons [56]: (i) Mn–Cr spinel oxide, MnCr2 O4 , is stable with respect to binary oxides at high temperature, (ii) Mn is soluble in chromia and (iii) Mn diffuses rapidly in Cr2 O3 but relatively slowly in the alloy. MnCr2 O4 spinel oxide overlaying chromia is evidenced after oxidation at 850 ◦ C (Raman spectra in Fig. 2b and EDX-map of Mn in Fig. 4). Concomitant oxidation of SiO2 and Cr2 O3 is observed in the case of external (EDX-SEM

map of Si in Fig. 4) and internal (EDX-SEM of Si map in Fig. 7) Cr2 O3 layers. The reaction of chromia formation is written as follows: 4/3 Cr + O2 (g) = 2/3 Cr2 O3 K =

1 4/3

acr × aO2

(1)

Where K is the equilibrium constant and a is the activity (activity of chromia is taken as unity). Since oxidation progresses, the Cr content (i.e. Cr activity) at the subsurface alloy begins to decrease. According the law of mass action, Eq. (1), the oxygen activity must increase to maintain the equilibrium condition. When the activity of oxygen reaches a value at which reaction with other elements becomes thermodynamically possible, 10−22 from Fig. 10a, the oxide is converted to a less protective Fe–Cr spinel oxide according to the following reaction: 2Fe + O2 (g) + 2Cr2 O3 = 2FeCr2 O4

(2)

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Fig. 10. Properties diagrams versus oxygen activity in gas for AISI 304L at 850 ◦ C calculated using Thermo-Calc with TCFE7database [54,55]: (a) Volume fraction of phases versus oxygen activity in gas, (b) Mole fraction of all elements in the spinel type structure (M3 O4 ) versus oxygen activity, (c) Mole fraction of all elements in the corundum type structure (M2 O3 ) versus oxygen activity.

If the oxygen involved in Eq. (2) is obtained only from the dissociation of chromia from Eq. (1), the overall reaction initiating the chemical failure can be obtained combining Eq. (1) and Eq. (2): Fe + 4/3Cr2 O3 = 2/3Cr + FeCr2 O4

(3)

Then, when Cr supply becomes insufficient to support the continuous growth of Cr-rich oxide scale, chromia becomes thermodynamically unstable with respect to the alloy composition (i.e.

low chromium activity) at the alloy/oxide interface and chemical failure occurs according to Eq. (3) [40]. From TEM-EDX mapping in Fig. 8, the OAZ core is characterized by intermixing of cavities, oxide grains (rich in Cr, Fe and Mn) and metallic phases (rich in Ni and depleted in Fe and Cr). From Raman spectral mapping (Fig. 7), the oxide in the OAZ core has a spinel structure. This oxide is probably chromite (Fe,Mn)(Cr,Fe)2 O4 . Since Cr2 O3 is more stable than the spinel oxide FeCr2 O4 (Fig. 10), local thermodynamic equilibrium would prevent internal spinel formation underneath the remaining primary Cr2 O3 layer. The inward oxygen diffusion to form the internal spinel oxide occurred only beneath the nodule centre, where the Cr-rich layer was no longer protective, and where the oxygen activity locally rises. The oxygen activity at the new metal/oxide interface increases because oxygen partial pressure is henceforth controlled by the dissociation pressure of FeCr2 O4 which is higher than the Cr2 O3 one (Fig. 10a). The alloy permeability to oxygen increases, outward diffusion of Cr becomes insufficient and internal oxidation of the surrounding Cr-depleted and Ni-enriched alloy as internal spinel oxide occurs (IOZ observed in Fig. 6). According to thermodynamic diagram in Fig. 10b, the partial pressure of oxygen below the Fe–Cr spinel oxide is insufficient to oxidize the remaining Ni-rich alloy. The remaining Ni–Fe islands acts as a source of metallic species for the outward growing oxide and are transformed into coarse pores as observed in OAZ in Fig. 8. Nevertheless, internal oxidation of Ni containing spinel oxide (Fe,Cr,Ni)3 O4 was already evidenced, for example in Refs [50,57], associated to stratified oxide scale probably related to a mechanically induced chemical failure [43,44]. Mechanically induced failure of the oxide scale occurs when the protective chromia layer spalls or cracks, allowing gas access (i.e. higher oxygen partial pressure) to the underlying Cr-depleted and Ni-enriched alloy. In the present study, as iron has faster diffusion than chromium in spinel lattice [58], the outward iron flux to form the outer oxide scale is enhanced and lead to the formation of hematite Fe2 O3 . Moreover, the interfaces network in the inner OAZ, oxide/metal or oxide/oxide, result in fast transport paths. Removal of iron from the underneath alloy enriches the remaining alloy in nickel and allows the formation of cavities observed in Figs. 6–8. According to the above results, the evolutions of the morphology and the composition of the OAZ with oxidation time are clear and a growth mechanism can be discussed. Figs. 11 and 12 are schematic representations of the nodular growth, of the IOZ formation and of the Cr-rich healing layer formation. In the early stage of fully-oxidized nodule growth, the counter flows of iron and oxygen are located at grain centre of the underlying substrate where the primary chromia scale has lost its protective properties (Fig. 11, step 2). Once the internal spinel oxide, FeCr2 O4 , has nucleated, it grows, thickens and spreads laterally from the grain centre towards the Cr depletion zone produced during the growth of the primary Cr-rich scale. Ahead of the slowly inward oxidation front, the chromium depleted zone is enhanced in the surrounding alloy since internal oxide growth requires Cr incorporation to form FeCr2 O4 . Chromium must be quickly supplied to the advancing oxidation front, otherwise this Cr-depleted zone is susceptible to rapid internal oxidation. When the Cr supply becomes insufficient, FeFe2-x Crx O4 is formed. Since this structure is less ordered, it results in the increase of diffusion coefficient in spinel [38] and then to an acceleration of the oxidation front. The inner OAZ formation is a rapid phenomenon since the underlying chromium depletion profile measured using EPMA is roughly unaffected. Removal of iron to the outer oxide and of chromium to the internal oxide leads to the formation of cavities and produces a significant enrichment in nickel of the remaining alloy since the partial pressure of oxygen below the Fe–Cr spinel oxide is insufficient to oxidize nickel. When the inwards oxidation front encounters a grain boundary

