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Effect of cerium implantation on the corrosion of alloy 800H in an S-O-C environment
Materials Science and Engineering, A l l 6 (1989) 103-110
103
Effect of Cerium Implantation on the Corrosion of Alloy 800H in an S - O - C Environment* M. E STROOSNIJDER, J. E NORTON and V. GUTTMANN
Commission of the European Communities, Joint Research Centre, Petten Establishment, Post Box 2, 1755 ZG Petten (The Netherlands') M. J. BENNETT
Materials Development Division, Atomic Energy Research Establishment, Harwell, Didcot (U.K.) J. H. W. DE WIT
Laboratory for Materials Science, Delft University of Technology, Postbox 5, 2600 AA Delft (The Netherlands) (Received September 16, 1988)
Abstract
The corrosion behaviour of a wrought austenitic Fe-20Cr-32Ni steel, Alloy 80OH, was studied in a simulated coal gasification atmosphere at 700 °C. Exposures ranged from as short as a few minutes to periods of up to 1000 h. The influence of cerium on the corrosion resistance was established by carrying out identical tests after ion implantation of the specimen surfaces. It was shown that the corrosive attack in the t12S-containing gas mixture can be reduced for the cerium-implanted material if the implantation dose is sufficiently high. In the case of both implanted and unimplanted material, sulphides and oxides were formed but the corrosion products of the implanted material showed a higher Cr.'Fe ratio relative to those of the unimplanted material. It is proposed that this feature is responsible for the observed improvement,
1. Introduction
In several technical high-temperature processes, such as during the gasification of coal, structural materials are exposed to atmospheres of low oxygen and high sulphur partial pressures. The ability of alloys to resist corrosion under these conditions depends on their potential to form and maintain a protective oxide scale, e . g .
*Paper presented at the Sixth International Conference on Surface Modification of Metals by Ion Beams, Riva del Garda, Italy, September 12-16, 1988. 0921-5093/89/$3.50
SiO 2 o r A1203. Unfortunately, in these atmospheres the simultaneous formation of sulphides as well as oxides is observed, although according to the thermodynamic equilibria, only oxides should be present. The exclusive formation of oxides occurs only beyond the so-called "kinetic boundary", which lies at a higher oxygen
Cr203,
activity than that predicted by thermodynamics [1-3]. In this context the role of rare earth elements is of interest. These elements may be incorporated in bulk alloys as metallic or oxide-dispersoid compounds or can be added to the surface layers by ion implantation. The addition of small amounts of rare earth elements to high-temperature alloys is known to improve the oxidation resistance and scale adherence [4-6], although the exact mechanism is still the subject of discussion. The early stages of nucleation and growth of the oxide scale seem to play an important role [4, 7-9]. Most of the work on the effect of rare earth elements has been undertaken in purely oxidizing environments. Only limited experimental work has been carried out in mixed oxidizing-sulphidizing environments. In this paper, the effect of cerium ion implantation on the corrosion behaviour of Alloy 800H (a wrought austenitic Fe-20Cr-32Ni alloy) in a simulated coal gasification atmosphere at 700 °C, is reported. The chemical potentials of the environment were selected so as to be of relevance to a moderately aggressive environment, i.e. a high tendency for oxide formation, without excluding sulphidation. © Elsevier Sequoia/Printed in The Netherlands
104
2. Experimental details The alloy investigated came from a commercial heat and by analysis contained 19.6 wt.% Cr, 32.2 wt.% Ni, 0.7 wt.% Mn, 0.45 wt.% Si, 0.46 wt.% Ti, 0.38 wt.% AI and 0.09 wt.% C. After solution annealing at 1150 °C for 0.5 h followed by water quenching, specimens in the form of coupons with dimensions 10 mm x 8 mm x 2.5 mm were machined. Prior to implantation they were abraded on 600 grit silicon carbide paper and degreased, The uniform implantation of cerium, at a nominal dose of either 1016 o r 1017 ions cm -2 into one of the principal coupon faces was undertaken using a 200 keV accelerating potential. The cerium distribution was determined experimentally using Rutherford backscattering spectrometry (RBS), secondary neutral mass spectroscopy (SNMS) and X-ray photoelectron spectroscopy(XPS). Corrosion tests were carried out in a hydrogen-based gas mixture with 7% CO, 1.2% H 2 0 and 0.2% H2S. At the test temperature of 700 °C and a total pressure of 1.4 atm, the sulphur partial pressure was 0.2 x 10 -9 atm, while the oxygen partial pressure was 0 . 7 x 10 -23 atm. The gas flowed continuously through the corrosion autoclaves at about 20 1 h- 1. Further details of the test facilities have been described elsewhere [10]. The specimens, each contained individually in a pure A I 2 0 3 crucible, were lowered into the reaction chamber at room temperature and then heated up in about 2 h to the test temperature in the test gas. The corrosion kinetics were established using discontinuous gravimetric measurements. For short-term tests a special autoclave system was used. This was fitted with a plunger device in order that samples could be rapidly inserted and retrieved from the hot gas atmosphere, thereby enabling the specimens to be exposed at temperature for periods of only a few minutes. To characterize the structure and composition of the corrosion products formed, a range of analysis techniques was used including optical metallography (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), electron probe microanalysis (EPMA), XPS and SNMS.
3. Results
3.1. Characterization of the implant distributions An estimate of the initial distribution of the implanted atoms in a material is possible using
the Lindhard-Scharff-Schott theory [11]. However, in the case of high implantation doses the observed implant distribution may deviate significantly from the theoretically predicted distribution, for example because of material loss due to sputtering. Computer modelling can take this into account. Distribution profiles calculated for the implantations of 1016 and 1017 Ce ions cm-2 in Alloy 800H are presented in Fig. 1. These profiles were obtained by the Monte Carlo computer program TRIM developed by Biersack and Haggmark [12]. It was calculated that the retained dose approximately equals the nominal dose in the situation of the low implantation dose. However, for the high implantation dose the retained dose (4 × 1016 C e cm -2) was only 40% of the nominal dose due to the sputtering loss of 80 nm. The actual implant distributions were determined experimentally using SNMS, RBS and XPS (Fig. 1). For the cerium distribution obtained by SNMS, standard sensitivity factors [13] were used. For SNMS the sputter rate was taken as 0.1 nm s- 1, whereas for XPS it was about 0.05 nm s- 1. The chemical state of cerium was difficult to define from the XPS data due to chemical reduction of the surface to a lower oxide state during ion etching.
