01374 Correlation of Oxidation Carburization and Metal Dusting Controlling Corrosion by Corrosion...

CORRELATION OF OXIDATION, CARBURIZATION AND METAL DUSTING; "CONTROLLING CORROSION BY CORROSION"

R. Kirchheiner Manager of Central R&D Schmidt & Clemens GmbH + Co. KG.

Edelstahlwerk Kaiserau, D-51779 Lindlar, Germany

J. L. Jim~nez Soler Manager of Local R&D Centracero S.A.

Carretera Estella-Vitoria, Km. 12, E-31280 Murieta / Navarra, Spain

ABSTRACT

Metal dusting has been observed in industrial high temperature process components such as waste heat boilers in reformers and gas pre-heaters in direct reduction plants. This corrosion phenomenon is a disintegration of metallic materials into carbon (graphite) and metal dust. The degradation takes place in strongly carburizing atmospheres with carbon activities greater 1 (ac>>l) and at intermediate temperatures between 400 °C - 900 °C. Slight modifications in the process conditions may affect significantly the occurrence and extent of metal dusting attack.

To correlate high temperature corrosion mechanisms, e. g. oxidation, carburization and metal dusting for alloy design purposes, results of laboratory investigations are compared with field observations and experience. Most of the work described herein is concerned with gas mixtures of hydrogen and carbon monoxide used in the direct reduction of iron ores (DRI). Considerations are given concerning the process variables as well as on the recent development of new high performance alloys which should form an inherent stable oxide layer formation.

A controlled formation of a dense, protective, self-healing alloy surface oxide gives reasonable protection of plant components under most aggressive metal dusting parameters.

Keywords: syngas plants; reformers; hydrogen; carbon monoxide, direct reduction of iron ores; gas pre-heaters; waste heat boilers; high temperature corrosion; carburization; oxidation; metal dusting; high performance alloys; nickel-base alloys; HK 40; HP 40; 45 Micro; G 4879; experimental alloys; coatings.

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INTRODUCTION

Metal Dusting Mechanisms of High Performance Alloys

Metal dusting is one of the most common high-temperature corrosion process observed in petrochemical and reforming industries and direct reduction plants. Metal dusting was also found in the heat treating industry and in coal gasification plants. This phenomenon occurs in strongly carburizing atmospheres mainly in a most critical temperature range between 450 °C - 750 °C. Corrosion products are carbonaceous deposits (coke) containing very fine metallic particles.

Actually there is no Fe- or Ni-base alloy for which one can be sure that it is absolutely resistant to metal dusting attack. Research work is focused on the reaction chain from the formation of alloy substrate oxide layers and deteriorating conditions. Accordingly carbon and coke deposition, carbide formation and decomposition are investigated.

Metal dusting as a catastrophic form of carburization has been investigated by Grabke et al. 1'2'3'4 in flowing CO-H2-H20 mixtures at carbon activities ac>>l.

For the nickel-base alloys, Grabke et al. proposed the following mechanism:

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3. 4.

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carbon transfer from the gas phase and dissolution of carbon into the metal phase at oxide defect sites formation of a supersaturated solution of carbon in the Ni-Fe matrix deposition of graphite on the alloy surface in different orientations growth of graphite into the metal phase by carbon atoms from the solid solution, attaching to graphite planes growing vertical to the metal surface destruction of the metal phase by the inward growing graphite under transfer of metal particles into the 'coke' layer further graphite deposition from the gas phase on these catalytically active metal particles (Fe, Ni)

The chemical reactions producing such high carbon activities, e. g. in reformer applications are

Hydrogen reformer: CH4 + H20 = CO + 3H2 (1)

in a temperature regime between 550 °C and 850 °C. This gas composition may change due the following reactions:

Reaction: CO + H2 = H20 + C (2) Boudouard reaction: 2 CO = CO2 + C (3)

In cases where carbon activities are calculated to be ac>> 1 for reactions (2) or (3) metal dusting may occur as a catastrophic reaction of carbon monoxide and carbon with metallic materials.

High performance alloys bear the potential to form a protective chromia scale (Cr203), even at high carbon activities and concurrently low oxygen partial pressures. On these high-alloy steels, attack develops locally, where oxide layers have failed. Then at first internal carbide formation occurs of the stable carbides M7C3, this carburization also retards the over-saturation and start of metal dusting somewhat. Metal dusting starts with the outgrowth of coke from shallow pits that widen with time into hemispherical pit geometry or grooves.

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Controlling Corrosion by Oxidation

Corrosion mechanisms like oxidation and carburization govern metal dusting. Uniform oxidation of the alloys definitely inhibits carburization and metal dusting temporarily through formation of protective scales. Scale formation is supported by high chromium contents, medium silicon contents, small grain sizes and induced surface deformation. All these factors make chromium and other stable oxide formers (Mn, Si, AI ...) diffuse faster to the surface and therefore increase the re-healing capacity of the dense protective chromia layer.

Controlling these most favored oxidation mechanisms therefore bears the potential to control metal dusting. 5

According to G. Lai, Figure 1 shows the central role of oxidation in directing all other forms of high temperature corrosion. Oxidation therefore plays a two-fold role in corrosion processes:

1. external and internal oxidation, corroding low chromium containing alloys 2. stable chromium oxide layer formation (Mn, Cr-spinel, chromia / Cr203) protecting high

chromium bearing alloys.

However, "real" metallic alloys contain a lot of microscopic defects which upon creep or thermal cycling tend to destroy formerly grown oxide layers. Figure 2 acc. to Engel shows the microscopic appearance of alloy structures with potential defect sites. These local defects, e. g. grain boundaries, brittle carbon films, non-metallic inclusions and lattice dislocations are able to disturb protective scale formation and cause its failure by cracking or spelling.

Figure 3 shows the potential reactions between gas and metal, whereas in phase I carburization under oxidizing conditions is delayed by formation of a substantial chromium-rich external scale. The initially formed dense oxide layers (phase I) are destroyed locally in phase II and give way for the access of high carbon activity gases.

The existing carbides are primarily chromium carbides. The phase CrzC3 has a much lower gas permeability than M7C3 which therefore protects the matrix a~ainst carbon ingress and diffusion if this phase remains in a fine dispersed form, see Figure 4 ~'.

In principle no material, able to dissolve carbon, is immune to disintegration by metal dusting, either by intermediate formation of a non-stable carbide M3C or by direct decomposition of the supersaturated solution of carbon into metal and graphite. Long-term corrosion resistance therefore is guaranteed by a stable oxide layer formation only. Retardation is possible by low carbon solubility and diffusivity and a concurrently high content of carbide forming elements.

Therefore the alloy composition is one of the major factors governing corrosion resistance under metal dusting atmospheres.

This paper is addressed to the very beginning of the reaction chain, the formation and interaction of protective oxide layers on high performance alloys. Within this research work it is demonstrated that the concerted action of an optimum alloy substrate composition with specifically controlled conditions for the formation of oxide layers is capable reducing metal dusting attack.

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EXPERIMENTAL PROCEDURE

Laboratory Investigations - Oxidation

Oxidation as a Function of Alloy Composition and Oxygen and Carbon Activities An investigation was made to elucidate the effects of alloy composition on oxidation, carburization / decoking and post-coking oxidation treatments B. The kinetics of air-steam oxidation, coking treatments at 1,000 °C and air-steam re-oxidation treatments after decoking were investigated at 900 °C. The materials investigated are listed in Table 1. The test gas atmosphere for oxidation tests of virgin materials was generated by passing an air-stream through a near boiling water reservoir producing a partial steam pressure of 0.046 MPa. The flow rate was adjusted at 400 ml/min air and 330 ml/min steam. Carburization and coking was executed at 1,000 °C in an 89 % hydrogen / 11% propylene atmosphere for two hours at a flow rate of 517 ml/min, followed by a decoking procedure at 900 °C again in an air-steam atmosphere. Post-coking air steam oxidation was performed at 900 °C for 24 hours to determine the different kinetics for the alloys. Specimens used in this investigations were cut from industrial scale centricast tubes. In the case of HP 40 a modified alloy was also made with an increased silicon content of 2.6 % Si, named alloy HP 40 hSi. Also an experimental Alloy "A" was included, based on a 30Ni-18Cr iron matrix bearing 1.7 % aluminum amongst other element additions.

