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Cyclic oxidation of yttria dispersed austenitic stainless steels
Corrosion Science 52 (2010) 3573–3576
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Short Communication
Cyclic oxidation of yttria dispersed austenitic stainless steels M.P. Phaniraj, Dong-Ik Kim *, Jae-Hyeok Shim, Young Whan Cho Materials/Devices Division, Korea Institute of Science and Technology, Wolsong-Gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea
a r t i c l e
i n f o
Article history: Received 21 January 2010 Accepted 23 June 2010 Available online 30 June 2010 Keywords: A. Stainless steel B. EPMA B. SEM B. Thermal cycling B. XPS C. Oxidation
a b s t r a c t Cyclic oxidation of austenitic steels (Fe–20Ni–14Cr–2.5Mo–2Mn–2.5Al-wt.%) dispersed with 0, 0.5 and 5 wt.% yttria was carried out at 800 °C in air. The scale surface and cross-section were characterized using X-ray diffraction, scanning electron microscopy (SEM), electron probe microanalysis (EPMA) and X-ray photoelectron spectroscopy (XPS). Yttria additions improve the resistance to spallation. The increase in the resistance to spallation appears to be related to the presence of mixed oxides between yttria and base metal oxides. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Austenitic stainless steels are conventionally used as heat resistant materials because of their good creep strength and oxidation resistance. The creep resistance can be improved further by microalloying with precipitate forming elements such as Ti, V and Nb [1,2]. Their oxidation resistance can also be improved by alloying with aluminum to form a protective layer of alumina. Brady et al. [3] studied the oxidation behavior of such steels with different amounts of microalloying elements and aluminum. Their results show that small changes in the alloying content can change the oxidation behavior from extensive spallation to that which forms a protective scale. Rare earth oxides, particularly yttria is known to increase the creep resistance and oxidation resistance. Several nickel-based and iron-based alloys dispersed with yttria having superior creep resistance and elevated temperature oxidation resistance are commercially available [4]. There are few reports [5,6] on the high temperature oxidation of austenitic steels dispersed with yttria. Lal and Upadhyaya [5] oxidized SS316L in pure oxygen at 550 °C and compared the effects of adding yttria, copper, phosphorus and silicon. They found that the steel containing yttria oxidized to a greater extent when compared with other additions. Bautista et al. [6] showed that addition of 0.5 wt.% yttria to SS304L increased the oxidation resistance and that the increase was higher at 800 °C than at 600 °C. Other reports pertain to the oxidation resistance of austenitic steels that were implanted [7–9] or alloyed
* Corresponding author. Tel.: +82 2 958 5432; fax: +82 2 958 5379. E-mail addresses: [email protected], [email protected] (D.-I. Kim). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.023
[10] with elemental yttrium. These studies show that while the adherence of the scale to the substrate always improves, the effect on oxidation rate depends on the yttrium content. In the present study the oxidation behavior of the Fe–20Ni– 14Cr–2.5Mn–2Mo–2.5Al-wt.% alloys dispersed with 0, 0.5 and 5 wt.% yttria is investigated. The higher yttria content was chosen because the nominal yttria content of 0.5 wt.% is too small to be detected and analyzed with confidence.