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Fig. 11. Schematic presentation of the nodular growth illustrated with SEM micrographs and EDX maps performed on AISI 304L oxidized for 110 h at 850 ◦ C in O2 .

Fig. 12. Schematic presentation of the oxidation affected zone (OAZ) with a dense Cr-rich layer illustrated using SEM micrographs, EDX maps and EBSD maps performed on AISI 304L oxidized for 110, 148 and 312 h at 850 ◦ C in O2 .

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of the alloy (e.g., vertical light grey layer denoted by an arrow in Fig. 6), Cr supply from the metal to the oxide scale is sufficient to maintain Cr2 O3 growth (Fig. 12, step 3). According to literature, the transition from internal to external oxidation depended on the dissolved oxygen, the bulk Cr concentration and its diffusion coefficient in the metallic matrix [59–61]. When the Cr concentration at the metal/oxide interface dropped below the Cr critical concentration to form a continuous layer of chromia, the oxide scale along the grain boundaries is gradually enriched in Fe and convert into Fe-Cr spinel oxide (Fig. 12, step 4). Then, internal oxidation of Cr and Fe takes place in the adjacent metal grain, leaving the Ni-rich metallic islands behind, up to encounter a new surface defects (grain boundaries, twin boundaries. . .) where Cr diffusion is enhanced, or up to a depth at which the local chromium content is sufficient to form a new protective chromia scale. This mechanism repeated several times leads to a banded inward-grown scale with alternated Cr-rich dense oxide scale (FeCr2 O4 or Cr2 O3 depending on the Cr content in the surrounding alloy) and IOZ (Fe–Cr spinel oxide and Ni–Fe metallic phase). 5. Conclusions The microstructural evolution during oxidation of AISI 304L austenitic stainless steel was investigated at 850 ◦ C up to 312 h in O2 flow using a combination of composional/elemental (TEM, Raman spectroscopy) and structural (EBSD) mapping techniques. The underneath steel supplies chromium required for chromia formation more rapidly at the steel grain boundaries than at the grain centres. After about 100 h at 850 ◦ C, for a few subsurface grains, the chromium supply at the centre of the grains becomes insufficient to support the continuously growing chromia scale. Chromia is then gradually enriched in Fe and eventually locally converted in less protective FeCr2 O4 spinel. It results in fast outward cationic diffusion of iron and inward anionic diffusion of oxygen leading to the formation of a two-layer oxide nodule separated by the initial metal/oxide interface (i.e. an outward growing Fe-rich oxide and an inward growing FeCr2 O4 spinel oxide). Depending on the composition of the surrounding alloy, an internal oxidation zone (IOZ), composed of Fe–Ni metallic phases and Fe–Cr spinel oxides, appears deeper in the alloy. The IOZ thickens, spreads laterally and eventually meets a grain boundary where Cr is present and easily supplied. A continuous chromia dense layer is then formed decorating the grain boundaries of the alloys and bordering the OAZ. The oxidation front is temporarily interrupted until this layer is gradually converted in spinel oxide FeCr2 O4 following the same mechanism as the primary chromia scale. Then, the inwards oxidation front progresses until a new surface defect is encountered where Cr supply is enhanced. It gives rise to a banded inward-grown microstructure with alternated Cr-rich dense oxide scale (FeCr2 O4 or Cr2 O3 depending on the Cr content in the surrounding alloy) and IOZ (Fe–Cr spinel oxide and Ni–Fe metallic phase). References [1] F. Riffard, H. Buscail, E. Caudron, R. Cueff, C. Issartel, S. Perrier, Yttrium implantation effect on 304L stainless steel high temperature oxidation at 1000 ◦ C, J. Mater. Sci. 37 (2002) 3925–3933. [2] X. Peng, J. Yan, Y. Zhou, F. Wang, Effect of grain refinement on the resistance of 304 stainlesssteel to breakaway oxidation in wet air, Acta Mater. 53 (2005) 5079–5088. [3] A.M. Huntz, A. Reckmann, C. Haut, C. Sévérac, M. Herbst, F.C.T. Resende, A.C.S. Sabioni, Oxidation of AISI 304 and AISI 439 stainless steels, Mater. Sci. Eng. A 447 (2007) 266–276. [4] N. Karimi, F. Riffard, F. Rabaste, S. Perrier, R. Cueff, C. Issartel, H. Buscail, Characterization of the oxides formed at 1000 ◦ C on the AISI 304 stainless steel by X-ray diffraction and infrared spectroscopy, Appl. Surf. Sci. 254 (2008) 2292–2299. [5] P. Kofstad, High Temperature Corrosion, Elsevier Science Publishing Co, New York, 1988.

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Please cite this article in press as: A. Col, et al., Oxidation of a Fe–18Cr–8Ni austenitic stainless steel at 850 ◦ C in O2 : Microstructure evolution during breakaway oxidation, Corros. Sci. (2016), http://dx.doi.org/10.1016/j.corsci.2016.10.029