Results are given in Table 2 and Figure 5. Table 2 shows the rate constants for air-steam oxidation and post-coke air-steam oxidation. The rate constants indicate that the standard reference Alloy HP 40 has the highest rate of reaction whereas the high silicon modified HP 40 shows a marked decrease in reaction speed. The lowest reaction rate, however, occurred with experimental Alloy "A" where reaction was hardly measurable after the test. The morphology of the scale formed on experimental Alloy "A" reveals a patchy sub-micron thick external scale with a certain degree of internal oxidation in near surface areas. Results of x-ray diffraction analysis (XRD) shows that oxidation in air-steam mixtures produced a duplex scale on virgin specimens at 900 °C. An external layer of chromia (Cr203) and an inner layer of iron-chromium spinel Fe Cr204 was found on all materials. These findings might be explained with ongoing oxidation of the initially formed spinels to chromium oxides under high partial pressures of oxygen. The coking procedure produced M7C3 and carbon precipitates in a near-surface region. Experimental Alloy "A" showed minor internal carburization at this temperature. After the post-coking air-steam oxidation at 900 °C traces of MTC3 were detected along with chromia and iron chromium spinels. Experimental Alloy "A" also showed a stable and thin layer of alumina (AI203). A preliminary conclusion is that alloys containing no aluminum are susceptible to internal carburization with parabolic weight gain kinetics. Experimental Alloy "A" showed, that the aluminum content was sufficient to suppress carburization almost completely in these short-time tests even at high carbon activities. High silicon contents were found slightly to reduce the weight gain during carburization / coking at 1,000 °C. Silicon acts in altering the solubility and diffusivity of carbon into the matrix metal.

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Laboratory Investigations- Carburization

Carburization as a Function of Oxygen Partial Pressures Cast austenitic steels were tested in the form of samples made from investment casting and centricast materials 1°. Prior to testing they received a #600 grid finish grinding. Tests were performed in flowing gas mixtures at temperatures of 900 °C, 1,050 °C and 1,150 °C. The two gas mixtures were mixed to form

1) a highly reducing and strongly carburizing atmosphere / 89 vol.-% hydrogen + 11 vol.-% propylene (pre-dried)

2) a mildly oxidizing gas mixture / 89 vol.-% hydrogen + 11 vol.-% propylene + steam (H20).

Carbon activity of both gases was unity (carbon deposits on the samples) but, in comparison to mixture 1, the oxygen potential in the second gas mixture was sufficient to oxidize both chromium and silicon. However, iron and nickel could not be oxidized at these low partial pressures (10 .24 atm / 900 °C, 10 .2o atm / 1,050 °C). At 900 °C the water content of gas mixture 2 amounted up to 3 x 10 .3 atm., whereas at 1,050 °C the steam partial pressure was 6 x 10 .2 atm.

Carburization kinetics for a centricast Alloy HP 40 Micro (C6) as a function of time is given in Figure 6 in reducing gas atmospheres as well as in oxidizing gases 1°. For temperatures between 900 °C and 1,150 °C the kinetics is parabolic and can be described by the expression

X + X' = 2 . k p - t (4)

X = carburization depth in micrometer t = time at temperature X '= constant k~ = parabolic, constant.

In the "reducing gas" atmospheres, it was found that specific alloying additions decrease carburization rates markedly. Comparing standard Alloy HP 40 Micro with 24 % Cr to an alloy with 30 % Cr in the matrix at 900 °C the carburization rate constant is decreased from

kp = 1.6 x 10 -3 (prn z x h -1) down to kp = 0.9 x 10 "3 (prn = x h "1)

(5) (6)

More or less the same reduction in carburization was found with alloys just differing in their silicon content, e. g. Alloy HP 40 1.8 % Silicon to Alloy HP 40 hSi with 2.6 %. Silicon is effective in altering the solubility and diffusivity of carbon in the fcc-matrix. Aluminum was found to be even more effective in forming a dense and stable external oxide layer.

Carburization in Pack Cementation Tests To compare the properties of centricast alloys, particularly concerning their carburization resistance, tests were performed in a carbon pack. Test bars were cut from the tube wall and packed in a carburizing powder at temperatures increasing stepwise (+ 25 °C) from 950 °C to 1,175 °C. The total test duration was 260 hours. Test results clearly differentiated between the standard HK 40 and the 45 % nickel - 35 % chromium containing Alloy 45 Micro, see Figure 7. Due to a better balanced composition Alloy 45 Micro with its optimum chromium content showed almost no visible carburization at 1,100 °C. Also the higher nickel content contributes to a lower rate of carburization.

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Laboratory Investigations - Metal Dusting

The Effect of Alloy Composition, Temperature, Pressure and Surface Rouqhness The intention of an ongoing study is to obtain additional information concerning the role of temperature and surface engineering on centricast high performance alloys. Four alloys were tested in a flowing hydrogen - carbon-monoxide - carbon-dioxide - methane - nitrogen gas mixture at 650 °C and 750 °C. Initial test periods were set to 180 hours. Mass gain or loss was measured accompanied by microstructural and surface analysis. The alloy compositions are given in Table 1, i. e. Alloys 45 MTZ (C12), 45 LC (C13) and HP 40 MW (C8), HP 40 MT (C9). These four alloys are grouped in two families, e. g. HP 40 and Alloy 45. Predominating differences are in the nickel and chromium contents (8 = 10 % Ni/Cr) as well as in minor alloying constituents.

The composition of the test gas was chosen close to the gas atmospheres produced in methane reforming plants and used in direct reduction of iron ore (DRI). The gas mixture is composed out of 70 % H2, 16 % CO, 3 % CO2, 8 % CH4, 1.3 % N, 2 % H20. The gas flow was maintained at I I/min.

Test pieces out of centricast tubes were cut in rectangular blocks (10 mm x 5 mm x 5 mm). The machined surfaces were subsequently finished with # 50 or # 600 grade silicon carbide paper under controlled conditions of grinding speed and force.

After heating the top-load autoclave (200 °C/h) to 650 °C and 750 °C respectively, the system pressure was adjusted to 0.1 Mpa (1 bar). Experiments were interrupted periodically for the purpose of inspection and weighing. One sample of each alloy was removed for analysis using XRD, SEM and optical methods.

To explain the intermediate results some preliminary remarks are essential. Thermodynamical calculations show that the carbon activity values obtained in the studied gas mixture varied, as expected with temperature and pressure. Up to 600 °C, there is a marked difference in carbon activities, yielding the highest ac values (ac ~ 230) for 1.0 MPa (10 bar) pressure, see Figure 8. Beyond a temperature limit of about 700 °C, carbon activities are obtained significantly less than 50. At around 800 °C, irrespective of the pressure, the carbon activities are below 25. Considering thermodynamics, the graphs of phase equilibrium of the main oxide-forming elements of the materials have been marked (asterisk) for the corresponding gas mixture of these experiments, see Figures 9, 10 and 11. For chromium, see Figure 9, the Cr203 stability field is reduced with the temperature. The balance given by the gas mixture is placed close to the Cr3C2 carbide. This supports the hypothesis that Cr3C2-carbides (surface) and partially Cr7C3 (eutective primary carbides) are transformed into chromium oxide Cr203 by a non-inhibited oxidation. CrTC3- carbides also will be transformed into 0r2306 (secondary carbides) during the ageing process. This is probably one of the deterioration mechanisms of the chromia scale, e. g. the oxidation alongside internal carbides. According to Figure 10 silicon dioxide is the stable phase in these experiments. In the vicinity of the metal dusting peak temperatures, e. g. between 600 °C and 800 °C the stability field of silicon dioxide is increased. This contributes to literature results, referring to a better metal dusting resistance with increasing silicon contents.

In Figure 11 it is shown that the gas mixture point is placed close to the borders of Fe3C (cementite) and Fe304 (hematite). Thermodynamically therefore it is possible that Fe3C carbides may be oxidized into Fe~O4 causing micro cracks in a protective chromia scale.

A deterioration mechanism is schematically referenced in Figure 12 with formation of unstable carbides at the surface followed by decomposition. These preliminary conclusions are in clear opposition to the literature referenced above. Further tests under controlled conditions are necessary to prove or disprove this hypothesis.

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In the course of this investigation grid # 50 and # 600 silicon carbide paper was used to check the impact of surface conditions on the metal dusting process. Preliminary results of weight change analysis with Alloy 45 M are given below. Figure 13 shows both alloy modifications with a weight gain of < 0.1 mg/cm 2. This indicates that the preformed scale is protective or if cracked, the re-healing tendency is rather rapid. Weight gain under these conditions therefore is related to the formation and growth of oxides. Contrarily at 650 °C Figure 14 shows a mass loss of about 0.2 mg/cm 2 at equal surface finish # 50. XRD spectrum analysis showed that amongst others, Cr23C~ carbides and traces of alpha phase (ferrite) were present. This true mass loss may be related to "real" pitting attack under peak metal dusting corrosion conditions. This is true just for the 750 °C test and supports the idea that a much faster chromium diffusion in the partially re- crystallized ferritic structure results in the formation of the more protective scale. These findings were proven so far also in other samples of the same investigation.

