Factors Affecting Chromium Carbide Precipitate ... - Springer Link

Oxidation of Metals, Vol. 50, Nos. 1/2, 1998

Factors Affecting Chromium Carbide Precipitate Dissolution During Alloy Oxidation R. N. Durham,* B. Gleeson,* and D. J. Young* Received July 17, 1997; revised November 11, 1997

Ferrous alloys containing signi® cant volume fractions of chromium carbides were formulated so as to contain an overall chromium level of 15% (by weight) but a nominal metal matrix chromium concentratio n of only 11% . Their oxidation at 850 °C in pure oxygen led to either protective Cr2 O 3 scale formation accompanied by subsurface carbide dissolution or rapid growth of iron-rich oxide scales associated with rapid alloy surface recession, which engulfed the carbides before they could dissolve. Carbide size was important in austenitic alloys: an as-cast Fe± 15Cr± 0.5C alloy contained relatively coarse carbides and failed to form a Cr2 O 3 scale, whereas the same alloy when hot-forged to produce very ® ne carbides oxidized protectively. In ferritic alloys, however, even coarse carbides dissolved suf® ciently rapidly to provide the chromium ¯ ux necessary to form and maintain the growth of a Cr2 O 3 scale, a result attributed to the high diffusivity of the ferrite phase. Small additions of silicon to the as-cast Fe± 15Cr± 0.5C alloy rendered it ferritic and led to protective Cr2 O 3 growth. However, when the silicon-conta ining alloy was made austenitic (by the addition of nickel ), it still formed a protective Cr2 O 3 scale, showing that the principal function of silicon was in modifying the scale± alloy interface. KEY WORDS: iron base alloys; multiphase alloy oxidation; reservoir effect; carbide dissolution.

INTRODUCTION It has been well established that for an iron± chromium alloy to oxidize in a protective manner, a certain minimum chromium concentration (which *School of Materials Science and Engineering, The University of New South Wales, Sydney, NSW 2052, Australia. 139 0030-770X y 98 y 0800-0139$15.00

y

0

Ó

1998 Plenum Publishing Corporation

140

Durham, Gleeson, and Young 1± 3

varies with reaction conditions ) is needed. The scale formed on such single-phase alloys is usually a dense, adherent layer of chromia. The extensive research performed to date has led to a good understand ing of the processes associated with single-phase alloy oxidation. However, since many commercial alloys used for high-temperature service application s are multiphase, a better understand ing of their oxidation behavior is important. The study of multiphase alloy oxidation presents a much greater prob4 lem than for single-phas e alloys. Gesmundo and Gleeson have recently developed the following semiquantitative criterion for the exclusive formation of a protective oxide scale BO, on a binary two± phase alloy AB (Eq. 1): [N B {N B 1 2fn (N B 2 N B )}] $ m

m

p

m

p

m

1/2

2V all V ox

1 2 kp

DB

1 /2

(1)

where N B and N B are the mole fractions of protective-scale-forming element B in the alloy matrix and precipitate , fv the precipitate volume fraction, Va ll and V ox are the molar volumes of alloy and oxide, k p the parabolic rate constant for BO growth, and D B is the diffusion coef® cient of B in a subsurface precipitate depletion zone of thickness x d . This relationship speci® es a number of important factors involved in multiphase alloy oxidation. However, it neglects others that may also be important: external scale nucleation , precipitate size and distributio n, and the dissolution kinetics of the precipitate phase. The purpose of this paper is to investigate the role these factors play in multiphase alloy oxidation. The effect of precipitate size distributio n on oxidation performance has 5 been considered by Wang et al., who developed a non-steady- state diffusional analysis for oxidation of a two-phase alloy. The main ® nding from this work was that for protective behavior of multiphase alloys to be maintained, the ratio of volume fraction of second phase precipitates, f v , to precipitate size, r, is critical. These conclusions have yet to be veri® ed experimentally. The ability of a ferrous alloy to form an external Cr2 O 3 scale also depends on the alloy crystal structure, as is shown by a comparison of the 6 7 6 results of Stott et al. and Kumar and Douglass. Stott et al. examined the oxidation of single-phase ferritic (bcc) Fe± Cr alloys at 1000 °C. They found that an Fe± 14Cr alloy oxidized in a protective manner at 1000 °C, forming an adherent Cr2 O 3 scale with localized areas of iron-rich oxide nodules. By 7 contrast, Kumar and Douglass conducted a series of oxidation experiments on a single-phase austenitic (fcc) alloy of composition Fe± 14Cr± 14Ni at 1000 °C. They found that instead of forming a thin, protective Cr2 O 3 scale, the austenitic alloy developed a thick oxide scale consisting of a thick and

