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The oxidation behaviour of heat resisting metallic fibres
Corrosion Science, Vol. 23, No. 9, pp. 1025-1043, 1983 Printed in Great Britain.
0010-938X/83 $3.00+0.00 Pergamon Press Ltd.
THE OXIDATION BEHAVIOUR OF HEAT RESISTING METALLIC FIBRES S. R. J. SAUNDERS and I. LEFEVER* National Physical Laboratory, Queens Road, Teddington, Middlesex, U.K. Abstract--The behaviour of 22 vm diameter alloy fibres of Type 310 stainless steel, Hastelloy X, Inconel 601 and Fecralloy has been studied at 600°C in combustion gases with and without HCI addition. Chromia-forming alloys exhibited higher rates of attack than would have been expected from a consideration of the properties of the bulk material whereas normal rates of attack were observed for alumina-forming alloys. The addition of HC1 to the combustion gases increased rates of attack for all alloys and generally this was associated with increased sulphur in the oxide layer. The results are discussed with reference to the unusual geometry of the samples, and it is suggested that the direction of motion of the ions in the growing oxide layer is an important factor governing its protective nature. Alloy depletion effects were also noted, particularly for some Fecralloy samples which were pre-oxidized before exposure, and this may cause eventual failure of the protective oxide. A common feature of the oxidation of iron--chromium alloys is the partitioning of elements in the scale; this wag especially marked for the Type 310 stainless steel samples exposed in this work and a stressassisted diffusion process is proposed to explain this effect. INTRODUCTION THIS paper describes the corrosion resistance of metallic fibres, with a diameter of 22 ~m, sintered into a high porosity structure. These fibres, Bekinox®, made by NV Bekaert SA, are available in Type 316L stainless steel, Hastelloy X, Inconel 601, nickel and titanium. The sintered products, the reinforced web and felts and Bekipor®, find applications in dust filtration, liquid filtration, electromagnetic shielding, sound suppression, flame arresting etc. These products can also be used for high temperature gas filtration, e.g. in automotive emission control, in power plants arid in other combustion processes. The requirement to use coal as an alternative fuel to oil has increased the need to develop suitable materials for hot-gas filters. In addition to resisting attack by hot gas, these filters must be capable of removing particles greater than about 10-20 ~tm and able to withstand mechanical shocks during a cleaning cycle. Generally, ceramic materials are not suitable because of inadequate mechanical properties. The filtration characteristics of the porous sintered fibre structures are determined by the fibre diameter and packing density and fibres with diameters between 4 and 22 ~tm can now be produced on a commercial scale. The purpose of the work described was to assess the high temperature oxidation/ sulphidation resistance of these thin fibres which may differ from that of bulk materials because of the large surface area-to-volume ratio. This could result in the rapid depletion of the alloy in additions such as chromium and aluminium that are required for oxidation resistance. Furthermore, the surface state and/or metallurgical structure of these fibres may also have a significant effect on their behaviour. In most cases, *Bekaert Research Centre, NV Bekaert SA, Leo Bekaertstraat 1, 8550 Zwevegem, Belgium. Manuscript received 15 February 1983. ~)1983 Crown copyright. 1025
1026
S.R.J. SAU~0EP.Sand I. LEF~WR
Type 316L stainless steel fibres perform satisfactorily for filtration purposes at temperatures up to about 350°C. The work reported in this paper describes the results of tests on a number of heat-resisting alloy fibres in combustion gas at 600°C with and without the presence of HC1 gas.
EXPERIMENTAL METHOD
Specimen preparation The specimens were made of fibres in four heat-resisting alloys: stainless steel Type 310, Incone1601, Hastelloy X and Fecralloy and the compositions are given in Table 1. The fibres were produced by NV Bekaert SA according to a proprietary process. TASL~ 1.
NOmNAL COMPOSITION OF THE MA1~RIAI.S TESTED (wt ~o)
Type 310 stainless steel
C 0.25max; S 0.03max; Ni 19-22;
Mn 0.20 max; Si 1.5 max; balance Fe
P 0.045max Cr 24-26
Inconel 601
C 0.05; Si 0.25; Cu 0.50;
Mn 0.5; Cr 23; Fe 14-15;
S 0.007 AI 1.35 balance Ni
Hastelloy X
C 0.10; Cr 22; W 0.6;
Mn 0.5; Co 1.5; Fe 18.5;
Si 0.5 Mo 9.0 balance Ni
Fecralloy
C 0.03; AI 4.8; balance Fe
Si Y
Cr 15.8 S 0.004
0.3; 0.3;
Sintered sheets of the four alloys, made from 22 ltm diameter fibres, with a nominal porosity of 80 ~. and a weight of 450-550 g m-L were cut into 25 x 70 mm samples. Some of the Fecralloy samples were given a special preoxidation treatment after sintering. All samples were then degreased in acetone and methanol and weighed prior to testing. Fibre shedding was observed during the first test and this caused difficulties in the interpretation of the gravimetric results, and thus, for subsequent work, loose fibres, particularly from the cut edges, were removed by gentle agitation.
Test procedure The specimens were exposed at 600°C to hot gases obtained by combustion of a 1 ~ S (by weight) fuel oil (BS 2869 Class D) at an air to fuel ratio of 30 : 1 using a burner rig. Complete combustion gave a SOt content of 280 ppm. Addition of 100 ppm HC1 gas to the combustion gas was achieved by atomizing a 0.1 M HCI solution flowing at 0.60 ml rain-a with a fuel flow rate of 0.77 ml min-a. The specimen holder, made from Type 310 stainless steel, which accepted 24 samples was rotated slowly to achieve uniform exposure of the samples to the test environment. Figure 1 shows typical specimens mounted in the specimen holder. Thermal cycles were imposed every 22 h by withdrawing the samples from the burner rig and allowingthem to cool to room temperature. Samples were weighed at approximately I00 h intervals during the test.
Specimen evaluation Some samples were examined by scanning electron microscopy and Auger spectroscopy. Auger analysis was carried out using a cylindrical mirror analyser with a coaxial electron gun at a mean spot diameter of 80 lain. Samples for optical metallography were mounted and polished using conventionalmetallographic techniques. A Cameca Camebax analyser operating at an accelerating voltage of 15 kV and a beam current of 10-g A was used for electron microprobe analysis.
FIG. 1. Partially loaded specimen holder.
9 LQ,
A
"8
FIo. 4. Optical micrographs of fibre samples after 528 h at 600°C in (A) clean combustion gases, and (]3)combustion gases q- 100 ppm HCI. (a) Type 310 stainless steel; (b) Inconel 601 ; (c) Hastelloy X; (d) Fecrailoy, untreated; (e) Fccralloy, treated.
FIG. 5. Scanning electron micrographs of fibre samples after exposure to clean combustion gases at 600°C (a) Type 310 stainless steel, 1056 h; (b) Incone1601, 1056 h; (c) Hastelloy X, 528 h; (d) Fecralloy, untreated, 1056 h; (e) Fecralloy, treated, 1056 h.
10 IJm L.._._.J
FIG. 6.
Microprobe analysis of a Type 310 stainless steel fibre after 1056 h at 600°C in (a) combustion gases; and (b) combustion gases + 100 ppm HCi.
J
oo
3~
-
.
-=
Cb)
10 pm I
FIG. 8.