Comparing results for HP 40 modified alloys similar mechanisms seem to be predominant. Figure 15 shows a clear weight gain of about 0.4 mg/cm 2 at 750 °C whereas Figure 16 shows significant weight loss for the "low" temperature 650 °C, finish # 50.

Microstructural investigation confirmed real re-crystallization at test temperatures of 750 °C and test duration of 180 hours.

XRD analysis of all investigated samples demonstrated the presence of Cr203, M2306, M3C as well as (Cr, Mn)304 for HP 40 and Alloy 45. A more in depth analysis of Alloy 45 tested at 750 °C (SIC # 50) demonstrates a superior dense and protective oxide layer after conclusion of the test.

Therefore it may be concluded so far that all theoretical considerations are fulfilled in the frame of this defined laboratory investigation.

Metal Dusting Investigations of Cast Nickel Alloys Carburization and metal dusting tests have been performed in the central laboratories of a DRI plant. The test conditions were adjusted to

• temperature 700 °C • pressure0.15 MPa • gas flow 55 I/min • gas composition reformer gas with 6 % methane (CH4). Rectangular samples were cut (15 mm x 8 mm x 4 mm) and ground with # 180 followed by a

chemical etching procedure. This special etching procedure was applied to reduce so-called incubation times by activating bare metal surfaces. The materials tested were HP 40 (C5), G 4879 (C14), G 4868 (C3) and Alloy 45 M (C11).

Weight gain or weight losses ranged in between -1 x 10 -4 up to +7 x 10 "4. With exception of Alloy 45 all other alloys have shown some local pits or uniform attack over the entire surface. Also randomly distributed carbon deposits were detected on almost any sample again with the exception of Alloy 45 Micro.

Metal Dusting Behavior of Wrought Nickel Base Alloys In Reference 11 H. J. Grabke, J. KlOwer et al. reported about the metal dusting behavior of several nickel based alloys at 650 °C in carbon monoxide (CO) and hydrogen (H2) environments with a carbon activity ac >> 1. Iron nickel chromium alloys showed susceptibility after rather short test times. Wrought nickel base alloys with chromium contents of about 25 % seem to be resistant in exposures even after 10,000 hours exposure. Table 3 shows a selection of alloys tested as well as their surface condition.

Specific consideration was given to Alloy 601 which was tested in three different conditions:

• ground • electro-polished and • black (scale from hot rolling)

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to study potential impacts of surface preparation. The activities of carbon were adjusted to ac = 14 for the first 5,000 hours and to ac = 19 for the second period of a further 5,000 hours. The oxygen partial pressure was sufficient to enable formation of chromium-oxides but not the oxidation of nickel and iron. Tests were performed at 650 °C with frequent interruptions to study corrosion effects over time. Results are given in Figure 17. Microscopic examinations revealed, that metal dusting attack starts locally by the formation of pits, in areas where oxide scales cracked or spalled. Under these severe test conditions several alloys had to be removed from the test prior to the intended limit of time (10,000 h). These alloys are Alloy 800 H, HP 40, 600H and Alloy 601 black and electro-polished. Samples of centricast HK 40 ranged in the same catastrophic metal dusting attack as HP 40 and not far from Alloy 800H.

The high nickel Alloy 601 (60 % nickel) reacted quite sensitively to different surface preparations. The as-ground condition was the most resistant against metal dusting, e. g. two orders of magnitude lower than for the same alloy composition in the oxidized condition. Alloys 602CA and 690 both bearing chromium contents of and above 25 wt.-% showed only minor mass loss rates after 10,000 hours exposure of between 10 .5 mg/crrF x h and 10 -6 mg/crn = x ho All high performance nickel alloys in the wrought condition, e. g. Alloy 602CA and Alloy 690 had developed dense oxide layers in the course of the exposure period. The pre-dominant mechanism of protection seems to be supported by minimum chromium contents of 25 % with Cr backed up by certain amounts of aluminum and rare elements addition.

FIELD EXPERIENCE AND PLANT DATA

Most pronounced cases of metal dusting of high alloyed materials were found in gas pre- heaters of direct reduction plants, see Figure 18. Pre-heaters are designed to heat up the reducing gas atmosphere, hydrogen (H2) plus carbon-monoxide (CO), to an optimum temperature range for the iron ore reduction. The extremely high carbon activities along with cyclic load operations result in significant metal dusting attack on lower alloyed stainless steels. A few cases from practical experience will be described below.

Direct Reduction Plants

Three different single trains of DRI plants have been checked due to their material performance in gas pre-heaters.

Direct Reduction Plant A / DRI A # 1: capacity:

operation:

pre-heater:

alloy: observations:

1,000 MTPD (metric tons per day) > 20 years with standard alloys / several years with Alloy 45 (sulfur inhibition not reported/known)

56 tubes OD = 152 mm, L = 10,000 mm + 14 u-bends tubes and bends in Alloy 45 (45-Ni/35Cr) No material failure / metal dusting since recent re-vamping.

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# 2: capacity: operation:

pre-heater:


2,000 MTPD > 15 years with standard alloys / several years with Alloy 45 (sulfur inhibition not reported/known)

96 tubes OD = 152 mm, L = 10,000 mm + 24 u-bends tubes and bends in Alloy 45 (45-Ni/35Cr) No reported material failure.

# 3: capacity: operation:

pre-heater:


2,000 MTPD < 5 years (auto-reforming reactor) (sulfur inhibition not reported/known) 68 tubes D = 152 mm, L = 10,000 mm + 68 tubes D = 100 mm, L = 10,000 mm + 68 u-bends tubes and bends in Alloy 45 (45-Ni/35Cr) No reported material failures since installation.

Direct Reduction Plant B / DRI B capacity: (not reported) operating conditions: acc. to Table 4 operation: < 12 months material: 34 pipe loops per box (4 boxes) alloy: material protection:

observations:

HK 40 / HK 40 mod pre-oxidation treatments and dimethyl-disulphide (DMDS) inhibition (max. 30 ppm) After a nine months' period of operation, in four out of 14 areas severe metal dusting attack was observed. Metal dusting attack occurred at weld protrusions of the root bead up to 8 mm in width and to a depth extension of 1.3 mm on the bare tube surface. In the high silicon modified HK 40 M (< 2 % Si) the frequency of attack was diminished but also present. Concerning the performance of HK 40 rood (<_ 2 % Si) several observations of weld cracking were made.

Along with weld cracking an increased void formation in the inner parts of the tube metal wall was reported. After only nine months of operation in a defined temperature regime this premature creep has shown up far to early.

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Direct Reduction Plant C / DRI C capacity: operation: gas atmosphere:

> 1,000 MTPD (metric tons per day) < 10 months H2 = 70 %, CO = 15 %, CH4 = 10 %, CO2 = 1.5 %, N2 = 2 %, H20 = 1.3 %

process gas temperature: 700 °C - 800 °C process pressure: 0.5 - 0.6 MPa, abs. carbon activity: >> 1 material: tubes OD = 168, L = 9,500 mm alloy: HP 40 Nb

Catastrophic forms of metal dusting attack were observed after just six months of operation. The inside surface was covered by multiple overlapping hemi-ellipsoidal pits with an approximate peak height of more than 5 mm and a diameter of 10 - 15 mm. The pit surface was partially covered by a dark grey appearing oxide in the center whereas the rim showed a metallic appearance. The remaining wall thickness was less than 4 mm in this specific section, see Figure 19. The rim of the pit contains a sponge-like surface with a glassy appearing surface structure, see Figure 20. The surface is covered with a high density of micro-sized pinholes with a maximum diameter of 6 pm.

In adjacent areas of the pits, a network of chop-stick-like carbides accumulated as shown in Figure 21.

A first hint to the synergistic action of oxidation, carburization and metal dusting process is given in Figure 22 and Figure 23. In these figures oxide infiltration is traced between the blocky carbides in the near-surface layers. Figure 23 clearly shows the pattern of oxygen / oxide distribution in that area. The explanation by Kleemann 12 was given as follows: "... in the top M703 carbide zone are clearly visible the inter-carbidic oxidation spikes extending from the gas metal interface downwards forming a hemi-ellipsoidal pit-like configuration. Visible in this micrograph is a thin layer of surface oxides and an infiltration pass of inter-carbidic oxidation spikes penetrating into this zone to a depth of < 100pm. The internal oxidation spikes engulfs an area forming particles. ''12 The total depth of carburization was found to be around 1,200 pm.

Figure 24 shows the significant loss of chromium in the profile of the surface area. This loss of elemental chromium corresponds with the massive carburization and carbide formation.

Taking into account the inter-carbidic attack mechanism, it can be stated, that the different carbide precipitation pattern of the tungsten alloyed G 4879 may be one of the reasonable explanations for the better behavior under metal dusting conditions.