Factors Affecting Chromium Carbide Precipitate Dissolution

141

porous outer layer and an internally oxidized zone. There was also delamination between the outer oxide and the internally oxidized zone. The difference in scaling behavior can be attributed to the signi® cantly lower 8± 10 diffusivity of chromium in austenite compared to that in ferrite. In each of these studies the authors also investigated the effect of 6 adding about 4% silicon to the alloys. Stott et al. found that the siliconcontaining Fe± Cr alloys oxidized to form a protective Cr2 O 3 scale, together with small discrete SiO 2 precipitate s located at the alloy± scale interface. There was a ® vefold reduction in oxide growth kinetics with the addition of 11 silicon. Wood et al. also observed the formation of a thin silica layer when 12 oxidizing ferritic Fe± Cr± Si alloys. Evans and Chatterji and Svedung and 13 Vannerberg examined the effect of silicon additions on the oxidation properties of iron. They both found that the silicon oxidized to form a silica sublayer adjacent to the metal surface. The conclusion drawn from these results was that the silica layer acted as a diffusion barrier, slowing the rate of alloy oxidation, particularly at very high temperatures. 7 Kumar and Douglass studied the oxidation behavior of an austenitic Fe± 14Cr± 14Ni± 4Si alloy. In contrast to the Fe± 14Cr± 14Ni alloy, the siliconbearing alloy oxidized to form a very thin, but adherent oxide scale. Electron probe microanalysis (EPMA) showed that the outer scale was chromia, with a very thin silica (or fayalite) layer adjacent to the metal surface. Simi14 lar results were obtained by Bennett et al. for an Fe± 20Cr± 25Ni niobiummodi® ed stainless steel. 7 In the work by Kumar and Douglass an as-cast austenitic alloy of 7 compositio n Fe± 14Cr± 14Ni failed to form a protective Cr2 O 3 scale. Studies 15,16 17 conducted by Bennett et al. and Allan and Dean were concerned with the effect of cold working of an austenitic alloy on oxidation performance. Allan and Dean found that when a Type 316 stainless steel was cold-worked using a surface-grinding technique and subsequently oxidized at 650 °C, it formed a protective Cr2 O 3 scale. The same alloy without mechanical treatment did not form such a protective scale. The variation in behavior was attributed to recrystallizat ion of the cold-worked steel during heating. The ® ne-grained, recrystallized structure allowed for short-circuit diffusion along grain boundaries, thus providing an increased ¯ ux of chromium to the surface. Bennett and co-workers extended this concept by considering the presence of ® ne NbC particles precipitated during cooling of the same steel. These particles pinned the grain boundaries and retarded grain growth, thus leading to a ® ner grain structure. It was surmized that this ® ner grain structure led to increased oxidation resistance. Alternative ly, 18,19 Giggins and Pettit concluded that a reduction in grain size enhances the nucleation rate of Cr2 O 3 during the early stages of oxidation of Ni± Cr alloys at 900 and 1100 °C in air. According to these authors, the grain boundaries

142

Durham, Gleeson, and Young

act as preferred nucleation sites for the Cr2 O 3 . As a consequence, the ® negrained alloys oxidized slower than coarse-grain ed alloys of similar compositio n. In summary, the ability of a ferrous alloy of intermediate chromium content to form a protective Cr2 O 3 scale has been shown to depend on whether it is austenitic or ferritic, on the alloy grain size, and on the presence of silicon. However, previous work has dealt predominan tly with singlephase materials and there is a need to investigate multiphase alloys in which chromium is enriched in the precipitate phase. A set of model alloys containing chromium carbides was designed for this purpose. Various alloying additions and mechanical treatments were used to alter the metal± matrix crystallograph ic structure and precipitate distributio n in order to examine their effects on oxidation at 850 °C. ALLOY DEVELOPMENT 20

The software package Thermo-Calc was used to predict how alloying additions would affect phase constitutions. In addition, it was used to predict alloy matrix composition and carbide precipitate weight fractions. The initial alloy examined (henceforth called the ``base alloy’ ’ ) was Fe± 15Cr± 0.5C (all compositions in weight percent), which is seen in the calculated phase diagram of Fig. 1a, to have a c 1 M 23 C 6 microstructure at 850 °C. To examine the effect of carbide size on oxidation performance, this alloy was oxidized in both the as-cast and hot-forged condition. All other alloys were oxidized in only the as-cast condition. Investigation of the silicon effect is complicated by the fact that this element is a ferrite stabilizer as well as a protective oxide former. A level of 21 1% silicon was chosen for its known effectiveness in suppressing iron oxide 21 scale growth. Orlova and Ipatyev found that at 800 °C, an addition of 1.43% silicon to a heat-resistant steel totally suppressed the formation of 22 wustite (FeO). Evans and co-workers, on the other hand, analyzed the effect of varying silicon content on oxidation peformance. They found that the total amount of oxidation tended to decrease with increasing silicon content to a minimum at around 0.9% , and thereafter, tended to increase. The exact mechanisms leading to these varying effects of silicon additions are still unclear. Adding 1% silicon to the base alloy alters its phase constitution to a -ferrite plus M 7 C 3 , as shown by the calculated isotherm in Fig. 1b. In order to study the chemical effect of silicon in isolation from its effect of altering the steel’ s crystal structure and diffusivity , it is necessary to maintain the austenitic matrix using an appropriate austenite stabilizer. Nickel was selected for this purpose and an alloy composition of Fe± 15Cr±


Fig. 1. Isothermal sections for alloys investigated. (a) Fe± 15Cr± 0.5C; (b) Fe± 15Cr± 0.5C± 1.0Si; (c) Fe± 15Cr± 0.4C± 1.0Si± 1.0Ni; (d) Fe± 15Cr± 0.5C± 3.0Mo; (e) Fe± 15Cr± 0.5C± 3.0Mo± 3.0Ni.

143

144

Durham, Gleeson, and Young Table I. Nominal Compositions and Microstructural Data M atrix

Alloy composition 1. 2. 3. 4. 5. 6. 7. 8.