.I
Microprobe analysis of an Inconel 601 fibre after 1056 h at 600°C in (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
FIG. 9. Microprobe analysis of uncorroded Fecralloy fibres (a) not sinterod; (b) sintered.
lOpm,,
(.)
lOpm,,
(b).
lOpmt,~
lOpm~,~ ,
(c)
(d)
FIG. 10. Microprobe analysis of Fecralloy fibres after 1056 h at 600°C. (a) Untreated, clean combustion gases; (b) treated, clean combustion gases; (c) untreated, combustion gases -k 100 ppm HCI; (d) treated, combustion gases ÷ 100 ppm HCI.
The oxidation behaviour of heat resisting metallic fibres
1035
EXPERIMENTAL RESULTS
Gravimetric data Figure 2 shows the gravimetric data obtained in the test without HCI addition. The curves dearly show the superiority of the Fecralloy and Inconel 601 materials compared with the Hastelloy X and Type 310 stainless steel. The weight loss observed for the Fecralloy specimens was due to fibre shedding during specimen handling. It is probable that all other specimens also lost weight by this process, but that the higher weight gain due to corrosion masked the effect, and, thus, all the gravimetric data slightly underestimated the corrosion rate. Figure 3 shows the influence of the 100 ppm HC1 addition which resulted in a greater than two-fold increase in the weight gain in the case of Type 310 stainless steel, a less marked increase for the Inconel 601 and Fecralloy samples but, somewhat unexpectedly, a lower weight gain for the Hastelloy X specimens.
Optical and scanning electron microscopy Optical microscopy of polished cross-sections of the corroded samples confirmed the overall trends observed from the gravimctric data. Figure 4 shows the behaviour of samples of all the alloys after 528 h in both types of test condition. It was evident that whereas the Hastdloy X and Type 310 stainless steel specimens had been extensively attacked, the Fecralloy and Inconel 601 samples were only slightly attacked, a result consistent with the gravimetdc data. These micrographs were taken from areas exhibiting the most severe attack and illustrate the quite variable behaviour of the individual fibres, particularly for Type 310 stainless steel and, to a lesser extent, for Incon¢l 601. For example, Fig. 4(a) shows a Type 310 stainless steel sample almost completely oxidized alongside a fibre showing virtually no attack. The internal attack observed in some Fecralloy samples (Figs. 12 r 10]/ |
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[] Fecralloy, untreated o Fecralloy, pre-oxidized • Inconel 601 •
Type 310 stainless steel
• Hastelloy X
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•
~
•
•
•
,
800
1000
0
-2
200
400
600 Time (h)
Fie. 2. Weight change as a function of time for 22 IJm diameter fibre samplesof Type
310 stainless steel, HastelloyX, Inconel 601 and Fecralloyexposedto combustiongases at 600°C.
S . R . J . SAUNDEP.Sand I. L~'EvER
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Fecralloy, untreated Fecralloy, pre-oxidized Inconel 601 Type 310 stainlesssteel Hastelloy X
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Time (h) FIG. 3. Weight change as a function of time for 22 pm diameter fibre samples of Type
310 stainless steel, Hastelloy X, Fecragoy and Incone1601 exlmsed to combustion gases
at 600°C with 100 ppm HCI addition. 4d and e) was also found in the uncorroded fibres and thus was not a consequence of exposure to the hot combustion gases; the origin of this will be discussed later. Scanning electron micrographs of the corroded fibres are shown in Fig. 5. Again it is evident that the Hastelloy X and Type 310 stainless steel samples have been severely attacked compared with those of Inconel 601 and Fecralloy. Of particular interest, however, is the appearance of the Type 310 stainless steel (Fig. 5a) which is consistent with the very heterogenous attack observed in cross-section. The structure of the fibre was such that an array of single crystals is formed ("bamboo" structure), and it is evident that attack was specific to certain grains, and also that grain boundary regions in the fibre were generally protected.
Microprobe analysis Microprobe analysis of some corroded fibres was carried out. Generally only the more severely attacked fibres were examined. Figure 6 compares the behaviour of Type 310 stainless steel after 1056 h in both test conditions. Virtually complete oxidation had occurred in both cases, leaving a compact cylinder of oxide with almost no trace of sulphur- or chlorine-containing compounds. Chromium generally was confined to the initial dimensions of the fibre but, in the case of the sample corroded in the absence of HCI, an outer layer of chromium-containing material was also present. An outer layer of iron-rich oxide formed on both specimens. Figure 7 shows the microprobe analysis of Hastelloy X samples after 528 h in the two test conditions. In both cases, a chromium-rich oxide containing molybdenum and sulphur formed at the metal/oxide interface. However, for the test in the absence of HC1, sulphur was also associated with the outer layer which contained nickel oxide. When HCI was added to the environment iron oxide was dominant in the outer layer.
The oxidation behaviour of heat resisting metallic fibres
1037
Inconel 601 and Fecralloy contain aluminium as an alloying addition and thus AlcOa formation is possible. Inconcl 601, with only 1.35% A1 formed both Cr~O8 and AlsOs (Fig. 8), when HC1 was absent, but only AlcOa could be detected on the sample from the test environment containing HCI. Sulphur was more evident after exposure in the test without HC1, but this could have been due to easier detection in the thicker scale. Microprobe analysis of untreated FecraUoy fibres that were uncorroded and unsintcred indicated the presence of an yttrium-rich intcrmetallic with the composition Fe6.8 Alz.5 Crl.s Y. During sintering in a vacuum furnace at 10 -6 torr and 1100°C, the intermetallic compound reacted with sulphur, as is indicated in Fig. 9. Figure 10 shows the behaviour of the FccraUoy samples after 1056 h in each test condition where it can be seen that some oxidation of the yttrium sulphide had occurred, especially for the untreated samples (Figs. 10a and c), and a thin layer of AI~O~ developed on the fibre surface in association with yttrium-containing material. The presence of HCI encouraged incorporation of sulphur into the scale but no chlorine-containing material was detected. The behaviour of the pre-oxidized samples (Figs. 10b and d) which had a very well-developed layer of AlcOa was generally similar, but there was less evidence of sulphur pick-up. The yttrium sulphide was also largely unoxidized.
Auger analysis Figure 11 shows the oxygen profiles obtained from samples of all alloys after testing in the two environments and this gives an indication of the extent to which the ion etching process proceeded relative to the oxide thickness. For example, on the Hastelloy X and Type 310 stainless steer samples, ion etching for 5 h was insufficient to penetrate the scale whereas that on the untreated Fecralloy sample was penetrated in about 3 h 20 rain. These profiles confirm the lower weight gain of Hastelloy X in the HCl-containing environment. Also, with the Fecralloy and Incon¢l 601 samples the Auger analysis indicates that the presence of HC1 resulted in slightly thinner scales OX~G[PI {o)
-- ~ ........
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FIG. 11. Oxygen concentration profiles through corroded fibres after exposure at 600°C to (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
1038
S.R.J. SAUNVeP.Sand I. LI~FEVER
whereas the gravimetric data suggested the opposite behaviour, but this was probably related to less fibre shedding of the samples. Little or no chloride was detected in the scales, a result in agreement with the electron microprobe analysis results. The presence of HCI generally caused an increase in the amount of sulphur in the scales, particularly for Inconel 601 and the preoxidized Fecralloy samples, as shown in Fig. 12. This does not agree with the microprobe results which show the opposite behaviour, but because these scales were thin the microprobe results may be in error. Hastelloy X fibres, however, had a deeper penetration of sulphur in the surface scale when HC1 was absent, a result in agreement with the microprobe results. ~0'
SULPHUR (0)
.....