Direct Reduction Plant D / DRI D capacity: operation: design:

materials: observations:

2,000 MTPD (metric tons per day) < 2 years gas pre-heater with 90 hairpin tube bundles 86 mm x 6 mm x 10,000 mm (Figure 25) Alloy 45 (Ni45Cr35) Due to a number of non-scheduled incidents, the reformer module along with the inter-connected gas pre-heater were exposed to numerous load cycles. Up to now there is no reported metal dusting problem in this plant. More in depth information will be generated soon.

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Ethylene Cracker Plants The plant experience described herein was selected to show the excellent stability of

protective oxide layers formation under conditions of cyclic carburization attack. Carbon activities are low (ac=l) compared to real metal dusting environments (at>>1).

Ethylene Cracker Plant E capacity: operation: design: temperature: decoking: materials:

(not reported) > 36 months ethylene cracker (naphtha-feed) 1,050 °C - 1,070 °C frequent (60 days' interval) HP 40, G 4879, Alloy 45

The materials have been tested under carburizing ethylene cracker conditions with frequent burn-off of coke deposits for more than 36 months 13.

Carburization proceeded into the materials HP 40 and G 4879. The tungsten bearing Alloy G 4879, compared to HP 40 showed just 40 % wall thickness penetration. HP 40, after 36 months of exposure was carburized up to 75 % wall thickness. It may be assumed that the slightly higher chromium content (+ 3 %) associated with the much higher nickel content of approximately 48 % was the reason for the good performance of G 4879.

In contrast, alloy Ni45Cr35 (Alloy 45) showed almost no signs of carburization. Comparing in Figure 26 the alloys HP 40 and Alloy 45 it is clearly eminent, that after 36 months of frequent alternating operations between strongly carburizing and oxidizing conditions HP 40 has no longer a protective barrier against carbon ingress. Alloy 45, bearing 35 % chromium in an austenitic matrix, showed an almost defect-free and therefore protective oxide layer. Also the close-to-ideal double-layer structure of chromia (Cr203) followed by a sublayer of silica (SiO2) is shown in the case of Alloy 45. The microscopic evaluation showed in a rather dense oxide layer of 40 pm total thickness. In Figure 27 the negative contrast pictures show the presence of only a few iron and nickel spots in the chromia layer (x-ray map). This "low density" pattern of catalytically active iron and nickel spots may be the reason for a decreased catalytic coke precipitation and minor carburization under frequent high carbon activities.

DISCUSSION

The discussion will follow the effects of oxidation and carburization as well as specific metal dusting parameters, e. g. impacts of temperature on microstructure of alloys, variation of diffusion speeds through cold deformation of fcc alloys and metal dusting peak intensity areas. In general, the impact of alloy composition and structure will be discussed as a measure to control metal dusting corrosion.

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Oxidation

Figure 28 demonstrates the controlled formation and alteration of protective oxide layers on high performance centricast tubes. In the course of a "tube life" several stages of high temperature - high oxygen partial pressure regimes are passed. In the very beginning this starts with the melting of refined metals and ferric-alloys in induction furnaces. Secondly, centricasting under air creates strongly oxidizing conditions. After the birth of the raw tube, inside machining and cleaning, transforms thick oxide layers into rather thin and very protective chromia and chrome spinel oxide films. Growth of these thin silvery shining layers into colorful yellow to deep blue heat tints is encountered under shop and field welding operations. Especially the brown and yellow colored heat tints must be removed prior to service because of the non-protective character of these low temperature oxide films. Most engineering companies issue operation manuals with a defined pre-oxidation condition before the "start-of-run" phase in high temperature gas reactors. In these controlled oxidation conditions a conversion of thin oxide films into thicker visible oxide layers is achieved. In the course of normal operation cracking carbon precipitation takes place under partial oxygen depletion in reducing gas atmospheres. Under these conditions a more uniform alteration of the protective oxide layers is achieved, e. g. general thinning. Under normal operation, e. g. creep conditions, local defect sites in oxide layers result in catalytically active spots exposing iron and nickel.

It is decisive that the substrate alloy enables the delivery of sufficient amounts of oxide layer forming elements by outward diffusion from the substrate alloy pool. Therefore the availability of elements like chromium, silicon and aluminum plays a major role in long-term protection of high performance centdcast tubes under metal dusting conditions.

Oxides, thermodynamically considered, are likely to form on a metal when the oxygen pressure of the atmosphere is greater than the oxygen partial pressure in equilibrium with the metal oxide under consideration, e. g. chromium oxide (chromia / Cr203). Lai s referenced the ~artial pressure of oxygen in equilibrium with chromium oxide (Cr203) at 1,000 °C to about 10

atmospheres, see Figure 29. If no molecular oxygen is available in the gas atmosphere, another potential source of oxygen is the water vapor dissociation:

H2 = = H2 + ½ 0 2 and the CO2 dissociation CO2 = CO + ½ 02 or from CO = _C + ½ 02.

For that reason in any chemical or petrochemical high temperature reaction sufficient amounts of oxygen are available from the atmosphere to form selected metal oxides.

Chromia and alumina forming alloys are considered as most resistant against high temperature corrosion attack after controlled initial formation of their oxide layers. These protective oxides are growing slowly. A ranking of the parabolic growth rates of chromia (Cr203), silica (SiO2) and alumina (AI203) shows chromia with the relatively fastest rate of formation and alumina and silica growing even by magnitudes slower. Despite the low rate of formation, aluminum oxide is the most protective compound on high temperature corrosion resistant alloys. The rather thin layers of alumina are atomistically highly ordered structures with little defects, rendering them nearly impermeable for carburization attack. Chromium oxide (Cr203) is converted and evaporates particularly in the temperature regime > 1,050 °C, therefore loosing its protective nature above this temperature. This is a limiting factor for high temperature resistant alloys in case that they rely just on chromium as a stable oxide former. Optimum compositions of high temperature resistant alloys for regimes over 1,100°C contain chromium, silicon and aluminum as well as minor doses of rare earth elements to improve oxide scale adherence (pegging effect). The above mentioned own investigations on oxidation behavior of HP 40 /

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HP 40 hSi and experimental Alloy "A" confirmed these statements. Based on a comparable chromium matrix alloy content of 24 % Cr in HP 40 and HP 40 hSi, the difference of approximately 1 % silicon content shows a marked decrease in the speed of oxide layer formation. This rate was further decreased when the experimental Alloy "A" with just 18 % Cr was supplemented with 1.7 % aluminum. A "rule of thumb" is, that under these specific conditions, a thin and therefore more "ductile" oxide layer is much more favorable than a thick, colored oxide layer with numerous micro-fissures resulting from the growth process. Although the above mentioned experiments were made to check different alloys' behavior versus load cycles in the ethylene cracker industry some of these mechanisms may be applied to components in reformers and gas pre-heaters which are, in general, operated at 100 °C - 200 °C lower temperature regimes. As high silicon contents are able to diminish solubility and diffusivity of carbon in austenitic matrix lattices, Alloy HP 40 hSi showed less weight gain in the intermediate carburization / coking step compared to the standard reference Alloy HP 40. However, the experimental Alloy "A", due to its balanced chromium, silicon and aluminum contents showed almost no visible signs of carburization. Overall, in this laboratory investigation near-surface carbide particles were oxidized, forming an external chromium-rich scale on top of a carbide- depleted zone in the subsurface.

Carburization

Thermodynamically considered, an alloy might be carburized when carbon activity of the environment is greater than the carbon activity at the bare metal surface. Different compounds in the corrosive high temperature atmosphere contribute to the carbon activity varying as a function of temperature 5.

According to Kleemann 7 and Mitchell et al 1°, high temperature corrosion under oxidizing and carburizing conditions (< 1,050 °C) will result in formation of a protective chromia scale and an interrelated carbide denuded zone below this scale.

In absence of a protective oxide layer the ingress of carbon is diffusion-controlled. With increasing carbon content in the alloy, chromium carbides of different compositions will be formed starting with M2306 which later converted to M703 on the alloy surface sometimes 0r302 is observed. In high performance alloys, reactive elements like niobium, titanium, tungsten and others also form carbides. These so-called primary carbides are formed intentionally during the casting and centricasting process of such high performance alloys to increase creep resistance.

Defect-free and dense protective oxide layers with tight adherence to the alloy matrix should completely suppress carbon ingress, see Figures 6 and 7. Alloys with sufficient chromium contents of more than 25 % Cr show a remarkably better carburization resistance under different load conditions. The optimum carburization resistance is shown nowadays by Alloy 45 (45 % Ni + 35 % Cr). In the absence of thermal shock conditions as well as in secondary creep regimes with minimum deformation rates, these layers should maintain their protective nature for years of operation. With reference to Grabke 8 there is no way that carbon diffuses through dense oxide layers. It was shown that carbon has no solubility in technically relevant oxides. Carburization profiles and gradients depend on numerous factors in a complex manner. The detrimental effect of carburization is secondary in nature, e. g. a loss in ductility due to carburization results in premature cracking, and therefore accelerated high temperature corrosion can start within the catalytically active crack tip areas. Interaction of creep and carburization was reported to be less critical in a creep regime of smaller than 10 -8 per second. Below these creep rates, the kinetics of oxide layer re-healing are fast enough to close microscopic fissures.