. ~*~ _~~ ~
.
.
. Attoy800H
1016Ce ions/cm 2
t
§2 § 6° 20
40
60
80
100
Depth. nrn
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°,
,'
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'
'
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Fig. 1. Experimentally determined and calculated cerium distributioon profiles for Alloy 800H: (a) 1016 ions cm-2;
(b) 1017ionscm-2.
105
3.2. Gravimetric determinations For mass gain measurements a set of three specimens (unimplanted, 1016 Ce ions cm -2 and 1017 Ce ions cm -2) was exposed for 200 h in the test gas, whilst another set was exposed intermittently for 25, 100, 250, 500, 750 and 1000 h. Because the ratio of implanted to unimplanted area was not precisely known it was not possible to obtain accurately the influence of implantation on mass gain and the measurements therefore only provided an approximate indication of the corrosion resistance. The situation was complicated further by the relatively large scatter observed in the case of low mass gains. The overall conclusion from these experiments was that the material implanted with 1016 Ce ions cm -2 exhibited a similar mass gain to the unimplanted material, whereas for the alloy implanted with 1017 Ce ions cm 2 alloy, the mass gain was reduced by a factor of about 3. The alloy implanted with 1017 Ce ions cm-2 which was successively exposed for durations up to 1000 h, also showed a lower mass gain compared with the Alloy 800H, although the difference became smaller with increasing exposure time.
3.3. Nucleation andgrowth studies In order to analyse the structural changes occurring during corrosion, additional specimens were exposed for 2, 5, 15 and 30 min, as well as 5 h. The behaviour of the implanted and unimplanted sides on the same specimen were directly compared. In all experiments the specimens implanted with 1016 Ce ions cm -2 showed identical corrosion behaviour to the unimplanted alloy and therefore will not be considered in further detail, The SEM micrographs in Fig. 2 depict the changes in scale morphology that occurred between exposures of 2 min and 200 h for the unimplanted alloy (Figs. 2(a)-2(c)) and the Alloy 800H implanted with 1017 Ce ions cm -2 (Figs. 2(d)-2(f)). XPS analyses showed that the corrosion products which formed initially on the unimplanted alloy were a mixture of mainly iron and chromium sulphides and small amounts of oxides, The lighter features in Fig. 2(b) initially became evident after 5 rain and increased significantly in number and size during the first hours of exposure. With continuing time they grew in size although their number decreased due to coalescence. These features were shown by XRD and
EDX to be sulphides of the types M3S4 and MS and to contain iron and chromium with the iron content increasing with the size of the sulphide. The sulphur content was higher than for the corrosion layer which formed initially. XRD analysis after exposures of 30 rain or longer indicated the presence of M304 for the unimplanted material. The lattice parameter of the M304 was consistent with a (Cr, Fe)304 spinel. The scale formed on Alloy 800H implanted with 1017 Ce ions cm -z after a few minutes showed a less rough morphology than that on the unimplanted alloy. From Fig. 2 it appears that the originally formed individual sulphide nuclei coalesce to a large extent thereby forming a fairly homogeneous flat scale covering the underlying alloy. This morphology was confirmed indirectly by XPS element profiling and comparing the scales formed on the implanted and unimplanted alloy after 5 min exposure. The implanted material gave a lower nickel signal (from the substrate) and a higher sulphur signal because of a more denser surface layer. Moreover, from observation of the sputter time necessary to reach the uncorroded alloy in the XPS analysis, it appeared that the corrosion layer formed on the implanted alloy was thinner than that on Alloy 800H. From the XPS data it may be concluded that the surface sulphides formed initially on the implanted alloy contained a higher Cr:Fe ratio than those on the unimplanted material. An oxygen enrichment under the predominantly sulphide scale, positioned slightly below the cerium peak, was evident for implanted material from the XPS analyses. The morphology of the original surface layer of the material implanted with 1017 Ce ions cm -2 appeared to change with exposure time from a relatively flat layer to a more crystalline morphology. Again, with increasing exposure time the formation of light features above the initial sulphide scale was observed on the implanted material (1017 Ce ions cm -2) but their appearance was clearly delayed and became evident only after 30 min exposure. EDX analysis showed that these features had a similar composition for both the implanted and the unimplanted alloy. The number of features increased with time, but for durations of up to 200 h their size was significantly smaller than for the unimplanted material and no coalescence was observed. XRD analyses indicated the same type of corrosion phases to be present on the implanted as on the unimplanted alloy, i.e. MS, M3S4 and M304
106
Fig. 2. Surface scale morphology of (a), (b), (c) Alloy 800H and (d), (e), (f) Alloy 800H implanted with 1017 Ce ions cm -2 after exposure for (a), (d) 2 rain, (b), (e) 15 rain and (c), (f) 200 h in H2-7voI.%CO-0.2voI.%H2S- 1.2vol.%H20 at 700 °C.