60
..... e
Fecrafloy.untrt~ led, 1056h FectaUoy, treated,
26&h
Fecrolloy, treoqed. AISI 3~0 Intone| 610
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FIG. 12. Sulphurconcentration profiles through corroded fibres after exposure at 600°C to (a) clean combustiongases; (b) combustiongases + 100 ppm HCI. The aluminium content of the oxide formed on the pre-oxidized Fecralloy sample decreased with increasing exposure time, and the addition of HCI to the test environment further decreased the aluminium content of the scale for the untreated as well as for the pre-oxidized sample (Fig. 13). In all cases aluminium in the scale was replaced by iron and chromium. However, HCI did not influence the aluminium content of the scale that formed on Inconel 601. The nickel content of the scale was significantly reduced by HCI gas in the case of the Hastelloy X samples, but there was little effect on Inconel 601 and Type 310 stainless steel. DISCUSSION
Corrosion behaviour of alloy fibres The results of corrosion testing of the four types of alloy indicated that resistance to attack increased in the following order; Type 310 stainless steel, Hastelloy X, Inconel 601 and Fecralloy. Pre-oxidation of the Fecralioy fibres further enhanced their corrosion resistance. Data on the oxidation rate of bulk Incone1601 and Fvcralloy are not readily available at this relatively low temperature of 600°C, but the observed
The oxidation behaviour of heat resisting metallic fibres
1039
ALUMtNiUM (o)
61Ii/"' X 50
"*
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t/
.....
Fecrolloy,
untreoted
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A LUMINIUM (b)
/-
I056h 26/* h 1056h
..... .....
50
* I .
..........
Fecralloy. untreated
0 h
o
Feerolloy, Ir eotee.
264h
Fecrotto y, tr tGtee.
IOS6h
Intoner 601
IOS6h
40.
" ¢"
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Fecrnlloy. untreated, tOSg h
"" ~o,
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35'
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100 E t c h t me rain ZOO u f i
300
01.6
1.'2
1.0
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10
DtstcJnce, p m
12
Distance,~m
FIG. 13. Aluminium concentration profiles through corroded fibres after exposure at 600°C to (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
low rate of attack on the fibre products indicated that the oxidation resistance was not significantly different to that expected from the bulk material. Again, data on the rates of attack of Type 310 stainless steel and Hastelloy X at 600°C are not extensive, but extrapolation of higher temperature data indicated that these alloy fibres were attacked at much higher rates than thicker material. For example, Hastelloy X after I000 h exposure in air at the reduced pressure of 8 torr at 760°C showed a metal loss of about 4 ~m, 1 that is a similar rate of attack to that represented by the present results which were obtained at a much lower temperature. Also Type 310 stainless steel samples were reported to lose only about 2 ~m in thickness when exposed to air for 1000 h at 750°C. ~ Clearly, the presence of small amounts of sulphur dioxide in the combustion gas used in the current work would increase the rate somewhat compared with attack in air, but it is unlikely that this would be the only factor contributing to the very high rates of attack observed on these fibres. Similarly, the thermal cycling could also increase rates of metal wastage, but there was only slight evidence of spalling, even when HCI was present. It is noteworthy that the fibres showing higher than normal rates of attack are from chromia-forming alloys whereas the fibres of alumina-forming alloys, viz. Inconel 601 and Fecralloy, had low rates of attack. Important considerations in understanding the behaviour of these fibres are the very large surface to volume ratio and the rapidly increasing surface area when expansion occurred due to incorporation of oxygen. Chromic oxide grows mainly by outward migration of cations s whereas alumina forms predominantly by inward diffusion of oxidant down the oxide grain boundaries, although it has also been proposed that lattice diffusion of cations is involved in the oxide growth process.4 Thus, for the fibres examined in this work, which have a very small radius of curvature, the formation of a protective scale of
1040
S.R.J. SAtn~DEKSand I. LErtWR
chromia by outward migration of cations would be very difficult because the rapidly increasing surface area would tend to "dilute" the chromium. Furthermore, with the alloy surface rapidly receding, adhesion of the scale to the alloy would be difficult. In the case of alumina growing largely by inward migration of anions however, a rapidly decreasing surface area would aid the formation of a coherent and compact protective layer, although, if a thick scale were to form, a large compressive stress would develop that would ultimately lead to failure. It is believed, therefore, that protective oxide formation on thin fibres is most likely where inward diffusion of anions is the dominant transport process in the growing oxide layer. This criterion would be satisfied when alumina-forming alloys are used.
The effect of HCl<~) additions Auger analysis indicated that addition of 100 ppm HCI~g)to the combustion gases decreased the scale thickness on Hastelloy X, Inconel 601 and possibly the Fecralloy samples, although, for the two latter alloys, the effect was very small and may have been related to errors resulting from simultaneous analysis of several randomlyoriented fibres. Gravimetric and metallographic evidence clearly demonstrated, however, that HCI encouraged attack on the Type 310 stainless steel. In cases where the oxide was thinner in the presence of HCI, this did not necessarily correspond to lower rates of attack because some of the reaction products may have been volatile. For example, with the Hastelloy X samples, the corrosion product formed in the presence of HC1 contained significantly less nickel than that found on the sample corroded in the absence of HCI. Presumably, volatile nickel chloride formed and was removed from the system, resulting in a thinner oxide layer. The rate of metal loss, even in the case of HasteUoy X, was insufficient to establish this mechanism because uncertainties in the initial dimensions of the fibres did not allow accurate assessment of residual metal thickness. Furthermore, chlorides are known to encourage oxide spalling6 which could also result in reduced oxide thickness for the samples exposed to the test environment containing HC1. For example, in the case of Type 310 stainless steel samples shown in Fig. 8, an outer chromium-containing layer was observed on the specimen taken from the test with no HCI present but was absent from the sample exposed to HC1. It is suggest that, because of the "dilution" effects described earlier, the initially formed layer never became protective and remained on the surface as oxidation continued. Possibly, spalling occurred to remove this chromium-containing layer when HC1 was present, or it did not form due to the more aggressive conditions. In addition to the spalling and volatilization effects mentioned above, HC! may encourage attack by altering the conductivity of the protective oxide layer, and the presence of sulphur in the scales indicated that an effect of this kind may have operated. However, the exact nature of the process involved is not completely dear. In the present case, either alumina, chromia, or a chromium-rich spinel was the ratecontrolling layer and these oxide systems are based upon a close packed oxide lattice with the cations located at octahedral and/or tetrahedral interstices. Traces of chlorine may have been in the scale, and, if present as chloride ions, could have had two effects. Firstly, when substituted for O ~- additional cation vacancies would be created in order to maintain a charge balance, and secondly, because of the somewhat larger siz? of the chloride ion compared with the oxide ion, an expansion of the lattice would occur.
The oxidationbehaviourof heat resisting metallicfibres
1041
An increase in the number of cation vacancies would not greatly alter cation diffusion rates because there already exists a large number of unoccupied sites in these oxides. The expanded oxide lattice would, however, allow easier incorporation of sulphur and more rapid oxide growth by encouraging movement between the various sites. In the case of alumina, which grows predominantly by inward diffusion of oxygen, the results showed little or no detrimental effect on the rate of scaling of the alumina-forming alloys when HCI was present. However, as will be discussed in the next section, there were indications that HC1 caused changes in oxide composition that might indicate formation of a less effective barrier to aggressive species, and thus encourage increased rates of attack at longer times.