2. Experimental The alloys (Fe–20Ni–14Cr–2.5Mn–2Mo–2.5Al-wt.%) were prepared by mechanical alloying (MA) from elemental powders mixed with 0, 0.5 and 5 wt.% Y2O3. X-ray diffraction (XRD) analysis of the Y2O3 powder showed that it consisted entirely of monoclinic yttria. The particle size of yttria was 20 nm. Mechanical alloying was carried out in a planetary ball mill (Retsch PM200) using steel balls and container. XRD and transmission electron microscopy (TEM) analysis of the milled powders was carried out after different intervals of milling to determine the progress of mechanical alloying. The mechanically alloyed powders were hot pressed in vacuum to full density. TEM analysis [11] showed that the hot pressed compacts contained nano-sized (<20 nm) oxide particles of Y2O3, YAlO3, and Y3Al5O12. Further details of alloy preparation and the development of microstructure during mechanical alloying and after hot compaction are published in Ref. [11]. The hot pressed samples were ground (up to 2000 grit SiC paper) and cleaned in acetone and methanol before carrying out oxidation studies. The ground samples are 10 mm in diameter and 4 mm in height. Oxidation experiments were carried out in air (humidity 80%, 15 °C) in a box furnace at 800 °C. The
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samples were kept in alumina crucibles and oxidized for 10 days. The weight of the samples was measured at the end of every cycle of 24 h in a Sartorius CP225D balance that is accurate to 10 2 mg. The samples were removed from the furnace and cooled in air for weight measurements. It took 60 min for cooling and measuring weight following which the sample was put back into the furnace for further oxidation. XRD measurements were carried out on the samples after oxidation. The surface and cross-section of oxide scale was observed using Hitachi 4300SE field emission gun (FEG) SEM equipped with energy dispersive spectroscopy (EDS: Horiba EMAX x-act). The cross-section samples were made after coating the surface (by sputtering) with gold followed by copper plating. The copper plated samples were mounted in conductive resin and then sectioned using a slow speed saw. The surface of the cross-section samples was then prepared by grinding on SiC (600–2000 grit) paper followed by polishing on the polishing cloth that was sprayed with colloidal silica. Composition analysis of the cross-section samples was also carried out using JEOL JXA8500F Field Emission EPMA. XPS measurements (monochromatic Al Ka source) were done on the cross-section samples to determine the chemical state of elements in the scale using PHI-5800 ESCA system. Multipak (PHI) software was used to curve fit the spectra. 3. Results and discussion The weight change per unit area of the specimen is plotted against the oxidation time in Fig. 1. The yttria dispersed alloys gain weight continuously with time whereas the alloy without yttria
Fig. 1. Specimen mass change measured every 24 h in air at 800 °C.
Fig. 2. XRD patterns measured from the sample surface after 10 cycles of oxidation in air at 800 °C. : Fe0.034Mn1.966O3, *: MnFe2O4, +: CrMn2O4, and o: Cr2FeO4.
dispersion continuously loses weight from the second cycle. The weight change in Fig. 1 reflects both the weight loss due to spallation and the weight gain due to oxidation. The weight loss in the alloy without yttria was mainly due to spallation as evidenced by the spalled oxide collected in the bottom of the crucible. There was no spalled oxide in the crucibles of yttria dispersed alloys which indicates that the weight gain was mainly due to oxidation. The weight gain of the 5 wt.% Y2O3 dispersed alloy is significantly higher than that of the 0.5 wt.% Y2O3 dispersed alloy at the end of the first cycle. However, in subsequent cycles the weight gains are similar indicating that the higher yttria content does not affect the oxidation rate of the alloy significantly. XRD analysis after 10 cycles of oxidation (Fig. 2) indicates that the oxides in the scale on all the alloys are mainly Fe0.034Mn1.966O3 (JCPDS 00-024-0507) and MnFe2O4 (JCPDS 01-073-1964). The oxides Mn2CrO4 (JCPDS 01-082-0663) and Cr2FeO4 (00-034-0140) are also detected in the scale particularly in the 0 wt.% Y2O3 alloy. The observation of the sample surface after 10 cycles of oxidation with naked eyes showed that the 0 wt.% Y2O3 alloy had spalled severely at the sample edges. SEM observation showed a number of micrometer sized spalled regions in the scale of the 0 wt.% yttria alloy (Fig. 3a). The scale on yttria dispersed alloys does not show such spalled regions (Fig. 3b). The scale surface in all the alloys consists of oxide crystals as shown in the high magnification SEM micrograph of the scale on the 5 wt.% Y2O3 alloy in Fig. 3c. The EDS analysis showed that the oxide crystals are Mn-rich, which is consistent with the XRD analysis.
Fig. 3. SEM micrographs of (a) oxide surface in 0 wt.% Y2O3 dispersed alloy, arrows indicate spalled regions, (b) oxide surface in 0.5 wt.% Y2O3 dispersed alloy, and (c) oxide surface at high magnification in 0.5 wt.% Y2O3 dispersed alloy.
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Fig. 4. (a) SEM micrograph of the cross-section of the oxidized 0.5 wt.% Y2O3 dispersed alloy. The arrows point to aluminum-rich oxide particles. (b) EDS line profile showing the composition of the scale in the 0.5 wt.% Y2O3 dispersed alloy. The vertical lines drawn above the element profiles show the parts of the oxide scale that are Mn-rich, Ferich or Cr-rich.