107
(obviously also (Cr,Fe)304). After 1000 h exposure the corrosion appeared to be similar for both unimplanted and implanted Alloy 800H. Element depth profiles obtained after 200 h exposure by SNMS of an area 3 mm in diameter through the scale for the unimplanted and implanted (1017 Ce ions cm 2) material are presented in Fig. 3. The sputter rate was about 0.5 nm s- i. It should be pointed out that although the relative intensity in Fig. 3 is indicative of the atomic percentage of the elements, average sensitivity factors [13] were taken for the calculations and no standards for calibration were applied, Hence the data cannot be used in a fully quantitative manner. The profiles (Fig. 3(a)) indicate for the unimplanted material only weakly pronounced peak positions for the various elements, The ion content was found to be higher than the chromium content throughout the entire scale. By comparison, Fig. 3(b) illustrates clearly separated
60
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peak positions for the implanted material. The high sulphur concentration is attributed to chromium and iron which were both present at similar levels. A region of oxygen enrichment observed beneath the sulphide layer was related to an increase of the chromium level with a concomitant reduction in iron. The cerium peak lies slightly above this oxygen enrichment. A marked oxygen concentration of the outermost layer cannot be due solely to impurity adsorption. For both materials the manganese content of the scale was found to be low. Cross-sectional examinations by OM and SEM (Fig. 4), revealed that the scale formed on the material implanted with 10 Iv Ce ions cm- 2 was more uniform in thickness than that on the unimplanted material. This was also found during the early stages of corrosion as well as after 200 h exposure. EDX spectra confirmed a higher Fe:Cr ratio for the scale on the unimplanted than on the implanted alloy. After 5 h exposure the untreated alloy exhibited a significant number of surface protrusions and also internal corrosion appeared. In contrast, the implanted material (1017 ions cm -2) after 5 h exposure showed almost no localized external nor localized internal corrosion. After 200 h both internal and external local corrosion was detected for the implanted material, although it was much less advanced than for the unimplanted sample. Attempts to perform quantitative EPMA failed due to the relatively bad lateral resolution and large analysis depth of this technique. The qualitative EPMA results of the outer scale formed after 200 h on the unimplanted material indicated essentially sulphur, iron and chromium with an increasing iron content in the larger nodules towards the outer surface. An oxide-rich zone appeared below this sulphide layer which contained chromium and iron. The internal corrosion products were mainly composed of Cr-Fe oxide and showed additionally an enrichment in silicon, aluminium and titanium. A low sulphur content, which increased towards the tip of the internal corrosion paths indicated the presence of small amounts of individual sulphides. EPMA of the scale formed
r 10
12
1/.
1(5
Sputter,ngtime.ks
Fig. 3. SNMS element distribution profiles of (a) Alloy 800H and (b) Alloy 800H implanted with 10 j7 Ce ions cm 2 after exposure for 200 h in H2-7vol.%CO-0.2vol.%H2S1.2vol.%H20 at 700 °C.
after 200
h exposure
confirmed
the SNMS results from the implanted alloy, i.e. the scale was rich in sulphur, chromium, iron and oxygen with an increased oxygen level towards both the gas-scale and the substrate-scale interfaces. The external nodules were Fe-Cr sulphides
with the few internal corrosion observed also containing manganese.
products
108
i:',:~'~
ii!:i:
~~! : ()~
...........
Fig. 4. Cross-sectional SEM examinations of (a), (b) Alloy 800H and (c), (d) Alloy 800H implanted with 1017 Ce ions cm -2 after exposure tot /a), (c) 5 h and (b), (d) 200 h in H2-7vol.%CO-0.2vol.%H2S-1.2vol.%H2O at 700 °C.
4. D i s c u s s i o n and c o n c l u s i o n s
The experimentally observed distributions and the total uptake of cerium in Alloy 800H (Fig. 1 ) for the high implantation dose were close to those expected from the Monte Carlo calculations. This agreement on the alloy implanted with 1017 Ce ions cm-2 shows clearly that the phenomenon of sputtering can significantly limit the maximum concentrations that can be achieved by implantation, especially in the case of high implantation doses. This agrees with the work of others [14, 15]. The investigations of the corrosion behaviour have shown that the attack in the S - O - C gaseous atmosphere can be reduced for the cerium implanted material if the implantation dose is high. The specimens with only 1016 Ce ions cm -2 show an identical behaviour, in terms of both the mass gain and scale formation, to the unimplanted material. Implantation with nominal 1017 Ce ions cm-2 leads to a decrease in mass gain by a factor of up to 3 and a clear retardation of the corrosion rate, at least for durations of up to 200 h. It should be noted that for a 1000 h test a very similar behaviour has been observed on the implanted and unimplanted alloy. However, the influence of temperature cycling, due to the interruptions of the test, may have contributed to this behaviour. Studies by SEM and various analytical techniques have shown a fairly classical corrosion pattern for the unimplanted alloy [1, 16-18]. Briefly, in mixed oxidizing and/or sulphidizing atmospheres, both oxides and sulphides may nucleate on the alloy surface even though thermodynamic considerations indicate that the oxide is the most stable phase. This is due to the fact that kinetic factors may dominate the overall corrosion process and sulphides may overgrow the oxide nuclei. No protective oxide will be formed in this case. This behaviour has also been observed in the present investigation. Although oxide stability would be predicted from thermodynamic considerations of the gas composition, it was found that after short exposures a high density of individual sulphides rich in iron and chromium were present but only a small amount of oxide could be detected by XPS. This behavtour shows that under the test conditions eraployed, sulphidation is the dominant process. The observation of internal corrosion indicates that the oxides do not form a protective layer.
109
Poor protection is also confirmed by the formation of fast growing iron-rich sulphides on the initially formed sulphide layer. Iron cation transport to the scale-gas interface appears to be facilitated locally which probably results from the formation of sulphide-rich channels which are known to be fast diffusion paths [19-22]. It is also known that the diffusion of iron through these channels is much faster than that of chromium [16, 23]. It should be noted that the internal corrosion products were primarily oxides and that for durations of up to 200 h exposure only a few sulphides have been observed at the base of the corrosive attack. This feature can be related to the release of sulphur caused by the transformation of sulphides into oxides, as has also been observed by others [18, 21, 24]. Regarding the corrosion behaviour of the material implanted with 10 ~7 Ce ions cm-2, it is somewhat surprising that although there was a high tendency for sulphidation to occur during the initial stage of exposure, the corrosion resistance in general was found to be improved. In comparison with the unimplanted alloy the sulphides initially formed appear to exhibit a great tendency to grow laterally with a consequent coalescence and a higher coverage of the underlying alloy. Apart from this difference in morphology, the sulphides formed on the material implanted with 1017 Ce ions cm -2 exhibited a higher Cr:Fe ratio than those on the untreated material. Similar observations derived from sulphidation studies of a Fe-15Cr-4AI-1Y alloy in comparison with Fe-15Cr-4A1 [25] and on the sulphidation of N i - C r - Y alloys [26]. In contrast to the unimplanted alloy, the nucleation and growth of local sulphides above the coherent sulphide layer which formed initially was strongly retarded for the alloy implanted with 10 Iv Ce ions cm-2. The Cr:Fe ratio of these local sulphides was also higher than for the unimplanted material. An oxide layer had again formed beneath the chromium-rich sulphide layer, comparable with the situation of the unimplanted material. From the SNMS investigations after 200 h exposure, it can be concluded that the oxide layer had a relatively high chromium content compared with that of iron. From the XPS and SNMS analyses it became apparent that the implanted cerium was concentrated on top of the oxide layer. This indicates that the oxide layer had grown preferentially by oxygen anion transport.