The behaviour of Fecralloy fibres Auger analysis indicated that the aluminium content of the scale formed on the pre-oxidized Fecralloy samples decreased with increasing test time and with exposure to HCI and that the aluminium was replaced by iron and chromium. These results would suggest that the pre-oxidized scale was slowly being degraded and that this process was accelerated by the presence of HCI. Figure 12 also shows that the sulphur content of the scale increased with time and exposure to HC1 gas, and, as suggested in the previous section, expansion of the alumina lattice by inclusion of sulphur and/or chlorine would encourage cation diffusion and could account for the presence of other cations. Possibly a more important consideration is that since the alloy fibre was very thin the average aluminium level in the alloy would be considerably reduced by the formation of a layer of alumina. For example, assuming the thickness of the pre-oxidized alumina layer was 1 ~m, then, with a fibre diameter of 22 ~m, this would reduce the aluminium content in the alloy from 4.8 to about 3.3 wt%. The untreated Fecralloy did not have such a marked reduction in the aluminium content of the scale when HCI was present since these fibres were presumably not so severely depleted in aluminium. Similarly, no aluminium depletion effects were observed in the lnconel 601 fibres. Diffusion processes in iron-chromium alloys A notable feature of the scales formed on the Type 310 fibres that had been completely oxidized (Fig. 8) was that there was little or no evidence of porosity or voidage, and that the chromium has been largely confined to the original dimensions of the sample while iron oxide formed external to the original fibre dimension. Exactly the same behaviour has been observed on completely oxidized flat samples of iron-chromium alloys,7 and for an Fe-9Cr alloy heated in CO~, where a similar partition of the iron and chromium occurs, Harrison et al. 8 have shown that the reaction rate is independent of the thickness of the iron oxide layer. An interpretation of the growth mechanism leading to this structure is that chromium remains immobile, with outward diffusion of iron and inward diffusion of oxidant. The newly formed iron oxide growing at the iron oxide/chromium-containing oxide interface and the chromium oxide at the alloy/oxide interface. The remarkable feature of the process is that rates of diffusion of iron and oxygen must be coupled and roughly equal in the chromium-containing layer, such that the volume expansion due to new oxide formation at the alloy/oxide interface is balanced
1042
S . R . J . SAtmDEr.S and I. LEF~VEX~
by the volume vacated by the removal of iron ions from the system, otherwise an approximately 100 ~o dense structure would not be maintained. Thus, if the diffusion rate of iron were greater than that of oxygen, voids or porosity would develop, or, if vice versa, large compressive stresses would be produced, resulting in a highly fractured or convoluted scale. It is well known that in the iron--chromium-oxygen system bulk diffusion rates of anions are many orders of magnitude smaller than those of cations and, therefore, in order to achieve an overall balance in the flux of anions and cations, factors other than simple concentration gradients must be involved. One possibility is, of course, that the rate of diffusion of iron is indeed higher than that of oxygen and this would result in the formation of many voids that could link up and provide a short circuit diffusion path for oxygen through the oxide. However, this would be a selfstifling process since the newly formed oxide would grow in the voids and thus eliminate the diffusion path. Alternatively, the diffusion processes may be influenced by a stress-assisted mechanism. If, for example, the chromium-rich oxide contained many short circuit diffusion paths so that oxygen diffusion were faster than iron diffusion and new oxide formation occurred at the alloy/oxide interface, large stresses would result. The chemical potential difference between iron in the unoxidized alloy and the iron in the iron oxide layer would be increased by the increase in pressure caused by the oxide growth stresses. This would then have the effect of increasing the iron diffusion rate, thereby accommodating the growing oxide. Further analysis of this mechanism will be presented in a later publication. SUMMARY
AND CONCLUSIONS
1. Alumina-forming alloys were better able to form protective oxide layers on fibres of very small radius because the predominant oxide growth process involved inward diffusion of oxidant. With chromia-forming alloys, outward diffusion of cations led to "dilution" of the protective alloy constituent at a rapidly expanding surface area, thus making formation of a compact oxide difficult. 2. Alloy depletion effects were observed, particularly in the case of the pre-oxidized Fecralloy fibres and, although failure of the oxide layer did not occur in the life time of the tests, there were indications that protection may breakdown at somewhat longer times, particularly in the tests with HCI present. It is suggested, therefore, that preoxidation at a high temperature may not be beneficial for these thin fibres. 3. HC1 generally caused higher rates of attack but chlorine-containing compounds were not detected in the corrosion product. However, HCI did cause an increase in the sulphur content of the scale and it is possible that incorporation of chlorine and/or sulphur in the oxide modified transport properties by creating additional cation vacancies or by expanding the oxide lattice. Both effects would tend to promote greater rates of attack in the chromia-forming alloys, but deleterious effects would be expected to be reduced somewhat for the alumina-forming alloys. 4. A mechanism of partitioning of alloy constituents during non-protective oxidation of iron-chromium alloys has been proposed. This involves stress-assisted diffusion resulting from formation of oxide at the alloy/oxide interface under compressive stresses. This has the effect of providing a mechanism of coupling the inward diffusion of oxidant and outward diffusion of cations so that the volume expansion required
The oxidation behaviour of heat resisting metallic fibres
1043
for new oxide formation is balanced by the volume vacated by the outward diffusion o f cations. 5. The performance o f Inconel 601 and Fecralloy fibre products was superior to that o f the Hastelloy X and Type 310 stainless steel fibres. However, since there are doubts a b o u t the reliability o f the processing and properties o f the experimental Fecralloy, Inconel 601 is the preferred material for most high temperature applications. Acknowledgements--Financial support from the Belgian Ministry of Science Policy is gratefully acknowledged. Auger spectroscopic measurements were undertaken by G. Haemers and L. Dewulf. T. G. Dye carried out the corrosion experiments and metallography. Helpful discussions on diffusion processes in oxide layers were held with B. Dyson.
REFERENCES D. M. Rovs~R and W. B. Lis^ooR, N.A.S.A. Tech. Note, TN D-6893 (1972). British Steel Corporation, Stainless steel plates for pressure vessels, SSD 793 (1973). V. R. Howv.s, Corros. Sei. 8, 221 (1968). F. A. GoLIon'rLY, F. H. S'ro'rr and G. C. WooD, Oxid. Metals 10, 163 (1976). O. KtmASCI-IL~St
1. 2. 3. 4. 5.
0010-938X/83 $3.00+0.00 Pergamon Press Ltd.