Fe± Fe± Fe± Fe± Fe± Fe± Fe± Ni±

15Cr± 15Cr± 15Cr± 15Cr± 15Cr± 15Cr± 15Cr 17Cr

0.5C 0.5C 0.5C± 0.4C± 0.5C± 0.5C±

(as-cast) (hot-forged) 1.0Si 1.0Si± 1.0Ni 3.0M o 3.0M o± 3.0Ni

M Cr (wt.%)

fv %

Phase constitution

11.2

7.5 (5.7)

c 1 M 2 3C 6

11.2 12.0 10.6 10.8 15.0 17.0

6.2 (6.7) 6.1 (5.4) 10.4 (9.0) 9.8 (7.8) Ð Ð

a1 c1 a1 c1

M 7C 3 M 2 3C 6 M 2 3C 6 M 2 3C 6 a c

0.4C± 1.0Si± 1.0Ni was selected on the basis of Thermo-Calc analysis. As shown in the calculated isotherm (Fig. 1c), the desired c 1 M 23 C 6 phase constitutio n is achieved by lowering the carbon content to 0.4% . To complete the alloy set, it was necessary to design an a 1 carbide steel in order to investigate the effect of changing the matrix to ferrite, without simultaneou sly introducin g the strong oxide former like silicon. The alloying element selected for this purpose was molybdenu m and a composition of Fe± 15Cr± 0.5C± 3.0Mo was chosen on the basis of the Thermo-Calc prediction shown in Fig. 1d. As seen, a phase constitutio n of a 1 M 23 C 6 is the result. Finally, to verify that molybdenu m had no major effect other than producing a ferrite matrix, a molybdenu m-bearing austenitic alloy was designed. The composition Fe± 15Cr± 0.5C± 3.0Mo± 3.0Ni was chosen to yield a c 1 M 23 C6 constitution on the basis of the Thermo-Calc prediction shown in Fig. 1e. All alloy compositions were adjusted so as to yield closely similar matrix concentrations. This information is summarized in Table I.

EXPERIMENTAL PROCEDURES Twenty-gram buttons of the alloys were prepared from the constituent materials (iron 99.98% , chromium 99.5% , graphite 99.95% , silicon 99.95% , molybdenum 99.5% , and nickel 99.95%) by nonconsumable electrode argon arc melting. The alloy compositions together with the predicted matrix chroM atrix mium composition, M Cr , and theoretical carbide volume fraction, fv , are given in Table I. The measured volume fractions, as determined by a pointcount method, are shown in brackets.


145

Table II. Actual Experimental Alloy Compositions (in wt.%) Nominal alloy composition 1± 2. Fe± 15Cr± 0.5C 3. Fe± 15Cr± 0.5C± 1.0Si 4. Fe± 15Cr± 0.4C± 1.0Si± 1.0Ni 5. Fe± 15Cr± 0.5C± 3.0M o 6. Fe± 15Cr± 0.5C± 3.0M o± 3.0Ni 7. Fe± 15Cr 8. Ni± 17Cr

%Cr

%C

%Si

15.0 15.0 15.0 15.2 15.1 15.3 17.1

0.47 0.51 0.31 0.46 0.47 Ð Ð

Ð 1.17 1.04 Ð Ð Ð Ð

%Mo Ð

%Ni Ð

Ð Ð 2.86 3.04 Ð Ð

Ð 1.12 Ð 2.75 Ð 82.9

Nominal alloy compositions were veri® ed using standard wet chemical analysis techniques. Table II summarizes the results of these analyses; there is reasonable agreement between the nominal and measured compositions. The alloys were subjected to hot-stage X-ray diffraction (XRD) at 850 °C using Cu± Ka radiation; the results are shown in Fig. 2. The range of 2h values shown encompasses the primary peaks for ferrite (44.66 °) and 23 austenite (43.53 °). The solid line is the room temperature trace and the broken line is the trace at 850 °C. Alloys 1, 2, 3, and 5 had the phase constitutions predicted. Alloys 4 and 6 consisted principally of the predicted austenite, but contained minority amounts of ferrite as well. The Fe± 15Cr± 0.5C alloy was hot-forged at 1150 °C to increase the carbide dispersion. The microstructure consisted exclusively of austenite at the hot-forging temperature of 1150 °C. This allowed for the homogeniza tion of elements that may have been segregated during the melting process. The alloy button was forged along a radial axis to ensure the maximum amount of deformation. All alloys were annealed for 24 hr at 1000 °C under ¯ owing argon, where the carbides were reprecipitat ed in the austenite matrix. At the oxidation temperature of 850 °C, the forged and as-cast alloys consisted of a metallic (ferrite or austenite) matrix plus carbides (M 23 C 6 or M 7 C 3 ), details of which are summarized in Table I. Oxidation samples approximate ly 10 3 10 3 1 mm were cut from the annealed alloys and abraded to a 1200-grit surface ® nish. The samples were then cleaned and degreased in acetone. Oxidation tests were carried out at 850 °C in pure oxygen (¯ ow rate 100 ml y min) in batch runs for a period of 2 weeks. To obtain oxidation kinetics, samples were oxidized in a conventional thermogravim etric analysis (TGA) unit with the same operating conditions as for the batch runs. The oxidation products were analyzed in situ using XRD to identify the surface scale product and then cold mounted in epoxy resin for crosssectional analysis using optical and scanning electron microscopy (SEM ). For optical identi® cation of the carbide precipitates, a solution of GroÈ sbeck’ s Reagent (4 gm KMnO 4 , 4 gm NaOH, and 120 mL water) was used

146


Fig. 2. X-ray diffraction traces for alloys investigated.