The mechanisms governing the improvement in corrosion resistance by the implantation of 1017 Ce ions cm -2, demonstrated first by the retardation in the development of individual large sulphides on top of the initially formed sulphide layer and second by the absence of internal corrosion at least up to 200 h, unfortunately could not be fully established. It is probable that the oxide layer which formed beneath the sulphides plays a decisive role. In general it is well known that chromium-rich oxide layers confer greater protection than that provided by iron-rich spinel phases [22, 27, 28]. Indeed, the present investigations have indicated a higher chromium content for the oxide on the implanted alloy compared with that on the unimplanted material. This is consistent with observations made in purely oxidizing atmospheres [4, 8, 29, 30] that materials with rare earth elements show a stronger tendency for chromia formation. Generally speaking, all the external corrosion products observed on the implanted alloy showed a relatively high chromium content compared with that of iron. It may be argued that, in addition, the morphology of the initial sulphide layer formed on the alloy implanted with 10 ~7 Ce ions cm -2 may have a beneficial influence on the oxide formation. Taking into account that the sulphur potential at the alloy-sulphide scale interface must be lower than that in the corrosive test gas, this will lead to a relative increase in the oxide-forming potential. Because of the improved coverage of the alloy surface by sulphides in the case of material implanted with 1017 Ce ions cm 2, this effect should probably be more pronounced than for thelesswell-coveredunimplantedmaterial. The oxygen enrichment observed at the outermost layer of the implanted material (Fig. 3) most likely arises from the transformation of sulphides into oxides as the specimen is cooled to room temperature. During cooling the oxygen and sulphur activities of the gaseous environment decrease. This, however, does not influence the overall observations significantly since the cooling period is small compared with the exposure period at the test temperature. The Cr: Fe ratio of the outermost sulphide layer of the unimplanted material is perhaps too low to permit its transformation to oxide. Further investigations are underway to elucidate in greater detail the reasons for the better quality of the scale in the case of material implanted with 1017 Ce ions cm-2.
110 Acknowledgments
The authors wish to thank Dr. H. Peters of Leybold, Cologne, for providing the SNMS analysis. Mr. S. Deckers of the Physics Depart-
11
ment of the University of Utrecht is acknowledged for providing the RBS analysis. They are grateful to M r . D. J. Chivers and Mr. R. Izzard, U K A E HarweU, for carrying out the ion implan-
13
12
14
tation and TRIM calculations, respectively. They also wish to acknowledge the assistance of the staff of JRC Petten in particular Mr. G. yon Birgelen (XPS), Miss E Gandrey (Corrosion Group), Mr. Ph. Glaude (XRD), Mr. A. E. Moulaert
17
( E P M A ) a n d Mr. K, S c h u s t e r ( S E M ) . I n t h e p r e p -
18
aration of this manuscript, thanks are due to Mr. J. Blom (illustrations) and Mrs. L. Gosselink (text).
15
17. 16 G.J. Yurekand K. Przybylski, Mater.
19 2o 21 22
References 1 T. C. Tiearny, Jr., and K. Natesan, J. Mater. Energy Systems, 1 (1980) 13. 2 D. S. Williams, R. M611er and H. J. Grabke, Oxid. Met., 16(1981)253. 3 M. F. Stroosnijder and W. J. Quadakkers, High Temp. TechnoL, 4(1986) 141. 4 D. P. Whittle and J. Stringer, Phil Trans. R. Soc. London, Ser. A, 295 (1980) 309. 5 J. G. Smeggil, A. W. Funkenbusch and N. S. Bornstein, Metall. Trans. A, 17(1986) 923. 6 K.N. Stratford, High Temp. TechnoL, 1 (1983) 307. 7 J. Stringer, B. A. Wilcox and R. I. Jaffee, Oxid. Met., 5 (1972) 11. 8 T. N. Rhys-Jones, H. J. Grabke and H. Kudielka, Corros. Sci., 27(1987)49. 9 M. J. Bennett, H. E. Bishop, P. R. Chalker and A . T . Tuson, Mater. Sci. Eng., 90 ( 1987) 177. 10 V. Guttmann, R. C. Hurst, J. E Norton, O. Van der Biest and M. Van de Voorde, in I. G. Wright (ed.), Proc. Symp.