THE OXIDATION BEHAVIOUR OF HEAT RESISTING METALLIC FIBRES S. R. J. SAUNDERS and I. LEFEVER* National Physical Laboratory, Queens Road, Teddington, Middlesex, U.K. Abstract--The behaviour of 22 vm diameter alloy fibres of Type 310 stainless steel, Hastelloy X, Inconel 601 and Fecralloy has been studied at 600°C in combustion gases with and without HCI addition. Chromia-forming alloys exhibited higher rates of attack than would have been expected from a consideration of the properties of the bulk material whereas normal rates of attack were observed for alumina-forming alloys. The addition of HC1 to the combustion gases increased rates of attack for all alloys and generally this was associated with increased sulphur in the oxide layer. The results are discussed with reference to the unusual geometry of the samples, and it is suggested that the direction of motion of the ions in the growing oxide layer is an important factor governing its protective nature. Alloy depletion effects were also noted, particularly for some Fecralloy samples which were pre-oxidized before exposure, and this may cause eventual failure of the protective oxide. A common feature of the oxidation of iron--chromium alloys is the partitioning of elements in the scale; this wag especially marked for the Type 310 stainless steel samples exposed in this work and a stressassisted diffusion process is proposed to explain this effect. INTRODUCTION THIS paper describes the corrosion resistance of metallic fibres, with a diameter of 22 ~m, sintered into a high porosity structure. These fibres, Bekinox®, made by NV Bekaert SA, are available in Type 316L stainless steel, Hastelloy X, Inconel 601, nickel and titanium. The sintered products, the reinforced web and felts and Bekipor®, find applications in dust filtration, liquid filtration, electromagnetic shielding, sound suppression, flame arresting etc. These products can also be used for high temperature gas filtration, e.g. in automotive emission control, in power plants arid in other combustion processes. The requirement to use coal as an alternative fuel to oil has increased the need to develop suitable materials for hot-gas filters. In addition to resisting attack by hot gas, these filters must be capable of removing particles greater than about 10-20 ~tm and able to withstand mechanical shocks during a cleaning cycle. Generally, ceramic materials are not suitable because of inadequate mechanical properties. The filtration characteristics of the porous sintered fibre structures are determined by the fibre diameter and packing density and fibres with diameters between 4 and 22 ~tm can now be produced on a commercial scale. The purpose of the work described was to assess the high temperature oxidation/ sulphidation resistance of these thin fibres which may differ from that of bulk materials because of the large surface area-to-volume ratio. This could result in the rapid depletion of the alloy in additions such as chromium and aluminium that are required for oxidation resistance. Furthermore, the surface state and/or metallurgical structure of these fibres may also have a significant effect on their behaviour. In most cases, *Bekaert Research Centre, NV Bekaert SA, Leo Bekaertstraat 1, 8550 Zwevegem, Belgium. Manuscript received 15 February 1983. ~)1983 Crown copyright. 1025
1026
S.R.J. SAU~0EP.Sand I. LEF~WR
Type 316L stainless steel fibres perform satisfactorily for filtration purposes at temperatures up to about 350°C. The work reported in this paper describes the results of tests on a number of heat-resisting alloy fibres in combustion gas at 600°C with and without the presence of HC1 gas.
EXPERIMENTAL METHOD
Specimen preparation The specimens were made of fibres in four heat-resisting alloys: stainless steel Type 310, Incone1601, Hastelloy X and Fecralloy and the compositions are given in Table 1. The fibres were produced by NV Bekaert SA according to a proprietary process. TASL~ 1.
NOmNAL COMPOSITION OF THE MA1~RIAI.S TESTED (wt ~o)
Type 310 stainless steel
C 0.25max; S 0.03max; Ni 19-22;
Mn 0.20 max; Si 1.5 max; balance Fe
P 0.045max Cr 24-26
Inconel 601
C 0.05; Si 0.25; Cu 0.50;
Mn 0.5; Cr 23; Fe 14-15;
S 0.007 AI 1.35 balance Ni
Hastelloy X
C 0.10; Cr 22; W 0.6;
Mn 0.5; Co 1.5; Fe 18.5;
Si 0.5 Mo 9.0 balance Ni
Fecralloy
C 0.03; AI 4.8; balance Fe
Si Y
Cr 15.8 S 0.004
0.3; 0.3;
Sintered sheets of the four alloys, made from 22 ltm diameter fibres, with a nominal porosity of 80 ~. and a weight of 450-550 g m-L were cut into 25 x 70 mm samples. Some of the Fecralloy samples were given a special preoxidation treatment after sintering. All samples were then degreased in acetone and methanol and weighed prior to testing. Fibre shedding was observed during the first test and this caused difficulties in the interpretation of the gravimetric results, and thus, for subsequent work, loose fibres, particularly from the cut edges, were removed by gentle agitation.
Test procedure The specimens were exposed at 600°C to hot gases obtained by combustion of a 1 ~ S (by weight) fuel oil (BS 2869 Class D) at an air to fuel ratio of 30 : 1 using a burner rig. Complete combustion gave a SOt content of 280 ppm. Addition of 100 ppm HC1 gas to the combustion gas was achieved by atomizing a 0.1 M HCI solution flowing at 0.60 ml rain-a with a fuel flow rate of 0.77 ml min-a. The specimen holder, made from Type 310 stainless steel, which accepted 24 samples was rotated slowly to achieve uniform exposure of the samples to the test environment. Figure 1 shows typical specimens mounted in the specimen holder. Thermal cycles were imposed every 22 h by withdrawing the samples from the burner rig and allowingthem to cool to room temperature. Samples were weighed at approximately I00 h intervals during the test.
Specimen evaluation Some samples were examined by scanning electron microscopy and Auger spectroscopy. Auger analysis was carried out using a cylindrical mirror analyser with a coaxial electron gun at a mean spot diameter of 80 lain. Samples for optical metallography were mounted and polished using conventionalmetallographic techniques. A Cameca Camebax analyser operating at an accelerating voltage of 15 kV and a beam current of 10-g A was used for electron microprobe analysis.
FIG. 1. Partially loaded specimen holder.
9 LQ,
A
"8
FIo. 4. Optical micrographs of fibre samples after 528 h at 600°C in (A) clean combustion gases, and (]3)combustion gases q- 100 ppm HCI. (a) Type 310 stainless steel; (b) Inconel 601 ; (c) Hastelloy X; (d) Fecrailoy, untreated; (e) Fccralloy, treated.
FIG. 5. Scanning electron micrographs of fibre samples after exposure to clean combustion gases at 600°C (a) Type 310 stainless steel, 1056 h; (b) Incone1601, 1056 h; (c) Hastelloy X, 528 h; (d) Fecralloy, untreated, 1056 h; (e) Fecralloy, treated, 1056 h.
10 IJm L.._._.J
FIG. 6.
Microprobe analysis of a Type 310 stainless steel fibre after 1056 h at 600°C in (a) combustion gases; and (b) combustion gases + 100 ppm HCi.
J
oo
3~
-
.
-=
Cb)
10 pm I
FIG. 8.
.I
Microprobe analysis of an Inconel 601 fibre after 1056 h at 600°C in (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
FIG. 9. Microprobe analysis of uncorroded Fecralloy fibres (a) not sinterod; (b) sintered.
lOpm,,
(.)
lOpm,,
(b).
lOpmt,~
lOpm~,~ ,
(c)
(d)
FIG. 10. Microprobe analysis of Fecralloy fibres after 1056 h at 600°C. (a) Untreated, clean combustion gases; (b) treated, clean combustion gases; (c) untreated, combustion gases -k 100 ppm HCI; (d) treated, combustion gases ÷ 100 ppm HCI.
The oxidation behaviour of heat resisting metallic fibres
1035
EXPERIMENTAL RESULTS
Gravimetric data Figure 2 shows the gravimetric data obtained in the test without HCI addition. The curves dearly show the superiority of the Fecralloy and Inconel 601 materials compared with the Hastelloy X and Type 310 stainless steel. The weight loss observed for the Fecralloy specimens was due to fibre shedding during specimen handling. It is probable that all other specimens also lost weight by this process, but that the higher weight gain due to corrosion masked the effect, and, thus, all the gravimetric data slightly underestimated the corrosion rate. Figure 3 shows the influence of the 100 ppm HC1 addition which resulted in a greater than two-fold increase in the weight gain in the case of Type 310 stainless steel, a less marked increase for the Inconel 601 and Fecralloy samples but, somewhat unexpectedly, a lower weight gain for the Hastelloy X specimens.