147

Fig. 2. Continued.

to stain the carbides. The alloy regions underneath the oxide scales were analysed using electron probe microanalysi s (EPMA). RESULTS Optical photomicrographs of the oxidized samples appear in Fig. 3. Two as-cast alloys, Fe± 15Cr± 0.5C and Fe± 15Cr± 0.5C± 3.0Mo± 3.0Ni, grew nonprotective scales while all others developed thin Cr2 O 3 scales. The microstructure of the as-cast Fe± 15Cr± 0.5C alloy was austenitic with interdendritic carbide particles of average size 3 mm. The oxide formed on this base alloy was a multilayere d scale with an outer layer composed essentially of hematite (Fe2 O 3 ), and a thick multiphase inner layer composed of FeO and FeCr2 O 4 (Fig. 3a). The second-phase carbides in the alloy were engulfed by the advancing oxidation front, precluding the formation of a carbide depletion zone. The Fe± 15Cr± 0.5C± 3.0M o± 3.0Ni alloy had a microstructure consisting of a ® ne array of interdendritic carbides of about 2 mm size, together with some secondary carbides of about 1 mm size in an austenitic matrix, similar to the base alloy. The alloy oxidized in a manner similar to the base alloy in that the scale was a thick, multiphase oxide. After 3 days of oxidation, the scale consisted of an outer layer of iron oxides (primarily Fe2 O 3 ) and an inner layer of mixed FeO and FeCr2 O 4 . Adjacent to the metal± scale interface was an oxide region enriched in molybdenum (Fig. 3b). Once again, the second-ph ase carbides in the alloy were enveloped by the fastgrowing scale. The hot-forged Fe± 15Cr± 0.5C alloy consisted of an austenite matrix with spheroidal carbides dispersed throughout the matrix. The average size of the carbides was about 1 mm, with some larger ones (in the order of 4 mm)

148


Fig. 3. Optical images of oxidized alloys. (a) Fe± 15Cr± 0.5C (as-cast); (b) Fe± 15Cr± 0.5C± 3.0Mo± 3.0Ni; (c) Fe± 15Cr± 0.5C (hot-forged); (d) Fe± 15Cr± 0.5C± 1.0Si; (e) Fe± 15Cr± 0.4C± 1.0Si± 1.0Ni; (f ) Fe± 15Cr± 0.5C± 3.0Mo; (g) Fe± 15Cr; (h) Ni± 17Cr.


Fig. 3. Continued.

149

150


Fig. 3. Continued.


Fig. 3. Continued.

151

152


Fig. 4. Particle size distribution for hot-forged alloy.

also present. Figure 4 shows a precipitate size distributio n; the majority of the precipitates lie in the 0.5± 1.0 mm range. However, about 3% of the precipitates have a size greater than 3 mm. The alloy oxidized to form a thin, protective Cr2 O 3 scale together with a subsurface depletion zone in which the carbides had dissolved (Fig. 3c). The interface between the bulk alloy and the depleted zone was found to be nonplanar, with islands of residual carbide precipitates left behind in the depleted zone. The cast Fe± 15Cr± 0.5C± 1.0Si alloy had a microstructure comprising large (around 5 mm) carbides, making up an almost continuou s interdendritic network, plus somewhat smaller secondary carbides within the ferrite dendrites. The alloy oxidized protectively, forming a thin Cr2 O 3 scale and a carbide-depleted zone. The interface between the bulk alloy and the depleted zone was approximate ly planar (Fig. 3d). The addition of nickel to the silicon-con taining alloy reestablished an austenitic matrix (Fe± 15Cr± 0.4Cr± 1.0Si± 1.0Ni), which contained very ® ne (less than 1 mm) spheroidal carbides plus some larger interdendritic carbides. The alloy oxidized to form a thin, protective Cr2 O 3 scale and associated with this protective behavior was the formation of a carbide depletion zone beneath the scale (Fig. 3e). The interface between the bulk alloy and the


153

Fig. 5. EPMA traces for binary, as-cast alloys.

depleted zone was nonplanar, as in the case of the austenitic hot-forged Fe± 15Cr± 0.5C alloy. The ferritic Fe± 15Cr± 0.5C± 3.0Mo alloy (Fig. 3f ) also oxidized in a protective manner. The ferrite contained carbides, which varied in size up to about 3 mm, together with coarse interdendritic carbides. This alloy formed a protective scale consisting of an inner Cr2 O 3 layer and an outer FeCr2 O 4 layer together with a precipitate-depletion zone in the subscale region. The interface between the bulk alloy and the depleted zone was also planar. As a comparison to the multiphase alloys, two as-cast single-phase binary alloys were also oxidized. The ferritic single-phase alloy (Fe± 15Cr) is seen in Fig. 3g to have oxidized to form an external Cr2 O 3 scale, as was found in the hot-forged base alloy. An EPMA trace of the underlying alloy region is seen in Fig. 5. The depth of the depletion zone in the alloy was found to be about 20 mm and to be associated with a step discontinui ty in chromium concentration. The austenitic single-phas e Ni± 17Cr alloy oxidized to form a thin multilayer scale, as shown in Fig. 3h. The outer layer was found to be NiO and the inner layer consisted of NiCr 2 O 4 . The corresponding EPMA trace is also shown in Fig. 5, where the depletion zone is seen to be about 10 mm deep.

154


Fig. 6. Oxidation kinetics for alloy systems.

Weight uptake kinetics are shown in Figs. 6 and 7. They followed the parabolic rate law (Eq. 2):

1 2 DW A

2

5 kp t

(2)

where DW y A is the weight change per unit surface area, t is the time, and k p the parabolic rate constant. Values of k p found from regression of the data in Eq. (2) are listed in Table III along with depletion zone depths after 2 weeks of oxidation. The values of alloy chromium concentrations (in weight percent) at the metal± scale interface as determined by EPMA are also listed. The differences in oxidation morphology of the alloys are re¯ ected in the values of the oxidation rate constant, k p . The value for the base alloy - 8 2 - 4 - 1 - 12 2 - 4 - 1 was 2.5 3 10 g cm s , compared with a vaue of 6.8 3 10 g cm s for the hot-forged alloy. The two molybdenu m-containing alloys also had quite - 11 2 - 4 - 1 different oxidation rate constants: a value of 1.4 3 10 g cm s for the - 9 - 4 - 1 ferritic alloy but 9.3 3 10 cm s for the austenitic alloy. The latter value is similar in magnitude to that found for the as-cast base alloy. The two