The EDS line scan analysis was carried out on the cross-section of oxide scale to determine the scale composition. The scale composition for the 0.5 wt.% yttria alloy is shown in Fig. 4. Typically, the scale consists of oxide layers that are in the order Cr-rich/ Mn-rich/Fe-rich/Mn-rich from the alloy–scale interface to the scale–air interface. Aluminum is present as discontinuous oxide in the scale and in the substrate near the scale/substrate interface. EDS element maps showed that the dark contrast regions, indi-
cated by arrows in Fig. 4, are aluminum-rich oxide particles. In order to determine the distribution of yttrium in the alloy and the scale, element maps were obtained using EPMA. The SEM micrographs of the cross-section of the 0.5 and 5 wt.% Y2O3 alloys and the corresponding EPMA element maps for yttrium are shown in Fig. 5. It can be seen that yttrium is present throughout the scale and substrate in the 0.5 and 5 wt.% Y2O3 alloys. The concentration of the yttrium in the scale along a line through the section is over-
Fig. 5. SEM micrograph and corresponding element map for yttrium measured using EPMA for (a) 0.5 wt.% and (b) 5 wt.% Y2O3 dispersed alloy.
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the interface and elimination of voids at the interface. The elimination of voids at the interface has been related to the formation of mixed oxides of yttria such as YCrO3 [20,21]. The latter perovskite type oxides allow the diffusion of oxygen but block the outward transport of cations [20]. This prevents the formation of vacancies at the alloy–scale interface and thereby reduces spallation. Although, this cannot be singled out as the prominent mechanism, the presence of mixed oxides of yttria (Figs. 5 and 6) in the present study makes it a possible mechanism. However, it is not clear how yttria or its mixed oxides improve the adherence between oxide layers within the scale. 4. Conclusions
Fig. 6. XPS Y(3d) spectra from the scale cross-section and the MA powder of the 5 wt.% yttria dispersed alloy. The data for the MA powder is from Ref. [11].
laid on the map. In the 0.5 wt.% Y2O3 alloy the concentration of yttrium in the scale and substrate is similar. However, in the 5 wt.% Y2O3 alloy, the concentration of yttrium in the scale is considerably less than that in the alloy indicating that the oxide layer at the scale–alloy interface acts as a barrier to the transport of yttrium. This observation is consistent with reports in literature [12–14] on segregation of reactive elements including yttrium to the alloy–scale interface. XPS measurements were carried out on the scale cross-section to determine the chemical state of yttrium. In Fig. 6, the Y(3d) spectrum from the scale cross-section of the 5 wt.% Y2O3 alloy is compared with the spectrum from the mechanically alloyed (MA) powder. The data for MA powder in Fig. 6 is from the authors’ previous work [11] where it was shown that nano-sized (<20 nm) mixed oxides of yttria and alumina such as Y3Al5O12 (YAG) and YAlO3 (YAP) are present in the MA powder as well as in the hot pressed compacts. It can be seen from Fig. 6 that most of the peaks and changes in gradient observed in the spectrum from the MA powder are also seen in the spectrum from the scale cross-section although with different relative intensities. The increase in intensity could be because the existing oxide phases have grown or increased in number. The binding energy values [15,16] indicating the presence of yttria (Y3d5/2 = 156.7 eV) and YAlO3 (Y3d5/2 = 157.4 eV) are marked in the figure. Oxidation studies on pure iron [17] and pure chromium [18] implanted with yttrium showed the formation of YFeO3 and YCrO3, respectively, during oxidation. The unidentified peaks in the spectra could be due to emissions from these mixed oxides, however, standard reference data for these oxides are not available. In Fe–Cr [13], Ni–Cr [19], Fe–Ni–Cr [8], alloys that form an inner chromia (or Cr-rich oxide) layer and outer spinel (or base metal oxide) layer, the addition of yttrium delays the formation of spinel. This greatly reduces the scale spallation because of the elimination of the oxide/oxide interface. However, in the present study the scale in yttria dispersed alloys is multilayered and adherent. Several other mechanisms have been proposed [12–14] to explain the improved alloy–scale adherence based on reduced growth stress, enhanced scale plasticity, enhanced chemical bonding at
1. The scale formed on austenitic steel during cyclic oxidation at 800 °C in air spalls severely. The addition of 0.5 and 5 wt.% Y2O3 to the steel improves the resistance to scale spallation. 2. The improvement in spallation resistance is possibly related to the presence of mixed oxides of yttria such as YAlO3, YCrO3 and YFeO3, in the scale. 3. The scale formed in the 5 wt.% Y2O3 dispersed steel shows that the Cr-rich oxide layer at the scale–alloy interface limits the transport of yttrium to the scale.