Optical and scanning electron microscopy Optical microscopy of polished cross-sections of the corroded samples confirmed the overall trends observed from the gravimctric data. Figure 4 shows the behaviour of samples of all the alloys after 528 h in both types of test condition. It was evident that whereas the Hastdloy X and Type 310 stainless steel specimens had been extensively attacked, the Fecralloy and Inconel 601 samples were only slightly attacked, a result consistent with the gravimetdc data. These micrographs were taken from areas exhibiting the most severe attack and illustrate the quite variable behaviour of the individual fibres, particularly for Type 310 stainless steel and, to a lesser extent, for Incon¢l 601. For example, Fig. 4(a) shows a Type 310 stainless steel sample almost completely oxidized alongside a fibre showing virtually no attack. The internal attack observed in some Fecralloy samples (Figs. 12 r 10]/ |
8~
[] Fecralloy, untreated o Fecralloy, pre-oxidized • Inconel 601 •
Type 310 stainless steel
• Hastelloy X
| 4 L-
•
~
•
•
•
,
800
1000
0
-2
200
400
600 Time (h)
Fie. 2. Weight change as a function of time for 22 IJm diameter fibre samplesof Type
310 stainless steel, HastelloyX, Inconel 601 and Fecralloyexposedto combustiongases at 600°C.
S . R . J . SAUNDEP.Sand I. L~'EvER
1036 (3 o • • •
14 12 -
.~
6
Fecralloy, untreated Fecralloy, pre-oxidized Inconel 601 Type 310 stainlesssteel Hastelloy X
•
~
•
~
•
~
•
2 -// /
v==~ ~,
0 ~
~.~-v
~ v
,_-
200
400
........
c
600
800
~ ~-~
1000
Time (h) FIG. 3. Weight change as a function of time for 22 pm diameter fibre samples of Type
310 stainless steel, Hastelloy X, Fecragoy and Incone1601 exlmsed to combustion gases
at 600°C with 100 ppm HCI addition. 4d and e) was also found in the uncorroded fibres and thus was not a consequence of exposure to the hot combustion gases; the origin of this will be discussed later. Scanning electron micrographs of the corroded fibres are shown in Fig. 5. Again it is evident that the Hastelloy X and Type 310 stainless steel samples have been severely attacked compared with those of Inconel 601 and Fecralloy. Of particular interest, however, is the appearance of the Type 310 stainless steel (Fig. 5a) which is consistent with the very heterogenous attack observed in cross-section. The structure of the fibre was such that an array of single crystals is formed ("bamboo" structure), and it is evident that attack was specific to certain grains, and also that grain boundary regions in the fibre were generally protected.
Microprobe analysis Microprobe analysis of some corroded fibres was carried out. Generally only the more severely attacked fibres were examined. Figure 6 compares the behaviour of Type 310 stainless steel after 1056 h in both test conditions. Virtually complete oxidation had occurred in both cases, leaving a compact cylinder of oxide with almost no trace of sulphur- or chlorine-containing compounds. Chromium generally was confined to the initial dimensions of the fibre but, in the case of the sample corroded in the absence of HCI, an outer layer of chromium-containing material was also present. An outer layer of iron-rich oxide formed on both specimens. Figure 7 shows the microprobe analysis of Hastelloy X samples after 528 h in the two test conditions. In both cases, a chromium-rich oxide containing molybdenum and sulphur formed at the metal/oxide interface. However, for the test in the absence of HC1, sulphur was also associated with the outer layer which contained nickel oxide. When HCI was added to the environment iron oxide was dominant in the outer layer.
The oxidation behaviour of heat resisting metallic fibres
1037
Inconel 601 and Fecralloy contain aluminium as an alloying addition and thus AlcOa formation is possible. Inconcl 601, with only 1.35% A1 formed both Cr~O8 and AlsOs (Fig. 8), when HC1 was absent, but only AlcOa could be detected on the sample from the test environment containing HCI. Sulphur was more evident after exposure in the test without HC1, but this could have been due to easier detection in the thicker scale. Microprobe analysis of untreated FecraUoy fibres that were uncorroded and unsintcred indicated the presence of an yttrium-rich intcrmetallic with the composition Fe6.8 Alz.5 Crl.s Y. During sintering in a vacuum furnace at 10 -6 torr and 1100°C, the intermetallic compound reacted with sulphur, as is indicated in Fig. 9. Figure 10 shows the behaviour of the FccraUoy samples after 1056 h in each test condition where it can be seen that some oxidation of the yttrium sulphide had occurred, especially for the untreated samples (Figs. 10a and c), and a thin layer of AI~O~ developed on the fibre surface in association with yttrium-containing material. The presence of HCI encouraged incorporation of sulphur into the scale but no chlorine-containing material was detected. The behaviour of the pre-oxidized samples (Figs. 10b and d) which had a very well-developed layer of AlcOa was generally similar, but there was less evidence of sulphur pick-up. The yttrium sulphide was also largely unoxidized.
Auger analysis Figure 11 shows the oxygen profiles obtained from samples of all alloys after testing in the two environments and this gives an indication of the extent to which the ion etching process proceeded relative to the oxide thickness. For example, on the Hastelloy X and Type 310 stainless steer samples, ion etching for 5 h was insufficient to penetrate the scale whereas that on the untreated Fecralloy sample was penetrated in about 3 h 20 rain. These profiles confirm the lower weight gain of Hastelloy X in the HCl-containing environment. Also, with the Fecralloy and Incon¢l 601 samples the Auger analysis indicates that the presence of HC1 resulted in slightly thinner scales OX~G[PI {o)
-- ~ ........
F¢¢roIIny, qlnf¢~*ot(,¢1 1 0 ~ h ~:ecr,-Iq¥. t f ea ed, Zig/~h
Fecroloy, reared.
1056h
~ intoner qi01 --~-Ho sf@iloy X 5~6h .......... Fecf~floy.untreotcd. Oh
'°I
'"' "°
7% / ~ ' ' ~ "
~
:;:::
30'
,o,
01
I00 Elchtime,mm
2~
3~
0'6
12
1'8
Dl~t Once. ~m
0I
Fc¢rolLny, ~reGled, FlcnlRoy, t r e a t e d ,
~
A151 310 Inccmel 601
2~ ~'~ )056h I(] 5C',
1056)1
\. ~
"6 15
........ =
\\
'\
,00 E . . . . . . . . . .
2;0
3~10
06
12
I$
0~',anc e. ,~m
i
FIG. 11. Oxygen concentration profiles through corroded fibres after exposure at 600°C to (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
1038
S.R.J. SAUNVeP.Sand I. LI~FEVER
whereas the gravimetric data suggested the opposite behaviour, but this was probably related to less fibre shedding of the samples. Little or no chloride was detected in the scales, a result in agreement with the electron microprobe analysis results. The presence of HCI generally caused an increase in the amount of sulphur in the scales, particularly for Inconel 601 and the preoxidized Fecralloy samples, as shown in Fig. 12. This does not agree with the microprobe results which show the opposite behaviour, but because these scales were thin the microprobe results may be in error. Hastelloy X fibres, however, had a deeper penetration of sulphur in the surface scale when HC1 was absent, a result in agreement with the microprobe results. ~0'
SULPHUR (0)
.....