A sequence of steps was reported by Grabke 8 in slow strain rate experiments with ~.=4.109/s

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1. grain boundaries sliding results in repeated cracking of oxide layers 2. carbon up-take while the crack is open (protective oxide layer not yet restored) 3. carbon diffusion into the grain boundaries 4. precipitation of carbides 5. depletion of the gradient area in chromium 6. grain boundary carbide conversion to oxides

Higher strain rates, e. g. L= 1.3.10-T/s resulted in a more uniform carburization. The carburization resistance of nickel-chromium alloys depends on alloy composition and

controlled surface engineering. As shown in Figure 30 there is an optimum content between 45 and 80 % nickel in austenitic matrix alloys. This function relies on the decreased carbon diffusivity in nickel alloyed materials. A decisive side effect in alloying with nickel is the increased ductility of the lattice. This results in a lower crack evolution of protective oxide scales. Stepwise, these increased nickel contents have been realized in cast alloys from 20 % Ni in alloy HK 40 to 35 % Ni in HP 40 up to 45 % Ni in high performance alloys like Alloy 45 Micro. Norton et a114. have shown that above the absolute content of nickel, a ratio of nickel to chromium plus iron is decisive for carburization resistance. Figure 31 shows the ranking of cast and wrought alloys according to their nickel chromium ratio after carburization at 1,100 °C in pack cementation tests (32 days). It is obvious, that high carbon, high chromium cast grades show satisfactory carburization results up to nickel contents of 50 % compared to wrought alloys. Hot forming properties of wrought (low carbon) nickel base alloys require a somewhat higher nickel content when chromium is fixed above 25 % due to metallurgical reasons. A local concentration of > 4 wt.-% alumina is required to suppress carburization effectively 8. Resistance against carburization actually is found in nickel base alloys with aluminum contents above 2 % or well defined additions of silicon and rare earth elements. A compromise regarding weldability is obtained with aluminum contents of 2.5 % in bulk matrix of the nickel alloy.

In case of re-oxidation with sufficient partial pressures of oxygen, silicon creates a sub-layer of silica (SiO2) under the preformed chromia structure. This substructure contributes to a higher carburization resistance in blocking carbon ingress and concurrently increasing the re-healing tendency of local defect sites.

Metal Dusting

Oxidation, carburization and metal dusting are linked high temperature corrosion phenomena in a "natural reaction chain". Metal dusting is catastrophic form of carburization that deteriorates alloys in high temperature regimes with high carbon activity gaseous environments reacting to elemental carbon and metal dust. This description relates metal dusting to a specific form of carburization. Generally, carburization of the metal matrix precedes the metal dusting process, in the case of chromium-containing alloys the formation of carbides such as M23C6 and M7C3 takes place in a zone of internal carburization 3. Precipitation of fine particles of M23C6 generates stresses which lead to an accumulation of dislocations surrounding the particles. Chromium is also consumed by oxidation, the resulting chromium depleted zone is easily attacked by metal dusting, see Figure 16.

As there is no intermediate metallic carbide being formed in the metal dusting of high alloyed nickel base materials there is another decisive mechanism in the reaction front. Graphite planes are growing in an oriented manner into the metal matrix and therefore cause the distortion of protective oxide layers. Further catalytical active iron and nickel spots will result in an increased rate of metal dusting attack. Pippel, et al.gpresented the mechanism at the workshop "Coking and Decoking" in Porto in May 1999."

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As is clearly visible from Figures 13 and 14 a shift of 100 °C may completely alter the performance of alloys in laboratory investigations. Depending either on thermodynamical considerations at that temperature or on the kinetics of faster element diffusion at 100 °C higher temperatures in a partially re-crystallized structure, metal dusting oxidation leads to a weight gain at Alloy 45 (750 °C / # 50) due to oxidation and a true weight loss for the same Alloy 45 at 650 °C due to pitting. Figure 15 and 16 more or less prove this trend with the lower alloyed HP 40. In the work of Grabke, KlOwer et al. 11 also the effect of "surface engineering", e. g. grinding and pre-oxidation was mentioned as acting positive against metal dusting attack. Much more pronounced was the effect of a balanced chemical composition of these wrought materials in terms of a minimum required chromium content of 25 % + additional elements, e. g. aluminum (> 2 % AI) and rare earth metal additions, see Figure 17.

Comparing laboratory investigations as described above with real plant experience, the general trend of a positive impact of a balanced alloy composition plus surface engineering is confirmed. Catastrophic forms of metal dusting could be traced back to a lack of chromium in the matrix of the alloys and / or a frequent deterioration of the non-stable chromia layers in thermocycling operation or creep regimes. Decisive is the combined action of internal oxidation and carburization as shown in Figures 21 - 23.

From results of laboratory investigations we know that metal dusting on high alloyed materials results in hemispherical pits which are expanding into the metal matrix. These pits are filled with fine coke and metallic powder which are removed after each test period. The origin of real metal wastage is therefore difficult to determine. Most authors report about metal dusting attack in pit depth and pit density according to ASTM specifications. Only at that point when all pits grow together a general metal wastage rate in mm/year can be given.

Surface preparation of the material is another decisive parameter. Surface grinding induces local stresses in face-centered cubic, austenitic alloys. This additional stress creates stacking faults which might act as promoters for fast element diffusion in both directions. Also re- crystallization is more probable at temperatures above 700 °C in austenitic alloys. Grinding therefore results in two major effects:

a) better adherence of protective oxide layers, if peaks and valleys are not too steep and b) the deformation which leads to many dislocations, sub-grains and a fine grain microstructure

therefore increased diffusion speed of chromium, silicon, and others outward from the metallic matrix.

In contrast, under the specific circumstances of the heater of direct reduction plants, neither measures of surface engineering (DMDS / pre-oxidation) nor minor alloy modifications, e. g. in silicon could suppress metal dusting attack to 100 %. These findings support the hypothesis that the best protection against metal dusting is probably a proper alloy selection with increased capacity of a self-re-healing protective oxide layer, e. g. Alloy 45 with 35 % chromium in the matrix.

Figure 24 shows the results of a microprobe analysis concerning the chromium profile of the attacked sample. This figure clearly shows the dramatic depletion of chromium in an area down to about 1,200 pm from the gas metal interface. This phenomenon of extreme carburization and chromium depletion may be explained by additional local overheating of the damaged tube section. The resulting "local alloy composition" points to a completely different nickel iron lattice compared to the austenitic Ni-35 % Cr-25 %. The relative volume difference between the virgin cast HP 40 and the resulting chromium depleted FeNi-alloy results in microscopic scale stresses. The relieve of these stresses by fracture at the gas metal interface will produce the forces required to eject the particles from the metallic matrix.

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Further investigations are required to clarify the mechanism of the formation of the distinctly located oxidation spikes in the metallic matrix between top carbide layers, see Figures 22 and 23. Preferential oxidation of chromium normally should not occur in this oxygen depleted atmospheres.

The interaction of heavy carburization, forming coarse carbides followed by selective internal oxidation alters the structure of the attacked matrix. Protective scales do not grow over carburized regions and are partially substituted by non-protective iron oxide rich scales. Therefore, there is no longer an effective barrier against carbon ingress. These factors may result in rapid growth and deterioration of the tube wall.

A reasonable correlation between alloy composition, e. g. chromium contents and the related formation and stability of oxide layers was shown after three years exposure in Plant E! Plant E is synonymous for an ethylene cracker furnace operated under alternating conditions with medium carbon activities (at=l), naphtha cracking followed by so-called decoking cycles representing tough steam oxidation conditions. Here, a standard reference Alloy HP 40 with 25 % chromium experienced much stronger carburization compared to Alloy 45 with about 10 % more chromium in the matrix. Although, operating conditions in ethylene crackers are not comparable to direct reduction plant gas pre-heaters, a principle of material protection by stable oxide layer formation and re-healing may be derived. Unexpectedly to the believed mechanism of dense stable oxide layers formation on Alloy 45, it was found, that the density of catalytically active iron and nickel particles embedded in such an oxide layer is diminished compared to standard reference alloys. This effect corresponds to observations made in CVD- coatings.

In general the effect of surface modifications, e. g. intentionally induced stresses and modified surface roughness should be investigated further. Up to now there is no final conclusion to be made about the best suited surface roughness or degree of cold deformation in surface areas. The initial surface state is very important for the initial protective oxide formation! From a practical standpoint, it should also be considered, that an initially engineered surface from the supplier of centricast tubes will be altered rapidly under frequent load conditions of reformers or DRI plants.