155

Table III. Oxidation Kinetics and Depletion Zone Depths

Alloy composition 1. 2. 3. 4. 5. 6. 7. 8. a

Fe± Fe± Fe± Fe± Fe± Fe± Fe± Ni±

15Cr± 15Cr± 15Cr± 15Cr± 15Cr± 15Cr± 15Cr 17Cr

0.5C a 0.5C 0.5C± 1.0Si 0.4C± 1.0Si± 1.0Ni 0.5C± 3.0M o 3.0M o± 3.0Ni

Protective behavior

xd (mm)

kp 2 - 4 - 1 (g cm s )

No Yes Yes Yes Yes No Yes Yes

Ð 35.2 22.0 24.7 45.1 Ð 22.0 12.0

2.5 3 6.8 3 1.3 3 1.4 3 1.4 3 9.3 3 3.0 3 1.4 3

- 8

10 10 10 10 10 10 10 10

Indicates hot-forged alloy; all others are as-cast.

Fig. 7. Oxidation kinetics for alloy systems.

12 12 12 11 9 10 10

Alloy± scale interface Cr concentration (%) Ð 5.9 9.6 10.1 6.8 Ð 4.8 4.6

156


Fig. 8. Diffusion path for hot-forged base alloy. - 12

2

- 4 - 1

silicon-con taining alloys had very similar k p values of 1.3 3 10 g cm s - 12 2 - 4 - 1 for the ferritic alloy and 1.4 3 10 g cm s for the austenitic alloy. The rate constant for the ferritic molybdenu m-contain ing alloy was 1.4 3 - 11 2 - 4 - 1 10 g cm s , about one order of magnitude faster than the corresponding silicon-con taining ferritic alloy. This result re¯ ects the fact that the molybdenum-containing alloy formed a spinel in conjunctio n with a chromia scale, as described below. The major difference between the protectively oxidized alloys was the size of the depletion zone. The hot-forged alloy developed a depletion zone depth of 35 mm in a 2-week oxidation period. The two silicon-con taining alloys once again showed comparable behavior, with depletion zone depths of 22 mm and 25 mm in the ferritic and austenitic alloys, respectively. The ferritic molybdenum -containin g alloy, which oxidized in a protective manner, formed a precipitate depletion zone 45 mm deep. The alloys forming protective scales were subjected to EPMA. Results for the hot-forge d base alloy are presented in the form of a diffusion path on the Fe± Cr± C isothermal section in Fig. 8. It is seen that the diffusion


157

path crosses tie lines in the two-phase c 1 M 2 3 C6 region. It ® nally exits into the single-phase c region, in accordance with the results shown in the photomicrograph in Fig. 3c. Diffusion path behavior for the other alloys could not be described, as carbon concentration measurements were irreproducib le. EPMA was also conducted on the scales of the hot-forged base alloy and the ferritic molybdenu m-containing alloy. These results are presented in Fig. 9, where it is seen that the former developed a scale of Cr2 O 3 and the latter grew a duplex scale of Cr2 O 3 overlayed by spinel. DISCUSSION The alloys examined fell into two distinct classes: those that oxidized slowly to grow a thin, protective Cr2 O 3 scale, and those that reacted fast, forming iron-rich oxide scales. The former group of alloys developed subsurface zones in which their chromium carbide precipitates had dissolved. The others experienced rapid surface recession, which engulfed the carbides before they could dissolve. The principal factors differentiating the two alloy classes were precipitate distributio n and alloy diffusivitie s. Precipitate distribution clearly has a dramatic in¯ uence on the oxidation performance of a multiphase alloy. The hot-forged base alloy oxidized to form a protective Cr2 O 3 scale (Fig. 3c), whereas the as-cast alloy of the same composition oxidized to form a thick, layered scale (Fig. 3a). This result con® rms the 5 suggestion by Wang et al. that a critical precipitate size and volume fraction are required for protective oxidation. The as-cast and hot-forged alloys had the same precipitate volume fraction, but the reduction in precipitate size produced by hot-forging led to preferential chromium oxidation and increased alloy oxidation resistance. To maintain protective behavior, the ¯ ux of chromium to the alloy surface must be great enough to balance the amount of chromium being consumed in the formation of the protective Cr2 O 3 scale. In the as-cast alloy, much of the chromium is bound up in large interdendritic carbides. The low surface area to volume ratio of these precipitates means that their dissolution provides a relatively low chromium ¯ ux, insuf® cient to sustain a continuou s, protective Cr2 O 3 scale. The much ® ner carbides produced by hot-forging have a higher total surface area and their dissolution provides a chromium ¯ ux suf® cient to grow a Cr2 O 3 scale. The interface between the bulk alloy and the depleted zone formed in the hot-forged base alloy was nonplanar, as shown in Fig. 3c, and it is suggested that this behavior was due to the distributio n of precipitate sizes. The precipitate distributio n in the hot-forged alloy, shown in Fig. 4, reveals that while most precipitates were of the order of 1 mm, a signi® cant number

158


Fig. 9. EPMA traces for oxide regions on alloys that formed protective scales.