Acknowledgement This research was supported by Korea Institute of Science and Technology (Grant No. 2E21652). References [1] P.J. Maziasz, J. Metals (July) (1989) 14–20. [2] Y. Yamamoto, M.P. Brady, Z.P. Lu, P.J. Maziasz, C.T. Liu, B.A. Pint, K.L. More, H.M. Meyer, E.A. Payzant, Science 316 (2007) 433–436. [3] M.P. Brady, Y. Yamamoto, M.L. Santella, B.A. Pint, Scripta Mater. 57 (2007) 1117–1120. [4] C. Capdevila, H.K.D.H. Bhadeshia, Adv. Eng. Mater. 3 (2001) 647–656. [5] S. Lal, G.S. Upadhyaya, Oxidat. Metals 32 (1989) 317–335. [6] A. Bautista, F. Velasco, J. Abenojar, Corros. Sci. 45 (2003) 1343–1354. [7] J.C. Pivin, D. Delaunay, C. Roques-Carmes, A.M. Huntz, P. Lacombe, Corros. Sci. 20 (1980) 351–373. [8] J.C. Pivin, C. Roques-Carmes, J. Chaumont, H. Bernas, Corros. Sci. 20 (1980) 947–962. [9] J.E. Antill, M.J. Bennett, R.F.A. Carney, G. Dearnaley, F.H. Fern, P.D. Goode, B.L. Myatt, J.B. Warburton, Corros. Sci. 16 (1976) 729–745. [10] M. Slater, W.A. Grant, G. Carter, Radiat. Effects 82 (1984) 239–248. [11] M.P. Phaniraj, D.-I. Kim, J.-H. Shim, Y.W. Cho, Acta Mater. 57 (2009) 1856– 1864. [12] D.P. Whittle, J. Stringer, Philos. Trans. Royal Soc. Lond. A 295 (1980) 309–329. [13] P.Y. Hou, J. Stringer, Mater. Sci. Eng. A 202 (1995) 1–10. [14] B.A. Pint, in: Proceedings of the John Stringer Symposium High Temperature Corrosion 5–8th November 2001, ASM International, Indianapolis, IN, 2003, pp. 9–19. [15] D.A. Pawlak, K. Wozniak, Z. Frukacz, T.L. Barr, D. Fiorentino, S. Hardcastle, J. Phys. Chem. B 103 (1999) 3332–3336. [16] J.F. Moulder, W.F. Stickel, P.E. Sobol, K.D. Bomben, Handbook of X-ray Photoelectron Spectroscopy, Physical Electronics Inc., USA, 1995. [17] C. Josse-Courty, H. Buscail, M.F. Stroosnijder, P. Dufour, J.P. Larpin, Oxidat. Metals 52 (1999) 321–336. [18] Y.P. Jacob, H. Buscail, E. Caudron, R. Cueff, M.F. Stroosnijder, J. Phys. IV France 12 (2002) 207–214. [19] Y. Zhang, D. Zhu, D.A. Shores, Acta Mater. 43 (1995) 4015–4025. [20] S. Chevalier, J.P. Larpin, Acta Mater. 50 (2002) 3105–3114. [21] Y. Saito, B. Onay, T. Maruyama, J. de Phys. IV 3 (1993) 217–230.