60
..... e
Fecrafloy.untrt~ led, 1056h FectaUoy, treated,
26&h
Fecrolloy, treoqed. AISI 3~0 Intone| 610
lOSSh |0§Eh
1056h
~m~ HQstelloy X 528h ........... Fecralloy. un tr.tt.¢l Oh
S.O I' ~
SULPHUR (b) /
....
~
e.Ot
'~ &.0 c" .o
~.
so
•-,"~ ~
t~
Feceolloy,untr~mted.10Sih Fecralloy. lreQted. 2Jth F.ralloy. tr.ated, lOS.h AI$1310601 Incono[
lOHh IOSgh
Ho|telloy X
SZ6h
o
-----....._.~~_x._.....,x
~
4.0
./
g .~ 3.0 2
0
1~)
",
~
\
\
~
~ zo
\
"%.,,,.
1.0 0 O i
I) ~ EtCh time*min 200 i
06
Distance. pm
1.2
i
300 i 18
0
O* ....
m._..,..~z_
~I~0 ittch Umermin ~ !0
0.6
~stmnce.I~m
I.z
300 i~6
FIG. 12. Sulphurconcentration profiles through corroded fibres after exposure at 600°C to (a) clean combustiongases; (b) combustiongases + 100 ppm HCI. The aluminium content of the oxide formed on the pre-oxidized Fecralloy sample decreased with increasing exposure time, and the addition of HCI to the test environment further decreased the aluminium content of the scale for the untreated as well as for the pre-oxidized sample (Fig. 13). In all cases aluminium in the scale was replaced by iron and chromium. However, HCI did not influence the aluminium content of the scale that formed on Inconel 601. The nickel content of the scale was significantly reduced by HCI gas in the case of the Hastelloy X samples, but there was little effect on Inconel 601 and Type 310 stainless steel. DISCUSSION
Corrosion behaviour of alloy fibres The results of corrosion testing of the four types of alloy indicated that resistance to attack increased in the following order; Type 310 stainless steel, Hastelloy X, Inconel 601 and Fecralloy. Pre-oxidation of the Fecralioy fibres further enhanced their corrosion resistance. Data on the oxidation rate of bulk Incone1601 and Fvcralloy are not readily available at this relatively low temperature of 600°C, but the observed
The oxidation behaviour of heat resisting metallic fibres
1039
ALUMtNiUM (o)
61Ii/"' X 50
"*
t
60
J
t/
.....
Fecrolloy,
untreoted
..... •
Fecrolloy. treated Fecratloy treated
A LUMINIUM (b)
/-
I056h 26/* h 1056h
..... .....
50
* I .
..........
Fecralloy. untreated
0 h
o
Feerolloy, Ir eotee.
264h
Fecrotto y, tr tGtee.
IOS6h
Intoner 601
IOS6h
40.
" ¢"
~
Fecrnlloy. untreated, tOSg h
"" ~o,
\
35'
-
30. 25.
k
i
',,,
~' 2o.
{.
15.
',.,
,
10'
O0 '
I00 Etchtime,mit~ 2 0
i
~
J 300
100 E t c h t me rain ZOO u f i
300
01.6
1.'2
1.0
06
10
DtstcJnce, p m
12
Distance,~m
FIG. 13. Aluminium concentration profiles through corroded fibres after exposure at 600°C to (a) clean combustion gases; (b) combustion gases + 100 ppm HCI.
low rate of attack on the fibre products indicated that the oxidation resistance was not significantly different to that expected from the bulk material. Again, data on the rates of attack of Type 310 stainless steel and Hastelloy X at 600°C are not extensive, but extrapolation of higher temperature data indicated that these alloy fibres were attacked at much higher rates than thicker material. For example, Hastelloy X after I000 h exposure in air at the reduced pressure of 8 torr at 760°C showed a metal loss of about 4 ~m, 1 that is a similar rate of attack to that represented by the present results which were obtained at a much lower temperature. Also Type 310 stainless steel samples were reported to lose only about 2 ~m in thickness when exposed to air for 1000 h at 750°C. ~ Clearly, the presence of small amounts of sulphur dioxide in the combustion gas used in the current work would increase the rate somewhat compared with attack in air, but it is unlikely that this would be the only factor contributing to the very high rates of attack observed on these fibres. Similarly, the thermal cycling could also increase rates of metal wastage, but there was only slight evidence of spalling, even when HCI was present. It is noteworthy that the fibres showing higher than normal rates of attack are from chromia-forming alloys whereas the fibres of alumina-forming alloys, viz. Inconel 601 and Fecralloy, had low rates of attack. Important considerations in understanding the behaviour of these fibres are the very large surface to volume ratio and the rapidly increasing surface area when expansion occurred due to incorporation of oxygen. Chromic oxide grows mainly by outward migration of cations s whereas alumina forms predominantly by inward diffusion of oxidant down the oxide grain boundaries, although it has also been proposed that lattice diffusion of cations is involved in the oxide growth process.4 Thus, for the fibres examined in this work, which have a very small radius of curvature, the formation of a protective scale of
1040
S.R.J. SAtn~DEKSand I. LErtWR
chromia by outward migration of cations would be very difficult because the rapidly increasing surface area would tend to "dilute" the chromium. Furthermore, with the alloy surface rapidly receding, adhesion of the scale to the alloy would be difficult. In the case of alumina growing largely by inward migration of anions however, a rapidly decreasing surface area would aid the formation of a coherent and compact protective layer, although, if a thick scale were to form, a large compressive stress would develop that would ultimately lead to failure. It is believed, therefore, that protective oxide formation on thin fibres is most likely where inward diffusion of anions is the dominant transport process in the growing oxide layer. This criterion would be satisfied when alumina-forming alloys are used.
The effect of HCl<~) additions Auger analysis indicated that addition of 100 ppm HCI~g)to the combustion gases decreased the scale thickness on Hastelloy X, Inconel 601 and possibly the Fecralloy samples, although, for the two latter alloys, the effect was very small and may have been related to errors resulting from simultaneous analysis of several randomlyoriented fibres. Gravimetric and metallographic evidence clearly demonstrated, however, that HCI encouraged attack on the Type 310 stainless steel. In cases where the oxide was thinner in the presence of HCI, this did not necessarily correspond to lower rates of attack because some of the reaction products may have been volatile. For example, with the Hastelloy X samples, the corrosion product formed in the presence of HC1 contained significantly less nickel than that found on the sample corroded in the absence of HCI. Presumably, volatile nickel chloride formed and was removed from the system, resulting in a thinner oxide layer. The rate of metal loss, even in the case of HasteUoy X, was insufficient to establish this mechanism because uncertainties in the initial dimensions of the fibres did not allow accurate assessment of residual metal thickness. Furthermore, chlorides are known to encourage oxide spalling6 which could also result in reduced oxide thickness for the samples exposed to the test environment containing HC1. For example, in the case of Type 310 stainless steel samples shown in Fig. 8, an outer chromium-containing layer was observed on the specimen taken from the test with no HCI present but was absent from the sample exposed to HC1. It is suggest that, because of the "dilution" effects described earlier, the initially formed layer never became protective and remained on the surface as oxidation continued. Possibly, spalling occurred to remove this chromium-containing layer when HC1 was present, or it did not form due to the more aggressive conditions. In addition to the spalling and volatilization effects mentioned above, HC! may encourage attack by altering the conductivity of the protective oxide layer, and the presence of sulphur in the scales indicated that an effect of this kind may have operated. However, the exact nature of the process involved is not completely dear. In the present case, either alumina, chromia, or a chromium-rich spinel was the ratecontrolling layer and these oxide systems are based upon a close packed oxide lattice with the cations located at octahedral and/or tetrahedral interstices. Traces of chlorine may have been in the scale, and, if present as chloride ions, could have had two effects. Firstly, when substituted for O ~- additional cation vacancies would be created in order to maintain a charge balance, and secondly, because of the somewhat larger siz? of the chloride ion compared with the oxide ion, an expansion of the lattice would occur.