The most reliable solution seems to be the balanced alloy matrix composition capable of a rapid delivery of strong oxide formers, e. g. chromium, silicon, aluminum accompanied by sufficient amounts of oxide strengthening rare earth metals. These alloys will be presented in short to the end-users and engineers of such plants.

Laboratory and plant experience are in faidy good correlation. Alloys with less than 25 % chromium are endangered by metal dusting attack in lab and field tests. The positive effect of silicon was found more pronounced in lab investigations. Under alternating load conditions of real plant cycles, silicon < 2 % was less effective, possibly due to negative impacts on creep life and weld ductility. The effect of AI in cast alloys as to date could not be evaluated under plant conditions. The excellent lab results of such cast and wrought alloys under severe metal dusting conditions are quite promising. Effects of surface engineering, e. g. stress induction by cold deformation, e. g. shot peening or grinding could not be followed over long exposure times. Also the durability of pre-oxidation treatment was not proven over long time. Lab results indicate a positive effect in first plant cycles.

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CONCLUSIONS

Metal dusting as a specific form of carburization deteriorates almost every known commercially available material in the field. Laboratory investigations are in good correlation to observations in real plants. It is the potential of the alloy system in forming and maintaining a stable passive oxide layer, that is a life guarantee for the components operated in metal dusting regimes. Concerning cast and centricast alloys, the following parameters are favorable preconditions for proper alloy performance:

1) Nickel content should be well above 40 %. 2) Chromium contents of minimum 25 - 30 % are required along with the aforementioned

nickel content. 3) A compromise between carburization / metal dusting resistance and formability / weldability

is a silicon content between 1.5 % Si and 2.5 % Si. 4) Certain amounts of aluminum and rare earth metal additions may lead to a completely new

generation of cast and centricast alloys replacing CVD-coated products in those applications.

At present the well designed Alloy 45, seems to have all required high temperature properties to withstand metal dusting attack. New alloy developments are close to commercialization with even better high temperature corrosion resistance.

Further investigations are required to clarify the mechanisms of surface engineering, e. g. grinding, shot peening and / or pre-oxidation treatments especially under the aspects of cyclic loads and long-term stability.

ACKNOWLEDGEMENTS

The extensive work of Professors D. Young (UNSW) and F. J. P6rez Trujillo (UCM) on laboratory investigation under carburizing and metal dusting conditions are acknowledged.

The contribution of L. Engel in microstructural characterization of samples after high temperature corrosion attack is worth to be named.

Finally we would like to express our thanks to Professor Grabke (MPI) for paper revision and discussion as well as to Mrs. B. Oehm-Schnippering and P. WOlpert for setting up the manuscript.

Part of this work was funded by the "Landeswirtschaftsministerium, NRW/Germany" under contract No. TPMW-335-48-18.1 (1993).

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REFERENCES

o

2.

3.

4.

5.

6.

.

o

9.

10.

11.

12.

13.

14.

H. J. Grabke, R. Krajak, E. M. M(311er-Lorenz, "Metal Dusting of High Temperature Alloys", Materials and Corrosion 44 (1993): p. 89 - 97 H. J. Grabke, "Metal Dusting of Low and High Alloyed Steels", Corrosion 51, 9 (1995): p. 711 - 720 H. J. Grabke, R. Krajak, E. M. MQIler-Lorenz, S. Straul~, "Metal Dusting of Nickel Base Alloys", Materials and Corrosion 47 (1996): p. 495 - 504 H. J. Grabke, C. B. Brancho-Troconis, E. M. MiJller-Lorenz, "Metal Dusting of Low Alloy Steels", Materials and Corrosion 45 (1995): p 215 -221 G. Lai, "High Temperature Corrosion of Engineering Alloys", ASM International (1990 - 1997) D. J. Hall, M. K. Hossain, R. F. Atkinson, "Carburizsation Behavior of HK 40 Steel in Furnaces used for Ethylene Production", High Temperatures- High Pressures 14 (1982): p. 527 W. Kleemann, "A Kinetic and Morphological Study of High Temperature Coking, Decoking and Oxidation Behavior of Heat Resistant Steels", non-published work September (1998) H. J. Grabke, "Carburization - A high Temperature Corrosion Phenomenon", MTI Publication 52 (1998) Q. Wei, E. Pippel, J. Woltersdorf, H. J. Grabke, "Microprocesses of Coke Formation in Metal Dusting", Materials and Corrosion 50 (1999), p. 6 2 8 - 633 D. R. G. Mitchell, D. J. Young, W. Kleemann, "Carburization of Heat Resistant Steels", Materials and Corrosion 49 (1998): p. 231 - 236 H. J. Grabke, E. M. M011er-Lorenz, J. KI6wer, D. C. Agarwal, "Metal Dusting of Nickel based Alloys", Materials Performance 7 (1998), p. 58 - 62 W. Kleemann, L. Engel, Investigation Report "Microstructural Analysis of a Centricast HP 40 Nb Tube from a DRI - Pre-Heater Suffering from Metal Dusting Corrosion", 01 (1999) P. W01pert, Investigation report "Carburization of four Centricast Alloys in the Ethylene Service" J. F. Norton and J. Barnes, in "Corrosion Fossil Fuel Systems, I. G. Wright, Ed., The Electrochemical Society (1983): p. 277

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TABLE 1 NOMINAL CHEMICAL COMPOSITION OF CAST ALLOYS FOR

PETROCHEMICAL, REFORMER AND DIRECT REDUCTION OF IRON ORE PLANTS

Alloy No.

C1

C2

C3

C4

C5

C6

C7

C8

C9

C10

Trade ACI Mark Design C

G 4848 HK40 0.40

G 4848S HK40 0.40

G 4868 • - - • 0.50

G 4857 HP40 0.45

G4852 HPmod. 0.40

G4852micro HPmicro 0.45

G4852micro HPmicro 0.45 high Si

HP 40 MW HPmod. 0.40

HP 40 MT HPmod. 0.4

Nominal Chemical Composit ion (mass%)

Al loy D1 • - - .0 .42

Si Mn Cr Ni Mo Nb W Zr Ti AI others

1,5 1.5 25 20 . . . . . . . . . . . . . .

2.0 1.5 25 20 . . . . . . . . . . . . . . .

2.0 1.5 30 30 . . . . . . . . . . .

1.5 0.7 25 35 . . . . . . . . . . . . . . .

1.5 1.5 25 35 . - - . 1.5 . . . . . . . . . . .

1.5 1.0 25 35 • - - • 0.8 • - - 0.08 0 .16. - - RE

2.6 1.0 25 35 0.02 0,8 . - - 0.04 0 . 1 2 . - RE

1.3 1.5 25 35 • - - . 1.5 . . . . . .

1.8 1.5 25 35 • - - • 1.5 . . . . . . . .

0.7 0.41 18 31 3.1 0.4 • - - 0.46 0.3 1.7 Exper imental 0.42 Al loy A

C l l Al loy 45 micro. - - • 0.45 1.6 1.0

C12 Alloy 45 MTZ • - - .0 .45 1.6 1.0

C13 Al loy 45 LC • - - ,0 .15 1.6 1.0

C14 G 4879 . - - .0 .45 1.5 1.5

C15 1.91 0.95

35 45 . - - , 1.0 . - - <1.0

35 45 • - - • 1.0 • - - <1.0

35 45 . - - . 1.0 . - - <1.0

28 48 . . . . . 5

28.2 19.5. - - . 0.04

< 1 . 0 . - - R E

< 1 . 0 • - - R E

< 1 . 0 • - - R E

0.05 . . . . . .

TABLE 2 RATE CONSTANTS FOR AIR-STEAM OXIDATION AND POST-COKE AIR-STEAM

OXIDATION FOR CENTRICAST ALLOYS HP 40 / HP 40 HSI AND EXPERIMENTAL ALLOY "A"

Alloy type

4852 micro

4852 micro, high Si

Al loy A

Air-steam oxidat ion at

900 °C kp*10 s

mg2cm-4min -1

7.0

2.0

0.4

Post-coke air- steam oxidat ion at

9O0 °C, kp* 105

mg2cm4min -1

8.6

0.6

0.1

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TABLE 3 NOMINAL CHEMICAL COMPOSITION OF WROUGHT ALLOYS IN LABORATORY METAL

DUSTING EXPERIMENTS (ACC. TO GRABKE, KLOWER, ET AL.)

A l l o y

No.

W 1

C 4

W 2

W 3

W 4

W 5

W 6

UNS No. Alloy Surfaces Design.