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Short Communication
Cyclic oxidation of yttria dispersed austenitic stainless steels M.P. Phaniraj, Dong-Ik Kim *, Jae-Hyeok Shim, Young Whan Cho Materials/Devices Division, Korea Institute of Science and Technology, Wolsong-Gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea
a r t i c l e
i n f o
Article history: Received 21 January 2010 Accepted 23 June 2010 Available online 30 June 2010 Keywords: A. Stainless steel B. EPMA B. SEM B. Thermal cycling B. XPS C. Oxidation
a b s t r a c t Cyclic oxidation of austenitic steels (Fe–20Ni–14Cr–2.5Mo–2Mn–2.5Al-wt.%) dispersed with 0, 0.5 and 5 wt.% yttria was carried out at 800 °C in air. The scale surface and cross-section were characterized using X-ray diffraction, scanning electron microscopy (SEM), electron probe microanalysis (EPMA) and X-ray photoelectron spectroscopy (XPS). Yttria additions improve the resistance to spallation. The increase in the resistance to spallation appears to be related to the presence of mixed oxides between yttria and base metal oxides. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Austenitic stainless steels are conventionally used as heat resistant materials because of their good creep strength and oxidation resistance. The creep resistance can be improved further by microalloying with precipitate forming elements such as Ti, V and Nb [1,2]. Their oxidation resistance can also be improved by alloying with aluminum to form a protective layer of alumina. Brady et al. [3] studied the oxidation behavior of such steels with different amounts of microalloying elements and aluminum. Their results show that small changes in the alloying content can change the oxidation behavior from extensive spallation to that which forms a protective scale. Rare earth oxides, particularly yttria is known to increase the creep resistance and oxidation resistance. Several nickel-based and iron-based alloys dispersed with yttria having superior creep resistance and elevated temperature oxidation resistance are commercially available [4]. There are few reports [5,6] on the high temperature oxidation of austenitic steels dispersed with yttria. Lal and Upadhyaya [5] oxidized SS316L in pure oxygen at 550 °C and compared the effects of adding yttria, copper, phosphorus and silicon. They found that the steel containing yttria oxidized to a greater extent when compared with other additions. Bautista et al. [6] showed that addition of 0.5 wt.% yttria to SS304L increased the oxidation resistance and that the increase was higher at 800 °C than at 600 °C. Other reports pertain to the oxidation resistance of austenitic steels that were implanted [7–9] or alloyed
* Corresponding author. Tel.: +82 2 958 5432; fax: +82 2 958 5379. E-mail addresses: [email protected], [email protected] (D.-I. Kim). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.023
[10] with elemental yttrium. These studies show that while the adherence of the scale to the substrate always improves, the effect on oxidation rate depends on the yttrium content. In the present study the oxidation behavior of the Fe–20Ni– 14Cr–2.5Mn–2Mo–2.5Al-wt.% alloys dispersed with 0, 0.5 and 5 wt.% yttria is investigated. The higher yttria content was chosen because the nominal yttria content of 0.5 wt.% is too small to be detected and analyzed with confidence.
2. Experimental The alloys (Fe–20Ni–14Cr–2.5Mn–2Mo–2.5Al-wt.%) were prepared by mechanical alloying (MA) from elemental powders mixed with 0, 0.5 and 5 wt.% Y2O3. X-ray diffraction (XRD) analysis of the Y2O3 powder showed that it consisted entirely of monoclinic yttria. The particle size of yttria was 20 nm. Mechanical alloying was carried out in a planetary ball mill (Retsch PM200) using steel balls and container. XRD and transmission electron microscopy (TEM) analysis of the milled powders was carried out after different intervals of milling to determine the progress of mechanical alloying. The mechanically alloyed powders were hot pressed in vacuum to full density. TEM analysis [11] showed that the hot pressed compacts contained nano-sized (<20 nm) oxide particles of Y2O3, YAlO3, and Y3Al5O12. Further details of alloy preparation and the development of microstructure during mechanical alloying and after hot compaction are published in Ref. [11]. The hot pressed samples were ground (up to 2000 grit SiC paper) and cleaned in acetone and methanol before carrying out oxidation studies. The ground samples are 10 mm in diameter and 4 mm in height. Oxidation experiments were carried out in air (humidity 80%, 15 °C) in a box furnace at 800 °C. The
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samples were kept in alumina crucibles and oxidized for 10 days. The weight of the samples was measured at the end of every cycle of 24 h in a Sartorius CP225D balance that is accurate to 10 2 mg. The samples were removed from the furnace and cooled in air for weight measurements. It took 60 min for cooling and measuring weight following which the sample was put back into the furnace for further oxidation. XRD measurements were carried out on the samples after oxidation. The surface and cross-section of oxide scale was observed using Hitachi 4300SE field emission gun (FEG) SEM equipped with energy dispersive spectroscopy (EDS: Horiba EMAX x-act). The cross-section samples were made after coating the surface (by sputtering) with gold followed by copper plating. The copper plated samples were mounted in conductive resin and then sectioned using a slow speed saw. The surface of the cross-section samples was then prepared by grinding on SiC (600–2000 grit) paper followed by polishing on the polishing cloth that was sprayed with colloidal silica. Composition analysis of the cross-section samples was also carried out using JEOL JXA8500F Field Emission EPMA. XPS measurements (monochromatic Al Ka source) were done on the cross-section samples to determine the chemical state of elements in the scale using PHI-5800 ESCA system. Multipak (PHI) software was used to curve fit the spectra. 3. Results and discussion The weight change per unit area of the specimen is plotted against the oxidation time in Fig. 1. The yttria dispersed alloys gain weight continuously with time whereas the alloy without yttria
Fig. 1. Specimen mass change measured every 24 h in air at 800 °C.