on Corrosion in Fossil Fuel Systems, Electrochemical Society, Princeton, 1983, p. 147. J. Lindhard, M. Scharff and H. E. Schott, Mat. Fys. Medd. Kgl. Dan. Vidensk. Selsk.,33(1963) 14. J. P. Biersack and L. G. Haggmark, Nucl. lnstrum. Methods, 174 (1980) 257. H. Oechsner, in H. Oechsner (ed.), Topics in Current Physics, Vol.37, Springer, Berlin, 1984, p. 63. Z. L. Liau and J. W. Mayer, in J. K. Hirvonen (ed,), Treatise on Materials and Technology, Vol. 18, Ion Implantation, Academic Press, New York, 1980, p. 17. D.I. Potter, M. Ahmed and S. Lamond, J. Met., 35 (1983)
23 24 25 26 27 28
29 30
Sci. Eng., 87(1987) 125. V. Nagarajan, R. G. Miner and A. V. Levy, J. Electrochem. Soc., 129(1982) 782. J. A. Kneeshaw, I. a . Menzies and J. F. Norton, Werkst. Korros., 38 (1987) 473. H.S. Hsu, Oxid. met., 28(1987) 213. s.F. Chu, E U Daniel, A. J. Blazewitz and R. E Dudek, J. Mater. Energy Systems, 7 (1986) 361. R. J. Hussey, P. Papaiacovou and H. J. Grabke, Werkst. Korros., 37(1986) 578. F. H. Scott, E M. E Chong and C. A. Stirling, in M. F. Rothman (ed.), High Temperature Corrosion in Energy Systems, Metallurgical Society of AIME, Warrendale, PA, 1985, p. 253, W. F. Chu and A. Rahmel, Oxid. Met., 16 ( 1981 ) 175. S. H. Choi and J. Stringer, Mater. Sci. Eng., 87 (1987) 237. D. Saxena, S. Prakash, M. L. Mehta and 1. P. Saraswat, Oxid. Met., 28(1987) 127. E.J. Vineberg and D. L. Douglass, Proc. Iron Steel Conference, Japan Institute of Metals, Sendai, 1983, p. 507. E. A. Polman, T. Fransen and P. J. Gellings, Mater. Sci. Eng., 88(1987) 157. H.J. Grabke, J. F. Norton and F. G. Casteels in W. Betz, R. Brunetaud, D. Coutsouradis, H. Fischmeister, T. B. Gibbons, I. Kvernes, Y. Lindblom, J. B. Marriott and D. B. Meadowcroft (eds.), Proc. Conf. High Temperature Alloys for Gas Turbines" and Other Applications, Liege, Reidel, Dordrecht, 1986, p. 245. P.Y. Hou and J. Stringer, Oxid. Met., 29 (1988) 45. S. Mrowec, A. Gil and A. Jedlifiski, Werkst. Korros., 38 (1987) 563.
103
Effect of Cerium Implantation on the Corrosion of Alloy 800H in an S - O - C Environment* M. E STROOSNIJDER, J. E NORTON and V. GUTTMANN
Commission of the European Communities, Joint Research Centre, Petten Establishment, Post Box 2, 1755 ZG Petten (The Netherlands') M. J. BENNETT
Materials Development Division, Atomic Energy Research Establishment, Harwell, Didcot (U.K.) J. H. W. DE WIT
Laboratory for Materials Science, Delft University of Technology, Postbox 5, 2600 AA Delft (The Netherlands) (Received September 16, 1988)
Abstract
The corrosion behaviour of a wrought austenitic Fe-20Cr-32Ni steel, Alloy 80OH, was studied in a simulated coal gasification atmosphere at 700 °C. Exposures ranged from as short as a few minutes to periods of up to 1000 h. The influence of cerium on the corrosion resistance was established by carrying out identical tests after ion implantation of the specimen surfaces. It was shown that the corrosive attack in the t12S-containing gas mixture can be reduced for the cerium-implanted material if the implantation dose is sufficiently high. In the case of both implanted and unimplanted material, sulphides and oxides were formed but the corrosion products of the implanted material showed a higher Cr.'Fe ratio relative to those of the unimplanted material. It is proposed that this feature is responsible for the observed improvement,
1. Introduction
In several technical high-temperature processes, such as during the gasification of coal, structural materials are exposed to atmospheres of low oxygen and high sulphur partial pressures. The ability of alloys to resist corrosion under these conditions depends on their potential to form and maintain a protective oxide scale, e . g .
*Paper presented at the Sixth International Conference on Surface Modification of Metals by Ion Beams, Riva del Garda, Italy, September 12-16, 1988. 0921-5093/89/$3.50
SiO 2 o r A1203. Unfortunately, in these atmospheres the simultaneous formation of sulphides as well as oxides is observed, although according to the thermodynamic equilibria, only oxides should be present. The exclusive formation of oxides occurs only beyond the so-called "kinetic boundary", which lies at a higher oxygen
Cr203,
activity than that predicted by thermodynamics [1-3]. In this context the role of rare earth elements is of interest. These elements may be incorporated in bulk alloys as metallic or oxide-dispersoid compounds or can be added to the surface layers by ion implantation. The addition of small amounts of rare earth elements to high-temperature alloys is known to improve the oxidation resistance and scale adherence [4-6], although the exact mechanism is still the subject of discussion. The early stages of nucleation and growth of the oxide scale seem to play an important role [4, 7-9]. Most of the work on the effect of rare earth elements has been undertaken in purely oxidizing environments. Only limited experimental work has been carried out in mixed oxidizing-sulphidizing environments. In this paper, the effect of cerium ion implantation on the corrosion behaviour of Alloy 800H (a wrought austenitic Fe-20Cr-32Ni alloy) in a simulated coal gasification atmosphere at 700 °C, is reported. The chemical potentials of the environment were selected so as to be of relevance to a moderately aggressive environment, i.e. a high tendency for oxide formation, without excluding sulphidation. © Elsevier Sequoia/Printed in The Netherlands
104
2. Experimental details The alloy investigated came from a commercial heat and by analysis contained 19.6 wt.% Cr, 32.2 wt.% Ni, 0.7 wt.% Mn, 0.45 wt.% Si, 0.46 wt.% Ti, 0.38 wt.% AI and 0.09 wt.% C. After solution annealing at 1150 °C for 0.5 h followed by water quenching, specimens in the form of coupons with dimensions 10 mm x 8 mm x 2.5 mm were machined. Prior to implantation they were abraded on 600 grit silicon carbide paper and degreased, The uniform implantation of cerium, at a nominal dose of either 1016 o r 1017 ions cm -2 into one of the principal coupon faces was undertaken using a 200 keV accelerating potential. The cerium distribution was determined experimentally using Rutherford backscattering spectrometry (RBS), secondary neutral mass spectroscopy (SNMS) and X-ray photoelectron spectroscopy(XPS). Corrosion tests were carried out in a hydrogen-based gas mixture with 7% CO, 1.2% H 2 0 and 0.2% H2S. At the test temperature of 700 °C and a total pressure of 1.4 atm, the sulphur partial pressure was 0.2 x 10 -9 atm, while the oxygen partial pressure was 0 . 7 x 10 -23 atm. The gas flowed continuously through the corrosion autoclaves at about 20 1 h- 1. Further details of the test facilities have been described elsewhere [10]. The specimens, each contained individually in a pure A I 2 0 3 crucible, were lowered into the reaction chamber at room temperature and then heated up in about 2 h to the test temperature in the test gas. The corrosion kinetics were established using discontinuous gravimetric measurements. For short-term tests a special autoclave system was used. This was fitted with a plunger device in order that samples could be rapidly inserted and retrieved from the hot gas atmosphere, thereby enabling the specimens to be exposed at temperature for periods of only a few minutes. To characterize the structure and composition of the corrosion products formed, a range of analysis techniques was used including optical metallography (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), electron probe microanalysis (EPMA), XPS and SNMS.