N 0 8 0 8 1 0 8 0 0 H Ground

H P 4 0

N 0 6 6 0 0 6 0 0 H Ground

N 0 6 6 0 1 6 0 1 Black

N 0 6 6 0 1 6 0 1 Polished

N 0 6 6 0 1 6 0 1 Ground

N 0 6 6 0 2 602CA Black

Nominal Chemical Composition (mass%) C Si Fe Cr Ni Ti AI others

0 . 1 0 0 . 5 4 6 2 0 3 2 0 . 4 0 . 4 A I + T i < 0 , 7

0 . 4 0 1 . 5 38 25 35 . . . . . . .

0 . 1 0 . - - . 9 1 6 7 2 . . . . . . .

0 . 1 0 • - - . 1 4 2 3 6 0 • - - . 1 . 4 . - - •

0 . 1 0 . - - . 1 4 2 3 6 0 . m . 1 . 4 • - - .

0 . 1 0 • m 1 4 2 3 6 0 • ~ . 1 . 4 . - - .

0 . 2 0 . - - . 9.5 25 6 0 • ~ . 2 . 3 0 . 1 Y

TABLE 4 DESIGN DATA FOR A PROCESS GAS HEATER IN A DRI PLANT

Gas Preheater operating conditions

Design rate Normal opration Mechanical rate Inlet conditions

Flow rate, Nm3/h 246.145 258.450 277.250 est. Temperature °C 280 280 280 Pressure, bar(g) 5.21 5.21 5.21

Outlet conditions

Temperature, °C 950 950 970 Pressure, bar(g) 4.28 4.28 est. 4.28 est. Gas composition, Vol %

Hydrogen 70 70 70 Carbon Monoxide 16 16 16

Carbon Dioxide 3 3 3

Methane 8 8 8

Nitrogen 1.3 1.3 1.3

Water Vapour 2 2 2

Hydrogen Sulfide ppm 20 - 30 20 - 30 20 - 30

CO / CO 2 rating 5.33

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F I G U R E _

a z 0

o

7"

uJ

o . J ne LLI "r" I.-- 0

"-SJ.. , o~.. " % "~'~Z':.

. ,

." ~.~.,~ "- u+,s/,/.

SULFUR

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5

I

I

LOW MELTING CAR BUR_IZA-EION ~ COKING ~. y N

co. oUNO ME. LOO$.,.G OZ

! I I

, eta" % "~ i

. .J '~ ~ , ~ ," 0 " ~ < 0 / -,% ,, ~,°, '4.~ , ~ , ,

O~,~.~X-"x" o o~' ~ " z .,Ooo "~O0~X ~

MOLTEN SALTS

O /

> - - - ~

~ t

The Central Role of Oxidation in High Temperature Corrosion Processes (acc. to G. Lai ~)

Inter-Lattice located Substituted Foreign Atoms Coherend High melting Foreign Atoms I Step Precipitations Foreign Phase

\ t . . . . Incoherend Lattice oriented / . . . . . . \ ~ | Umpjacement / b11mng t-aces . . e ~ \ ~ l l ~ l Precipitaions / _ J / 1

Void

~ E ! e m e n t a r y cell ~ i i ~ l ~ / ' ~ j ~ ~ . ~ i l ~ ~ b ~ / / ~

2.86 A -~ 0.28 nm Screw-type Scale shaped !~" Grain diam. ~ ' Dislocation Precipitations p.e. 50t~rn

on Grain = 50,000 nm Boundaries

FIGURE 2 - Microscopic Section of a Multi-Grain Structure of Face Centered Cubic Alloys (fcc = Austenitic Structure); Crystallographic Defects in Real Alloy Lattices

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Spinel (Mn, Fe) Cr204 + M7C3 and Cr3C2

'3

~ersed IC6

Or

i 203 + Cr3C 2

Cr depleted

M23 c 6

dispersed MTC 3

Initial Oxidation Oxidation in a carburizing Environment

FIGURE 3- Oxidation and Carburization on a Microscopic Scale; Potential Reactions Between Gas and Metal Surface (Left: Layer formation after 1 st start-up / initial oxidation right: Oxidation in a carburizing environment)

TUBE A TUBE B

oxidesM7C3 ( 0 3 ) ~ [ (2.0)

M703 ÷ M2306

M7C 3 ÷ M230 6 + M2(C,N )

M230 6 + M2(C,N )

oxides (0.2)

......~H (0.4)

. . . . . . ~ l (o.8)

IV (4.5)

! (2.5)

I1 (2.8)

m (o.6)

oxides (0.3)

M 7 03

M7C3 + M2 (C,N)

MTC 3 + M23C 6 + M2(C,N )

IV (4.5) ~M2306 + M2(C,N )

oxides (0.4)

FIGURE 4 - Primary and Secondary Carbide Precipitation of two Alloy Modifications of HK 40 (acc. to Hall, et al. 6)

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0.30

0.25

E o 0.20 o~

E

o~ 0.15 t -

t -

O

"" 0.10 t - o~ . m

0.05

0

FIGURE 5 -

G 4842 micro

10 20

Sqare root of time ~/min

2 micro high Si

erimental Alloy A

30 40

Air-Steam Oxidation Kinetics at 900 °C for Alloys HP 40 M, HP 40 hSi and Experimental Alloy "A" (acc. to Kleemann)

FIGURE 6 -

700

E -

.c 500 -

"10 c- O

= 300 N

E L

o 100 -i

,,,0oc/ /

I I I I I I

1 2 3 4 5 6 7 S q u a r e root of t ime (,g-if)

Carburization Kinetics as a Function of Time for Alloy HP 40 M (C6) (acc. to Kleemann)

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e ' -"

e - "

o

.

.

0

/

J

.K40J G 4852

/

~ 2 ~ ~ ~ ~/

mi,:ro7 G4 ~

950 1000 1050 1100 1150 1200 °C

F I G U R E 7 - Comparison of Relative Carburization Resistance as a Function of Temperature; Pack Cementation Tests for 260 Hours

250

200

%

%

%

150 % %

(J co %

.> 100

• - %

2 50 'l

• p = 1 bar

• p = 10 bar

600 700 800 900 1000

Temperature °C

F I G U R E 8 - Carbon Activity Versus Temperature and Pressure in a Laboratory Investigation (acc. to Perez)

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FIGURE 9-

log pCO(g) 25 L ,

20

15

I ( I

Cr(CO) 6

10

5

0

-5

-10

-15

-20 -50 -40 -30 -20 -10 0

I I I

C r O 3

r [ J

1 0 2 0 ~30 4 0

log PO2(g)

Phase Diagrams at 650 °C in Contact with a Defined DRI Gas Mixture; Allocation of Thermodynamic Balance is Marked by a Star (*) Stability Regions for Chromium Oxide / Chromium Carbides

log pCO(g) 0 I I

-2

-4

-6

- 8

-10

-12

-14

-16 I .

SiO 2

,,, I

-45 -40 -35 -30 -25 -20

log PO2(g )

FIGURE 10 - Phase Diagrams at 650 °C in Contact with a Defined DRI Gas Mixture; Allocation of Thermodynamic Balance is Marked by a Star (*) Stability Regions for Elemental Silicon, Silicon Oxide and Silicon Carbide

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log pCO(g) 8 i ~ \ 1 " ' 1 i i i i J

4

2o Fe3C, / ~ FeCO3

-2 Fe304 ~101.5(W ) -4 Fe~, Fe°'9470 I ' ~ - 6 I I I I I I r , r I"-.

- 3 5 - 2 5 - 1 5 - 5 5 1 5

log PO2(g ) FIGURE 11 - Phase Diagrams at 650 °C in Contact with a Defined DRI Gas Mixture;

Allocation of Thermodynamic Balance is Marked by a Star (*) Stability Regions for Metallic Iron, Iron Oxides and Iron Carbides (Cementite)

I

~ ~ i ~ . ~ - ~ , ~ ' ~ Cr203

Alloy II

Alloy

III Graphite / Cr20 3

Alloy

IV Graphite Cr203

Alloy V Fe3C

C r2%

Alloy ' ~

Metastable Carbide = (Fe,X) 3 C VI X= Ni, Cr

Alloy ~ F e , Cr, Ni +C

FIGURE 12 - Metal Dusting Process Proposed by the Investigators for High Performance Centricast Alloys

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0.4

0.3

E 0.2 o o )

E 0.1 O3 '- 0 ¢ -

o -0.1 ffl

:~ -0.2

-0.3

-0.4

Alloy MTZ Alloy LC

Temperature 750 °C, Surface Finish #50

FIGURE 13 - Laboratory Controlled Metal Dusting Tests as a Function of Temperature and Alloy Composition; Surface Finish: Grid # 50 Weight Gain for Alloy 45, Low (LC) and High (MTZ) Carbon Version at 750 °C

0.4

0.3

0.2 E o

o3 0.1 E o3 0 EE

" -0.1 O " "

¢/}

- 0 . 2

.0.3 i

-0.4

Alloy MTZ

i

Alloy LC


FIGURE 14 - Laboratory Controlled Metal Dusting Tests as a Function of Temperature and Alloy Composition; Surface Finish: Grid # 50 Mass Loss for Alloy 45, Low (LC) and High (MTZ) Carbon Version at 650 °C

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£ M

E tO

E a) o ) t -

t -

O t,/) t,t}

0.4

0.3

0.2

0.1-

0

-0.1

-0.2

-0.3

-0.4

I

!