159

were of size 3 mm or more. The precipitates functioned as chromium reservoirs that dissolved, supplying chromium to the surface. The larger ones took longer to dissolve and their residues were left behind in the denuded zone while the depletion zone front advanced further into the alloy, thus creating a nonplanar interface. This is consistent with the observation that carbides of 3 mm size in the as-cast base alloy did not dissolve fast enough to supply an adequate chromium ¯ ux. Implicit in the preceding analysis is the fact that the rate of carbide dissolution is not much greater than the diffusion process and that precipitate size is critically important in the micron range. The diffusion path for the hot-forged base alloy was determined using EPMA. The diffusion path traversed the c 1 M 23 C 6 phase ® eld, exiting into the single-phase c ® eld. Evidently the alloy subsurface region was depleted in both chromium and carbon, that is to say, the alloy decarburized while oxidizing. The alloy± scale interface operated as a carbon sink, allowing carbon to escape through an apparently protective Cr2 O 3 scale. Carbon solu24 bility in Cr2 O 3 scales is negligible, as determined by Wolf and Grabke, therefore solid-state diffusion of carbon through the scale is not a possible mechanism. Gas diffusion of carbon oxides is also impossible, if the Cr2 O 3 24,25 26 scale is truly protective. Wolf et al. and Zheng et al. have examined carbon transport mechanisms in oxide scales, however, the carbon was diffusing inward, from a carbon-con taining gas. Both groups reported that carbon entered the metal by being transported across the scale via submicroscopic defects, presumably as carbon monoxide molecules. It is suggested that the reverse of this process is operative in the present case. Decarburization results from carbon diffusing to the alloy surface, undergoing oxidation to form carbon monoxide, and then diffusing through the scale, along the above mentioned defects. Diffusion paths for the other alloys could not be obtained due to poor results for carbon analysis. However, analysis of the other elements revealed that concentration gradients in molybdenu m, nickel, and silicon were negligible. This indicates that a silica subscale, if present, was extremely thin, involving very little consumption of silicon. Field-emission scanning electron microscopy (FESEM ) and EPMA were used in an attempt to locate such a scale. No indication s of a subscale were found, con® rming that it 27,28 must be very thin. Bennett et al. characterized such thin silicon-con taining subscales using transmission electron microscopy (TEM) and found them to be in the order 0.03± 0.08 mm thick and discontinu ous. The effect of silicon additions on oxidation was very pronounced. The addition of 1% silicon to the base alloy produced an alloy with a ferrite matrix plus M 7 C 3 carbides, which oxidized to produce a protective Cr2 O 3 scale. Even with an austenitic matrix (induced by virtue of a 1% nickel

160


addition), the silicon-bear ing alloy formed a protective Cr2 O 3 scale. It is, therefore, concluded that the effect of silicon on oxidation is related not to the change it produces in alloy diffusivity, but rather its ability to slow scale growth by forming a silica sublayer that acts as a partial barrier to chromium transport. This is re¯ ected in the higher interface chromium concen29 trations seen in Table III. Work by Johnston further supports this conclusion. This author examined the effect of increasing silicon content on the diffusion of chromium in nickel± chromium± silicon alloys. It was found that doubling the silicon content of the alloy produces a 20% increase in the diffusion coef® cient, a value not signi® cant enough to produce protective oxidation behavior. In the absence of an additional scale-forming element like silicon, the alloy diffusion coef® cient is nonetheless important. Given the difference a 2 - 1 - 11 between values for ferrite and austenite [D Cr ~ 2.8 3 10 cm s (refs. 8, 9) c 2 - 1 - 13 compared with DC r ~6.8 3 10 cm s [refs. 9, 10], at 850 °C], it is expected that a ferritic alloy would oxidize protectively at a lower chromium level than would an austenitic alloy. The behavior of the cast alloys 1 and 5 qualitativel y con® rms this expectatio n. Alloy 1 was austenitic and oxidized nonprotectively, whereas the ferritic alloy 5 formed a protective Cr2 O 3 scale. The chosen ferrite stabilizer in alloy 5, molybdenu m, does not form a protective scale when oxidized, and at the 3 wt.% concentration added, is not expected to alter the scaling rate. This fact is con® rmed by the behavior of the austenitic molybdenu m-containing alloy 6, which oxidized in a nonprotective manner, comparable to the as-cast base alloy. Because molybdenum is of no protective value in an austenitic chromium steel and because it forms no protective scale layer in a ferritic chromium steel, it is concluded that the sole value (from an oxidation point of view) of molybdenum is as a ferrite stabilizer. As such, it promotes chromium diffusion and protective Cr2 O 3 scale formation. To examine the effect of the various parameters associated with the supply and consumption of chromium within multiphase alloys, a mass bal30 ance expression was employed. The ¯ ux of the protective scale-forming element transported through the single-phase precipitate depletion zone equals the ¯ ux of that element transported through the oxide away from the alloy± scale interface. Using the assumption of local equilibrium at the carbide dissolution front, one ® nds for the chromium mass balance (Eq. 3): N Cr 2 N Cr 5 o

o

i

u 2

Vall c exp( c )

1 Ö p u erf (c ) i

(3)

where N C r is the bulk alloy concentration, N Cr is the concentration of chromium in the alloy at the alloy± scale interface, V all is the molar volume of


161

Table IV. Calculated D C r Values for Cr2O 3-Forming Alloys

Alloy system Fe± Fe± Fe± Fe±

15Cr± 15Cr± 15Cr± 15Cr±

Alloy matrix type

0.5C (hot-forged) 0.5Cr± 1.0Si (as-cast) 0.4C± 1.0Si± 1.0Ni (as-cast) 0.5C± 3.0Mo (as-cast)

Austenite Ferrite Austenite Ferrite

Calculated diffusion coef® cient, 2 D Cr (cm s) 5.8 3 4.4 3 1.3 3 3.8 3

- 12

10 10 10 10

11 11 11

the alloy, erf is the error function, and u and c are terms related to the 2 alloy surface recession rate constant, k c (in cm y s) and depletion zone depth x d , respectively (Eqs. 4 and 5): u5