3. Results
3.1. Characterization of the implant distributions An estimate of the initial distribution of the implanted atoms in a material is possible using
the Lindhard-Scharff-Schott theory [11]. However, in the case of high implantation doses the observed implant distribution may deviate significantly from the theoretically predicted distribution, for example because of material loss due to sputtering. Computer modelling can take this into account. Distribution profiles calculated for the implantations of 1016 and 1017 Ce ions cm-2 in Alloy 800H are presented in Fig. 1. These profiles were obtained by the Monte Carlo computer program TRIM developed by Biersack and Haggmark [12]. It was calculated that the retained dose approximately equals the nominal dose in the situation of the low implantation dose. However, for the high implantation dose the retained dose (4 × 1016 C e cm -2) was only 40% of the nominal dose due to the sputtering loss of 80 nm. The actual implant distributions were determined experimentally using SNMS, RBS and XPS (Fig. 1). For the cerium distribution obtained by SNMS, standard sensitivity factors [13] were used. For SNMS the sputter rate was taken as 0.1 nm s- 1, whereas for XPS it was about 0.05 nm s- 1. The chemical state of cerium was difficult to define from the XPS data due to chemical reduction of the surface to a lower oxide state during ion etching.
. ~*~ _~~ ~
.
.
. Attoy800H
1016Ce ions/cm 2
t
§2 § 6° 20
40
60
80
100
Depth. nrn
(a) 30 . . . . ~ AHoy800x "~ 10'7ce~°°s/em2 ~2 0 ~ ~ Ps ; ~0 / % , L . , C o,o,,0
°,
,'
"7"6'o Depth, n m
'
'
1oo
Fig. 1. Experimentally determined and calculated cerium distributioon profiles for Alloy 800H: (a) 1016 ions cm-2;
(b) 1017ionscm-2.
105
3.2. Gravimetric determinations For mass gain measurements a set of three specimens (unimplanted, 1016 Ce ions cm -2 and 1017 Ce ions cm -2) was exposed for 200 h in the test gas, whilst another set was exposed intermittently for 25, 100, 250, 500, 750 and 1000 h. Because the ratio of implanted to unimplanted area was not precisely known it was not possible to obtain accurately the influence of implantation on mass gain and the measurements therefore only provided an approximate indication of the corrosion resistance. The situation was complicated further by the relatively large scatter observed in the case of low mass gains. The overall conclusion from these experiments was that the material implanted with 1016 Ce ions cm -2 exhibited a similar mass gain to the unimplanted material, whereas for the alloy implanted with 1017 Ce ions cm 2 alloy, the mass gain was reduced by a factor of about 3. The alloy implanted with 1017 Ce ions cm-2 which was successively exposed for durations up to 1000 h, also showed a lower mass gain compared with the Alloy 800H, although the difference became smaller with increasing exposure time.
3.3. Nucleation andgrowth studies In order to analyse the structural changes occurring during corrosion, additional specimens were exposed for 2, 5, 15 and 30 min, as well as 5 h. The behaviour of the implanted and unimplanted sides on the same specimen were directly compared. In all experiments the specimens implanted with 1016 Ce ions cm -2 showed identical corrosion behaviour to the unimplanted alloy and therefore will not be considered in further detail, The SEM micrographs in Fig. 2 depict the changes in scale morphology that occurred between exposures of 2 min and 200 h for the unimplanted alloy (Figs. 2(a)-2(c)) and the Alloy 800H implanted with 1017 Ce ions cm -2 (Figs. 2(d)-2(f)). XPS analyses showed that the corrosion products which formed initially on the unimplanted alloy were a mixture of mainly iron and chromium sulphides and small amounts of oxides, The lighter features in Fig. 2(b) initially became evident after 5 rain and increased significantly in number and size during the first hours of exposure. With continuing time they grew in size although their number decreased due to coalescence. These features were shown by XRD and
EDX to be sulphides of the types M3S4 and MS and to contain iron and chromium with the iron content increasing with the size of the sulphide. The sulphur content was higher than for the corrosion layer which formed initially. XRD analysis after exposures of 30 rain or longer indicated the presence of M304 for the unimplanted material. The lattice parameter of the M304 was consistent with a (Cr, Fe)304 spinel. The scale formed on Alloy 800H implanted with 1017 Ce ions cm -z after a few minutes showed a less rough morphology than that on the unimplanted alloy. From Fig. 2 it appears that the originally formed individual sulphide nuclei coalesce to a large extent thereby forming a fairly homogeneous flat scale covering the underlying alloy. This morphology was confirmed indirectly by XPS element profiling and comparing the scales formed on the implanted and unimplanted alloy after 5 min exposure. The implanted material gave a lower nickel signal (from the substrate) and a higher sulphur signal because of a more denser surface layer. Moreover, from observation of the sputter time necessary to reach the uncorroded alloy in the XPS analysis, it appeared that the corrosion layer formed on the implanted alloy was thinner than that on Alloy 800H. From the XPS data it may be concluded that the surface sulphides formed initially on the implanted alloy contained a higher Cr:Fe ratio than those on the unimplanted material. An oxygen enrichment under the predominantly sulphide scale, positioned slightly below the cerium peak, was evident for implanted material from the XPS analyses. The morphology of the original surface layer of the material implanted with 1017 Ce ions cm -2 appeared to change with exposure time from a relatively flat layer to a more crystalline morphology. Again, with increasing exposure time the formation of light features above the initial sulphide scale was observed on the implanted material (1017 Ce ions cm -2) but their appearance was clearly delayed and became evident only after 30 min exposure. EDX analysis showed that these features had a similar composition for both the implanted and the unimplanted alloy. The number of features increased with time, but for durations of up to 200 h their size was significantly smaller than for the unimplanted material and no coalescence was observed. XRD analyses indicated the same type of corrosion phases to be present on the implanted as on the unimplanted alloy, i.e. MS, M3S4 and M304
106
Fig. 2. Surface scale morphology of (a), (b), (c) Alloy 800H and (d), (e), (f) Alloy 800H implanted with 1017 Ce ions cm -2 after exposure for (a), (d) 2 rain, (b), (e) 15 rain and (c), (f) 200 h in H2-7voI.%CO-0.2voI.%H2S- 1.2vol.%H20 at 700 °C.