Fig. 2. XRD patterns measured from the sample surface after 10 cycles of oxidation in air at 800 °C. : Fe0.034Mn1.966O3, *: MnFe2O4, +: CrMn2O4, and o: Cr2FeO4.
dispersion continuously loses weight from the second cycle. The weight change in Fig. 1 reflects both the weight loss due to spallation and the weight gain due to oxidation. The weight loss in the alloy without yttria was mainly due to spallation as evidenced by the spalled oxide collected in the bottom of the crucible. There was no spalled oxide in the crucibles of yttria dispersed alloys which indicates that the weight gain was mainly due to oxidation. The weight gain of the 5 wt.% Y2O3 dispersed alloy is significantly higher than that of the 0.5 wt.% Y2O3 dispersed alloy at the end of the first cycle. However, in subsequent cycles the weight gains are similar indicating that the higher yttria content does not affect the oxidation rate of the alloy significantly. XRD analysis after 10 cycles of oxidation (Fig. 2) indicates that the oxides in the scale on all the alloys are mainly Fe0.034Mn1.966O3 (JCPDS 00-024-0507) and MnFe2O4 (JCPDS 01-073-1964). The oxides Mn2CrO4 (JCPDS 01-082-0663) and Cr2FeO4 (00-034-0140) are also detected in the scale particularly in the 0 wt.% Y2O3 alloy. The observation of the sample surface after 10 cycles of oxidation with naked eyes showed that the 0 wt.% Y2O3 alloy had spalled severely at the sample edges. SEM observation showed a number of micrometer sized spalled regions in the scale of the 0 wt.% yttria alloy (Fig. 3a). The scale on yttria dispersed alloys does not show such spalled regions (Fig. 3b). The scale surface in all the alloys consists of oxide crystals as shown in the high magnification SEM micrograph of the scale on the 5 wt.% Y2O3 alloy in Fig. 3c. The EDS analysis showed that the oxide crystals are Mn-rich, which is consistent with the XRD analysis.
Fig. 3. SEM micrographs of (a) oxide surface in 0 wt.% Y2O3 dispersed alloy, arrows indicate spalled regions, (b) oxide surface in 0.5 wt.% Y2O3 dispersed alloy, and (c) oxide surface at high magnification in 0.5 wt.% Y2O3 dispersed alloy.
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Fig. 4. (a) SEM micrograph of the cross-section of the oxidized 0.5 wt.% Y2O3 dispersed alloy. The arrows point to aluminum-rich oxide particles. (b) EDS line profile showing the composition of the scale in the 0.5 wt.% Y2O3 dispersed alloy. The vertical lines drawn above the element profiles show the parts of the oxide scale that are Mn-rich, Ferich or Cr-rich.