The oxidationbehaviourof heat resisting metallicfibres
1041
An increase in the number of cation vacancies would not greatly alter cation diffusion rates because there already exists a large number of unoccupied sites in these oxides. The expanded oxide lattice would, however, allow easier incorporation of sulphur and more rapid oxide growth by encouraging movement between the various sites. In the case of alumina, which grows predominantly by inward diffusion of oxygen, the results showed little or no detrimental effect on the rate of scaling of the alumina-forming alloys when HCI was present. However, as will be discussed in the next section, there were indications that HC1 caused changes in oxide composition that might indicate formation of a less effective barrier to aggressive species, and thus encourage increased rates of attack at longer times.
The behaviour of Fecralloy fibres Auger analysis indicated that the aluminium content of the scale formed on the pre-oxidized Fecralloy samples decreased with increasing test time and with exposure to HCI and that the aluminium was replaced by iron and chromium. These results would suggest that the pre-oxidized scale was slowly being degraded and that this process was accelerated by the presence of HCI. Figure 12 also shows that the sulphur content of the scale increased with time and exposure to HC1 gas, and, as suggested in the previous section, expansion of the alumina lattice by inclusion of sulphur and/or chlorine would encourage cation diffusion and could account for the presence of other cations. Possibly a more important consideration is that since the alloy fibre was very thin the average aluminium level in the alloy would be considerably reduced by the formation of a layer of alumina. For example, assuming the thickness of the pre-oxidized alumina layer was 1 ~m, then, with a fibre diameter of 22 ~m, this would reduce the aluminium content in the alloy from 4.8 to about 3.3 wt%. The untreated Fecralloy did not have such a marked reduction in the aluminium content of the scale when HCI was present since these fibres were presumably not so severely depleted in aluminium. Similarly, no aluminium depletion effects were observed in the lnconel 601 fibres. Diffusion processes in iron-chromium alloys A notable feature of the scales formed on the Type 310 fibres that had been completely oxidized (Fig. 8) was that there was little or no evidence of porosity or voidage, and that the chromium has been largely confined to the original dimensions of the sample while iron oxide formed external to the original fibre dimension. Exactly the same behaviour has been observed on completely oxidized flat samples of iron-chromium alloys,7 and for an Fe-9Cr alloy heated in CO~, where a similar partition of the iron and chromium occurs, Harrison et al. 8 have shown that the reaction rate is independent of the thickness of the iron oxide layer. An interpretation of the growth mechanism leading to this structure is that chromium remains immobile, with outward diffusion of iron and inward diffusion of oxidant. The newly formed iron oxide growing at the iron oxide/chromium-containing oxide interface and the chromium oxide at the alloy/oxide interface. The remarkable feature of the process is that rates of diffusion of iron and oxygen must be coupled and roughly equal in the chromium-containing layer, such that the volume expansion due to new oxide formation at the alloy/oxide interface is balanced
1042
S . R . J . SAtmDEr.S and I. LEF~VEX~
by the volume vacated by the removal of iron ions from the system, otherwise an approximately 100 ~o dense structure would not be maintained. Thus, if the diffusion rate of iron were greater than that of oxygen, voids or porosity would develop, or, if vice versa, large compressive stresses would be produced, resulting in a highly fractured or convoluted scale. It is well known that in the iron--chromium-oxygen system bulk diffusion rates of anions are many orders of magnitude smaller than those of cations and, therefore, in order to achieve an overall balance in the flux of anions and cations, factors other than simple concentration gradients must be involved. One possibility is, of course, that the rate of diffusion of iron is indeed higher than that of oxygen and this would result in the formation of many voids that could link up and provide a short circuit diffusion path for oxygen through the oxide. However, this would be a selfstifling process since the newly formed oxide would grow in the voids and thus eliminate the diffusion path. Alternatively, the diffusion processes may be influenced by a stress-assisted mechanism. If, for example, the chromium-rich oxide contained many short circuit diffusion paths so that oxygen diffusion were faster than iron diffusion and new oxide formation occurred at the alloy/oxide interface, large stresses would result. The chemical potential difference between iron in the unoxidized alloy and the iron in the iron oxide layer would be increased by the increase in pressure caused by the oxide growth stresses. This would then have the effect of increasing the iron diffusion rate, thereby accommodating the growing oxide. Further analysis of this mechanism will be presented in a later publication. SUMMARY
AND CONCLUSIONS
1. Alumina-forming alloys were better able to form protective oxide layers on fibres of very small radius because the predominant oxide growth process involved inward diffusion of oxidant. With chromia-forming alloys, outward diffusion of cations led to "dilution" of the protective alloy constituent at a rapidly expanding surface area, thus making formation of a compact oxide difficult. 2. Alloy depletion effects were observed, particularly in the case of the pre-oxidized Fecralloy fibres and, although failure of the oxide layer did not occur in the life time of the tests, there were indications that protection may breakdown at somewhat longer times, particularly in the tests with HCI present. It is suggested, therefore, that preoxidation at a high temperature may not be beneficial for these thin fibres. 3. HC1 generally caused higher rates of attack but chlorine-containing compounds were not detected in the corrosion product. However, HCI did cause an increase in the sulphur content of the scale and it is possible that incorporation of chlorine and/or sulphur in the oxide modified transport properties by creating additional cation vacancies or by expanding the oxide lattice. Both effects would tend to promote greater rates of attack in the chromia-forming alloys, but deleterious effects would be expected to be reduced somewhat for the alumina-forming alloys. 4. A mechanism of partitioning of alloy constituents during non-protective oxidation of iron-chromium alloys has been proposed. This involves stress-assisted diffusion resulting from formation of oxide at the alloy/oxide interface under compressive stresses. This has the effect of providing a mechanism of coupling the inward diffusion of oxidant and outward diffusion of cations so that the volume expansion required
The oxidation behaviour of heat resisting metallic fibres
1043
for new oxide formation is balanced by the volume vacated by the outward diffusion o f cations. 5. The performance o f Inconel 601 and Fecralloy fibre products was superior to that o f the Hastelloy X and Type 310 stainless steel fibres. However, since there are doubts a b o u t the reliability o f the processing and properties o f the experimental Fecralloy, Inconel 601 is the preferred material for most high temperature applications. Acknowledgements--Financial support from the Belgian Ministry of Science Policy is gratefully acknowledged. Auger spectroscopic measurements were undertaken by G. Haemers and L. Dewulf. T. G. Dye carried out the corrosion experiments and metallography. Helpful discussions on diffusion processes in oxide layers were held with B. Dyson.
REFERENCES D. M. Rovs~R and W. B. Lis^ooR, N.A.S.A. Tech. Note, TN D-6893 (1972). British Steel Corporation, Stainless steel plates for pressure vessels, SSD 793 (1973). V. R. Howv.s, Corros. Sei. 8, 221 (1968). F. A. GoLIon'rLY, F. H. S'ro'rr and G. C. WooD, Oxid. Metals 10, 163 (1976). O. KtmASCI-IL~St
1. 2. 3. 4. 5.