107
(obviously also (Cr,Fe)304). After 1000 h exposure the corrosion appeared to be similar for both unimplanted and implanted Alloy 800H. Element depth profiles obtained after 200 h exposure by SNMS of an area 3 mm in diameter through the scale for the unimplanted and implanted (1017 Ce ions cm 2) material are presented in Fig. 3. The sputter rate was about 0.5 nm s- i. It should be pointed out that although the relative intensity in Fig. 3 is indicative of the atomic percentage of the elements, average sensitivity factors [13] were taken for the calculations and no standards for calibration were applied, Hence the data cannot be used in a fully quantitative manner. The profiles (Fig. 3(a)) indicate for the unimplanted material only weakly pronounced peak positions for the various elements, The ion content was found to be higher than the chromium content throughout the entire scale. By comparison, Fig. 3(b) illustrates clearly separated
60
.
0<~0
.
.
.
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~ .
.
.~ .
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.
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8
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.
700oC/200h
.~ 30 _g 20 10 0
.
Alloy 800H
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1~
16
. . . .r . . AL{oy8QOH÷ 1017Ceions/cm2 700°CnOOh S-O-C gas
so
Fe
~0 30 20
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;
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8
peak positions for the implanted material. The high sulphur concentration is attributed to chromium and iron which were both present at similar levels. A region of oxygen enrichment observed beneath the sulphide layer was related to an increase of the chromium level with a concomitant reduction in iron. The cerium peak lies slightly above this oxygen enrichment. A marked oxygen concentration of the outermost layer cannot be due solely to impurity adsorption. For both materials the manganese content of the scale was found to be low. Cross-sectional examinations by OM and SEM (Fig. 4), revealed that the scale formed on the material implanted with 10 Iv Ce ions cm- 2 was more uniform in thickness than that on the unimplanted material. This was also found during the early stages of corrosion as well as after 200 h exposure. EDX spectra confirmed a higher Fe:Cr ratio for the scale on the unimplanted than on the implanted alloy. After 5 h exposure the untreated alloy exhibited a significant number of surface protrusions and also internal corrosion appeared. In contrast, the implanted material (1017 ions cm -2) after 5 h exposure showed almost no localized external nor localized internal corrosion. After 200 h both internal and external local corrosion was detected for the implanted material, although it was much less advanced than for the unimplanted sample. Attempts to perform quantitative EPMA failed due to the relatively bad lateral resolution and large analysis depth of this technique. The qualitative EPMA results of the outer scale formed after 200 h on the unimplanted material indicated essentially sulphur, iron and chromium with an increasing iron content in the larger nodules towards the outer surface. An oxide-rich zone appeared below this sulphide layer which contained chromium and iron. The internal corrosion products were mainly composed of Cr-Fe oxide and showed additionally an enrichment in silicon, aluminium and titanium. A low sulphur content, which increased towards the tip of the internal corrosion paths indicated the presence of small amounts of individual sulphides. EPMA of the scale formed
r 10
12
1/.
1(5
Sputter,ngtime.ks
Fig. 3. SNMS element distribution profiles of (a) Alloy 800H and (b) Alloy 800H implanted with 10 j7 Ce ions cm 2 after exposure for 200 h in H2-7vol.%CO-0.2vol.%H2S1.2vol.%H20 at 700 °C.
after 200
h exposure
confirmed
the SNMS results from the implanted alloy, i.e. the scale was rich in sulphur, chromium, iron and oxygen with an increased oxygen level towards both the gas-scale and the substrate-scale interfaces. The external nodules were Fe-Cr sulphides
with the few internal corrosion observed also containing manganese.
products
108
i:',:~'~
ii!:i:
~~! : ()~
...........
Fig. 4. Cross-sectional SEM examinations of (a), (b) Alloy 800H and (c), (d) Alloy 800H implanted with 1017 Ce ions cm -2 after exposure tot /a), (c) 5 h and (b), (d) 200 h in H2-7vol.%CO-0.2vol.%H2S-1.2vol.%H2O at 700 °C.