The EDS line scan analysis was carried out on the cross-section of oxide scale to determine the scale composition. The scale composition for the 0.5 wt.% yttria alloy is shown in Fig. 4. Typically, the scale consists of oxide layers that are in the order Cr-rich/ Mn-rich/Fe-rich/Mn-rich from the alloy–scale interface to the scale–air interface. Aluminum is present as discontinuous oxide in the scale and in the substrate near the scale/substrate interface. EDS element maps showed that the dark contrast regions, indi-
cated by arrows in Fig. 4, are aluminum-rich oxide particles. In order to determine the distribution of yttrium in the alloy and the scale, element maps were obtained using EPMA. The SEM micrographs of the cross-section of the 0.5 and 5 wt.% Y2O3 alloys and the corresponding EPMA element maps for yttrium are shown in Fig. 5. It can be seen that yttrium is present throughout the scale and substrate in the 0.5 and 5 wt.% Y2O3 alloys. The concentration of the yttrium in the scale along a line through the section is over-
Fig. 5. SEM micrograph and corresponding element map for yttrium measured using EPMA for (a) 0.5 wt.% and (b) 5 wt.% Y2O3 dispersed alloy.
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the interface and elimination of voids at the interface. The elimination of voids at the interface has been related to the formation of mixed oxides of yttria such as YCrO3 [20,21]. The latter perovskite type oxides allow the diffusion of oxygen but block the outward transport of cations [20]. This prevents the formation of vacancies at the alloy–scale interface and thereby reduces spallation. Although, this cannot be singled out as the prominent mechanism, the presence of mixed oxides of yttria (Figs. 5 and 6) in the present study makes it a possible mechanism. However, it is not clear how yttria or its mixed oxides improve the adherence between oxide layers within the scale. 4. Conclusions
Fig. 6. XPS Y(3d) spectra from the scale cross-section and the MA powder of the 5 wt.% yttria dispersed alloy. The data for the MA powder is from Ref. [11].
laid on the map. In the 0.5 wt.% Y2O3 alloy the concentration of yttrium in the scale and substrate is similar. However, in the 5 wt.% Y2O3 alloy, the concentration of yttrium in the scale is considerably less than that in the alloy indicating that the oxide layer at the scale–alloy interface acts as a barrier to the transport of yttrium. This observation is consistent with reports in literature [12–14] on segregation of reactive elements including yttrium to the alloy–scale interface. XPS measurements were carried out on the scale cross-section to determine the chemical state of yttrium. In Fig. 6, the Y(3d) spectrum from the scale cross-section of the 5 wt.% Y2O3 alloy is compared with the spectrum from the mechanically alloyed (MA) powder. The data for MA powder in Fig. 6 is from the authors’ previous work [11] where it was shown that nano-sized (<20 nm) mixed oxides of yttria and alumina such as Y3Al5O12 (YAG) and YAlO3 (YAP) are present in the MA powder as well as in the hot pressed compacts. It can be seen from Fig. 6 that most of the peaks and changes in gradient observed in the spectrum from the MA powder are also seen in the spectrum from the scale cross-section although with different relative intensities. The increase in intensity could be because the existing oxide phases have grown or increased in number. The binding energy values [15,16] indicating the presence of yttria (Y3d5/2 = 156.7 eV) and YAlO3 (Y3d5/2 = 157.4 eV) are marked in the figure. Oxidation studies on pure iron [17] and pure chromium [18] implanted with yttrium showed the formation of YFeO3 and YCrO3, respectively, during oxidation. The unidentified peaks in the spectra could be due to emissions from these mixed oxides, however, standard reference data for these oxides are not available. In Fe–Cr [13], Ni–Cr [19], Fe–Ni–Cr [8], alloys that form an inner chromia (or Cr-rich oxide) layer and outer spinel (or base metal oxide) layer, the addition of yttrium delays the formation of spinel. This greatly reduces the scale spallation because of the elimination of the oxide/oxide interface. However, in the present study the scale in yttria dispersed alloys is multilayered and adherent. Several other mechanisms have been proposed [12–14] to explain the improved alloy–scale adherence based on reduced growth stress, enhanced scale plasticity, enhanced chemical bonding at
1. The scale formed on austenitic steel during cyclic oxidation at 800 °C in air spalls severely. The addition of 0.5 and 5 wt.% Y2O3 to the steel improves the resistance to scale spallation. 2. The improvement in spallation resistance is possibly related to the presence of mixed oxides of yttria such as YAlO3, YCrO3 and YFeO3, in the scale. 3. The scale formed in the 5 wt.% Y2O3 dispersed steel shows that the Cr-rich oxide layer at the scale–alloy interface limits the transport of yttrium to the scale.
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