Î

kc 4D Cr

(4)

and,

c5 o

i

xd

Ö 4D Cr t

(5)

Values for N Cr and N Cr were determined from EPMA. Values of k c were calculated from thermogravimetric kinetics on the basis that the scales were pure Cr2 O 3 . The k c value for the ferritic molybdenum -containin g alloy was weighted to take into account the fact that there was a mixed scale of Cr2 O 3 and FeCr2 O 4 , as shown in Fig. 9b. Values for the chromium diffusion coef® cient in the depleted zone, D C r , were then calculated from Eq. (3); the results appear in Table IV. Diffusion coef® cients for chromium in the ferritic alloys are in the order of those reported in the literature. The austenitic alloys, however, were found to have unexpectedly high diffusion coef® cients. The value for the hot-forged base alloy was around one order of magnitude higher than the reported value, while the value for the silicon-con taining austenitic alloy was around thirty times faster than reported. The high diffusion coef® cients deduced for austenitic alloys are dif® cult 29 to understand. It is known that the presence of silicon in austenite leads to an increase in D Cr , but only by about 20% . While it is clear that the two ferrite stabilizers, chromium and silicon, would interact thermodyn amically with each other in the austenite solution, the effect of this interaction on the apparent diffusion coef® cient cannot be predicted without knowledge of 31 the kinetic interaction between the two solutes. As the solutes are both substitutional, their diffusion is inevitably correlated through their shared use of lattice vacancies and the kinetic interaction effect cannot be ignored.

162


The remaining possible explanation for the unusually high D Cr values is microstructural. At the temperature of oxidation, grain-bound ary diffusion may be playing a dominant role in the diffusion of chromium to the alloy surface. As the alloy is hot-forged, the carbide morphology is changed, while at the same time, the alloy grain size is greatly reduced when compared to the as-cast alloy. The alloy matrix, of course, remained austenitic and short-circuit diffusion through grain boundaries could dominate. Work 10 by Smith and Gibbs examined the grain-bound ary and lattice-diffusion characteristics of an austenitic stainless steel. They derived the following relationship s for grain-bound ary and lattice diffusion, respectively (Eqs. 6 and 7): D Cr 5 2.5 3 10 grain

- 6

exp(- 44700 y RT )

(6)

and, D Cr 5 1.9 3 10 latt

- 1

exp(- 58800 y RT )

(7)

At 850 °C, the calculated values for the two diffusion coef® cients are D 5 2 - 1 latt 2 - 1 - 15 - 13 5.0 3 10 cm s and D Cr 5 6.8 3 10 cm s ; lattice diffusion is seen to be predominan t over grain-bound ary diffusion. The reasons for the unusually rapid diffusion of chromium in this hot-forged ternary alloy remain unknown. The oxidation behavior of the single- and multiphase alloys was quite different. First, the chromium pro® les developed in the two single-phas e alloys were quite different. The austenitic single-phase Ni± 17Cr alloy shows a pro® le that is indicative of diffusion control. In contrast, the single-phase ferritic Fe± 15Cr alloy, shows an abrupt step in chromium concentration. This can be understood in terms of the phase transformation occurring as chromium is preferentially removed by oxidation. This is made evident by the alloy diffusion path shown in Fig. 10. The path starts in the ferrite ® eld and, as the chromium is removed, it moves toward and through the austenite loop, then back into the chromium-depleted ferrite ® eld. A steady-state diffusion pro® le representing this path would contain step discontinu ities at each of the two-phase boundaries. However, they are small (about 1 wt.% each) and are dif® cult to detect. More signi® cant is the appearance of an austenite zone spanning a chromium concentration range of approximate ly 5 to 8.5 wt.% . Because diffusion through this zone occurs at the same rate as in the surrounding ferrite, but is accomplish ed with a much lower diffusivity, the concentration pro® le within the zone is necessarily steeper. The shape of the pro® le in Fig. 5 is thereby explained. This 32,33 pro® le will control the rate of chromium delivery to the alloy surface. Similar effects are available in the carbide-co ntaining alloys. Reference to Fig. 1c for the austenitic Fe± 15Cr± 0.4C± 1.0Si± 1.0Ni alloy shows that if grain Cr


163

Fig. 10. Diffusion path for as-cast, single-phase Fe± 15Cr alloy.

selective chromium removal is, as expected, accompanied by decarburization, the formation of a subsurface ferrite zone is highly likely. In this event, chromium diffusion will be faster and the high average value of D C r calculated for this alloy by mass balance would thereby be accounted for. The remaining dif® culty is the base Fe± 15Cr± 0.5C alloy which underwent both chromium depletion and decarburization, evidenced an anomalously high value for D Cr , but according to EPMA measurements did not develop a ferrite zone. These observation s are inconsisten t. Given the substantial measurement error in the microprobe analysis of carbon and the availability of a ferrite phase in the reacting system (Fig. 1a), it is therefore suggested that a subsurface ferrite region may, in fact, have formed. CONCLUSIONS The oxidation behavior of a set of model Fe± Cr alloys containing chromium carbides has been used to de® ne the factors that determine whether or not chromium is selectively oxidized. The alloys were designed using Thermo-Calc and the prediction s of this model as to alloy phase constitution have been veri® ed.