4. D i s c u s s i o n and c o n c l u s i o n s
The experimentally observed distributions and the total uptake of cerium in Alloy 800H (Fig. 1 ) for the high implantation dose were close to those expected from the Monte Carlo calculations. This agreement on the alloy implanted with 1017 Ce ions cm-2 shows clearly that the phenomenon of sputtering can significantly limit the maximum concentrations that can be achieved by implantation, especially in the case of high implantation doses. This agrees with the work of others [14, 15]. The investigations of the corrosion behaviour have shown that the attack in the S - O - C gaseous atmosphere can be reduced for the cerium implanted material if the implantation dose is high. The specimens with only 1016 Ce ions cm -2 show an identical behaviour, in terms of both the mass gain and scale formation, to the unimplanted material. Implantation with nominal 1017 Ce ions cm-2 leads to a decrease in mass gain by a factor of up to 3 and a clear retardation of the corrosion rate, at least for durations of up to 200 h. It should be noted that for a 1000 h test a very similar behaviour has been observed on the implanted and unimplanted alloy. However, the influence of temperature cycling, due to the interruptions of the test, may have contributed to this behaviour. Studies by SEM and various analytical techniques have shown a fairly classical corrosion pattern for the unimplanted alloy [1, 16-18]. Briefly, in mixed oxidizing and/or sulphidizing atmospheres, both oxides and sulphides may nucleate on the alloy surface even though thermodynamic considerations indicate that the oxide is the most stable phase. This is due to the fact that kinetic factors may dominate the overall corrosion process and sulphides may overgrow the oxide nuclei. No protective oxide will be formed in this case. This behaviour has also been observed in the present investigation. Although oxide stability would be predicted from thermodynamic considerations of the gas composition, it was found that after short exposures a high density of individual sulphides rich in iron and chromium were present but only a small amount of oxide could be detected by XPS. This behavtour shows that under the test conditions eraployed, sulphidation is the dominant process. The observation of internal corrosion indicates that the oxides do not form a protective layer.
109
Poor protection is also confirmed by the formation of fast growing iron-rich sulphides on the initially formed sulphide layer. Iron cation transport to the scale-gas interface appears to be facilitated locally which probably results from the formation of sulphide-rich channels which are known to be fast diffusion paths [19-22]. It is also known that the diffusion of iron through these channels is much faster than that of chromium [16, 23]. It should be noted that the internal corrosion products were primarily oxides and that for durations of up to 200 h exposure only a few sulphides have been observed at the base of the corrosive attack. This feature can be related to the release of sulphur caused by the transformation of sulphides into oxides, as has also been observed by others [18, 21, 24]. Regarding the corrosion behaviour of the material implanted with 10 ~7 Ce ions cm-2, it is somewhat surprising that although there was a high tendency for sulphidation to occur during the initial stage of exposure, the corrosion resistance in general was found to be improved. In comparison with the unimplanted alloy the sulphides initially formed appear to exhibit a great tendency to grow laterally with a consequent coalescence and a higher coverage of the underlying alloy. Apart from this difference in morphology, the sulphides formed on the material implanted with 1017 Ce ions cm -2 exhibited a higher Cr:Fe ratio than those on the untreated material. Similar observations derived from sulphidation studies of a Fe-15Cr-4AI-1Y alloy in comparison with Fe-15Cr-4A1 [25] and on the sulphidation of N i - C r - Y alloys [26]. In contrast to the unimplanted alloy, the nucleation and growth of local sulphides above the coherent sulphide layer which formed initially was strongly retarded for the alloy implanted with 10 Iv Ce ions cm-2. The Cr:Fe ratio of these local sulphides was also higher than for the unimplanted material. An oxide layer had again formed beneath the chromium-rich sulphide layer, comparable with the situation of the unimplanted material. From the SNMS investigations after 200 h exposure, it can be concluded that the oxide layer had a relatively high chromium content compared with that of iron. From the XPS and SNMS analyses it became apparent that the implanted cerium was concentrated on top of the oxide layer. This indicates that the oxide layer had grown preferentially by oxygen anion transport.
The mechanisms governing the improvement in corrosion resistance by the implantation of 1017 Ce ions cm -2, demonstrated first by the retardation in the development of individual large sulphides on top of the initially formed sulphide layer and second by the absence of internal corrosion at least up to 200 h, unfortunately could not be fully established. It is probable that the oxide layer which formed beneath the sulphides plays a decisive role. In general it is well known that chromium-rich oxide layers confer greater protection than that provided by iron-rich spinel phases [22, 27, 28]. Indeed, the present investigations have indicated a higher chromium content for the oxide on the implanted alloy compared with that on the unimplanted material. This is consistent with observations made in purely oxidizing atmospheres [4, 8, 29, 30] that materials with rare earth elements show a stronger tendency for chromia formation. Generally speaking, all the external corrosion products observed on the implanted alloy showed a relatively high chromium content compared with that of iron. It may be argued that, in addition, the morphology of the initial sulphide layer formed on the alloy implanted with 10 ~7 Ce ions cm -2 may have a beneficial influence on the oxide formation. Taking into account that the sulphur potential at the alloy-sulphide scale interface must be lower than that in the corrosive test gas, this will lead to a relative increase in the oxide-forming potential. Because of the improved coverage of the alloy surface by sulphides in the case of material implanted with 1017 Ce ions cm 2, this effect should probably be more pronounced than for thelesswell-coveredunimplantedmaterial. The oxygen enrichment observed at the outermost layer of the implanted material (Fig. 3) most likely arises from the transformation of sulphides into oxides as the specimen is cooled to room temperature. During cooling the oxygen and sulphur activities of the gaseous environment decrease. This, however, does not influence the overall observations significantly since the cooling period is small compared with the exposure period at the test temperature. The Cr: Fe ratio of the outermost sulphide layer of the unimplanted material is perhaps too low to permit its transformation to oxide. Further investigations are underway to elucidate in greater detail the reasons for the better quality of the scale in the case of material implanted with 1017 Ce ions cm-2.
110 Acknowledgments
The authors wish to thank Dr. H. Peters of Leybold, Cologne, for providing the SNMS analysis. Mr. S. Deckers of the Physics Depart-
11
ment of the University of Utrecht is acknowledged for providing the RBS analysis. They are grateful to M r . D. J. Chivers and Mr. R. Izzard, U K A E HarweU, for carrying out the ion implan-
13
12
14
tation and TRIM calculations, respectively. They also wish to acknowledge the assistance of the staff of JRC Petten in particular Mr. G. yon Birgelen (XPS), Miss E Gandrey (Corrosion Group), Mr. Ph. Glaude (XRD), Mr. A. E. Moulaert
17
( E P M A ) a n d Mr. K, S c h u s t e r ( S E M ) . I n t h e p r e p -
18
aration of this manuscript, thanks are due to Mr. J. Blom (illustrations) and Mrs. L. Gosselink (text).
15
17. 16 G.J. Yurekand K. Przybylski, Mater.
19 2o 21 22
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