Alloy MW Alloy MT


FIGURE 1 5 - Laboratory Controlled Metal Dusting Tests as a Function of Temperature and Alloy Composition; Surface Finish: Grid # 50 Significant Mass Gain for Alloy HP 40 (MW / MT) at 750 °C

E O

E

E: (.-

(o ¢/) ¢/) 0~

0.4

0.3

0,2

0.1

0

-0.1

-0.2

-0.3

-0.4

I i I

Alloy MW

, I I I

Alloy MT


FIGURE 16- Laboratory Controlled Metal Dusting Tests as a Function of Temperature and Alloy Composition; Surface Finish: Grid # 50 Mass Loss for Alloy HP 40 (MW / MT) at 650 °C

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0.25

t - £,,I E 0.2O O O) E

0.15 (If oC

O ) (I I --, 0.10 03

00

0.05

0

5 -....

""; ] 6,697 1,988

1.0E-0.31 1.1E-0.5

Ground HP40 Ground Black Polished Ground Black 800H 600H 601 601 601 602CA

FIGURE 1 7 - Metal Wastage Rates for Centricast and Wrought Nickel Base Alloys and Stainless Steels Alloy compositions acc. to Table 3 (acc. to Grabke, Klbwer, et al.)

Hydrogen Reformer

Iron Ore Pellets

opt.~O2 H2+CO+CO2+H20 L ~ i

. . . . . opt. H2 Gas Preheater ] O ~ ~ - ~ ~crubbe r

650 C- ~ / ^ H ~+ CO ,= { ~ L.r- "1 < - y ,. j y ,0oc "1 M~v~ngl ~ J

~" " "= , , r , , , I Reactor l \ ~ ~ I Danger of [ ,,J

2 0 ° ,=. ~ , Z ~ ', Metal Dusting \ / 8 C AE - ' , ' ~ I

/ ~ ' I ,, Outlet Coil \ / I " " %-" ET 45 Micro ~ Iron I VPellets

- - - ~ ~ . 1 " ~ - Top Gas Recycling Line Steam a~orT H H^O Gener z In let/( ,oil

Danger of Embrittlement ET 45 Micro (LC?)

FIGURE 18 - Flow Sheet of a Direct Reduction Plant for Iron Ores (DRI); Gas Pre-Heater: T Inlet < 650 °C, T Outlet = 950 °C

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FIGURE 19 - Failed Tube Section from a DRI Pre-Heater After six Months of Operation

FIGURE 20 - Microscopic Evaluation of Structures Resulting on a Tube Surface of Centricast Alloy HP 40; DRI Pre-Heater After six Months of Operation Glassy Surface Scale Morphology Close to the Rim of the Pit

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FIGURE 21 - Microscopic Evaluation of Structures Resulting on a Tube Surface of Centricast Alloy HP 40; DRI Pre-Heater After six Months of Operation Localized Accumulation of Interwoven Carbide Fibers in Selected Areas

Chain o Oxides

Oxides

FIGURE 22 - Microscopic Evaluation of Structures Resulting on a Tube Surface of Centricast Alloy HP 40; DRI Pre-Heater After six Months of Operation Inter-Carbidic Oxide Infiltration and Carbide Morphology in Surface Layers

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Ch~ Oxiq

Oxides

FIGURE 23 - Microscopic Evaluation of Structures Resulting on a Tube Surface of Centricast Alloy HP 40; DRI Pre-Heater After six Months of Operation X-Ray Map of Oxygen Distribution from FIGURE 15C

25 t . -

20 x"

I11

~ 15 0 f -

r "

• '-" 10 ( - (D

IZ 0

? 5 o

0

e~ <

0

J

i , i

300 1000 1500 2000 Distance from Surface, pm

FIGURE 24 - Chromium Content in a Centricast HP 40 Nb Matrix After Exposure to Metal Dusting Corrosion in a DRI Pre-Heater (acc. to Kleemann, Engel)

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F I G U R E 2 5 - DRI Plant / Gas Pre-Heater Hairpin Tube Section made of Ni45Cr35 (Alloy 45); in Frequent Operation for two Years

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Some spots of CrO + CrC

Heavy secondary carbides on the grain boundaries

SiO dispersed in Cr-depleted matrix

with pores and coke Inclusions

No visible protective layer on the inner tube surface

HP40 (C4)

Very small

Cr-depleted zone with voids where primary carbides had been located

Precipitations of M7C3 TiC and NbC on grain boundaries and fine dispersed in the matrix

No carbon diffusion beyond the depleted zone.

Continuous layer of Cr203 followed by SiO2

Alloy 45 (C11)

FIGURE 26 - Sections of Inner Tube Wall (100 x)

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Decarburized area

Black = Sil icon Oxides

Dark Grey = Chromium Oxides (mixed with Si-oxides)

ID Surface

X-ray Mapping of Iron in Cr-oxide Layer

X-ray Mapping of Nickel in Cr-oxide Layer

FIGURE 27 - Oxide Layer Formation Under Frequent Alternating Cycles of Strongly Carburizing and Oxidizing Cycles in an Ethylene Cracker Plant; t = 1.050 ° - 1.100 °C

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PRODUCTION PROCESS OXIDE FORMATION

O V o ° 0

0 0 0 0 0 0 0 0

Slag formation in liquid phases

Risk of nonmetallic inclusions (oxides)

MELTING in induction furnace

I

i

CENTRICASTI NG in metallic molds

with refractory liner

INSIDE MACHINING

Slag formation and strong oxidation

in the turbulent atmosphere

during pouring and

at the inner surface of the tube

Removal of heavy, porous

oxide layers and controlled

formation of thin and

protective "Spinel-Oxide-Layers"

Automatic Manual

i i i i i v i i i i l l i

Root protrusion is removed by grinding

ASSEMBLY OF TUBE SECTIONS BY WELDING

Partial alteration of thin (light silvery)

oxide layers into coloured "heat tints"

must be removed after welding

from inside and outside surfaces.

APPL ICAT ION OXIDE FORMATION

Conversion of the

thin glass-like oxide film

into thicker oxide layers

with defined

@ growth conditions

START OF RUN in high temperature gas-reactor under defined preoxidation conditions

~'~:'~!!t~ ;! J (C )

Carbon precipitation and partial oxide decomposition in the reducing atmospheres

Uniform alteration of the

protective oxide layer

(thinning)

Local defects in the oxide layers

resulting in catalytically active

spots of Ni / Fe under creep and

thermal shock conditions

Weight Creep conditions under high temperature and temperature gradients through the wall

Reheating of defect-sites

under oxygen depletion with

strong oxide forming elements

from the substrate alloy pool

Decoking or steam cleaning during plant shut downs

F IGURE 28 - The central Role of Oxidation in the Life Cycle of a Centr icast High Al loyed Tube

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,og Pcod Pco log PH20/PH2

0

-4

m

o - 8

-12

-16

-16 -12 -8 I I l I I

-16 -12 -8 I I I I I

Cr3C2 "o

Cr7C 3

/P 0r230 6

Cr

-4 0 4 8 12 I I I I I I I I I I

-4 0 4 8 12 I I I i I I I I l I

c (s)

C r 2 0 3

I I I I I I I I I I

-50 -40 -30 -20 -10 0

log P02 (ata)

0 4 N

-r-

O 3 o

4

-8

12

16 +10

~ 0

0

4 0

..(

" - 4

1 -12

-16

F I G U R E 2 9 - Stability Diagrams of Cr-C-O Systems at 620 °C (acc. to Lai 5)

2

I n I n

" " 1 I I I I

2 0 4 0 6 0 8 0 1 0 0

F I G U R E 3 0 - Carburization as a Function of Nickel Content in FCC Alloys

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100-

E

L) 5 0 -

,~ ._~ _ ___ , . H P M i c r o

10- Alloy 4 ~ '

601,602

35/25 60/23 72/16 45/35 60/25 Ni / Cr Ratio

Cast alloys Wrought alloys

FIGURE 31 - Alloy Ranking According to Typical Carbon Pick-up at 1.100 °C in Solid Carburization Cement / 32 Days

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01374 Correlation of Oxidation Carburization and Metal Dusting Controlling Corrosion by Corrosion...

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Transcript of 01374 Correlation of Oxidation Carburization and Metal Dusting Controlling Corrosion by Corrosion...