164


Austenitic Fe± Cr± C alloys were oxidized in the as-cast and hot-forged condition. Hot-forging reduced the average carbide size from about 3 mm to about 1 mm and led to a change from nonprotec tive iron-rich oxide formation to protective Cr2 O 3 scaling accompanied by subsurface carbide dissolution. The change is due to the alteration in carbide morphology and it is concluded that precipitate dissolution is too slow, in the case of coarse carbides, to supply enough chromium for selective Cr2 O 3 formation. Addition of silicon to the otherwise unprotective Fe± Cr± O alloy led to selective Cr2 O 3 formation. As no re® nement in carbide size resulted, it is concluded that the silicon functions by converting the alloy matrix to ferrite (with more rapid chromium diffusion) and y or by forming a thin sublayer of SiO 2 at the scale base. Since an austenitic alloy (with slow chromium diffusion) containing the same level of silicon also formed a protective Cr2 O 3 scale, it is concluded that SiO 2 formation is the critical factor. In the absence of silicon, the alloy phase and its diffusivity are nonetheless important. A ferritic alloy, Fe± Cr± Mo± Cr, oxidized protectively. Since this alloy had rather coarse carbides and since the molybdenu m did not form a protective oxide itself, it is concluded that the stabilization of ferrite and the resulting high chromium diffusivity was the mechanism of the molybdenum effect. The lack of any direct effect on scale properties was con® rmed by the observation that an austenitic Fe± Cr± Mo± Ni± C alloy failed to produce a protective Cr2 O 3 scale. Values of D Cr were calculated from a chromium mass balance and were in qualitative agreement with the above conclusions . However, the values calculated for austenitic alloys were surprisingly high. It is thought that higher effective D Cr values resulted from the formation of a subsurface ferrite zone due to the simultaneo us chromium depletion and decarburization of the alloy surface region.

REFERENCES 1. G. C. Wood, I. G. Wright, T. Hodgkiess, and D. P. Whittle, Werkst. Korros. 21, 900 (1970). 2. I. G. Wright, Oxidation of Iron-, Nickel-, and Cobalt-base Alloys, MCIC Report, M CIC 72-07, (Metals and Ceramics Information Centre, 1972). 3. G. Wallwork, Rept. Progr. Phys. 39, 401 (1976). 4. F. Gesmundo and B. Gleeson, Oxid. Met. 44, 211 (1995). 5. G. Wang, B. Gleeson, and D. L. Douglass, Oxid. M et. 35, 333 (1991). 6. F. H. Stott, G. J. Gabriel, F. I. Wei, and G. C. Wood, Werkst. Korros. 38, 521 (1987). 7. A. Kumar and D. L. Douglass, Oxid. Met. 10, 1 (1976). 8. A. W. Bowen and G. M. Leak, Met. Trans. 1, 2767 (1970). 9. C. J. Smithells, Metals Reference Book (Butterworths, London, 1967), Vol. II, p. 664. 10. A. F. Smith and G. B. Gibbs, Met. Sci. J. 3, 93 (1969). 11. G. C. Wood, J. A. Richardson, M. G. Hobby, and J. Boustead, Corros. Sci. 9, 659 (1969). 12. J. W. Evans and S. K. Chatterji, J. Electrochem . Soc. 106, 860 (1959).


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13. I. Svedung and N. G. Vannerberg, J. Less-Common Met. 14, 391 (1975). 14. M. J. Bennett, J. A. Desport, and P. A. Labun, Oxid. Met. 22, 291 (1984). 15. M. J. Bennett, G. W. Horsley, and M . R. Houlton, Proc. Fundamental Aspects of Corrosion Protection by Surface Modi® cation (The Electrochemi cal Society, Pennington, NJ, 1984), p. 282. 16. M. J. Bennett, J. A. Desport, and P. A. Labun, Oxid. Met. 22, 291 (1984). 17. S. J. Allan and M. J. Dean, Proc. Behavior of High Temperature Alloys in Aggressive Environments (The Metals Society, Warrendale, PA, 1980), p. 319. 18. C. S. Giggins and F. S. Pettit, Trans. Met. Soc. AIME 245, 2495 (1969). 19. C. S. Giggins and F. S. Pettit, Trans. Met. Soc. AIME 245, 2509 (1969). 20. B. Sundman, B. Jansson, and J. O. Andersson, CALPHAD 9, 153 (1985). 21. G. M. Orlova and V. V. Ipatyev, J. App. Chem. USSR, 29, 889 (1956). 22. H. E. Evans, D. A. Hilton, R. A. Holm, and S. J. Webster, Oxid. Met. 19, 1 (1983). 23. M. King, (ed.), Powder Diffraction Files, Inorganic Phases (Sets 1± 43) (International Centre for Diffraction Data, Newtown Square, Pennsylvania, 1993). 24. I. Wolf and H. J. Grabke, Solid State Commun. 54, 5 (1985). 25. I. Wolf, H. J. Grabke, and P. Schmidt, Oxid. M et. 29, 289 (1988). 26. X. G. Zheng and D. J. Young, Oxid. Met. 42, 163 (1994). 27. M. J. Bennett, G. W. Horsley, and M . R. Houlton, Proc. Fundamental Aspects of Corrosion Protection by Surface Modi® cation (The Electrochemi cal Society, Pennington, NJ, 1984), p. 282. 28. M. J. Bennett, J. A. Desport, and P. A. Labun, Oxid. Met. 22, 291 (1984). 29. G. R. Johnston, High Temp. High Press. 14, 695 (1982). 30. P. Carter, B. Gleeson, and D. J. Young, Acta Metall. 44, 4033 (1996). 31. L. C. Brown and J. S. Kirkaldy (1964), Trans. Met. Soc. AIME 230, 223 (1964). 32. D. P. W hittle, G. C. Wood, D. J. Evans, and D. B. Scully, Acta Metall. 15, 1747 (1967). 33. B. D. Bastow, D. P. Whittle, and G. C. Wood, Oxid. Met. 12, 413 (1978).